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The Primary Creep Behavior of Single Crystal, Nickel Base Superalloys PWA 1480 and PWA 1484

Permanent Link: http://ufdc.ufl.edu/UFE0021544/00001

Material Information

Title: The Primary Creep Behavior of Single Crystal, Nickel Base Superalloys PWA 1480 and PWA 1484
Physical Description: 1 online resource (297 p.)
Language: english
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2008

Subjects

Subjects / Keywords: creep, gamma, nickel, processing, rhenium, superalloy, tensile, xrd
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: Primary creep occurring at intermediate temperatures (650C to 850C) and loads greater than 500 MPa has been shown to result in severe creep strain, often exceeding 5-10%, during the first few hours of creep testing. This investigation examines how the addition of rhenium and changes in aging heat treatment affect the primary creep behavior of PWA 1480 and PWA 1484. To aid in the understanding of rhenium?s role in primary creep, 3wt% Re was added to PWA 1480 to create a second generation version of PWA 1480. The age heat treatments used for creep testing were either 704C/24 hr. or 871C/32hr. All three alloys exhibited the presence of secondary gamma prime confirmed by scanning electron microscopy and local electrode atom probe techniques. These aging heat treatments resulted in the reduction of the primary creep strain produced in PWA 1484 from 24% to 16% at 704C/862 MPa and produced a slight dependence of the tensile properties of PWA 1480 on aging heat treatment temperature. For all test temperatures, the high temperature age resulted in a significant decrease in primary creep behavior of PWA 1484 and a longer lifetime for all but the lowest test temperature. The primary creep behavior of PWA 1480 and PWA 1480+Re did not display any significant dependence on age heat treatment. The creep rupture life of PWA 1480 is greater than PWA 1484 at 704C, but significantly shorter at 760C and 815C. PWA 1480+Re, however, displayed the longest lifetime of all three alloys at both 704C and 815C (PWA 1480+Re was not tested at 760C). Qualitative TEM analysis revealed that PWA 1484 deformed by large dislocation 'ribbons' spanning large regions of material. PWA 1480, however, deformed primarily due to matrix dislocations and the creation of interfacial dislocation networks between the ? and ?? phases. PWA 1480+ contained stacking faults as well, though they acted on multiple slip systems generating work hardening and forcing the onset of secondary creep. X-ray diffraction and JMatPro calculations were also used to gain insight into the cause of the differences in behaviors.
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Thesis: Thesis (Ph.D.)--University of Florida, 2008.
Local: Adviser: Fuchs, Gerhard E.

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2008
System ID: UFE0021544:00001

Permanent Link: http://ufdc.ufl.edu/UFE0021544/00001

Material Information

Title: The Primary Creep Behavior of Single Crystal, Nickel Base Superalloys PWA 1480 and PWA 1484
Physical Description: 1 online resource (297 p.)
Language: english
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2008

Subjects

Subjects / Keywords: creep, gamma, nickel, processing, rhenium, superalloy, tensile, xrd
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: Primary creep occurring at intermediate temperatures (650C to 850C) and loads greater than 500 MPa has been shown to result in severe creep strain, often exceeding 5-10%, during the first few hours of creep testing. This investigation examines how the addition of rhenium and changes in aging heat treatment affect the primary creep behavior of PWA 1480 and PWA 1484. To aid in the understanding of rhenium?s role in primary creep, 3wt% Re was added to PWA 1480 to create a second generation version of PWA 1480. The age heat treatments used for creep testing were either 704C/24 hr. or 871C/32hr. All three alloys exhibited the presence of secondary gamma prime confirmed by scanning electron microscopy and local electrode atom probe techniques. These aging heat treatments resulted in the reduction of the primary creep strain produced in PWA 1484 from 24% to 16% at 704C/862 MPa and produced a slight dependence of the tensile properties of PWA 1480 on aging heat treatment temperature. For all test temperatures, the high temperature age resulted in a significant decrease in primary creep behavior of PWA 1484 and a longer lifetime for all but the lowest test temperature. The primary creep behavior of PWA 1480 and PWA 1480+Re did not display any significant dependence on age heat treatment. The creep rupture life of PWA 1480 is greater than PWA 1484 at 704C, but significantly shorter at 760C and 815C. PWA 1480+Re, however, displayed the longest lifetime of all three alloys at both 704C and 815C (PWA 1480+Re was not tested at 760C). Qualitative TEM analysis revealed that PWA 1484 deformed by large dislocation 'ribbons' spanning large regions of material. PWA 1480, however, deformed primarily due to matrix dislocations and the creation of interfacial dislocation networks between the ? and ?? phases. PWA 1480+ contained stacking faults as well, though they acted on multiple slip systems generating work hardening and forcing the onset of secondary creep. X-ray diffraction and JMatPro calculations were also used to gain insight into the cause of the differences in behaviors.
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Thesis: Thesis (Ph.D.)--University of Florida, 2008.
Local: Adviser: Fuchs, Gerhard E.

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2008
System ID: UFE0021544:00001


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THE PRIMARY CREEP BEHAVIOR OF SINGLE CRYSTAL, NICKEL BASE
SUPERALLOYS PWA 1480 AND PWA 1484




















By

BRANDON CHARLES WILSON


A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY

UNIVERSITY OF FLORIDA

2008




































O 2008 Brandon Charles Wilson



































In memory of Richard "Doc" Connell, Ph.D.









ACKNOWLEDGMENTS

This work would not have been successful without the help of many great and wonderful

people. First, I must thank my family and friends that have supported and loved me through this

process. These are the people that were there to give me advice when I needed it, to help me

focus when I needed to, to help me relax when I needed to, and to pray for me always. So much

of what brings a Ph.D. to completion happens outside of the lab and away from a computer and I

am incredibly thankful for the love and care directed towards my wife and me while we were in

graduate school.

My research hinged on the kindness of several strangers (at the time) during the early

stages of my work. Three engineers at Pratt & Whitney Aircraft Engines (East Hartford, CT)

deserve to be mentioned for their gift of a combined total of 47 single crystal test bars. Alan

Cetel and Dilip Shaw supplied the original 12 bars along with some journal articles of note to get

me started. Later, Samuel Krotzer supplied 1 1 bars of PWA 1480, 12 bars of PWA 1484, and

even agreed to make a special heat of PWA 1480 with 3wt.% rhenium added free of charge.

Without their help, this proj ect would not have gone very far. I also need to recognize Rick

Black and Bill Kumrow of Satec (a division of Instron, Norwood, MA) for their help in

diagnosing and repairing electrical and software failures on the M-3 creep system used for most

of the data presented here.

Additionally, the hard work of several members of the Maj or Analytical Instrumentation

Center (MAIC) at the University of Florida (Gainesville, FL) and the Advanced Materials

Processing and Analysis Center (AMPAC) at the University of Central Florida (Orlando, FL)

contributed to several aspects of this investigation. Specifically, Wayne Acree, Ben Pletcher,

Jung Hun Jang, Valentine Craciun, Gerald Bourne, and Kerry Siebein of MAIC and Kirk

Scammon and Helge Heinrich of AMPAC were instrumental in training and helping me with the









various analytical techniques utilized for this research. Mike Kaufman and Anantha Puthicode,

formerly of the University of North Texas (Denton, TX), deserve to be recognized for their help

with generating the Local Electrode Atom Probe (LEAP) data presented herein as well.

Within the University of Florida, I would like to acknowledge my advisor and mentor, Dr.

Gerhard Fuchs. Without his help and direction I would have been lost in the enormity of the task

at hand. Former graduate student Slade Stotlz also deserves to be mentioned for his here-to-fore

unacknowledged SEM work during my undergraduate research. In addition to Slade's help, I am

indebted to my fellow students in the High Temperature Alloys Laboratory (HTAL) for their

advice, encouragement, and humor while I have been a student. From the Particle Engineering

Research Center (PERC) at UF, I thank Nate Stevens, Ph.D. who is a great friend that

volunteered to proof read this dissertation in various stages of completion. I am also thankful for

the quick and responsive work of several support staff members between the academic, payroll,

finance, and secretary offices. Without their help nothing that is done in the department would

ever be completed and I am thankful for their help in spite of the difficulties I often presented

them. Finally, I need to recognize my wife, Krista Renner Wilson, Ph.D., for her constant

support and encouragement. Whether I was frustrated or nervous or unsure of what to do next,

she would gently, lovingly guide my hand and my thoughts in the right direction. Without her

help, I would not have completed this project.












TABLE OF CONTENTS


page

ACKNOWLEDGMENTS .............. ...............4.....


LIST OF TABLES ............_...... ...............9...


LIST OF FIGURES .............. ...............10....


AB S TRAC T ......_ ................. ............_........2


CHAPTER


1 INTRODUCTION ................. ...............25.......... ......


2 BACKGROUND .............. ...............29....


Overview ................... .......... ...............29.......
The Rhenium Effect .............. ...............30....

Secondary y' Precipitates............... ..............3
Lattice Mismatch ................. ...............34.................

Modern Superalloys............... .. ..............3
The Challenge of High Temperature Service ................ ...............37...............
Tensile Behavior............... ...............38

Creep Behavior............... ...............41
Summary ................. ...............43.................


3 EXPERIMENTAL PROCEDURES............... ...............4


M material s .............. ...............44....
Heat Treatment .............. .... ...............46

Heat Treatment Development............... ..............4
Differential Thermal Analysis............... ...............48
Furnaces............... ...............49
Characterization ................. ............... .. ...............50......

Preparing Samples for Metallography ...._.._.._ ..... .._._. ...._.._ ...........5
Preparing Samples for TEM ................. ...............52................
Preparing Samples for LEAP .............. ...............53....
Preparing Samples for X-Ray Diffraction............... ..............5
JMatPro Thermodynamic Prediction............... ...............5
M echanical Behavior ................. ...............57........ ......
Tensile Testing .............. ...............58....

Creep Testing............... ...............60

4 RESULTS: ALLOY MICRO STRUCTURES............... ............6


Phase Descriptions ................. ...............62..__..........













The y Phase ................. ...............62........... ....
The y' Phase............... ...............64.
Prim ary y' ............. ...............65.....
Secondary y' ............. ...............66.....

The y/y' Eutectic ................ ...............68........... ....
The Carbides............... ............ ... ..... .......6

The Topologically Close Packed (TCP) Phases ......___ ...... .. __ ............... .....70
Changes Following Primary Creep ................. ...............72......... .....


5 RESULTS: TENSILE BEHAVIOR ................. ...............91......... .....


6 RESULTS: CREEP BEHAVIOR ................. ...............113......... ....


Full Length Tests ..............._ ...............114........_......
PW A 1480 ................. ...............114......... ......
PWA 1480+ ............... ...............116................
PW A 1484 ................. ...............118......... ......

Interrupted Tests ................. .... ............ ...............120......
Transmission Electron Microscopy (TEM) ................. ...............121...............
PW A 1480 .............. ...............121....
PWA 1480+ ............... ...............122................
PW A 1484 .............. ...............123....


7 RESULTS: ADDITIONAL CHARACTERIZATION............... ...........14


Local Electrode Atom Probe (LEAP) ................. ...............140........... ...
PW A 1484 LT .................. ......... ............14

Reconstruction (solution HT) ................. ...............143................
Composition profie (solution HT) ................. ...............143...............
PW A 1484 HT ............... ... ........... ...............145......

Reconstruction (age HT) ................. ...............145...............

Composition profie (age HT) ................. ...............145...............
Reconstruction (solution HT) ................. ...............146................

Composition profie (solution HT) ................. ...............146...............
Secondary y' Concentrations ................ ...............147................
X -Ray Diffraction (X RD) ................. ................. 148........ ....


8 DI SCUS SSION ................. ................. 163........ ....


M i cro structure ................. ...............164................

y/y' M orphology .............. ............... 164...
Carbides ................ .................. ...............168......

Topologically Close Packed Phases ................. ...............168...............
Tensile Behavior............... ...............16

Secondary y' ............. ...............170....

y Channel Thickness ................. ...............172................
Lattice M i sf t ................. ................. 173........ ....












Stacking Fault Energy .............. ...............173....
Anti-Phase Boundary Energy ................. ...............174...............
Tensile Results............... ...............175

Creep Behavior ................. ...............176......... ......
Tertiary Creep............... ...............177.
Rafting ................. ...............178......... ......
Primary Creep............... ...............180.
Creep Results ................. ...............182................
Modeling Primary Creep .............. ...............186....
Lattice Mi sfit ................. ................. 189........ ....

Second ary y' ............. ............... 19 1...
Composition .............. ...............193....
Concluding Remarks .............. ...............195....


9 CONCLU SION................ ..............19


Conclusions............... ..............19
Future Directions .............. ...............200....


APPENDIX


A DIFFERENTIAL THERMAL ANALYSIS .............. ...............201....


B XRD PEAK DECONVOLUTION ................. ...............208......... .....


LIST OF REFERENCES ................. ...............292................


BIOGRAPHICAL SKETCH .............. ...............297....










LIST OF TABLES


Table page

2-1 Compositions in weight percent of several common Nickel-based superalloys................36

3-1 Laue orientation data and original heat treatments for PWA 1480 .............. ..................44

3-2 Laue orientation data and original heat treatments for PWA 1484 .............. ..................45

3-3 Laue orientation data and original heat treatments for PWA 1480+ ............... ...............45

3-4 Heat treatments used in this study .............. ...............47....

3-5 DTA results for all three alloys............... ...............49.

3-6 Tensile test matrix with sample identification ................. ...............59........... ..

3-7 Creep test matrix with sample identification .............. ...............61....

5-1 Tensile results at 7000C and 8150C for all three alloys............... ...............101

5-2 Elastic modulus calculations from creep loads ...._ ................. ............... .....10

5-3 Creep loads vs. yield strength for all three alloys ................. ....___ ................ .11

6-1 Primary creep and creep rates from interrupted creep tests ................. ........._.._.......125

6-2 Rupture lives and total creep elongation from full-length creep tests .............................127

7-1 Composition (wt%) of secondary y' precipitates ................. ..............................16

7-2 Net changes in secondary y' concentration with increasing precipitate size ................... 160

7-3 Lattice misfit (%) for the (002) plane following heat treatments at 10800C ...................161

B-1 PWA 1480 (002) peak lattice misfit calculations ................. ............... ......... ...211

B-2 PWA 1480+ (002) peak lattice misfit calculations ................. ................. ......... 21

B-3 PWA 1484 (002) peak lattice misfit calculations .............. ...............213....

B-4 PWA 1480 (004) peak lattice misfit calculations .............. ...............214....

B-5 PWA 1480+ (004) peak lattice misfit calculations ................. ................ ......... .21

B-6 PWA 1484 (004) peak lattice misfit calculations .............. ...............216....










LIST OF FIGURES


Figure page

2-1 Idealized creep curve with the three creep stages labeled .............. .....................3

3-1 Flow-chart of heat treatments for LEAP samples ................. .............. ...............54

3-2 SEM micrograph of the tip of a LEAP specimen ................. ...............55.............

3-3 Creep and tensile specimen geometry and dimensions............... ...............5

3-4 Three Type K thermocouples attached to a creep specimen............... ...............61

4-1 y/y' microstructure of PWA 1480 HT3 .............. ...............73....

4-2 y/y' microstructure of PWA 1480+ HT3 ......... ........_____ ......... ..........7

4-3 y/y' microstructure of PWA 1484 HT3 .............. ...............74....

4-4 Composition of the y phase of PWA 1480 as a function of temperature ................... ........74

4-5 Composition of the y phase of PWA 1480+ as a function of temperature. ................... .....75

4-6 Composition of the y phase of PWA 1484 as a function of temperature ................... ........75

4-7 Composition of the y' phase of PWA 1480 as a function of temperature ................... ......76

4-8 Composition of the y' phase of PWA 1480+ as a function of temperature ................... ....76

4-9 Composition of the y' phase of PWA 1484 as a function of temperature ................... ......77

4-10 The y' phase in PWA 1480 near a retained eutectic .............. ...............77....

4-11 Irregular y' phase in PWA 1480 near a partially solutioned eutectic region .....................78

4-12 y/y' eutectic region in PWA 1480+ during the early stages of solutioning .......................78

4-13 The y' structure of as-cast PWA 1484 .............. ...............79....

4-14 The y' volume fraction vs. temperature for all three alloys ................ ............ .........79

4-15 Secondary y' in the y matrix of PWA 1480 following an interrupted creep test ...............80

4-16 Secondary y' in the y matrix of PWA 1480 following an interrupted creep test ...............80

4-17 Secondary y' in the matrix of PWA 1480+ following an interrupted creep test. ...............81











4-18 Three dimensional LEAP compositional map (18wt.% Al iso-surface).. ................... .......82

4-19 y/y' eutectics in as-cast PWA 1480 ................. ...............83..............

4-20 y/y' eutectics in the vicinity of primary carbides in PWA 1480+ ............... ................. 83

4-21 Close-up of a eutectic in as-cast PWA 1480+ ............ ...............84.....

4-22 Close-up of a retained eutectic in PWA 1480+ ............ ...............84.....

4-23 Carbide phase in PWA 1480+ (HTIA)............... ...............85.

4-24 Carbide phase in PWA 1480+ (as-cast, longitudinal section) .............. .....................8

4-25 Local carbide network in PWA 1480............... ...............86..

4-26 Carbide phase in PWA 1484............... ...............86..

4-27 A possible carbide that dissolved during solution heat treatment ................. ................. 87

4-28 TCP phase formation in PWA 1480+ during interrupted creep testing.............................87

4-29 TCP phase formation in PWA 1480+ during interrupted creep testing.............................88

4-30 y phase elongation in the [110] direction in PWA 1484 ................ ................. ...._.88

4-31 y phase elongation in the [110] direction in PWA 1484 ................ ............... ...._...89

4-32 y' phase shear along (111) planes in PWA 1484 .............. ...............89....

4-33 y' phase shear along (111) planes in PWA 1484 .............. ...............44....

5-1 Tensile results for PWA 1480............... ...............100.

5-2 Tensile results for PWA 1480+ ............ ...............100.....

5-3 Tensile results for PWA 1484............... ...............101.


5-4 Comparison of tensile results for all three alloys with the LT age (7000C) ................... .102

5-5 Comparison of tensile results for all three alloys with the LT age (8150C) ................... .102

5-6 Comparison of tensile results for all three alloys with the HT age (7000C) ........... ......103

5-7 Comparison of tensile results for all three alloys with the HT age (8150C) ........... ......103

5-8 Plastic deformation behavior of PWA 1480 LT at 7000C ................ ............ .........104

5-9 Plastic deformation behavior of PWA 1480 HT at 7000C ....._____ ... ... .....__..........105











5-10 Plastic deformation behavior of PWA 1484 LT at 7000C ........_..._... ......._._. .........105

5-11 Plastic deformation behavior of PWA 1484 HT at 7000C ....._____ ...... ....__ ..........106

5-12 Plastic deformation behavior of PWA 1480+ LT at 7000C ........._._. ...... .._.._..........106

5-13 Plastic deformation behavior of PWA 1480+ HT at 7000C ............___ .........__ ......107

5-14 Tensile behavior of PWA 1480 LT at both temperatures ........._.. ........ .._..............107

5-15 Tensile behavior of PWA 1480 HT at both temperatures. ....._____ ...... ....__..........108

5-16 Tensile behavior of PWA 1480+ LT at both temperatures ........._.. ....... ..._.. .........108

5-17 Tensile behavior of PWA 1480+ HT at both temperatures ...........__... .......__.........109

5-18 Tensile behavior of PWA 1484 LT at both temperatures ........._.. ........ .._..............109

5-19 Tensile behavior of PWA 1484 HT at both temperatures. ....._____ ...... ....__..........1 10

5-20 Yield strength as a function of temperature for all three alloys ...........__...................110

5-21 Ultimate Tensile Strength as a function of temperature for all three alloys...................11 1

5-22 True Failure Stress as a function of temperature for all three alloys.........._.._.. ..............1 12

6-1 Creep at 7040C/862 MPa of all three alloys .............. .....................125

6-2 Creep at 7600C/690 MPa of all three alloys .............. .....................126

6-3 Creep at 8150C/621 MPa of all three alloys ................. ...............126...........

6-4 Primary creep of PWA 1480............... ...............128.

6-5 Primary creep of PWA 1480+ ............ ...............128.....

6-6 Primary creep of PWA 1484............... ...............129.

6-7 Creep at 7040C/862 MPa of PWA 1484 ................. ...............129.............

6-8 Primary creep comparison at 7040C/862 MPa for all three alloys .............. .................130

6-9 Primary creep comparison at 7040C/862 MPa for all three alloys .............. .................130

6-10 Primary creep comparison at 8150C/621 MPa for all three alloys ................. ...............131

6-11 Bright Hield(a)/Dark Hield(b) pair of deformation in PWA 1480 .............. ....................132

6-12 Bright Hield(a)/Dark Hield(b) pair of deformation in PWA 1480 .............. ....................133











6-13 Bright Hield TEM image of dislocation networks in PWA 1480 ................. .................1 34

6-14 Stacking fault and dislocation shear of PWA 1480 ................ ............... ........ ...134

6-15 A stacking fault in PWA 1480 ................. ...............135........... ..

6-16 Secondary y' precipitates (marked by arrows) in PWA 1480............... ...................3

6-17 Stacking fault interactions following primary creep in PWA 1480+ ............ ................136

6-18 Stacking fault interactions and a dislocation network in PWA 1480+ ............ ...... .........136

6-19 Short range stacking fault shear of PWA 1480+ ................ ................ ......... .. 137

6-20 Bright Hield TEM image of PWA 1484 LT (7040C) ................. ................. ........ 137

6-21 Stacking fault shear of y' precipitates in PWA 1484 ................ ...........................13 8

6-22 Interfacial dislocation networks present in PWA 1484 .............. ....................13

7-1 Schematic illustrating the basic function of the LEAP system ................. ................. 15 1

7-2 LEAP specimens before, (a), and after, (b), the final polishing step ............... .... ...........152

7-3 Iso surfaces created with the LEAP system (PWA 1484 LT with solution HT).............153

7-4 Magnified (SEM) view of the tip analyzed in Figure 7-3 ......____ ... ... ....__ ............154

7-5 Composition profie from the specimen in Figure 7-3 ......____ ... ......_ ...............154

7-6 The distribution of all recorded ions for PWA 1484 HT ....._____ ... ... ..._ ............155

7-7 Iso surfaces created with the LEAP system (PWA 1484 HT no solution HT) ..............155

7-8 Composition profie from the specimen in Figure 7-7................ ...............156.

7-9 Iso surface (18% Aluminum) created with the LEAP system ................ ................ ...157

7-10 LEAP results with only Al, Ta, Cr, and Mo ions represented ................ ........._.._.. ...158

7-11 Composition profie from the specimen in Figures 7-9 and 7-10............... ..................15

7-12 Illustration of the data selected for the profie shown in Figure 7-1 1................... ...........159

7-13 An example of the deconvolution process ................ ...............161........... ..

7-14 Lattice misfit vs. heat treatment from from both the (002) and (004) peaks ................... 162

A-1 DTA trace for PWA 1480 in the HT1 condition............... ...............20










DTA trace for PWA 1484 in the HT1 condition............... ...............20

DTA trace for PWA 1480+ in the HTO condition .............. ...............204....

DTA trace for PWA 1480 in the HT2 condition............... ...............20

DTA trace for PWA 1484 in the HT2 condition............... ...............20

DTA trace for PWA 1480+ in the HT2 condition .............. ...............207....


A-2

A-3

A-4

A-5

A-6


B-1 XRD deconvolution of PWA 1480 (4hr. 10800C, skewed, 6.28% error)..


....................217


B-2

B-3

B-4

B-5

B-6

B-7

B-8

B-9

B-10

B-11

B-12

B-13

B-14

B-15

B-16

B-17

B-18

B-19


XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA


1480 (4hr. 10800C, skewed, 14.81% error) ................... ...218

1480 (4hr. 10800C, unskewed, 6.63% error) ................... .219

1480 (4hr. 10800C, unskewed, 12.58% error) ................. .220

1480 (10hr. 10800C, unskewed, 7.94% error) ................. .221

1480 (10hr. 10800C, skewed, 5.63% error) ................... ...222

1480 (100hr. 10800C, unskewed, 6.94% error) ................223

1480 (100hr. 10800C, skewed, 6.82% error) ................... .224

1480 (100hr. 10800C, unskewed, 5.20% error) ................225

1480 (1000hr. 10800C, unskewed, 5.93% error) .............226

1480 (1000hr. 10800C, skewed, 4.61% error) ........._.......227

1480 (1000hr. 10800C, skewed, 5.45% error) ........._.......228

1480 (4hr. 10800C, unskewed, 5.58% error) .........._.......229

1480 (4hr. 10800C, unskewed, 7.77% error) .........._.......230

1480 (4hr. 10800C, skewed, 4.99% error)..............._._.....231

1480 (4hr. 10800C, skewed, 1.70% error).............._._......232

1480 (10hr. 10800C, unskewed, 5.79% error) ........._.......233

1480 (10hr. 10800C, skewed, 1.41% error).............._._....234

1480 (1000hr. 10800C, unskewed, 3.17% error) ..............235










B-20

B-21

B-22

B-23

B-24

B-25

B-26

B-27

B-28

B-29

B-30

B-31

B-32

B-33

B-34

B-35

B-36

B-37

B-38

B-39

B-40

B-41

B-42

B-43


XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA


1480 (1000hr. 10800C, unskewed, 1.77% error) ..............236

1480 (1000hr. 10800C, unskewed, 1.75% error)..........._...237

1480 (1000hr. 10800C, skewed, 3.14% error) ................. .238

1480 (1000hr. 10800C, skewed, 1.71% error) ................. .239

1480 (1000hr. 10800C, skewed, 1.71% error) ................. .240

1484 (4hr. 10800C, unskewed, 13.3 6% error) ........._.......241

1484 (1000hr. 10800C, unskewed, 6.30% error) .............242

1484 (10hr. 10800C, unskewed, 4.61% error) ........._.......243

1484 (1000hr. 10800C, unskewed, 5.85% error) .............244

1484 (10hr. 10800C, unskewed, 4.45% error) ........._.......245

1484 (4hr. 10800C, skewed, 11.71% error).............._._....246

1484 (1000hr. 10800C, skewed, 3.71% error) ........._.......247

1484 (1000hr. 10800C, skewed, 3.20% error) ........._.......248

1484 (10hr. 10800C, skewed, 4.07% error).............._._....249

1484 (10hr. 10800C, skewed, 4.03% error).............._._....250

1484 (1000hr. 10800C, unskewed, 4.15% error) .............251

1484 (1000hr. 10800C, unskewed, 3.24% error) ..............252

1484 (1000hr. 10800C, unskewed, 4.3 5% error)............. .253

1484 (1000hr. 10800C, skewed, 3.09% error) ................. .254

1484 (1000hr. 10800C, skewed, 2.07% error) ................. .255

1480+ (4hr. 10800C, unskewed, 5.07% error) .................256

1480+ (4hr. 10800C, unskewed, 10.42% error) ...............257

1480+ (4hr. 10800C, unskewed, 10.33% error) ...............258

1480+ (10hr. 10800C, unskewed, 9.55% error) ...............259










B-44

B-45

B-46

B-47

B-48

B-49

B-50

B-51

B-52

B-53

B-54

B-55

B-56

B-57

B-58

B-59

B-60

B-61

B-62

B-63

B-64

B-65

B-66

B-67


XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA


1480+ (10hr. 10800C, unskewed, 5.73% error) ...............260

1480+ (10hr. 10800C, unskewed, 4.87% error) ...............261

1480+ (1000hr. 10800C, unskewed, 4.41% error) ...........262

1480+ (1000hr. 10800C, unskewed, 4.37% error) ...........263

1480+ (4hr. 10800C, skewed, 5.05% error) .....................264

1480+ (4hr. 10800C, skewed, 6.91% error) .....................265

1480+ (4hr. 10800C, skewed, 6.85% error) .....................266

1480+ (10hr. 10800C, skewed, 9.54% error) ...................267

1480+ (100hr. 10800C, skewed, 5.42% error) .................268

1480+ (100hr. 10800C, skewed, 4.13% error) .................269

1480+ (1000hr. 10800C, skewed, 4.30% error) ...............270

1480+ (1000hr. 10800C, skewed, 4.28% error) ...............271

1480+ (4hr. 10800C, unskewed, 2.38% error) .................272

1480+ (10hr. 10800C, unskewed, 9.32% error) ...............273

1480+ (10hr. 10800C, unskewed, 7.96% error) ...............274

1480+ (10hr. 10800C, unskewed, 3.84% error) ...............275

1480+ (10hr. 10800C, unskewed, 3.89% error) ...............276

1480+ (10hr. 10800C, unskewed, 3.37% error) ...............277

1480+ (10hr. 10800C, unskewed, 3.38% error) ...............278

1480+ (1000hr. 10800C, unskewed, 2.83% error) ...........279

1480+ (1000hr. 10800C, unskewed, 1.49% error) ...........280

1480+ (1000hr. 10800C, unskewed, 1.49% error) ...........281

1480+ (4hr. 10800C, skewed, 2.36% error) .....................282

1480+ (10hr. 10800C, skewed, 3.77% error) ...................283










B-68

B-69

B-70

B-71

B-72

B-73

B-74

B-75


XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA

XRD deconvolution of PWA


1480+ (10hr. 10800C, skewed, 7.19% error) ...................284

1480+ (10hr. 10800C, skewed, 8.42% error) ...................285

1480+ (10hr. 10800C, skewed, 7.04% error) ...................286

1480+ (10hr. 10800C, skewed, 3.28% error) ...................287

1480+ (10hr. 10800C, skewed, 3.31% error) ...................288

1480+ (1000hr. 10800C, skewed, 2.79% error) ...............289

1480+ (1000hr. 10800C, skewed, 1.18% error) ...............290

1480+ (1000hr. 10800C, skewed, 1.18% error) ...............291









LIST OF ABBREVIATIONS

Air cooled (following a heat treatment)

Advanced Insertion of Materials (a DARPA program)

Advanced Materials Processing and Analysis Center (a unit of the
University of Central Florida)

Anti phase boundary (can be formed in ordered phases)

American Society for Metals

d spacing (the spacing between planes of atoms)

Defense Advanced Research Proj ects Agency (a division of the United
States Department of Defense)

Full Width at Half Maximum (a parameter that describes the shape of an
intensity peak used for XRD deconvolution)

Gas furnace quench (used for solution heat treatment, Helium gas was
inj ected into a vacuum furnace for rapid cooling)

Generic term used to indicate the miller indices of a plane

Used to designate the high temperature age (8710C/32 hr./AC), also a
generic abbreviation for "heat treatment"

Solution heat treatment designations. All heat treatments are defined in
Table 3-4

Industrial gas turbine (usually used in the power generation industry)

Used to designate the low temperature age (7040C/24 hr./AC)

Linear Variable Differential Transducer

Maj or Analytical Instrumentation Center (a unit of the University of
Florida)

Designates the commercial heat treatment for the alloy in question (Table
3-4)

Type of X-ray diffractometer

The angle that a diffractometer measures during a scan

Denotes a single crystal superalloy developed by Canon Muskegon


AC:

AIM:

AMPAC :


APB :

ASM:



DARPA :


FWHM:


GFQ:


hkl :

HT:


HT1, HT2, etc.:


IGT:

LT:

LVDT:

MAIC:


STD HT:


6/26:

26:

CMSX:










First generation single crystal superalloy developed by Canon Muskegon,
Table 2-1

Second generation single crystal superalloy developed by Canon
Muskegon, Table 2-1

Denotes an alloy developed by Pratt & Whitney

First generation single crystal, nickel base superalloy made by Pratt &
Whitney. The composition is shown in Table 2-1

An experimental single crystal, nickel base superalloy created by adding
3wt.% rhenium to PWA 1480. The composition is shown in Table 2-1

Same as PWA 1480+Re

Second generation single crystal, nickel base superalloy made by Pratt &
Whitney. The composition is shown in Table 2-1

Differential Thermal Analysis

Energy dispersive spectroscopy

Focused ion beam

Local electrode atom probe

Scanning electron microscope (or microscopy)

Transmission electron microscope (or microscopy)

X-ray diffraction

Electrical discharge machining

Hydrochloric acid

Nitric acid

Molybdenum oxide

Aluminum

Carbon

Cobalt

Chromium


CMSX-2 :


CMSX-4 :


PWA:

PWA 1480:


PWA 1480+Re:


PWA 1480+:

PWA 1484:


DTA:

EDS:

FIB:

LEAP:

SEM:

TEM:

XRD :

EDM:

HCl:

HNO3 :

MoO3:

Al:

C:

Co:

Cr:









Hf: Hafnium

Mo: Molybdenum

Ni: Nickel

Nb: Niobium (also called Columbium, Cb)

Re: Rhenium

Ta: Tantalum

Ti: Titanium

W: Tungsten

Y: Gamma phase (fcc Nickel solid solution, matrix)

Y': Gamma prime phase (L12 Ordered Ni3Al, precipitates)

y/y':Used for discussion of both phases as a system (eg. the y/y' interface)

fcc: Face centered cubic

L12: An ordered fcc-like structure

MC: Type of carbide (M represents the metal atom, C represents the carbon
atom)

M6C: Type of carbide (M represents the metal atom, C represents the carbon
atom)

M23 6: Type of carbide (M represents the metal atom, C represents the carbon
atom)

TCP: Topologically Close Packed (deleterious phases that precipitate in some
alloys: CL, o, P, and Laves phases)

a: Lattice parameter

aY: Lattice parameter of the y phase

a : Lattice parameter of the y' phase

a angle: The error (in degrees) of a bar from the [001] orientation

at.%: Atomic percent

8: Lattice mismatch









Er: Failure strain

FS: Failure strength

kal: The wavelength of the X-ray radiation caused by kml electrons

ku2: The wavelength of the X-ray radiation caused by ku2 electrons

of: Failure strength

Oy: Yield strength

RIA: Reduction in area

T: Temperature

Tm: Temperature of melting (absolute temperature units only)

UTS: Ultimate tensile strength

wt.%: Weight percent

YS: Yield strength

A+: Angstrom (10-10 m)

cm: Centimeter (10-3 m)

GPa: Gigapascal (109 kg-m ~s-2)

hr: Hour

in: Inch

ksi: 1000 pounds per square inch (1000 psi)

L: Liter

lb.: Pound

m: Meter

Cpm: Micrometer (10-6 m)

min: Minute

mL: Milliliter (10-3 L)

MPa: Megapascal (106 kg-m ~s-2)










Nanometer (10-9 m)

Second

Volt (m2 kg s-3 A- )


nm:









Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy

THE PRIMARY CREEP BEHAVIOR OF SINGLE CRYSTAL, NICKEL BASE
SUPERALLOYS PWA 1480 AND PWA 1484

By

Brandon Charles Wilson

May 2008

Chair: Gerhard Fuchs
Major: Materials Science and Engineering

Primary creep occurring at intermediate temperatures (6500C to 8500C) and loads greater

than 500 MPa has been shown to result in severe creep strain, often exceeding 5-10%, during the

first few hours of creep testing. This investigation examines how the addition of rhenium and

changes in aging heat treatment affect the primary creep behavior of PWA 1480 and PWA 1484.

To aid in the understanding of rhenium' s role in primary creep, 3wt% Re was added to PWA

1480 to create a second generation version of PWA 1480. The age heat treatments used for

creep testing were either 7040C/24 hr. or 8710C/32hr. All three alloys exhibited the presence of

secondary y' confirmed by scanning electron microscopy and local electrode atom probe

techniques. These aging heat treatments resulted in the reduction of the primary creep strain

produced in PWA 1484 from 24% to 16% at 7040C/862 MPa and produced a slight dependence

of the tensile properties of PWA 1480 on aging heat treatment temperature.

For all test temperatures, the high temperature age resulted in a significant decrease in

primary creep behavior of PWA 1484 and a longer lifetime for all but the lowest test

temperature. The primary creep behavior of PWA 1480 and PWA 1480+Re did not display any

significant dependence on age heat treatment. The creep rupture life of PWA 1480 is greater










than PWA 1484 at 7040C, but significantly shorter at 7600C and 8150C. PWA 1480+Re,

however, displayed the longest lifetime of all three alloys at both 7040C and 8150C (PWA

1480+Re was not tested at 7600C). Qualitative TEM analysis revealed that PWA 1484 deformed

by large dislocation "ribbons" spanning large regions of material. PWA 1480, however,

deformed primarily due to matrix dislocations and the creation of interfacial dislocation networks

between the y and y' phases. PWA 1480+ contained stacking faults as well, though they acted on

multiple slip systems generating work hardening and forcing the onset of secondary creep. X-

ray diffraction and JMatPro calculations were also used to gain insight into the cause of the

differences in behaviors.









CHAPTER 1
INTTRODUCTION

Conventional wisdom regarding the first and second generation superalloys states that

they are well known systems with little left to learn about their behavior. Blanket statements are

often used to generalize about these older alloy systems while current research and development

continues to refine third, fourth, and fifth generation alloys. Recently, however, renewed interest

has been given to this class of alloys due to a curious phenomenon during intermediate

temperature creep. Within the temperature range of about 650oC to 850oC and under high stress,

some second generation superalloys, bearing 1 atomic percent or about 3 weight percent

rhenium, experience excessive primary creep. In these cases primary creep can be as high as 28-

30% in as little as 12 hours of creep testing.

These stress and temperature conditions are important as they are present near the root, or

attachment point, of turbine blades as well as within the internal support structure directing

internal air cooling paths between the airfoil surfaces. If these regions deform at an excessive

rate, internal stresses can result and failure can occur. It is also important to note that a turbine

blade may fail through multiple methods. The most obvious is failure due to fracture. Another

method is by exceeding dimensional tolerances. Gas turbine engines are designed with very

tight tolerances. With turbine blades, the allowed expansion due to creep and other processes is

often only a few percent. These tolerances, therefore, become threatened if some regions within

a turbine blade can creep several percent in the first few hours of operation.

There have been numerous attempts at identifying the causes and controlling factors of

primary creep in these alloys. Two suspected causes are rhenium content and the presence of

secondary y' precipitates in the y matrix channels. Secondary y' precipitates are produced

following the last aging heat treatment. These precipitates are usually about 10-20 nm in size










and populate the y matrix channels that are about 100-200 nm in diameter. It is widely

documented that dislocations in superalloys are found predominately in the y matrix. If this

same region is populated by very fine, densely spaced y' then dislocation slip in these channels

may be greatly impeded. It has been reported that the presence of these secondary particles

raises the difficulty of shear in the y phase to such a level that shearing of the primary y'

precipitates becomes the preferred method of dislocation motion. Once a dislocation pair enters

the y' phase there is little barrier to glide until the opposite y/y' interface is reached. These long

glide paths, then, are possible causes of large primary creep strains.

Another source of y matrix strengthening takes the form of solid solution strengthening

brought about by additions of Re. Rhenium has been used for solid solution strengthening of the

y phase beginning in the 1980's. How Re achieves its strengthening effect has long been a topic

for debate in the superalloy community. Recent research has focused on changes in mechanical

behavior and the elemental segregation of Re relative to the y/y' interface. The strengthening

effect brought by the addition of Re was so significant that the addition of Re alone was enough

to define the second and third generations of superalloys. The most striking difference in

behavior can be seen by comparing first and second generation alloys, or alloys with no Re to

alloys with 3wt.% (lat.%) Re. The second generation alloys demonstrated greater than 250C

improvements in creep and tensile strength capabilities.1-4 After over 20 years of research, the so

called "Rhenium Effect" is still not fully understood.

Adding impetus to the need to understand the Rhenium Effect, is the sudden climb of the

cost of rhenium. Additions of rhenium are found in most of the alloys used for critical

applications in aerospace gas turbines as well as some industrial gas turbines (IGT). Over the

last 18 months, the price of rhenium has climbed from $600/pound to $2500/pound. This drastic









increase in cost has many companies investigating alternative methods to strengthen their alloys.

One possible route includes replacing rhenium with another refractory element like tungsten, W.

Also, if investigators could determine the effect rhenium has on dislocation behavior, perhaps

new strategies for strengthening could be developed.

Interestingly, excessive primary creep and the Rhenium Effect may be related. The alloys

most commonly associated with excessive primary creep are rhenium bearing superalloys (i.e.

2nd, 3rd, 4th, and next generation alloys). Typical creep behavior for these alloys will display

large primary creep strains in the first few hours of testing under high loads and low

temperatures. Following primary creep is a sharp transition to secondary, or steady-state, creep.

During secondary creep in these alloys the creep rate is rather low and unchanging. First

generation alloys, however, tend to not display much primary creep and the secondary creep rate

is higher and slowly increases for the duration of the creep test. Therefore, while rupture times

for second generation alloys may be a full order of magnitude longer, time to 1% creep may be

significantly shorter than their first generation counterparts. This behavior relates directly to the

issue of dimensional tolerance in modern aerospace engines: while second generation alloys will

rupture at longer times, they may, in fact, fail dimensional controls significantly earlier than their

rupture lives might suggest. Alternatively, this behavior might require unique design

considerations during turbine run-in to account for excessive deformation.

The current investigation focuses on commonly used first and second generation alloys:

PWA 1480 and PWA 1484, respectively. A third, experimental alloy was also produced to

isolate the effect of rhenium during intermediate temperature creep by adding 3 weight percent

Re to PWA 1480. The approach used to test the three alloys begins with a solution and

homogenizing heat treatment for all three alloys to reduce the effects of segregation during









solidification. Two different aging heat treatments were then used with the intention of creating

specimens with and without secondary y' in the y matrix channels in order to observe changes in

primary creep effected by these precipitates. Full length creep testing and interrupted creep

testing was conducted to correlate microstructure with mechanical properties. Additionally, X-

ray diffraction studies, JMatPro thermodynamic predictions, and TEM techniques were used to

investigate the primary creep behavior of the three alloys. Finally, the problem of secondary y'

and the Rhenium Effect with regard to excessive primary creep is discussed.









CHAPTER 2
BACKGROUND

Overview

Modern superalloy technology has a long history of developments in processing, alloy

chemistry, and fundamental metallurgical knowledge. The advent of directionally solidified,

followed by single crystal, blades and vanes led to maj or improvements in the high temperature

tensile, creep, and fatigue properties of superalloys. Alloy refinement through advanced

remelting processes and alloy chemistry has led to consistency between master heats and

improvements in mechanical properties and environmental resistance.2, 5 These developments

occurred over several decades; however, the pace of research has occasionally exceeded the rate

of application of these new alloys. For instance, a typical development cycle from inception to

service for a new alloy can take 7-10 years. By the time that alloy is put into service, newer

alloys with improved properties are already in development. Adding to this, the United States

Department of Defense has instituted the Accelerated Insertion of Materials (AIM) program, a

Defense Advanced Research Proj ects Agency (DARPA) initiative. The goal of this program is

to decrease the time to active service of new alloys used for aerospace turbine engine

applications.6 Similarly, the United States Department of Energy instituted a similar program in

the early 1990's to improve performance of industrial gas turbines. The Advanced Turbine

Systems (ATS) program also created a push for new materials with higher temperature

capabilities.7 These government programs coupled with the already strong drive to increase

performance has led to rapid development cycles for new materials.

While the high pace of development in the superalloy industry is beneficial for engine

manufacturers, customers, and national defense alike, new alloys are often developed before a

firm understanding of previous performance gains could be achieved. A prime example is the









addition of rhenium to second and third generation superalloys. Since the first implementation

of Re to these alloys in the early to mid 1980's, the superalloys community has seen the

development of 2nd generation (3 wt% Re), 3rd generation (6 wt% Re), 4th generation (with

platinum group additions), and now early development of the next generation of superalloys. As

the momentum of research continues to advance alloy performance, the method by which Re

improves the strength of superalloys is still debated. Early studies examined microstructural

effects, changes in y' stacking fault energy, changes in y/y' lattice mismatch, and Re segregation

to the y phase during precipitation and coarsening of y' precipitates. "1 Entering the 1990's,

interest in exploring this effect was replaced by the need to boost engine performance. Thus the

third generation of superalloys was created through even greater rhenium additions, though the

exact strengthening mechanism was still not entirely clear.

The Rhenium Effect

Twenty years later, there has been an increase in the interest in the so called "Rhenium

Effect". Part of this rekindled interest is due simply to the lag in development between military

aircraft engines to commercial engines to industrial gas turbines, IGTs. Alloys that are new to

commercial aircraft applications were first used in military aircraft engines 10 years earlier. The

same gap exists between commercial engines and industrial gas turbines. These technological

divides are due partly to market pressure and partly to processing challenges. New, high

performance alloys are often very difficult to cast into a properly oriented single crystal. As the

engine increases in size, its turbine blades must increase in size as well. Turbine blades on

military aircraft may be 8 to 15 cm in length while commercial engines require 10 to 20 cm and

IGTs require blades significantly greater than 30 cm in length. The consequence of these size

differences is an increased rejection rate using conventional processing technologies. Thus,










improvements in casting processes were often necessary before an established alloy could be

applied or scaled up to new markets.

Now that second generation single crystal alloys have made their way into commercial

aircraft and IGT applications, interest is growing to understand their behavior primarily due to

cost pressures and increasing loads at intermediate temperatures. These new applications require

larger components and, consequently, larger volumes of Re bearing superalloys. The recent

rapid rise in the cost of Re has the superalloy community investigating alternative strengthening

additions. Additionally, some second generation alloys during intermediate temperature creep

(6500C to 8500C) at high loads (greater than 500 MPa) will exhibit severe creep anisotropy that

has been linked to stacking fault shear of the y' precipitates.12-15 Creep anisotropy has also been

linked to large primary creep strains during testing under these same conditions.

By convention, creep behavior is separated into three stages as shown in Figure 2-1. The

primary creep stage of interest to the current investigation occurs first. Primary creep is typically

marked by a relatively large initial creep rate due to the initially undeformed nature of the

microstructure. As deformation occurs, and hardening takes place, the creep rate is reduced.

Eventually, a balance between the active deformation processes (dislocation slip) and recovery

processes is reached and the creep rate stabilizes. This regime of creep is called secondary creep.

Near the end of the life of a specimen, the total accumulation of deformation within the

microstructure increases the rate of deformation such that the active recovery processes can no

longer maintain a steady-state balance. As a result, the creep rate begins to increase.

Approaching failure, the creep rate is continually increasing and this behavior marks tertiary

creep. The idealized behavior shown in Figure 2-1, however, is not always applicable. Often,

usually at high temperatures, alloys will exhibit brief primary creep that transitions immediately










into tertiary creep with no steady-state secondary creep stage. Another variation, which is

discussed in this investigation, is creep dominated by the primary creep stage. In specimens

exhibiting this behavior, the primary creep stage is extended to greater levels of strain before

secondary creep can occur. Additionally, with alloys exhibiting large primary creep stages,

incubation periods of near zero creep usually preceded the primary creep stage. It is thought that

these incubation periods mark the formation of stacking fault ribbons prior to y' shear.16, 17



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of lower temperatures similar to those seen in the attachment and internal support regions of

turbine blades, it may also have been deemed less important than high temperature creep

strength.

Secondary y' Precipitates

A microstructural feature that may impact the primary creep behavior of single crystal

nickel base superalloys is the presence of secondary y' precipitates. While it has been long

known that primary y' precipitate size has a maj or impact on mechanical properties,21, 22 TOSearch

has also shown the importance of controlling the secondary y', or cooling y', precipitateS20, 23

The final heat treatment step before a turbine blade is released for service is typically a y' aging

heat treatment. This final step significantly impacts the performance of the component. First,

the time and temperature for the age heat treatment controls the coarsening behavior of the

primary y' precipitates. Following solution heat treatment, these precipitates are mostly coherent

in nature.

Aging serves to increase the size of the precipitates to the optimum 0.3-0.5 Cpm size range.

Additionally, the coherency of the precipitates is reduced adding a misfit strengthening

component to the overall strength characteristic of the alloy. Besides time and temperature, the

cooling rate following aging has been found to create microstructural differences that can impact

mechanical properties. Rapid cooling from the aging heat treatment temperature has been shown

to produce a fine dispersion of secondary y' precipitates of the size range 10-50 nm in diameter.

These precipitates reside in the y matrix channels between primary y' precipitates. Because

dislocations tend to initiate in the y matrix, it is expected that interactions between dislocations

and these precipitates would be common. Furnace cooling, however, from the aging temperature

is capable of creating a y matrix that is devoid of secondary y' precipitates.20










While it is expected that a significant interaction exists between the secondary y' and the

active deformation mechanism, exactly how the secondary y' effects the deformation process is

not fully understood. These precipitates have even been linked to large primary creep strains as

they may prohibit deformation in the y matrix channels, forcing shear of primary y' precipitates

resulting in creep anisotropy. Alternatively, it is thought that these precipitates may stabilize the

stacking fault ribbons within the y matrix, leading to enhanced y' shear and large primary creep

strains.16' 17, 20, 24

Lattice Mismatch

Another microstructural feature that is thought to impact primary creep is lattice mismatch

(misfit). While it is known that lattice mismatch can alter behavior for high temperature creep,

specifically in the case of y' rafting,25' 26 lattice misfit also bears a significant impact on

dislocation motion in the y matrix and at the y/y' interfaces.27 Lattice mismatch between the

matrix and precipitates in superalloys results in coherency strains along these interfaces. If the

coherency strain is large enough, misfit dislocations will nucleate and/or congregate at these

interfaces to relieve the strain. Alloys with large lattice misfit values have shown a propensity to

form dense networks of interfacial dislocations during creep testing. Conversely, alloys with

reduced y/y' misfit, exhibiting creep anisotropy, are characterized by relatively few interfacial

dislocations and large regions of stacking faults in the precipitates. This difference is possibly

the result of compositional changes (which will cause lattice mismatch modification). In the

case of CMSX-2 and CMSX-4, for instance, the most notable changes were the additions of

rhenium and hafnium, an increase in cobalt, and a decrease in tungsten and chromium. Between

the alloys PWA 1480 and PWA 1484 similar changes were made except for an increase in

tungsten, a significant decrease in tantalum, and the removal altogether of titanium. Since lattice









mismatch has been shown to be significantly impacted by composition, and mechanical

properties as a result, these changes from the first to second generation alloys can also impact the

underlying deformation processes that occur during creep.22 Therefore, any significant change in

composition increases the risk of changes in mechanical properties such as the increased primary

creep strains discussed in the current investigation.

Each of the aforementioned examples of alloy developments from the last 20 years results

in changes in the intermediate temperature mechanical properties of these alloys. Besides the

individual changes brought about by increasing rhenium content, secondary y' precipitation, and

lattice mismatch variation, these three variations may all be interrelated. As is often the case in

alloy development, changes in behavior are the results of multiple factors. While significant

gains in performance may be made without a complete understanding of these behaviors, future

gains may rely on applying lessons learned from investigations into the Rhenium Effect, lattice

misfit, and aging heat treatment (among many other aspects of alloy design).

Modern Superalloys

Superalloys are known to combine a mix of high temperature tensile and creep strength

and environmental resistance. Their microstructure consists of two phases with a face-centered

cubic (fcc) matrix known as y, and a fcc-like L12 cuboidal precipitate known as y'. They are

strengthened through the use of solid solution and precipitation hardening techniques resulting in

high strength over a wide temperature range. Typically they are also directionally solidified to

produce a single grain through the Bridgman process, thus the application of the phrase "single

crystal" to describe these alloys when processed in this manner. Following solidiaication, these

alloys require long, high temperature heat treatments for proper homogenization and aging.

Several common alloy compositions representing the first three generations of superalloy










Table 2-1. Compositions in weight recent of several common Nickel-based supealos3.
Alloy Generation Co Cr Mo W Ta Re Al Ti Nb Hf Ni

Pratt &
White
PWA 1480 1st 5 10 4 12 5 1.5 Bal.
PWA 1484 2nd 10 5 2 6 9 3 5.6 0.1 Bal.

General
Electric
Rene N4 1st 8 9 6 6 4 3.7 4.2 0.5 Bal.
Rene N5 2nd 8 7 2 5 7 3 6.2 0.2 Bal.
Rene N6 3rd 12.5 4.2 1.4 6 7.2 5.4 5.8 0.2 Bal.

Cannon-
Muskeg~on
CMSX-2 1st 5 8 0.6 8 6 5.6 1 Bal.
CMSX-4 2nd 9 6.5 0.6 6 6.5 3 5.6 1 -0.1 Bal.
CMSX-10 3rd 3 2 0.4 5 8 6 5.7 0.2 0.1 0 Bal.

development are given in Table 2-1. Due to the extreme demand of service conditions, the

compositions of modern superalloys are quite complex, requiring 10 or more elemental

additions. Additionally, a wide variety of elements not shown in Table 2-1 are commonly added

to superalloys. These include, but are not limited to, carbon, boron, nitrogen, ruthenium, and

some rare earth elements.

The elemental addition that produced the greatest impact in properties during the early

stages of superalloy development was rhenium. Beginning with the second generation of alloys,

Re became an important factor for high temperature creep strength and microstructural stability.

The second and third generations of superalloys contain 3wt.% and 6wt.% Re, respectively. The

first generation of superalloys contain no Re. They do, however, contain the y' strengtheners Al,

Ti, and Ta and the refractory elements W and Mo for solid solution strength and microstructural

stability. As development continued, the need to restrict y' coarsening due to thermal exposure

became apparent. This led to the addition of Re to superalloys for it' s low diffusivity that









restricts the growth of primary y' precipitates. The diffusivity of rhenium in Nickel is so low that

it significantly retards the diffusion controlled coarsening of the y' precipitates. Additionally, Re

was found to be a very potent solid solution strengthener, thus the second generation of

superalloys was created.s, 10, 11

Continued development saw the addition of greater Re additions resulting in greatly

improved temperature capabilities and strength. The third generation contains 6wt.% and marks

a significant improvement in properties relative to the generations that preceded it. This increase

in Re concentration, however, came with a cost. Castability, as measured by incidence of casting

defects, was greatly decreased and the propensity for the formation of deleterious topologically

close packed (TCP) phases was greatly increased.1, 2, 8 Current research in fourth generation

alloys and beyond is focused on continuing to improve strength while minimizing the harmful

side effects of large Re additions through additions of platinum group metals, such as

ruthenium.28 New research was also initiated to find ways to maintain strength while reducing

Re concentrations due to the drastic increase in the cost of Re over the last 2 years.29

The state of the art of modern superalloys relies on a variety of alloy strengthening and

processing techniques developed over the last 30 years. As the high performance and low cost

demands continue to grow, the technology will need to grow. The next discussion will review

the challenges of these goals from the alloy and processing viewpoints.

The Challenge of High Temperature Service

The service environment experienced by turbine blades in modern gas turbine engines

demands high strength and environmental resistance at high temperatures. Normal alloy systems

are not capable of withstanding this type of environment. Structural alloys are typically

strengthened by a mix of solid solution strengthening, precipitation hardening, and strain









hardening. At temperatures above half the melting point (T>0.5Tm) strain hardening is less

effective due to recovery and recrystallization processes.30 Additionally, normal alloy systems

that are precipitation hardened lose strength at the temperatures experienced in gas turbine

engines because the precipitate solvus is exceeded, yielding a single phase material. As a result,

most normal alloy systems are insufficient for these applications.

In addition to tensile strength are several other properties to consider: creep strength, high

cycle and low cycle fatigue, ductility/toughness, thermo-mechanical fatigue, oxidation

resistance, and corrosion resistance. Each of these properties will impact the usefulness of an

alloy during high temperature service. Because these alloys are subjected to high temperatures

and high stresses for long durations during service, creep has been the subj ect of numerous

investigations over the past 40 years. At low temperatures, deformation processes are controlled

by dislocation motion (whole and partial dislocations). The movement of dislocations along

glide planes is governed by a variety of material properties like elastic modulus, dislocation

friction sources (peierls stresses), the presence of solute atmospheres (Cottrell atmospheres),

cross-slip difficulty, stacking fault energy, and anti-phase boundary energy. Add to these the

presence of secondary reinforcing phases and interfacial dislocation networks and the

intermediate temperature deformation processes of single crystal nickel base superalloys become

quite complex.31

Tensile Behavior

Low temperature deformation of superalloys is dominated by dislocation motion and

stacking fault propagation. At low temperatures, there is insufficient energy for diffusion

controlled processes and cross-slip is also more difficult, though not impossible. At high

temperatures, however, diffusion controlled processes become active. These include recovery,









vacancy motion, solute motion solutee drag), and phase instabilities (y' coarsening, TCP

precipitation, and carbide transitions).32-35

Despite all of the advances in superalloy technology, the basic microstructure employed

for single crystal turbine blades is quite simple. Generally, superalloys consist of an fcc matrix

with coherent, stress-free ordered L12 precipitates. Most first and second generation alloys

contain about 60-70% y': the optimum value for excellent creep and tensile properties.

Maintaining low misfit through alloying has also been shown to effect change in the morphology

of the y' precipitates. Even alloys with "high" lattice misfits (181>0.5%) will exhibit less than

one percent misfit between the y and y' phases. Changing misfit values can produce y'

precipitate shapes varying from spherical to cuboidal to dendritic. The so-called cuboidal

morphology is typically desired for high temperature creep resistance with the ideal precipitate

edge length between 0.35Cpm and 0.45Cpm.36 The cuboidal structure consists of uniform cubes of

y' with rounded edges and corners. The resulting interstices between y' particles consists of the y

matrix. These y "channels" are relatively large in length and width, but narrow in thickness

(typically <50nm).

Strengthening of superalloys can be achieved through alloying and heat treatment. The

most common goals involved in strengthening superalloys through alloying can be broken down

into several categories:

* Solid solution strengthening of y matrix
* Modifying misfit to produce cuboidal y' precipitates
* Increasing y' volume fraction to the ideal limit
* Solid solution strengthening of y' precipitates
* Modifying the anti-phase boundary, APB, energy of the y' phase
* Stabilizing the microstructure with low diffusivity elements (W, Mo, and Re for example)









The addition of carbide phases may also help with fracture strength, but is not likely to

affect yield strength due to their large spacing compared to the dislocation spacing.37 Another

consideration is the prevention of TCP phases. Alloys with large refractory element

concentrations are at risk of developing TCP phases during long exposure to high temperatures.

While the prevention of TCP phases does not necessarily lead to improved strength, the

formation of TCP phases is suspected to result in premature failure of superalloys and should be

avoided.35

Solid solution strengthening of the matrix is typically accomplished by increasing

additions of Cr, W, Mo, and/or Re. While these are suitable for solid solution strengthening,

secondary concerns may dictate their use. Increasing the W, Mo, and Re content of an alloy may

result in decreased TCP resistance and an increase in the y/y' misfit as these elements are mostly

rej ected from the y' phase.8, 35 Increasing the volume fraction and strength of the y' precipitates

is typically achieved by additions of Al, Ti, and Ta. Controlling the y/y' misfit and APB energy

is a difficult and not fully understood process; however. They are usually maintained through

slight adjustment to the entire set of elemental additions until the desired value is achieved as

virtually every element can affect misfit and/or APB energy. Finally, one of the benefits of

adding slow-diffusing elements like W and Re for solid solution strengthening is that they also

serve to stabilize the y/y' microstructure. With a diffusivity for self-diffusion within nickel that

is much lower than any other elements, y' coarsening kinetics are reduced allowing for

precipitate coherency to be maintained longer.8

Tensile strain of superalloys usually begins in the y matrix phase. Dislocations are created

on {111}<110>)'; slip oplae which are~, thenalet pr~ moopagt through the material. As in other

fcc materials, these dislocations will often dissociate into Shockley partial dislocations of the










a/6 {111}<112> type with a stacking fault formed in between. As the strain increases, these

dislocations will squeeze through the y matrix channels as described by the orowan bowing

model. As this process progresses, dislocations begin to lie along the surfaces of the cuboidal y'

precipitates eventually forming interfacial networks of dislocations. Continued strain gradually

reduces the distance between dislocations in the interfacial networks. As the distance gets

smaller, the strength increases due to work-hardening.24, 27, 33, 37 The tensile behavior of single

crystal, nickel base superalloys may also display a yield-point on an engineering stress vs.

engineering strain graph. This phenomenon is not unexpected as it is commonly linked to

interstitial or substitutional impurities and these alloys contain both in the form of alloying

additions.31

Creep Behavior

When superalloys are subj ected to tensile loads at temperatures greater than 5000C, creep

can occur. Creep is a time-dependent plastic deformation process that occurs at loads below the

yield stress of the material in question. Creep occurs because the thermal energy that is available

at high temperatures allows for thermally activated processes that do not occur at low

temperatures to become viable deformation mechanisms. In the case of single crystal

superalloys, vacancy and solute diffusion, dislocation climb, cross-slip, and activation of

secondary slip systems are all processes that are either difficult or non-existent at low

temperatures. Additionally, microstructural evolution may occur resulting in changes to the

creep behavior over time. Superalloys are prone to the following microstructural changes in

particular: y' coarsening and loss of coherency, topological inversion (rafting) of the y' phase,

precipitation of additional phases (especially topologically close packed, TCP, phases), and

carbide transition/precipitation (MC, M6C, and M23 6).35, 38-41 It is the activation of these










processes that allow for metals to deform over time at loads significantly lower than their yield

stes25, 31, 33, 42, 43

Mechanically, creep tests can be conducted in one of two ways. To obtain a true measure

of creep requires that the specimen be subj ected to a constant stress throughout the duration of

the test. Consequently, as the specimen deforms and the cross-sectional area decreases, the load

would be reduced accordingly. Unfortunately, though, creep tests must be run at high

temperatures within a furnace so accurate measurements of the specimen dimensions are not

feasible during testing. Because of this problem, another test procedure has been adopted for use

in most cases. This is the constant load creep test (contrasting with the constant stress test

discussed above). A constant load creep test is performed by applying the correct amount of

weight to the specimen to produce the desired initial stress. As the specimen deforms and the

cross-section is reduced, no changes are made to the load. The consequence of this design is that

the stress level increases throughout the creep test. This investigation uses constant load creep

tests because of the simplicity of data acquisition and commercial relevance, so the following

discussion refers to data collected with this type of creep test.31

For single crystal, nickel base superalloys, there are primarily two temperature regimes of

interest that involve creep. At high temperatures, greater than 8500C, creep tests are run with

relatively low stresses and are used to simulate the behavior of the thin, exposed parts of a

turbine blade. This environment allows for thermally activated processes to occur at a rapid

pace. At temperatures below 8500C, the loads are significantly greater. The performance of

alloys under intermediate temperatures and high loads depends largely on the tensile strengths of

the alloys being investigated. Because of the link between high loads and intermediate










temperatures to large primary creep strains, this set of conditions is the focus of this

investigation.

Summary

In order to conduct an investigation into the primary creep behavior of the superalloys

PWA 1480 and PWA 1484, specimens will be creep tested at temperatures between 7000C and

8150C and loads greater than 621 MPa. Tensile testing, metallographic examination, X-ray

diffraction (XRD), and transmission electron microscopy (TEM) have been conducted to gain an

understanding of the underlying factors involved in creating large primary creep strains. Three

alloys have been investigated including the first generation PWA 1480, the second generation

PWA 1484, and the experimental second generation PWA 1480+Re. These three alloys helped

to shed light on several aspects of primary creep behavior and the effect of age heat treatment.

Finally, the current investigation may lead to an improved understanding of the effect of rhenium

on single crystal nickel base superalloys.










CHAPTER 3
EXPERIMENTAL PROCEDURES

Materials

All three alloys used in this investigation were provided by Pratt & Whitney Aircraft

Engines in East Hartford, CT. Test bars were prepared by Pratt & Whitney in investment casting

cluster molds of 12 bars measuring approximately 20 cm (8 in.) long with a 1.6 cm (0.625 in.)

diameter. Single crystals were prepared using Bridgman style directional solidification furnaces

and single crystal selectors to produce an orientation in the [001] direction. Following

solidification, the bars were examined through the use of Laue X-Ray Diffraction to verify that

all the bars were near the desired [001] orientation. Laue orientation data and original heat

treatments for all the material is presented in Tables 3-1, 3-2, and 3-3. Test material was

Table 3-1. Laue orientation data and original heat treatments for PWA 1480.*
Master Bar oc Solution Heat Coating Heat
Heat Number angle Treatment Treatment
P9976 0401 4.2 12880C /2hr. 10800C /4hr.
P9976 0402 5.2 12880C /2hr. 10800C /4hr.
P9976 0403 6.2 12880C /2hr. 10800C /4hr.
P9976 0404 7.3 12880C /2hr. 10800C /4hr.
P9976 0405 8.8 12880C /2hr. 10800C /4hr.
P9976 0406 9.5 12880C /2hr. 10800C /4hr.
P1083 0901 2.5 As-Cast As-Cast
P1083 0902 2.1 As-Cast As-Cast
P1083 0903 19.3 As-Cast As-Cast
P1083 0904 4.0 As-Cast As-Cast
P1083 0905 3.0 As-Cast As-Cast
P1083 0906 0.9 As-Cast As-Cast
P1083 0907 4.8 As-Cast As-Cast
P1083 0908 4.5 As-Cast As-Cast
P1083 0909 3.1 As-Cast As-Cast
P1083 0910 3.9 As-Cast As-Cast
P1083 0911 2.3 As-Cast As-Cast
P1083 0912 4.5 As-Cast As-Cast
*Bar number 0903 was rejected due to an excessive oc value.










Table 3-2. Laue orientation data and original heat treatments for PWA 1484.
Master Bar a Solution Heat Coating Heat
Heat Number angle Treatment Treatment
P1096 0407 0.5 13100C /0.5hr. 10800C /4hr.
P1096 0408 3.8 13100C /0.5hr. 10800C /4hr.
P1096 0409 5.7 13100C /0.5hr. 10800C /4hr.
P1096 0410 7.2 13100C /0.5hr. 10800C /4hr.
P1096 0411 9.4 13100C /0.5hr. 10800C /4hr.
P1096 0412 10.1 13100C /0.5hr. 10800C /4hr.
P1086 0913 4.0 As-Cast As-Cast
P1086 0914 7.0 As-Cast As-Cast
P1086 0915 5.7 As-Cast As-Cast
P1086 0916 7.5 As-Cast As-Cast
P1086 0917 8.5 As-Cast As-Cast
P1086 0918 7.0 As-Cast As-Cast
P1086 0919 5.6 As-Cast As-Cast
P1086 0920 8.1 As-Cast As-Cast
P1086 0921 5.9 As-Cast As-Cast
P1086 0922 3.1 As-Cast As-Cast
P1086 0923 6.5 As-Cast As-Cast
P1086 0924 5.9 As-Cast As-Cast

Table 3-3. Laue orientation data and original heat treatments for PWA 1480+.
Master Bar a Solution Heat Coating Heat
Heat Number angle Treatment Treatment
P1106 1201 12.6 As-Cast As-Cast
P1106 1202 15.5 As-Cast As-Cast
P1106 1203 5.7 As-Cast As-Cast
P1106 1204 1.6 As-Cast As-Cast
P1106 1205 9.3 As-Cast As-Cast
P1106 1206 10.2 As-Cast As-Cast
P1106 1207 9.0 As-Cast As-Cast
P1106 1208 0.3 As-Cast As-Cast
P1106 1209 9.1 As-Cast As-Cast
P1106 1210 12.6 As-Cast As-Cast
P1106 1211 4.2 As-Cast As-Cast
P1106 1212 3.1 As-Cast As-Cast

provided in two stages. First, a supply of 6 bars of PWA 1480 and 6 bars of PWA 1484 was

delivered having been subj ected to the standard commercial solution heat treatments (alloy










specific) used by Pratt & Whitney as well as a coating heat treatment cycle. The bars were not,

however, subjected to an aging heat treatment. The second delivery of material expanded the

testing possibilities significantly. First, another 1 1 bars of PWA 1480 and 12 bars of PWA 1484

were provided in the as-cast condition. Then, 12 bars of PWA 1480 with 3 wt% Re added were

produced and delivered in the as-cast condition. Altogether, 47 single crystal test bars, allowing

for the production of up to 94 creep/tensile specimens, were available for experimentation.

Heat Treatment

The first step in the preparation of test specimens was to heat treat all the bars to an

appropriate condition. As stated above, the first batch ofPWA 1480 and PWA 1484 had already

been given solution heat treatments (HT1), Table 3-4. Additionally, this first batch had already

received a coating simulation cycle as these alloys are commonly used in the first and second

turbine stages of gas turbine engines where coatings are necessary for thermal and environmental

protection. Because one aim of this study is to determine the cause of changes in creep behavior

due to age temperature, two different aging heat treatments were given to each alloy. Half of the

PWA 1480 bars were given a low temperature age (LT, 7040C/24 hr.) and half were given a high

temperature age (HT, 8710C/32 hr.) as shown in Table 3-4. The same aging heat treatments

were also given to the PWA 1484 bars. Following aging, this first batch of material was ready

for sample machining and subsequent testing.

Heat Treatment Development

The remaining test bars, as-cast PWA 1480, as-cast PWA 1484, and as-cast PWAl480+,

still required a complete heat treatment cycle. Following early results and metallography of the

first batch with the standard solution heat treatments, it was determined that a longer time, higher

temperature solution heat treatment might be helpful to reduce the effects of chemical









Table 3-4. Heat treatments used in this study. Solution heat treatment designations (HT#) are given.*
First Batch
Alloy Solution Heat Treatment Coating Heat Treatment LT Age HT Age
PWA 1480 HT1 1290oC/2hr./GFQ 1080oC/4hr./GFQ 704oC/24hr./AC 87 1oC/3 2hr./AC
PWA 1484 HT1 1 3 10OC/0.5Shr./GFQ 1080oC/4hr./GFQ 704oC/24hr./AC 87 1oC/3 2hr./AC

Second Batch
Alloy Solution Heat Treatment Coating. Heat Treatment LT Age HT Age
PWA 1480 HTO None
HT1 1290oC/4hr./GFQ (STD HT)
HT2 1290oC/8hr./GFQ
HT3 1295oC/1hr.- 12990C/1hr. 1080oC/4hr./GFQ 704oC/24hr./AC 87 1oC/3 2hr./AC
-13 02oC/10Ohr./GFQ
PWA 1484 HTO None
HT1 1315oC/4hr./GFQ (STD HT)
HT2 1315oC/8hr./GFQ
HT3 1 325oC/1hr. 13 3 2C/1 hr. 1080oC/4hr./GFQ 704oC/24hr./AC 87 1oC/3 2hr./AC
-133 80C/10Ohr./GFQ
PWA 1480 + HTO None
HTIA 1290oC/4hr./GFQ
HTIB 1315oC/4hr./GFQ
HT2 1280oC/1hr.- 1290oC/7hr./GFQ
HT3 1293o"C/1 hr. 1296oC/1hr. 1080oC/4hr./GFQ 704oC/24hr./AC 87 1oC/3 2hr./AC
-12990C/2hr. -13 02oC/1 hr.
-13080C/4hr./GFQ
*All samples were heated from room temperature to 12000C for 0.5hr. at a rate of 200C/minute, then samples were heated to the first soak temperature at a rate
of 100C/minute. All ramping following the first soak temperature is at a rate of loC/minute. GFQ is Gas Furnace Quench (ultra high purity He) and AC is Air
Cool. Solution and Coating heat treatments were performed under vacuum and Age heat treatments were performed in air environments.









segregation and the formation of eutectic regions during solidification. Previous research has

also shown that incomplete y' solutioning and dendrite homogenization can lead to substantial

reductions in creep life.36, 44 To counter this effect, a series of heat treatment trials beginning

with the standard Pratt & Whitney developed heat treatments was performed.

The first step taken to design a new solution heat treatment was to increase the amount of

time of the original heat treatment. Increasing the maximum temperature hold time from 2 hours

to 8 hours for PWA 1480 and 0.5 hours to 8 hours for PWA 1484 did not significantly improve

the homogeneity of the samples; however, it proved a useful heat treatment for Differential

Thermal Analysis (DTA). With samples in the as-cast condition DTA is only moderately useful

as the solidus and solvus temperatures tend to be suppressed and transition temperature peaks

tend to be broad due to high segregation. Following even a brief heat treatment, however,

meaningful data can be obtained as described below.45

Differential Thermal Analysis

Differential Thermal Analysis was performed on all three alloys following the extended

version (HT2) of the standard solution heat treatment (HT 1). DTA was performed by Dirats

Laboratories in Westfield, MA using a DuPont 9000 Thermal Analyzer. Data was collected at a

rate of 200C/minute on heating only to eliminate the effects of undercooling. Samples were cut

from the single crystal bars to produce a disk approximately 0.5 cm (0.188 in.) thick by 1.6 cm

(0.625 in.) in diameter. The HT2 treatment was applied to all three alloy groups for DTA. All

samples were run with pure nickel standards. Additional discussion of DTA practices is

provided in Appendix A along with the complete DTA scan profies generated for each sample.

The resulting data allowed for determination of the solidus, liquidus, and y' solvus temperatures

for the three alloys and are given in Table 3-5. Clearly, the increased homogeneity of the









Table 3-5. DTA results for all three alloys as-received (HT1/HTO) and heat treated (HT2).*
Alloy PWA 1480 PWA 1484 PWA 1480+
Heat Treatment HT1 HT2 HT1 HT2 HTO HT2
y' Solvus (oC) 1288 1298 1291 1300 N/F 1298
Solidus (oC) 1304 1307 1338 1346 1291 1311
Liquidus (oC) 1333 1344 1378 1392 1339 1352
*The y' solvus was not found for PWA 1480+, HTO.

samples following the HT2 heat treatment resulted in an increase in all of the transformation

temperatures. Using these temperatures as a guide, a new heat treatment (HT3) scheme was

developed to improve the degree of solutioning and homogenization.

The third heat treatment included brief hold steps at a few temperatures prior to the final

soak temperature. Lower temperature steps were used in the early stages of the heat treatment to

prevent incipient melting in the interdendritic regions of the alloy where the solidus temperature

is suppressed. The final soak was designed to be as close to the solidus of each of the alloys as

possible in order to give the best possible degree of homogenization within the alloy. The final

heat treatment (HT3) significantly reduced the presence of y/y' eutectics and increased the degree

of homogenization within the alloys.

Furnaces

Solution and coating simulation heat treatments were given to all three alloys using an

Elatec high temperature vacuum furnace. The furnace is capable of temperatures up to 1400 oC

under a vacuum greater than 10-4 Torr. The heating elements, hearth plate, and j ail are fabricated

from graphite. Alumina trays were used to separate the test bars from the graphite hearth plate.

Temperature control is provided by three type C (W + 5% Re / W + 26% Re) thermocouples all

sheathed in individual molybdenum jackets. Two thermocouples were lowered very near the

surface of the test bars in the center of the hearth plate. These were used for controlling the










furnace temperature to within approximately 1.5 OC of the setpoint. The third thermocouple is

a survey thermocouple used for monitoring the temperature near the front of the hot-zone.

Following all heat treatments in the vacuum furnace, samples and test bars were cooled by

inj section of ultra high purity helium and the circulation of the helium through a water cooled

copper heat exchanger within the furnace by a fan. This rapid quenching technique allows for

high cooling rates greater than 250 oC/minute down to approximately 500 oC. This rapid cooling

rate minimizes the precipitation and growth of y' during cooling. Below 500 oC, cooling slows

to between 100-200 oC/minute, however, diffusion in the samples at this point has slowed

sufficiently to prevent significant y' coarsening. Furnace control is managed by a Honeywell

controller that manages temperature, vacuum, gas quenching, and all valves.

Aging heat treatments (both LTA and HTA) were performed in Carbolite box furnaces

with maximum operating temperatures of 1300 oC. The atmosphere is not controlled for these

heat treatments as the temperature is typically low enough and the time is short enough that

oxidation is not problematic. Temperature control was maintained through the use of two Type

K thermocouples placed in direct contact with the bars inside the furnace. Temperatures in the

box furnaces were maintained within 13 OC of the setpoint as verified by handheld digital

thermometers. Following aging heat treatments, samples and test bars were removed and

allowed to air cool on alumina racks. Cooling rates exceeded 100 oC per minute to prevent

undesirable y' coarsening.

Characterization

Throughout this investigation, sample characterization was crucial to understanding what

changes had taken place in the alloys. A variety of characterization methods were used to

examine samples including the already mentioned Differential Thermal Analysis. Other methods









include metallographic techniques involving optical and electron microscopes, fractography

following creep and tensile testing, and transmission electron microscopy (TEM) on interrupted

creep specimens to examine deformation structures.

Preparing Samples for Metallography

Samples were taken from test bars and creep/tensile specimens for metallographic

examination. The same basic process can be applied to a variety of sample shapes and

geometries. Metallographic preparation was conducted according to the ASM Metals Handbook

recommendations. Samples were first sectioned from the bulk by the use of either a Leco

abrasive cut-off saw or an Allied slow-cut diamond sectioning saw. Both saws are liquid cooled

to keep cutting temperatures low. Once sectioned, the specimens were ground flat and polished.

Specimens were first leveled using silicon carbide grinding papers on an 8 inch Leco

metallography wheel, then polished using alumina powder and water suspensions. All

metallography specimens were polished to 0.3 Cpm alumina polishing media. Following

polishing, the specimens were etched with a y' etchant developed by Pratt & Whitney (100 mL

HC1, 100 mL HNO3, 10g MoO3, 100 mL H20). The etchant was applied with cotton tipped

applicators and was swabbed evenly about the surface until the surface of the specimen appeared

hazy. Specimens were then rinsed with water, then methanol, and were finally dried under a jet

of compressed air.

Additionally, a second etchant was used to etch away the y matrix phase. This etchant was

an electrolytic etch utilizing a 20% oxalic acid solution at 20 V. The specimens were mounted

on a metal stub to maintain the necessary electrical conductivity from the specimen to the power

supply and were connected as the anode. The counter electrode (cathode) was a 500 mL

stainless steel beaker. The electrolyte was placed in the beaker and the steel beaker was placed









in an ice water bath to keep the electrolyte from heating. The voltage was applied and the

specimens were dipped for no longer than 5 seconds. The specimens were immediately rinsed

and dried and were ready for metallographic inspection.

All Specimens were then observed on a Leco optical metallograph, a JEOL 6400 tungsten

filament scanning electron microscope, or a JEOL 6335F field emission scanning electron

microscope. Both electron microscopes were equipped with Energy Dispersive Spectroscopy

(EDS) detectors as well. Images taken on these microscopes were used for examination of

microstructures during the heat treatment development cycle. Additionally, post test

microstructures and fracture surfaces were observed to aid in the understanding of

microstructural changes during high temperature testing.

Preparing Samples for TEM

Making specimens for transmission electron microscopy (TEM) involves the same

metallographic techniques described above. Interrupted and full-length creep tests were

sectioned using the Allied slow-cut diamond saw in the transverse and longitudinal directions.

These samples were mounted on aluminum stubs with mounting wax and polished to 0.3 Cpm

alumina powder suspension. The samples were then transferred to special mounts for use with

the Focused lon Beam (FIB) or were thinned for use with the twin j et electropolisher.

The FIB is a scanning electron microscope with an attached Gallium ion beam for milling

the sample. Using the FIB, TEM liftout specimens were cut perpendicular, parallel, and at 450 to

the stress axis. On PWA 1484 samples, it was possible to obtain TEM specimenss near [001],

[01 1], and [1 11] zone axes from a single sample due to the orientation based distortion of PWA

1484 during creep testing described in later chapters. Lift-outs were placed on 200 mesh Carbon

coated Copper girds using micromanipulators.









Alternatively, specimens were also prepared using a twin jet electropolisher. After

samples were sectioned from the interrupted creep specimens, as described above, they were

thinned by hand to between 100 and 200 pm. Disks of 3 mm diameter were then removed from

the thinned material with a TEM punch. The samples were further thinned using a die polisher

with a micrometer to between 30 and 70 Cpm thick. Electropolishing was performed on the 3 mm

disks with a solution of 90% methanol and 10% perchloric acid and a voltage of 20V. Liquid

nitrogen was added to maintain the temperature of -250C. Transmission electron microscopy

was conducted using JEOL 200 CX. The TEM was used to for qualitative observation of

microstructural changes, dislocation and stacking fault behavior, and variations in y/y' misfit.

Preparing Samples for LEAP

Local Electrode Atom Probe (LEAP) specimens were prepared from a single bar of PWA

1484 that had already received a solution heat treatment (HT3) and a coating simulation heat

treatment. The bar was first cut in four equal pieces on a water cooled Leco abrasive cut-off

wheel. Then, two sections (out of four) were subjected to either the LTA or HTA aging heat

treatments. One piece from both groups was then separated and subj ected to a brief heat

treatment above the y' solvus (30 minutes at 1310 oC) in order to redissolve the y' and quickly

cool before the slow diffusing rhenium atoms have enough time to redistribute, Figure 3-1. This

heat treatment was performed in an attempt to form a better understanding of rhenium

segregation behavior at the y/y' interfaces. Once the samples were heat treated, they were sent to

AMT, Clifton Park, NY to be sectioned into specimens by electrical discharge machining

(EDM). Specimens were 5 cm long cylinders with a diameter of 1.8 mm. Ten specimens were

obtained from each of the four samples mentioned above.




























Figure 3-1. Flow-chart of heat treatments for LEAP samples.

The LEAP specimens were then electropolished to reduce the diameter of the cylinder and

form a point on one end. The electropolishing unit consisted of a TekPower HY 3005D-3 power

supply with 2 variable outputs each capable of 30 V and 5 A. If linked in series or parallel it was

capable of 60 V at 5 A or 30 V at 10 A, respectively. The cathode was a 500 mL stainless steel

beaker in an ice bath to control the electrolyte temperature. The electrolyte was 90% ethanol and

10% perchloric acid. The specimen was connected to the power supply as the anode and was

held in the beaker by a brace such that the cylinder was centered along its length to reduce the

effect of varying anode to cathode distance. The applied voltage was varied between 5 V and 20

V until stable polishing took place (typically betweenl0 V and 15 V). Polishing was stopped

once the specimens reached a diameter around 500 Cpm to 700 Cpm, which correlates to a tip

diameter of approximately 100 pm.

The specimens were then shipped to the Materials Science and Engineering Department at

the University of North Texas (Denton, TX). Anantha Puthucode, Ph.D., and Michael Kaufman,

Ph.D., of the University of North Texas provided final sample preparations using an Imago









Electropointer. The electropointer used a 3 mm Platinum loop as a cathode to refine the tip

radius to less than 75 nm, Figure 3-2. Additional sharpening was occasionally performed using

FIB based techniques as well. Once the tip was formed, specimens were run on the Imago LEAP

3000X. Data was then analyzed and manipulated using the Imago IVAS analysis software

package. Specimens were run multiple times by reforming the tip following the experiment.

Data from these experiments were returned in the form of 3-dimensional and 2-dimensional

composition maps with a lateral resolution of 0. 1 nm. Additionally, 1-dimensional line scans

were also simulated using the same data sets.




















Figure 3-2. SEM micrograph of the tip of a LEAP specimen.

Preparing Samples for X-Ray Diffraction

Specimens were also prepared for X-Ray Diffraction (XRD) to study the y/y' misfit in all

three alloys. Following the coating heat treatment cycle, bars from each alloy were sectioned to

create 4 disks of each alloy 16 mm in diameter by 10 mm thick. These disks were then separated

into 4 groups. The first group was tested with no further heat treatment resulting in a total heat

treatment time at 10800C of 4 hours. The other three groups were heat treated an additional 6,










96, or 996 hours to produce samples with total heat treatment times at 10800C of 10, 100, and

1000 hours. Specimens were cut from the disks at a thickness of 2 mm. The specimens were

then thinned and polished on both sides to create a Einal thickness between 0.3 and 0.7 mm. The

polished surfaces were both made parallel to the (001) plane.

All samples were run on Rigaku 6/26 powder diffractometers at the University of Central

Florida' s Advanced Materials Processing and Analysis Center (AMPAC). Routine scans were

made from 100 to 1300 at a scan step of 0.020 and a 1 second dwell time. Local scans of the

(002) peak, centered around 26 = 510, and the (004) peak, centered around 26 = 1180, were taken

with a smaller scan step of 0.010 and a longer dwell time ranging from 3 to 5 seconds. Some

specimens were also run at 0.010 steps and 30 second dwell times in an attempt to improve

signal clarity. The resulting intensity data were then indexed.

The local scans of the (002) and (004) peaks were used for lattice mismatch calculation.

Due to the structure factors and lattice parameters of the fcc y phase and the L12 y' phase, the

peaks produced from both phases overlap at the (002) and (004) positions. In order to separate

the contributions, peak deconvolution was performed utilizing the MDI Jade (ver. 7) software

package and in agreement with published work.46 Deconvolution practices included fitting

skewed and unskewed Gaussian profies to account for both phases as well as for the Cu km1 and

Cu ku2 wavelengths. A complete discussion of the deconvolution procedure and results are given

in Appendix B. Once the centers of the kml peaks for both phases are determined, lattice

parameters and mismatch were determined. Lattice mismatch values in this investigation were

calculated using Equation 3-1. Where 6 is lattice misfit (multiply by 100 to covert to percent), aY

is the lattice parameter of the y phase, and aY is the lattice parameter of the y' phase.










2(ar' r)
Equation 3-1 -=


JMatPro Thermodynamic Prediction

The thermodynamic modeling and prediction software package, JMatPro (United

Kingdom), was used to predict phase compositions, lattice misfit, and volume fractions of phases

as a function of temperature. While these predictions are useful and fast, they are not guaranteed

to be accurate for all alloy compositions. To enhance usability and accuracy, the models used for

JMatPro have been calibrated through the use of several common benchmark alloys. PWA 1480

is one such benchmark alloy. PWA 1484 and PWA 1480+ predictions, as a consequence of not

being used to calibrate the software results, will require some degree of extrapolation to make

thermodynamic predictions such as cooling curves, expected compositions, and lattice misfit.

Once an alloy composition has been selected, thermodynamic data can be viewed as a function

of temperature or calculated at specific temperatures. Data from these predictions were used as a

helpful guide and as potential indicators of interesting behaviors. Conclusions based on software

results alone, however, should be verified through experimental testing and, for this reason,

characterization was conducted to verify several aspects of the JMatPro predictions.

Mechanical Behavior

Following the Einal aging heat treatment, the test bars were sent to Joliet Metallurgical

Labs in Joliet, IL for Einal machining. Each bar was machined into two creep/tensile specimens

according to the drawing in Figure 3-3. The same specimen geometry is suitable for either high

temperature tensile testing or creep testing. The specimen gauge length was 2.60 cm and the

gauge diameter was 4.5 mm. Specimens were machined from the test bars using low stress

grinding techniques to prevent strain hardening along machined surfaces. Following machining,












all specimens were measured to verify all dimensions of interest, most notably gauge length and


gauge diameter.


Tensile Testing


High temperature tensile testing was conducted at 7040C and 8150C to compare the


strengths of the three alloys with both age heat treatments. The tensile test matrix can be seen in


Table 3-6. Fixtures were lubricated with boron nitride high temperature lubricant and the


specimens were threaded into the grips. An Instron servo-hydraulic load frame with a 20,000 lb.


load cell was used in coordination with the Merlin control and data acquisition software package.






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Figure 3-3.


i 11111


Creep and tensile specimen geometry and dimensions.









Table 3-6. Tensile test matrix with sample identification.
Age HT LT LT HT HT
Test Temp. (oC) 704 815 704 815

PWA 1480 0402-2 0403-2 0405-2 0406-2
0905-1 0901-2 0906-1 0906-2

PWA 1484 0408-1 0409-2 0411-2 0411-1
0915-1 0915-2 0919-1 0919-2

PWA 1480+ 1205-1 1203-2 1207-1 1209-2


A clam-shell style furnace capable of temperatures up to 10000C was used to control the

temperature to within 1.70C of the setpoint throughout the test. Temperatures were maintained

by two Type K thermocouples with Nextel high temperature ceramic insulation for thermal

protection. Both thermocouples were tied to the gauge section of the specimen with 24 gauge

80Ni-20Cr wire. Prior to starting the test, all specimens were subjected to a 15 minute soak

period to ensure uniform temperature of the sample and fixtures. All tensile tests were

performed in air.

During tensile testing, a constant cross-head speed of 0.25 cm/min. was used for all tests.

This cross-head speed corresponds with an initial strain rate of 0.25 cm/cm/min (0. 1 in./in./min.).

Data acquisition was made possible by the use of a high temperature extensometer frame that

attached to the specimen through the use of knife edges. This extensometer frame then extended

down out of the furnace to an easily accessible area where a digital extensometer connected to

the computer could be attached. The Merlin software package (by Instron) was used to record

pre-test sample dimensions, control the servo-hydraulic system, and record load, position, and

extensometry data during the test. Data recorded during a test were subsequently analyzed to

determine yield strength, ultimate tensile strength, percent elongation, and failure strength.










Creep Testing

Creep tests were conducted on Satec M-3 style creep frames. The tests conducted for this

study were constant load creep tests. Each frame has a weight pan for applying the load and a 16

to 1 lever arm ratio to load the specimen. As in the tensile tests, all threaded connections were

lubricated with boron nitride high temperature lubricant. Specimens were also affixed into high

temperature extensometers as in tensile testing through the use of either knife edges on the gage

sections or set screws on the shoulders. The bases of the creep extensometers, though, included

fixtures for Linear Variable Differential Transducers (LVDT) for measurement of displacement

during the test. Again, clam-shell furnaces capable of 12000C or 13000C (depending on which

frame was used) heated the specimens to the desired temperature for the duration of the creep

test. Temperature control is maintained by the use of 3 Type K thermocouples attached to the

gauge section of the specimens by 24 gauge 80Ni-20Cr wire as shown in Figure 3-4.

Temperatures were maintained to within +1.70C. Following heating to the test temperature, all

specimens enter a soak period for 1 hour prior to loading and the commencement of the creep

test. Hot loading was performed in multiple steps and displacement measurements were taken

allowing for calculation of the elastic modulus.

The NuVision Mentor Creep Controller software (Satec) allowed for control of the creep

frames, furnaces, and data collection. Creep data was collected at a rate of 5 times per minute

for the first hour, then 1 time per minute for the remainder of the test. Additionally, times to

0.1%, 0.2%, 0.5%, 1%, 5%, and 10% creep were recorded automatically. Test specifications

were written to conduct creep tests at four conditions: 7040C / 862 MPa, 7040C / 758 MPa,

7600C / 690 MPa, and 8150C / 621 MPa. These conditions were chosen because of their

similarity to real-world service conditions and to generate the large primary creep strains as



























Figure 3-4. Three Type K thermocouples are attached to the gauge length of creep specimens
with 80Ni-20Cr wire.

reported in the literature.16, 20, 47 The creep test matrix used for this study can be seen in Table 3-

7. Full length and Interrupted creep tests were conducted to obtain a variety of useful data. Full

length tests are those that were allowed to run to failure. Interrupted tests are those that were

stopped prior to failure in one of two conditions: after 0.5% secondary creep and tests running

longer than 1200 hours with no sign of imminent failure.

Table 3-7. Creep test matrix with sample identification. Creep tests were run to failure or
terminated following 0.5% secondary creep.
LT Age HT Age
Tem (C) 704 760 815 704 760 815
Tests run to PWA 1480 0401-1 0401-2 0403-1 0404-1 0405-1 0406-1
failure 0402-1 0403-2 0907-1 0907-2
0902-1 0902-2
PWA 1484 0407-1 0408-2 0409-1 0410-2 0410-1 0412-1
0913-1 0913-2 0412-2 0917-2
0917-1
0920-1
PWA 1480+ 1203-1 1205-2 1209-1 1207-2


PWA 1480 0903-1 0903-2 0908-1 0908-2
PWA 1484 0914-1 0914-2 0918-1 0918-2
PWA 1480+ 1204-1 1204-2 1210-1 1210-2


Tests stopped
after 0.5%
secondary creep









CHAPTER 4
RESULTS: ALLOY MICROSTRUCTURES

The three alloys used in this study represent first and second generation superalloys.

Typically, nickel based superalloys may contain two or more phases including: y (fcc Ni solid

solution, matrix), y' (L12 Ni3Al), Carbides (MC, M6C, Or M23 6), and/or Topologically Close

Packed (Sigma, Mu, Laves, and/or P phases). The occurrence of several of these phases in the

three alloys investigated here is discussed as a result of heat treatment and mechanical testing.

All three alloys contain primary y phase as a matrix with a high volume percent of y' precipitates

following solution heat treatment and aging. The presence of a low amount of Carbon results in

the formation of both script and blocky carbide phases for all three alloys. PWA 1480 and PWA

1480+ contain the greatest amount of carbide phases possibly as a result of a lower solubility for

Carbon in the alloys. This effect is particularly prominent in PWA 1480+ with the Re addition

resulting in an increase in carbide amounts when compared to PWA 1480 without Re.

Additionally, the alloy PWA 1480+ displayed poor resistance to the formation of TCP phases.

The following discussion will focus on the properties of the phases themselves and their

incidence as a result of casting and subsequent heat treatment as well as any changes during

mechanical testing. Due to the large number of tensile graphs, creep graphs, micrographs, and

other figures presented in the current and following results chapters, all figures have been placed

at the end of the respective chapters (Chapters 4-7) to aid reading speed and comprehension.

Phase Descriptions

The y Phase

The matrix of single crystal nickel base superalloys consists of the fcc solid solution y

phase. Due to the large volume fraction of y' used for high strength and creep resistance in these

alloys, the y phases only constitutes around 25-3 5% of the alloy by volume. Topographically,










the y phase forms in the long narrow regions between the cuboidal (cubes with rounded edges) y'

precipitates. These "channels" are typically 50-150 nm thick by 0.3-0.5 Cpm in length and width.

The length and width dimensions, obviously, are controlled by the local y' size and distribution.

The y phase from all three alloys can be seen in Figures 4-1 through 4-3 as the light gray regions

between the darker squares (y' precipitates).

Compositionally, there are significant differences between the y and y' phases. For as long

as superalloys have been in development there have been research investigations into the

compositions of the two phases. In general terms, the refractory elements added for solid

solution strengthening partition to the y matrix and the so called "y' strengtheners" partition to

the y' phase. Specifically, Re has been shown to partition especially strong to the y phase,

however the addition of W can result in up to 20% of the added Re partitioning to the y' phase,

effectively reducing the Re content of the y matrix.35

The concentration of the y phase can be predicted utilizing the JMatPro thermodynamic

prediction software. The calculated equilibrium concentration of the y phase as a function of

temperature is given in Figures 4-4 through 4-6. The nickel concentration for the three alloys

varies from about 50% to 65% depending on the temperature. The enrichment of the gamma

phase by Cr is most prominent in PWA 1480, followed by PWA 1480+ and PWA 1484,

respectively. Within the temperature range of interest to the present investigation (7000C to

8150C), the refractory content of the y phase (with the exception of W) of PWA 1480 and PWA

1480+ is low. The prediction indicates that the y phase would contain 2% or less of Re and Ta.

The W content is around 5% for the PWA 1480 based alloys. The Cr content of PWA 1480 and

PWA 1480+ is much larger than that of PWA 1484, while the Co content of PWA 1484 is nearly









double that of PWA 1480 and PWA 1480+. The y phase of PWA 1484 is also predicted to

contain a greater amount of rhenium.

The y' Phase

Because nickel base superalloys are precipitation hardened alloys, the nature of the

reinforcing phase is critical to the performance of the alloy. Very early during superalloy

development it became clear that y' (Ni3Al) would be a good candidate for the strengthening

phase.37, 48 Single phase y' has been shown to exhibit a large increase in strength as the

temperature is increased to 8000C. Traditional (non-superalloy) alloys exhibit a slow decrease in

strength with increasing temperature. When an alloy containing both phases is tested, the

Critical Resolved Shear Stress (CRSS) is much larger than the CRSS of either alloy individually

and is stable with increasing temperature.37 The predicted compositions of the y' phase in all

three alloys is presented in Figures 4-7 to 4-9.

Examples of the y' morphologies present in the three alloys tested here can be seen in

Figures 4-10 to 4-13. PWA 1480 exhibits a large variability in y' size and shape based on the

location within the sample. For instance, dendrite core regions contain fine, cuboidal y' and the

interdendritic regions near eutectics contain large, irregular y' precipitates (Figures 4-10 and 4-

11). Both PWA 1480 and PWA 1480+ contain retained eutectics due to the very small heat

treatment window (90C for PWA 1480 and 130C for PWA 1480+, see Table 3-5). A retained

eutectic is shown for PWA 1480+ in Figure 4-12. Additionally, segregation remains after the

solution heat treatment and this leads to differences in y' characteristics by location. To counter

these effects, as much as was practical, the solution heat treatments for PWA 1480 and PWA

1480+ ended with final hold temperatures only 50C and 30C below the solidus for each alloy,

respectively. PWA 1484 has a much larger heat treatment window (defined as the difference in









OC of the y' solvus and solidus temperatures). As a result, PWA 1484 exhibits very little retained

eutectic following casting, Figure 4-13. Solution heat treatment was able to eliminate eutectic

regions entirely for this alloy.

Primary y'

The cuboidal primary y' precipitates for all three alloys are between 0.3 Cpm and 0.5 Cpm in

edge length. In the case of the single crystal nickel superalloys PWA 1480 and PWA 1484, a

high volume fraction of primary y' is precipitated partially from cooling following solution heat

treatment and partially during the aging heat treatments. As the temperature is increased, the

equilibrium volume fraction of y' decreases as shown in Figure 4-14. Comparing the three alloys

in question it becomes clear that PWA 1480 and PWA 1480+ have a significantly greater volume

fraction of y' as predicted by JMatPro. A larger y' volume fraction can lead to differences in

mechanical behavior.49

The nature of the interaction between the y' phase and Re has been the subj ect of several

investigations spanning the past 30 years. It is generally accepted that Re is rej ected from the y'

phase during precipitation and so partitions to the y matrix. As a result, Re additions are often

implicated as the probable cause of large negative lattice misfit values in the vicinity of the y/y'

interface.8 The local enrichment of the matrix side of the y/y' interface occurs over a relatively

short distance due to the low diffusivity of rhenium. Two theories regarding the nature of the

enriched layer have been argued for the last 20 years. One states that the rejected solute atoms

form a hardened "shell" around the precipitates. A contrasting study published in 1988 suggests

that the rej ected Re forms hard clusters roughly 10 A+ in size.10. 11 Similarly, clustering has been

shown to occur with Cr forming enriched regions less than 4 nm in size. These Cr clusters are

thought to be due to the creation of the ordered structure Ni3Cr (DO22).50










The y' phase is of particular interest for the present investigation due to the correlation of

large primary creep strains with y' shearing mechanisms. The nature of the y' phase and the y/y'

interface may be related to the primary behavior described herein. Additionally, while most

investigations that pursue an understanding of the y' phase focus on the primary y' precipitates,

the secondary y' precipitates are also of significant value due to their location in the y matrix

channels. As discussed in Chapter 6, TEM analysis revealed that dislocations in PWA 1480 are

limited to the y matrix where they are likely to frequently interact with the secondary y'. PWA

1484, however, is not so constrained and so an understanding of both the primary and secondary

y' precipitates will aid in understanding the active deformation mechanisms.

Secondary y'

Metallography. Secondary y' precipitates, Figures 4-15 to 4-17, are much smaller (less

than 50 nm in diameter) and spherical. These secondary precipitates lie in the narrow y channels

and in the larger primary y' free regions near eutectics and minor phases. Figures 4-15 and 4-16

are from PWA 1480 following an interrupted creep test (LTA, 7040C/862 MPa). Figure 4-17

shows the secondary y' in PWA 1480+ also following an interrupted creep test (LTA, 7040C/862

MPa). Not much is really understood about these ultra-fine precipitates. One key study,

(Kakehi, 1999), demonstrated the ability to eliminate secondary y' precipitates by utilizing a

slow furnace cool following the final age temperature.20 It is generally agreed that secondary y'

form during the rapid cooling commonly employed in standard heat treatment practices 47, 51

The small size of the secondary precipitates coupled with the complex environment in which

they occur makes it difficult study the structure and composition of these precipitates. Viewing

the secondary y' precipitates via SEM techniques was unsuccessful with PWA 1484; however.









Local electrode atom probe (LEAP). The LEAP system is a new characterization

technique developed for high resolution compositional analysis coupled with high spatial

resolution. This new system was employed in the present investigation to study the rejection of

Re from y' during aging and solution heat treatments and it allowed for examination of

secondary y' in PWA 1484 where metallographic techniques failed due to the much better

solution and homogenization that was achieved in the alloy. In order to interpret LEAP results, a

brief discussion of its use is necessary (see Chapter 7 for a complete discussion). Because the

LEAP utilizes compositional data, two and three dimensional map representations can be created

by assigning colors to individual atoms. To make the images useful, maps are created by

applying thresholds to limit the number of atoms appearing in the image. For example, a limit of

18% aluminum would eliminate almost all Al dots in the y phase while allowing Al dots to appear

in the y' phase. In this way, the limits of the y and y' phases can be mapped.52 To examine the y'

phase, these limits were applied to the Al dots such that the y phase is transparent and the y'

phase appears in stark contrast. An example of this type of image for PWA 1484 can be seen in

Figure 4-18.

The micrographs, coupled with the LEAP data, revealed the presence of secondary y' in

both age conditions of all three alloys. This result is consistent with those reported by Kakehi

due to the use of rapid cooling from all heat treatments in the current investigation.20 Despite the

similarities in secondary y' morphology between alloys, though, significant differences in

primary creep occurred during mechanical testing. This behavior will be discussed in further

detail in Chapter 8, however, it is clear that primary creep is controlled by many mechanisms and

not primarily by the presence of secondary y' precipitates.










The y/y' Eutectic

All three alloys formed y/y' eutectics during solidification. Examples can be seen in

Figures 4-19 through 4-20. These two-phase features form during solidification due to the

complex nature of the alloys. During dendritic growth of superalloys, the first material to

solidify is enriched in the high melting point refractory elements. These first-to-solidify regions

eventually comprise the dendrite cores and have the highest solidus temperature compared to

other regions in a sample (excluding, of course, minor phases such as high melting point

carbides/TCP phases).44, 53 As solidification progresses the liquid becomes depleted in these

refractory elements and, consequently, enriched in lower melting point elements. The

solidification temperature of the remaining liquid continues to decline until the eutectic

temperature is reached. At this lowest temperature, a two-phase eutectic region is formed with

the remaining liquid. This range of solidification temperatures ahead of an advancing

solidification front leads to the creation of the so-called "mushy zone" in directional

solidification.

Because the eutectic regions solidified at the lowest temperature, these regions limit the

thermal capability of the alloy. To improve the thermal capability and reduce the effect of

eutectics, solution heat treatments have been developed to provide enough thermal energy for

significant diffusion to take place to allow the enriched regions of the dendrite cores and

eutectics to approach the original alloy composition. In order for this to occur, solution heat

treatments must exceed the y' solvus but remain below the solidus temperature. One danger,

however, is the risk of incipient melting. Increasing the temperature of the material too quickly

to a temperature above the y' solvus may cause inadvertent melting because the solidus may be

depressed to the same level as the y' solvus itself. For this reason, solution heat treatments of










these alloys often takes the form of a multi-step heat treatment that approaches the y' solvus and

the solidus slowly (the HT3 heat treatments, for example, follow this idea, Table 3-4).2, 3

The y/y' eutectics that formed during solidification were retained in PWA 1480 and PWA

1480+ following solution heat treatment. The already mentioned narrow solution heat treatment

windows for these two alloys prevented the eutectics from being eliminated entirely. Following

the HT3 solution heat treatment, however, they were significantly reduced due to the longer time

and higher temperature than the HT1 heat treatments. Figures 4-19 and 4-20 contain examples

of eutectics in PWA 1480 and PWA 1480+. As seen in Figure 4-20, eutectic regions often occur

in close proximity to carbides. Figures 4-21 and 4-22 are close-up views of eutectics in which

both eutectic phases (y and y') can be seen. Figure 4-21 shows an as-cast example and Figure 4-

22 shows a eutectic following an interrupted creep test. PWA 1484, however, exhibited no

appreciable eutectics following heat treatment due to the much larger solution heat treatment

window (460C in the HT2 condition, Table 3-5).

The Carbides

The presence of a small amount of carbon leads to the formation of carbide phases in most

single crystal nickel superalloys. The three alloys discussed herein contain between 0.02 and

0.04 wt.% C. This small addition is enough to produce a low volume fraction of carbide phases.

The carbide phases most commonly encountered in superalloys are the primary MC type and the

transition, or secondary, M6C and M23 6. Many other carbides are possible as well in addition to

the formation of boride, nitride, and carbo-nitride phases (in the presence of carbon, nitrogen,

and/or boron).38, 50 There has been a recent increase in interest in the many perceived benefits

and/or detriments of adding carbon to single crystal superalloys; however, they are beyond the

scope of this investigation. Originally added as a grain boundary strengthener, carbon was









removed from early single crystal superalloys. Recently, though, carbon has been added back

into single crystal superalloys in low amounts to lower casting defects, to capture tramp

elements, and to improve defect tolerance.

The carbides found in PWA 1480 and PWA 1484 are found in interdendritic regions and

take on script like morphologies, Figures 4-23 to 4-25. Some blockier shapes are possible as

well depending on the local conditions. Larger amounts of carbon are required to produce

dendritic carbides and so the carbides appear to be localized and not networked as shown in the

longitudinal section in Figure 4-24. Both PWA 1480 and PWA 1480+ exhibit local networks of

script like carbides throughout. PWA 1484, however, exhibits isolated clusters of small blocky

carbides as shown in Figure 4-26. This difference is likely primarily due to the lower Carbon

concentration of PWA 1484 (roughly half of the Carbon concentration of PWA 1480 and PWA

1480+). Additionally, some carbides in the heat treated PWA 1480 and PWA 1480+ samples

appear to be in the processes of dissolving as solutioning and homogenizing processes are

occurring, Figure 4-27. The carbide phases, while subject of many investigations, are not

considered to be cause for concern during primary creep. Because primary creep occurs well

before failure, the issue of premature crack initiation is not critical and so these phases are not

likely to play much of a role in primary creep.

The Topologically Close Packed (TCP) Phases

Topologically close packed phases are deleterious phases that form during high

temperature exposure of many nickel base superalloys. Topologically close packed structures

differ from geometrically close packed (GCP) structures (like fcc y and L12 y') by having planes

of close packed atoms separated by relatively large planar spacings, while GCP phases are close

packed in all directions. The large interplaner distances are caused by large diameter solute









atoms such as the refractory elements Re, W, and Mo. Several related phases fall under the

category of "Topologically Close Packed" including o, CL, P, and Laves. Due to their brittle

nature, their often needle and plate-like morphology, and their ability to rob the surrounding

material of solid solution strengtheners, TCP phases have been the focus of great concern within

the superalloy community. The drive to eliminate TCP phases from commercial alloys has lead

to the development of computer modeling and prediction methodologies like PHACOMP

(PHAse COMPutation) utilizing the number of unpaired electrons, Nv, for each element in an

alloy. There have been several attempts to modify this model with varying levels of success

(including atomic size factors, Md); however, the basic theory remains the same.

By assigning specific Nv numbers to each alloy and building assumptions about the typical

concentration of phases in an alloy, an average Nv value can be determined. It has been shown

that for many alloys, an Nv value that exceeds a critical value of 2.45-2.5 will lead to the

formation of a phase. Using the traditional Nv method of PHACOMP, PWA 1480 has an

average Nv value of 2.54 while PWA 1484 has an average Nv value of 2.50 (reported in Durrand-

Charre, 1997).5o These values would indicate that PWA 1480 is likely to produce TCP phases

during long-term, high temperature exposure, while PWA 1484 is potentially safe from TCP

formation. Adding 3 wt.% Re, however significantly increases the already high Nv ofPWA

1480, leading to early and rapid precipitation of TCP phases during high temperature exposure.

This prediction has been proven during creep testing as shown in Figures 4-28 and 4-29. The

TCP phases shown were produced during primary creep of PWA 1480+ and were prevalent in

specimens tested at 7040C and 8150C and with either age heat treatment. While TCP phases are

linked to premature failure of superalloys, it should be noted that PWA 1480+ exhibited the

longest creep lifetime and the lowest creep rate and primary creep strain as discussed in Chapter









6. These results indicate that TCP formation is not necessarily guaranteed to cause a detriment

to mechanical strength properties, however, fracture toughness and ductility may be reduced.

Changes Following Primary Creep

The most significant changes to the microstructures of the three alloys was exhibited by

PWA 1484. Interrupted creep tests of PWA 1484 were stopped following as much as 28% creep

in 15 hours or less of testing. As will be shown in Chapter 6, PWA 1484 deformed by massive

stacking fault shear of the y' phase. When these specimens were observed after testing, all the

PWA 1484 specimens exhibited elliptical cross-sections rather than the original circular cross-

sections. This behavior was consistent with work published on both PWA 1484 and another

second generation superalloy, CMSX-4.16, 17, 47 Metallographically, these specimens also

exhibited elongation of the y phase in the [110] direction (observed with a y' etch), Figures 4-30

and 4-31i. Following the use of the electrolytic etch, the primary y' precipitates appeared to be

cut along planes consistent with (111) planes, Figures 4-32 and 4-33. These figures indicate that

the y' shear processes active in PWA 1484 must be severe due to the significant changes to the

primary y' precipitates following interrupted creep testing.




























Figure 4-1. y/y' microstructure of PWA 1480 HT3 (y' etch). The y phase is the light gray phase
between the cuboidal y' phase (dark grey).


Figure 4-2. y/y' microstructure of PWA 1480+ HT3 (y' etch). The y phase is the light gray
phase between the cuboidal y' phase (dark grey).









































Figure 4-3. y/y' microstructure of PWA 1484 HT3 (y' etch). The y phase is the light gray phase
between the cuboidal y' phase (dark grey).



Gamma Phase
70

65-

60-

55-

50 _--N
-= tAl
45 -- -- Co
40 -- --- Cr -

35 _- -R
-*-Ta
S30-
-- cTi
0 25 -- -- W

20 -C --- -C

15-

10




600 700 800 900 1000 1100 1200 1300 1400
Temperature (C)

Figure 4-4. Composition of the y phase of PWA 1480 as a function of temperature as predicted
by JMatPro.
























































_ _ _ ~ __ U_


Gamma Phase


70 -

65-

60-

55-

50-

45-

40-

35-

30-

25-

20-

15-

10-

5-


600


700 800 900 1000 1100
Temperature (C)

Composition of the y phase of PWA 1480+ as
by JMatPro.


1200 1300 1400



a function of temperature as predicted


Figure 4-5.


Gamma Phase


70

65

60

55

50

45 -

40 -

35

30 -

25

20

15

10

5


600


-*-Ni
-=-Al
Co
Cr


-*-Re
STa
-W


700 800 900 1000 1100 1200 1300 1400
Temperature (C)

Composition of the y phase of PWA 1484 as a function of temperature as predicted
by JMatPro.


Figure 4-6.




































































1200 1300 1400


Temperature (C)

Figure 4-8. Composition of the y' phase of PWA 1480+ as a function of temperature as
predicted by JMatPro.


75
70
65
60 -- - -*--Ni
55~~~~ ~~ -- -- --Al -
S50 -- -- co
~~45 ----- cr
~tRe
0 40
-*- Ta
S35 i T
430 -----
0 25
20
15
10

5 -----------------;


600 700 800


Figure 4-7. Composition of the y' phase of PWA 1480 as a function of temperature as predicted
by JMatPro.


Gamma Prime Phase


----- -*-N i
-. Al
Co
cr
- - R e
------~ ~ -- Ta
- - - T i
w


S55
50
-
45
. 40
.-i 35
03 30
S25
C. 2 0


IJ
5-
0 ttt~FF-~~FFf~~4


J


15 -I


I


Gamma Prime Phase


900 1000 1100 1200 1300 1400
Temperature (C)


600 700 800 900 1000 1100












Gamma Prime Phase

75
70
65
60 --------~Ni
55 -- -- .A l
50 -- -- -- co -
45 -- cr
S40 _I Hf
: 35 -*- R
O 30
E 25
C. 2 0
15 -
10



600 700 800 900 1000 1100 1200 1300 1400

Temperature (C)

Figure 4-9. Composition of the y' phase of PWA 1484 as a function of temperature as predicted
by JMatPro.


























Figure 4-10. The y' phase in PWA 1480 near a retained eutectic (lower right) as revealed by the
electrolytic y etch.




























Figure 4-11. Irregular y' phase in PWA 1480 near a partially solutioned eutectic region (y etch).


Figure 4-12. y/y' eutectic region in PWA 1480+ during the early stages of solutioning (y' etch).
Fine primary y' precipitates can be seen in the lower right quadrant of the micrograph.
The large, hard phase in the middle is a carbide (most likely MC type).

































Figure 4-13. The y' structure of as-cast PWA 1484 (y' etch).


80
+ PWA 1480
78 -- -- --- PWA 1480+
tPA 1484




.p 72



S68





62

60
680 700 720 740 760 780 800 820 840
Temperature (C)

Figure 4-14. The y' volume fraction vs. temperature for all three alloys.




























Figure 4-15. Secondary y' in the y matrix of PWA 1480 following an interrupted creep test (y
etch).


Figure 4-16. Secondary y' in the y matrix of PWA 1480 following an interrupted creep test (y
etch).




























Figure 4-17. Secondary y' in the matrix of PWA 1480+ following an interrupted creep test (y
etch). A TCP precipitate can be seen positioned diagonally from top to bottom-right.









































Figure 4-18. Three dimensional LEAP compositional map (18wt.% Al iso-surface) The green
surfaces represent areas of Aluminum concentrations averaged at lease 18wt.%.
These regions are predominately y' due to the highly ordered of the y' phase. The y
phase is typically very low in Al concentration. The fine precipitates can be seen in
the transparent y matrix channels (near z = 100 nm). Displacement scales are in units
of nanometers (nm).





























figure 4-19. y/y' eutectics In as-cast PWA 1480 (y' etch).


Figure 4-20. y/y' eutectics in the vicinity of primary carbides in PWA 1480+ (HTIA solution
heat treatment, y' etch). Solutioning of the eutectics is already underway with the y
phase infiltrating the mostly y' eutectic (center right).





























Figure 4-21. Close-up of a eutectic in as-cast PWA 1480+ (y' etch). Small y "precipitates" can
be seen throughout the eutectic region.




















Figure 4-22. Close-up of a retained eutectic in PWA 1480+ following an interrupted creep test
(HTA, 7040C/862 MPa, y etch).





























Figure 4-23. Carbide phase In PWA 1480+ (HTI'A, y' etch).


Figure 4-24. Carbide phase in PWA 1480+ (as-cast, longitudinal section, y' etch). Carbides do
not appear to be dendritic in nature, but do wrap around dendrite arms (center-right).




























Figure 4-25. Local carbide network in PWA 1480 following an interrupted creep test (HT,
8150C/621 MPa, y' etch).


Figure 4-26. Carbide phase in PWA 1484 (y' etch). Due to the lower wt.% Carbon, PWA 1484
exhibits isolated clusters of small blocky carbides.




























Figure 4-27. A possible carbide that dissolved during solution heat treatment leaving behind y'
(PWA 1480 LT, 7040C/862 MPa, y' etch).


Figure 4-28. TCP phase formation in PWA 1480+ during interrupted creep testing (y' etch).




























Figure 4-29. TCP phase formation in PWA 1480+ during interrupted creep testing (y etch).


Figure 4-30. y phase elongation in the [110] direction in PWA 1484 following interrupted creep
testing (HT age, 7040C/862 MPa, y' etch).




























Figure 4-31. y phase elongation in the [110] direction in PWA 1484 following interrupted creep
testing (HT age, 7040C/862 MPa, y' etch).


Figure 4-32. y' phase shear along (111) planes in PWA 1484 following interrupted creep testing
(HT age, 7040C/862 MPa, y' etch).




























Figure 4-33. y' phase shear along (111) planes in PWA 1484 following interrupted creep testing
(LT age, 7040C/862 MPa, y' etch)









CHAPTER 5
RESULTS: TENSILE BEHAVIOR

Tensile tests of PWA 1480, PWA 1480+, and PWA 1484 were conducted to evaluate the

yield and tensile strengths of the three alloys. When designing creep experiments, it is important

to understand the relative differences in yield strength between the three alloys. For instance,

alloys with different yield strengths subj ected to the same creep load may experience different

creep behaviors as a result. It is possible that an alloy that is stressed to a greater fraction of its

yield stress may experience a shorter creep life because of an increase in deformation

(alternatively, a decrease in resistance against creep. Additionally, tensile testing is useful to

evaluate work hardening and ductility.

Tensile testing was conducted at two temperatures, 7000C and 8150C, set at the lowest and

highest temperatures used for creep testing. The following discussion will focus on the effect of

the two different age heat treatments on the tensile strengths of PWA 1480, PWA 1484, and

PWA 1480+. As mentioned earlier, the age heat treatment temperatures were 7040C and 8710C

and all testing was conducted at or above the LT age temperature, but below the HT age

temperature. Additionally, the effect of rhenium on the tensile properties of PWA 1480 was

investigated by the creation of a second generation version of PWA 1480 called PWA 1480+.

Because microstructural evolution is not a maj or concern during the short duration of a

tensile test, the relationship between test temperature and age temperature is minor. For

instance, if the time required for tensile test was sufficient to change the heat treated condition,

then testing at 8150C might be expected to reduce or eliminate entirely any benefit of the low

temperature (LT) age heat treatment. In this case (8150C), the thermal exposure during tensile

testing following the HT age will act as a secondary age heat treatment much like the multi-step

age heat treatments already employed in the processing of some second and third generation










superalloys. Additionally, the HT age would be expected to increase coarsening (size) of the y'

phase and, subsequently, reduce coherency of the precipitates. Rhenium additions have been

shown to significantly increase solid solution strengthening of the y phase. Because Re

preferentially segregates to the y phase, lattice misfit is significantly increased as well.

The tensile behavior of PWA 1480 is very similar for both age heat treatments. At 7000C,

both samples exhibit a large yield stress followed by slight work-hardening until failure, as

shown in Figure 5-1. The HT aged sample has a yield strength slightly lower than the LT age

and the ductility is about half of the ductility of the LT aged samples. At 8150C, both age heat

treatments have a yield point that is lower than the yield strength of the samples tested at 7000C.

Again, the HT aged samples have slightly lower yield strength, failure strength, and ductility.

The Re modified PWA 1480+ performed similarly to unmodified PWA 1480 with a few

notable differences, Figure 5-2. The general shape of the stress-strain curves remained the same,

potentially indicating no significant change in deformation mechanisms. The two most notable

changes are the resulting increase in yield strength and corresponding reduction in ductility. The

LT aged PWA 1480+ sample at 700 OC has a much lower yield strength than the PWA 1480 LT

specimen as well as significantly less ductility (11% for PWA 1480 vs. 2% for PWA 1480+). It

should be noted, however, that this result may not be valid due to its significant departure from

the trend established by the other test results for PWA 1480+. The HT aged PWA 1480+

specimen at 7000C showed a great improvement in yield strength (1333 MPa for PWA 1480 vs.

1376 MPa for PWA 1480+). The ductilities for the HT age specimens of both alloys were

similar with a difference of less than 2% elongation.

Both alloys exhibit yield points at 8150C and display very similar strain hardening

behavior after yielding. For three of the four conditions, PWA 1480+ has a greater yield stress









than conventional PWA 1480, which is likely due to the solid solution strengthening effect of the

Re addition. For both test temperatures, the HT age appears to decrease ductility in PWA 1480

(11.6% to 6.3% and 18.3% to 15.2) while simultaneously reducing the yield stress (1357 MPa to

1333 MPa and 1281 MPa to 1212 MPa, LT and HT respectively). PWA 1480+, though, had

increased ductility and yield strength in the HT age when compared to the LT age condition.

Additionally, the usually stronger PWA 1480+ does not show as much ductility as PWA 1480 in

the LT condition but does in the HT condition. Tensile strength and elongation measurements

from all three alloys are presented in Table 5-1.

The tensile behavior of PWA 1484 was somewhat different from PWA 1480 (and PWA

1480+), Figure 5-3. PWA 1484 was characterized by a significantly lower yield point, greater

ductility, and greater amount of plastic work hardening. PWA 1484 also has greater strength at

the higher temperature test condition while PWA 1480 and PWA 1480+ show a decrease in

strength at 8150C. In general, nickel base superalloys have excellent high temperature strength

to near 8000C. Continuing to increase the temperature will result in a decrease in strength in

nickel base superalloys because of the decreasing strength of the y' phase at these

temperatures.37' 48 The continuing increase in strength of PWA 1484 at 8150C is noteworthy as

the total refractory content (Mo+W+Re+Ta) for PWA 1484 is similar to that of PWA 1480+

(20wt.% vs. 19wt.%, respectively). The total refractory content of PWA 1480, however, is less

at 16wt.% with the absence of Re. Despite the similarity in solid solution strengthener content of

PWA 1484 and PWA 1480+, PWA 1484 shows an increase in strength at 8150C relative to

7000C while PWA 1480+ does not. At 7000C, the strength of PWA 1484 continues to increase

until failure. At 8150C, the strength increases to the ultimate tensile strength, UTS, following 2-

4% plastic deformation. All of the PWA 1484 tests produced between 14 and 23% elongation









compared with 5-11% for PWA 1480 and 2-10% for PWA 1480+, Table 5-1. For PWA 1480

and PWA 1484, the LT age condition is slightly stronger than the HT age condition. For PWA

1480+, however, the HT age produces greater yield strength values at both test temperatures.

Direct comparison between the three alloys also reveals the relative differences in

performance. Figures 5-4 and 5-5 show direct comparisons of the three alloys tested in the LT

age condition at 7000C and 8150C respectively. Figures 5-6 and 5-7 are similar except the HT

age is presented (rather than the LT age shown in Figures 5-4 and 5-5). Specimens with the LT

age at 7000C exhibited a small amount of work hardening after yielding as shown in Figure 5-4.

Of the three alloys, PWA 1484 displayed the greatest potential for work hardening. PWA 1480

had the largest yield strength value and substantial ductility, but little work hardening. PWA

1480+ produced little ductility (less than 2%) but experienced slight hardening after yielding.

PWA 1484 had the lowest yield point at this combination of age heat treatment and test

temperature.

At 8150C, the LT aged samples all experienced a sharp yield point. In the case of PWA

1480 and PWA 1480+, the yield point was followed by an immediate reduction in tensile stress.

PWA 1484, however, maintained the yield stress level briefly before the onset of plastic

hardening. The engineering stress is reduced 100 to 150 MPa during this period of the test

possibly as a result of the onset of necking. No further strengthening takes place before failure

occurs. Contrasting this behavior, PWA 1484 displays a significant increase in strength during

the first few percent plastic elongation. The applied engineering stress slowly decreases

following the work hardening (the transition occurs around 4-5% elongation). The decrease in

stress may be a result of slight necking, though necking was slight. While the yield strength of

PWA 1480 is 373 MPa greater than that of PWA 1484, the reduction in measured flow stress of









PWA 1480 and the increase in flow stress of PWA 1484 lead to a larger failure stress for PWA

1484 by 78 MPa. Compared to PWA 1480+, the UTS of PWA 1484 is 152 MPa below that of

PWA 1480+; however, the true failure stress of PWA 1484 (corrected by post test measurement)

is 57 MPa greater than the failure stress of PWA 1480+. PWA 1484 has the highest failure

strength despite a 445 MPa and 374 MPa disadvantage in yield strength to PWA 1480+ and

PWA 1480, respectively.

The HT age heat treated specimens share several similarities to the LT age specimens. At

7000C, PWA 1480 and PWA 1480+ exhibited the same high yield strength with a nearly flat

plastic hardening region, Figure 5-6. PWA 1480 experienced a significant reduction in ductility

with the HT age. PWA 1480+ when tested with the LT age at 7000C had the highest yield

strength of the group. The ductility for PWA 1480+ is increased with the HT age from 2.8% to

5.1%. The behavior of PWA 1484 in the HT age condition closely matches its behavior in the

LT condition. The yield points and amount of work hardening are very similar as are the

ductilities of both alloys. PWA 1480, however, experienced a decrease in yield strength, UTS,

and ductility in the HT age condition. Most notable is the decrease in ductility of PWA 1480

from 11.8% to 5.5%.

The HT age when tested at 8150C brought about similar tensile performances to the LT age

for all three alloys. Again, PWA 1480+ had the highest yield strength followed by PWA 1480

and PWA 1484. Both PWA 1480 and PWA 1480+ exhibited a sharp yield point followed by a

drop in strength of nearly 100 MPa. The measured strength continued to decline until failure.

Conversely, PWA 1484 displayed a sharp rise in strength following yielding to a maximum

(1141 MPa) at 3.96% elongation. When compared with the results from the 7000C testing, there

was little difference in properties for PWA 1480+. PWA 1480 had lower yield strength and less









ductility in the HT age condition and PWA 1484 had similar yield strengths in both conditions

but less work hardening in the HT age condition. The total ductility of PWA 1484 was also

slightly lower for the HT age condition.

The elastic modulus for all three alloys was calculated during the step loading procedure

immediately prior to the beginning of a creep test. These values are used here as the creep

loading system is more elastically rigid and the results fit a linear model with much lower error

than those generated on the servo-hydraulic tensile test system. The elastic modulus values were

obtained at 7040C and 8150C for all three alloys and at 7600C for PWA 1480 and PWA 1484.

These values are reported in Table 5-2.

Additionally, the three alloys can be differentiated by their plastic deformation behavior.

PWA 1480 displays relatively smooth plastic deformation until failure. The LT condition

transitions from elastic to plastic behavior smoothly as shown in Figure 5-8. The HT condition

behaves similarly; however, the elastic-plastic transition is marked by a brief spike in stress

shown in Figure 5-9. PWA 1484 produced a steady increase in flow stress until an instability

point is reached. After this point, the engineering stress is initially reduced and wavy until

failure. The UTS may occur at failure or at the point of instability as in the cases of the LT age

and the HT age, respectively: Figures 5-10 and 5-11. The point of instability may be caused by

necking or a similar behavior; however, only slight necking could be observed on all specimens

tested herein. PWA 1480+ exhibited a mix of behaviors from the LT age to the HT age. In the

LT aged condition, PWA 1480+ displayed a smooth elastic-plastic transition that was followed

by a rapid increase in flow stress to a potential point of instability at which the slope of the

plastic deformation region decreased sharply, Figure 5-12. Overall, the LT condition exhibited a

smooth plastic response from yielding until failure. The HT age, however, exhibited wavy









behavior with several spikes in flow stress from yielding until failure, Figure 5-13. The spikes

produced in the HT condition of PWA 1480 and PWA 1480+ are possibly a result of dynamic

strain aging processes such as the formation of solute atmospheres along dislocation cores during

the age heat treatment.

Comparing the behaviors of the three alloys between the two test temperatures, it can be

seen that two alloys, PWA 1480 and PWA 1480+, exhibit a significant change in behavior over

the temperature range in question while PWA 1484 does not. At 7000C, PWA 1480 and PWA

1480+ transition from elastic to plastic deformation with no drop in stress caused by a yield

point, Figures 5-14 and 5-16. At 8150C, however, both alloys exhibit a sharp upper yield point

followed by a constantly decreasing flow stress, Figures 5-15 and 5-17. This behavior continues

until failure for both alloys. The true failure stress of PWA 1480 is only slightly greater than the

yield stress and UTS values (both recorded at yielding). The true failure stress of PWA 1480+,

however, shows a decrease in strength of 27 MPa following yielding in the LT condition and 2

MPa following yielding in the HT age condition, Table 5-1. PWA 1484, however, does not

show a significant change in behavior from 7000C to 8150C. For both age conditions of PWA

1484, the yield point is nearly identical at both temperatures with a difference of 17 MPa at

7000C and 10 MPa at 8150C, Figures 5-18 and 5-19. The primary difference recorded for PWA

1484 is the slope of the plastic deformation region immediately following yielding. For both age

conditions, the 8150C sample exhibits a more rapid increase in flow stress due to increased work

hardening and a larger value for the ultimate tensile strength. The improvement in tensile

properties of PWA 1484 is not unexpected as it is a second generation alloy designed for higher

temperature capability than PWA 1480.54









Finally, changes with temperature in common tensile properties like yield strength, UTS,

and failure strength can also be recognized for each of the three alloys. The yield strength of the

three alloys as a function of temperature is given in Figure 5-20. PWA 1480 and PWA 1484

exhibit reductions in yield strength as a result of the increase in temperature from 7000C to

8150C. PWA 1480 showed the largest drop in yield strength (76 MPa for the LT condition and

121 MPa for the HT condition), while PWA 1484 exhibited a reduction of 41 MPa in the LT

condition and 34 MPa in the HT condition. PWA 1480+, however, exhibited an increase in yield

strength with temperature for both age conditions. The LT condition of PWA 1480+ produced

the greatest increase in yield strength with temperature (118 Mpa) while the HT age produced a 3

MPa rise in yield strength at 8150C. Additionally, comparing the measured yield strengths to the

applied initial creep loads provides a useful description of the percentage of the yield strength

required to support the applied stress, Table 5-.

The ultimate tensile strength of the three alloys as a function of test temperature is given in

Figure 5-21. PWA 1480 with both age conditions and PWA 1480+ HT exhibit a substantial

reduction in UTS of between 140 and 160 MPa at 8150C when compared with the values

produced at 7000C. As the temperature is increased further, it would be expected that the tensile

strength of these alloys would continue to decline due to the decreasing strength of the y matrix

and the y' precipitates.37, 48 An increase in the measured UTS was exhibited by PWA 1484 with

both age conditions and by PWA 1480+ LT. Again, the increase in strength of PWA 1484 at the

higher temperature can be attributed to its higher temperature performance capability produced

as a result of a greater refractory content than either of the PWA 1480 alloys. The increase in

strength for the PWA 1480+ LT sample may be an artifact as a result of the significant departure

by the sample as indicated in Figure 5-2.









When true failure stress is graphed for each alloy as a function of temperature, only two

specimens exhibited an increase in strength with increasing temperature, Figure 5-22. PWA

1480+ LT and PWA 1484 LT both produced a substantial increase in failure stress at 8150C (90

MPa for PWA 1480+ LT and 83 MPa for PWA 1484 LT). PWA 1480+ HT, PWA 1484 HT, and

both conditions of PWA 1480 all exhibited a reduction in failure stress at the higher temperature

test with PWA 1480 LT showing the greatest reduction in strength.





















-LT Ag~e


0 2 4 6 8 10 12 14 16 18 20 22
Tensile Strain (%u)


PWVA 1480


16001
1-C0
140



S800~

400'

E 200'


700 C


815 C


Figure 5-1. Tensile results for PWA 1480 at both age heat treatments and test temperatures.


PWVA 1480+


1600
1400~

S12001


8 00i
c 600'
4001
E 200'


815 C


-LT Age
---- HT Age


0 2 4 6 8 10 12 14 16 18 20 22
Tensile Strain (%)


Figure 5-2. Tensile results for PWA 1480+ at both age heat treatments and test temperatures.



























































-LT Age

---- HT Age


0 2 4 6 8 10 12 14 16 18 20 22

Tensile Strain (%)


Table 5-1. Tensile results at 7000C and 8150C for all three alloys and both ag~e heat treatments.

UTS (1VPa) Elongation at
Temperature Age o, (1VPa) 7 of (1VPa) Faiur (%)RA %


1357.44 1429.33 (1) 1604.29
1333.45 1371.15 (p) 1459.98

1281.39 1287.39 (y) 1321.31
1212.24 1213.76 (y) 1220.33

1149.77 1215.07 (f) 1250.34
1375.57 1516.20 (f) 1606.16

1267.39 1352.09 (y) 1340.15
1378.33 1385.63 (y) 1384.00

948.72 1112.69 (f) 1314.27
931.76 1116.55 (p) 1328.08

907.90 1199.62 (p) 1397.51
897.90 1141.44 (p) 1203.85


11.76
5.49

12.36
7.68

2.58
5.09

7.12
8.99

13.46
16.58


11.58 10.93
6.28 6.09

18.27 16.70
15.18 14.09

2.85 2.81
6.86 6.63

12.17 11.46
17.04 15.67

17.26 15.86
18.30 16.72


700oC



815oC



700oC


815oC



700oC


PWA 1480


PWA 1480+


PWA 1484


18.22 23.21 20.71
19.26 13.84 12.92
if the UTS occurred during plastic


815oC
HT


*Note: UTS values are designated with a "y" if the UTS occurred at a yield point, a "p"
deformation before failure, and an "f" if the UTS occurred at failure.


815 C


Figure 5-3. Tensile results for PWA 1484 at both age heat treatments and test temperatures.


PWA 1484










LT Age Tensile (700 C)
1600-
9i 1400-



S800-

.26001 ~(1480/LT/700)
S400- --- (1480+/LTP/700)
1- 200- ---- (1484/LT/700)

0 2 4 6 8 10 12 14 16 18 20 22
Tensile Strain (%)

Figure 5-4. Comparison of tensile results for all three alloys with the LT age (7040C/24hr.)
tested at 7000C with an air environment.


LT Age Tensile (815 C)
1600-
S1400-
S1200- ...
1000-
800-

~- (1480/LT/815)
400-- (40/T85
200- ---- (1484/ILT/815)


0 2 4 6 8 10 12 14 16 18 20 22
Tensile Strain (%/)

Figure 5-5. Comparison of tensile results for all three alloys with the LT age (7040C/24hr.)
tested at 8150C with an air environment.










H]T Age Tensile (700 C)
1600-

S1400-
S1200-
J;1000-i Itl""
800-
or600-
(1 480/HT, 700)
400-
--- (1480+/HT/700)
200-
---- (1484/H-T/700)

0 2 4 6 8 10 12 14 16 18 20 22
Tensile Strain (%)

Figure 5-6. Comparison of tensile results for all three alloys with the HT age (8710C/32hr.)
tested at 7000C with an air environment.


H]T Age Tensile (815 C)
1600-

S14001


-1000-i
800-

4r 00-
--(1480+/HT/815)
200-
---- (1484~/HT/815)
20-

0 2 4 6 8 10 12 14 16 18 20 22
Tensile Strain (%/)

Figure 5-7. Comparison of tensile results for all three alloys with the HT age (8710C/32hr.)
tested at 8150C with an air environment.









Table 5-2. Elastic modulus calculations from creep, loads.
Elastic Modulus (GPa)
700 C 815 C
LT HT LT HT
PWA 1480 81.48 95.75 87.44 64.30
PWA 1480+ 95.01 86.63 100.78 79.96
PWA 1484 85.15 86.39 106.44 87.41


PWA 1480/LT/700
Plastic Deformation


'J; 142







140


Figure 5-8. Plastic deformation behavior of PWA 1480 LT at 7000C.


2 3 4 5 6 7 8 9 10 11 12
Tensile Strain (%/)




























F1 s


PWA 1480/HT/700
Plastic Deformation


1375-


1 365-

S13551
k-


S1345-
E-

S13351


O 1


J


132i


2 3 4

Tensile Strain (%/)


Figure 5-9. Plastic deformation behavior of PWA 1480 HT at 7000C.


PWA 1484/LT/700
Plastic Defonnationl


h
C*
Fn
v
V1
c~
t*r
CI
e~R

81
m
a
t~l


16 7 8 9 10 1112

Tensile Strain (%/)


13 14


Figure 5-10. Plastic deformation behavior of PWA 1484 LT at 7000C.










PWA 1484/HT/700
Plastic Deformation


h

E1
v
vl
Fn
a
L1
CI
cC1

"1
E
ar
IE-l


4 5 6 7 8 9 10111~213145161718

Tensile Strain (%/)


Figure 5-11. Plastic deformation behavior of PWA 1484 HT at 7000C.


PWVA 1480+-/LT/700
Plastic Deformation


PI122
S121
~120

.~ 117
ua 116

S115
114


0.0 0.5 1.0 1.5

Tensile Strain (%/)


2.0


Figure 5-12. Plastic deformation behavior of PWA 1480+ LT at 7000C.









PWA 1480+/-iHT/700
Plastic Deformation


0 1 2 3 4
Tensile Strain (%/)


Figure 5-13. Plastic deformation behavior of PWA 1480+ HT at 7000C.


P'WA 1480 LT Tensile


1600
1400
120 0
1000
800


400
200


- (1480/LT/700)
---- (1480/LT/815)


---....


0 2 4 6 8 10 12 14 16 18 20) 22
TFensile Strain (%/)

Figure 5-14. Tensile behavior of PWA 1480 LT at both temperatures.









PWVA 1480 HFT Tensile


1600
1400
1200
1000
800


400
200


-- (1480/HT/700)
---- (1480/HT/815)


0 2 4 6 8 10 12 14 16 18 20 22
"Tensile Strain (%)

Figure 5-15. Tensile behavior of PWA 1480 HT at both temperatures.

PWA 1480+ LT Tensile


- (1480+/LT/700)
---- (1480+/LT/815)


,A.....


0 2 4 6 8 10 12 14 16 18 210 22
Tensile Strain (%/)

Figure 5-16. Tensile behavior of PWA 1480+ LT at both temperatures.










P'WA 1480+ HT~ Tensile


1480+/HT/700)
p "" ---.....~- ---- (1480+t/HT/815)
S1200 "---.
1000-
800-
cu600-
400-
E 200-

0 2 4 6 8 10 12 14 16 18 20 22
Tensile Strain (%)

Figure 5-17. Tensile behavior of PWA 1480+ HT at both temperatures.


PWJA 1484 LT Tensile

1600-
(1484/LT/700)
1400-o
a ---- (1484/LT/815)
S1200- .------.....
gl1000-
800-
S600-
S400-
200-


0 2 4 6 8 10 12 14 16 18 20 22
Tensile Strain (%/)


Figure 5-18. Tensile behavior of PWA 1484 LT at both temperatures.













-(1484/HT/700)
---- (1484/HT/815)


0 2 4 6 8 10 12 14 16 18 20 22
Te~nsile Strain (%)


Figure 5-19. Tensile behavior of PWA 1484 HT at both temperatures.


+ 1480 LT
-Ea-1480 HT
+1480+ LT
n1480+ HT~
+ 1484 LT
-e-1484 HT


1700
160
100
140
130
120
110
100



675


700 725


750 775 800 825


Temperature (C)
Figure 5-20. Yield strength as a function of temperature for all three alloys.


PWA, 1484 HT Tensile










Table 5-3. Creep, loads vs. yield strength for all three alloys.

Age OY (MPa) Creep Load (MPa) % cy
LT 1357.44 862 63.5
PWA 1480
HT 1333.45 862 64.6

u LT 1149.77 862 75.0
o PWA 1480+
0 HT 1375.57 862 62.7

LT 948.72 862 90.9
PWA 1484
HT 931.76 862 92.5


1281.39
1212.24

1267.39
1378.33

907.90
897.90


48.5
51.2

49.0
45.1


PWA 1480


PWA 1480+


PWA 1484


68.4
69.2


+l480 LTS
3- -1480 HT
+~1480+t LT:
-a- 1480+ HT
+c1484 LT
"---- ----.--... -.4 4H


170-
160

1500









675


700 725 750 775 800 825
Temperature (C)


Figure 5-21. Ultimate Tensile Strength as a function of temperature for all three alloys.













-e1480 LT
1480 HT
S1480+ LT
1480+ HT
-C1484 LT
**.. '--... -+ 1484 HT













700 750 800 850
Temperature (C)

True Failure Stress as a function of temperature for all three alloys.


650


Figure 5-22.









CHAPTER 6
RESULTS: CREEP BEHAVIOR

Creep testing was initially performed at three combinations of temperature and initial

stress as follows: 7040C/758 MPa, 7600C/690 MPa, and 8150C/621 MPa. After running several

tests at 7040C/758 MPa the initial applied stress was increased to 862 MPa due to the excessive

failure lives of PWA 1484 specimens (greater than 1700 hours). As discussed earlier in Chapter

3, creep testing consisted of two phases of testing. First, full-length testing was performed to

establish lifetime and performance expectations. The original goal of this investigation was to

study the effect of secondary y' precipitates and (related) age heat treatments. Following the first

complete round of testing, the unique primary creep behavior became obvious and the

investigation took up a new focus (on primary creep behavior). The second batch of single

crystal bars of PWA 1480 and PWA 1484 was acquired to investigate primary creep further.

Additionally, the third alloy (PWA 1480+) was created due to the prevalence of primary creep in

rhenium bearing alloys. Full length creep tests continued with the second batch of material.

The second phase of testing was performed after observing the primary creep behavior of

PWA 1484. In order to view a similar condition for all three alloys, it was decided to stop

samples following 0.5% secondary creep. This level of creep ensured that primary creep

mechanisms were still obvious while allowing for useful comparison between alloys. Stopping

primary creep specimens at a specific amount of primary creep (e.g. 0.3%) would not have been

very useful due the significant difference in primary creep behaviors expressed by the alloys.

Additionally, ending the interrupted creep tests during the maximum primary creep rate would

only allow useful examination of PWA 1484. PWA 1480 and PWA 1480+ exhibited

continuously declining creep rates beginning almost immediately following loading of the

specimens. As a result of these concerns, the deformation mechanisms were observed for










specimens just entering the secondary creep stage. Finally, another benefit of interrupting creep

early in secondary creep is that data from the entire primary creep regime are preserved for all

three alloys as will be presented below (for example, total primary creep strain, maximum

primary creep rate, minimum secondary creep rate).

Full Length Tests

PWA 1480

The first generation PWA 1480 exhibits a brief primary creep stage followed by a

continuously increasing creep rate. None of the four creep conditions produced true steady-state

secondary creep behavior in PWA 1480. This is especially obvious when viewing the creep rate

change with temperature. The minimum creep rate is usually recorded soon after the end of the

primary creep stage. The creep rate then slowly increases throughout the remainder of the creep

test. As a result, a true secondary creep stage is not observed. The initial creep rates of PWA

1480 reflect the early onset of primary creep and are typically larger than all creep rates to follow

until late in the tertiary creep stage. Immediately after primary creep the creep rates are at their

minimum values and significantly lower than the maximum primary creep rates earlier in the life

of the specimens, Table 6-1.

The complete lifetime of the PWA 1480 specimens at 7040C/862 MPa, 7600C/690 MPa,

and 8150C/621 MPa can be seen in Figures 6-1, 6-2, and 6-3. Additionally, creep test results are

given in Table 6-2 for all three alloys. A number of points can be made about the creep behavior

of PWA 1480. First and as already mentioned, tertiary creep dominates the life of the alloy at all

three test conditions. Second, as the temperature is lowered from 8150C to 7040C, and the load

increased, the performance of PWA 1480 improves significantly. Third, the amount of primary

creep is virtually unchanged for all three test conditions and both age heat treatments. And










finally, the creep behavior of PWA 1480 does not seem to have a strong dependence on age heat

treatment temperature.

While tertiary creep behavior is dominant for PWA 1480 at all test conditions, changes in

creep rates and lifetimes are clearly evident. For instance, time to 1% creep (tm9) at 7040C/862

MPa and 7600C/690 MPa are similar; however, time to 2% creep (t2%/) begins to show some

variation in performance with the higher temperature test resulting in a shorter t2%~ time. When

the temperature is raised still further to the 8150C test condition the t19 and t2%~ VaUeS are

decreased significantly. The creep rupture life of PWA 1480 declines from the lowest

temperature test condition to the highest temperature test condition. As the creep life decreases

with increasing temperature, the minimum creep rates increase with increasing temperature. The

creep rates at 8150/621 MPa are more than double the creep rates at 7040C/862 MPa. The

reduction in lifetime and creep rate performance has not affected the amount of primary creep,

however.

The primary creep strains produced by PWA 1480 at all three conditions are unchanged.

The amount of primary creep is between 0.3% and 0.5% for all test conditions even though the

maximum creep rate during primary creep is significantly increased at higher temperatures,

Table 6-1. Changing the age heat treatment temperature does not change these behaviors in a

predictable manner. In fact, no obvious dependence on age heat treatment was found. One

exception to this was found during interrupted creep testing, Figure 6-4. Here, the LT aged

specimens experienced less creep than their HT age counterparts. The amount of primary creep,

however, was the same. Overall, the performance of both heat treatments was similar enough to

be within the expected range of scatter for creep testing. Additionally, there is no change in the

amount of primary creep strains due to age heat treatment.









PWA 1480+

The addition of Re to PWA 1480 that improved the tensile properties discussed in Chapter

5 also significantly improved the creep behavior of the alloy. The already low primary creep of

PWA 1480 was reduced in PWA 1480+ and the creep rate under all heat treatment and test

conditions was significantly lower than any other alloy/heat treatment combinations in this

investigation, Figures 6-1 to 6-3 and Tables 6-1 and 6-2. Also, the rupture lifetimes of PWA

1480+ specimens were much longer than the same of PWA 1480 specimens. The overall

lifetime and creep rate performance of the experimental alloy were significantly improved;

however, creep ductility was reduced for all conditions. Most PWA 1480+ samples failed with

less than 5% creep elongation. This reduction in ductility is likely a direct consequence of the

large strengthening effect of the Re addition to PWA 1480. The alloy may have been hardened

to such a point that ductility is greatly reduced. Though not tested during this investigation, the

fracture toughness of PWA 1480+ is potentially reduced as well.

Similarly to PWA 1480, the age heat treatments did not create any obvious differences in

creep performance. The primary creep strains produced with both age heat treatments were

similar. While PWA 1480 exhibited similar primary creep strains at all test temperatures, PWA

1480+ exhibited lower primary creep strains at higher temperatures, Table 6-1. The primary

creep strains of PWA 1480+ specimens at 7040C/862 MPa are in the same range as those

produced in PWA 1480 without Re. At 8150C/621 MPa, PWA 1480+ exhibits less than half the

primary creep of the lower temperature specimens.

The creep behavior of PWA 1480+ still maintained some similarities to the behavior of

PWA 1480. For instance, the minimum creep rate slowly climbed during the test and the overall

creep behavior appeared to be dominated by tertiary creep. The rise in creep rates of PWA









1480+ were much less pronounced than the rise in minimum creep rates of PWA 1480.

Additionally, primary creep for both alloys was similar, with the exception of 8150C. This

remains true despite a significant increase in t19 and t2%/ times, indicating that the amount of

primary creep that is attained is not necessarily linked to the amount of time required for the

completion of primary creep. For example, the interrupted creep tests of PWA 1480 and PWA

1480+ were terminated around 0.75% creep elongation for TEM analysis. The PWA 1480 tests

were terminated between 10 and 25 hours, Figure 6-4. The PWA 1480+ samples, however,

finished between 250 and 450 hours for the same elongation, Figure 6-5.

The amount of primary creep produced was the same in this case at both temperatures and

for both age heat treatments. If it was only a matter of time to complete primary creep, PWA

1480+ would likely have produced much less primary creep than PWA 1480. In actuality, it

appears that the total primary creep strain is the determining factor. Another way of stating this

is to say that the onset of secondary creep terminates primary creep. This reasoning may sound

redundant; however, the onset of secondary creep is related to work hardening processes in the

alloys. Once significant work hardening is produced to slow the primary creep deformation

processes, secondary creep begins. The primary contributor to the onset of secondary creep,

then, is not time, but strain. The amount of work hardening necessary for an alloy to enter

secondary creep will vary based on microstructural characteristics and composition. As will be

shown later, substantial work hardening takes place in PWA 1480 and PWA 1480+ during the

early stages of primary creep leading to rapid work hardening and only 0.3% to 0.5% primary

creep strain. PWA 1484, however, does not exhibit this behavior and instead demonstrates very

little dislocation and stacking fault interaction during primary creep leading to reduced work

hardening. Finally, because strain is the controlling factor in the onset of secondary creep, it is









not surprising that PWA 1480 and PWA 1480+ can produce the same amount of primary creep

with large differences in the time required.

PWA 1484

The creep behavior of the second generation alloy PWA 1484, however, displays a brief

incubation period followed by a large primary creep strain. Primary creep typically ends very

soon after it starts by transitioning to secondary creep with a constant, low creep rate. Despite

the large primary creep strains, most of the rupture lives of the PWA 1484 specimens occurred

during secondary creep. Plotted as creep elongation (%) vs. run time (hr.), the secondary creep

stage of these alloys is very nearly linear. As the test nears the eventual failure life of the alloy,

the creep rate slowly begins to increase. The rising creep rate continues to increase until failure,

and thus comprises the tertiary stage of creep for PWA 1484.

Compared to PWA 1480 for the highest temperature test, PWA 1484 has a rupture life that

is roughly 10 times (1000%) the rupture life of PWA 1480, Table 6-2. As the temperature is

lowered (and the stress increases), the relative difference decreases. At 7600C, PWA 1484 has a

lifetime of about 4 times (400%) the lifetime of PWA 1480. At 7040C, however, PWA 1480 has

the greater rupture life. In the LT age condition, for example, the rupture life of PWA 1484 is

now only 1/10 (10%) of the life of PWA 1480, Figures 6-1 through 6-3.

The prominence of the primary creep stage of PWA 1484 is evident for all test conditions

used in this investigation. Generally, as the temperature is increased and the load decreased, the

magnitude of primary creep decreases, Figure 6-6. Alternatively, as the temperature is reduced

and the load is increased, the magnitude of primary creep in PWA 1484 increases. By

comparison, there does not appear to be a significant difference in primary creep as a result of

test condition in PWA 1480 and PWA 1480+, Figures 6-4 and 6-5. PWA 1484 also shows a










fairly strong primary creep dependence on age heat treatment. For all PWA 1484 creep tests, the

LT aged specimens produced significantly more primary creep than the HT aged specimens.

The effect of test temperature/stress and age heat treatment can be readily seen in Figures 6-1 to

6-3, and 6-6 to 6-7. The HT age reduced the primary creep strain of PWA 1484 by up to 10% of

the LT primary creep strain at all temperatures. Alternatively, the primary creep strains of the

HT aged specimens have been reduced to as low as 50% of the primary creep strains of the LT

aged counterparts.

Post-test measurements revealed that the PWA 1484 specimens deformed non-uniformly

during creep testing. Inhomogeneous creep has been shown to occur frequently in second

generation superalloys like CMSX-4 and PWA 1484.13, 47, 55 All of the PWA 1484 specimens

were elongated from the original circular cross-section to an elliptical cross-section. The major

and minor axes of the ellipse were parallel to <1 10> directions. It has been reported that the

elliptical cross-section is formed mostly during primary creep if few slip systems are active.

After the initiation of secondary creep, work-hardening should result in greater uniformity of

deformation due to the activation of multiple slip systems.

Following this reasoning, the difference in creep rate between the two age heat treatments

of PWA 1484 at 8150C/621 MPa, Figure 6-3, may be due to changes in cross-sectional area

during primary creep. With a smaller cross-section the LT age sample will deform faster than

the HT age, giving the impression of a higher creep rate. Contrasting this reasoning are the data

produced at 7600C/690 MPa, Figure 6-2. Here, an even greater difference in primary creep

strain yielded nearly identical secondary creep rates between the two age heat treatments. More

testing is needed before the cause of the varied secondary creep rate can be stated with

confi dence.









Interrupted Tests

Interrupted creep results are presented in Figures 6-4 to 6-6 for all three alloys and both

age heat treatments. From these results it is immediately clear that PWA 1484 exhibits unusually

large primary creep strains. Additionally, the amount of time that elapses during primary creep

changes significantly for each alloy. In order of increasing time to complete primary creep:

PWA 1484 (5-15 hours), PWA 1480 (12-22 hours), and PWA 1480+ (230-410 hours). This

order also applies to decreasing amounts of primary creep produced during interrupted testing.

Comparing the age heat treatments for the three alloys reveals a few correlations of

interest. First, the primary creep behavior of PWA 1480 may have slight age heat treatment

dependence; however, as already mentioned the overall behavior of PWA 1480 does not appear

to follow a dependence on age heat treatment temperature. Second, the primary creep of PWA

1484 has a strong dependence on age heat treatment temperature, Figure 6-6. At both

temperatures, the HT age specimens resulted in the lowest primary creep which is consistent

with full length creep testing. By comparison, at 7040C the HT age reduced the primary creep of

PWA 1484 by 32% and at 8150C reduced the primary creep by 67%. Finally, PWA 1480+

shows no correlation to age heat treatment, which is also consistent with full length testing.

The test temperature also brought about some behavioral changes. Both PWA 1480 and

PWA 1480+ exhibited less time to reach 0.7% creep at 8150C. The minimum and maximum

creep rates were slightly higher at 8150C than at 7040C for both alloys, Table 6-1. PWA 1484

exhibited a much more dramatic change in properties with increasing temperature. While PWA

1480 and PWA 1480+ produced approximately the same amount of primary creep at both

temperatures, the primary creep of PWA 1484 was greatly reduced at 8150C for both age heat

treatments. For the LT age, the primary creep strain of PWA 1484 at 7040C was 24.44% and









3.68% at 8150C. For the HT age, the results were 16.66% at 7040C and 1.20% at 8150C. For

the LT age, this represents an 85% reduction in primary creep. For the HT age it is a 93%

reduction.

Transmission Electron Microscopy (TEM)

Transmission electron microscopy was performed to produce a qualitative understanding

of the active deformation mechanisms because the two reported deformation mechanisms are

distinct in appearance. First, the more common deformation mechanism reported in superalloys

consists of matrix dislocations bowing to fill the y matrix channels and creating interfacial

dislocation networks.19, 27, 56, 57 This mechanism applies to most superalloys due to the growing

incoherency of the y/y' interface during a creep test. Deformation usually begins in the y' matrix

and this leads to the formation of interfacial networks to one degree or another in all nickel base

superalloys. The second deformation mechanism commonly associated with primary creep

occurs by stacking fault shear of the y' precipitates and has been reported in several alloys.16, 17,

23, 24, 33, 58 The alloys most commonly linked to this behavior are second generation and later

alloys like PWA 1484 and CMSX-4; however, one first generation alloy has been reported to

exhibit this method of creep deformation as well.33 Following the interrupted creep tests it

would be expected that all of the specimens would exhibit sufficient formation of y/y' interfacial

dislocation networks to cause the onset of secondary creep, which was found to be true for the

specimens examined here.16, 19

PWA 1480

The deformation of PWA 1480 during interrupted creep consisted primarily of matrix

dislocations bowing to fill the y matrix channels, Figures 6-11 and 6-12. Most of the

deformation appears to be limited to the matrix phase; however, a low amount of stacking fault









formation was also found within the specimens, Figures 6-13 to 6-15. The stacking faults that

were found tended to be limited to a range of one or two y' precipitates only. It will be shown

that the limited nature of stacking fault formation stands in contrast to the prevalent nature of the

stacking faults in PWA 1484. The limiting of dislocations to the y matrix during primary creep

is also consistent with low amounts of primary creep as reported by several sources.16, 17, 19, 20, 27,

33, 56, 57 Additionally, the presence of secondary y' precipitates was confirmed by TEM, Figure 6-

16.

PWA 1480+

The second generation version ofPWA 1480, PWA 1480+Re, exhibited a significant

change in behavior (relative to PWA 1480). Most notably was an increase in the number of

stacking faults within the material, Figure 6-17. Additionally, stacking faults were often seen

interacting with stacking faults lying on different slip planes, Figures 6-17 to 6-19. While

stacking fault formation is tied to large primary creep strains, they are typically limited to a

single slip system in alloys exhibiting large primary creep strains.16, 17, 33 PWA 1480+, however,

was shown to exhibit low amounts of primary creep. Additionally, the fact that the stacking

faults appear to be interacting with other stacking faults, indicating slip on more than one slip

system, is worth exploring. These interactions are common in superalloys displaying stacking

fault shear, but not until the final stages of creep leading to failure.33 These interactions in PWA

1480+ occurred early in the life of the specimens during primary creep and failure is not

imminent. These interactions, coupled with the active matrix dislocations, may be generating

relatively large amounts of strain hardening that brings about the onset of secondary creep earlier

than alloys, such as PWA 1484, that exhibit stacking fault shear on a single slip system during

primary creep. Additionally, stacking faults in PWA 1480+ are limited to regions spanning only









2 to 3 y' precipitates (or fewer). The localized nature of the stacking fault shear may be a result

of these stacking fault interactions restricting the propagation of the shear bands as they form. In

this way, the ability to produce strain is reduced and the expected primary creep strains would be

lower. Despite the formation of larger numbers of stacking faults, deformation is limited and the

secondary creep stage occurs early (in terms of total strain, not time).

PWA 1484

The PWA 1484 specimens, when observed on the TEM, appear to share the same

deformation mechanisms widely reported among second generation alloys at low temperatures

and high loads as used in this investigation. Wide-spread stacking fault shear that appears to act

on a single slip system was apparent over large regions within the specimens, Figures 6-20 to 6-

21. This highly planar deformation mechanism, as reported by othersl6, 17, Occurs by the passage

of two a/2<1 10> matrix dislocations (same slip plane with burger' s vectors at 600 to each other)

into the y' precipitates. These dislocations then dissociate into a/3<112> and a/6<112> partial

dislocations with stacking faults in between. The two pairs of partial (with stacking faults in

between) are separated by an anti phase boundary (APB) due to the ordered nature of the y'

precipitates. It is this complex system of dislocations, stacking faults, and the associated anti

phase boundary that is able to shear large distances with relatively little impedance, generating

large primary creep strains." The numerous rows of stacking fault ribbons in PWA 1484 can be

seen in Figures 6-20 and 6-21. Additionally, a large density of matrix dislocations has formed in

PWA 1484 following primary creep, Figure 6-22. Again, this is consistent with the onset of

secondary creep. These shear bands are also conservative in nature and leave no dislocation

debris at the y/y' interfaces (assuming no other dislocations are interacting with the stacking fault

"ribbons"). As a result, the ribbons are free to expand and glide until they interact with other










matrix dislocations. These matrix dislocations will eventually form interfacial networks;

however, the amount of primary creep that has been produced is already quite large by this point.

Therefore, by the time enough work hardening is present to force the onset of secondary creep in

PWA 1484, the amount of primary creep that has been conferred can be quite large. The matrix

dislocations shown in Figure 6-22 are the likely cause of the beginning of the secondary creep

stage.










Table 6-1. Primary creep, and creep, rates from interrupted creep, tests.
704oC/862 MPa 815oC/621 MPa
Primary Max. Primary Minimum Primary Max. Primary Minimum
Age Creep Creep Rate Creep Rate Creep Creep Rate Creep Rate
Alloy HT (%) (%/hr.) (%/hr.) (%) (%/hr.) (%/hr.)


PWA
1480


0.35
0.49


0.1632
0.4015

0.0073
0.0296

10.704
4.3877


0.0116
0.0166

0.0007
0.0011

1.8697
0.7010


0.34
0.32

0.22
0.11

3.68
1.20


0.6001
1.0377

0.0339
0.0304

2.4563
0.6751


0.0277
0.0418

0.0021
0.0022

0.0768
0.0454


PWA LT
1480+ HT


0.36
0.34

24.44
16.66


PWA
1484


Creep (704 C/862 MPa)


PWVA 1484




PWA 1480


PWA 1480+


0 300 600 900 1200 1500 1800
Run Time (hr)


Figure 6-1. Creep at 7040C/862 MPa of all three alloys.





















- LT Age
---- HIT Age


Creep (76i0 C/690 MPa)


PWA 14180


1600


0 400 800 1200
Run Time (hr)
Figure 6-2. Creep at 7600C/690 MPa of all three alloys.

Creep (815 C/621 FMPa)


PWNA1480


- LT Age
---- H~T Affe


0 400 800 1200
Run Time (hr)
Figure 6-3. Creep at 8150C/621 MPa of all three alloys.


1600










Table 6-2. Rupture lives and total creep elongation from full-length creep tests. Also included is
time to 1% and time to 2% creep. Note: The PWA 1480+ HT sample at 8150C/621
MPa failed before 2% creep, was achieved.
Alloy Age tm9 (hr.) t2% h. tutr (hr.) Elongation (%)
LT 55.2 199 842 16.77
1480
HT 52.0 135 424 14.52

ud 180+ LT 63.0 427 1223 5.52
ON HT 619 1225 1721 3.32

LT 7.19 8.33 90.3 38.81
1484
HT 10.1 11.7 23.7 19.64

LT 56.6 117 311 10.78
HT 47.4 124 428 13.90
io
r- o\LT 3.27 4.10 1465 20.08
S1484
HT 5.96 13.9 1525 14.69

LT 11.2 28.8 99.3 15.01
1480
HT 14.4 26.8 77.2 14.93

uv, 140+ LT 486 797 1239 8.94
HT 372 N/A 505 1.45

LT 1.15 1.18 391 13.40
1484
HT 4.92 49.8 914 19.06









Interrupted Creep: PW~A. 1480


I~Ic704 C





-LT ~ge
---- H~T Alge


0 5 10 15 20 25
Run Time (hr)

Figure 6-4. Primary creep of PWA 1480.


Interrupted Creep: PWA 1480+


LT


704 C


' LT


- LT Age
---- HT Age


0 100 200 300 400
Run Timae (hr)


500


Figure 6-5. Primary creep of PWA 1480+.
















704 C


815 C


-LT Age
---- H-T Age


0 5 1
Run Time (hr)


Figure 6-6. Primary creep of PWA 1484.


Creep (704 C/862 MPa)
nn, PWA 1484 Closeup


40AII8 L~rIIO


0 10 20 30 40 50 60 70 SO 90 100
Run Time (hr)




PWA 1480+


-'3


a


1


PWA 1484


PWA 1480


m ,


0 300 600 900 1200 1500 1800
Run TFimae (hr)


Figure 6-7. Creep at 7040C/862 MPa of PWA 1484 compared to PWA 1480 and PWA 1480+.
Inset: magnified view of the creep behavior of PWA 1484.


Interrupted Creep: ]PWA. 1484

















a
Ei 201

E
o
W 101
P,

U


P WA 1434


PWA 1480


Figure 6-8. Primary creep comparison at 7040C/862 MPa for all three alloys. Note: large strain
range was used better viewing of the behavior of PWA 1484.


PErimtary Creep (704 C/862: MPa)


2.0











~-0.0
U


PWA 1484






; PWA 1480


PWIA 1480+


0 100 200 300 400
RIun "Time (hr)


Figure 6-9. Primary creep comparison at 7040C/862 MPa for all three alloys. Note: long time
range was used for better viewing of the behavior of PWA 1480 and PWA 1480+.


Primary Creep (704 C/862 MPa)


10 20
Run ~Time (hr)









Primary Creep (815 C/621 MPa)


1484


PWA 1480


PWA 148(H-


0 10


]Run Timte (hr)


Figure 6-10. Primary creep comparison at 8150C/621 MPa for all three alloys.


























































Figure 6-11. Bright field(a)/Dark field(b) pair of deformation in PWA 1480 (HT age, 7040C).
Creep deformation is primarily limited to the y matrix and interfacial dislocation
networks have already formed.


132


























































Figure 6-12. Bright field(a)/Dark field(b) pair of deformation in PWA 1480 (HT age, 7040C)
revealing the formation of few stacking faults.




133


































Figure 6-13. Bright field TEM image of dislocation networks in PWA 1480 following primary
creep.


Figure 6-14. Stacking fault and dislocation shear of PWA 1480 is limited to small regions within
the specimens (bright field).





































r figure 6-13. A staclcing tault In F wA 14rru (Garke tela).


Figure 6-16. Secondary y' precipitates (marked by arrows) in PWA 1480 following interrupted
creep.

































Figure 6-17. Stacking fault interactions following primary creep in PWA 1480+.


Figure 6-18. Stacking fault interactions and a dislocation network in PWA 1480+.


































r figure 6-19. Ynort range stacialng tault snear or rwH 14ru+.


Figure 6-20. Bright field TEM image of PWA 1484 LT (7040C) following interrupted creep.
Inhomogeneous deformation by stacking fault shear of the y' precipitates is apparent.



























Figure 6-21. Stacking fault shear ofy' precipitates inPWA 1484 (bright field).



























































Figure 6-22. Bright field(a)/Dark field(b) pair showing the interfacial dislocation networks
present in PWA 1484. Note: stacking fault shear is also present, but out of contrast
(marked by arrows).


139









CHAPTER 7
RESULTS: ADDITIONAL CHARACTERIZATION

In addition to common metallographic characterization, mechanical testing, and

transmission electron microscopy, two additional approaches were used to characterize the three

alloys in this investigation. The first is a recently developed method called the Local Electrode

Atom Probe (LEAP). This technique improves upon ideas established with the Scanning Atom

Probe (SAP) and the 3 Dimensional Atom Probe (3DAP) that were developed in the early

1990's. The LEAP allowed for high resolution compositional characterization in three

dimensions and was used to observe the secondary y' in PWA 1484 as well as the segregation

behavior of Re to the y matrix. The results presented in the LEAP section of this chapter were

conducted at the University of North Texas by Anantha Puthicode and Mike Kaufman in

conjunction with the University of Florida.

The second characterization method included in this chapter is X-ray diffraction (XRD).

X-ray diffraction was used to study the lattice misfit of all three alloys and to observe how misfit

changes with heat treatment. Additionally, the data collected by XRD needed to be processed to

separate the contributions of the y and y' phases that otherwise overlap. This extra step involved

deconvolution of the intensity data produced by the (002) and (004) planes within the alloys.

Both of these techniques were used in an attempt to gain an understanding of the fundamental

differences between the alloys as a result of chemistry and processing.

Local Electrode Atom Probe (LEAP)

The Local Electrode Atom Probe is capable of near atomic resolution in both three

dimensional space and mass number. As a result, these instruments are capable of rendering

three dimensional representations of the distribution of atoms within a specimen. Due to the

relatively recent development of the LEAP, a brief discussion of its operation will be included









below. A full review of the development and capability of the LEAP can be found in (Kelly and

Larson, 2000).52 At it's most basic, the LEAP functions by extracting ions (atoms) from a

specimen and accelerating the ions towards a detector. With the system highly calibrated, the

time of flight of an ion is used to determine the mass and the location of the ion on the detector

determines the location in two dimensions of the point of origin for the ion at the tip of the

specimen. The third dimension (along the axis of the specimen) is controlled by careful

adjustment of the specimen and rate of extraction. A two dimensional side-view schematic of

the device is given in Figure 7-1.

As shown in the figure, the specimen is positioned beneath the extraction electrode. The

extraction electrode resembles a hollow cone with the tip removed to create an opening at the tip.

It is through this opening that the extracted ions are accelerated towards the detector. The

secondary electrode is a disk (with a concentric hole in the center) of larger diameter than the

base of the extraction electrode. The presence of the secondary electrode is required because the

LEAP requires a relatively low extraction potential, thus the ions must be accelerated towards

the detector. Additionally, the extraction electrode can be pulsed with a high frequency to

control the rate of extraction events.52

Specimens for the LEAP take the form of tall narrow cones resembling spikes or needles.

These spikes need to be very small in diameter with a tip radius less than 10 to 50 nm. Smaller

tip radii aid in focusing the applied field at the tip, improving the ability to extract ions from the

specimen. The specimen size utilized for this investigation consisted of a diameter of less than

50 nm and a length (of analyzed volume) of at least 150 nm. The complete method of sample

preparation is given in Chapter 3. Creating the specimens, however, involved two stages of

electropolishing in order to achieve the dimensions required for the LEAP system, Figure 7-2.









The first step was a "bulk" electrothinning operation that resulted in tip radii between 250 nm

and 500 nm. The second step used a small platinum loop to refine the tip to a radius below 50

nm. All LEAP specimens in this investigation are oriented parallel to the [001] direction (also

parallel to the applied stress axis). Additionally, the compositions of the individual LEAP

specimens may vary slightly based on location. Due to incomplete homogenization, a specimen

near a dendrite will be more enriched in Re, Mo, and W and depleted of Al and Ta, while a

specimen near the interdendritic region will express the opposite. PWA 1484, however, exhibits

relatively little segregation and a large solution heat treatment window so these effects should be

small.

Four sets of LEAP specimens were created from a single bar of PWA 1484. The four

sections of PWA 1484 received heat treatments to induce changes in the microstructure for

observation of secondary y' and the segregation behavior of rhenium. Two sections received the

LT age and two received the HT age. One section from each of the two age heat treated groups

was then subj ected to another brief solution heat treatment. This extra solution heat treatment

was conducted with the goal of quickly dissolving the y' precipitates but not allowing enough

time for the enriched Re layer around the prior y' interface boundaries time to homogenize. The

samples were then quenched rapidly. The goal was to create a condition that would allow the Re

shells (or clusters) to be examined by the LEAP method. The results of the LEAP analysis are

presented below and separated by heat treatment. It should be noted that a successful analysis of

PWA 1484 LT (age heat treatment only) could not be accomplished due to losses during

specimen preparation.









PWA 1484 LT

Reconstruction (solution HT)

The first of the LEAP reconstructions is given in Figure 7-3. Figure 7-3a is an 18%

aluminum iso-surface and Figure 7-3b is an 18% chromium iso-surface from PWA 1484 LT with

the extra solution heat treatment following aging. An iso-surface is a concept used to develop an

understanding of the distribution of specific elements within the specimen. In the case of Figure

7-3a, any region with a concentration of at least 18% Al will create a surface that encompasses

the region. In this case, the areas that are colored in green are most likely y' precipitates, while

the areas with no color are most likely the y matrix. The y' precipitate with it' s thickness fully

contained within the LEAP specimen is potentially a secondary y' precipitate based on the

measured thickness of 75 nm. Primary precipitates, by comparison, in PWA 1484 measure

between 300 and 500 nm in thickness. The small, Al-rich regions in Figure 7-3 in the y matix

between the larger secondary y' precipitates measure less than 5 nm in diameter. These locations

may mark clustering of Al atoms leading to the development of secondary y' precipitates.

Without further characterization, no specific conclusion can be made as to the structure of these

regions. Figure 7-3b is an 18% Cr iso-surface. Because Cr partitions more to the y matrix, the

the y phase becomes apparent. It should be noted, however, that the y' phase still contains a

significant Cr concentration and, as a result, blue surfaces appear in the y' phase as well.

Composition profile (solution HT)

The image in Figure 7-4 is an SEM image of the specimen tip used to create the

reconstructionn in Figure 7-3. The line along the specimen axis is the direction and length of the

analyzed volume for both the reconstructionn in Figure 7-3 and for the composition profile in

Figure 7-5. The composition profile in Figure 7-5 is very similar in appearance to a composition









line scan that can be produced on an Electron Probe Micro Analysis (EPMA) microscope except

with greater accuracy. In actuality, the profile produced by the leap is created from a cylinder of

user defined length, diameter, and orientation within the collected data. This cylinder can be

oriented parallel to the axis of the specimen or perpendicular to the specimen or at any other

angle that is deemed useful. In the case of the profile in Figure 7-5, the cylinder of analyzed

volume was oriented within the center of the specimen (concentric) and parallel to the long axis.

From the compositional profile, it can be seen that fine y' precipitates are present between

the distances of 20 and 95 nm and 125 and 170 nm. Studying the profile reveals several distinct

partitioning behaviors among the elements added to PWA 1484. First, aluminum and tantalum

partition strongly to the y' phase. As a result, the composition of these elements is much greater

in the y' phase than in the y matrix. Additionally, it can be seen that Al partitions so strongly to

the y' phase that a significant depletion of Al exists for the y matrix in the vicinity of the y'

precipitate. The second behavior of interest results from those elements that partition to the y

matrix. These elements are rhenium, molybdenum, chromium, and cobalt. Rhenium and

chromium appear to partition the most the y phase with a slight enrichment layer on the y side of

the y/y' interface. Molybdenum only slightly partitions to the y phase as it expresses significant

solubility in both the y and y' phases. The third behavior is exhibited by tungsten. Tungsten

does not partition particularly strong to either phase and, as a result, maintains relatively uniform

composition throughout the analyzed volume. This behavior is interesting as previous research

has indicated that W additions enable increased Re solubility in the y' phase.35 If W does not

partition strongly to the y phase like the other solid solution strengtheners, then a significant

portion of the added W is present in the y' phase. This allows for the possibility that W additions










increase the average lattice parameter in the y' resulting in a synergistic effect between W and

Re.

PWA 1484 HT

Reconstruction (age HT)

The next reconstruction is presented in Figure 7-6 for a PWA 1484 HT specimen in the age

heat treated condition. This reconstruction is an example with all of the recorded ions present.

Each ion species is recorded as a different color dot placed within the volume at its place of

origin. Figure 7-7 is another iso-surface construction to illustrate the partitioning behavior of Cr

(Figure 7-7a) and Al (Figure 7-7b). From these figures it is apparent that the lower phase is most

likely y' due to the high Al concentration. The upper phase is most likely y. No fine precipitates

were observed for this specimen, however, it should be noted that only 30 nm in the length of the

specimen were analyzed due to excessive specimen thickness.

Composition profile (age HT)

The composition profile given in Figure 7-8 was produced from the analyzeld volume

presented in the reconstruction in Figure 7-6. This profile is similar in nature to the one in

Figure 7-5 except for a higher scale on the composition axis to allow for the Ni profile. Based on

the composition profiles it is apparent that the y phase is present over the distances 0 to 20 nm

and the y' phase is present over the distances 20 to 30 nm. All of the partitioning behaviors

discussed above are present with the addition Ni partitioning to the y' phase (due to the Ni3Al

formula). Unfortunately, few conclusions can be drawn from this specimen due to the small

volume that was analyzed. Additionally, significant enrichment and/or depletion were not

observed with this specimen.









Reconstruction (solution HT)

The final reconstruction was produced with a PWA 1484 HT specimen exposed to the

extra solution heat treatment following aging. The two images in Figure 7-9 are two different

views of the same reconstruction. Figure 7-9 is an 18% Al iso-surface, again coloring regions of

y' green. The y' precipitates all appear to be secondary y' with thicknesses of 20 nm or less

(except for the relatively large precipitate in the bottom corner). Additionally, as with Figure 7-

3, there are many small aluminum rich regions of less than 5 nm in thickness. Again, it is

difficult to conclude if these are ultra-fine y' precipitates or just Al rich clusters that are

precursors to y' precipitate formation. Another representation of the same reconstructionn is

provided in Figure 7-10. Here only the Al, Ta, Cr, and Mo ions are displayed. Because Al and

Ta partition to the y' phase and Cr and Mo partition to the y phase, the contrast between the two

phases is still evident. From both Figure 7-3 and Figure 7-9, it is apparent that PWA 1484 is

capable of producing secondary y' precipitates. It has already been shown that PWA 1480 and

PWA 1480+ are capable as well via SEM techniques, Chapter 4.

Composition profile (solution HT)

Finally, Figure 7-11 is a composition profile generated from the reconstruction in Figure 7-

9. Figure 7-12 illustrates the location and orientation of the cylinder of material selected to

produce the composition profile. This region spans two secondary y' precipitates between the

distances 8 and 15 nm and 36 and 47 nm. The measured thickness of these two precipitates is

then 7 and 11 nm, respectively. Slight local enrichment by Re and depletion of Al in the y matrix

in the vicinity of the 7 nm precipitate can be seen on the profile. Additionally, the partitioning

behaviors already discussed were also present for the PWA 1484 HT (solution HT) specimen as

expected.










Secondary y' Concentrations

One benefit of the LEAP system is the ability to determine the concentrations of the

secondary y' precipitates. These concentrations were determined from the "line" scan data

generated with cylinders of data selected from the reconstructions. Of the most interest to the

current investigation are the differences in secondary y' concentration as a function of precipitate

size (thickness). These precipitate composition data were collected from the composition

profiles in Figures 7-5, 7-8, and 7-11. Contained within these profiles are the entire thicknesses

of 3 y' precipitates and partial thicknesses of 2 y' precipitates. The compositions, listed in order

of thickness, are given in Table 7-1. The aluminum concentration is similar in all five

precipitates (roughly 18 wt%). The elements chromium, molybdenum, and tungsten were also

maintained at the same composition regardless of precipitate size (approximately 2wt% for Cr,

3wt% for Mo, and 3wt% for W). The elements nickel, cobalt, tantalum, and rhenium, however,

all varied with precipitate size. As the y' precipitates increased in size, the concentrations of Ni

and Ta increased while the concentrations of Co and Re decreased. The relative changes in

composition for these four elements are given in Table 7-2. Compared to the smallest y'

precipitate, the largest precipitate (in this set of measurements) exhibited an increase in Ni

concentration of 10wt%, an increase in Ta concentration of nearly 5wt%, a decrease in Co

concentration of 7wt%, and a decrease in Re concentration of nearly 2wt%. These results appear

to be consistent with normal precipitate coarsening behavior. As the precipitate grows, the

concentration of the y' former would be expected to grow (Ni and Ta) while the elements that

are rej ected from the y' phase would be expected to decrease in concentration (Co and Re). The

elements Cr, Mo, and W, however, do not display this behavior as they maintain uniform

composition with precipitate size.









X-Ray Diffraction (XRD)

The X-ray diffraction study was initiated to determine the lattice misfit of the three alloys

in this investigation: PWA 1480, PWA 1480+, and PWA 1484. Sections were taken from the

single crystal bars of each alloy and heat treated to 4 hr., 10 hr., 100 hr., and 1000 hr. at 10800C.

This temperature was selected as a continuation of the coating heat treatment cycle that these

alloys were subj ected to prior to service which is consistent with experimental procedures

utilizing over-aging heat treatments between 9500C and 11000C used in other sources.27' 59-62

Additionally, this temperature was selected, rather than the LT or HT age temperature, because

the higher temperature allows faster coarsening and a better approximation of the equilibrium

structure after 1000 hours.

Following the over-aging heat treatments, the specimens were thinned and polished to

produce the best results. Upon recording the intensity data, peak deconvolution was employed to

differentiate the contributions of the y and y' phases in agreement with other sources working

with nickel base superalloys.46, 59 Figure 7-13 illustrates the deconvolution process. Two peaks

are assigned to both phases creating a total of 4 peaks. The two peaks for each phase represent

the Cu kml and Cu ku2 wavelengths of X-ray radiation to which the specimens were exposed.

Due to the slight difference in wavelength between the two, the diffracted peak produced by each

is slightly displaced. This effect necessitates the separation of both wavelengths for each phase

accounted for in a given peak. The MDI Jade (ver. 7) software is programmed to account for the

difference in Cu kml and Cu ku2 radiation and is calibrated to the diffractometer. Creating four

peaks from the original is a mathematic operation and takes place utilizing an iterative algorithm.

Upon completion of a deconvolution routine, a report is generated detailing critical information

about the peaks and the fit of the model. These reports have been included in full in Appendix









B. Repeated deconvolution of the same data set does not guarantee the same result every time.

Depending on the selected starting conditions and the choices made among seven peak fit

options, these results can vary. As a result, it is critical that each report is observed to verify if

the model is useful or if it has strayed too far from the expected result. An example of this is

Figure B-57 in which the contribution of the y phase has been almost entirely eliminated due to

the ability of the y' phase to fit the original intensity data without any further contribution by the

y phase. In this case, the deconvolution was repeated producing Figure B-59. Due to these

concerns, repeated deconvolution was necessary to generated accurate predictions that could be

averaged to produce a valid lattice parameter measurement.

The results of the deconvolution procedures were then processed using a spreadsheet to

calculate the lattice parameters and lattice misfit values from each of the alloys. These

calculations are also presented in Appendix B. The results presented here are the most reliable

data from the deconvolution process. The lattice misfit of the three alloys was found for the 4

hr., 10 hr., and 1000 hr. specimens. The 100 hour specimens produced low peak intensities and

were not useful for deconvolution and subsequent misfit calculation. The calculated lattice

parameters for the (002) plane are provided in Table 7-3. These results are also plotted in Figure

7-14 along with the limited results from the (004) plane. From observation of Figure 7-14, it is

apparent that PWA 1484 has the largest, negative misfit value of the three alloys. This is

expected as PWA 1484 contains more solid solution strengtheners that partition strongly to the y

matrix, including Re which is known to partition to the y phase particularly potently. The lattice

misfit of PWA 1480, however, varies on the slightly negative side of zero. The magnitude of the

misfit of PWA 1480 is low (less than 0. 15%). The addition of Re to PWA 1480 resulted in a









continually decreasing (from positive to negative, but increasing in magnitude) lattice misfit

from an initially positive value of +0.034% to a final value of -0. 114%.

The results from the (004) plane were more difficult to obtain due to the lower intensity

produced through diffraction. As a result, fewer specimens yielded strong enough intensity

measurements over the 26 range 1170 to 1200. For this reason, fewer results are reported for the

(004) plane in Figure 7-14. Additionally, the greater spacing achieved at this high angle proved

more difficult for the deconvolution process. Increased difficulty in data acquisition coupled

with decreased analysis accuracy resulted in fewer reportable results. The results shown in

Figure 7-14 consistently indicated more positive values for lattice misfit for all three alloys.

PWA 1480, for instance, is calculated to have a positive misfit of roughly the same magnitude

that was calculated from the (002) data. PWA 1480+ is also shown to have positive misfit of the

same range in magnitude. Finally, the (004) result for PWA 1484 after 1000 hours is over 0.1%

more positive representing greater than 50% difference in calculated misfit values from the same

specimen. The result of these inconsistencies is an apparent need to continue examining the

lattice misfit of these alloys. Alternative methods (TEM based) also likely need to be employed

to verify the results from the XRD deconvolution process.















Q
~
~

s
~
t


Specimen


Figure 7-1. Schematic illustrating the basic function of the LEAP system. The specimen is
positioned beneath the cone of the extractor anode and ions are extracted using a
pulsing voltage. The ions are then accelerated towards the detector by a secondary
el ectrode.52


M Detectlor




/; Extracted Ions from Samplc



Secondary Electrode









































Figure 7-2. LEAP specimens before, (a), and after, (b), the final polishing step with the Imago
Electropointer.











V ~jr b "5


Figure 7-3. Iso surfaces created with the LEAP system (PWA 1484 LT with solution HT). (a)
18% Aluminum surface, (b) 18% Chromium surface.



























Figure 7-4. Magnified (SEM) view of the tip analyzed in Figure 7-3.

25 ,


O



o

C.)


0 20 40 60 80 100 120 140 160


Figure 7-5. Composition profile from the specimen in Figure 7-3 (PWA 1484 LT with solution
HT). Fine y' precipitates are present at distances between 20 and 95 nm and between
125 and 170 nm (Corresponds with Figure 7-3a).


Distance




































Figure 7-6. The distribution of all recorded ions for PWA 1484 HT with no additional solution
heat treatment.


I
~--~--


III b


E~~B)~-
r
I


Figure 7-7. Iso surfaces created with the LEAP system (PWA 1484 HT no solution HT). (a)
18% Chromium surface, (b) 18% Aluminum surface. These results are from the same
specimen shown in Figure 7-6.











-- Ni % -Co % -Cr % Al %


Mo % Ta % -W % Re %


70


60

O
50

O
0.. 40

O
S30

20


10



0 5 10 15 20 25 30

Distance (nm)
Figure 7-8. Composition profile from the specimen in Figure 7-7 (PWA 1484 HT no solution
HT). A primary y' precipitate is present at distances between 20 nm and 30 nm
(corresponds with Figure 7-7a)




























H I 1 ''r ;J- i, r
~d~.~J~f: ~Stii~t J
4i)C.
rrl~Le S
-*

r: c~;~Y*l;~S, ~k,
d

Iruit
~ k"~-
-a~


4' 1: I





Figure 7-9. Iso surface (18% Aluminum) created with the LEAP system (PWA 1484 HT with
solution HT). Secondary y' precipitates are apparent in the y matrix.





4
! dl


as~ ~


nl


OIe
~ i


Figure 7-10. LEAP results with only Al, Ta, Cr, and Mo ions represented (PWA 1484 HT with
solution HT). The results shown here are from the same specimen reported in Figure
7-9.















O

On

o


0 5 10 15 20 25 30 35 40 45 50


Distance (nm)
Figure 7-11. Composition profile from the specimen in Figures 7-9 and 7-10 (PWA 1484 HT
with solution HT). An illustration of the volume of material used for this
composition profile is given in Figure 7-12.













Figure 7-12. Illustration of the data selected for the compositional profile shown in Figure 7-11.
The cylinder is user selected and only the data contained within this region is counted
in the profile.










Table 7-1. Composition (wt%) of secondary y' precipitates of varying diameter in PWA 1484.
Diameter: 6nm
Ni Co Cr Al Mo Ta W Re
53 17 3 18 3 <1 3 2
Diameter: 11nm
Ni Co Cr Al Mo Ta W Re
54 16 2 19 4 1 3 1
Diameter: 74nm
Ni Co Cr Al Mo Ta W Re
58 14 2 18 3 2 3 <1
Diameter: 46+nm
Ni Co Cr Al Mo Ta W Re
62 9 2 18 3 3 3 <1
Diameter: 10+nm
Ni Co Cr Al Mo Ta W Re
63 10 2 18 2 5 2 <1

Table 7-2. Net changes in secondary y' concentration with increasing precipitate size.
Particle size (nm): 6 11 74 46+ 10+ Net Change (wt%) % change
aNickel 53 54 58 62 63 +10 +19

O-2 Cobalt 17 16 14 9 10 -7 -41
~~Tantalum <1 1 2 3 5 +5 +1000
U Rhenium 2 1 <1 <1 <1 -2 -75
































49 50 51 52
26 (degrees)
Figure 7-13. An example of the deconvolution process used to separate the contributions of the y
and y' phases. Notice, both phases have two peaks associated with them from the Cu
kml and Cu ku2 wavelengths.

Table 7-3. Lattice misfit (%) for the (002) plane following heat treatments at 10800C.
Heat treatment time at 1080 C
4 hr. 10 hr. 1000 hr.
PWA 1480 -0.075 +0.002 -0.138
PWA 1480+ +0.034 -0.052 -0.114
PWA 1484 -0.228 -0.251 -0.201


Y' peak


Y peak












0.30
-*-911-002
o- 0911-004
0.20 --- --- -- -- 922-002
E- 0922-004
1212-002

0- .10 -i 212-004



S0.00



j- -0.10



-0.20



-0.30
1 2 3

Heat Treatment

Figure 7-14. Lattice misfit vs. heat treatment from measurements of both the (002) and (004)
peaks. Note: The heat treatments for each point are given by numbers (1: 4 hr. at
10800C, 2: 10 hr. at 10800C, and 3: 1000 hr. at 10800C). The 100 hr. at 10800C heat
treatment is not shown due to the unreliable nature of the data produced from samples
given the 100 hour heat treatment. Also the sample names are as follows: 0911 is
PWA 1480, 0922 is PWA 1484, and 1212 is PWA 1480+.









CHAPTER 8
DISCUSSION

Through the course of the investigation into the primary creep behavior of PWA 1480 and

PWA 1484 several relationships between processing and creep behavior were observed. The

composition and processing an alloy has received plays a vital role in determining the nature of

the creep behavior that should be expected. Numerous studies have been performed to

incorporate the many inherent differences in microstructure into a material based modeling

scheme. The following list includes several material conditions that are impacted by chemistry

and/or processing often included in creep models:

* Stacking fault energy (both y and y' phases)13, 63, 64
* Anti-phase boundary energy (y' phase only)13, 63, 64
* y/y' morphology
0 Precipitate dimensions and distributionS21, 23, 63
o The absence/presence of secondary y' precipitateS16, 20
0 y' volume fraction63, 65
o y channel thickneSS66, 67
o y' composition (relates to the strength of the y' phase)11, 68, 69
o y composition (relates to the strength of the y phase)10, 11, 68, 69
o Microstructural stability (y' coarsening, rafting)38, 70-73
* Chemistry effects
o Segregation (eg. Re segregation to the y matrix~io, 11
0 Short range order (eg. presence of DO22 Ordered clusters of W and/or Cr)69
o Formation of secondary phases carbidess and TCP phases)35, 38
* Active deformation mechanism
0 Dislocation shear of the matrix21, 74
o Stacking fault shear of the y' phasel2, 13, 16, 17, 23, 64
o Mixed-mode deformation63

Each of the above points can impact the behavior of a single crystal superalloy during

creep. Consequently, it is necessary to investigate any relationships that may exist between these

material attributes in order to simplify any models that are derived. The current investigation

represents an attempt to understand the effect of secondary y' precipitates and the element

rhenium on the primary creep behavior of first and second generation superalloys PWA 1480 and









PWA 1484, respectively. A third alloy, named PWA 1480+ for this investigation, was also

produced by adding 3wt% Re to PWA 1480 to make a second generation version of PWA 1480.

The following discussion will focus on the different behaviors among the three alloys produced

by two different aging heat treatment schemes and alloy chemistry. Special attention will be

given to several of the material attributes listed above and their effect on the creep behavior of all

three alloys at low temperatures (7000C-8150C) and high stresses (0>500 MPa) where primary

creep dominates the creep behavior of PWA 1484. Finally, methods of applying the knowledge

gained from this investigation towards creep behavior modeling and future alloy development

will be discussed.

Microstructure

The three alloys used for this investigation all share several microstructural similarities.

First, the typical cuboidal y/y' microstructure for high y' volume fraction superalloys extends to

all three alloys. Second, all three contain a low volume fraction of carbide phases due to the

addition of between 0.02wt% and 0.04wt% C. Additionally, secondary y' precipitates were

found in all three alloys with both age heat treatments. All of these similarities were impacted by

alloy chemistry and processing.

y/y' Morphology

The cuboidal microstructure common among modern single crystal nickel base superalloys

is produced due to a combination of the high volume fraction of y' and a negative lattice misfit.40

In the case of PWA 1480, the lattice misfit is small and only slightly negative at room

temperature. The addition of Re to create PWA 1480+ did not significantly reduce the lattice

misfit as expected. For both PWA 1480 and PWA 1480+, the measured values of lattice misfit

at room temperature would be expected to become more negative as the temperature is raised









due to thermal expansion effects. The low misfit values have led to slightly rounded y'

precipitates and slightly larger y channels (than PWA 1484). A benefit of low lattice misfit, and

rounded y' precipitates, is increased stability against coarsening and coherency loss. As the

misfit is increased, the internal stress near the y/y' interfaces is increased. Under an applied

stress, dislocations will be attracted to the interfaces to alleviate the high stresses. Additionally,

diffusion-based processes will occur to reduce the misfit stresses still further. The second

generation PWA 1484, however, has a larger misfit resulting in sharper edges and corners on the

cuboidal y' precipitates.

Rhenium additions have also been tied to changes in the y/y' morphology.8 '" In particular,

the addition of Re has been linked to smaller y' precipitate size and an increase in rounded edges

(to the point of spherical shapes in some lower volume fraction superalloys).8 Within the present

investigation, the effect of Re on the y/y' microstructure is small. Following the HT3 solution

heat treatment, there is little difference between PWA 1480 and PWA 1480+ in size and shape of

the y' phase. The effect of Re, as reported in the literature, is caused due to the low diffusivity of

the element in both y and y'. During cooling from the solution heat treatment, the super-

saturated nature of the y phase brings about nucleation of fine primary y' precipitates. The low

diffusivity of Re prevents these precipitates from growing significantly during cooling. Upon

subsequent aging heat treatments, the growth of the y' precipitates is slower due, again, to the

rej section of the slow diffusing Re from the precipitates into the matrix. Additionally, the

rounded corners of y' precipitates were caused by Re due to reduced growth kinetics. These

results were produced in alloys with lower y' volume fractions than are present in PWA 1480

and PWA 1484.8 '5









Precipitate coarsening is also driven by the Gibbs-Thompson effect by which the reduction

in interfacial energy drives the growth of larger precipitates at the expense of smaller

precipitates.76 The Oswald coarsening of precipitates in a matrix has been shown to be impacted

by a number of factors including composition, coherency, precipitate size, radii of curvature near

corners (morphology), and applied stresses.76 In the case of single crystal superalloys, both

primary and secondary y' precipitates are subject to the Gibbs-Thompson effect. This effect can

usually be seen in the y matrix near the y/y' interface where the population of secondary y'

precipitates is reduced. The growth of the much larger primary y' precipitate depletes the matrix

near the y/y' of y' former and causes the elimination of the ultra-fine secondary y' precipitates in

this region. These growth processes occur during isothermal aging heat treatments in order to

obtain optimal mechanical properties shown to be maximized with a primary y' size between

0.30 Cpm and 0.45 Cpm for high volume fraction nickel-base superalloys.36 Modelling the

compositions of precipitates with varied sizes and shapes can be performed based on this effect.

Another application of the Gibbs-Thompson effect can be found with the application of an

external stress at elevated temperature on the primary y' precipitates. The Gibbs-Thompson

equation can be modified to account for applied stress on coherent, or partially coherent,

precipitates in a matrix. The application of stress can create directional growth of normally

spherical or cuboidal precipitates do to the superposition of interfacial stresses that vary by

location around the precipitates. As these stresses are relaxed, the microstructure can become

directionally oriented. This has been shown experimentally as well as theoretically and has been

called y' rafting or topological inversion.70 71, 76

For the alloys in this investigation, these effects are limited by the high equilibrium volume

fraction of y'. The predicted equilibrium volume fractions for PWA 1480, PWA 1480+, and









PWA 1484 are shown in Figure 4-14. The super-saturated y phase, upon cooling, nucleated a

Eine dispersion of coherent y' precipitates. Because the equilibrium volume fraction of y' was so

high, it is expected that the this thermodynamic driving force caused growth of the precipitates,

possibly with slightly rounded corners, until the narrow y channels became depleted with y'
former elements. As two y' precipitates grew wit;h parallel {0C1}\ facesp approahing each other,


the depletion of the y channel adj acent to the centers of the y' precipitate surfaces in y' former

slowed the local growth of y' precipitates. The y channel regions near what would become the

corners of the y' precipitates would have contained a greater amount of y' former and precipitate

growth would have continued until these regions also became depleted. The result would be

cuboidal y' precipitates with fairly well defined corners and edges even if Re additions would

dictate slightly rounded corners in lower volume fraction alloys. Comparing the y/y'

microstructure of PWA 1480 and PWA 1480+ revealed nearly identical y' shapes and sizes

despite the addition of Re. The y channel widths also shared the same size and shape within both

alloys.

PWA 1484, however, exhibited sharper y' corners and slightly narrower y channels. This

effect can be described by comparing the alloy chemistry with that of PWA 1480. Aside from

the addition of Re which would be expected to increase misfit and create rounded y' cubes, an

increase in Co and reduction in Ta were made in the newer alloy." The presence of both of

these elements has been shown to reduce the lattice misfit which would lead to further rounding

of the y' precipitate edges.8, 32, 75 COuntering these effects, though, is the increased diffusion

coefficient produced by increasing Co and reducing Ta. Reported results have shown that

increased cobalt concentrations can reduce or even eliminate any adverse effects caused by the

addition of Re." The increase in diffusion caused by these alloy modifications are possibly









responsible for the slight improvement in y' shape and the reduction in y channel thickness in

PWA 1484 when compared to PWA 1480.

Carbides

Carbides were found to be present for all three alloys. PWA 1480 and PWA 1480+ both

contained about 0.04wt% C and PWA 1484 contained about 0.02wt% C. These small C

additions were sufficient to bring about precipitation of carbide phases in the form of localized

carbide networks. The level of carbon in the three alloys was insufficient to cause dendritic

carbide formations.78, 79 All three alloys contained script carbide networks with a few small

blocky carbides associated with the local networks. Carbon additions have been shown to reduce

casting defect formation and reduce y/y' eutectic formation.79 Because the carbide content for all

three alloys was low, the carbides did not impact the primary creep process significantly.

Research has shown that carbide interfaces are often the site of void and crack formation leading

to failure.39 It was not found that the carbides present in PWA 1480, PWA 1480+, and PWA

1484 played a significant role during primary creep.

Topologically Close Packed Phases

The addition of rhenium to PWA 1480+, while improving creep properties, brought about

precipitation of plate-like TCP phases produced during high temperature exposure consistent

with the formation of a phase.35, 38 The TCP phases present in PWA 1480+ grew throughout the

heat treated microstructure with the length parallel to <1 10> directions. The presence of these

phases during primary creep indicates that the alloy is very unstable with regard to TCP phase

formation. A direct consequence of this is that PWA 4180+ is not a suitable alloy for

commercial use. These phases are expected to be the cause of the reduced ductility exhibited

during creep and tensile testing. It is also expected that the fatigue life of PWA 1480+ would be









reduced as a result of the formation of the TCP precipitates. As a consequence, the applicability

of PWA 1480+ is primarily limited to experimental testing and investigation of the rhenium

effect. The formation of TCP phases during primary creep occurred over a much longer time

period than the time for completion of primary creep in PWA 1480 and PWA 1484 due to the

reduction of creep rate caused by the addition of rhenium. Increased rhenium content has been

tied to increased TCP phase content as a result of the increase in overall refractory element

content.35, 38, 50 While TCP phases have received a great amount of attention due to the acicular

morphology that is often associated with their presence, PWA 1480+ demonstrated the greatest

creep lifetime at every test condition. The PWA 1480+ specimens failed with low ductility

which might be tied to the formation of cracks and voids associated with TCP precipitates.

Interesting future work could pursue the link between TCP precipitation and the early failure of

PWA 1480+. Additionally, slight modifications in alloy chemistry might be made that preserve

the excellent creep behavior of PWA 1480+ while reducing the TCP content to increase lifetime

and/or ductility. While all three alloys share several similarities in microstructure, differences in

performance are apparent as a result of the impact the microstructures of the three alloys in this

investigation had on the active deformation mechanisms during tensile and creep testing.

Tensile Behavior

The tensile strength properties of the three alloys of interest are necessary for a complete

understanding of their respective creep behavior. The strain controlled nature of a tensile test

reveals a different material response than a load controlled creep test. For example, superalloys

most often exhibit <110> type deformation in the y phase.so, sl Depending on a variety of

factors, including lattice misfit, dislocations will either shear or bypass the y' precipitates during

tensile testing. Shear of the y' phase often occurs with pairs of a/2<1 10> dislocations. Similarly









to creep testing, the a/2<110> super-dislocation pair can further dissociate into partial

dislocations separated by a complex fault.so Deformation during creep, however, has been

shown to occur through a variety of methods: dislocation shear of the y matrix by <1 10> and

<112> dislocations, y' shear by <112> super-dislocation pairs and stacking fault pairs, and

<010> cube slip.12, 13, 16, 17, 24, 56, 63, 74, 82, 83 Both test methods bring about complex deformation

mechanisms. Tensile testing, therefore, is complementary to creep testing and can aid in

generating a more complete understanding of the uniaxial properties of these alloys.

As with creep deformation, tensile deformation primarily begins within the y matrix

through the generation of a/2<1 10 {)1 11}\ dil locations. Initially, the micrmostrctur e is relativerly

free of dislocations. Upon yielding, these dislocations begin moving through the microstructure

as more dislocations are created by the available sources. During the early stages of plastic

deformation, dislocations begin bowing between y' precipitates, lining the y/y' interfaces and

filling the channels rather than entering and shearing the relatively hard y' precipitates. It would

be expected that this process will be affected by the presence (or absence) of secondary y'

precipitates within the narrow y channels. The ultra-fine secondary y' may play a role in tensile

deformation by a number of possible methods.

Secondary y'

First, because secondary y' precipitates are very small in size (spherical and 10-50 nm in

diameter), they would be highly coherent, increasing the probability of y' shear. In this case, the

<112> dislocations would be likely to shear the y' without dissociating into partial due to the

small size of the precipitates. Anti phase boundaries would be created and so additional

dislocations would be drawn into these precipitates to relieve this energy. While shear of the

secondary y' precipitate would be more difficult for the first dislocation, the second dislocation









would experience softening while eliminating APB.20 The second potential interaction between

secondary y' precipitates and matrix dislocations is due to the possibility of dislocation bypass

mechanisms. At high temperatures, dislocation climb would be expected to play a vital role in

bypass operations. Within the lower temperature range of 7000C to 8150C, however, climb is

kinetically slow. As a result, cross-slip is more likely to account for precipitate bypass by

dislocations. Both mechanisms would require added stress to accomplish.

The composition of the secondary y' precipitates is also useful while considering either

method of secondary y'-dislocation interaction. The composition of the secondary y' precipitates

(determined by LEAP analysis, Tables 7-1 and 7-2) revealed that the precipitates contained

lowered levels of Ta than is present in the base alloy (1-5 wt% vs. 9 wt% nominal composition).

Additionally, in the smaller secondary y' the Mo content was nearly double the nominal alloy

composition (3-4 wt%/ vs. 2 wt%), the Co content was higher (14-17 wt% vs. 10 wt%), the W

content was reduced to half (3 wt%/ vs. 6 wt%/), and the Cr content was lower (2 wt% vs. 5 wt%).

With reduced y' strengthener content, namely Ta, and lower solid solution strengthener content

these precipitates may be lower in strength than the larger primary y' precipitates. As a result,

the strength benefit due to the presence of secondary y' would be expected to be relatively small

and interactions would be frequent due to the dense, fine dispersion within the y matrix.

The serrations that appeared in the plastic deformation regime during tensile testing of

PWA 1480, PWA 1480+, PWA 1484 are likely due to the presence of secondary y'. Of the

possible interactions described above, y' shear is the most likely due to the coherent nature (and

potentially lower strength) of the secondary y'. Similar interactions would be expected during

creep with some exceptions. The differences that are observed during creep are due to









differences in deformation mechanism and will be discussed further in the Creep Mechanism

discussion.

y Channel Thickness

Another microstructural feature that impacts tensile properties is the thickness of the y

channels. This feature is primarily important for the early stages of tensile deformation because

as precipitate shear becomes more common in the later stages, the channel widths become less

important as dislocation bowing becomes less preferred. The importance of the y channel width

can be seen in Equation 8-1 that describes the stress increment to cause bowing of a straight

dislocation between two precipitates:

Gb
Equation 8-1: B Acr a
(A 2r)

Where the stress increment due to bowing is Ao, G is the elastic modulus, b is the burgers vector

for the dislocation, h is the center-to-center particle spacing, and r is the particle radius

(assuming the particles are the same size). In this case, it can be seen that the strengthening

effect is increased if the inter-particle distance is reduced.84 Strengthening due to dislocation

bowing is likely to play a role in the tensile deformation of PWA 1480, PWA 1480+ and

PWAl484; however, there are several interactions that occur that complicate the problem. For

example, the already mentioned secondary y' would serve to drastically reduce the inter-particle

spacing (if no shear was assumed and bypass was not active).

It is likely that a combination of these effects takes place during tensile deformation. For

example, when a dislocation is stretched across a y channel ("pinned" by the primary y'

precipitates), the first strength increment to be accounted for should be the bowing stress from

the primary precipitates. As the stress is increased, the dislocation will come against many fine

secondary y' precipitates. The effective inter-particle spacing becomes significantly reduced and










the bowing strength effect is increased. As the applied stress continues to increase, either the

stress to shear the secondary y' or the stress to cross-slip and bypass the secondary y' is reached

and the dislocation begins to glide. These interactions continue until the stress to shear the

primary y' precipitates is reached and large scale shear becomes dominant leading to eventual

failure .

Lattice Misfit

Additional considerations are related to the effect of alloy composition and, more

specifically, the equilibrium concentration of the y' phase. Three aspects related to composition

and heat treatment are lattice misfit, stacking fault energy, and anti-phase boundary energy. The

magnitude of the lattice misfit can impact the resolved shear stresses within the y channels. High

misfit alloys experience large misfit stresses that can create compressive stresses in the y

channels parallel to the applied stress direction. These misfit stresses can be quite large. During

high temperature, low stress creep testing, for instance, these compressive stresses can exceed

the applied external stress creating a very large driving force for microstructural change and,

subsequently, drives rafting. During tensile testing, these compressive stresses can relieve the

applied stress from a dislocation and reduce the actual shear stress felt by the dislocation creating

a strengthening effect. In order to relieve misfit stresses, dislocations can be attracted to the y/y'

interfaces to create dislocation networks. The formation of these networks increases the stress in

the y matrix leading to y' shear.65, 80

Stacking Fault Energy

Shear of the y' phase usually involves the formation of stacking faults as matrix

dislocations enter the y'. Stacking faults are regions of atoms shifted from normal lattice sites by

the passage of a partial dislocation. Stacking faults are contained between two partial









dislocations that can be combined, with enough stress, to form a perfect dislocation. The

stacking fault energy, SFE, is composition dependent and is related to the energy required to

create the fault. The larger the SFE the smaller the spacing between partial dislocations. With

large SFE alloys recombination and cross-slip is easier, but with low SFE alloys, the partial

dislocations can spread out further and recombination becomes more difficult and interactions

with other stacking faults and/or dislocations becomes more likely. The SFE of the y' phase can

also impede deformation of the y' phase. If the SFE of the primary y' precipitates is high, more

resolved shear stress will be needed to force a dislocation to enter the precipitate. Shear of the y'

precipitates is slightly easier with lower stacking fault energy alloys.65, 80

Anti-Phase Boundary Energy

Related to SFE is the anti-phase boundary energy. While stacking faults displace atoms a

partial atomic spacing, anti-phase boundaries are created by displacing atoms in an ordered

precipitate by a whole atomic spacing. The result is a change in nearest neighbor species type.

For the highly ordered y' phase, an APB will position Al atoms next to Al atoms and Ni atoms

next to Ni atoms across the boundary. The result is increased energy and resistance to

deformation. To relieve the energy produced by the creation of an APB, a second dislocation

must shear along the same plane to shift the atoms into "proper" atomic positions. Like the SFE

discussed above, the magnitude of the APB relates to the energy required to create it. Alloys

with large APB energies will have short dislocation spacing and vise-versa for low APB energy

alloys. There is a strength increment required to produce an anti-phase boundary that prevents

dislocations from entering the y' phase. Once a single dislocation has entered (either in whole

form or dissociated into partial with a stacking fault), though, a second dislocation will be

attracted into the y' phase to relieve the APB. The consequence of this behavior is the









observation of stacking fault pairs that shear the y' phase. As a result of each of these

strengthening effects, the tensile deformation of single crystal superalloys can become

complex.65, 80

Tensile Results

The results of the tensile testing conducted on PWA 1480, PWA 1480+, and PWA 1484

revealed several differences between the alloys. The first generation superalloy PWA 1480

provides a basis of comparison for both the Re modified PWA 1480+ and the second generation

PWA 1484. Both second generation alloys are descendant from the original PWA 1480 alloy.

As discussed earlier, the transition from PWA 1480 to PWA 1484 during the development cycle

was brought about by several key changes in composition including Cr, Co, Ta, Ti, and Re. The

creation of PWA 1480+ was performed by adding 3 wt% Re (the same amount added to PWA

1484) to PWA 1480 to create a second generation version of the alloy.

The second generation superalloy, PWA 1484, has demonstrated significant improvements

in high temperature creep capability compared to PWA 1480.77 A consequence of the changes in

composition that led to improved creep strength was a decrease in tensile strength for PWA 1484

relative to the original PWA 1480. The yield strength of PWA 1484 is over 400 MPa lower at

7000C and over 300 MPa lower at 8150C. PWA 1480 exhibited high yield strength, UTS, and

failure strength values for both aging heat treatments and test temperatures used in this study,

Table 5-1. PWA 1484 exhibited significantly lower yield strength, as mentioned earlier, but

demonstrated a greater ability to work harden following yielding. In fact, at 8150C the failure

strength of PWA 1484 LT exceeds the failure strength of PWA 1480 LT. Additionally, the

strength of the HT version of both alloys at 8150C is only separated by 16 MPa at failure despite









a 214 MPa disadvantage in yield strength. Clearly the alloy design approach taken for PWA

1484 significantly changed the mechanical behavior of the alloy relative to PWA 1480.

While the behavior of PWA 1484 is very different from the original PWA 1480 alloy, the

performance ofPWA 1480+ was quite similar to PWA 1480. The modified alloy is marked by a

greater yield strength (in the HT condition) and less ductility. This response is not unexpected

due to the addition of rhenium. Rhenium has long been known to be a "potent" solid solution

strengthener as well as a retardant of y' coarsening.n7 Simply adding Re raised the yield strength

of PWA 1480+ HT by 42 MPa at 7000C andl62 MPa at 8150C. The LT condition, however,

exhibited lower yield strength than HT condition at both test temperatures.

Another effect of the age heat treatments that were given to the three alloys can be seen in

PWA 1480 with which the LT age produced greater than 10% ductility during tensile testing.

The HT age, however, significantly reduced the ductility, yield strength, and UTS values. The

HT age does not, however, produce this same effect in PWA 1484. The HT aged specimens of

PWA 1484 exhibited a more rapid increase in flow stress during work hardening, but the

ductility and UTS values were similar to the LT aged specimens. Additionally, it is noteworthy

that while the tensile behavior of PWA 1480 was significantly altered by age heat treatment and

PWA 1484 was not, creep testing, as discussed below, produced creep behavior of PWA 1484

that was significantly altered by age heat treatment while PWA 1480 was not.

Creep Behavior

Primary creep at low temperatures and high stresses has become an important topic as

models of creep behavior are being developed for modern superalloys. The early superalloys

required relatively simple to design creep models. These first generation superalloys, such as

PWA 1480, exhibit tertiary creep behavior across a broad range of temperatures and stresses.









Modern superalloys, beginning with the second generation superalloys, however, exhibit a range

of creep behaviors depending on the stress and temperature regime to which they are subj ected.

For CMSX-4, for example, the three creep regimes that are displayed are primary creep

(T<8500C, 0>500 MPa), tertiary creep (7000C
(T>9500C, 0<200 MPa).63 Each of these creep regimes are related to the microstructure and,

consequently, the processing that the alloy was subj ected to prior to service. Additionally the

behaviors at these three temperature/stress regimes have been shown to be highly dependent on

the chemistry of the alloy.

Tertiary Creep

The most common creep regime that is encountered in testing of single crystal nickel-base

superalloys is tertiary creep. For example, the alloy PWA 1480, a first generation superalloy,

exhibits increasing creep rate as creep strain accumulates (tertiary creep), over a wide range of

temperatures and stresses including those used for this investigation. Tertiary creep behavior is

has been linked to a continuously increasing dislocation density implying a creep softening

process.63 Often, the tertiary creep rate can be modeled with the relatively simple equation given

below:

Equation 8-2: E = Eo (1 + Cs)

Where E is the creep strain rate, E, is the initial creep rate, C is a fitting constant and E is creep

strain.43 The disPlcations responsible for tertiaory crepp are primarily a/2<110> { 111) tyrpe


dislocations that are mostly contained within the y matrix. These dislocations, under the applied

stress, are forced to bend between the primary y' precipitates and stretch such that they line both

sides of the y "channels". As dislocations from different sources begin to meet along the y/y'









interfaces, nodal networks are formed. As the dislocation density continues to rise, the spacing

between dislocations at the interfacial networks is reduced.

For alloys that display tertiary creep in the absence of rafting, the y/y' microstructure is

relatively stable for the duration of the creep test. This is largely due to the fact that shearing of

primary y' precipitates is low and migration of vertical y channels to horizontal y channels is

slow. Failure during creep of single crystal superalloys typically involves the condensation and

coalescing of vacancies to create voids. Often, cracks and voids are nucleated in conjunction

with casting porosity. In alloys containing hard, brittle carbides and/or TCP phases, voids

typically form at the interface between the hard particle and the more ductile matrix. For cases

where these hard interfaces are not present (or present in low amounts), it is thought that void

formation may also occur through the development of dense dislocation cells near y/y' interfaces

where dislocation annihilation can lead to a local increase in vacancy concentration.63 Changing

the temperature and/or stress such that another creep regime becomes active reduces the

accuracy of the basic creep rate model shown above. As rafting or primary creep become

dominant, more complex models are required.

Rafting

Rafting in single crystal nickel base superalloys has become commonplace at low stresses

and high temperatures. Rafting, also known as topological inversion, is a process by which the

original cuboidal y' precipitate morphology evolves into long plate-like "rafts". The final

orientation of the rafts depends on the lattice misfit between the y and y' phases. Most often this

misfit value is negative, leading to y' rafts oriented perpendicular to the applied stress. The

rafting response occurs through diffusion of matter from thinner rafts through the y channels to

the thicker rafts. The result is a widening of the perpendicular y channels. This morphological










shape change results in the elimination of y channels parallel to the applied stress direction. The

<110> { 11) matrix dislocations begin gathering along the y/y' interfaces creating dense

dislocation networks.

Because matter is primarily diffusing across the y channels during rafting, the y channel

widening can be described by a parabolic rate law:

Equation 8-3: Aw1 = c, J

Where Aw is the change in thickness, C1 is a material and temperature dependent constant (for

isothermal creep), and t is time. Notice, equation 8-2 describes widening of the y channels using

a t1 2 time dependence. Coarsening of the y' precipitates, or Ostwald ripening, typically follows a

t1 3 time dependence.32 As matter diffuses from the y channels that are parallel to the applied

stress into y channels perpendicular to the applied stress, matter from the cube surfaces of the y'

precipitates parallel to the applied stress is simultaneously diffusing into the parallel y channels,

effectively closing the y channels oriented parallel to the applied stress. The change in y channel

width due to this morphology change effect has been described by the following:

(2was~ +w oso)
Equation 8-4: Aw, = 2
'IC 2(wo +so)2

Where wo is the initial y channel thickness and so is the initial y' edge length. The morphology

effect is used to calculate the amount of change in channel width as a result, simply, of the

formation of y' rafts (ie. the change in shape necessary to create rafts). Finally, the parabolic rate

constant from equation 8-3 can be found by:

Equation 8-5: Aw = w(t) wo Aw,,









Plotting Aw versus 1~allows for determination of the rate constant (slope of the line). Modeling

the y channel widening kinetics of superalloys experiencing rafting is necessary for the

development of constitutive equations to predict the creep behavior at these conditions.66' 67

Primary Creep

While numerous studies have focused on developing accurate descriptions of tertiary creep

behavior and y/y' rafting, the origin of primary creep has relatively little attention until recently.

Large primary creep strains produced at low temperatures and high loads are typically found in

second generation and later single crystal superalloys. While creep deformation in the tertiary

creep regime is governed primarily by <110> {111) dislocatio n sea~nrn in the y matrix


deformation during primary creep under these conditions is produced by pairs of stacking faults

that cooperatively shear both the y and y' phases. Specifically, stacking faults in the y' are

formed by reactions of dislocations of the type a/2<1 10> to form a/3<1 12> and a/6<1 13>

dislocations. An example reaction is given below:

Equatin 8-6:a/2[10 l]+a/2[011] 4a/3[11 2]+ a/ 6[112]"

Between the a/3<112> and s/6<112> dislocations lies a stacking fault that exist in both the

y and y' phases. As the pair of dislocations (with the stacking fault between) enter an ordered y'

precipitate an anti-phase boundary is created in the y' phase, but not in the y phase. To reduce

the energy required to create the APB, a second pair of <1 12> dislocations enters the y' in the

opposite configuration. As they progress, the APB is eliminated. An example of this

configuration is given below:

Equation 8-7: a/3<112>+SISF+a/6<1 12>+APB+a/6<1 12>+SESF+a/3<112>









Where SISF and SESF are intrinsic and extrinsic stacking faults, respectively, and APB is anti-

phase boundary. It is this configuration that is referred to as a dislocation or stacking fault

"ribbon."14

These dislocation ribbons, two stacking faults separated by an APB cooperatively shearing

both phases, are capable of traveling relatively long distances within the material without leaving

dislocation segments behind at the y/y' interfaces. It should also be noted that while the first pair

of dislocations is impeded by the formation of the APB, the second pair is aided by the

elimination of the APB. Large creep strains can be produced by this mechanism in the absence

of forest dislocations at the y/y' interfaces. As common a/2<1 10> matrix dislocations expand to

fill the y channels and create interfacial networks, the difficulty of cooperative shear increases.

Eventually, the networks that are formed are sufficient to reduce the rate of shear of the y/y'

microstructure and a steady-state condition is reached where the dislocation density would

remain constant under constant stress creep conditions.16, 17, 63

Adding primary creep behavior to existing creep behavior models is a challenge due to the

complexity of the behavior. For example, as shown in Figure 8-5, the dislocation pairs required

to form the stacking fault pairs form from a/2<110> type dislocations. This implies that before

cooperative shear can take place, a population of a/2<1 10> dislocations must already be present

in the microstructure.16 For this reason, it is believed that this population is produced during the

incubation period that often precedes primary creep. In this way, alloys that deform by y' shear

exhibit the same deformation mechanisms prior to the start of primary creep that are found in

alloys that display tertiary creep behavior only. Developing a model that can predict when the

transition to y/y' shear will take place has proven difficult.63









Additionally, primary creep of superalloys does not always occur by y' shear due to

stacking fault ribbons. Alloys exhibiting tertiary creep do experience primary creep in the form

of a higher initial creep rate that is reduced during the first stages of exposure to creep

conditions. For conditions governed by tertiary creep, primary creep typically occurs by <110>

dislocations bowing to fill the y channels. This process is similar to the deformation process

exhibited during the later stages of tertiary creep; however, prior to the start of a creep test the

native dislocation density is low. Due to this low initial density, the creep rate is more rapid than

in later stages when dislocation-dislocation interactions are common. Primary creep under these

conditions occurs as the small dislocation population is increased until the dislocation

interactions bring about the transition to tertiary creep (ie. the creep rate is reduced to the

minimum creep rate immediately following primary creep).74

Creep Results

The primary creep behavior of single crystal nickel-base superalloys has proven to be

controlled by multiple factors. Previous research has shown a dependence of primary creep on

orientation, magnitude of the load applied, secondary y' precipitates, lattice misfit, and possibly

rhenium and/or ruthenium content, all of which may lead to non-uniform deformation (only 1 or

2 slip systems) and stacking fault shear (of the y' phase).16, 17, 19, 20, 23, 24, 27, 33, 47, 56-58 COmparing

yield strength to the applied initial creep stress, Table 5-3, it is apparent that the large primary

creep strains of PWA 1484 at 7040C were produced at greater than 90% of the yield strength.

This high load condition produced primary creep strains of 17% for the LT age and 24% for the

HT age. Such a high stress level, however, was not necessary to produce large primary creep

strains in PWA 1484. Also at 7040C and a reduced load of 758 MPa, primary creep strains of

5% for the LT age and 10-14% for the HT age were observed. This clearly indicates that the









magnitude of the applied initial stress contributes to the amount of primary creep that results, as

reported by (Rae and Reed, 2006) and (Shah et al., 2004).16 47 Under creep testing at conditions

that yield large primary creep strains, raising the initial stress results in larger primary creep

strains and shorter times to complete primary creep.

This lower stress ratio for PWA 1484 at 8150C is similar to the stress ratios for creep

testing of PWA 1480 and PWA 1480+ at 7040C. Even still, the primary creep of PWA 1484 is

significantly larger than the other two alloys. Additionally, a PWAl480 HT specimen was tested

at 7040C/1200 MPa which represents a creep stress to yield stress ratio of 0.9 (90%). The PWA

1480 HT specimen failed in 0.9 hours with no discernable primary creep stage because the creep

rate exhibited continuous acceleration beginning with the earliest measurements. These results

demonstrate two related points. First, the ratio of creep stress to yield stress does not cause

primary creep nor do high ratios increase the primary creep seen in PWA 1480. The second

point is that the magnitude of the applied stress, while increasing the primary creep strain

produced in alloys prone to large primary creep strains, does not cause large primary creep

strains in alloys resistant to primary creep.

Related to the previous discussion are reports that there may be a stress threshold that,

above which, stacking faults in the y' phase may form due to the entry of pairs of dislocations

that dissociate into pairs of stacking faults, separated by an APB, that then shear the y' phase. It

is thought that at stresses below this threshold matrix dislocations do not shear precipitates but,

instead, exhibit cross-slip or climb to bypass barriers such as the fine secondary y' in the y

channels. During this process work hardening is more likely due to the distribution of slip on

multiple planes. If the stress exceeds this threshold, however, climb and cross-slip do not occur

and deformation is able to proceed without generating significant work hardening and secondary









creep is delayed.16, 17, 20 IHCOrporating the Eindings of the current investigation, these results are

likely to only apply to alloys prone to primary creep. The distinguishing factors that differentiate

alloys exhibiting primary creep are not low yield strengths or the magnitudes of the applied

loads, but microstructural differences as a result of alloy chemistry and processing that resist

primary creep strain production.

When dislocations are confined to the y matrix, primary creep is low due to the interaction

of dislocations within the narrow spaces of the y channels. During y' shear, dislocations can

move longer distances without interacting with an obstacle in the y' precipitates resulting in large

creep strains and few dislocation interactions.16 Creep behavior following the HT age heat

treatments supports these Eindings. The longer time and higher temperature of the HT age allows

the y' to coarsen, decreasing coherency. The greater degree of incoherency of the y' precipitates

after the HT age may be responsible for the decrease in primary creep at all test conditions for

PWA 1484 while no significant change in primary creep can be found in PWA 1480 and PWA

1480+. Increasing incoherency in PWA 1484 following heat treatment may reduce the ability of

dislocations to enter the y' phase to form stacking faults. PWA 1480 shows little effect with heat

treatment because the formation of stacking faults is already difficult, Figures 6-11 to 6-15.

PWA 1480+, while able to produce stacking faults, deforms by multiple deformation systems so

that work hardening is rapid and secondary creep starts soon after the initiation of primary creep.

Because both PWA 1480 and PWA 1480+ are prone to wide-spread dislocation and stacking

fault interactions, leading to hardening, they show little change in primary creep strain with age

heat treatment.

Research into primary creep mechanisms has yielded interesting findings into the

efficiency of these two shear mechanisms. Calculating dislocation densities and predicting the









expected amount of strain produced for each mechanism, differences in the ability of the

different processes to confer shear can be seen. For alloys that deform primarily by the

movement of matrix dislocations, the matrix channels become filled with dislocations and

dislocation interactions become common following 0.3% to 0.5% creep deformation. The large

amount of dislocation interactions at this point is related to the onset of secondary creep, where a

balance exists between deformation processes (slip) and recovery processes. As a result, it is

predicted that deformation provided primarily the movement of matrix dislocations will only

yield about 0.5% primary creep at the maximum. If the ability to form stacking faults is

included, the dislocation density in the y matrix increases more slowly. The result is a

deformation process that can operate much longer before work hardening in the matrix causes

secondary creep. For this scenario, primary creep strains greater than 5% can be expected.16, 17

These results are consistent with the current investigation and those of other researchers.19, 20, 57

In the narrow sense of comparing overall rupture lives only, PWA 1484 out-performed

PWA 1480 at all but highest load at the lowest temperature (7040C/862 MPa). Simply adding

Re to PWA 1480, however, remarkably improved the rupture life of PWA 1480. Creep ductility

and toughness were reduced, but the minimum creep rate was decreased by over an order of

magnitude and the same low primary creep behavior was maintained. The time to 1% creep,

however, shows the difference in primary creep behaviors very clearly, Table 6-2. While the

lifetime of PWA 1484 is the longest at several conditions, the time to 1% creep is the shortest,

indicating rapid deformation early in the life of the specimens. Continuing the discussion to time

from 1% to 2% creep strain, it can be seen that PWA 1480 and PWA 1480+ experience a much

longer time from 1 to 2% than from 0 to 1%. For several PWA 1484 samples, however, the time









from 1% to 2% creep was achieved faster than the first 1% of creep strain because the creep rate

was still increasing through the first several percent primary creep.

If the usual benchmarks of rupture life, time to 1% creep, and minimum creep rate are

utilized in turbine engine design then an alloy might be selected that could exhibit rapid creep

during the early stages of it' s service life. Thus it is necessary to incorporate an understanding of

the nature of the active primary creep behavior into the usual design schemes. Within the

context of current investigation, a simple change in the aging heat treatment temperature of a

turbine component may be enough to manage the primary creep that will be produced. Another

method of avoiding large primary creep strains is to pre-creep the alloy prior to service. Some

reports indicate that specimens that usually yield large primary creep strains were crept to small

strains (<0. 1%) at 9500C before being subj ected to low temperature, high stress creep testing.

The primary creep produced during the second test was reduced to low values similar to those

produced by PWA 1480 and PWA 1480+. It is thought that this reduction in the expected

primary creep strain is related to the formation of matrix dislocations that then interfere with the

passage of stacking faults and bring about the onset of secondary creep earlier than in specimens

that were not pre-crept. These changes in processing are both examples of simple techniques

that can dramatically improve the usefulness of alloys that are prone to excessive primary creep

by stacking fault shear.

Modeling Primary Creep

There have been several models developed to describe the creep process for single crystal

nickel-base superalloys since the introduction of PWA 1480 in the early 1980's. Most of these

were made to describe the tertiary creep regime that is most common among superalloys. PWA

1480, for example, exhibits tertiary creep across a very wide range of stress and temperature










including loads as high as 90% of the yield strength of PWA 1480 at 7040C. The second

generation superalloy PWA 1484, however, exhibits a more complex combination of creep

behaviors. During high temperature, low stress creep PWA 1484 displays tertiary creep behavior.

At low temperatures and low stresses, tertiary creep still persists. If the stress is raised above a

critical threshold, however, PWA 1484 (and other second generation superalloys) produces large

primary creep strains followed by steady-state creep.16 It is this transition that has proven

difficult to predict with conventional models often leading to large inconsistencies between

models and experimental data.63

The alloy that has received the most attention for primary creep modeling is CMSX-4.16, 17,

63 CMSX-4 and PWA 1484 share fairly similar compositions, Table 2-1, and the reported creep

behavior and deformation substructures are similar in nature to those found for PWA 1484 in the

current investigation so modeling approaches for both alloys are expected to be similar.3, 16, 47

Deformation modeling of creep of superalloys usually begins by separating the deformation

gradient into constituents representing the y matrix and the y' precipitates using an equation

similar to the one given below:

12 12
Equation 8-8: P y aOi)r r aOi)
a=1 a=1

where j is the shear strain rate, d is the unit vector representing the slip direction, ii is the slip


plane normal unit vector, a designates a slip system (12 systems included in this calculation), and

fis volume fraction of each phase denoted by y and y' subscripts. Because the model ignores any

contribution by TCP phases or carbides, the sum of the volume fractions of y and y' are set equal

to one.









The next step is to define the shear strain rate (j ). This step typically begins with the

Orowan equation:

Equation 8-9: fjz = 9,, b -va

where p is the mobile dislocation density, b is the burgers vector, and v is the velocity of mobile

dislocations. Note that Equation 8-8 is specific to a given slip system. Also, shear strain rate

calculations are performed for each phase independently. Substituting an expression for the

mobile dislocation velocity gives:


Equation 8-10: jFCC = FCC b. FCC attack expl +e I
k T k,T



E~udoB-l Lk T kT

Where the sub scripts FCC and L12 denote the y and y' phase, respectively. Also, h denotes

dislocation jump distance, F denotes dislocation jump frequency, kb is Boltzmann's constant, z is

the resolve shear stress. Also note, the misfit stress, zmis, passing stress, z,ass, and Orowan stress,

zoro, are all accounted for in the y phase, but just the passing stress is involved in the calculation

for the y' phase. The Orowan stress as defined in the model is based on the y channel width and

does not account for secondary y' within the channels. The strengthening contribution from

secondary y' was not included as a separate contribution within any description of the creep

behavior so its contribution was most likely accounted for in the more general terms like passing

stress. Each of these terms are described in detail as is the rest of the model in the work by A.

Ma et al.63 As the model was developed, additional terms were added to account for additional

behaviors to be described.









The above described model goes on to account for <110> dislocation shear within the y

matrix as well as the generation of <1 12> stacking fault ribbons. To account for tertiary creep,

the model ignores dislocation ribbons and y' shear because tertiary creep occurs almost

exclusively by deformation contained within the y matrix. The distribution of slip between y

channels was found to be critical to the success of the model. The simplest microstructural

constituents. are a cube shaped y' phase, and three rectangular y channels (two sides equal to the

edge length of the y precipitate and one side equal to the specified y channel thickness) oriented

in three directions with the normal to the "plate" parallel to the cube directions [100], [010], and

[001].63 With the applied stress axis parallel to the [001] direction, one y channel is

perpendicular to the stress [001], and two are parallel [100] and [010]. Misfit stresses were

found to interact quite strongly with the active deformation mechanisms. Initially, deformation

takes place almost exclusively in the perpendicular, [001], channels due to the superposition of

the misfit stress and applied stress. As long as the y/y' interfaces remain coherent, the [001]

channels are preferred. As the misfit stresses are relieved, however, the preference is reduced

and deformation increases within the [100] and [010] channels.63

Lattice Misfit

Lattice misfit plays a significant role during both tertiary creep and primary creep. Alloys

with lower values for lattice misfit are reportedly more prone to the formation of the dislocation

ribbons responsible for y' shear and, therefore, large primary creep strains.59-61, 71, 85 Large,

negative lattice misfit values have been shown to produce better creep resistance at high

temperatures and are thought to be beneficial to the prevention of large primary creep strains.

Alloys with large lattice misfit have been shown to form dislocation networks along the y/y'

interfaces more quickly due the large misfit stresses that result. Dislocations then build up at the










y/y' interfaces to relieve the misfit. As the interface becomes lined with dislocations forming the

interfacial network, shear of the y' phase becomes more difficult, the creep rate is reduced, and

secondary creep begins. Alloys with low misfit are more coherent and it has been suggested that

low misfit alloys are prone to shear of the y' phase. In addition, the dislocation ribbons that form

glide relatively easily and leave little, if any, dislocation segments at y/y' interfaces. These two

factors combine to delay the formation of the interfacial networks that herald the start of the

secondary creep stage. As a result, the primary creep stage (y' shear) occurs over a longer time

allowing for the production of greater primary creep strains.16

In the current investigation, PWA 1484 was shown to have the largest, negative lattice

misfit (-0.20% to -0.25%) of the three alloys tested in this investigation. This finding would tend

to contradict conventional wisdom regarding the effect of lattice misfit due to the large degree of

stacking fault shear in PWA 1484 during primary creep. PWA 1480 exhibited low lattice misfit

values ranging from 0.00% to -0.14% and the Re addition to PWA 1480 increased the lattice

misfit values to a slightly more positive range (-0.05% to +0.034%) for PWA 1480+. These

results suggest that the emphasis placed on lattice misfit in creep modeling may need to be

reevaluated. While it is clear, that lattice misfit plays a large role in tensile and creep

deformation of superalloys, the more recent second and third generation superalloys have larger

magnitudes for lattice misfit than the older first generation alloys. The first generation alloys,

however, do not deform by large-scale y' shear unlike the second generation alloys PWA 1484

and CMSX-4. Additional work is clearly necessary to clarify the effect of lattice misfit while

modeling primary creep behavior these alloys.









Secondary y'

In addition to accounting for the nature of the y/y' interface during creep deformation, the y

channels play a vital role in creep (both primary and tertiary). The y channels are the location

where <1 10> dislocations nucleate and propagate leading to the nucleation of dislocation ribbons

in the y' phase or y/y' interfacial networks. In addition to y channel thickness, discussed earlier,

the secondary y' that forms within the matrix interact with dislocations as they expand to fill the

matrix. Despite these interactions, the secondary y' precipitates are omitted from most models

even though y channel thickness is nearly unanimously utilized for creep and tensile modeling.

There is some disagreement as to which stage of heat treatment is responsible for

precipitating secondary y'. The presence of secondary y' has been shown to occur following

solution heat treatment with coarsening occurring during the ensuing aging heat treatment and

secondary y' has been reported to appear during rapid cooling from the final aging heat

treatment.20 In either case, the secondary y' are likely to always be present in the y channels

prior to service for most high volume fraction superalloys. For PWA 1480, PWA 1480+, and

PWA 1484, secondary y' was shown to be present in all three alloys following both age heat

treatments. These precipitates are expected to be highly coherent with the matrix due to their

small size resulting in an increased probability of shear by matrix dislocations bowing within the

y channels.

Using the LEAP technique, compositions for a small sample of secondary y' precipitates in

PWA 1484 were determined. While no statements can be made with high statistical certainty,

four composition variation trends could be seen with increasing size. The larger secondary y'

precipitates contained higher levels of Ni and Ta and lower levels of Co and Re. This indicates

that as these precipitates, or potentially clusters, grow in size, strengthening is likely to occur due










to the increased y' former content. The larger secondary y' precipitates reside towards the

centers of the y channels with smaller secondary precipitates near the relatively large primary

precipitates. As a result, the interaction between secondary y' and matrix dislocations is

complex. Despite the differences in strength and size distribution across the y channel, it is

likely that matrix dislocations shear these ultra-fine precipitates during tensile testing. Creep

testing, however, allows for thermally activated processes to occur that might alter the paths of

dislocations when they meet these precipitates. For alloys that exhibit tertiary creep, such as

PWA 1480 and PWA 1480+, matrix dislocations remain in the <110> form and likely shear

through or cross-slip past the secondary y'. Alloys that produce large primary creep strains, such

as PWA 1484, deform largely by <112> dislocations separated by stacking faults. While these

dislocations and stacking faults typically nucleate y/y' interfaces prior to shearing the y' phase,

this configuration is relatively stable in the y matrix as well. As a result, these dislocation

ribbons likely shear the secondary y' in a similar manner to shear of the primary y'. It has even

been suggested that the secondary y' may act to stabilize the dislocation ribbons through the y

channels between primary precipitates. Additionally, it should be noted that there is no APB

formed in the y matrix so the dislocation pairs (and associated stacking faults) tend to separate

further apart in the y channels.16, 17, 63 Clearly the secondary y' present in the y channels plays a

role in creep deformation during both primary and tertiary creep. The ability of PWA 1484 to

produce large primary creep strains while PWA 1480 does not, though, must be linked to another

material attribute as the uniform presence of secondary y' in all three alloys did not yield similar

behaviors.









Composition

Alloy composition clearly controls several aspects of mechanical behavior. From

controlling the strengths of phases to the lattice misfit, composition plays a direct role in

deformation. Several key compositional changes were made from PWA 1480 to PWA 1484.

While the effect of composition has already been discussed, a few points will be readdressed.

Rhenium additions to superalloys have been praised for their potency in solid solution

strengthening of the y phase. This effect was so significant that the first three generations of

superalloys are defined by their Re content (0 wt%, 3 wt%, and 6 wt%, respectively). Despite

the clear benefit in strength, the addition of Re has been identified as potential contributor to

large primary creep strains due to the large number of second and third generation alloys that

show this behavior. While it is possible that Re has some sort of contributory effect on primary

creep in second generation and later alloys, this investigation showed an improvement in creep

properties and a slight decrease in primary creep when Re was added to PWA 1480. It was

found, however, that PWA 1480+ appeared to produce more stacking faults than PWA 1480

indicating a possible decrease in the stacking fault energy of the y' phase. A lower stacking fault

energy would allow dislocations shearing the y' phase to spread further leading to an increased

potential of dislocation-dislocation interactions. In fact, the stacking faults observed in PWA

1480+ exhibited many interactions similar to the deformation found in specimens of other alloys

interrupted near failure. As a result of this and other investigations, the effect of Re on the

microstructure and properties of single crystal nickel base superalloys has been regarded as

positive. Specifically, Re additions increase lattice misfit, improve creep life, improve tensile

strength, reduce the y' size, and reduce the rate of y' coarsening.2-4, 8, 9, 77 One disadvantage with









the use of large Re additions, though, is decreased stability against TCP phase formation as

shown by the precipitation of TCP phases during primary creep of PWA 1480+.

There is also agreement within the superalloys community that Re segregates to the y

matrix during nucleation and growth of y' precipitates. This effect is expected to cause an

increase in lattice misfit near the y/y' interface as Re is rej ected from the y' precipitates result in

an enriched region in the y matrix. Research has also shown, however, that in the presence of W,

up to 20% of the Re that has been added will segregate to the y' precipitates.35 As rhenium

additions modify the compositions of the y and y' phases, changes in deformation mechanisms

might be expected as a consequence of the potent strengthening effect.

Aside from the well known solid solution strengthening effect are reports of the formation

of Re-rich clusters in the y phase near the y/y' interface. Understanding the nature of the

distribution (in the y matrix) of the Re that has been rej ected from the y' phase could lead to

improved model accuracy and potentially aid in the search for a Re replacement. Reported

research is still contradictory in this area. The two leading theories, currently, indicate that Re

may either form clusters around 10 nm in diameter or may form hardened shells in the y matrix

around the y' precipitates. Recent studies, including the present one, have attempted to utilize

the Local Electrode Atom Probe (LEAP) to examine the local concentration distribution of Re in

the y matrix with limited results. While publications have yet to be produced, a recently

presented collaboration between Cambridge University (Cambridge, UK) and Oak Ridge

National Lab (Oak Ridge, TN) pointed to possible evidence of the formation of Re clusters. The

formation of clusters by solid solution strengtheners is not a new phenomenon and has been

shown to occur with Cr and W as well.8, 35, 50 Chromium clusters have been found to have the

DO22 CryStal structure and are also formed due to rejection from the y' phase. It is expected that










these clusters would make dislocation shear of the y phase in the vicinity of the y/y' interface

more difficult. Dislocations that encounter these regions are likely to be disrupted by changes in

crystal structure and strength leading to greater buildup of dislocations at the y/y' interface.

How Re clusters or shells interact with dislocations and stacking faults during primary

creep is certain to be the subj ect of many future investigations as, clearly, additional work is

necessary. It still remains unclear what the interactions are between the rej ected Re in the y

phase and the matrix dislocations. Additionally, the high price of Re has forced both industry

and academia to begin researching viable alternative strengtheners to replace or reduce Re in

future superalloys. Continued research into the rhenium effect as it relates to these interactions is

likely to grow in the coming years.

The difference between PWA 1480 and PWA 1484, however, is more than just the

addition of rhenium. PWA 1484 has increased solid solution strengthener/refractory content and

decreased y' hardening. Based on the present investigation, Re alone does not appear to cause

increased primary creep. In fact, the opposite effect was realized in PWA 1480 when Re was

added. With the many changes in alloy composition between the first generation PWA 1480 and

the second generation PWA 1484, it is likely that the stacking fault energy and anti-phase

boundary energy were also modified. These two properties very strongly influence shear of the

y' phase during creep leading to the possibility of reduced resistance to large primary creep

strains caused by y' shear.

Concluding Remarks

The large primary creep strains exhibited by PWA 1484 were shown to occur through

deformation processes consistent with y' shear reported among some second generation

superalloys. The first generation alloy PWA 1480, however, exhibited tertiary creep behavior









that is commonly reported among superalloys. The goal of this investigation was to study the

effect of Re and secondary y' precipitates on the primary creep behavior of PWA 1480 and PWA

1484. In order to see the effect of Re additions on PWA 1480, a third alloy was created by

adding 3 wt% Re to PWA 1480 (the new alloy was named PWA 1480+ for this study).

Through this investigation, several aspects of the primary creep behavior of the PWA 1480

and PWA 1484 alloy systems have become clear. First, while all three alloys contained

secondary y' precipitates within the y channels, only PWA 1484 exhibited large primary creep

strains. While secondary y' has been linked to large primary creep strains in CMSX-4, primary

creep was limited for both PWA 1480 and PWA 1480+ despite the presence of secondary y' in

the matrix. It is reasonable to expect that these ultra-fine precipitates are rather coherent in

nature due to their small size and, as such, would be susceptible to shear by matrix dislocations

during tertiary creep as well as by dislocation ribbons (pairs of stacking faults) during primary

creep.

Second, the addition of rhenium did not increase the primary creep strains produced in

PWA 1480. This alloy addition (PWA 1480+) was responsible for decreased primary creep,

lower creep rates, and longer creep lifetimes. While PWA 1480+ is clearly not suitable for

service due to the early formation of TCP phases and reduced ductility, the potency of the solid

solution strengthening effect of Re was again proven. Significant increases in yield strength and

temperature capability of PWA 1480+ also accompanied the improved creep properties when

compared to PWA 1480. The exact nature of the "Rhenium Effect" still eludes researchers to

this day; however, it can be seen to offer substantial benefits over the current alternative

strengtheners. The LEAP data presented herein affirmed the expected rej section of Re from the y'










precipitates into the y matrix. An enriched layer was found along the y side of the y/y' interface,

though evidence proving the formation of clusters or shells was not produced.

Third, the HT age was reasonably effective at reducing the primary creep strains of PWA

1484 without reducing the lifetime or increasing the secondary stage creep rate. This effect has

been attributed to a decrease in coherency at the y/y' interface. Decreased coherency may have

increased the difficulty of nucleating dislocation ribbons in order to shear the y' phase. This

effect would result in increased hardening and an earlier termination of primary creep due to the

generation of more interfacial dislocations. PWA 1480 showed no response to age heat

treatment because the already difficult ability to produce stacking faults within the alloy

prevented significant primary creep. As a result, PWA 1480 only exhibited the amount of

primary creep that would be expected from an alloy deforming primarily by dislocations

confined to the y matrix. The Re modification of PWA 1480 was observed to increase stacking

fault production slightly; however, the stacking fault energy may have been lowered leading to

increased dislocation interactions and the formation of "locks" leading to low primary creep

strains. More work is necessary to fully understand the nature of the primary creep deformation

mechanism of PWA 1480+ due to the appearance of deformation structures more often reported

in specimens terminated near failure.

The origin of large primary creep strains is likely not the result of a single cause. Several

investigations have attempted to point out the primary cause only to reveal several underlying

issues that contribute to the primary creep behavior. The current investigation approached the

problem of primary creep by attempting to illustrate the effects of rhenium additions and age

heat treatment. Both of these processing changes can impact primary creep. The Re addition

reduced primary creep and improved creep performance of PWA 1480. Utilizing a higher age










temperature produced less primary creep in PWA 1484 specimens. Several other factors,

however, should be considered when designing alloys to avoid this behavior. Alloy composition,

for instance, involves several synergistic effects among alloying additions. Even small changes

in composition can impact microstructural properties like lattice misfit, stacking fault energy,

APB energy, and the strengths of the y and y' phases (among others). The primary creep

behavior of PWA 1484 is most likely due to an "ideal" combination of the following: lower

volume fraction of y', changes in SFE and APB, slightly narrower y channel widths, stabilization

of stacking faults across y channels by secondary y', higher misfit stresses superimposed on the

applied stress, and an increased ease of nucleation of <1 12> dislocation ribbons in the y' phase.

Future development of single crystal alloys for gas turbine engines will continue to pursue

higher temperature capability; however, additional emphasis has been placed on lower

temperature properties as the applied loads are increasing in the lower temperature regions of

turbine blades. While the increased concern regarding large primary creep strains will continue

to yield a deeper understanding of this behavior, the effect of rhenium will bear greater scrutiny

in the future. Rhenium related research will focus on developing new alternative strengthening

approaches in order to reduce Re consumption. As a result, renewed investigation into how Re

achieves its strengthening effect, such as this study, will aid future alloy development strategies.









CHAPTER 9
CONCLUSION

Conclusions

This investigation into the primary creep behavior of the single crystal nickel base

superalloys PWA 1480 and PWA 1484 has led to the following conclusions:

* The ability of PWA 1484 to exhibit large primary creep strains during testing at low
temperatures and high stresses appears to be linked to a combination of many factors
including secondary y', composition, y channel width, stacking fault energy, anti-phase
boundary energy, and y' volume fraction.

* The first generation superalloy PWA 1480 did not produce large primary creep strains
when subj ected to a wide range of creep loads (including the application of stress equal to
90% of the yield strength of PWA 1480). From these results, it is concluded that that large
primary creep strains the result of a combination of microstructural and compositional
attributes (such as listed above for PWA 1484).

* Stated another way, it is unlikely that a single cause is responsible for the production of
large primary creep strains. The first generation superalloy, PWA 1480, appears to be
incapable of exhibiting large primary creep strains because the deformation during creep is
contained within the y matrix as <110> dislocations that shear the secondary y', but not the
primary y' precipitates. The second generation superalloy, PWA 1484, allows the matrix
dislocations generated during the incubation period to enter the y' as <112> dislocation
ribbons that are capable of shearing the y' phase on a large scale. The specific
microstructural attributes responsible for the entry of dislocations into the y' phase are still
not entirely clear and are likely the result of increased sensitivity due to many factors.

* The rhenium addition to PWA 1480+ resulted in lowered primary creep and secondary
creep rate as well as increased rupture life (compared with PWA 1480). This effect is most
likely due to the already known solid solution strengthening effect and a lowered stacking
fault energy. The former effect increased load bearing capability and resistance to
dislocation motion, while the latter effect resulted in wider separations of dislocations in
the y' phase that led to dislocation interactions and lock formation. Both effects aided in
improving the creep life of PWA 1480+.

* The use of the high temperature aging heat treatment reduced the primary creep strain
exhibited by PWA 1484 by nearly half at all conditions. This effect is attributed to
reduced coherency at the y/y' interface that resulted in an increase in interfacial
dislocations. These "forest" dislocations added resistance to the y' shear mechanisms
active in PWA 1484 as well as interacted with matrix dislocations to form interfacial
networks of dislocations earlier during primary creep. The earlier formation of dislocation
networks would lead to the earlier onset of secondary creep and a reduction in the primary
creep strain exhibited by the alloy.









Future Directions


As a result of this work, several future paths of research have been opened. These future

directions are listed in brief below:

* Additional work involving alloy modification and the resulting effects on overall and
primary creep behavior: Suggested alloy modifications are PWA 1484 without Re and
alloys that substitute W for Re at various ratios. Removing Re from PWA 1484 would be
useful to determine if the stacking fault formation in the alloy is indeed linked to Re
content. Substitutions of W (or another refractory solid solution strengthener) for Re could
be used with a slightly different goal. The rising cost of Re warrants investigation in
strengthening alternatives that are lower in cost and are more readily available. As with
any alloy modification work, these proj ects would require broad characterization and
testing to examine the effects of the modifications.

* Continued X-ray diffraction studies of the change in lattice misfit as a function of heat
treatment: Additional XRD studies would be useful to aid in determination of the
correlation between lattice misfit and primary creep. Because primary creep appears to be
impacted directly by the y/y' interface, any technique that sheds light on the nature of the
interface would be useful in generating a complete understanding the primary creep
process.

* Additional creep testing with specimens prepared without secondary y' could eliminate
confusion regarding the effect these precipitates have on the primary creep behavior of
these alloys. While a simple furnace cool following aging is the only processing change
necessary to create this condition, extra coarsening and decreased coherency of the y' will
result. Care would need to be directed towards proper heat treatment development to
ensure that the end product maintains similar values for y' size and coherency. This is to
eliminate the possibility that the decrease in primary creep associated with secondary y'
free matrix channels is actually due to an enhanced ability to work harden due to increased
incoherency.

* Renewed work on the Local Electrode Atom Probe could serve to answer the question of
Re shell or cluster formation. As Re is rej ected from the precipitates into the matrix,
exactly how the element is distributed in the y matrix is in question. If ordered clusters or
disordered shells are found, then the use of other strengthening elements that behave
similarly might improve the strength of alloys and/or allow for further reductions in Re
content.

* Detailed TEM analysis of the deformation mechanism in PWA 1480+ could yield
information on how the alloy achieves such a low creep rate and long lifetime. The early
appearance of deformation structures commonly tied to failure requires investigation to
understand how PWA 1480+ is capable of superior creep properties to PWA 1484 over the
temperature range in question (7000C-8150C)


200









APPENDIX A
DIFFERENTIAL THERMAL ANALYSIS

Differential thermal analysis, or DTA, was performed on six specimens for the purpose of

designing new solution heat treatments for all three alloys: PWA 1480, PWA 1484, and PWA

1480+. In the case of single crystal superalloys, the solution heat treatment is necessary to

dissolve the y/y' eutectics that remain after solidification as well as to promote homogenization

of the alloy due to the high degree of segregation exhibited between the dendrites and the

interdendritic regions of the alloy. Thus a solution heat treatment for these alloys must not

simply exceed the y' solvus temperature long enough to dissolve the remaining eutectics. These

heat treatments must also approach the solidus temperature in order to deliver enough thermal

energy into the alloy to allow for diffusion of slow diffusing elements like tungsten, tantalum,

and rhenium. The final heat treatments selected for this study stop just 50C, 80C, and 30C short

of the solidus temperatures for each of the aforementioned alloys, respectively. It should also be

noted that the final hold temperature for the PWA 1484 HT3 heat treatment equals the solidus

temperature following the HT1 heat treatment. The final temperature for PWA 1480+ exceeds

the solidus of the as-cast condition by 190C. The fact that the HT3 heat treatment can be

performed without incipient melting in the alloy is an indication of just how much the solidus of

these alloys can be suppressed as a result of segregation. Throughout the course of the heat

treatment, the solidus temperature of the interdendritic region climbs as homogenization

proceeds. The solidus values after the HT2 heat treatment are likely within a few degrees of the

actual values for these alloys44, 45, 53





" .n ~d


1347


1303 92"C

1279.33"C


O
- 0


1340.50*C


1316.53"C


1000
Exo Up


11'00


13~00


1400


1500Y.
Universal V3.28 TA Instruments


'1200


Temperature ("C)


Figure A-1. DTA trace for PWA 1480 in the HT1 condition.


1287.52"C













1297.61*C 1337.98"C



















1290.68*C

1362 57 C


0.4 O
U


O

C3


02












-0 2


1375,82*C


,~ -~r-U h~-----.l.-


Eixo Up


"TOImperature ("C) Universal V3.29 TA Instruments


Figure A-2. DTA trace for PWA 1484 in the HT1 condition.










I


-- --1 1 I I I I I 1 I r I I r T


1353 22'C








~C


1266.71"C


1290


- 0.4 O








CL


E
a,


`\~~


1341.13"r




1311.49*C


13~00


1400


1000
Exro Up


1100


1200


1500
unwaersal'~ va TA mInstruments


Figure A-3. DTA trace for PWA 1480+ in the HTO condition.


Temperatures ("C











1307.19' C .


1290 54 C










1298.11*C


1357.14'C


-0.4





-0.2





-0.0 0.
E


1-=




04



1600


1349.57*C


Exo Up


1100


"12b00


13100

Temperature ("C)


14700


1500


Universal V3.2B TA Instnrments


Figure A-4. DTA trace for PWA 1480 in the HT2 condition.











I ?


1314.76"C 1345.79*C



1405.57'C


-04


2,

C1
E:
0 0

a,

0.


-02 4
160
Unvsl 328TAistuent


1384.
1299.62*C



-- ---


" -L

E


10900
Exo Up


1100


1200


13~00

Terrperature ("C)


1400 0


1500


Figure A-5. DTA trace for PWA 1484 in the HT2 condition.


1291.30*C





1310 97"C


-0.4
O







O

- 0.2



1600


1288,27"C


1298.1


~Y I.~ :~~*S~I~~J~-


1358.65"C


19000
Exo Up


1100


12b00


13~00
Te operate ("C


1400


1500


Universal V3.2B TA instruments


Figure A-6. DTA trace for PWA 1480+ in the HT2 condition.









APPENDIX B
XRD PEAK DECONVOLUTION

The use of X-ray diffraction to determine the lattice parameters for the y and y' phases for

misfit determination requires the use of peak deconvolution to separate the contribution of both

phases from the total intensity of the peak. The ordered L12 Space group of the y' precipitates

allows for all simple cubic reflections to appear due to the difference in structure factor of the Ni

and Al atoms in the unit cell. The fcc space group of the y matrix, however, only allows for

planes in which hkl from the miller index of the plane in question is all even or all odd. For

example, the (002), (222), (333) planes show positive intensity while the (001), (011), (321)

planes will show no intensity.

In single crystal specimens oriented with a [001] surface normal under Cu k, radiation as

used in this study, the y phase will only yield the following peaks between the 26 angles of 100

and 1400 (using the diffractometers described in Chapter 3): (111), (002), (022), (113), (222),

and (004). The y' precipitate phase will yield peaks at the same planes as the y phase with the

addition of mixed index phases like (001) and (0 11). From the above list of peaks produced by

the y phase, the (002) and (004) peaks show the greatest intensity and as a result are the most

useful for quantitative XRD analysis. The (002) peak can be found for all three alloys between a

26 angle of 50.50 and 51.50 while the (004) peak can be found between 117.50 and 119.00

While both peaks can be used for lattice parameter determination (and subsequent misfit

calculation), the higher angle peak, (004), will produce the more accurate result because at

higher 26 angles the contributions from the y and y' phase have a greater separation than at lower

angles. Additionally, since the radiation used contains both Cu kni and Cu ku2, the separate

contribution from both wavelengths of radiation is also displaced further.


208









Peak deconvolution was accomplished using the MDI Jade (ver. 7) software package on a

workstation at the University of Central Florida's Advanced Materials Processing and Analysis

Center (AMPAC). The software package allows for easy manipulation and indexing of

diffraction patterns as well as peak deconvolution. Utilizing 6 constraints including peak height,

Full Width at Half Maximum (FWHM), location, skew (symmetry), and type of distribution

(Gaussian, Lorentzian, etc.), XRD scans of (002) and (004) peaks could be mathematically

separated into as many peaks as necessary. For this investigation, scans were taken of both the

(002) and (004) peaks. Each peak was then separated to produce a total of four peaks as follows:

y phase Cu kal, y phase Cu ku2, y' phase Cu kni, and y' phase Cu ku2. All peaks were assumed to

have a Gaussian distribution in agreement with published work.46, 59, 60, 85 Deconvolution was

performed by first selecting the approximate peak locations with a special "peak fit cursor" and

the peak fit window was opened. After selecting the approximate location of the Cu kni peaks

for the two phases (the software automatically accounts for the presence of Cu ku2), the

deconvolution algorithm allows for further modifying of the starting conditions. The Gaussian

distribution was selected and the skew was set to zero initially. The FWHM was not specified to

allow the software to find the best fit. The first deconvolution was then performed.

The result of the deconvolution appears as a table of data describing the fit as well as a plot

of the raw data and the model fit. The calculated centers of the Cu kml peaks for both the y and y'

phases were used for lattice parameter determination. To further reduce the error of fit,

deconvolution was run a second time with the option to skew the peaks in use. This allows for a

"better" fit (or reduction in residual error of fit), but is not necessarily better for analysis. Each

deconvolution should be evaluated to ensure that the algorithm did not skew the peaks too much

resulting in unrealistic peak shifts.


209










The tables presented in Tables B-1 through B-8 are the resulting lattice misfit values

calculated from the results of the deconvolution of the XRD data that were collected. The

figures that follow are the results of the deconvolution algorithm. They have been included to

aid in the understanding of the deconvolution process and to compare the quality of the fit for

each peak. Each peak was analyzed with this method at least once, while several were analyzed

several times. The degree of peak skew is highlighted at the top of each figure. Blue designates

a result with zero skew while orange designates nonzero skew values. Also given are the Full

Width at Half Maximum values, which help to define how broad or narrow a peak may be. For

the purposes of lattice parameter calculation, the tall, narrow peak is assumed to be the y'

contribution while the shorter, broader peak is assumed to represent the y contribution. This

assumption is made because the y' phase is highly ordered (therefore narrower) and though it is a

precipitate phase it is present in volume fractions greater than 70% (therefore taller, more

intense). The solid solution strengthened y phase would be expected to be significantly broader

in width and of less intensity due to a smaller volume fraction and the random nature of the solid

solution.46, 59

Note: Two methods were employed to codify the data presented in Tables B-1 through B-

8. A number code was used to differentiate the averaging heat treatments and a color code was

used to differentiate attempts at deconvolution. These two methods are explained below:

* Number code: The last digit in the specimen name indicates which of the four averaging
heat treatments to which the specimen was exposed: xxxx3 for 4 hr., xxxx4 for 10 hr.,
xxxx5 for 100 hr., and xxxx6 for 1000 hr. at 10800C
* Color code: Tan for the first attempt, Blue for skewed deconvolution, and yellow for
unskewed deconvolution


210











Table B-1. PWA 1480 (002) peak lattice misfit calculations.
Ir+k +12 tall/short? 29 9 h2 sin 9 a2 a 6
4 t 50.81 25.41 2.373 0.1841 12.8943 3.5909 0.0918
4 s 50.86 25.43 2.373 0.1844 12.8707 3.5876
4 t 50.82 25.41 2.373 0.1841 12.8924 3.5906 0.0790
tC) 4 s 50.86 25.43 2.373 0.1844 12.8721 3.5878
4 07 53 .33 013 2957 352 007
4 s 50.78 25.39 2.373 0.1839 12.9076 3.5927 007
4 08 54 .33 014 2895 351 011
0\ 4 s 50.72 25.36 2.373 0.1835 12.9361 3.5967
,, 4 t 50.85 25.42 2.373 0.1843 12.8778 3.5886 -0.0092
4 s 50.84 25.42 2.373 0.1843 12.8801 3.5896
4 t 50.77 25.38 2.373 0.1838 12.9142 3.5936 -0.1039
4 s 50.83 25.41 2.373 0.1842 12.8858 3.5897
4 t 50.79 25.40 2.373 0.1839 12.9038 3.5922 0.1873
d 4 s 50.89 25.45 2.373 0.1846 12.8556 3.5855
4 t 50.81 25.40 2.373 0.1840 12.8972 3.5913 0.01
0\ 4 s 50.81 25.40 2.373 0.1840 12.8967 3.5912
h 4 t 50.74 25.37 2.373 0.1836 12.9728 3.5956 -0.0202
4 s 50.73 25.36 2.373 0.1835 12.9333 3.5963
4 t 50.74 25.37 2.373 0.1836 12.9276 3.5955 -0.1020
4 s 50.67 25.33 2.373 0.1831 12.9638 3.5600
4 t 50.73 25.37 2.373 0.1835 12.9309 3.5960 -0.38400
0\ 4 s 50.53 25.26 2.373 0.1821 13.0953 3.6096
,, 4 t 50.72 25.36 2.373 0.1835 12.9371 3.5968 -0.1416
4 s 50.80 25.40 2.373 0.1840 1.2.905 3.0917
4 t 50.68 25.34 2.373 0.1832 12.9552 3.5993 0.1289
4 s 50.75 25.38 2.373 0.1837 12.9218 3.5947
4 t 50.72 25.36 2.373 0.1835 12.9366 3.5967 0.0699
4 s 50.76 25.38 2.373 0.1837 12.92185 3.5942
~O 4 t 50.72 25.36 2.373 0.1834 12.9380 3.5969 0.0847
4 s 50.77 25.38 2.373 0.1837 12.9161 3.5939
0\ 4 t 50.71 25.36 2.373 0.1834 12.9438 3.5975 -0.1364
,, 4 s 50.64 25.32 2.373 0.1829 12.9777 3.6025
4 t 50.71 25.35 2.373 0.1834 12.9433 3.5977 -0. 1340
4 s 50.63 25.32 2.373 0.1829 12.9796 3.6027











Table B-2. PWA 1480+ (002) peak lattice misfit calculations.
Ir+k +12 tall/short? 29 9 h2 sin 9 a2 a 6
4 t 50.75 25.38 2.373 0.1836 12.9233 3.5949 0.1324
4 s 50.82 25.41 2.373 0.1841 12.8891 3.5901
4 t 50.72 25.36 2.373 0.1835 12.9356 3.5966 0.0552
4 s 50.75 25.38 2.373 0.1837 12.9214 3.5946
4 t 50.70 25.35 2.373 0.1833 12.9495 3.5985 0.0037
4 s 50.70 25.35 2.373 0.1833 12.9485 3.5984
4 t 50.70 25.35 2.373 0.1833 12.9495 3.5985 -0.0166
tC) 4 s 50.69 25.34 2.373 0.1832 12.9538 3.5991
4 t 50.73 25.37 2.373 0.1835 12.9328 3.5962 0.0405
C~ 4 s 50.75 25.38 2.373 0.1837 12.9223 3.5948
r( 4 t 50.73 25.37 2.373 0.1835 12.9318 3.5961 0.0239
4 s 50.75 25.37 2.373 0.1836 12.9257 3.5952
4 t 50.69 25.35 2.373 0.1833 12.9504 3.5987 -0.0682
4 s 50.66 25.33 2.373 0.1830 12.9681 3.6011
4 t 50.75 25.37 2.373 0.1836 12.9247 3.5951 0.0460
4 s 50.77 25.39 2.373 0.1838 12.9128 3.5934
4 t 50.75 25.37 2.373 0.1836 12.9238 3.5950 0.0276
4 s 50.76 25.38 2.373 0.1837 12.9166 3.5940
4 t 50.72 25.36 2.373 0.1834 12.9395 3.5971 -0.0607
d 4 s 50.75 25.37 2.373 0.1836 12.9238 3.5950
4 t 50.73 25.36 2.373 0.1835 12.9342 3.5964 -0.0516
C~ 4 s 50.70 25.35 2.373 0.1833 12.9476 3.5983
r( 4 t 50.73 25.36 2.373 0.1835 12.9342 3.5964 -0.0534
4 s 50.70 25.35 2.373 0.1833 12.9480 3.5983
4 t 50.67 25.34 2.373 0.1831 12.9600 3.6000 -0.1381
4 s 50.75 25.37 2.373 0.1836 12.9242 3.5950
4 t 50.70 25.35 2.373 0.1833 12.9471 3.5982 -0.1012
4 s 50.76 25.38 2.373 0.1837 12.9209 3.5946
C~ 4 t 50.53 25.26 2.373 0.1821 13.0300 3.6097 -0.4095
r( 4 s 50.75 25.37 2.373 0.1836 12.9238 3.5950
4 t 50.52 25.26 2.373 0.1821 13.0324 3.6100 -0.4206
4 s 50.75 25.38 2.373 0.1836 12.9233 3.5949
4 t 50.77 25.38 2.373 0.1837 12.9161 3.5939 -0.0129
4 s 50.76 25.38 2.373 0.1837 12.9195 3.5944
4 t 50.76 25.38 2.373 0.1837 12.9180 3.5942 -0.0129
4 s 50.75 25.38 2.373 0.1837 12.9214 3.5946
~O 4 t 50.76 25.38 2.373 0.1837 12.9171 3.5940 -0.0405
4 s 50.74 25.37 2.373 0.1836 12.9276 3.5955
C~l 4 t 50.76 25.38 2.373 0.1837 12.9190 3.5943 -0.0423
r( 4 s 50.74 25.37 2.373 0.1836 12.9299 3.5958
4 t 50.76 25.38 2.373 0.1837 12.9180 3.5942 -0.1160
4 s 50.70 25.35 2.373 0.1833 12.9480 3.5983
4 t 50.76 25.38 2.373 0.1837 12.9195 3.5944 -0.1123
4 s 50.70 25.35 2.373 0.1833 12.9485 3.5984


212











Table B-3. PWA 1484 (002) peak lattice misfit calculations.
Ir+k +12 tall/short? 29 9 h2 sin 9 a2 a 6
4 t 50.79 25.39 2.373 0.1839 12.9062 3.5925 0.1836
tC) 4 s 50.89 25.44 2.373 0.1846 12.8589 3.5859
C~ 4 t 50.86 25.43 2.373 0.1844 12.8702 3.5875 0.0349
0\ 4 s 50.88 25.44 2.373 0.1845 12.8612 3.5863
,, 4 t 50.83 25.41 2.373 0.1842 12.8872 3.5899 -0.2281
4 s 50.70 25.35 2.373 0.1833 12.9461 3.5981
4 t 50.77 25.39 2.373 0.1838 12.9123 3.5934 -0.0920
4 s 50.72 25.36 2.373 0.1835 12.9361 3.5967
4 t 50.66 25.33 2.373 0.1831 12.9643 3.6006 -0.2049
d 4 s 50.55 25.28 2.373 0.1823 13.0175 3.6080
C~ 4 t 50.66 25.33 2.373 0.1830 12.9671 3.6010 -0.1901
0\ 4 s 50.56 25.28 2.373 0.1823 13.0165 3.6078
h 4 t 50.64 25.32 2.373 0.1829 12.9767 3.6023 -0.2605
4 s 50.50 25.25 2.373 0.1819 13.0445 3.6117
4 t 50.63 25.32 2.373 0.1829 12.9786 3.6026 -0.2420
4 s 50.50 25.25 2.373 0.1820 13.0416 3.6113
4 t 50.49 25.24 2.373 0.1819 13.0493 3.6124 0.0740
4 s 50.53 25.26 2.373 0.1821 13.0300 3.6097





4 t 50.72 25.36 2.373 0.1834 12.9385 3.5970 0.0865
4 s 50.77 25.38 2.373 0.1837 12.9161 3.5939
4 t 50.82 25.41 2.373 0.1841 12.8896 3.5902 -0.2650
4 s 50.68 25.34 2.373 0.1832 12.9581 3.5997
~O 4 t 50.77 25.39 2.373 0.1838 12.9133 3.5935 -0. 1436
4 s 50.69 25.35 2.373 0.1833 12.9504 3.5987
0\ 4 t 50.77 25.39 2.373 0.1838 12.9123 3.5934 -0.0478
h 4 s 50.75 25.37 2.373 0.1836 12.9247 3.5951
4 t 50.73 25.37 2.373 0.1835 12.9318 3.5961 -0.2027
4 s 50.62 25.31 2.373 0.1828 12.9844 3.6034
4 t 50.73 25.37 2.373 0.1835 12.9309 3.5960 -0.1990
4 s 50.63 25.31 2.373 0.1828 12.9825 3 .6031


213





16 t 118.01 59.00 2.373 0.7348 12.9198 3.5944 0.2120
16 s 118.41 59.21 2.373 0.7379 12.8651 3.5868
16 t 118.01 59.00 2.373 0.7348 12.9198 3.5944 0.2156
r(16 s 118.42 59.21 2.373 0.7380 12.8642 3.5867
r( 16 t 117.98 58.99 2.373 0.7346 12.9229 3.5948 0.0519
16 s 118.08 59.04 2.373 0.7354 12.9095 3.5930
16 t 117.97 58.99 2.373 0.7345 12.9246 3.5951 0.2090
16 s 118.37 59.19 2.373 0.7376 12.8707 3.5876


Table B-4. PWA 1480 (004) peak lattice misfit calculations.


Ir+k +12 tall/short?


29 9 h2 sin 9


t 118.05 59.03 2.373
s 118.53 59.26 2.373
t 118.07 59.03 2.373
s 118.47 59.23 2.373

s 118.10 59.05 2.373
t118.04 59.02 2.373
t 118.02 59.01 2.373
s 118.47 59.23 2.373
t 118.04 59.02 2.373
s 118.59 59.29 2.373
t 118.01 59.00 2.373
s 118.53 59.27 2.373


0.7351
0.7388
0.7353
0.7383
0.7355
0.7351
0.7349
0.7383
0.7350
0.7392
0.7348
0.7388


12.9139 3.5936 0.2487
12.8499 3.5847
12.9116 3.5933 0.2092
12.8577 3.5858
12.9074 3.5927 0.0293

12.9150 3.5937
12.9177 3.5941 0.2332
12.8576 3.5857
12.9160 3.5939 0.2871
12.8420 3.5836
12.9195 3.5944 0.2733
12.8491 3.5846


12.9415 3.5974 0.1719
12.8971 3.5912
12.9259 3.5953 0.0541
12.9398 3.5972
12.9446 3.5979 0.1118
12.9157 3.5938
12.9443 3.5978 0.1055
12.9170 3.5940
12.9143 3.5937 -0.1376
12.9499 3.5986
12.9206 3.5945 -0.1963
12.9714 3.6016
12.9211 3.5946 0.0740
12.9403 3.5973
12.9442 3.5978 0.1275
12.9112 3.5932
12.9084 3.5928 0.1643
12.9509 3.5987
12.9193 3.5943 -0.1831
12.9667 3.6009


117.85 58.92 2.373
118.18 59.09 2.373
S117.96 58.98 2.373
S117.86 58.93 2.373
117.82 58.91 2.373
118.04 59.02 2.373
117.83 58.91 2.373
118.03 59.01 2.373
118.05 59.02 2.373
117.79 58.89 2.373
118.00 59.00 2.373
117.63 58.81 2.373
1118.00 59.00 2.373
S117.86 58.93 2.373
117.83 58.91 2.373
118.07 59.04 2.373
S118.09 59.05 2.373
S117.78 58.89 2.373
118.01 59.01 2.373
117.66 58.83 2.373


0.7336
0.7361
0.7344
0.7336
0.7334
0.7350
0.7334
0.7349
0.7351
0.7331
0.7347
0.7319
0.7347
0.7336
0.7334
0.7353
0.7354
0.7330
0.7348
0.7321


214











Table B-5. PWA 1480+ (004) peak lattice misfit calculations.
Ir+k +12 tall/short? 29 9 h2 sin 9 a2 a 6
16 t 118.10 59.05 2.373 0.7355 12.9069 3.5926 -0.0597
16 s 117.99 58.99 2.373 0.7346 12.9223 3.5948
C 16 t 117.96 58.98 2.373 0.7344 12.9264 3.5953 0.0084
16 s 117.97 58.99 2.373 0.7345 12.9242 3.5950
1~ 6 t 117.96 58.98 2.373 0.7345 12.9256 3.5952 -0.0718
r( 16 s 118.10 59.05 2.373 0.7355 12.9070 3.5926
16 t 117.98 58.99 2.373 0.7346 12.9238 3.5950 -0.0550
16 s 118.08 59.04 2.373 0.7354 12.9096 3.5930
16 t 117.98 58.99 2.373 0.7346 12.9233 3.5949 0.2833
16 s 118.52 59.26 2.373 0.7388 12.8503 3.5847
16 t 117.98 58.99 2.373 0.7346 12.9230 3.5949 0.0441
16 s 117.90 58.95 2.373 0.7340 12.9344 3.5964
16 t 117.99 58.99 2.373 0.7347 12.9222 3.5947
16 s 118.66 59.33 2.373 0.7398 12.8326 3.5823
16 t 117.98 58.99 2.373 0.7346 12.9237 3.5950 0.0304
16 s 118.04 59.02 2.373 0.7350 12.9158 3.5939
16 t 117.98 58.99 2.373 0.7346 12.9230 3.5949
S16 s 118.37 59.19 2.373 0.7376 12.8707 3.5876
16 t 117.98 58.99 2.373 0.7346 12.9238 3.5950 0.0288
16 s 118.03 59.02 2.373 0.7350 12.9164 3.5939
S16 t 117.97 58.99 2.373 0.7345 12.9245 3.5951 0.0210
16 s 118.01 59.01 2.373 0.7348 12.9191 3.5943
16 t 117.98 58.99 2.373 0.7346 12.9229 3.5948 0.0231
16 s 118.03 59.01 2.373 0.7350 12.9169 3.5940
16 t 117.98 58.99 2.373 0.7346 12.9235 3.5949 0.0194
16 s 118.02 59.01 2.373 0.7349 12.9185 3.5942
16 t 117.98 58.99 2.373 0.7346 12.9230 3.5949 0.0231
16 s 118.03 59.01 2.373 0.7349 12.9170 3.5940
16 t 118.02 59.01 2.373 0.7349 12.9183 3.5942 0.0205
16 s 117.98 58.99 2.373 0.7346 12.9235 3.5949
16 t 118.11 59.05 2.373 0.7356 12.9062 3.5925 0.1055
16 s 118.31 59.15 2.373 0.7371 12.8790 3.5887
16 t 118.06 59.03 2.373 0.7352 12.9130 3.5935 0.1603
16 s 118.36 59.18 2.373 0.7375 12.8716 3.5877
16 t 118.13 59.07 2.373 0.7358 12.9026 3.5920 0.0267
16 s 118.08 59.04 2.373 0.7354 12.9095 3.5930
~O16 t 118.10 59.05 2.373 0.7355 12.9074 3.5927 0.0236
S16 s 118.05 59.03 2.373 0.7351 12.9135 3.5935
C 16 t 118.10 59.05 2.373 0.7355 12.9076 3.5927 0.0230
r( 16 s 118.05 59.03 2.373 0.7351 12.9135 3.5935
16 t 118.13 59.06 2.373 0.7357 12.9034 3.5921 0.0277
16 s 118.08 59.04 2.373 0.7353 12.9105 3.5931
16 t 118.06 59.03 2.373 0.7352 12.9126 3.5934 0.0100
16 s 118.04 59.02 2.373 0.7351 12.9151 3.5938
16 t 118.06 59.03 2.373 0.7352 12.9126 3.5934 0.0100
16 s 118.04 59.02 2.373 0.7351 12.9151 3.5938


215











Table B-6. PWA 1484 (004) peak lattice misfit calculations.
Ir+k +12 tall/short? 29 9 h2 sin 9 a2 a 6
16 t 118.33 59.16 2.373 0.7373 12.8765 3.5884 0.2839
Cc 16 s 118.88 59.44 2.373 0.7415 12.8036 3.5782








09224




09225


118.43 59.22 2.373
118.13 59.06 2.373
118.50 59.25 2.373
118.47 59.23 2.373
118.52 59.26 2.373
118.89 59.44 2.373
118.50 59.25 2.373
118.76 59.38 2.373
118.49 59.24 2.373
118.96 59.48 2.373
118.32 59.16 2.373
118.51 59.25 2.373
118.30 59.15 2.373
118.48 59.24 2.373
118.26 59.13 2.373
118.47 59.23 2.373
118.24 59.12 2.373
118.46 59.23 2.373
118.45 59.22 2.373
118.01 59.01 2.373
118.25 59.12 2.373
118.46 59.23 2.373


0.7381 12.8623 3.5864
0.7357 12.9037 3.5922
0.7385 12.8541 3.5853
0.7383 12.8579 3.5858
0.7387 12.8508 3.5848
0.7415 12.8021 3.5780
0.7386 12.8539 3.5852
0.7405 12.8193 3.5804
0.7385 12.8549 3.5854
0.7421 12.7922 3.5766
0.7372 12.8771 3.5885
0.7386 12.8527 3.5851
0.7371 12.8797 3.5888
0.7385 12.8556 3.5855
0.7367 12.8861 3.5897
0.7383 12.8576 3.5857
0.7366 12.8888 3.5901
0.7382 12.8593 3.5860
0.7382 12.8607 3.5862
0.7348 12.9192 3.5943
0.7366 12.8876 3.5899
0.7383 12.8588 3.5859


-0.1606

-0.0145

0.1898

0.1346

0.2445

-0.0952

-0.0936


-0.1109

-0.1146

-0.2271

0.1120


216





SProfile Fitting Repor


;[091134102.MDI] Scan Data
SCAN: 49 0/52 99/0 01/1 (sec), Cu, i max r-99055, 01/29108 1 1:09a
Residual Error of Fit = 6.28% (Two-Tnefa Range of Fit = 49.0/52.99), 27101=0 0j(deg)
#l~--- 2 :-Thela diA) Arealall Area% Skew
1 50 78310 0231 1 7964 ((0 00151 7164 8 (288 21 44 1 0.2d6 0 6:
2? 50 787 10 0021 1 7962 (0 00011 16236 7 258 51 100 0 0.179 0 2;
Total Area = 23401 4 1387 1)


FWHM XS(A)
39 (O 009) 139 131
22 t0 003) 443 iT)


49.0 49.5 50.0 50.5 51.0 51 5 52.0 52.5

Twvo-Theta (deg)
Figure B-1. XRD deconvolution of PWA 1480 (4hr. 10800C, skewed, 6.28% error)


217











( 09113402,MDI] Scan Data Proflie Fitting Report
SCAN: 47 OiS2 96/0 02/5 Residual Error of Fit = 14.81% (Two-Theta Range of Fit= 47 0/52.98), 2T1OJ=0 O(deg)
# 2-Theta dtA) Area(ali Area%~ Skew FWrlHM XS(AI
1 50 858 0(O 14 1 7;939(0 00091 24339 5 33F3 7) 35 1 0 000 0 289 (0 033) 324 138)
21 50.815 (0.004) 1 7953(0 00031 69440 713508 4) 100 0 -0 755 0 223 r0 009) 442 120r
Total Area =93780.2 (4860.4)


47 48 49 50 51 52
Two-Theta (deg)

Figure B-2. XRD deconvolution of PWA 1480 (4hr. 10800C, skewed, 14.81% error)


218











[o91134X)2.MDI Scan Data Profle Fitting Report
SCAN: 49 0/52 99/0 01/11sec1 Cu, lunax)=990j55 01/29/08 11 09a
Residual Error of Fit= 6 63% (Two-Theta Range of Fit = 49 0/52.99), 2T(0)=0 0(deg)
#i 2-Theta drA) Area(al) AreaB Skew FWHM X8(A)
1 50 "23 (0 006) 1 7984 10 0003) 6136 2 (215 31 35 S 0 000l 0 692 (0 009) 128 (31
2 50 800(0 001) 1 795810 0001) 17298 3(200 21 100 0 0 000 0 233 (0 0021 419 (51
Total Area =23434.5 (294.0)


49.0 49.5 50 0 50.5 51.0 51.5 52 0 52.5
Two-Theta (deg)

Figure B-3. XRD deconvolution of PWA 1480 (4hr. 10800C, unskewed, 6.63% error)


219











[ 09113-002.L~M DI) Scan Dat a~~ ~- Profile Fitting Report
SCAN. 47 0/52 98/0 02151secl. Cu, llmawr=462488. 01/09108 10:06a
Residual Error of Fit = 12 58%/ ( Two-Theta Range of FR = 47 0/52 98), 2T(0)=0 0(deg)
# 2-Trheta dtA) Arealall AreaE( Skew FWHM XS(A)
1 50 846 (0 001) 1 7943(0 00011 66332.5 (1857.3) 100.0 0.000 0.175 (0 004) 1 (5
2 50 841 (0.006) 1.7945 (0.0004) 28406.7 (2008.6) 42.8 0,000 0.523 (0 012) 171(5)
Total Area = 94739.2 (2735 7)


47 48 49 50 51 52
Two-Theta (deg)

Figure B-4. XRD deconvolution of PWA 1480 (4hr. 10800C, unskewed, 12.58% error)


220












SCAN: 49 0/52 99/0 01/5(sec), Cu, lrmax1=452401 0109/08 10:46a
Residual Error of FIt = 7 94%6 (Two-Theta Range of Fit = 49 OtS2 991. 2T(0)=0 Oldegl


I~ofleitig Report


Total Area = 95833 1 (1787 9)


49.0 49.5 SO.0 50.5 51.0 51.5 52.0 52.5
Twvo-Theta (deg)

Figure B-5. XRD deconvolution of PWA 1480 (10hr. 10800C, unskewed, 7.94% error)











[0911rs 02.DI] San ~~ Da Profile Fitting Report
SCAN: 49 0/52 4990 01/Selsec,. Cu, l(max)=452401. 01/09108 10:46a
ieResidual Error of Fit = 5 63% iTwo-~Theta Range of Fit 49 0/52 99), 2T(0)=0.0(deg)
# 2-thela d(Al Area(al) Area% Skew FWVHM XS(A)
1 50 791 0 0011 1 7961 (0 00011 59656 2(787 7) 100 0 -0199 0 171 r0 002) 632 481
21 I 0 893(0 0071 179;8 (0 00051 36588 711016 2, 61 3 0.495 0 494 (0 005, 182 (3)


49.0 49.5 50.0 50.5 51 0 51 5 52.0 52 5
Two-Theta (deg)

Figure B-6. XRD deconvolution of PWA 1480 (10hr. 10800C, skewed, 5.63% error)


222











[09115-002.MDI] Scan Data Profile Fitting Report
S~CAN: 49 0/52 99/ 01/5(sec), Cu, Itmax)=935 01/09/0811 28a
SResidiual Error of Fit = 6 94%6 (Two-Theta Range of Frt= 49.0/52.99), 2T(0)=0 0(deg)
#I 2-Theta d(A, Arealal ) Area% Skew FWHMXSA
1 50 528 (0 022) 1 8049 (0 00141 86 1 (4 4) ;3 3 0 000 0 941 (0 0311 94 (4)
2 50 734 (0 002) 1 7980 rO 0001, 1 17 S 3 6) 1000 0.000 0 230 (0 006, 425 1121
Total Area = 203 6 (5 Til


49.0 49;.5 50.0 50.5 51.0 51.5 52.0 52.5
Two-Theta (deg)

Figure B-7. XRD deconvolution of PWA 1480 (100hr. 10800C, unskewed, 6.94% error)


223










[09115-4)02.MOIg scan Data Po~ uigRpr
SCAN: 49 0/52.9910 01/5(sec). Cu, l(maxl=936 01/09108 11 28a1
1sdsmnResidual Error of Fit = 6.82% !Two-Thera Range of Fit = 49 0152 99), 2T(0)=0.0(deg)
# 2-Theta dlA) Areacal) Area% Sitew FWHM XS(AI
1 I 50 66510 0761 1 8003(0 0051r 93 8 (7 41 85 5 0407 0 915(0 0541 97 (7)
2 ii:50 741 .0 004r 1.7978 (O 00031 10 7 14 ) 100 0 0 152 0 219 (0 008) 450 (171
Total Area =203 5 (8.7)


49.0 49.5 50.0 50.5 51.0 51.5 52.0
Two-Theta (deg)

Figure B-8. XRD deconvolution of PWA 1480 (100hr. 10800C, skewed,


52.5


6.82% error)


224


ICI~Jkr~3~Y~bYI~











[ 09116-002b.MDI] Scan Data
SCAN: 49,0/51 9910 01 3(sec). Cu. Ilroaxl=24020. 01/14/08 12:29p
Residual Error of Fit = 5.20% (Two-Theta Range of Fit = 49 0/51 99), 2T(0)=0l Olaegr


P~rofi Fin eortl


14,4 4 153 11 35 6
4140 1 45 11 100 0


49 0 40 5 50.0 50.5 51,0 51,5

Two-Theta (deg)
Figure B-9. XRD deconvolution of PWA 1480 (100hr. 10800C, unskewed, 5.20% error)


225











[0916-00a.MDI] Scan Data Profile Fitting Report
SCAN: 49 0/51 98/0 02/1(sec), Cu, I(max)=7828, 01/14108 12.12p
Resdual Error of Fit = 5 93%4 (Two-Theta Range of Fit= 49.0/51 98), 2Tl01=0 Otdeg1
# 2-Theta d(A) Area(al) Areab Skew FWHM XS(A
1 I/_ 50636 (O 008] 1 80313 (0 0005, 515 4 130 01 39 0 0.000 0 641 10 Old; 139 (4)
2 r1 50 71040 001) 1 7988 10 0001) 1322 812501 100 0 0 000 0 232 !0 004) 419 (81
Total Area = 1838.2 (39.0)


).o 49.5 50.(


Figure B-10. XRD deconvolution of PWA 1480 (1000hr. 10800C, unskewed, 5.93% error)


226


Two-Theta (deg)











[01602bMI can Deain~-~~-- Profile Fitting Rpo
SCAN: 49 0/51 99/O 01/3secl Cu, I(rnax)=24020. 01/41408 12 29p
Residual Error of F1t= 4.61%/ (TwoTheta Range of Fit=; 49 0/51 991 2Tl01=0 Otdeg)
#'~~ ~ 2-Thera a(A() Arealall Areai' SKew FWHM XS(A)
1 50.765 (0.016) 1.7970 10 0011) 1613.7 (59.5) 40 3 0.486 0.658 (0.011) 135 (3r
2 50 :19 10 00121 1 7985 10 0001 r 4006 3 148 9) 100 0 0 227 0 232 (0 0021 421 (51
Total Area = 5620.0 (77.0)


49.0 49~.5 50.0 50 5 51.0 51.5
Two-Theta (deg)

Figure B-11i. XRD deconvolution of PWA 1480 (1000hr. 10800C, skewed, 4.61% error)


227











[01164)0a.MDI Scan Data Profile Fittng Re~port
SCAN. 49 0/51 98/0 02/1 (sec), Cu, I(max)=7828, 0114/08 12:12p
Residual Error of Fit = 5 45%6 (Two-Theta Range of Fit = 49.0/51 98), 2T(0)=0 Oldegl
I# 2-Theta d(A) Area(all Areal Skew FWHM XS(A)
1 50.760 (0.027) 1.7972 10 0018) 546 9 (34 2) 42 2 0 466 0 645 (0 0171 139 (5)
'2 50 722 (O003r 1 7984(0 0002) 129r4 4281) 100 0 0 266 0 228 (O 004r 429 19)
Total Area = 1841.2 (44.3)


49.0 49.5 50.0 50 5 51.0 51.5
TwEo-Theta (deg)

Figure B-12. XRD deconvolution of PWA 1480 (1000hr. 10800C, skewed, 5.45% error)


228











[O913404.M110 scan Data Profle Ftthng Report
SCAN: 117 0/119 99/0 01/1 (sec), Cu, Ilmaw)=4727, 01/29/08 11 21a
Residual Error of Fit = 5 58%/i (Two- Thera Range of Fit = 117 Gil 19 99), 2T(0)=0.0(deg)
#I 2-Theta d(A) Area(al) Area% Skew FWHM XS(A)
1 '118 00 (0 0021 0 8986 (0 0000) 1849 1 1o199 100 0 0.000 0 450 (O 006) 351 (4 I
2 1' 18 533 !0 014) 0 8962 (0 00011 273 9 11r7 3) 14 8 0.000 0 442 (0 024) 361 (21)
Total Area =2123.0 (26.4)


117.0 117 5 118.0 1 TO 5 119.0 119.5

Two-T;heta (deg)

Figure B-13. XRD deconvolution of PWA 1480 (4hr. 10800C, unskewed, 5.58% error)


229











[09)113M.004MI] Scan Data Promie Fitting Report
SCAN: 117 01119 99/0 0115(secr, Cu, Icmaxl= 1 74699. 01/09/08 11:05a
Restdual Error of Fit = 7.77%/ (Two-Theta Range of Fit = 117 0,11 99 g), 2T(O =0 0(degl
# 2-Theta d(A) Area(al) Areag Skew FWNHM XS(A)
1 118 035 (0 002) 0.8985 10 0000) 75146 6 (909 3) 100 0 0.000 0.474 !0 004 1 333 (4)
2 118 586(0 015r 0 8959 (0 00011 8973 2 (608 81 11 9 0.000 0 431 10 024 I 37i1 21r
Total Area = 84119.9 (1094.3)


117.0 117.5 118.0 118.5 119.0 119.5
Two-Theta (deg)

Figure B-14. XRD deconvolution of PWA 1480 (4hr. 10800C, unskewed, 7.77% error)


230













SCAN: 117 01119 9910 01/1tsec). Cu, l(max1-4727, 01729/08 11 21a
Residual Error of Fit = 4 99%0 (Two-Theta Range of Fit = 117 0,1119 991. 2T(0) 0 0(deg)
#I 2 Theta d(A) Areatal) Area% Skew FWHM X(A
I1 l 118 022 (0 003) 0 8986 (0 00001 1846 4145 2) 100 0 0 113 0 447 (O 011) 354 (10)
2 r 18 d69 (0 010) 0 8965 (0 0001I 289 3 (48.0) 15 7 -0 479 0 417 (D 068) 384 (64)
Total Area 2135 7 (66 0!


Proftie Fitting Report|


[09113-004MDJI] Scan Data


) 118.5 119.0 119 5

Twvo-Theta (deg)


117 5


Figure B-15. XRD deconvolution of PWA 1480 (4hr. 10800C, skewed, 4.99% error)












SCAN: 117 0/119 4990 0115(sec), Cu, I(maxr)= 174699 01/09108 11 05a
Residual Error of Fit = 1 70%~ (Two-Theta Range of Fit = 117 0/1 19 991. 2T(0)=0.0(deg)
# 2-'Theta d(A) Area(al1) Aree%6 Sk~ew FWJHM _XS(A)
S1 118 042 (0 0011 0 8985 (0 00001 40223 2 (413 6) 84 5 0 066 0 350 (0 0021 460 (41
2 1118 0098(0 006) 0 8982 (0 00011 47591.2 (1050.7) 100.0 5)095 0.862 (0.010) 180 (3)
Total Area = 87814 5 (1129 2r


SProllie Fitting Report


[Og13)04.MD] Scan Data


117.0 197 5 118.0 198.5 119.0
Two-Theta (deg)


119 5


Figure B-16. XRD deconvolution of PWA 1480 (4hr. 10800C, skewed, 1.70% error)


232











[09114-00tMDI] Scan Data Profle Fitting Report
SCAN: 117.0/119,99.10 015(sec), Cu, grna1i=260 171, 01/09/08 11:35a
Residual Error of Fit = 5.79%o ( Two-Thera Range of Fit = 117 0,1 19 99), 2T(Op=0 Oldeg)
# ~2-Theta di8A) Area~al) Areab Skew FWHM XC~
:1 117 971 (0003) 0 8988 10 0000) 105119 1 (20590) 100O 0 .000 0 464 10 0014) 340 (4r
'2 118 37110 041r 0896910 0004) 29535 8(3354 3) 28 1 0 000 0 761 (0 0311 206 (19
Total Area =134654.9 (3935.8)


117,0 117.5 118.0 118.5 119.0 119.5
Two-Theta (deg)

Figure B-17. XRD deconvolution of PWA 1480 (10hr. 10800C, unskewed, 5.79% error)


233











[09114.004 MDI] Scan Data Profile Fitting Rpr
SCAN: 117 0/11S999O10 011(see), Cu, (max)=260171. 01/09/08 11 35a
Residual Error of Fit= 1 41% (Two-Theta Rangle of Fit= 1 17 0,1 19 99), 2T(0)=0 0(deg)
#/ 2-Theta diAi Areaiali Area% Skew FWHM XS(A)
1 117 984 (O 001) 0 8987 (0 0000) 70017 1 4716 Sp 100 0 0 038 0 387 r0 002) 413 13
2 118 083 10 00; 0 8983 (0 0001) a6850 3(1828 41 98 3 -0 126 0 900 (0 0131 173 13)
~Total Area = 1388637.5 1963 8 I


7 17.0 117.5 118.0 118.5 119.0 119 S
Two-Theta (deg)

Figure B-18. XRD deconvolution of PWA 1480 (10hr. 10800C, skewed, 1.41% error)


234










[09116-004la.MDI] Scan Data Profile Fitting Report
SCAN: 117 01119 98/0 02/1(sec), Cu, .(rnax)=25 10, 01/14/0N8 10:12a
SResidual Error of Fit = 3 17% (Two-Theta Range of Fitl =T 01 l 19 98).~ 2T(0)-0 0(deg)
/# 2-Theta d(A) Area~all Areab Skew FWVHM XS(A)
S1 1~17 856 (0 027) 0 6993 (0 0003! 562 & 03 58 5 0 00 0.796 10 0661 195 (17)
2 117 997 0O 108) 0 8987 (0 0)010, 9617 01 100 0 0 000 1 020 !0 120) 152 (19)
Total Area = 1524 1 10 01


1 97 0 117.5 118.0 118 5 119.0 119.5
Two-Theta (deg)

Figure B-19. XRD deconvolution of PWA 1480 (1000hr. 10800C, unskewed, 3.17% error)


235










































































117.0 117.5 118.0 118 5 119.0 119.5

Two-Theta (deg)

Figure B-20. XRD deconvolution of PWA 1480 (1000hr. 10800C, unskewed, 1.77% error)


236


i_ ~1


SCAN, 117 0/119 99.10 01/(sec), Cu, I(max)=7061 01/14/08 10:08a
Residual Error of Fit= 1 77% (Two-Theta Range of Frt = 117.0/119 99) 2TL0)=0 0(degi
#~~ 2-Theta d(A) Area(al) Area%~ Skew


I


r


1 17 778 r0 0191 D 8997 10 00021 20:9 3 (261i 2) 88 4 0 000 0 807 (0 0191 192 161
118.0901 O0431 0.8982 (0 00041 2351 6 364 8i 100 0 0 000 0 945 (0 038) `164 (8)


Profle Fittng Reportl


[O9116-004b.MDI] Scan Data


WHM XS(A4)


Total Area = 4430 9 (448 6r


















































































237


SCAN: 117 011 19 9910 01/3 1secl Cu, Isr ax)=7061, 01/14/08 10 08a
Reeldual Error of Fit = 1 75%~ (Two-Theta Range of Fit = 117 0,1119 991 2T(0)=0 0(dteg)
# 2~~!Tht -The d(A)I Area(al) Area% Skeaw FV


I


I


117 827(0 0131 0 8996 (0 00011 2371 6(370 6) 100 0 0000 0 851 r0 017) 182 (5)
118 070 1O 0611 0 8983 (0 00061 2114 8 (479 9, 69 2 0 000 1 013 r0 052) 153 491


1Profle Fittlng Report


[091164-04b.MDI] Scan Data


WHM KS(A)


Total Area =4486.4 (606.4)


117.0 117.5 1980 118.511019.

Two-Theta (deg)

Figure B-21. XRD deconvolution of PWA 1480 (1000hr. 10800C, unskewed, 1.75% error)












SCAN: 117 0/119 98/0 02/1(sec). Cu. Ilmaxl=2510 01/14/08 10:12a
SResidual Error of Fi= 3 14% (Tw*o Th~etRage of Fit= 117 0119 l981. 2T1(0)=0.0(deg
# 2-Thete dlA) Area(al) AreaC Skewl FWHM XS(A)
1 117 859 !0 068r 0 8993 (0 0D006 578 7 ID, 61 2 -0 100 0 824 r0 1821 188 (42)
2 I;~117.962 (0 163) 0 8988 (0 00151 945 1 17 100 0 -0 OS6 1 016 (0.202) 162 (31)
Total Area = ~523 8 (0 01


09 6404a.MDDI] Scan Data


IProfle Fiting Reportl


117,0 117.5 118 0 178.5 119.0 119.5
Twvo-Theta (deg)

Figure B-22. XRD deconvolution of PWA 1480 (1000hr. 10800C, skewed, 3.14% error)


238











[091164504b.~MDI] Scaln ata~- --- Profile F ith~ng Repor
ISCAN 117 0"119 9/0 01/3(ecl Cu, I(max)=7061l, 01/14/08 10 08a
Residual Error of FRt= 1 71%/ (Two-Thera Range of Fit= 117 0,1 19 99), 2T(0:=0 Oldeg,
7----------- 2-Tht atA) Area(ali Areab Skew FWHM X(A
1 117 826 10028; 0 8995 (0 0003) 2385 3 (76 01 100 0 -0157 0 873 (0 107; 17(3
2 118.027 (0.065) 0 8965 (O 0006) 2086 6 (996 1) 87 5 0 027 1 035 40 0981 150 r15)
Total Area = 4471.8 (1262 7)


117.0 117.5 178.0 1185 119.0 199.5
Two-Theta (deg)

Figure B-23. XRD deconvolution of PWA 1480 (1000hr. 10800C, skewed, 1.71% error)


239











[09116404b MDil Scan Data Profile Fitting Report
SCAN: 117 01119 9910 0~13(sec), Cu, l(mrnax=7061. 01/14108 10:08a
Reeadual Error of Fit = 1 71% (Two-Theta Range of Fit = 117 011 19 99), 2T(01=0 0(degl
#i 2 -Thata ~ ~ d(Al Area(al) Area% Skew FWI-W XS(A)
1 !~117 824 (0026) 0 8995 (0 0002) 2380.5 (751 1) 100.0 -0.130 0 869 (0.108) 178 (231
;2 118 037 t0 062) 0 8985 (0 0006) 20392 5 (958 61 87 9 0.033 1 032 (0 0981 150 115)
Total Area = 4473 0 (1217 81


117.0 117.5 118.0 118.5 199.0 119,5
Two-Theta (deg)

Figure B-24. XRD deconvolution of PWA 1480 (1000hr. 10800C, skewed, 1.71% error)


240











[ 09223-002.MDI] Scan Data Profte Fitting Reporrtl
SCAN: 47 0/52 9810 02/5(seel Cu, l(rnax)=790242, 01/09/08 10 05a
Residual Error of Fit= 13 36% (Two-Theta Riange of Fit= 47 0/52 98), 2T(0)=0 0(deg1)
# 2-Theta dlA) Area(all AreaB Skew FWHM XS(A)
1 50 702 (0 0141 1 7991 to00091 37146 8(2652 21 27 3 0.000 0 669 t0 0181 133 15)
2 50 826 (0002 1 7i950 (00001 136201 9(3248 1) 100 0 0000 0 213(0 034 i 468 001
Total Area = '173348 6 (4193 31


47 48 4)9 SO1 51 52
Twco-Theta (deg)

Figure B-25. XRD deconvolution of PWA 1484 (4hr. 10800C, unskewed, 13.36% error)











i


1Pr le fitting Repor


[09~g22-002a.MDI) S~can Data
SCAN. 49 0/52 98/0 02/1 (sec). Cu, 1(rnax)=14422, 01/14/08 10:07a
Residual Error of Fit = 6 .30%1 ('Two-Theta Range of Fit = 49.0/52.98), 2T(0)=0.0(dleg)


FWHM


Foral Area = 4326 8 (80 81


49.0 49.5 50.0 50,5 51 0 51.5 52.0 52.5

Two-Theta (deg)

Figure B-26. XRD deconvolution of PWA 1484 (1000hr. 10800C, unskewed, 6.30% error)


242











P Gk Ftng Rep


i


[D9224-002b.MDI] Scan Data
SCAN, 49.0/51 9910 01131secr Cu, I(rnax)=41'11, 01/14/08 01 42p
Residual Error of Fit= 4.61%8 (Two-Theta Range of Fa = 49.0/51.99), 2T(0 ;=0 O( aeg!
#` 2-Theta d(A) Area(al) Area% Skew

2:: I 50.38 ( 02 1 8012(00011 746 1011 On 9100 0 m 0.00 0 40


FWHM XS(A)
10 010) 93 12)
(0 004) 271 (41


(20.2)


40.5


50 0 50.5
Tlwo-Theta (deg)


Figure B-27. XRD deconvolution of PWA 1484 (10hr. 10800C, unskewed, 4.61% error)


243











[092264102b MDoj Scan Data Profile Fitting Report
SCAN: 49 0/51 .99/0 01/3(sec). Cu. itrnaxl=44143 01/14108 10:24a
I Reeldual Error of Fit= 5 i% (Two-Theta Range of Fit = 49.0/51.99), 2110I=0 Olaegi
I#1 2-Theta dlA) Area(al) Area% Skew FWlHM XS(A4)
1 i:50 626 (0 006) 1 8016 (0 00041 4322 3 1~6r 8! 49 8 0.000 0 703 (0 008) 126 (2)
S50 734 (0 0011 1 7980 ro 00011 B668 9 1r8 6) 100 00.000 0 305 (0 003) 305 14)
Total Area =13001.3 :200 6)


49 0 49.5 50 0 50 5 51.0 51.5
TwNo-Theta (deg)

Figure B-28. XRD deconvolution of PWA 1484 (1000hr. 10800C, unskewed, 5.85% error)


244












SSCAN: 49 0151 996/0 004t3(sec), Cu. imrnax=4232. 01/14/08 02 22p3
SResidual Error of Fit = 4 45%A Tw~o-Theta Range of Fit = 49.0/51 996). 2T(0)=0 0(deg)
#' 2-Theta d(A) Areacall Area%~ Skew FWVHM XS(A)
1 50 50)3 (0 0041 1 805- 0 0002, 755 2 !10 71 99 0 0.000 0 95500 006) 92 1
2j 50 634 1o 001~ 1 8013 (0 0001, 762 516 8) 100 0 0.000 0 341 (0 0021 270 131
Total Area =1517 7 (12 7)


FrfieFittng creportl


[9224-002c~.MDI) Scan Data


49 0 ~ ;49;.5 50.0 50 5 51.0 51-5
Two-Theta (deg)

Figure B-29. XRD deconvolution of PWA 1484 (10hr. 10800C, unskewed, 4.45% error)


245











[0224)2.DI] a aScan Data P~RofeFltmg Reor
SCANU 47 Oi52 98i0 0r2/5tsecr. Cu, I(maxr=790242. 01/09108 10 05a
IResidualError of Fnt 171% (Two-Theta Range ofFit=47.0/52.98), T0=0Oog
#I 2-Theta d(A) Area(al l Area%~ Skew FWI-IM XS(A)
1l~ 50 881 (0 049r 1 7932 (0 0032) 30441 2 2490 7) 21 0 0 630 0 864 (0 043) 103 (6)
2 50 862 003r 1 7938 40 0002r 144712 7 3193 9) 100 0 0 566 0 224 (0 004) 438 (91
Total Area = 175154 0 (4050 3,


47 48 49 50 51 52
Two-Theta (deg)

Figure B-30. XRD deconvolution of PWA 1484 (4hr. 10800C, skewed, 1 1.71% error)


246











[OS 6-002a.M0f] Scan Data
SCAN: 49.0/52 98/0 02/1(sec), Cu, I~max)=14422, 01/14/08 10 07a
Residual Error of Fit = 3 71%u 1Two-Theta Range of F1t= 49 0/52 981. 2T(0)=0 0(deg)
#; 2-Theta d(A) Area(ai) Area%4 Skew
1 50 693 (0 024) 1 7994 (0 0016) 1049 5 (41 1) 31 7 0 171 0 8'
I 2l 50 771 (0.002) 1.7968 (0 00011 3307 6131 71 100 0 0.454 0 3.
Total Area = 4357.1 (51,9)


Prle Fittin g R epor


49.0 49 5 50.0 50.5 51 0 51,5 52.0 52.5
Twvo-Theta (deg)

Figure B-31i. XRD deconvolution of PWA 1484 (1000hr. 10800C, skewed, 3.71% error)


247











S[09226-002b.MDOI] Scan Data Profile Fitting Report
SCAN' 49.0/51 9910 01/3lsecl. Cu, itmakl=44143, 01/14108 10 24a
Residual Error of Fit = 3.20% (Two-Theta Range of Fit = 49 0/51 99), 2TIO)-0 Otdegr
# ~ 2-The~ta d(A) Area(al) Area% Skew FWHM XS(AI)
1 50 747 (0 016) 1 7976 (0 001r1I 3452 3 (102 41 35 7 0 353 0 823 (0 0101 108 (2)
2 1: 50.773 (0 0011 1.7967 (0 00011 98:670476 01 100 0 0 50B 0 328 (0 0021 282 (31
Total Area =13122.7 (127.5)


49.0 49.5 50 0 50.5 51.0 51 5
Twvo-Theta (deg)

Figure B-32. XRD deconvolution of PWA 1484 (1000hr. 10800C, skewed, 3.20% error)


248





j


ProfAe Fitting Report


[09224-002b.MDI] Scan Data
SCAN: 49 0/51 .99/O 01/(sec), Cu, Itrnax)=4111. 01/14/08 01 42p
Residual Error of Fit = 4 07% (Two-Theta Range of Fit = 49.0/51 99), 2T(0)=0 0(deg)
#i 2Theta d(A Area(al) Areab SkewF F
1 S 553 (0 0251 1 80d0 (0 0017) 736 5 120 8) 96 5 0 146 0 987
2 1/i50 6641 (0 003) 1 8003 (0 00021 763 4113 21 100 0 0.344 0 348
Total Area = 1499 9 124 4)


49.0 49 5 50,0 50.5 51.0 51.5

Two-Theta (deg)

Figure B-33. XRD deconvolution of PWA 1484 (10hr. 10800C, skewed, 4.07% error)


249












SCAN: 49 0/51 99610 004/3(sec), Cu, I(rnax)=4232, 01/14/08 02 22p
Residual Error of Fit = 4 03%6 (Twvo-Theta Range of Fit = 49 Oi51 996). 2T(0)=0 0(deg)
# 2-Theta d(A) Area~all Area% Skew FWM XSA
1 50 555 40 015) 1 80e0 (O 00r10 754 5 112 8r 97 6 0 133 0 984 r0 0081 90 (21
2 50 658 r0 002) 1 8005 10 0001) 772 9 (O21 100 0 0.307 0 346 r0 003) 265 (31
Total Area = 1527 4 (15 2)


IProfile Fitting Report1


[0922~4402c.MDI] Scan Data


49 0 49.5 50 0 50 5 51.0 51.5
T~wo-Theta (deg)

Figure B-34. XRD deconvolution of PWA 1484 (10hr. 10800C, skewed, 4.03% error)


250











[09226-400a.MDI] Scan Data Profile Fitting Report
SCAN: 117 0,1119 98/0 02,1lisecl Cu, I(max)=366~6, 01/14/08 10:47a
Residual Error of Fit = 4 15"A (Two-Thata Range of Fit = 117 0,1119 98), 2Tl01=0 0(degl
# *----2-ThThet d(A)1 Area(al) Area% Skew FWHM XS(A)
1 118 256 (024r 0 8975 100002) 1155 7 60 0) 100 0 0.000 0 926 10 017) 168 td)
12 118 469 (0 003) 0 8965 r0 0000) 813 5 ld9 91 70 4 0 000 0 426 10 013) 375 (12r
Total Area = 1969.2 (78.0)


117.0 117.S 118.0 118.5 119.0 119.5

Two-Theta (deg)
Figure B-35. XRD deconvolution of PWA 1484 (1000hr. 10800C, unskewed, 4.15% error)











|09226-400b.MDI] Scan Data Profile Fitting Report
z~SCAN: 117 0,1119 9910 01731secr. Cu, l(max)=10386 01/14/08 11 05a
Residual Error of Fit = 3 24% (Two-Theta Range of Fit= 117 01119 99), 2T(0)=0.0(deg)
# 2-Thete d(A) Area(al l Area%~ Skewr FWHIM XS(A)
1 l 118 236 (0 0131 0 8975 (0 0001) 3281 2 (91 81100 0 0.000 0 916 (0 010; 170 (3)
2 1iC:i18 456 (0 002) 0 8965 <0 00001 2301 9 177 7) 70 2 0.000 0 418 (0 007) 382 17,
Total Area =5583 2 (120 3)


117.0 117.5 118.0 118.5 119.0 119.5
TwNo-Theta (deg)

Figure B-36. XRD deconvolution of PWA 1484 (1000hr. 10800C, unskewed, 3.24% error)


252












SCAN: 117 011 19 9ul0 01t3(secl, Cu. !(raxl=10386. 01/14/08 11:05a
Residual Error of Fit = 4 35%4 (Two-Theta Range of Fit = 117 0/11 9.99), 2T(01=0.0(deg)
#i 2~~-Teta ~ d(A) ~ Area(al) A~rea% Skew FWI-M XS(A)
1 Irj118 011 (0 021) 0 8986 (0 00021 1648 9 (84 31 43 1 0.000 0 757 (0 037) 206 111;
2 118 446 (0 003r 0 8966 (0 0000) 3826 3(62 0r 100 0 0.000 0 501 (0 00~4) 316 14;
Total Area = 5475 2 (117.6)


[09226400b.MDI] Scan Data


PrfleFttn Report


117.0 117.5 1180 118.5 11990 1195
Two-Theta (deg)

Figure B-37. XRD deconvolution of PWA 1484 (1000hr. 10800C, unskewed, 4.35% error)


253












SCAN: 117 0,1 19 /0 02/1ssecr, Cu, Imaxi=36j6. 01/1408 10.47a
Residual Error of Fit = 3 09%k (Two-Theta Range of Fit = 117 0,1119 98r. 2T(0)=0.0(deg)
# 2-Theta d(A)~ Area(al) Area% Skew FWVHM _XS(A)
1 118 323 (0 0361 0 8971 (0 00031 1088 4173 61 100 0 0 1168 1 062 (0 03di 147 (61
2 118 506 (0 0041 0 8963 r0 0000, 932 6 (38 Of 85 7 0.547 0 463 (0 009) 343 (71
Total Area =2021,1(82.9)


[09226400a.MDI] Scan Data


SProfle Fitting Report


117.0 117.5 118 0 116.5 t 19.0 119.5
Tw~o-Theta (deg)

Figure B-38. XRD deconvolution of PWA 1484 (1000hr. 10800C, skewed, 3.09% error)


254











[O9226400b.MDI] ScanR Data Profile Fitting Report
SCAN: 117 0/119999/0 01 3(sec), Cu, I(max)=10386,. 01/11 40811 05a
Rtesidlual Error of Fit= 2 07%6 (Two- Theta Range of Fit = 117 0111 ~9 99), 2T(0)=0 0(deg)
#rt 2-Theta dA Area~al) Area% Skew FWHM X8(A)
1j 118 304 (0 015) 0 8972 (0 0001) 3224 1 (95 2) 100 0 0 121 1 007 (0 014) 155 (3)
2 118 484 10 0021 0 8964 (0 0000) 2456 9 r53 5) 76 2 0 465 0 443 (0 0031 360 (4)
Total Area = 5681.1 (109.1)


11. -117.5 1 10 1.110195
Two-Theta (deg)
Figure B-39. XRD deconvolution of PWA 1484 (1000hr. 10800C, skewed, 2.07% error)


255











[12~123-002.M B1~DI] Scan Data Proftie FIttng Re~port
SCAN: 49 0/52 9910 011(sec), Cu, I(max =6029, 01/09/08 02:21p
Res dual Error of Fit = 5.07%6 (Two-Theta Range of Fit = 49 0/52 991, 2T(0)=0.0(deg)
#/ 2-Theta d(A) Arealali Area%/ Sktew FWHM XS(A)
1 50 693 (0 001 ) 1 7994 (0 000 1) 1032 8(10 9) 100 0 0.000 0 251 (0 002) 381 (41
2 50 656 (0 005) 1 8006 (0 0003r 429 4 (14 31 41 6 0.000 0 827 !0 013) 107 (31
Total Area = 1462.2 (17.9)


49.0 49.5 50.0 50.5 51.0- 515 2 5

Twvo-Theta (deg)

Figure B-40. XRD deconvolution of PWA 1480+ (4hr. 10800C, unskewed, 5.07% error)


256











[121234502 b.MD I Scan Data Profil Fitting Report
SCAN: 49 0/51 98/0 0211(seei, Cu, l(max)=8585, 01/14/08 12 58p
/Residual Error ofFit = 10 42% (Two-Thete Range of Fit= 4 051 98), TI O)=0 (deg)
# 2-Theta d(A) Area(a'l) Area%~ Skew FWHM XS(A)
1 50 747 (0 001' 1 79-'6 10 0001) 1067 5 29 9) 100 0 0.000 0 136 (0 0031 965 1221
2 50 772 (0 007 1 1 7968 (0 00051 376 3 (37 2) 35 2 0 000 0 d44 (0 0181 201 (91
Total Area = 1443.8 (47.7)


49 0 49.5 50 0 50 5 51.0 51.5
Two-Theta (deg)

Figure B-41. XRD deconvolution of PWA 1480+ (4hr. 10800C, unskewed, 10.42% error)


257











[12123402c.MD] Scan Data Poftle Fitting Report
SCAN: 49 0/51 .99/0 01/3(sec), Clr I(max)=30596, 01/14/08 01 16p
Residual Error of Fit = 10 33%A ('Two-Theta Range of Fit = 49 0/51 99), 2T(0)=0 0(dedg)
#/ 2-Theta d(A) Area(all Area% Skew FWHM XS(A)
S50 749 (0 001,r 1 7975 (0 0001 r 3980 3 (736) 100 0 O0 00 143 (0 002) 859 (13,
2 50 764 (0 006) 1 7970 (0 0004) 1149 2 (81 6) 28 9 O0090 0 485 (0 014) 185 (6r
Total Area = 5129 5 r 104 91


49.0 49.5 50.0 50.5 51.0 51.5

Twvo-Theta (deg)

Figure B-42. XRD deconvolution of PWA 1480+ (4hr. 10800C, unskewed, 10.33% error)


258











[~12124-002 1MD] Scan Daa Profile Fitting Rpr
SCAN: 49 0/52 9910 01/5(sec), Cu, I(max)=88971. 01/09108 01 42p
Residual Error of Fit = 9 55%6 (Two-Theta Range of Fit = 49 0/52 99), 2TO1=0 Oldega
#i 2-Theta d(A) Area(al) Areab Skew FWHIMXSA
1IC 50 727(0 001l 1 7982 (0 0000) 10966 B~r d86 4)00 0 0 000 0 137 r0 001) 935 r11)
2 I:: 50 698 (0 005) 1 7992 (0 0003) 3088 6 (133 6) 28 2 0 000 0 533 (0 010) 168 (4)
Total Area = 14055 5 (199 7)


49.0 49.5 50.0 50.5 51.0 51.5 52.0 52 5
Twvo-Theta (deg)

Figure B-43. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 9.55% error)


259











Profle Fitting Report


[1~2125-002a.MDI Scan Data
SCAN: 49 0/52 99/0 01lirsec,. Cu, r(marnexi 35. 01/29/8 ;10:04a
ResidualI Error of Fit = 5.73% (Two-Thela Range of Fit= 49 0/52 991. 2T(OI=0 0 deg;


= 63.7 (14.5)


49,0 49.5 50 0 SD 5 51 0 51.5 52.0 52 5
Two-Theta (deg)

Figure B-44. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 5.73% error)


260











---~ ~-TPronie Fitung Reportl


L


[1226-02b.D q scan Data
SCAN: 49 0/51 9910 01131secl, Cu, Ilmax1210223. 01/29/08 10:20a
Residual Error of Fbt= 4 87% (Two-Theta Range of Fit = 49 06j1 99), 2T(0)=0.0(deg)


ea(alr Areab Skew FWHM XS(A)
r29 6, 100 0 0.000 0 789 10 0101 1 r12 (2
(28 6, 90 3 0.000 0 212 i0 004) 470 (10)


49,0 49 .5 50 0 50.5 51.0 51,5

Two-Theta (deg)

Figure B-45. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 4.87% error)











SProfle Fittingl Report


[12126-002a MD1] Scan Data
SCAN: 49 0/$1 981/0 02/1(sec), Cu, Itmax1=92436. 01/14/08 11 11a
iResidual Error of Fit = 4 41%6 (Two-Thefa Range of Fit = 49 0151 98). 2T(0)=0 0(degd)
#/ 2-Theta d(A) Alrea(al) Areab Skew
11C 50 698 (0 005) 1 7992 (0 0003) 5907 7 1281 6) 38 2 0.000 0 51
2 50 761 !0 0011 1 79:1 (0 00011 15479 7 (223 9i 100 0 0.000 0 2:
Total Area = 21387 4 (350 8)


FWVHM XS(A)
59 r0 00~8) 16 1:3)
29 r0 003) 426 16)


49.0 49.5 50.0 50.5 51.0 51.5

Two-Theta (deg)
Figure B-46. XRD deconvolution of PWA 1480+ (1000hr. 10800C, unskewed, 4.41% error)


262












I 1


PraeFitting Report


[2126-002b.MDq Scan Deata
SCAN: 49.0/51 99/0O 1/3secl. Cu, l(maxy-288040, 01/14108 11 30a
Resildual Error of Fit = 4.37% (Two-Theta RangeofFnt= 49 0/51 991 2T(0;=0 0(celeg
#i 2-Thata d(A) Area(all Area%/ Skew
1 50 697(0 0031 1 7992(0 00021 18502 & (621 51 38 2 0.000 0
2 50 756 10 0011 1 79;2 (0 00001 48416 7(468 0) 100 0 0000 0
Total Area = 66919 1t 1790 21


FWHM XS(Al
557 (0 005! 161 !3!
230 (0 0021 424 515


40.0 49 5 50.0 50.5 51.0 51.5

Twvo-Theta (deg)

Figure B-47. XRD deconvolution of PWA 1480+ (1000hr. 10800C, unskewed, 4.37% error)


263












SCANrl 49 0/S2 99/0 01/5(sec), Cu, I(max)=6029. 01/09/08 02:21p
Residual Error of Fit = 5 OS%~ (Two-Theta Range of Fit = 49 OtS2 99). 2T(0)=0 0(degl
# 2-Theta d(A4) Area(al) AreaB SkewN FWHM XS(AI)
1 C 50 695 (0 0021 1.7993 (0.00011 102 0 (12 3) 100 0 0 034 0) 250 10 0021 384 1S)
2 1 50 686 (O0221 1.7996 40 00151 435.7 (14 8) 42 5 0 107 0 820 (0 013, 108 13)
Total Area =1461 6(19 2)


[r12123402.MDI] anSca Dta


Prfle Fitting Report


49.0 49.5 50.0 SOS5 51,0 51.5 52.0 52.5
Two-Theta (deg)
Figure B-48. XRD deconvolution of PWA 1480+ (4hr. 10800C, skewed, 5.05% error)


264












| Prome Fating Report


j


[ii21~n23-002b.MS can Data
SCAN: 49.0/51.98/0.02/1(sec), Cu, IlrnaK)=8585, 01/14/08 12 58p
Residual Error of Fit= 6.91%6 (Two- Thera Range of Fit = 49.0/51.98), 2T(0)=0.0(deg)
# 2-Theta (MA) Area(al) Area%, Skew
1 50 73010 001r 1 7982 (0 00011 1107 7 (201 100 0 -0 46 0 41
2 50 752 r0 027r 1 7974 10 0018) 34: d (24 3r 31 4 -0.026 0 469
Total Area = 1455 2 (32,1)


FWHM XS{AI
(0 0021 885 (141
10 0r13 192 (61


49.0 49.S 50.0 50.5 51.0 51.5

Two-Theta (deg)

Figure B-49. XRD deconvolution of PWA 1480+ (4hr. 10800C, skewed, 6.91% error)


265











[ 121234)02c;.MDI) Scan Data Profile Fitting Report
SCAN: 49 0/51 9910 0173(secl. Cu, l(rnax:,-30596 01/14/08 01 163p
SRe,,~i ~ssidul r~,rorofFt = 85 To-TeaRag fFI = 49.0/51 99), 2T(0)=0 0(deg)
# 2.Theta d(A) Area(all Area%/ Skew FNHM XS(A)
1Z 50.732 (0.001) 1,7981(0 0001 I 4039.1 (51.1) 100.0 -0.427 0.14540 0011 833 (9)
2 50.74510 020) 1 7977 !0 00131 1123 7(54 41 27 8 -0028 0 489 (0 009! 184 (4)
Total Area =5162.9 (74.7)


49 0


49.5


90.0 50 5
Two-Theta (deg)


Figure B-50. XRD deconvolution of PWA 1480+ (4hr. 10800C, skewed, 6.85% error)


266















































































Figure B-51i. XRD deconvolution of PWA 1480+ (10hr. 10800C, skewed, 9.54% error)


267


50 5 51.0 51.5

Two-Theta (deg)


SCAN: 49 0/52 9910 01 15(sec), Cu, I(rnax)=889371, 01/09/08 01 42p
Res8idual Error of Fit = 9 54% (Twho-Theta Range of Fit = 49 0/52.99), 2T(0)=0.0(deg)
# 2-Theta diAl Area(al) Areat Skrew FW~iVI XS(A)
1 50 727 (0 001) 1 7982(0 0001, 1096i3 152 61 100 0 0 010 0 137 r0 001) 936 (111
2 50 699 (O 021) 1 7992 (0 0014) 3091 8 (139 0) 28 2 0 005 0 533 (0 010) 168 (41
Total Area = 14055 0 (206 4;


_ _____


52 0 82 8


490 49. 5 50.0


[224-002.MDI| Scan Data


SProfile Fitting Report


V
















I,










[1215410acaMDI] Scan Dat roueik Fitting Report
SSCAN: 49 0/52 9910 01 /11sel. Cu, II sr av-34 15 01/29/08 10 04a
IResidual Error of Fit = 542%L (Two-Theta Range of Ft = 49 Oi5299r 2T()=0 (deg
# 2-Thata d(A) Area(al) Area%/ Skew FWHM XS(A)
1 50 673 10 0361 1 8000 (0 00124) 529 1(20 2, 100 0 0 480 0 757 (0 0211 117 (4]
21 50 74810 0031 1 7976 (00002i 331 9 (13 4) 62 7 -0 011 0 186 (0 0061 562 (19
Total Area = 861,0 (24.2)


49.0 49550.0 50.5 51.0 51.5 52.0 52.5
Two-Theta (deg)

Figure B-52. XRD deconvolution of PWA 1480+ (100hr. 10800C, skewed, 5.42% error)


268











S[12125002b.Mog] scan Data Promle Fitting Report
SCAN: 49 0/51 99/.0 0 1/31sec r Cu, It rnar)=10223, 01/29/108 10 20a
Residual Error of Fit = 4 13% (Two-Theta Range of Fit = 49 Oi51 99,, 2T(0)=0 0(deg)
# 2-The- diA) Area~all AreaX Skewu FW~I-M XS(A)
1 50 700 (O028) 1 7991 10 0019) 162 2 (54 1) 100 0 0 571 0 748 (0 019) 11r914)
2 50 755 (0 002r 1 7973 10 0001) 1000 8 (34 9) 61 5 0 134 0 184 (0 005) 570 r11)
Total Area =2628.0 (64.4)


49.0 49 S 50 0 50 5 51.0 51,5

Twvo-Theta (deg)

Figure B-53. XRD deconvolution of PWA 1480+ (100hr. 10800C, skewed, 4.13% error)


269











(121lo26-02a.DI]an~ ScnDaaP~rolle FtigR epor
SCAN: 49 0/51 98/0 01 1(sec), Cu, I(rax)=92436, 01/14/08 11 11a
Residual Error of Fit 4 30%6 (Twvo-Theta Range of FIl = 49 0/51 98), 2T(OI=0 0(deg)
#/ 2-Thera a(Ai Arealal) Area% Skew FWHM XS(A)
1 SO 741 (0 024) 1 79,8 10 0016) 6239 3 1359 81 41 2 0 181 0 554 10 009) 161 (4)
2 50 763 (0 002) 1 79"'0 10 0001, 15149 5 (287 81 100 0 0 074 0 227 10 004> 431 (8>
Total Area = 21388,8 (460.7)


51.0 51.5


Figure B-54. XRD deconvolution of PWA 1480+ (1000hr. 10800C, skewed, 4.30% error)


270


Two-Theta (deg)












I~


1PIroBfiitting Report


[12126-002bMDI] Scan Data
SCAN: 49 0/5' 9910 01/3(se r Cu, luna* r288040, 01/14/08 11:30a
~kB-~Residual Error of Fit = 4 28%~ ('wo-Thera Range of Fit = 49.0/51.99), 2T(0)-=0.0(deg)
#2-Theta d(Al Arealal, Aream Skew
1 SO 736(0 01;1 1 79-'9 0 00111 19504 5 1788 81 41 1 0.164 0
2 50.759 (0 001) '1 7972 (0.00011 47415 81629 9) 100 0 0.050 0
Total Area = 66920.4 (1009.5)


FWHM XS(A)
551 (0 006) 162 13r
228 (0 0021 429 161


49.0 49 S 50.0 50 5 51 0 51.5

Twvo-Theta (deg)

Figure B-55. XRD deconvolution of PWA 1480+ (1000hr. 10800C, skewed, 4.28% error)










SCAN: 117.0019 99,00115lsec r,Cu, Ilmaxl=2707,01/09/08 0127p
Realdual Error of Fit = 2 38% (Two-Theta Range of Fit=+ 117 0:1l19 991. 2T(0)=0.0(dieg)
#; 2-Theta d(Al Areafal) Area% Skew FWHM X(S(A)
1 li_ 117.977 (0 0141 086988\0 00011 All5125On 11000 0000 08531 (0029) 187181
2 1!' 118 082 (0 005) 0 8983 (0 00001 143 4 (12 3r 34 3 0.000 0 279 (0017) 594 (371
Total Area =561 0 (27 9)


~'ke~b~i~49~O~84)dr(


IProfile Fitting Report~


[12123-004.MDI] Scan Data


V


Ui
1170 117.5 118.0 118.5 119.0 11995
Twvo-Theta (deg)
Figure B-56. XRD deconvolution of PWA 1480+ (4hr. 10800C, unskewed, 2.38% error)


272


~"4pJZ~Yni~?~


4VIJ~E~b~blY111 I.._ .Lhh~~~~ ~~IW










(12Z124-004a.MDI] Scan Data Profile Fitting ReportI
SCAN: 117 0/119 98/O 02/11secl. Cu I(rnax)=20048, 01/14/08 11 30a
Residual Error of FIt= 9 32%/ (Two-Theta Range of Fd = 117 0/119 98), 2T(0:=0 0(deg!
# 2-Theta d(A) Area(al) AreaQ Skew FWHM XS(A)
1 117 986 10 002) 0 8987 (0 0000) 6539 5 108 9) 100 0 0 000 0 336 !O0 004 481 16)


117.0 117.5 118.0 118.5 11990 1195
Two-Theta (deg)

Figure B-57. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 9.32% error)


273












SCAN: 117 011194 9610 0)2/1(sec). Cu, 14 raxl=20048 01/14108 11:30a
Residual Error of Fit = 7 96% (Two-Theta Range of Fit = 117 01119.98), 2TlD)=0 0(rlegr
#; 2-Thata dtA) Arealall AreaI Skew FWHM XS(AI
1 10117 978 (0 002) 0 8988 (0 0000) 5809 3 106 5, 100 0 0.000 0 317 (0 0041 512 (7r
2 11'81" 250 (0 0401l 0 89:5 (0 0004 r 935 2 (185j 2i 16 1 0 000 0 5"'2 (0 0651 275 (32i
Total Area = 674 5 (213 6:


Profile Fitting Report


[1;2124-004a.MD] Scran Data


117.0 117.5 118.0 118.5 119.0 119.5
TwNo-Theta (deg)

Figure B-58. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 7.96% error)


274










[12124004a.MDI)l Scan Data Profile Fiting Report
SCAN: 117 01119 98/0 02/1(sec), Cu, l(maxi=0048I 01/14/08 11:30a
Residual Error of Fit = 3 84% I~wo-Theta Range of Fit = 117 01119,98), 2T(0)=0 Il(degl
# ~~ 2-Theta d(A)dl~ Area(al) Area%4 Skew FWHM XS(A)


117.0 117.5 118.0 118.5 119.0 119: r l.S
Two-Theta (deg)

Figure B-59. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 3.84% error)


275










[12124410aeMDI] Scan Datai Profle Fitting Report'
SCAN: 117 0111r9 9810 02/1(sec), Cu, I(max)=20048, 01/14/08 11 30a
Residual Error of Fit = 3 89%6 (Two-Thera Range of Fit = 117 011 19 981. 2T(0)=0 0(deg)
#2-Theta d(Al~~ A~reacal ) Area% Skew~ FWI-M XS(As


Total Area = 6915.2 (226.8)


117.0 117.5 118.0 118.5 119.0 119 5
Two-Theta (deg)

Figure B-60. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 3.89% error)


276











[12124004b.MOq Scan Data Proflea Fitting Report'
SCAN: 117 O/19 899/0 01;3(secl Cu. l(rnax)=62262, 01/14/08 12 03p
Residual Error of Fit = 3,37% (Two-Theta Range of Fit = '117 011 19 99), 2T(0)=0 Oldeg I
# 2-Thera d(Ai Arealal) Area%/ Skew FWHM XS(A)
1 1 17.984 (0.001) 0.8987 :0 0000, 12538 4 (160 91 100 0 0.000 0 261 (0 002) 639 17)
2 118 028 (0.003) 0.8985 l0 00001 9730 7 1353 01 77 6 0 000 0 650 t0 010) 240 (5)


ii


~------- ---- -------- -;--- ------------------- -- -- -r--------- -I------------r-------- ---i- ----- -- '


117.0 117 5 118.0 198.5 1190 119.5

Two-Theta (deg)

Figure B-61. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 3.37% error)


277












[121244)Q4b.MDI] Scan Data Profile Fitting Rpr
SCAN: 117 01119 99/0 01 3(sec), Cu, 1(max)=62262, 01/14/08 12:03p
Residual Error of Fit = 3 38% (Two-Theta Range of Fit=r 117 0/119,99), 2Tl0,-0 0(degr
#2-Theta d(A) Aree(ai) Area%~ Skew FWH)-M XS(A)
1 117 983 (0 001 0i 8987 (0 0000) 12151 6 1163 41 100 0 0.000 0 257 (0 002, 650 (7)
2 118 027 10 0031 0 8985 t0 00001 10052 3 1357 9) 82 7 0 000 0 633 10 010) 247 (51
Total Area =22203,9 (393.4)


__
_ij ~7 _T~


i


117.0 117.5 118.0 118.5 119.0 119.5

Two-Theta (deg)

Figure B-62. XRD deconvolution of PWA 1480+ (10hr. 10800C, unskewed, 3.38% error)


278











[1~2126-004a.MDI] Scana Data~ ~--~-~ Profile Fitting Report
SCAN: 117 00r 19 98/0 02/1 Residual Error of Fit = 2 83%. ITwo-Theta Ran~ge of Fit = 117,0/119.98). 2T(0)=0 0(deg)
#/ 2-Theta d(A) Area~all Areab Skew FWHM XS(Al)
1 118 128(0 019) 0 8980 (O0002) 2472 41686 Op 100 0 0.000 0 852 10 06 1) 183 (141
2 118 075 (0 004, O 8983 (0 0000) 2146 9 (435 51 8 6 80.000 0 S63 (0 027) 279 (141
Total Area = 4619.3 (812 61


117.0 117.5 118.0 118.5 11990 119.5
~Two-Theta (deg)

Figure B-63. XRD deconvolution of PWA 1480+ (1000hr. 10800C, unskewed, 2.83% error)


279










[1t2126-004b MDI] Scan Data~ Profile Fittng Report
SCAN: 117 0/119 99/00 01i31secl Cu, I~max)-27376, 01/14/08 12:361p
/Residual Eor ofFit = 1a 49 Two-Theta Range ofFt=R 117.0119199), 2T(0)=0 0deg)


117.0 '1 7 5 118.0 118.5 119.0 110.5
Two-Theta (deg)
Figure B-64. XRD deconvolution of PWA 1480+ (1000hr. 10800C, unskewed, 1.49% error)


280










[12126-004b.MDI] Scan Data PrdBl Fitting Report
SCAN: 117 0,1119 9910 Ott31secl. Cu, Isrnaxl=27376, 01/14/08 12 36p
Resdual Error of Fit 1 49 (Two-Theta Range of Fit = 117 0/119 99), 2T(0)]=0 0(deg)
#i 2-Thete d(A) Area(all Area% Skew FWHM XS(A)
1 _118 041 (0 001)I 0 8985 (O 00001 7219 2 ri264 1) 76 9 0 000 0 527 (0 00'I 298 (51
2 118 060! (0 0021 0 8981 (0 0000) 9386 91457 3, 100 0 0.DDD 0 939 (0 0151 165 (4)
Total Area =16606.1 (528.0)


1177.0 117,5 1'18.0 1185 119.0 179.5
TIwo-Theta (deg)

Figure B-65. XRD deconvolution of PWA 1480+ (1000hr. 10800C, unskewed, 1.49% error)











[12123-004.MDI] Scan Data Promie Fitting Report
SCAN: 117.0/119s.9910 ')1151sceCu. lmnaxr=2707, 01/09/08 01 27p
Residual Error of Fit = 2 36% tTwho-Theta Range of Fit = 117 0,1 19 99r, 2T(0)=0 0)(degl
# 2-Theta d(Al Area(al) Area% Sktwu FWHM XS(A)
1 117 964 10 031) 0 8988 (o 0003) 406 4(36 7 100 0 -0 02 0 864 (0 052, 180 (12)
2 1 18 101 !0 009) 0 8982 (0 0001) 158 6 r14 51 39 0 0 311 0 292 !0 016J 564 (31]


117.5 118.0 178.5
Two-Theta (deg)


1190 119.5


Figure B-66. XRD deconvolution of PWA 1480+ (4hr. 10800C, skewed, 2.36% error)


282











[21244~a04aMo scan Data Profie Fittng Report
SCAN: 117 0,1 19 981[0 02/1(seci, Cu, I(max)=20048, 01/14/08 11 3Da
Residual Error of Fit = 3.77% (Two-Theta Range of FIt= 117 01119 98), 2T(0)=0 0(desg)
#1 2-Theta d(A~) Area(al) Area% Skew FNHM X(A
1 117 972 (O002) 0 898810 00001 3222 41105 1, 86 8 -0.141 0 239 10 0051 711 (171
2 118 012 (0 0051 0 8986 r0 00001 3 11 31217 0, 100 0 -0.022 0 532 10 012) 295 r8r
Total Area = 6933.6 (241.1)


117.0 177.5 118.0 118.5 119.0 119.5
Two-Theta (deg)

Figure B-67. XRD deconvolution of PWA 1480+ (10hr. 10800C, skewed, 3.77% error)


283










[11450aMS San Data ~rofile Fitting Report
SCAN: 117.0/119.98/0 02/1 (sec), Cu, I(rnax)=20048, 01/14/08 11:30a
SResidual Error of Fit= 7 19% (Two-Theta Riange of F1= 117 0/1119.98), 2TIO)=0 0(degl


Total A~rea = 6;74 1 1,446 5)


117.0 117 5 198.0 118 5 119.0 119.5
Two-Theta (deg)

Figure B-68. XRD deconvolution of PWA 1480+ (10hr. 10800C, skewed, 7.19% error)


284











[~12124-004a.MDI] Scan Data Profile Fitting Report
SCAN: 117 0/11 810 SM02/1(see), Cu, (max)=20048. 01/14/08 11:30a
Residual Error of Fit = 8 42%/ (Two-Theta Range of Fit = 117r.0/ 19.98), 2Tc01=0 Oldegl
#l 2-Theta d(A,) Areatal) Area% Skew FWHM XS(A)
1 117 989 (0 003) 0 8987,0 00I00; 6305 4 (05 31 100 0 0 018 0 329 (D300~4) 9(7
2 118 657 :o 074l a0 8956 (0 0007) 633.6 (155 1) 10 0 0.9000 1.143 (0.1~66) 137 (21)
Total Area = 6938 9 (187 5)


118.0 118 5 179.0
Two-Theta (deg)


119 5


Figure B-69. XRD deconvolution of PWA 1480+ (10hr. 10800C, skewed, 8.42% error)


285











[12124-004arMDI Scn Data Profilhg e Fttn eotl
SCAN: 117.0/119.98/0.02, Ilsecl. Cu, Ilmax)=20048, 01/14/08 11 30a
SResidual Error of Fit= 7.04% (Twvo-Theta Riange of Fit = 117 01119 198 2aT(0)=0 0(deg)
# 2-Theta d(A) Area(alr AreaO% Skew FWHM XSIAI
1 I. 1 17.98:! l0 004) 0.8987 (0 0000) 5747.1 (88.9) 100 0 0.006 0.314 (0.003) 517 (6)
2 1: ;118.371 (0.043) 0 8969 1o 0004 r 1123 9 (122 6) 19 6 0.711 0.748 (0 051 1 209 (15)
Total Area= 687' 0 (1r51


1170 117,5 1180 178.5 199.0 179.5
Two-Theta (deg)

Figure B-70. XRD deconvolution of PWA 1480+ (10hr. 10800C, skewed, 7.04% error)


286










[I12124-04b.MDI] Scan Data~ Pmfile Fitting Report
SCAN: 117 011 19 9910 0 1i3ts~ec, Cu, 10nax)=62262, 01/14/08 12 03p
Residual Error of Fit = 3 28% (Twvo-Theta Range of Fit = 117 0,1119 993. 2T(0)=0 0(deg)
# 2-Theta atAl Areatail A~rea%: SIkew FWYHM XS(A)
1 1~17 978 (0 001) 0 6968 r0 00001 1249T7 1162 31 100 -0 001 0 261 10 002) 640 17,
2 118 036 r0 005r 0 6985 r0 00001 97;4 2 (351 31 78 2 0.086 0 649 10 010) 240 (51


117 0 117.5 1 ta0 118 5 119.0 119 5
Two-Theta (deg)

Figure B-71. XRD deconvolution of PWA 1480+ (10hr. 10800C, skewed, 3.28% error)


287











[1212~4-04b.MDI Scan Data Profile Fittng Report
SCAN: 117 0/119 9.9/30 01/(sec), Cu, IsrnaK)=62262, 01/14/08 12:03p
Residual Error of Fit= 3 31%6 (Twyo-Theta Range of Fit = 11'7 0/119.99), 2T(0)=0 0(deg)
# 2-Theta dtAl Arealall Area%/ Sltew FWHM XS(A)
1 I:: 117 977 100011 0 8988 (0.0000) 12139 1 (166 5) 100.0 -0.092 0.258 (0.003) 650 17 I
2 I 118 032 10 0051 0 898510! 00001 10)0746 8 357 6] 83 0 0 048 0 63310 0101 4 5
Total Area =22213,7 (394 5)


117 0 117.5 118.0 118.5 119.0 119.5
Two-Theta (deg)

Figure B-72. XRD deconvolution of PWA 1480+ (10hr. 10800C, skewed, 3.31% error)


288










[121264)O4a.MDI Scan Data Proffl Fitting Report
SCAN: 117 0/119 98410D2tlisecl, Cu, I(max)=7600, 01/14/08 12 19p
Residual Error of Fit = 2 79% (Two Theta Range of Fit = 117.0/119.98), 2Tl0)=0 Oldegl
#I 2-Theta d(A) Area(al) AreaA Skew FWH X(A
1 I 118 134 (0 029) 0 8980 r0 0)003) 2465 6 1741 51 100 0 -0 002 0 852 (0 067, 183 (15r
2 118 083 (0 006) 0.8983 r0 0001) 2162 8 1471 5) 87 3 -0078 0 562 I0 0291 2791161
Total Area = 4618 41878.7)


117.0 117.5 118.0 118.5 119.0 119 5
TwJo-Theta (deg)


Figure B-73. XRD deconvolution of PWA 1480+ (1000hr. 10800C, skewed, 2.79% error)


289














2T(0)=0.0(deg)
Stewv FWHM XS(A)l
0 173 0 527 (0 006) 299 (4) -
0 068 0 939 (0 017) 185 (4)


yPrfle Fitting fRepr


[121264104b.MDI] Scan Data
SCAN: 117 0,1119 9910 01tr3stse Cu, 1(maxl=27376, 01/140)8 12:36p
Residual Error of Fit = 1 18%P (Two-Theta Range of FR = 1'17 011 19 991
# 2-Thete dtA) Areatall A~ree%
1 118 053 10 0011 0 8984 l0 00001 7308 1 r255 6) 78 6
2 IC118 098 (0 0091 0 8982 (0 0001) 92961 (1476 51 100 0
Tota Area = 16604.2 (540 7)


11770 117 5 118.0 118 5 119.0 119.5

Two-Theta (deg)

Figure B-74. XRD deconvolution of PWA 1480+ (1000hr. 10800C, skewed, 1.18% error)


290











[12126-04b.MDI) Scan Data Profile Fitting ReportI
SCAN: 117 0/i11 999001/3lsec i.Cu, :I mrnan27376. 01/14/08 12:36p
Residual Error of Fit = 1 18%o (Two-Theta Range of Fit = 117 0/119_99), 2T!0I;=0 oegs
# 2-Theta d(A) A~reaffi) Area% Skew FWHM XS(A)
1 118 097 (0 009) 0 8982 10 0001) 9300 0 I78 4) 100 0 0 065 0) 939 r0 017) 165 (4!
21 118.053 (0.001; 0.8984 (0 0000) 7307 1 1256 61 78 6 0 171 0 527 (0 00-) 299 (41
Total Area = 16607.1 1542 91


117.0 1;7.5 11 0 17 S 1. 1.
Two-Theta (deg)

Figure B-75. XRD deconvolution of PWA 1480+ (1000hr. 10800C, skewed, 1.18% error)










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BIOGRAPHICAL SKETCH

Brandon Wilson was born and raised in the heart of horse country around Ocala, FL. At a

young age he j oined the Cub Scouts of America and later the Boy Scouts of America and

achieved the highest rank of Eagle at the age of 17. Brandon graduated with honors from

Vanguard High School (Ocala, FL) in 1998 and enrolled in college at Central Florida

Community College (Ocala, FL). Originally intending to attend film school at the University of

Central Florida (Orlando, FL), he made the decision to attend the University of Florida

(Gainesville, FL) instead and entered the mechanical engineering program. Following his first

materials course, Introduction to Materials, taught by Dr. Dow Whitney and Dr. John Ambrose,

he made the decision to enter the materials science and engineering program with a metals

specialization.

Shortly after entering the materials science and engineering department, he applied for the

University Scholars Program organized by the University of Florida. The University Scholars

Program is an opportunity for undergraduate students to learn about research in a laboratory

setting. Upon being accepted into the program, he was partnered with Dr. Gerhard Fuchs who

became his eventual graduate advisor and mentor. Graduating cum laude with his bachelors

degree in May 2003, Brandon went on to marry Krista Renner three weeks later. He then

entered the graduate program in materials science and engineering in August 2003. He

continued working with Dr. Fuchs on the research presented herein. Upon graduating he plans to

take an engineering position working with nickel and titanium alloys for aerospace and power

generation applications.


297





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1 THE PRIMARY CREEP BEHAVIOR OF SINGLE CRYSTAL, NICKEL BASE SUPERALLOYS PWA 1480 AND PWA 1484 By BRANDON CHARLES WILSON A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLOR IDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2008

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2 2008 Brandon Charles Wilson

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3 In memory of Richard Doc Connell, Ph.D.

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4 ACKNOWLEDGMENTS This work would not have been successful without the help of ma ny great and wonderful people. First, I must thank my family and frie nds that have supported and loved me through this process. These are the people that were there to give me advice when I needed it, to help me focus when I needed to, to help me relax when I needed to, and to pray for me always. So much of what brings a Ph.D. to completion happens outs ide of the lab and away from a computer and I am incredibly thankful for the love and care dire cted towards my wife and me while we were in graduate school. My research hinged on the kindness of severa l strangers (at the time) during the early stages of my work. Three engineers at Pratt & Whitney Aircraft Engines (East Hartford, CT) deserve to be mentioned for their gift of a comb ined total of 47 single crystal test bars. Alan Cetel and Dilip Shaw supplied the original 12 bars along with some journal articles of note to get me started. Later, Samuel Krotzer supplie d 11 bars of PWA 1480, 12 bars of PWA 1484, and even agreed to make a special heat of PWA 1480 with 3wt.% rhenium added free of charge. Without their help, this project would not have gone very far. I also need to recognize Rick Black and Bill Kumrow of Satec (a division of Instron, Norwood, MA) for their help in diagnosing and repairing electrical and software failures on the M3 creep system used for most of the data presented here. Additionally, the hard work of several members of the Major Analytical Instrumentation Center (MAIC) at the University of Florida (Gainesville, FL) and the Advanced Materials Processing and Analysis Center (AMPAC) at the University of Central Florida (Orlando, FL) contributed to several aspects of this investigation. Specifica lly, Wayne Acree, Ben Pletcher, Jung Hun Jang, Valentine Craciun, Gerald Bourne and Kerry Siebein of MAIC and Kirk Scammon and Helge Heinrich of AMPAC were inst rumental in training and helping me with the

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5 various analytical techniques uti lized for this research. Mike Kaufman and Anantha Puthicode, formerly of the University of North Texas (Dento n, TX), deserve to be recognized for their help with generating the Local Electrode Atom Probe (LEAP) data presented herein as well. Within the University of Florida, I would like to acknowledge my advisor and mentor, Dr. Gerhard Fuchs. Without his help and direction I would have been lo st in the enormity of the task at hand. Former graduate student Slade Stotlz also deserves to be mentioned for his here-to-fore unacknowledged SEM work during my undergraduate re search. In addition to Slades help, I am indebted to my fellow students in the High Temperature Alloys Laboratory (HTAL) for their advice, encouragement, and humor while I have been a student. From the Particle Engineering Research Center (PERC) at UF, I thank Nate Stevens, Ph.D. who is a great friend that volunteered to proof read this disse rtation in various stages of comp letion. I am also thankful for the quick and responsive work of several support staff members between the academic, payroll, finance, and secretary offices. Without their help nothing that is done in the department would ever be completed and I am thankful for their he lp in spite of the difficulties I often presented them. Finally, I need to recognize my wife, Krista Renner Wilson, Ph.D., for her constant support and encouragement. Whether I was frustrat ed or nervous or unsure of what to do next, she would gently, lovingly guide my hand and my thoughts in the right direction. Without her help, I would not have completed this project.

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6 TABLE OF CONTENTS page ACKNOWLEDGMENTS...............................................................................................................4 LIST OF TABLES................................................................................................................. ..........9 LIST OF FIGURES.......................................................................................................................10 ABSTRACT...................................................................................................................................23 CHAPTER 1 INTRODUCTION................................................................................................................. .25 2 BACKGROUND................................................................................................................... .29 Overview....................................................................................................................... ..........29 The Rhenium Effect........................................................................................................30 Secondary Precipitates.................................................................................................33 Lattice Mismatch.............................................................................................................34 Modern Superalloys................................................................................................................35 The Challenge of High Temperature Service.........................................................................37 Tensile Behavior..............................................................................................................38 Creep Behavior................................................................................................................41 Summary.................................................................................................................................43 3 EXPERIMENTAL PROCEDURES.......................................................................................44 Materials.................................................................................................................................44 Heat Treatment.......................................................................................................................46 Heat Treatment Development..........................................................................................46 Differential Thermal Analysis.........................................................................................48 Furnaces...........................................................................................................................49 Characterization......................................................................................................................50 Preparing Samples for Metallography.............................................................................51 Preparing Samples for TEM............................................................................................52 Preparing Samples for LEAP..........................................................................................53 Preparing Samples for X-Ray Diffraction.......................................................................55 JMatPro Thermodynamic Prediction...............................................................................57 Mechanical Behavior............................................................................................................ ..57 Tensile Testing................................................................................................................58 Creep Testing...................................................................................................................60 4 RESULTS: ALLOY MICROSTRUCTURES.......................................................................62 Phase Descriptions............................................................................................................. .....62

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7 The Phase......................................................................................................................62 The Phase.....................................................................................................................64 Primary .................................................................................................................65 Secondary .............................................................................................................66 The / Eutectic..............................................................................................................68 The Carbides....................................................................................................................69 The Topologically Close Packed (TCP) Phases..............................................................70 Changes Following Primary Creep.........................................................................................72 5 RESULTS: TENSILE BEHAVIOR......................................................................................91 6 RESULTS: CREEP BEHAVIOR........................................................................................113 Full Length Tests.............................................................................................................. ....114 PWA 1480.....................................................................................................................114 PWA 1480+...................................................................................................................116 PWA 1484.....................................................................................................................118 Interrupted Tests.............................................................................................................. .....120 Transmission Electron Microscopy (TEM)..........................................................................121 PWA 1480.....................................................................................................................121 PWA 1480+...................................................................................................................122 PWA 1484.....................................................................................................................123 7 RESULTS: ADDITIONAL CHARACTERIZATION........................................................140 Local Electrode Atom Probe (LEAP)...................................................................................140 PWA 1484 LT...............................................................................................................143 Reconstruction (solution HT).................................................................................143 Composition profile (solution HT).........................................................................143 PWA 1484 HT...............................................................................................................145 Reconstruction (age HT)........................................................................................145 Composition profile (age HT)................................................................................145 Reconstruction (solution HT).................................................................................146 Composition profile (solution HT).........................................................................146 Secondary Concentrations.........................................................................................147 X-Ray Diffraction (XRD).....................................................................................................148 8 DISCUSSION................................................................................................................... ....163 Microstructure.......................................................................................................................164 / Morphology............................................................................................................164 Carbides.........................................................................................................................168 Topologically Close Packed Phases..............................................................................168 Tensile Behavior ...................................................................................................................169 Secondary ..................................................................................................................170 Channel Thickness......................................................................................................172 Lattice Misfit.................................................................................................................173

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8 Stacking Fault Energy...................................................................................................173 Anti-Phase Boundary Energy........................................................................................174 Tensile Results...............................................................................................................1 75 Creep Behavior.....................................................................................................................176 Tertiary Creep................................................................................................................177 Rafting...........................................................................................................................178 Primary Creep................................................................................................................18 0 Creep Results.................................................................................................................182 Modeling Primary Creep......................................................................................................186 Lattice Misfit.................................................................................................................189 Secondary ..................................................................................................................191 Composition..................................................................................................................19 3 Concluding Remarks............................................................................................................1 95 9 CONCLUSION................................................................................................................... ..199 Conclusions...........................................................................................................................199 Future Directions..................................................................................................................200 APPENDIX A DIFFERENTIAL THERMAL ANALYSIS.........................................................................201 B XRD PEAK DECONVOLUTION.......................................................................................208 LIST OF REFERENCES.............................................................................................................292 BIOGRAPHICAL SKETCH.......................................................................................................297

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9 LIST OF TABLES Table page 2-1 Compositions in weight percent of several common Nickel-based superalloys................36 3-1 Laue orientation data and orig inal heat treatments for PWA 1480...................................44 3-2 Laue orientation data and orig inal heat treatments for PWA 1484...................................45 3-3 Laue orientation data and orig inal heat treatments for PWA 1480+.................................45 3-4 Heat treatments used in this study.....................................................................................47 3-5 DTA results for all three alloys..........................................................................................49 3-6 Tensile test matrix w ith sample identification...................................................................59 3-7 Creep test matrix with sample identification.....................................................................61 5-1 Tensile results at 700 C and 815 C for all three alloys...................................................101 5-2 Elastic modulus calculations from creep loads................................................................104 5-3 Creep loads vs. yield st rength for all three alloys............................................................111 6-1 Primary creep and creep rates from interrupted creep tests.............................................125 6-2 Rupture lives and total creep elonga tion from full-length creep tests.............................127 7-1 Composition (wt%) of secondary precipitates.............................................................160 7-2 Net changes in secondary concentration with increasing precipitate size...................160 7-3 Lattice misfit (%) for the (002) pl ane following heat treatments at 1080 C...................161 B-1 PWA 1480 (002) peak lattice misfit calculations............................................................211 B-2 PWA 1480+ (002) peak lattice misfit calculations..........................................................212 B-3 PWA 1484 (002) peak lattice misfit calculations............................................................213 B-4 PWA 1480 (004) peak lattice misfit calculations............................................................214 B-5 PWA 1480+ (004) peak lattice misfit calculations..........................................................215 B-6 PWA 1484 (004) peak lattice misfit calculations............................................................216

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10 LIST OF FIGURES Figure page 2-1 Idealized creep curve with th e three creep stages labeled.................................................32 3-1 Flow-chart of heat treatments for LEAP samples..............................................................54 3-2 SEM micrograph of the tip of a LEAP specimen..............................................................55 3-3 Creep and tensile specimen geometry and dimensions......................................................58 3-4 Three Type K thermocouples attached to a creep specimen..............................................61 4-1 / microstructure of PWA 1480 HT3..............................................................................73 4-2 / microstructure of PWA 1480+ HT3............................................................................73 4-3 / microstructure of PWA 1484 HT3..............................................................................74 4-4 Composition of the phase of PWA 1480 as a function of temperature...........................74 4-5 Composition of the phase of PWA 1480+ as a function of temperature.........................75 4-6 Composition of the phase of PWA 1484 as a function of temperature...........................75 4-7 Composition of the phase of PWA 1480 as a function of temperature.........................76 4-8 Composition of the phase of PWA 1480+ as a function of temperature.......................76 4-9 Composition of the phase of PWA 1484 as a function of temperature.........................77 4-10 The phase in PWA 1480 near a retained eutectic..........................................................77 4-11 Irregular phase in PWA 1480 near a partia lly solutioned eutectic region.....................78 4-12 / eutectic region in PWA 1480+ during the early stages of solutioning.......................78 4-13 The structure of as-cast PWA 1484...............................................................................79 4-14 The volume fraction vs. temperat ure for all three alloys...............................................79 4-15 Secondary in the matrix of PWA 1480 following an interrupted creep test...............80 4-16 Secondary in the matrix of PWA 1480 following an interrupted creep test...............80 4-17 Secondary in the matrix of PWA 1480+ fo llowing an interrupted creep test................81

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11 4-18 Three dimensional LEAP compositional map (18wt.% Al iso-surface)............................82 4-19 / eutectics in as-cast PWA 1480....................................................................................83 4-20 / eutectics in the vicinity of primary carbides in PWA 1480+......................................83 4-21 Close-up of a eutec tic in as-cast PWA 1480+...................................................................84 4-22 Close-up of a retained eutectic in PWA 1480+.................................................................84 4-23 Carbide phase in PWA 1480+ (HT1A)..............................................................................85 4-24 Carbide phase in PWA 1480+ (as-cast, longitudinal section)...........................................85 4-25 Local carbide network in PWA 1480.................................................................................86 4-26 Carbide phase in PWA 1484..............................................................................................86 4-27 A possible carbide that dissolved during solution heat treatment......................................87 4-28 TCP phase formation in PWA 1480+ during interrupted creep testing.............................87 4-29 TCP phase formation in PWA 1480+ during interrupted creep testing.............................88 4-30 phase elongation in the [ 110] direction in PWA 1484....................................................88 4-31 phase elongation in the [ 110] direction in PWA 1484....................................................89 4-32 phase shear along (111) planes in PWA 1484...............................................................89 4-33 phase shear along (111) planes in PWA 1484...............................................................44 5-1 Tensile results for PWA 1480..........................................................................................100 5-2 Tensile results for PWA 1480+.......................................................................................100 5-3 Tensile results for PWA 1484..........................................................................................101 5-4 Comparison of tensile results for all three alloys with the LT age (700 C)....................102 5-5 Comparison of tensile results for all three alloys with the LT age (815 C)....................102 5-6 Comparison of tensile results for a ll three alloys with the HT age (700 C)....................103 5-7 Comparison of tensile results for a ll three alloys with the HT age (815 C)....................103 5-8 Plastic deformation be havior of PWA 1480 LT at 700 C...............................................104 5-9 Plastic deformation beha vior of PWA 1480 HT at 700 C...............................................105

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12 5-10 Plastic deformation be havior of PWA 1484 LT at 700 C...............................................105 5-11 Plastic deformation beha vior of PWA 1484 HT at 700 C...............................................106 5-12 Plastic deformation behavior of PWA 1480+ LT at 700 C.............................................106 5-13 Plastic deformation behavior of PWA 1480+ HT at 700 C............................................107 5-14 Tensile behavior of PW A 1480 LT at both temperatures................................................107 5-15 Tensile behavior of PWA 1480 HT at both temperatures................................................108 5-16 Tensile behavior of PWA 1480+ LT at both temperatures..............................................108 5-17 Tensile behavior of PWA 1480+ HT at both temperatures.............................................109 5-18 Tensile behavior of PW A 1484 LT at both temperatures................................................109 5-19 Tensile behavior of PWA 1484 HT at both temperatures................................................110 5-20 Yield strength as a function of temperature for all three alloys.......................................110 5-21 Ultimate Tensile Strength as a functi on of temperature for all three alloys....................111 5-22 True Failure Stress as a function of temperature for all three alloys...............................112 6-1 Creep at 704 C/862 MPa of all three alloys....................................................................125 6-2 Creep at 760 C/690 MPa of all three alloys....................................................................126 6-3 Creep at 815 C/621 MPa of all three alloys....................................................................126 6-4 Primary creep of PWA 1480............................................................................................128 6-5 Primary creep of PWA 1480+.........................................................................................128 6-6 Primary creep of PWA 1484............................................................................................129 6-7 Creep at 704 C/862 MPa of PWA 1484..........................................................................129 6-8 Primary creep comparison at 704 C/862 MPa for all three alloys..................................130 6-9 Primary creep comparison at 704 C/862 MPa for all three alloys..................................130 6-10 Primary creep comparison at 815 C/621 MPa for all three alloys..................................131 6-11 Bright field(a)/Dar k field(b) pair of deformation in PWA 1480.....................................132 6-12 Bright field(a)/Dar k field(b) pair of deformation in PWA 1480.....................................133

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13 6-13 Bright field TEM image of di slocation networks in PWA 1480.....................................134 6-14 Stacking fault and disl ocation shear of PWA 1480.........................................................134 6-15 A stacking fault in PWA 1480.........................................................................................135 6-16 Secondary precipitates (marked by arrows) in PWA 1480..........................................135 6-17 Stacking fault interactions fo llowing primary creep in PWA 1480+..............................136 6-18 Stacking fault interactions and a dislocation network in PWA 1480+............................136 6-19 Short range stacking fault shear of PWA 1480+.............................................................137 6-20 Bright field TEM im age of PWA 1484 LT (704 C)........................................................137 6-21 Stacking fault shear of precipitates in PWA 1484.......................................................138 6-22 Interfacial dislocation ne tworks present in PWA 1484...................................................139 7-1 Schematic illustrating the basic function of the LEAP system........................................151 7-2 LEAP specimens before, (a), and af ter, (b), the final polishing step...............................152 7-3 Iso surfaces created with the LEAP system (PWA 1484 LT with solution HT).............153 7-4 Magnified (SEM) view of the tip analyzed in Figure 7-3................................................154 7-5 Composition profile from the specimen in Figure 7-3.....................................................154 7-6 The distribution of all re corded ions for PWA 1484 HT.................................................155 7-7 Iso surfaces created with the LEAP system (PWA 1484 HT no solution HT)................155 7-8 Composition profile from the specimen in Figure 7-7.....................................................156 7-9 Iso surface (18% Aluminum) created with the LEAP system.........................................157 7-10 LEAP results with only Al, Ta, Cr, and Mo ions represented.........................................158 7-11 Composition profile from the specimen in Figures 7-9 and 7-10....................................159 7-12 Illustration of the data selected for the profile shown in Figure 7-11..............................159 7-13 An example of the deconvolution process.......................................................................161 7-14 Lattice misfit vs. heat treatment from from both the (002) and (004) peaks...................162 A-1 DTA trace for PWA 1480 in the HT1 condition..............................................................202

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14 A-2 DTA trace for PWA 1484 in the HT1 condition..............................................................203 A-3 DTA trace for PWA 1480+ in the HT0 condition...........................................................204 A-4 DTA trace for PWA 1480 in the HT2 condition..............................................................205 A-5 DTA trace for PWA 1484 in the HT2 condition..............................................................206 A-6 DTA trace for PWA 1480+ in the HT2 condition...........................................................207 B-1 XRD deconvolution of PWA 1480 (4hr. 1080 C, skewed, 6.28% error)........................217 B-2 XRD deconvolution of PWA 1480 (4hr. 1080 C, skewed, 14.81% error)......................218 B-3 XRD deconvolution of PWA 1480 (4hr. 1080 C, unskewed, 6.63% error)....................219 B-4 XRD deconvolution of PWA 1480 (4hr. 1080 C, unskewed, 12.58% error)..................220 B-5 XRD deconvolution of PWA 1480 (10hr. 1080 C, unskewed, 7.94% error)..................221 B-6 XRD deconvolution of PWA 1480 (10hr. 1080 C, skewed, 5.63% error)......................222 B-7 XRD deconvolution of PWA 1480 (100hr. 1080 C, unskewed, 6.94% error)................223 B-8 XRD deconvolution of PWA 1480 (100hr. 1080 C, skewed, 6.82% error)....................224 B-9 XRD deconvolution of PWA 1480 (100hr. 1080 C, unskewed, 5.20% error)................225 B-10 XRD deconvolution of PWA 1480 (1000hr. 1080 C, unskewed, 5.93% error)..............226 B-11 XRD deconvolution of PWA 1480 (1000hr. 1080 C, skewed, 4.61% error)..................227 B-12 XRD deconvolution of PWA 1480 (1000hr. 1080 C, skewed, 5.45% error)..................228 B-13 XRD deconvolution of PWA 1480 (4hr. 1080 C, unskewed, 5.58% error)....................229 B-14 XRD deconvolution of PWA 1480 (4hr. 1080 C, unskewed, 7.77% error)....................230 B-15 XRD deconvolution of PWA 1480 (4hr. 1080 C, skewed, 4.99% error)........................231 B-16 XRD deconvolution of PWA 1480 (4hr. 1080 C, skewed, 1.70% error)........................232 B-17 XRD deconvolution of PWA 1480 (10hr. 1080 C, unskewed, 5.79% error)..................233 B-18 XRD deconvolution of PWA 1480 (10hr. 1080 C, skewed, 1.41% error)......................234 B-19 XRD deconvolution of PWA 1480 (1000hr. 1080 C, unskewed, 3.17% error)..............235

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15 B-20 XRD deconvolution of PWA 1480 (1000hr. 1080 C, unskewed, 1.77% error)..............236 B-21 XRD deconvolution of PWA 1480 (1000hr. 1080 C, unskewed, 1.75% error)..............237 B-22 XRD deconvolution of PWA 1480 (1000hr. 1080 C, skewed, 3.14% error)..................238 B-23 XRD deconvolution of PWA 1480 (1000hr. 1080 C, skewed, 1.71% error)..................239 B-24 XRD deconvolution of PWA 1480 (1000hr. 1080 C, skewed, 1.71% error)..................240 B-25 XRD deconvolution of PWA 1484 (4hr. 1080 C, unskewed, 13.36% error)..................241 B-26 XRD deconvolution of PWA 1484 (1000hr. 1080 C, unskewed, 6.30% error)..............242 B-27 XRD deconvolution of PWA 1484 (10hr. 1080 C, unskewed, 4.61% error)..................243 B-28 XRD deconvolution of PWA 1484 (1000hr. 1080 C, unskewed, 5.85% error)..............244 B-29 XRD deconvolution of PWA 1484 (10hr. 1080 C, unskewed, 4.45% error)..................245 B-30 XRD deconvolution of PWA 1484 (4hr. 1080 C, skewed, 11.71% error)......................246 B-31 XRD deconvolution of PWA 1484 (1000hr. 1080 C, skewed, 3.71% error)..................247 B-32 XRD deconvolution of PWA 1484 (1000hr. 1080 C, skewed, 3.20% error)..................248 B-33 XRD deconvolution of PWA 1484 (10hr. 1080 C, skewed, 4.07% error)......................249 B-34 XRD deconvolution of PWA 1484 (10hr. 1080 C, skewed, 4.03% error)......................250 B-35 XRD deconvolution of PWA 1484 (1000hr. 1080 C, unskewed, 4.15% error)..............251 B-36 XRD deconvolution of PWA 1484 (1000hr. 1080 C, unskewed, 3.24% error)..............252 B-37 XRD deconvolution of PWA 1484 (1000hr. 1080 C, unskewed, 4.35% error)..............253 B-38 XRD deconvolution of PWA 1484 (1000hr. 1080 C, skewed, 3.09% error)..................254 B-39 XRD deconvolution of PWA 1484 (1000hr. 1080 C, skewed, 2.07% error)..................255 B-40 XRD deconvolution of PWA 1480+ (4hr. 1080 C, unskewed, 5.07% error).................256 B-41 XRD deconvolution of PWA 1480+ (4hr. 1080 C, unskewed, 10.42% error)...............257 B-42 XRD deconvolution of PWA 1480+ (4hr. 1080 C, unskewed, 10.33% error)...............258 B-43 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 9.55% error)...............259

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16 B-44 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 5.73% error)...............260 B-45 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 4.87% error)...............261 B-46 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, unskewed, 4.41% error)...........262 B-47 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, unskewed, 4.37% error)...........263 B-48 XRD deconvolution of PWA 1480+ (4hr. 1080 C, skewed, 5.05% error).....................264 B-49 XRD deconvolution of PWA 1480+ (4hr. 1080 C, skewed, 6.91% error).....................265 B-50 XRD deconvolution of PWA 1480+ (4hr. 1080 C, skewed, 6.85% error).....................266 B-51 XRD deconvolution of PWA 1480+ (10hr. 1080 C, skewed, 9.54% error)...................267 B-52 XRD deconvolution of PWA 1480+ (100hr. 1080 C, skewed, 5.42% error).................268 B-53 XRD deconvolution of PWA 1480+ (100hr. 1080 C, skewed, 4.13% error).................269 B-54 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, skewed, 4.30% error)...............270 B-55 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, skewed, 4.28% error)...............271 B-56 XRD deconvolution of PWA 1480+ (4hr. 1080 C, unskewed, 2.38% error).................272 B-57 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 9.32% error)...............273 B-58 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 7.96% error)...............274 B-59 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 3.84% error)...............275 B-60 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 3.89% error)...............276 B-61 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 3.37% error)...............277 B-62 XRD deconvolution of PWA 1480+ (10hr. 1080 C, unskewed, 3.38% error)...............278 B-63 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, unskewed, 2.83% error)...........279 B-64 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, unskewed, 1.49% error)...........280 B-65 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, unskewed, 1.49% error)...........281 B-66 XRD deconvolution of PWA 1480+ (4hr. 1080 C, skewed, 2.36% error).....................282 B-67 XRD deconvolution of PWA 1480+ (10hr. 1080 C, skewed, 3.77% error)...................283

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17 B-68 XRD deconvolution of PWA 1480+ (10hr. 1080 C, skewed, 7.19% error)...................284 B-69 XRD deconvolution of PWA 1480+ (10hr. 1080 C, skewed, 8.42% error)...................285 B-70 XRD deconvolution of PWA 1480+ (10hr. 1080 C, skewed, 7.04% error)...................286 B-71 XRD deconvolution of PWA 1480+ (10hr. 1080 C, skewed, 3.28% error)...................287 B-72 XRD deconvolution of PWA 1480+ (10hr. 1080 C, skewed, 3.31% error)...................288 B-73 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, skewed, 2.79% error)...............289 B-74 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, skewed, 1.18% error)...............290 B-75 XRD deconvolution of PWA 1480+ (1000hr. 1080 C, skewed, 1.18% error)...............291

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18 LIST OF ABBREVIATIONS AC: Air cooled (following a heat treatment) AIM: Advanced Insertion of Materials (a DARPA program) AMPAC: Advanced Materials Processing and Analysis Center (a unit of the University of Central Florida) APB: Anti phase boundary (can be formed in ordered phases) ASM: American Society for Metals d: d spacing (the spacing between planes of atoms) DARPA: Defense Advanced Research Proj ects Agency (a division of the United States Department of Defense) FWHM: Full Width at Half Maximum (a para meter that describes the shape of an intensity peak used for XRD deconvolution) GFQ: Gas furnace quench (used for solution heat treatment, Helium gas was injected into a vacuum furnace for rapid cooling) hkl: Generic term used to indicat e the miller indices of a plane HT: Used to designate the high temperature age (871 C/32 hr./AC), also a generic abbreviation for heat treatment HT1, HT2, etc.: Solution heat treatment designati ons. All heat treatments are defined in Table 3-4 IGT: Industrial gas turbin e (usually used in the power generation industry) LT: Used to designate the low temperature age (704 C/24 hr./AC) LVDT: Linear Variable Differential Transducer MAIC: Major Analytical Instrumentation Center (a unit of the University of Florida) STD HT: Designates the commercial heat tr eatment for the alloy in question (Table 3-4) /2 : Type of X-ray diffractometer 2 : The angle that a diffractometer measures during a scan CMSX: Denotes a single crystal s uperalloy developed by Canon Muskegon

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19 CMSX-2: First generation single crysta l superalloy developed by Canon Muskegon, Table 2-1 CMSX-4: Second generation single crystal superalloy developed by Canon Muskegon, Table 2-1 PWA: Denotes an alloy developed by Pratt & Whitney PWA 1480: First generation single crystal, nickel base superall oy made by Pratt & Whitney. The composition is shown in Table 2-1 PWA 1480+Re: An experimental single crystal, nickel base superalloy created by adding 3wt.% rhenium to PWA 1480. The co mposition is shown in Table 2-1 PWA 1480+: Same as PWA 1480+Re PWA 1484: Second generation single crystal, nickel base superalloy made by Pratt & Whitney. The composition is shown in Table 2-1 DTA: Differential Thermal Analysis EDS: Energy dispersive spectroscopy FIB: Focused ion beam LEAP: Local electrode atom probe SEM: Scanning electron microscope (or microscopy) TEM: Transmission electron microscope (or microscopy) XRD: X-ray diffraction EDM: Electrical discharge machining HCl: Hydrochloric acid HNO3: Nitric acid MoO3: Molybdenum oxide Al: Aluminum C: Carbon Co: Cobalt Cr: Chromium

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20 Hf: Hafnium Mo: Molybdenum Ni: Nickel Nb: Niobium (also called Columbium, Cb) Re: Rhenium Ta: Tantalum Ti: Titanium W: Tungsten : Gamma phase (fcc Nickel solid solution, matrix) : Gamma prime phase (L12 ordered Ni3Al, precipitates) / : Used for discussion of both phases as a system (eg. the / interface) fcc: Face centered cubic L12: An ordered fcc-like structure MC: Type of carbide (M represents the metal atom, C represents the carbon atom) M6C: Type of carbide (M represents the metal atom, C represents the carbon atom) M23C6: Type of carbide (M represents th e metal atom, C represents the carbon atom) TCP: Topologically Close Packed (deleter ious phases that precipitate in some alloys: P, and Laves phases) a: Lattice parameter a: Lattice parameter of the phase a : Lattice parameter of the phase angle: The error (in degrees) of a bar from the [001] orientation at.%: Atomic percent : Lattice mismatch

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21 f: Failure strain FS: Failure strength k 1: The wavelength of the X-ray radiation caused by k 1 electrons k 2: The wavelength of the X-ray radiation caused by k 2 electrons f: Failure strength y: Yield strength RIA: Reduction in area T: Temperature Tm: Temperature of melting (absol ute temperature units only) UTS: Ultimate tensile strength wt.%: Weight percent YS: Yield strength : Angstrom (10-10 m) cm: Centimeter (10-3 m) GPa: Gigapascal (109 kgm-1s-2) hr: Hour in: Inch ksi: 1000 pounds per square inch (1000 psi) L: Liter lb.: Pound m: Meter m: Micrometer (10-6 m) min: Minute mL: Milliliter (10-3 L) MPa: Megapascal (106 kgm-1s-2)

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22 nm: Nanometer (10-9 m) s: Second V: Volt (m2 kg s-3 A-1)

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23 Abstract of Dissertation Pres ented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy THE PRIMARY CREEP BEHAVIOR OF SINGLE CRYSTAL, NICKEL BASE SUPERALLOYS PWA 1480 AND PWA 1484 By Brandon Charles Wilson May 2008 Chair: Gerhard Fuchs Major: Materials Science and Engineering Primary creep occurring at intermediate temperatures (650 C to 850 C) and loads greater than 500 MPa has been shown to result in severe creep strain, often ex ceeding 5-10%, during the first few hours of creep testing. This investig ation examines how the addition of rhenium and changes in aging heat treatment affect the primary creep behavior of PWA 1480 and PWA 1484. To aid in the understanding of rheniums role in primary creep, 3wt% Re was added to PWA 1480 to create a second generation version of PW A 1480. The age heat treatments used for creep testing were either 704 C/24 hr. or 871 C/32hr. All three alloys ex hibited the presence of secondary confirmed by scanning electron micros copy and local electrode atom probe techniques. These aging heat treatments resulted in the reduction of the primary creep strain produced in PWA 1484 from 24% to 16% at 704 C/862 MPa and produced a slight dependence of the tensile properties of PWA 1480 on aging heat treatment temperature. For all test temperatures, the high temperatur e age resulted in a significant decrease in primary creep behavior of PWA 1484 and a l onger lifetime for all bu t the lowest test temperature. The primary creep behavior of PWA 1480 and PWA 1480+Re did not display any significant dependence on age heat treatment. The creep rupture life of PWA 1480 is greater

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24 than PWA 1484 at 704 C, but significantly shorter at 760 C and 815 C. PWA 1480+Re, however, displayed the longest lifeti me of all three alloys at both 704 C and 815 C (PWA 1480+Re was not tested at 760 C). Qualitative TEM analysis re vealed that PWA 1484 deformed by large dislocation ribbons spanning large regions of material. PWA 1480, however, deformed primarily due to matrix dislocations an d the creation of interfacial dislocation networks between the and phases. PWA 1480+ contained stackin g faults as well, though they acted on multiple slip systems generating work hardening and forcing the onset of secondary creep. Xray diffraction and JMatPro calculations were also used to gain insight into the cause of the differences in behaviors.

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25 CHAPTER 1 INTRODUCTION Conventional wisdom regarding the first a nd second generation supe ralloys states that they are well known systems with li ttle left to learn about their be havior. Blanket statements are often used to generalize about these older alloy systems while cu rrent research and development continues to refine third, fourt h, and fifth generation alloys. Recently, however, renewed interest has been given to this class of alloys due to a curious phenome non during intermediate temperature creep. Within the temperature rang e of about 650C to 850C and under high stress, some second generation superalloys, bearing 1 atomic percent or a bout 3 weight percent rhenium, experience excessive primary creep. In these cases primary creep can be as high as 2830% in as little as 12 hou rs of creep testing. These stress and temperature conditions are impor tant as they are pres ent near the root, or attachment point, of turbine blades as well as within the internal suppo rt structure directing internal air cooling paths between the airfoil surf aces. If these regions deform at an excessive rate, internal stresses can result and failure can occur. It is also important to note that a turbine blade may fail through multiple methods. The most obvious is failure due to fracture. Another method is by exceeding dimensional tolerances. Gas turbine engines are designed with very tight tolerances. With turbine blades, the allowe d expansion due to creep and other processes is often only a few percent. These tolerances, theref ore, become threatened if some regions within a turbine blade can creep several percent in the first few hours of operation. There have been numerous attempts at identif ying the causes and controlling factors of primary creep in these alloys. Two suspected causes are rhenium content and the presence of secondary precipitates in the matrix channels. Secondary precipitates are produced following the last aging heat treatment. These precipitates are usually about 10-20 nm in size

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26 and populate the matrix channels that are about 100200 nm in diameter. It is widely documented that dislocations in supe ralloys are found predominately in the matrix. If this same region is populated by ve ry fine, densely spaced then dislocation slip in these channels may be greatly impeded. It has been reported th at the presence of th ese secondary particles raises the difficulty of shear in the phase to such a level that shearing of the primary precipitates becomes the preferre d method of dislocation motion. Once a dislocation pair enters the phase there is little barrier to glide until the opposite / interface is reached. These long glide paths, then, are possible cause s of large primary creep strains. Another source of matrix strengthening ta kes the form of solid solution strengthening brought about by additions of Re. Rhenium has been used for solid soluti on strengthening of the phase beginning in the 1980s. How Re achieves its strengthening effect has long been a topic for debate in the superalloy community. Recent research has focused on changes in mechanical behavior and the elemental segregation of Re relative to the / interface. The strengthening effect brought by the addition of Re was so signi ficant that the addition of Re alone was enough to define the second and third generations of superalloys. The most striking difference in behavior can be seen by comparing first and s econd generation alloys, or alloys with no Re to alloys with 3wt.% (1at.%) Re. The second generation alloys demonstr ated greater than 25 C improvements in creep and tensile strength capabilities.1-4 After over 20 years of research, the so called Rhenium Effect is still not fully understood. Adding impetus to the need to understand the Rhenium Effect, is the sudden climb of the cost of rhenium. Additions of rhenium are f ound in most of the alloys used for critical applications in aerospace gas turbines as well as some industria l gas turbines (IGT). Over the last 18 months, the price of rhenium has climbed from $600/pound to $2500/pound. This drastic

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27 increase in cost has many companies investigating alternative methods to st rengthen their alloys. One possible route includes replacing rhenium w ith another refractory element like tungsten, W. Also, if investigators could determine the eff ect rhenium has on disloc ation behavior, perhaps new strategies for strengthening could be developed. Interestingly, excessive primary creep and the Rhenium Effect may be related. The alloys most commonly associated with excessive primar y creep are rhenium bearing superalloys (i.e. 2nd, 3rd, 4th, and next generation alloys). Typical cr eep behavior for these alloys will display large primary creep strains in the first fe w hours of testing under high loads and low temperatures. Following primary creep is a sharp transition to secondary, or steady-state, creep. During secondary creep in these alloys the creep rate is ra ther low and unchanging. First generation alloys, however, tend to not display mu ch primary creep and th e secondary creep rate is higher and slowly increases fo r the duration of the cr eep test. Therefore, while rupture times for second generation alloys may be a full order of magnitude longer, time to 1% creep may be significantly shorter than their firs t generation counterparts. This be havior relates directly to the issue of dimensional tolerance in modern aerospace engines: wh ile second generation alloys will rupture at longer times, they may, in fact, fail dimensional controls significantly earlier than their rupture lives might suggest. Alternatively, this behavior might require unique design considerations during turbine run-in to account for excessive deformation. The current investigation focuses on commonly used first and second generation alloys: PWA 1480 and PWA 1484, respectively. A third, e xperimental alloy was also produced to isolate the effect of rhenium during intermedia te temperature creep by adding 3 weight percent Re to PWA 1480. The approach used to test the three alloys begins with a solution and homogenizing heat treatment for all three alloys to reduce the effects of segregation during

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28 solidification. Two differ ent aging heat treatments were then used with the intention of creating specimens with and without secondary in the matrix channels in order to observe changes in primary creep effected by these precipitates. Full length creep testing and interrupted creep testing was conducted to correlate microstructure with mechanical properties. Additionally, Xray diffraction studies, JMatPro thermodynamic predictions, and TEM techniques were used to investigate the primary creep behavior of the th ree alloys. Finally, th e problem of secondary and the Rhenium Effect with regard to excessive primary creep is discussed.

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29 CHAPTER 2 BACKGROUND Overview Modern superalloy technology has a long history of developments in processing, alloy chemistry, and fundamental metallurgical knowledge The advent of directionally solidified, followed by single crystal, blades and vanes led to major improvements in the high temperature tensile, creep, and fatigue properties of supe ralloys. Alloy refinement through advanced remelting processes and alloy chemistry has le d to consistency between master heats and improvements in mechanical properties and environmental resistance.2, 5 These developments occurred over several decades; however, the pace of research has occasionally exceeded the rate of application of these new all oys. For instance, a typical deve lopment cycle from inception to service for a new alloy can take 7-10 years. By the time that alloy is put into service, newer alloys with improved properties ar e already in development. Adding to this, the United States Department of Defense has instituted the Acceler ated Insertion of Materials (AIM) program, a Defense Advanced Research Projects Agency (DARPA ) initiative. The goal of this program is to decrease the time to active service of new alloys used for aerospace turbine engine applications.6 Similarly, the United States Department of Energy instituted a similar program in the early 1990s to improve performance of indu strial gas turbines. The Advanced Turbine Systems (ATS) program also created a push for new materials with higher temperature capabilities.7 These government programs coupled with the already strong drive to increase performance has led to rapid development cycles for new materials. While the high pace of development in the su peralloy industry is beneficial for engine manufacturers, customers, and na tional defense alike, new alloys are often developed before a firm understanding of previous performance gains could be achieved. A prime example is the

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30 addition of rhenium to second and third generatio n superalloys. Since the first implementation of Re to these alloys in the early to mid 1980s, the superalloys community has seen the development of 2nd generation (3 wt% Re), 3rd generation (6 wt% Re), 4th generation (with platinum group additions), and now early developmen t of the next generation of superalloys. As the momentum of research continues to advance alloy performance, the method by which Re improves the strength of superall oys is still debated. Early st udies examined microstructural effects, changes in stacking fault energy, changes in / lattice mismatch, and Re segregation to the phase during precipitation and coarsening of precipitates.8-11 Entering the 1990s, interest in exploring this effect was replaced by the need to boost engine performance. Thus the third generation of superalloys was created through even greater rhen ium additions, though the exact strengthening mechanism was still not entirely clear. The Rhenium Effect Twenty years later, there has been an increase in the interest in the so called Rhenium Effect. Part of this rekindled interest is due simply to the lag in development between military aircraft engines to commercial engines to industria l gas turbines, IGTs. Alloys that are new to commercial aircraft applications were first used in military aircraft engines 10 years earlier. The same gap exists between commercial engines and industrial gas turbines. These technological divides are due partly to market pressure a nd partly to processing challenges. New, high performance alloys are often very difficult to cast into a properly oriented single crystal. As the engine increases in size, its turbine blades must increase in size as well. Turbine blades on military aircraft may be 8 to 15 cm in length while commercial engines require 10 to 20 cm and IGTs require blades significantly greater than 30 cm in length. The consequence of these size differences is an increased rejection rate us ing conventional processi ng technologies. Thus,

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31 improvements in casting processes were often ne cessary before an established alloy could be applied or scaled up to new markets. Now that second generation single crystal al loys have made their way into commercial aircraft and IGT applications, in terest is growing to understand their behavior primarily due to cost pressures and increasing loads at intermediate temperatures. These new applications require larger components and, consequently, larger vol umes of Re bearing superalloys. The recent rapid rise in the cost of Re has the superalloy community inves tigating alternat ive strengthening additions. Additionally, some s econd generation alloys during in termediate temperature creep (650 C to 850 C) at high loads (greater than 500 MPa) will exhibit severe creep anisotropy that has been linked to stacking fault shear of the precipitates.12-15 Creep anisotropy has also been linked to large primary creep strains during testing under these same conditions. By convention, creep behavior is separated into three stages as shown in Figure 2-1. The primary creep stage of interest to the current investigation occurs fi rst. Primary creep is typically marked by a relatively large initi al creep rate due to the init ially undeformed nature of the microstructure. As deformation occurs, and hard ening takes place, the cr eep rate is reduced. Eventually, a balance between th e active deformation processes (d islocation slip) and recovery processes is reached and the creep rate stabilizes. This regime of creep is called secondary creep. Near the end of the life of a specimen, the total accumulation of deformation within the microstructure increases the rate of deformati on such that the active recovery processes can no longer maintain a steady-state balance. As a result, the creep rate begins to increase. Approaching failure, the creep rate is continually increasing and this behavior marks tertiary creep. The idealized behavior shown in Figure 2-1, however, is not always applicable. Often, usually at high temperatures, alloys will exhibit brief primary creep that transitions immediately

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32 into tertiary creep with no steady-state seconda ry creep stage. Anot her variation, which is discussed in this investigation, is creep dominated by the primar y creep stage. In specimens exhibiting this behavior, the prim ary creep stage is extended to gr eater levels of strain before secondary creep can occur. Additionally, with alloys exhibiting large primary creep stages, incubation periods of near zero creep usually preced ed the primary creep stage. It is thought that these incubation periods mark the forma tion of stacking fault ribbons prior to shear.16, 17 Figure 2-1. Idealized creep curve with the three creep stages labeled. The box marks the region of most interest to th e current investigation. The primary creep regime in some second genera tion alloys may be so severe that it is the dominant feature of their respec tive creep curves. As a result, primary creep has become the focus of several recent investigations because intermediate temperature, high stress creep conditions are common in IGT and co mmercial aerospace applications.16-20 This phenomenon may have gone unnoticed in the early part of the development cycle because the emphasis was placed on increasing service temperatures beyond 1000 C. Since this behavior is characteristic

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33 of lower temperatures similar to those seen in the attachment and internal support regions of turbine blades, it may also have been deemed less important than high temperature creep strength. Secondary Precipitates A microstructural feature that may impact the primary creep behavior of single crystal nickel base superalloys is the presence of secondary precipitates. While it has been long known that primary precipitate size has a major impact on mechanical properties,21, 22 research has also shown the importance of controlling the secondary or cooling precipitates20, 23. The final heat treatment step before a turbine blade is released for service is typically a aging heat treatment. This final step significantly im pacts the performance of the component. First, the time and temperature for the age heat treatm ent controls the coarsening behavior of the primary precipitates. Following solution heat treatment, these precipitates are mostly coherent in nature. Aging serves to increase the size of the precip itates to the optimum 0.3-0.5 m size range. Additionally, the cohere ncy of the precipitates is redu ced adding a misfit strengthening component to the overall strength characteristic of the alloy. Besides time and temperature, the cooling rate following aging has b een found to create microstructura l differences that can impact mechanical properties. Rapid cooling from the aging heat treatment temperature has been shown to produce a fine dispersion of secondary precipitates of the size ra nge 10-50 nm in diameter. These precipitates reside in the matrix channels between primary precipitates. Because dislocations tend to initiate in the matrix, it is expected that in teractions between dislocations and these precipitates would be common. Furnace cooling, however, from the aging temperature is capable of creating a matrix that is devoid of secondary precipitates.20

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34 While it is expected that a significant interaction exists between the secondary and the active deformation mechanism, exactly how the secondary effects the deformation process is not fully understood. These precipitates have even been linked to large pr imary creep strains as they may prohibit deformation in the matrix channels, forcing shear of primary precipitates resulting in creep anisotropy. A lternatively, it is thought that th ese precipitates may stabilize the stacking fault ribbons within the matrix, leading to enhanced shear and large primary creep strains.16, 17, 20, 24 Lattice Mismatch Another microstructural feature that is thought to impact prim ary creep is lattice mismatch (misfit). While it is known that lattice mismat ch can alter behavior for high temperature creep, specifically in the case of rafting,25, 26 lattice misfit also bears a significant impact on dislocation motion in the matrix and at the / interfaces.27 Lattice mismatch between the matrix and precipitates in supera lloys results in coherency strains along these interf aces. If the coherency strain is large enough, misfit dislocations will nucleat e and/or congregate at these interfaces to relieve the strain. Alloys with large lat tice misfit values have shown a propensity to form dense networks of interfaci al dislocations during creep tes ting. Conversely, alloys with reduced / misfit, exhibiting creep anisotropy, are characterized by relatively few interfacial dislocations and large regions of stacking faults in the precipitat es. This difference is possibly the result of compositional changes (which will ca use lattice mismatch m odification). In the case of CMSX-2 and CMSX-4, for instance, the most notable changes were the additions of rhenium and hafnium, an increase in cobalt, a nd a decrease in tungsten and chromium. Between the alloys PWA 1480 and PWA 1484 similar change s were made except for an increase in tungsten, a significant decrease in ta ntalum, and the removal altogether of titanium. Since lattice

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35 mismatch has been shown to be significan tly impacted by compos ition, and mechanical properties as a result, these changes from the firs t to second generation alloys can also impact the underlying deformation processe s that occur during creep.22 Therefore, any si gnificant change in composition increases the risk of changes in mechanical properties such as the increased primary creep strains discussed in the current investigation. Each of the aforementioned examples of alloy developments from the last 20 years results in changes in the intermediate temperature mech anical properties of these alloys. Besides the individual changes brought about by increasing rhenium content, secondary precipitation, and lattice mismatch variation, these three variations may all be inte rrelated. As is often the case in alloy development, changes in behavior are the results of multiple factors. While significant gains in performance may be made without a comp lete understanding of these behaviors, future gains may rely on applying lessons learned from i nvestigations into the Rhenium Effect, lattice misfit, and aging heat treatment (among ma ny other aspects of alloy design). Modern Superalloys Superalloys are known to combine a mix of hi gh temperature tensile and creep strength and environmental resistance. Their microstructu re consists of two phases with a face-centered cubic (fcc) matrix known as and a fcc-like L12 cuboidal precipitate known as They are strengthened through the us e of solid solution and precipitation hardening t echniques resulting in high strength over a wide temperature range. Typical ly they are also direc tionally solidified to produce a single grain through the Bridgman process, thus the application of the phrase single crystal to describe these alloys when processe d in this manner. Follo wing solidification, these alloys require long, high temper ature heat treatments for pr oper homogenization and aging. Several common alloy compositions representing the first thr ee generations of superalloy

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36 Table 2-1. Compositions in weight percent of several common Nickel-based superalloys3. Alloy Generation Co Cr Mo W Ta Re Al Ti Nb Hf Ni Pratt & Whitney PWA 1480 1st 5 10 4 12 5 1.5 Bal. PWA 1484 2nd 10 5 2 6 9 3 5.6 0.1 Bal. General Electric Rene N4 1st 8 9 6 6 4 3.7 4.2 0.5 Bal. Rene N5 2nd 8 7 2 5 7 3 6.2 0.2 Bal. Rene N6 3rd 12.54.2 1.4 6 7.2 5.4 5.8 0.2 Bal. CannonMuskegon CMSX-2 1st 5 8 0.6 8 6 5.6 1 Bal. CMSX-4 2nd 9 6.5 0.6 6 6.5 3 5.6 1 0.1 Bal. CMSX-10 3rd 3 2 0.4 5 8 6 5.7 0.2 0.1 0 Bal. development are given in Table 2-1. Due to the extreme demand of service conditions, the compositions of modern superalloys are quite complex, requiring 10 or more elemental additions. Additionally, a wide variety of elements not shown in Table 2-1 are commonly added to superalloys. These include, but are not limited to, carbon, boron, nitrogen, ruthenium, and some rare earth elements. The elemental addition that produced the grea test impact in properties during the early stages of superalloy development was rhenium. Beginning with the second generation of alloys, Re became an important factor for high temperatur e creep strength and microstructural stability. The second and third generations of superalloys contain 3wt.% and 6wt.% Re, respectively. The first generation of superalloys contai n no Re. They do, however, contain the strengtheners Al, Ti, and Ta and the refractory elements W and Mo for solid solution strength and microstructural stability. As development cont inued, the need to restrict coarsening due to thermal exposure became apparent. This led to the addition of Re to superalloys for its low diffusivity that

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37 restricts the growth of primary precipitates. The diffusivity of rhenium in Nickel is so low that it significantly retards the diffusi on controlled coarsening of the precipitates. Additionally, Re was found to be a very potent solid solution strengthener, thus the second generation of superalloys was created.8, 10, 11 Continued development saw the addition of gr eater Re additions re sulting in greatly improved temperature capabilities and strength. The third gene ration contains 6wt.% and marks a significant improvement in proper ties relative to the ge nerations that preceded it. This increase in Re concentration, however, came with a cost. Castability, as measured by incidence of casting defects, was greatly decreased a nd the propensity for the formati on of deleterious topologically close packed (TCP) phases was greatly increased.1, 2, 8 Current research in fourth generation alloys and beyond is focused on continuing to improve strength while minimizing the harmful side effects of large Re a dditions through additions of pl atinum group metals, such as ruthenium.28 New research was also initiated to fi nd ways to maintain strength while reducing Re concentrations due to the drastic increas e in the cost of Re over the last 2 years.29 The state of the art of modern superalloys relies on a variety of alloy strengthening and processing techniques developed over the last 30 years. As th e high performance and low cost demands continue to grow, the technology will need to grow. The next discussion will review the challenges of these goals from the alloy and processing viewpoints. The Challenge of High Temperature Service The service environment experienced by turbin e blades in modern gas turbine engines demands high strength and environmental resistance at high temperatures. Normal alloy systems are not capable of withstanding th is type of environment. Structural alloys are typically strengthened by a mix of solid solution strengthening, prec ipitation hardening, and strain

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38 hardening. At temperatures a bove half the melting point (T>0.5Tm) strain hardening is less effective due to recovery a nd recrystallization processes.30 Additionally, nor mal alloy systems that are precipitation hardened lose strength at the temperatures experienced in gas turbine engines because the precipitate solvus is exceede d, yielding a single phase ma terial. As a result, most normal alloy systems are insufficient for these applications. In addition to tensile strength are several othe r properties to consider : creep strength, high cycle and low cycle fatigue, ductility/toughne ss, thermo-mechanical fatigue, oxidation resistance, and corrosion resistance. Each of these properties will impact the usefulness of an alloy during high temperature service. Because th ese alloys are subjected to high temperatures and high stresses for long durations during servi ce, creep has been the subject of numerous investigations over the past 40 years. At low temperatures, deformation processes are controlled by dislocation motion (whole and pa rtial dislocations). The m ovement of dislocations along glide planes is governed by a variety of materi al properties like elastic modulus, dislocation friction sources (peierls stresses), the presence of solute atmospheres (Cottrell atmospheres), cross-slip difficulty, stacking fault energy, and anti-phase boundary energ y. Add to these the presence of secondary reinforcing phases a nd interfacial dislocation networks and the intermediate temperature deformation processes of single crystal nickel base superalloys become quite complex.31 Tensile Behavior Low temperature deformation of superalloys is dominated by di slocation motion and stacking fault propagation. At low temperatures, there is insufficient energy for diffusion controlled processes and cross-slip is also more difficult, though not impossible. At high temperatures, however, diffusion c ontrolled processes become active. These include recovery,

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39 vacancy motion, solute motion (solute drag), and phase instabilities ( coarsening, TCP precipitation, and carbide transitions).32-35 Despite all of the advances in superalloy technology, the basic mi crostructure employed for single crystal turbine blades is quite simple. Generally, superalloys c onsist of an fcc matrix with coherent, stress-free ordered L12 precipitates. Most first and second generation alloys contain about 60-70% : the optimum value for excellent creep and tensile properties. Maintaining low misfit through alloying has also b een shown to effect change in the morphology of the precipitates. Even alloys with high lattice misfits ( >0.5%) will exhibit less than one percent misfit between the and phases. Changing misfit values can produce precipitate shapes vary ing from spherical to cuboidal to dendritic. The so-called cuboidal morphology is typically desired for high temperatur e creep resistance with the ideal precipitate edge length between 0.35m and 0.45m.36 The cuboidal structure cons ists of uniform cubes of with rounded edges and corners. The resulting interstices between particles consists of the matrix. These channels are relatively large in le ngth and width, but narrow in thickness (typically <50nm). Strengthening of superalloys can be achieved through alloying and heat treatment. The most common goals involved in strengthening superalloys through alloying can be broken down into several categories: Solid solution strengthening of matrix Modifying misfit to produce cuboidal precipitates Increasing volume fraction to the ideal limit Solid solution strengthening of precipitates Modifying the anti-phase boundary, APB, energy of the phase Stabilizing the microstructure with low diffusi vity elements (W, Mo, and Re for example)

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40 The addition of carbide phases may also help with fracture strength, but is not likely to affect yield strength due to their large spacing compared to the dislocation spacing.37 Another consideration is the preventi on of TCP phases. Alloys with large refractory element concentrations are at risk of developing TCP phases during long exposure to high temperatures. While the prevention of TCP phases does not ne cessarily lead to im proved strength, the formation of TCP phases is suspected to result in premature failure of s uperalloys and should be avoided.35 Solid solution strengthening of the matrix is typically accomplished by increasing additions of Cr, W, Mo, and/or Re. While these are suitable for solid solution strengthening, secondary concerns may dictate their use. Increa sing the W, Mo, and Re content of an alloy may result in decreased TCP resistance and an increase in the / misfit as these elements are mostly rejected from the phase.8, 35 Increasing the volume fraction and strength of the precipitates is typically achieved by additions of Al, Ti, and Ta. Controlling the / misfit and APB energy is a difficult and not fully unde rstood process; however. They are usually maintained through slight adjustment to the entire set of elementa l additions until the desi red value is achieved as virtually every element can affect misfit and/or APB energy. Finally, on e of the benefits of adding slow-diffusing elements like W and Re for so lid solution strengthening is that they also serve to stabilize the / microstructure. With a diffusivity for self-diffusion within nickel that is much lower than any other elements, coarsening kinetics are reduced allowing for precipitate coherency to be maintained longer.8 Tensile strain of superall oys usually begins in the matrix phase. Dislocations are created on {111}<110> slip planes which are then able to propagate through the material. As in other fcc materials, these dislocations will often disso ciate into Shockley part ial dislocations of the

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41 a/6{111}<112> type with a stacking fault formed in between. As the strain increases, these dislocations will squeeze through the matrix channels as described by the orowan bowing model. As this process progresses, dislocatio ns begin to lie along the surfaces of the cuboidal precipitates eventually forming interfacial networks of dislocations. Cont inued strain gradually reduces the distance between disloc ations in the interfacial netw orks. As the distance gets smaller, the strength increases due to work-hardening.24, 27, 33, 37 The tensile behavior of single crystal, nickel base superalloys may also display a yield-point on an engineering stress vs. engineering strain graph. This phenomenon is not unexpected as it is commonly linked to interstitial or substitutional impurities and th ese alloys contain both in the form of alloying additions.31 Creep Behavior When superalloys are subjected to tensile loads at temperatures greater than 500 C, creep can occur. Creep is a time-dependent plastic de formation process that occurs at loads below the yield stress of the material in que stion. Creep occurs because the thermal energy that is available at high temperatures allows for thermally ac tivated processes that do not occur at low temperatures to become viable deformation mechanisms. In the case of single crystal superalloys, vacancy and solute diffusion, disl ocation climb, cross-s lip, and activation of secondary slip systems are all processes that are either difficult or non-existen t at low temperatures. Additionally, microstructural evolution may occur resulting in changes to the creep behavior over time. Superalloys are pron e to the following microstructural changes in particular: coarsening and loss of coherency, t opological inversion (rafting) of the phase, precipitation of additional phases (especially t opologically close packed, TCP, phases), and carbide transition/precipitation (MC, M6C, and M23C6).35, 38-41 It is the activation of these

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42 processes that allow for metals to deform over ti me at loads significantly lower than their yield stress.25, 31, 33, 42, 43 Mechanically, creep tests can be conducted in one of two ways. To obtain a true measure of creep requires that the speci men be subjected to a constant stress throughout the duration of the test. Consequently, as the specimen deforms and the cross-sectional ar ea decreases, the load would be reduced accordingly. Unfortunatel y, though, creep tests must be run at high temperatures within a furnace so accurate meas urements of the specimen dimensions are not feasible during testing. Because of this problem, another test pr ocedure has been adopted for use in most cases. This is the constant load creep test (contrasting with the constant stress test discussed above). A constant lo ad creep test is performed by applying the correct amount of weight to the specimen to produce the desired initial stress. As the specimen deforms and the cross-section is reduced, no changes are made to the load. The cons equence of this design is that the stress level increases throughout the creep test. This investigation uses constant load creep tests because of the simplicity of data acquisition and commercial relevance, so the following discussion refers to data collected with this type of creep test.31 For single crystal, nickel base superalloys, there are primarily two temperature regimes of interest that involve creep. At high temperatures, greater than 850 C, creep tests are run with relatively low stresses and are used to simulate the behavior of the thin, exposed parts of a turbine blade. This environment allows for th ermally activated processes to occur at a rapid pace. At temperatures below 850 C, the loads are significantly greater. The performance of alloys under intermediate temperatures and high loads depends largely on the tensile strengths of the alloys being investigated. Because of the link between high loads and intermediate

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43 temperatures to large primary creep strains, this set of conditions is the focus of this investigation. Summary In order to conduct an investig ation into the primary creep behavior of the superalloys PWA 1480 and PWA 1484, specimens will be creep tested at temperatures between 700 C and 815 C and loads greater than 621 MPa. Tensile testing, metallographic examination, X-ray diffraction (XRD), and transmission electron microscopy (TEM) have been conducted to gain an understanding of the underlying factors involved in creating large primary creep strains. Three alloys have been investigat ed including the first genera tion PWA 1480, the second generation PWA 1484, and the experimental second generation PWA 1480+Re. These three alloys helped to shed light on several aspects of primary creep behavior and the effect of age heat treatment. Finally, the current investigation may lead to an improved understand ing of the effect of rhenium on single crystal nickel base superalloys.

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44 CHAPTER 3 EXPERIMENTAL PROCEDURES Materials All three alloys used in this investigati on were provided by Pratt & Whitney Aircraft Engines in East Hartford, CT. Test bars were prepared by Pratt & Whitney in investment casting cluster molds of 12 bars measuring approximately 20 cm (8 in.) long with a 1.6 cm (0.625 in.) diameter. Single crystals were prepared using Bridgman style directional solidification furnaces and single crystal selectors to produce an or ientation in the [001] direction. Following solidification, the bars were exam ined through the use of Laue X-Ra y Diffraction to verify that all the bars were near the desi red [001] orientation. Laue orientation data and original heat treatments for all the material is presented in Tables 3-1, 3-2, and 3-3. Test material was Table 3-1. Laue orientation data and original heat treatments for PWA 1480.* Master Heat Bar Number angle Solution Heat Treatment Coating Heat Treatment P9976 0401 4.2 1288 C / 2hr. 1080 C / 4hr. P9976 0402 5.2 1288 C / 2hr. 1080 C / 4hr. P9976 0403 6.2 1288 C / 2hr. 1080 C / 4hr. P9976 0404 7.3 1288 C / 2hr. 1080 C / 4hr. P9976 0405 8.8 1288 C / 2hr. 1080 C / 4hr. P9976 0406 9.5 1288 C / 2hr. 1080 C / 4hr. P1083 0901 2.5 As-Cast As-Cast P1083 0902 2.1 As-Cast As-Cast P1083 0903 19.3 As-Cast As-Cast P1083 0904 4.0 As-Cast As-Cast P1083 0905 3.0 As-Cast As-Cast P1083 0906 0.9 As-Cast As-Cast P1083 0907 4.8 As-Cast As-Cast P1083 0908 4.5 As-Cast As-Cast P1083 0909 3.1 As-Cast As-Cast P1083 0910 3.9 As-Cast As-Cast P1083 0911 2.3 As-Cast As-Cast P1083 0912 4.5 As-Cast As-Cast *Bar number 0903 was rej ected due to an excessive value.

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45 Table 3-2. Laue orientation data and original heat treatments for PWA 1484. Master Heat Bar Number angle Solution Heat Treatment Coating Heat Treatment P1096 0407 0.5 1310 C / 0.5hr. 1080 C / 4hr. P1096 0408 3.8 1310 C / 0.5hr. 1080 C / 4hr. P1096 0409 5.7 1310 C / 0.5hr. 1080 C / 4hr. P1096 0410 7.2 1310 C / 0.5hr. 1080 C / 4hr. P1096 0411 9.4 1310 C / 0.5hr. 1080 C / 4hr. P1096 0412 10.1 1310 C / 0.5hr. 1080 C / 4hr. P1086 0913 4.0 As-Cast As-Cast P1086 0914 7.0 As-Cast As-Cast P1086 0915 5.7 As-Cast As-Cast P1086 0916 7.5 As-Cast As-Cast P1086 0917 8.5 As-Cast As-Cast P1086 0918 7.0 As-Cast As-Cast P1086 0919 5.6 As-Cast As-Cast P1086 0920 8.1 As-Cast As-Cast P1086 0921 5.9 As-Cast As-Cast P1086 0922 3.1 As-Cast As-Cast P1086 0923 6.5 As-Cast As-Cast P1086 0924 5.9 As-Cast As-Cast Table 3-3. Laue orientation data and original heat treatments for PWA 1480+. Master Heat Bar Number angle Solution Heat Treatment Coating Heat Treatment P1106 1201 12.6 As-Cast As-Cast P1106 1202 15.5 As-Cast As-Cast P1106 1203 5.7 As-Cast As-Cast P1106 1204 1.6 As-Cast As-Cast P1106 1205 9.3 As-Cast As-Cast P1106 1206 10.2 As-Cast As-Cast P1106 1207 9.0 As-Cast As-Cast P1106 1208 0.3 As-Cast As-Cast P1106 1209 9.1 As-Cast As-Cast P1106 1210 12.6 As-Cast As-Cast P1106 1211 4.2 As-Cast As-Cast P1106 1212 3.1 As-Cast As-Cast provided in two stages. First, a supply of 6 bars of PWA 1480 and 6 ba rs of PWA 1484 was delivered having been subjected to the standard commercial solution heat treatments (alloy

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46 specific) used by Pratt & Whitney as well as a coating heat treatme nt cycle. The bars were not, however, subjected to an aging heat treatment. The second delivery of material expanded the testing possibilities sign ificantly. First, another 11 bars of PWA 1480 and 12 bars of PWA 1484 were provided in the as-cast c ondition. Then, 12 bars of PWA 1480 with 3 wt% Re added were produced and delivered in the as-cast condition. Altogether, 47 single crys tal test bars, allowing for the production of up to 94 creep/tensile spec imens, were available for experimentation. Heat Treatment The first step in the preparation of test spec imens was to heat treat all the bars to an appropriate condition. As stated above, the first batch of PWA 1480 and PWA 1484 had already been given solution heat treatments (HT1), Tabl e 3-4. Additionally, this first batch had already received a coating simulation cycle as these allo ys are commonly used in the first and second turbine stages of gas turbine engines where coa tings are necessary for thermal and environmental protection. Because one aim of th is study is to determine the cause of changes in creep behavior due to age temperature, two differe nt aging heat treatments were gi ven to each alloy. Half of the PWA 1480 bars were given a low temperature age (LT, 704 C/24 hr.) and half were given a high temperature age (HT, 871 C/32 hr.) as shown in Table 3-4. The same aging heat treatments were also given to the PWA 1484 bars. Following aging, this first batch of material was ready for sample machining and subsequent testing. Heat Treatment Development The remaining test bars, as-cast PWA 1480, as-cast PWA 1484, and as-cast PWA1480+, still required a complete heat treatment cycle. Following early results and metallography of the first batch with the standard solution heat treatments, it was determined that a longer time, higher temperature solution heat treatment might be helpful to reduce the effects of chemical

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47Table 3-4. Heat treatments used in this study. Solution heat treat ment designations (HT#) are given.* First Batch Alloy Solution Heat Treatment Coating Heat Treatment LT Age HT Age PWA 1480 HT1 1290C/2hr./GFQ 1080C/4hr./ GFQ 704C/24hr./AC 871C/32hr./AC PWA 1484 HT1 1310C/0.5hr./GFQ 1080C/4hr ./GFQ 704C/24hr./AC 871C/32hr./AC Second Batch Alloy Solution Heat Treatment Coating Heat Treatment LT Age HT Age PWA 1480 HT0 None HT1 1290C/4hr./GFQ (STD HT) HT2 1290C/8hr./GFQ HT3 1295C/1hr. 1299C/1hr. 1080C/4hr./GFQ 704C/24hr./AC 871C/32hr./AC 1302C/10hr./GFQ PWA 1484 HT0 None HT1 1315C/4hr./GFQ (STD HT) HT2 1315C/8hr./GFQ HT3 1325C/1hr. 1332C/1hr. 1080C/4hr./GFQ 704C/24hr./AC 871C/32hr./AC 1338C/10hr./GFQ PWA 1480 + HT0 None HT1A 1290C/4hr./GFQ HT1B 1315C/4hr./GFQ HT2 1280C/1hr. 1290C/7hr./GFQ HT3 1293C/1hr. 1296C/1hr. 1080C/4hr./GFQ 704C/24hr./AC 871C/32hr./AC 1299C/2hr. 1302C/1hr. 1308C/4hr./GFQ *All samples were heated from room temperature to 1200 C for 0.5hr. at a rate of 20 C/minute, then samples were heated to the first soak temperature at a rate of 10 C/minute. All ramping following the first soak temperature is at a rate of 1 C/minute. GFQ is Gas Furnace Quench (u ltra high purity He) and AC is Air Cool. Solution and Coating heat treatments were performed under vacuum and Age heat treatments were performed in air environme nts.

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48 segregation and the formation of eutectic regions during solidifi cation. Previous research has also shown that incomplete solutioning and dendrite homogeniza tion can lead to substantial reductions in creep life.36, 44 To counter this effect, a series of heat treatment trials beginning with the standard Pratt & Whitney devel oped heat treatments was performed. The first step taken to design a new solution he at treatment was to increase the amount of time of the original heat treat ment. Increasing the maximum temperature hold time from 2 hours to 8 hours for PWA 1480 and 0.5 hours to 8 hours for PWA 1484 did not significantly improve the homogeneity of the samples; however, it prov ed a useful heat treatment for Differential Thermal Analysis (DTA). With samples in the as-cast condition DTA is only moderately useful as the solidus and solvus temperatures tend to be suppressed and transi tion temperature peaks tend to be broad due to high segregation. Fo llowing even a brief heat treatment, however, meaningful data can be obtained as described below.45 Differential Thermal Analysis Differential Thermal Analysis was performed on all three alloys following the extended version (HT2) of the standard solution heat treatment (HT1). DTA was performed by Dirats Laboratories in Westfield, MA using a DuPont 9000 Thermal Analyzer. Data was collected at a rate of 20 C/minute on heating only to eliminate the effects of undercooling. Samples were cut from the single crystal bars to produce a disk approximately 0.5 cm (0.188 in.) thick by 1.6 cm (0.625 in.) in diameter. The HT2 treatment was a pplied to all three all oy groups for DTA. All samples were run with pure nickel standards. Additional discussion of DTA practices is provided in Appendix A along with the complete DTA scan profiles generated for each sample. The resulting data allowed for determ ination of the solidus, liquidus, and solvus temperatures for the three alloys and are gi ven in Table 3-5. Clearly, the increased homogeneity of the

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49 Table 3-5. DTA results for all three alloys as -received (HT1/HT0) and heat treated (HT2).* Alloy PWA 1480 PWA 1484 PWA 1480+ Heat Treatment HT1 HT2 HT1 HT2 HT0 HT2 Solvus (C) 1288 1298 1291 1300 N/F 1298 Solidus (C) 1304 1307 1338 1346 1291 1311 Liquidus (C) 1333 1344 1378 1392 1339 1352 *The solvus was not found for PWA 1480+, HT0. samples following the HT2 heat treatment resulted in an increase in all of the transformation temperatures. Using these temperatures as a guide, a new heat treatment (HT3) scheme was developed to improve the degree of solutioning and homogenization. The third heat treatment include d brief hold steps at a few temp eratures prior to the final soak temperature. Lower temperature steps were us ed in the early stages of the heat treatment to prevent incipient melting in the in terdendritic regions of the alloy where the solidus temperature is suppressed. The final soak was designed to be as close to the solidus of each of the alloys as possible in order to give the be st possible degree of homogenizati on within the alloy. The final heat treatment (HT3) signifi cantly reduced the presence of / eutectics and increased the degree of homogenization within the alloys. Furnaces Solution and coating simulation heat treatments were given to all three alloys using an Elatec high temperature vacuum furnace. The furnace is capable of temperatures up to 1400 C under a vacuum greater than 10-4 Torr. The heating elements, hearth plate, and jail are fabricated from graphite. Alumina trays were used to separate the test bars from the graphite hearth plate. Temperature control is provided by three type C (W + 5% Re / W + 26% Re) thermocouples all sheathed in individual molybdenum jackets. Tw o thermocouples were lowered very near the surface of the test bars in the center of the hear th plate. These were used for controlling the

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50 furnace temperature to within approximately .5 C of the setpoint. The third thermocouple is a survey thermocouple used for monitoring the te mperature near the front of the hot-zone. Following all heat treatments in the vacuum fu rnace, samples and test bars were cooled by injection of ultra high purity helium and the circulation of th e helium through a water cooled copper heat exchanger within the furnace by a fa n. This rapid quenching technique allows for high cooling rates greater than 250 C/minute down to approximately 500 C. This rapid cooling rate minimizes the precipitation and growth of during cooling. Below 500 C, cooling slows to between 100-200 C/minute, however, diffusion in the sa mples at this point has slowed sufficiently to prevent significant coarsening. Furnace control is managed by a Honeywell controller that manages temperature, vacuum, gas quenching, and all valves. Aging heat treatments (both LTA and HTA) were performed in Carbolite box furnaces with maximum operating temperatures of 1300 C. The atmosphere is not controlled for these heat treatments as the temper ature is typically low enough and the time is short enough that oxidation is not problematic. Te mperature control was maintained through the use of two Type K thermocouples placed in direct contact with the bars inside the furnace. Temperatures in the box furnaces were maintained within C of the setpoint as verified by handheld digital thermometers. Following aging heat treatments, samples and test bars were removed and allowed to air cool on alumina r acks. Cooling rates exceeded 100 C per minute to prevent undesirable coarsening. Characterization Throughout this investigation, sample characte rization was crucial to understanding what changes had taken place in the alloys. A variety of characterization methods were used to examine samples including the already mentioned Differential Thermal Analysis. Other methods

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51 include metallographic techniques involving op tical and electron microscopes, fractography following creep and tensile testing, and transm ission electron microscopy (TEM) on interrupted creep specimens to examine deformation structures. Preparing Samples for Metallography Samples were taken from test bars and creep/tensile specimens for metallographic examination. The same basic process can be applied to a variety of sample shapes and geometries. Metallographic pr eparation was conducted accord ing to the ASM Metals Handbook recommendations. Samples were first sectione d from the bulk by the use of either a Leco abrasive cut-off saw or an Allied slow-cut diamond sectioning saw. Both saws are liquid cooled to keep cutting temperatures low. Once secti oned, the specimens were ground flat and polished. Specimens were first leveled using silicon carbide grinding papers on an 8 inch Leco metallography wheel, then polished using al umina powder and water suspensions. All metallography specimens were polished to 0.3 m alumina polishing media. Following polishing, the specimens were etched with a etchant developed by Pratt & Whitney (100 mL HCl, 100 mL HNO3, 10g MoO3, 100 mL H2O). The etchant was applied with cotton tipped applicators and was swabbed evenly about the su rface until the surface of the specimen appeared hazy. Specimens were then rinsed with water, then methanol, and were finally dried under a jet of compressed air. Additionally, a second etchant was used to etch away the matrix phase. This etchant was an electrolytic etch utilizing a 20% oxalic acid solution at 20 V. The specimens were mounted on a metal stub to maintain the necessary electri cal conductivity from the specimen to the power supply and were connected as the anode. The counter electrode (cathode) was a 500 mL stainless steel beaker. The electrolyte was placed in the beaker and the steel beaker was placed

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52 in an ice water bath to keep the electrolyte from heating. The voltage was applied and the specimens were dipped for no longer than 5 sec onds. The specimens were immediately rinsed and dried and were ready fo r metallographic inspection. All Specimens were then observed on a Leco optical metallograph, a JEOL 6400 tungsten filament scanning electron microscope, or a JEOL 6335F field emission scanning electron microscope. Both electron microscopes were equipped with Energy Di spersive Spectroscopy (EDS) detectors as well. Images taken on thes e microscopes were used for examination of microstructures during the heat treatment development cycl e. Additionally, post test microstructures and fracture surfaces were observed to aid in the understanding of microstructural changes duri ng high temperature testing. Preparing Samples for TEM Making specimens for transmission electron microscopy (TEM) involves the same metallographic techniques described above. Interrupted and full-length creep tests were sectioned using the Allied slow-c ut diamond saw in the transverse and longitudinal directions. These samples were mounted on aluminum stubs with mounting wax and polished to 0.3 m alumina powder suspension. The samples were then transferred to special mounts for use with the Focused Ion Beam (FIB) or were thinned fo r use with the twin je t electropolisher. The FIB is a scanning electron microscope with an attached Gallium ion beam for milling the sample. Using the FIB, TEM liftout specimen s were cut perpendicular, parallel, and at 45 to the stress axis. On PWA 1484 samples, it was possible to obtain TEM specimenss near [001], [011], and [111] zone axes from a single sample due to the orientation based distortion of PWA 1484 during creep testing described in later chapters. Lift-outs were placed on 200 mesh Carbon coated Copper girds usin g micromanipulators.

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53 Alternatively, specimens were also prepared using a twin jet el ectropolisher. After samples were sectioned from the interrupted cr eep specimens, as described above, they were thinned by hand to between 100 and 200 m. Disks of 3 mm diameter were then removed from the thinned material with a TEM punch. The samp les were further thinned using a die polisher with a micrometer to between 30 and 70 m thic k. Electropolishing was performed on the 3 mm disks with a solution of 90% methanol and 10% perchloric acid and a vol tage of 20V. Liquid nitrogen was added to mainta in the temperature of -25 C. Transmission electron microscopy was conducted using JEOL 200 CX. The TEM wa s used to for qualitative observation of microstructural changes, disl ocation and stacking fault be havior, and variations in / misfit. Preparing Samples for LEAP Local Electrode Atom Probe (LEAP) specimens were prepared from a single bar of PWA 1484 that had already received a solution heat tr eatment (HT3) and a coating simulation heat treatment. The bar was first cut in four equa l pieces on a water cooled Leco abrasive cut-off wheel. Then, two sections (out of four) were subjected to either the LTA or HTA aging heat treatments. One piece from both groups was then separated and subjected to a brief heat treatment above the solvus (30 minutes at 1310 C) in order to redissolve the and quickly cool before the slow diffusing rhenium atoms have enough time to redistribute, Figure 3-1. This heat treatment was performed in an attempt to form a better unde rstanding of rhenium segregation behavior at the / interfaces. Once the samples were heat treated, they were sent to AMT, Clifton Park, NY to be sectioned into specimens by electrical discharge machining (EDM). Specimens were 5 cm long cylinders wi th a diameter of 1.8 mm. Ten specimens were obtained from each of the four samples mentioned above.

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54 Figure 3-1. Flow-chart of heat treatments for LEAP samples. The LEAP specimens were then electropolished to reduce the diameter of the cylinder and form a point on one end. The electropolishing unit consisted of a Te kPower HY 3005D-3 power supply with 2 variable outputs each capable of 30 V and 5 A. If linked in series or parallel it was capable of 60 V at 5 A or 30 V at 10 A, respectively. The cathode was a 500 mL stainless steel beaker in an ice bath to contro l the electrolyte temperature. Th e electrolyte was 90% ethanol and 10% perchloric acid. The specimen was connect ed to the power supply as the anode and was held in the beaker by a brace such that the cy linder was centered along its length to reduce the effect of varying anode to cathode distance. Th e applied voltage was va ried between 5 V and 20 V until stable polishing took place (typically between10 V and 15 V). Polishing was stopped once the specimens reached a diameter around 50 0 m to 700 m, which correlates to a tip diameter of approximately 100 m. The specimens were then shipped to the Mate rials Science and Engineering Department at the University of North Texas (Denton, TX). Anantha Puthucode, Ph.D., and Michael Kaufman, Ph.D., of the University of North Texas provided final sample preparations using an Imago

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55 Electropointer. The electropointer used a 3 mm Pl atinum loop as a cathode to refine the tip radius to less than 75 nm, Figure 3-2. Additional sharpening was occasionally performed using FIB based techniques as well. Once the tip wa s formed, specimens were run on the Imago LEAP 3000X. Data was then analyzed and manipulated using the Imago IVAS analysis software package. Specimens were run multiple times by reforming the tip following the experiment. Data from these experiments were returned in the form of 3-dimensional and 2-dimensional composition maps with a lateral resolution of 0.1 nm. Additionally, 1-dimensional line scans were also simulated using the same data sets. Figure 3-2. SEM micrograph of the tip of a LEAP specimen. Preparing Samples for X-Ray Diffraction Specimens were also prepared for X-Ray Diffraction (XRD) to study the / misfit in all three alloys. Following the coating heat treatment cycle, bars from each alloy were sectioned to create 4 disks of each alloy 16 mm in diameter by 10 mm thick. These disks were then separated into 4 groups. The first group was tested with no further heat treatment resulting in a total heat treatment time at 1080 C of 4 hours. The other three groups were heat treated an additional 6,

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56 96, or 996 hours to produce samples with total heat treatment times at 1080 C of 10, 100, and 1000 hours. Specimens were cut from the disks at a thickness of 2 mm. The specimens were then thinned and polished on both sides to creat e a final thickness between 0.3 and 0.7 mm. The polished surfaces were both made parallel to the (001) plane. All samples were run on Rigaku /2 powder diffractometers at the University of Central Floridas Advanced Materials Processing and Analysis Center (AMPAC). Routine scans were made from 10 to 130 at a scan step of 0.02 and a 1 second dwell time Local scans of the (002) peak, centered around 2 = 51 and the (004) peak, centered around 2 = 118 were taken with a smaller scan step of 0.01 and a longer dwell time ranging from 3 to 5 seconds. Some specimens were also run at 0.01 steps and 30 second dwell times in an attempt to improve signal clarity. The resulting inte nsity data were then indexed. The local scans of the (002) a nd (004) peaks were used for lattice mismatch calculation. Due to the structure factors and lattice parameters of the fcc phase and the L12 phase, the peaks produced from both phases overlap at the (002) and (004) positions. In order to separate the contributions, peak deconvolution was perfor med utilizing the MDI Jade (ver. 7) software package and in agreement with published work.46 Deconvolution practices included fitting skewed and unskewed Gaussian profiles to account for both phases as well as for the Cu k 1 and Cu k 2 wavelengths. A complete discussion of the deconvolution procedure and results are given in Appendix B. Once the centers of the k 1 peaks for both phases ar e determined, lattice parameters and mismatch were determined. Latti ce mismatch values in this investigation were calculated using Equation 3-1. Where is lattice misfit (multiply by 100 to covert to percent), a is the lattice parameter of the phase, and a is the lattice parameter of the phase.

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57 Equation 3-1 )( )(2 aa aa JMatPro Thermodynamic Prediction The thermodynamic modeling and prediction software package, JMatPro (United Kingdom), was used to predict phase compositions, lattice misfit, and volume fractions of phases as a function of temperature. While these predic tions are useful and fast, they are not guaranteed to be accurate for all alloy compositions. To enhance usability and accuracy, the models used for JMatPro have been calibrated through the use of several common benchmark alloys. PWA 1480 is one such benchmark alloy. PWA 1484 and PW A 1480+ predictions, as a consequence of not being used to calibrate the softwa re results, will require some degree of extrapolation to make thermodynamic predictions such as cooling curves expected compositions, and lattice misfit. Once an alloy composition has been selected, th ermodynamic data can be viewed as a function of temperature or calculated at specific temperatur es. Data from these predictions were used as a helpful guide and as potential indicators of intere sting behaviors. Conclusions based on software results alone, however, should be verified thr ough experimental testi ng and, for this reason, characterization was conducted to verify se veral aspects of the JMatPro predictions. Mechanical Behavior Following the final aging heat treatment, the test bars were sent to Joliet Metallurgical Labs in Joliet, IL for final machining. Each bar was machined into two creep/tensile specimens according to the drawing in Figure 3-3. The same specimen geometry is suitable for either high temperature tensile testing or creep testing. The specimen gauge length was 2.60 cm and the gauge diameter was 4.5 mm. Specimens were m achined from the test bars using low stress grinding techniques to prevent strain hardeni ng along machined surfaces. Following machining,

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58 all specimens were measured to verify all dimens ions of interest, most notably gauge length and gauge diameter. Tensile Testing High temperature tensile te sting was conducted at 704 C and 815 C to compare the strengths of the three alloys with both age heat treatments. The tens ile test matrix can be seen in Table 3-6. Fixtures were lubricated with boron nitride high temperature lubricant and the specimens were threaded into the grips. An Instron servo-hydraulic load frame with a 20,000 lb. load cell was used in coordination with the Merlin control and data acquisition software package. Figure 3-3. Creep and tensile sp ecimen geometry and dimensions.

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59 Table 3-6. Tensile test matrix with sample identification. Age HT LT LT HT HT Test Temp. ( C) 704 815 704 815 PWA 1480 0402-2 0403-2 0405-2 0406-2 0905-1 0901-2 0906-1 0906-2 PWA 1484 0408-1 0409-2 0411-2 0411-1 0915-1 0915-2 0919-1 0919-2 PWA 1480+ 1205-1 1203-2 1207-1 1209-2 A clam-shell style furnace capable of temperatures up to 1000 C was used to control the temperature to within .7 C of the setpoint throughout the test. Temperatures were maintained by two Type K thermocouples with Nextel hi gh temperature ceramic insulation for thermal protection. Both thermocouples were tied to th e gauge section of the specimen with 24 gauge 80Ni-20Cr wire. Prior to starti ng the test, all specimens were subjected to a 15 minute soak period to ensure uniform temperature of the sample and fixtures. All tensile tests were performed in air. During tensile testing, a constant cross-head speed of 0.25 cm/mi n. was used for all tests. This cross-head speed corresponds with an initial strain rate of 0.25 cm/cm/min (0.1 in./in./min.). Data acquisition was made possible by the use of a high temperature extensometer frame that attached to the specimen through the use of knife edges. This extensometer frame then extended down out of the furnace to an easily accessible ar ea where a digital extensometer connected to the computer could be attached. The Merlin soft ware package (by Instron) was used to record pre-test sample dimensions, control the servo-hydraulic system and record load, position, and extensometry data during the test. Data recorded during a test were subsequently analyzed to determine yield strength, ultimate tensile strength, percent elongation, and failure strength.

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60 Creep Testing Creep tests were conducted on Sa tec M-3 style creep frames. The tests conducted for this study were constant load creep te sts. Each frame has a weight pan for applying the load and a 16 to 1 lever arm ratio to load the specimen. As in the tensile tests, all threaded connections were lubricated with boron nitride high temperature lubricant. Specimen s were also affixed into high temperature extensometers as in tensile testing through the use of either knife edges on the gage sections or set screws on the shoulders. The bases of the creep extens ometers, though, included fixtures for Linear Variable Differential Transducers (LVDT) for measurement of displacement during the test. Again, clam-shell furnaces capable of 1200 C or 1300 C (depending on which frame was used) heated the specimens to the desired temperature for the duration of the creep test. Temperature control is maintained by the use of 3 Type K thermocouples attached to the gauge section of the specimens by 24 gauge 80Ni-20Cr wire as shown in Figure 3-4. Temperatures were maintained to within .7 C. Following heating to the test temperature, all specimens enter a soak period for 1 hour prior to loading and the commencement of the creep test. Hot loading was performed in multiple st eps and displacement measurements were taken allowing for calculation of the elastic modulus. The NuVision Mentor Creep Controller software (Satec) allowed for control of the creep frames, furnaces, and data collection. Creep data was collected at a rate of 5 times per minute for the first hour, then 1 time per minute for the remainder of the test. Additionally, times to 0.1%, 0.2%, 0.5%, 1%, 5%, and 10% cr eep were recorded automatica lly. Test specifications were written to conduct creep tests at four conditions: 704 C / 862 MPa, 704 C / 758 MPa, 760 C / 690 MPa, and 815 C / 621 MPa. These conditions were chosen because of their similarity to real-world servic e conditions and to generate the large primary creep strains as

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61 Figure 3-4. Three Type K thermo couples are attached to the ga uge length of creep specimens with 80Ni-20Cr wire. reported in the literature.16, 20, 47 The creep test matrix used for this study can be seen in Table 37. Full length and Interrupted creep tests were conducted to obtain a variety of useful data. Full length tests are those that were a llowed to run to failure. Interrupted tests are those that were stopped prior to failure in one of two conditions : after 0.5% secondary creep and tests running longer than 1200 hours with no si gn of imminent failure. Table 3-7. Creep test matrix with sample iden tification. Creep tests we re run to failure or terminated following 0.5% secondary creep. LT Age HT Age Temp ( C) 704 760 815 704 760 815 PWA 1480 0401-1 0401-20403-1 0404-1 0405-1 0406-1 0402-1 0403-2 0907-1 0907-2 0902-1 0902-2 PWA 1484 0407-1 0408-20409-1 0410-2 0410-1 0412-1 0913-1 0913-2 0412-2 0917-2 0917-1 0920-1 Tests run to failure PWA 1480+ 1203-1 1205-2 1209-1 1207-2 PWA 1480 0903-1 0903-2 0908-1 0908-2 PWA 1484 0914-1 0914-2 0918-1 0918-2 Tests stopped after 0.5% secondary creep PWA 1480+ 1204-1 1204-2 1210-1 1210-2

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62 CHAPTER 4 RESULTS: ALLOY MICROSTRUCTURES The three alloys used in this study repres ent first and second ge neration superalloys. Typically, nickel based superalloys may contain two or more phases including: (fcc Ni solid solution, matrix), (L12 Ni3Al), Carbides (MC, M6C, or M23C6), and/or Topologically Close Packed (Sigma, Mu, Laves, and/or P phases). Th e occurrence of several of these phases in the three alloys investigated here is discussed as a result of heat tr eatment and mechanical testing. All three alloys contain primary phase as a matrix with a high volume percent of precipitates following solution heat treatment and aging. The pr esence of a low amount of Carbon results in the formation of both script and blocky carbide phases for all three alloys. PWA 1480 and PWA 1480+ contain the greatest amount of carbide phases possibly as a result of a lower solubility for Carbon in the alloys. This effect is particul arly prominent in PWA 1480+ with the Re addition resulting in an increase in carbide amounts when compared to PWA 1480 without Re. Additionally, the alloy PWA 1480+ di splayed poor resistance to th e formation of TCP phases. The following discussion will focus on the properties of the phases themselves and their incidence as a result of casti ng and subsequent heat treatment as well as any changes during mechanical testing. Due to the large number of tensile graphs, creep graphs, micrographs, and other figures presented in the current and following results chapters, all figures have been placed at the end of the respective ch apters (Chapters 4-7) to aid reading speed and comprehension. Phase Descriptions The Phase The matrix of single crystal nickel base s uperalloys consists of the fcc solid solution phase. Due to the large volume fraction of used for high strength and creep resistance in these alloys, the phases only constitutes ar ound 25-35% of the alloy by volume. Topographically,

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63 the phase forms in the long narrow regions betw een the cuboidal (cubes with rounded edges) precipitates. These channels are typically 50-150 nm thick by 0.3-0.5 m in length and width. The length and width dimensions, o bviously, are controlled by the local size and distribution. The phase from all three alloys can be seen in Figures 4-1 through 4-3 as the light gray regions between the darker squares ( precipitates). Compositionally, there are significant differences between the and phases. For as long as superalloys have been in development there have been research investigations into the compositions of the two phases. In general te rms, the refractory elements added for solid solution strengthening partition to the matrix and the so called strengtheners partition to the phase. Specifically, Re has been shown to partition especially strong to the phase, however the addition of W can result in up to 20% of the added Re partitioning to the phase, effectively reducing the Re content of the matrix.35 The concentration of the phase can be predicted util izing the JMatPro thermodynamic prediction software. The calculated equilibrium concentration of the phase as a function of temperature is given in Figures 4-4 through 4-6. The nickel con centration for the three alloys varies from about 50% to 65% depending on the temperature. The enrichment of the gamma phase by Cr is most prominent in PWA 1480, followed by PWA 1480+ and PWA 1484, respectively. Within the temperature range of interest to the present investigation (700 C to 815 C), the refractory content of the phase (with the exception of W) of PWA 1480 and PWA 1480+ is low. The predic tion indicates that the phase would contain 2% or less of Re and Ta. The W content is around 5% for the PWA 1480 ba sed alloys. The Cr content of PWA 1480 and PWA 1480+ is much larger than that of PWA 14 84, while the Co content of PWA 1484 is nearly

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64 double that of PWA 1480 and PWA 1480+. The phase of PWA 1484 is also predicted to contain a greater amount of rhenium. The Phase Because nickel base superalloys are precip itation hardened alloys, the nature of the reinforcing phase is critical to the performance of the all oy. Very early during superalloy development it became clear that (Ni3Al) would be a good candidate for the strengthening phase.37, 48 Single phase has been shown to exhibit a larg e increase in strength as the temperature is increased to 800 C. Traditional (non-s uperalloy) alloys exhib it a slow decrease in strength with increasing temperature. When an alloy containing both phases is tested, the Critical Resolved Shear Stress (CRSS) is much la rger than the CRSS of e ither alloy individually and is stable with increasing temperature.37 The predicted compositions of the phase in all three alloys is presented in Figures 4-7 to 4-9. Examples of the morphologies present in the three a lloys tested here can be seen in Figures 4-10 to 4-13. PWA 1480 ex hibits a large variability in size and shape based on the location within the sample. For instance, dendrite core regions contain fine, cuboidal and the interdendritic regions near eut ectics contain large, irregular precipitates (Figures 4-10 and 411). Both PWA 1480 and PWA 1480+ contain retain ed eutectics due to the very small heat treatment window (9 C for PWA 1480 and 13 C for PWA 1480+, see Table 3-5). A retained eutectic is shown for PWA 1480+ in Figure 4-1 2. Additionally, segregation remains after the solution heat treatment and this leads to differences in characteristics by location. To counter these effects, as much as was practical, th e solution heat treatments for PWA 1480 and PWA 1480+ ended with final hold temperatures only 5 C and 3 C below the solidus for each alloy, respectively. PWA 1484 has a much larger heat treatment window (defined as the difference in

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65 C of the solvus and solidus temperatures). As a result, PWA 1484 exhibits very little retained eutectic following casting, Figure 4-13. Solution h eat treatment was able to eliminate eutectic regions entirely for this alloy. Primary The cuboidal primary precipitates for all th ree alloys are between 0.3 m and 0.5 m in edge length. In the case of the single cr ystal nickel superalloys PWA 1480 and PWA 1484, a high volume fraction of primary is precipitated partially from cooling following solution heat treatment and partially during the aging heat trea tments. As the temperature is increased, the equilibrium volume fraction of decreases as shown in Figure 414. Comparing the three alloys in question it becomes clear that PWA 1480 and PWA 1480+ have a significantly greater volume fraction of as predicted by JMatPro. A larger volume fraction can lead to differences in mechanical behavior.49 The nature of the interaction between the phase and Re has been the subject of several investigations spanning the past 30 years. It is generally accepte d that Re is rejected from the phase during precipitation a nd so partitions to the matrix. As a result, Re additions are often implicated as the probable cause of large negativ e lattice misfit values in the vicinity of the / interface.8 The local enrichment of the matrix side of the / interface occurs over a relatively short distance due to the low diffu sivity of rhenium. Two theories regarding the nature of the enriched layer have been argued for the last 20 year s. One states that the rejected solute atoms form a hardened shell around the precipitates. A contrasting study published in 1988 suggests that the rejected Re forms ha rd clusters roughly 10 in size.10, 11 Similarly, clustering has been shown to occur with Cr forming enriched regions less than 4 nm in size. These Cr clusters are thought to be due to the creati on of the ordered structure Ni3Cr (DO22).50

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66 The phase is of particular inte rest for the present investigation due to the correlation of large primary creep strains with shearing mechanisms. The nature of the phase and the / interface may be related to the primary behavior described herein. Additionally, while most investigations that purs ue an understanding of the phase focus on the primary precipitates, the secondary precipitates are also of significant value due to their location in the matrix channels. As discussed in Chapter 6, TEM analysis revealed that disl ocations in PWA 1480 are limited to the matrix where they are likely to fre quently interact with the secondary PWA 1484, however, is not so constrained and so an un derstanding of both the primary and secondary precipitates will aid in understanding the active deformation mechanisms. Secondary Metallography. Secondary precipitates, Figures 4-15 to 4-17, are much smaller (less than 50 nm in diameter) and spherical. Th ese secondary precipitates lie in the narrow channels and in the larger primary free regions near eutectics and mi nor phases. Figures 4-15 and 4-16 are from PWA 1480 following an in terrupted creep test (LTA, 704 C/862 MPa). Figure 4-17 shows the secondary in PWA 1480+ also following an interrupted creep test (LTA, 704 C/862 MPa). Not much is really understood about th ese ultra-fine precipita tes. One key study, (Kakehi, 1999), demonstrated the ability to eliminate secondary precipitates by utilizing a slow furnace cool following the final age temperature.20 It is generally agreed that secondary form during the rapid cooling commonly employed in standard heat treatment practices.20, 47, 51 The small size of the secondary precipitates c oupled with the complex environment in which they occur makes it difficult study the structure and composition of these precipitates. Viewing the secondary precipitates via SEM techniques was unsuccessful with PWA 1484; however.

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67 Local electrode atom probe (LEAP). The LEAP system is a new characterization technique developed for high re solution compositional analysis coupled with high spatial resolution. This new system was employed in the present investigation to study the rejection of Re from during aging and solution heat treatments and it allowed for examination of secondary in PWA 1484 where metallographic tech niques failed due to the much better solution and homogenization that wa s achieved in the alloy. In orde r to interpret LEAP results, a brief discussion of its use is necessary (see Chap ter 7 for a complete discussion). Because the LEAP utilizes compositional data, two and three dimensional map representations can be created by assigning colors to individua l atoms. To make the images useful, maps are created by applying thresholds to limit the nu mber of atoms appearing in the image. For example, a limit of 18% alumnum would eliminate almost all Al dots in the phase while allowing Al dots to appear in the phase. In this way, the limits of the and phases can be mapped.52 To examine the phase, these limits were applied to the Al dots such that the phase is transparent and the phase appears in stark contrast. An example of this type of image for PWA 1484 can be seen in Figure 4-18. The micrographs, coupled with the LEAP da ta, revealed the presence of secondary in both age conditions of all three alloys. This resu lt is consistent with those reported by Kakehi due to the use of rapid cooling from all he at treatments in the current investigation.20 Despite the similarities in secondary morphology between alloys, thou gh, significant differences in primary creep occurred during mechanical testing. This behavior will be discussed in further detail in Chapter 8, however, it is clear that pr imary creep is controlled by many mechanisms and not primarily by the presence of secondary precipitates.

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68 The / Eutectic All three alloys formed / eutectics during solidification. Examples can be seen in Figures 4-19 through 4-20. These two-phase f eatures form during solidification due to the complex nature of the alloys. During dendritic growth of superalloys, the first material to solidify is enriched in the high melting point refr actory elements. These fi rst-to-solidify regions eventually comprise the dendrite cores and have the highest solidus temperature compared to other regions in a sample (excluding, of c ourse, minor phases such as high melting point carbides/TCP phases).44, 53 As solidification pr ogresses the liquid becomes depleted in these refractory elements and, consequently, enrich ed in lower melting point elements. The solidification temperature of the remaining liquid continues to decline until the eutectic temperature is reached. At this lowest temperat ure, a two-phase eutectic region is formed with the remaining liquid. This range of solidific ation temperatures ahead of an advancing solidification front leads to the creation of the so-called mushy zone in directional solidification. Because the eutectic regions so lidified at the lowest temperature, these regions limit the thermal capability of the alloy. To improve th e thermal capability and reduce the effect of eutectics, solution heat treatments have been developed to provide enough thermal energy for significant diffusion to take place to allow the enriched regions of the dendrite cores and eutectics to approach the origin al alloy composition. In order fo r this to occur, solution heat treatments must exceed the solvus but remain below the solidus temperature. One danger, however, is the risk of incipien t melting. Increasing the temperatur e of the material too quickly to a temperature above the solvus may cause inadvertent melting because the solidus may be depressed to the same level as the solvus itself. For this reas on, solution heat treatments of

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69 these alloys often takes the form of a multi-step heat treatment that approaches the solvus and the solidus slowly (the HT3 heat treatments, for example, follow this idea, Table 3-4).2, 3 The / eutectics that formed during solidifi cation were retained in PWA 1480 and PWA 1480+ following solution heat treatment. The already mentioned narrow solution heat treatment windows for these two alloys prev ented the eutectics from being eliminated entirely. Following the HT3 solution heat treatment, however, they we re significantly reduced due to the longer time and higher temperature than the HT1 heat treatm ents. Figures 4-19 and 4-20 contain examples of eutectics in PWA 1480 and PWA 1480+. As seen in Figure 4-20, eutec tic regions often occur in close proximity to carbides. Figures 4-21 a nd 4-22 are close-up views of eutectics in which both eutectic phases ( and ) can be seen. Figure 4-21 shows an as-cast example and Figure 422 shows a eutectic following an interrupted creep test. PWA 1484, however, exhibited no appreciable eutectics following h eat treatment due to the much larger solution heat treatment window (46 C in the HT2 condition, Table 3-5). The Carbides The presence of a small amount of carbon leads to the formation of carbide phases in most single crystal nickel superalloys The three alloys discussed herein contain between 0.02 and 0.04 wt.% C. This small addition is enough to pr oduce a low volume fraction of carbide phases. The carbide phases most commonly encountered in superalloys are the primary MC type and the transition, or secondary, M6C and M23C6. Many other carbides are possible as well in addition to the formation of boride, nitride, and carbo-nitr ide phases (in the presen ce of carbon, nitrogen, and/or boron).38, 50 There has been a recent increase in in terest in the many perceived benefits and/or detriments of adding carbon to single cr ystal superalloys; however, they are beyond the scope of this investigation. Originally adde d as a grain boundary strengthener, carbon was

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70 removed from early single crystal superalloys Recently, though, carbon has been added back into single crystal superalloys in low amounts to lower casting defects, to capture tramp elements, and to improve defect tolerance. The carbides found in PWA 1480 and PWA 1484 are found in interdendritic regions and take on script like morphologies, Figures 4-23 to 4-25. Some blockier shapes are possible as well depending on the local conditions. Larg er amounts of carbon ar e required to produce dendritic carbides and so the carb ides appear to be localized a nd not networked as shown in the longitudinal section in Figure 4-24. Both PW A 1480 and PWA 1480+ exhibi t local networks of script like carbides throughout. PWA 1484, however exhibits isolated clusters of small blocky carbides as shown in Figure 4-26. This differenc e is likely primarily due to the lower Carbon concentration of PWA 1484 (r oughly half of the Carbon conc entration of PWA 1480 and PWA 1480+). Additionally, some carbides in the heat treated PWA 1480 and PWA 1480+ samples appear to be in the processe s of dissolving as solutioning and homogenizing processes are occurring, Figure 4-27. The carbi de phases, while subject of many investigations, are not considered to be cause for concern during primary creep. Because primary creep occurs well before failure, the issue of prem ature crack initiation is not cri tical and so these phases are not likely to play much of a role in primary creep. The Topologically Close Packed (TCP) Phases Topologically close packed phases are de leterious phases that form during high temperature exposure of many nickel base supera lloys. Topologically close packed structures differ from geometrically close pa cked (GCP) structures (like fcc and L12 ) by having planes of close packed atoms separated by relatively large planar spacings, while GCP phases are close packed in all directions. The large interplane r distances are caused by large diameter solute

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71 atoms such as the refractory elements Re, W, and Mo. Several related phases fall under the category of Topologically Close Packed including P, and Laves. Due to their brittle nature, their often needle a nd plate-like morphology, and their ability to rob the surrounding material of solid solution strengtheners, TCP phase s have been the focus of great concern within the superalloy community. The drive to eliminat e TCP phases from commercial alloys has lead to the development of computer modeli ng and prediction methodologies like PHACOMP (PHAse COMPutation) utilizing the number of unpaired electrons, Nv, for each element in an alloy. There have been several attempts to modi fy this model with varying levels of success (including atomic size factors, Md); however, the basic theory remains the same. By assigning specific Nv numbers to each alloy and building assumptions about the typical concentration of phases in an alloy, an average Nv value can be determined. It has been shown that for many alloys, an Nv value that exceeds a critical value of 2.45-2.5 will lead to the formation of phase. Using the traditional Nv method of PHACOMP, PWA 1480 has an average Nv value of 2.54 while PWA 1484 has an average Nv value of 2.50 (reported in DurrandCharre, 1997).50 These values would i ndicate that PWA 1480 is likely to produce TCP phases during long-term, high temperat ure exposure, while PWA 1484 is potentially safe from TCP formation. Adding 3 wt.% Re, however si gnificantly increase s the already high Nv of PWA 1480, leading to early and rapid precipitation of TCP phases duri ng high temperature exposure. This prediction has been proven during creep test ing as shown in Figures 4-28 and 4-29. The TCP phases shown were produced during primary creep of PWA 1480+ and were prevalent in specimens tested at 704 C and 815 C and with either age heat treatment. While TCP phases are linked to premature failure of superalloys, it s hould be noted that PW A 1480+ exhibited the longest creep lifetime and the lowe st creep rate and primary creep strain as discussed in Chapter

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72 6. These results indicate that TCP formation is not necessarily guaranteed to cause a detriment to mechanical strength properties, however, fracture toughness and ductility may be reduced. Changes Following Primary Creep The most significant changes to the microstruc tures of the three all oys was exhibited by PWA 1484. Interrupted creep test s of PWA 1484 were stopped follo wing as much as 28% creep in 15 hours or less of testing. As will be shown in Chapter 6, PWA 1484 deformed by massive stacking fault shear of the phase. When these specimens we re observed after testing, all the PWA 1484 specimens exhibited elliptical cross-sections rather than the original circular crosssections. This behavior wa s consistent with work published on both PWA 1484 and another second generation superalloy, CMSX-4.16, 17, 47 Metallographically, these specimens also exhibited elongation of the phase in the [110] dire ction (observed with a etch), Figures 4-30 and 4-31. Following the use of the electrolytic etch, the primary precipitates appeared to be cut along planes consistent with (111) planes, Fi gures 4-32 and 4-33. These figures indicate that the shear processes active in PWA 1484 must be severe due to the significant changes to the primary precipitates following in terrupted creep testing.

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73 Figure 4-1. / microstructure of PWA 1480 HT3 ( etch). The phase is the light gray phase between the cuboidal phase (dark grey). Figure 4-2. / microstructure of PWA 1480+ HT3 ( etch). The phase is the light gray phase between the cuboidal phase (dark grey).

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74 Figure 4-3. / microstructure of PWA 1484 HT3 ( etch). The phase is the light gray phase between the cuboidal phase (dark grey). Gamma Phase0 5 10 15 20 25 30 35 40 45 50 55 60 65 70 60070080090010001100120013001400 Temperature (C)Composition (wt.%) Ni Al Co Cr Re Ta Ti W C Figure 4-4. Composition of the phase of PWA 1480 as a function of temperature as predicted by JMatPro.

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75 Gamma Phase0 5 10 15 20 25 30 35 40 45 50 55 60 65 70 60070080090010001100120013001400 Temperature (C)Composition (wt.%) Ni Al Co Cr Re Ta Ti W C Figure 4-5. Composition of the phase of PWA 1480+ as a functi on of temperature as predicted by JMatPro. Gamma Phase0 5 10 15 20 25 30 35 40 45 50 55 60 65 70 60070080090010001100120013001400 Temperature (C)Composition (wt.%) Ni Al Co Cr Hf Re Ta W C Figure 4-6. Composition of the phase of PWA 1484 as a function of temperature as predicted by JMatPro.

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76 Gamma Prime Phase0 5 10 15 20 25 30 35 40 45 50 55 60 65 70 75 60070080090010001100120013001400 Temperature (C)Composition (wt.%) Ni Al Co Cr Re Ta Ti W Figure 4-7. Composition of the phase of PWA 1480 as a function of temperature as predicted by JMatPro. Gamma Prime Phase0 5 10 15 20 25 30 35 40 45 50 55 60 65 70 75 60070080090010001100120013001400 Temperature (C)Composition (wt.%) Ni Al Co Cr Re Ta Ti W Figure 4-8. Composition of the phase of PWA 1480+ as a function of temperature as predicted by JMatPro.

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77 Gamma Prime Phase0 5 10 15 20 25 30 35 40 45 50 55 60 65 70 75 60070080090010001100120013001400 Temperature (C)Composition (wt.%) Ni Al Co Cr Hf Re Ta W Figure 4-9. Composition of the phase of PWA 1484 as a function of temperature as predicted by JMatPro. Figure 4-10. The phase in PWA 1480 near a retained eute ctic (lower right) as revealed by the electrolytic etch.

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78 Figure 4-11. Irregular phase in PWA 1480 near a partially solutioned eutectic region ( etch). Figure 4-12. / eutectic region in PWA 1480+ during the early stages of solutioning ( etch). Fine primary precipitates can be seen in the lowe r right quadrant of the micrograph. The large, hard phase in the middle is a carbide (most likely MC type).

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79 Figure 4-13. The structure of as -cast PWA 1484 ( etch). 60 62 64 66 68 70 72 74 76 78 80 680700720740760780800820840Temperature (C)Volume Fraction (%) PWA 1480 PWA 1480+ PWA 1484 Figure 4-14. The volume fraction vs. temperat ure for all three alloys.

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80 Figure 4-15. Secondary in the matrix of PWA 1480 following an interrupted creep test ( etch). Figure 4-16. Secondary in the matrix of PWA 1480 following an interrupted creep test ( etch).

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81 Figure 4-17. Secondary in the matrix of PWA 1480+ follo wing an interrupted creep test ( etch). A TCP precipitate can be seen positioned diagonally from top to bottom-right.

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82 Figure 4-18. Three dimensional LEAP compositiona l map (18wt.% Al iso-surface) The green surfaces represent areas of Aluminum con centrations averaged at lease 18wt.%. These regions are predominately due to the highly ordered of the phase. The phase is typically very low in Al concentra tion. The fine precipitates can be seen in the transparent matrix channels (near z = 100 nm). Displacement scales are in units of nanometers (nm).

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83 Figure 4-19. / eutectics in as-cast PWA 1480 ( etch). Figure 4-20. / eutectics in the vicinity of primar y carbides in PWA 1480+ (HT1A solution heat treatment, etch). Solutioning of the eutect ics is already underway with the phase infiltrating the mostly eutectic (center right).

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84 Figure 4-21. Close-up of a eutectic in as-cast PWA 1480+ ( etch). Small precipitates can be seen throughout the eutectic region. Figure 4-22. Close-up of a reta ined eutectic in PWA 1480+ follo wing an interrupted creep test (HTA, 704 C/862 MPa, etch).

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85 Figure 4-23. Carbide phase in PWA 1480+ (HT1A, etch). Figure 4-24. Carbide phase in PWA 1480+ (as-cast, longitudinal section, etch). Carbides do not appear to be dendritic in nature, but do wrap around de ndrite arms (center-right).

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86 Figure 4-25. Local carbide network in PWA 1480 following an interrupted creep test (HT, 815 C/621 MPa, etch). Figure 4-26. Carbide phase in PWA 1484 ( etch). Due to the lower wt.% Carbon, PWA 1484 exhibits isolated clusters of small blocky carbides.

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87 Figure 4-27. A possible carbide that dissolved during solution heat treatment leaving behind (PWA 1480 LT, 704 C/862 MPa, etch). Figure 4-28. TCP phase formation in PW A 1480+ during interrupted creep testing ( etch).

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88 Figure 4-29. TCP phase formation in PW A 1480+ during interrupted creep testing ( etch). Figure 4-30. phase elongation in the [110] direction in PWA 1484 following interrupted creep testing (HT age, 704 C/862 MPa, etch).

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89 Figure 4-31. phase elongation in the [110] direction in PWA 1484 following interrupted creep testing (HT age, 704 C/862 MPa, etch). Figure 4-32. phase shear along (111) planes in PWA 1484 following interrupted creep testing (HT age, 704 C/862 MPa, etch).

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90 Figure 4-33. phase shear along (111) planes in PWA 1484 following interrupted creep testing (LT age, 704 C/862 MPa, etch)

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91 CHAPTER 5 RESULTS: TENSILE BEHAVIOR Tensile tests of PWA 1480, PWA 1480+, and PWA 1484 were conducted to evaluate the yield and tensile strengths of th e three alloys. When designing cr eep experiments, it is important to understand the relative differences in yield strength between the three alloys. For instance, alloys with different yi eld strengths subjected to the same creep load may experience different creep behaviors as a result. It is possible that an alloy that is st ressed to a greater fraction of its yield stress may experience a shorter creep lif e because of an increase in deformation (alternatively, a decrease in resistance against cr eep. Additionally, tensile testing is useful to evaluate work hardening and ductility. Tensile testing was conducte d at two temperatures, 700 C and 815 C, set at the lowest and highest temperatures used for cr eep testing. The following discus sion will focus on the effect of the two different age heat treatments on th e tensile strengths of PWA 1480, PWA 1484, and PWA 1480+. As mentioned earlier, the age heat treatment temperatures were 704 C and 871 C and all testing was conducted at or above th e LT age temperature, but below the HT age temperature. Additionally, the effect of rhenium on the tens ile properties of PWA 1480 was investigated by the creation of a second gene ration version of PWA 1480 called PWA 1480+. Because microstructural evolution is not a major concern during the short duration of a tensile test, the relationship between test te mperature and age temperature is minor. For instance, if the time required for tensile test wa s sufficient to change the heat treated condition, then testing at 815 C might be expected to reduce or elim inate entirely any benefit of the low temperature (LT) age heat treatment. In this case (815 C), the thermal exposure during tensile testing following the HT age will act as a second ary age heat treatment much like the multi-step age heat treatments already employed in the pr ocessing of some second and third generation

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92 superalloys. Additionally, the HT age would be expected to increase co arsening (size) of the phase and, subsequently, reduce coherency of the precipitates. Rhenium additions have been shown to significantly increase so lid solution strengthening of the phase. Because Re preferentially segregates to the phase, lattice misfit is significantly increased as well. The tensile behavior of PWA 1480 is very sim ilar for both age heat treatments. At 700 C, both samples exhibit a large yield stress followed by slight work-hardening until failure, as shown in Figure 5-1. The HT aged sample has a yield strength slightly lower than the LT age and the ductility is about half of the ductility of the LT aged samples. At 815 C, both age heat treatments have a yield point that is lower than the yield strength of the samples tested at 700 C. Again, the HT aged samples have slightly lowe r yield strength, failure strength, and ductility. The Re modified PWA 1480+ performed sim ilarly to unmodified PWA 1480 with a few notable differences, Figure 5-2. The general shape of the stress-strain curves remained the same, potentially indicating no signifi cant change in deformation mech anisms. The two most notable changes are the resulting increase in yield stre ngth and corresponding reduction in ductility. The LT aged PWA 1480+ sample at 700 C has a much lower yield strength than the PWA 1480 LT specimen as well as significantly less ductility (1 1% for PWA 1480 vs. 2% for PWA 1480+). It should be noted, however, that this result may not be valid due to its sign ificant departure from the trend established by the other test resu lts for PWA 1480+. The HT aged PWA 1480+ specimen at 700 C showed a great improvement in yield strength (1333 MPa for PWA 1480 vs. 1376 MPa for PWA 1480+). The ductilities for the HT age specimens of both alloys were similar with a difference of less than 2% elongation. Both alloys exhibit yield points at 815 C and display very similar strain hardening behavior after yielding. For th ree of the four conditions, PWA 1480+ has a greater yield stress

PAGE 93

93 than conventional PWA 1480, which is likely due to the so lid solution strengthening effect of the Re addition. For both test temperatures, the HT age appears to decr ease ductility in PWA 1480 (11.6% to 6.3% and 18.3% to 15.2) while simultan eously reducing the yield stress (1357 MPa to 1333 MPa and 1281 MPa to 1212 MPa, LT and HT respectively). PWA 1480+, though, had increased ductility and yield strength in the HT age when compared to the LT age condition. Additionally, the usually stronge r PWA 1480+ does not show as much ductility as PWA 1480 in the LT condition but does in the HT condition. Tensile strength and elongation measurements from all three alloys are presented in Table 5-1. The tensile behavior of PWA 1484 was somewhat different from PWA 1480 (and PWA 1480+), Figure 5-3. PWA 1484 was characterized by a significantly lower yield point, greater ductility, and greater amount of pl astic work hardening. PWA 1484 al so has greater strength at the higher temperature test condition while PWA 1480 and PW A 1480+ show a decrease in strength at 815 C. In general, nickel base superalloys have excellent high temperature strength to near 800 C. Continuing to increase the temperature will result in a decrease in strength in nickel base superalloys because of the decreasing strength of the phase at these temperatures.37, 48 The continuing increase in strength of PWA 1484 at 815 C is noteworthy as the total refractory content (Mo+W+Re+Ta) fo r PWA 1484 is similar to that of PWA 1480+ (20wt.% vs. 19wt.%, respectivel y). The total refractory cont ent of PWA 1480, however, is less at 16wt.% with the absence of Re. Despite the similarity in solid solution strengthener content of PWA 1484 and PWA 1480+, PWA 1484 shows an increase in strength at 815 C relative to 700 C while PWA 1480+ does not. At 700 C, the strength of PWA 1484 continues to increase until failure. At 815 C, the strength increases to the ulti mate tensile strength, UTS, following 24% plastic deformation. All of the PWA 1484 tests produced between 14 and 23% elongation

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94 compared with 5-11% for PWA 1480 and 210% for PWA 1480+, Table 5-1. For PWA 1480 and PWA 1484, the LT age condition is slightly stronger than the HT age condition. For PWA 1480+, however, the HT age produces greater yield st rength values at both te st temperatures. Direct comparison between the three alloys also reveals the rela tive differences in performance. Figures 5-4 and 5-5 show direct co mparisons of the three alloys tested in the LT age condition at 700 C and 815 C respectively. Figures 5-6 and 5-7 are similar except the HT age is presented (rather than th e LT age shown in Figures 5-4 and 5-5). Specimens with the LT age at 700 C exhibited a small amount of work harden ing after yielding as shown in Figure 5-4. Of the three alloys, PWA 1484 displayed the grea test potential for work hardening. PWA 1480 had the largest yield strength valu e and substantial ductility, but little work hardening. PWA 1480+ produced little ductility (less than 2%) bu t experienced slight hardening after yielding. PWA 1484 had the lowest yield point at this co mbination of age heat treatment and test temperature. At 815 C, the LT aged samples all experienced a sharp yield point. In the case of PWA 1480 and PWA 1480+, the yield point was followed by an immediate reduction in tensile stress. PWA 1484, however, maintained the yield stress level briefly before the onset of plastic hardening. The engineering stress is reduced 100 to 150 MPa during this period of the test possibly as a result of the onset of necking. No further strengthening take s place before failure occurs. Contrasting this behavior, PWA 1484 displa ys a significant increase in strength during the first few percent plastic elongation. The applied engineering stre ss slowly decreases following the work hardening (the transition occu rs around 4-5% elongation). The decrease in stress may be a result of slight necking, though necking was slight While the yield strength of PWA 1480 is 373 MPa greater than that of PWA 1484, the reduction in measured flow stress of

PAGE 95

95 PWA 1480 and the increase in flow stress of PWA 1484 lead to a larger failure stress for PWA 1484 by 78 MPa. Compared to PWA 1480+, the UTS of PWA 1484 is 152 MPa below that of PWA 1480+; however, the true failure stress of PW A 1484 (corrected by post test measurement) is 57 MPa greater than the failure stress of PWA 1480+. PWA 1484 has the highest failure strength despite a 445 MPa and 374 MPa disadvan tage in yield strengt h to PWA 1480+ and PWA 1480, respectively. The HT age heat treated specimens share severa l similarities to the LT age specimens. At 700 C, PWA 1480 and PWA 1480+ exhibited the same high yield strength with a nearly flat plastic hardening region, Figure 5-6. PWA 1480 experienced a significant reduction in ductility with the HT age. PWA 1480+ when tested with the LT age at 700 C had the highest yield strength of the group. The ductility for PWA 1480+ is increased with the HT age from 2.8% to 5.1%. The behavior of PWA 1484 in the HT age condition closely matches its behavior in the LT condition. The yield points and amount of wo rk hardening are very similar as are the ductilities of both alloys. PWA 1480, however, expe rienced a decrease in yield strength, UTS, and ductility in the HT age condition. Most notable is the decrease in ductility of PWA 1480 from 11.8% to 5.5%. The HT age when tested at 815 C brought about similar tensile performances to the LT age for all three alloys. Again, PWA 1480+ had th e highest yield strength followed by PWA 1480 and PWA 1484. Both PWA 1480 and PWA 1480+ e xhibited a sharp yield point followed by a drop in strength of nearly 100 MPa. The measur ed strength continued to decline until failure. Conversely, PWA 1484 displayed a sharp rise in strength following yielding to a maximum (1141 MPa) at 3.96% elongation. When co mpared with the results from the 700 C testing, there was little difference in properties for PWA 1480+. PWA 1480 had lower yield strength and less

PAGE 96

96 ductility in the HT age condition and PWA 1484 ha d similar yield strengths in both conditions but less work hardening in the HT age condition. The total ductility of PWA 1484 was also slightly lower for the HT age condition. The elastic modulus for all th ree alloys was calculated during the step loading procedure immediately prior to the beginning of a creep test. These values are used here as the creep loading system is more elastically rigid and the results fit a linear model with much lower error than those generated on the servo-hydraulic tensile test system. The elastic modulus values were obtained at 704 C and 815 C for all three alloys and at 760 C for PWA 1480 and PWA 1484. These values are reported in Table 5-2. Additionally, the three alloys can be differentiated by their pl astic deformation behavior. PWA 1480 displays relatively smooth plastic de formation until failure. The LT condition transitions from elastic to plasti c behavior smoothly as shown in Figure 5-8. The HT condition behaves similarly; however, the elastic-plastic transition is marked by a brief spike in stress shown in Figure 5-9. PWA 1484 produced a steady in crease in flow stress until an instability point is reached. After this point, the engineering stress is initially reduced and wavy until failure. The UTS may occur at failu re or at the point of instability as in the cases of the LT age and the HT age, respectively: Figures 5-10 and 5-11. The point of instability may be caused by necking or a similar behavior; however, only slight necking coul d be observed on all specimens tested herein. PWA 1480+ exhibited a mix of beha viors from the LT age to the HT age. In the LT aged condition, PWA 1480+ displayed a smooth elastic-plastic transi tion that was followed by a rapid increase in flow stress to a potential point of instability at which the slope of the plastic deformation region decrea sed sharply, Figure 5-12. Overall, the LT condition exhibited a smooth plastic response from yielding until fa ilure. The HT age, however, exhibited wavy

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97 behavior with several spikes in flow stress fr om yielding until failure, Figure 5-13. The spikes produced in the HT condition of PWA 1480 and PWA 1480+ are possibly a result of dynamic strain aging processes such as the formation of solute atmospheres along dislocation cores during the age heat treatment. Comparing the behaviors of the three alloys between the two te st temperatures, it can be seen that two alloys, PWA 1480 and PWA 1480+, e xhibit a significant cha nge in behavior over the temperature range in questi on while PWA 1484 does not. At 700 C, PWA 1480 and PWA 1480+ transition from elastic to plastic deform ation with no drop in stress caused by a yield point, Figures 5-14 and 5-16. At 815 C, however, both alloys exhi bit a sharp upper yield point followed by a constantly decreasing flow stress, Figures 5-15 and 5-17. Th is behavior continues until failure for both alloys. The true failure stre ss of PWA 1480 is only slightly greater than the yield stress and UTS values (both recorded at yielding). The tr ue failure stress of PWA 1480+, however, shows a decrease in strength of 27 MPa following yielding in the LT condition and 2 MPa following yielding in the HT age condi tion, Table 5-1. PWA 1484, however, does not show a significant change in behavior from 700 C to 815 C. For both age conditions of PWA 1484, the yield point is nearly id entical at both temperatures w ith a difference of 17 MPa at 700 C and 10 MPa at 815 C, Figures 5-18 and 5-19. The primary difference recorded for PWA 1484 is the slope of the plastic deformation region immediately following yielding. For both age conditions, the 815 C sample exhibits a more rapid increase in flow stress due to increased work hardening and a larger value for the ultimate tensile strength. The improvement in tensile properties of PWA 1484 is not unexpected as it is a second generation alloy designed for higher temperature capability than PWA 1480.54

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98 Finally, changes with temperat ure in common tensile propertie s like yield strength, UTS, and failure strength can also be recognized for each of the three alloys. The yield strength of the three alloys as a function of temperature is given in Figure 5-20. PWA 1480 and PWA 1484 exhibit reductions in yield strength as a result of the increase in temperature from 700 C to 815 C. PWA 1480 showed the largest drop in yiel d strength (76 MPa for the LT condition and 121 MPa for the HT condition), while PWA 1484 exhibited a reduction of 41 MPa in the LT condition and 34 MPa in the HT condition. PWA 1480+, however, exhibited an increase in yield strength with temperature for both age condi tions. The LT condition of PWA 1480+ produced the greatest increase in yield strength with temp erature (118 Mpa) while the HT age produced a 3 MPa rise in yield strength at 815 C. Additionally, comparing the m easured yield strengths to the applied initial creep loads provide s a useful description of the pe rcentage of the yield strength required to support the applied stress, Table 5-. The ultimate tensile strength of the three alloys as a function of test temperature is given in Figure 5-21. PWA 1480 with both age conditions and PWA 1480+ HT exhibit a substantial reduction in UTS of between 140 and 160 MPa at 815 C when compared with the values produced at 700 C. As the temperature is increased furthe r, it would be expect ed that the tensile strength of these alloys would continue to decline due to th e decreasing strength of the matrix and the precipitates.37, 48 An increase in the measured UTS was exhibited by PWA 1484 with both age conditions and by PWA 1480+ LT. Again, the increase in strength of PWA 1484 at the higher temperature can be attrib uted to its higher temperature performance capability produced as a result of a greater refractory content than either of the PWA 1480 allo ys. The increase in strength for the PWA 1480+ LT sample may be an ar tifact as a result of the significant departure by the sample as indicated in Figure 5-2.

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99 When true failure stress is graphed for each a lloy as a function of temperature, only two specimens exhibited an increase in strength wi th increasing temperature, Figure 5-22. PWA 1480+ LT and PWA 1484 LT both produced a substantial increase in fa ilure stress at 815 C (90 MPa for PWA 1480+ LT and 83 MPa for PWA 1484 LT). PWA 1480+ HT, PWA 1484 HT, and both conditions of PWA 1480 all exhibited a reducti on in failure stress at the higher temperature test with PWA 1480 LT showing the greatest reduction in strength.

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100 Figure 5-1. Tensile results for PWA 1480 at both age heat treatm ents and test temperatures. Figure 5-2. Tensile results for PWA 1480+ at both age heat treat ments and test temperatures.

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101 Table 5-1. Tensile results at 700 C and 815 C for all three alloys and both age heat treatments. Temperature Age y (MPa) UTS (MPa) y/p/f f (MPa) Elongation at Failure (%) f (%) RIA (%) LT 1357.44 1429.33 (f) 1604.29 11.76 11.58 10.93 700C HT 1333.45 1371.15 (p) 1459.98 5.49 6.28 6.09 LT 1281.39 1287.39 (y) 1321.31 12.36 18.27 16.70 PWA 1480 815C HT 1212.24 1213.76 (y) 1220.33 7.68 15.18 14.09 LT 1149.77 1215.07 (f) 1250.34 2.58 2.85 2.81 700C HT 1375.57 1516.20 (f) 1606.16 5.09 6.86 6.63 LT 1267.39 1352.09 (y) 1340.15 7.12 12.17 11.46 PWA 1480+ 815C HT 1378.33 1385.63 (y) 1384.00 8.99 17.04 15.67 LT 948.72 1112.69 (f) 1314.27 13.46 17.26 15.86 700C HT 931.76 1116.55 (p) 1328.08 16.58 18.30 16.72 LT 907.90 1199.62 (p) 1397.51 18.22 23.21 20.71 PWA 1484 815C HT 897.90 1141.44 (p) 1203.85 19.26 13.84 12.92 *Note: UTS values are designated with a y if the UTS occurred at a yield point, a p if the UTS occurred during plastic deformation before failure, and an f if the UTS occurred at failure. Figure 5-3. Tensile results for PWA 1484 at both age heat treatm ents and test temperatures.

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102 Figure 5-4. Comparison of te nsile results for all three alloys with the LT age (704 C/24hr.) tested at 700 C with an air environment. Figure 5-5. Comparison of te nsile results for all three alloys with the LT age (704 C/24hr.) tested at 815 C with an air environment.

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103 Figure 5-6. Comparison of tens ile results for all three a lloys with the HT age (871 C/32hr.) tested at 700 C with an air environment. Figure 5-7. Comparison of tens ile results for all three a lloys with the HT age (871 C/32hr.) tested at 815 C with an air environment.

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104 Table 5-2. Elastic modulus calculations from creep loads. Elastic Modulus (GPa) 700 C 815 C LT HT LT HT PWA 1480 81.48 95.75 87.44 64.30 PWA 1480+ 95.01 86.63 100.78 79.96 PWA 1484 85.15 86.39 106.44 87.41 Figure 5-8. Plastic deformati on behavior of PWA 1480 LT at 700 C.

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105 Figure 5-9. Plastic deformation behavior of PWA 1480 HT at 700 C. Figure 5-10. Plastic deformation behavior of PWA 1484 LT at 700 C.

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106 Figure 5-11. Plastic deformati on behavior of PWA 1484 HT at 700 C. Figure 5-12. Plastic deformati on behavior of PWA 1480+ LT at 700 C.

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107 Figure 5-13. Plastic deformation behavior of PWA 1480+ HT at 700 C. Figure 5-14. Tensile behavior of PWA 1480 LT at both temperatures.

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108 Figure 5-15. Tensile behavior of PWA 1480 HT at both temperatures. Figure 5-16. Tensile behavior of PWA 1480+ LT at both temperatures.

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109 Figure 5-17. Tensile behavior of PWA 1480+ HT at both temperatures. Figure 5-18. Tensile behavior of PWA 1484 LT at both temperatures.

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110 Figure 5-19. Tensile behavior of PWA 1484 HT at both temperatures. Figure 5-20. Yield strength as a functi on of temperature for all three alloys.

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111 Table 5-3. Creep loads vs. yiel d strength for all three alloys. Age y (MPa) Creep Load (MPa) % y LT 1357.44 862 63.5 PWA 1480 HT 1333.45 862 64.6 LT 1149.77 862 75.0 PWA 1480+ HT 1375.57 862 62.7 LT 948.72 862 90.9 700C PWA 1484 HT 931.76 862 92.5 LT 1281.39 621 48.5 PWA 1480 HT 1212.24 621 51.2 LT 1267.39 621 49.0 PWA 1480+ HT 1378.33 621 45.1 LT 907.90 621 68.4 815C PWA 1484 HT 897.90 621 69.2 Figure 5-21. Ultimate Tensile Strength as a function of temperature for all three alloys.

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112 Figure 5-22. True Failure St ress as a function of temperature for all three alloys.

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113 CHAPTER 6 RESULTS: CREEP BEHAVIOR Creep testing was initially performed at three combinations of temperature and initial stress as follows: 704 C/758 MPa, 760 C/690 MPa, and 815 C/621 MPa. After running several tests at 704 C/758 MPa the initial applied stress was increased to 862 MPa due to the excessive failure lives of PWA 1484 specime ns (greater than 1700 hours). As discussed earlier in Chapter 3, creep testing consisted of two phases of testing. First, full-length testing was performed to establish lifetime and performance expectations. The original goal of this investigation was to study the effect of secondary precipitates and (related) age heat treatments. Following the first complete round of testing, the unique primary creep behavior became obvious and the investigation took up a new focus (on primary cr eep behavior). The second batch of single crystal bars of PWA 1480 and PWA 1484 was acquired to investigate primary creep further. Additionally, the third alloy (PWA 1480+) was created due to the prevalence of primary creep in rhenium bearing alloys. Full length creep test s continued with the s econd batch of material. The second phase of testing was performed af ter observing the primary creep behavior of PWA 1484. In order to view a similar condition for all three alloys, it was decided to stop samples following 0.5% secondary creep. This level of creep ensure d that primary creep mechanisms were still obvious while allowing for useful comparison be tween alloys. Stopping primary creep specimens at a specific amount of primary creep (e.g. 0.3%) would not have been very useful due the significant difference in prim ary creep behaviors expressed by the alloys. Additionally, ending the in terrupted creep tests during the ma ximum primary creep rate would only allow useful examination of PW A 1484. PWA 1480 and PWA 1480+ exhibited continuously declining creep ra tes beginning almost immediat ely following loading of the specimens. As a result of these concerns, the deformation mechanisms were observed for

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114 specimens just entering the secondary creep stage. Finally, another benefit of interrupting creep early in secondary creep is that data from the entire primary creep regime are preserved for all three alloys as will be presented below (for example, total primary creep strain, maximum primary creep rate, minimum secondary creep rate). Full Length Tests PWA 1480 The first generation PWA 1480 exhibits a brief primary creep stage followed by a continuously increasing creep rate. None of the four creep conditions pr oduced true steady-state secondary creep behavior in PWA 1480. This is especially obvious when viewing the creep rate change with temperature. The minimum creep rate is usually recorded so on after the end of the primary creep stage. The creep rate then slowly increases thr oughout the remainder of the creep test. As a result, a true sec ondary creep stage is not observed. The initial cree p rates of PWA 1480 reflect the early onset of prim ary creep and are typically larger than all creep rates to follow until late in the tertiary creep stage. Immediatel y after primary creep the creep rates are at their minimum values and significantly lower than the maximum primary creep rates earlier in the life of the specimens, Table 6-1. The complete lifetime of the PWA 1480 specimens at 704 C/862 MPa, 760 C/690 MPa, and 815 C/621 MPa can be seen in Figures 6-1, 6-2, and 6-3. Additionally, creep test results are given in Table 6-2 for all three a lloys. A number of points can be made about the creep behavior of PWA 1480. First and as alrea dy mentioned, tertiary creep dominates the life of the alloy at all three test conditions. Second, as the temperature is lowered from 815 C to 704 C, and the load increased, the performance of PWA 1480 improves significantly. Third, the amount of primary creep is virtually unchanged for all three test conditions and bot h age heat treatments. And

PAGE 115

115 finally, the creep behavior of PWA 1480 does not seem to have a strong dependence on age heat treatment temperature. While tertiary creep behavior is dominant fo r PWA 1480 at all test conditions, changes in creep rates and lifetimes are clearly evident. For instance, time to 1% creep (t1%) at 704 C/862 MPa and 760 C/690 MPa are similar; however, time to 2% creep (t2%) begins to show some variation in performance with the higher temperature test resulting in a shorter t2% time. When the temperature is raised still further to the 815 C test condition the t1% and t2% values are decreased significantly. The creep rupture life of PWA 1480 declines from the lowest temperature test condition to the highest temper ature test condition. As the creep life decreases with increasing temperature, the minimum creep ra tes increase with increasing temperature. The creep rates at 815 /621 MPa are more than doub le the creep rates at 704 C/862 MPa. The reduction in lifetime and creep rate performance has not affected the amount of primary creep, however. The primary creep strains produced by PWA 1480 at all three conditions are unchanged. The amount of primary creep is between 0.3% an d 0.5% for all test conditions even though the maximum creep rate during primary creep is sign ificantly increased at higher temperatures, Table 6-1. Changing the age heat treatment temp erature does not change these behaviors in a predictable manner. In fact, no obvious dependence on age heat treatment was found. One exception to this was found during interrupted creep testing, Figure 6-4. Here, the LT aged specimens experienced less creep than their HT age counterparts. The amount of primary creep, however, was the same. Overall, the performan ce of both heat treatments was similar enough to be within the expected range of scatter for cree p testing. Additionally, th ere is no change in the amount of primary creep strain s due to age heat treatment.

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116 PWA 1480+ The addition of Re to PWA 1480 that improved the tensile properties discussed in Chapter 5 also significantly improved the creep behavior of the alloy. The already low primary creep of PWA 1480 was reduced in PWA 1480+ and the creep rate under all heat treatment and test conditions was significantly lowe r than any other alloy/heat tr eatment combinations in this investigation, Figures 6-1 to 6-3 and Tables 61 and 6-2. Also, the ru pture lifetimes of PWA 1480+ specimens were much longer than the same of PWA 1480 specimens. The overall lifetime and creep rate performance of the e xperimental alloy were significantly improved; however, creep ductility was reduced for all co nditions. Most PWA 1480+ samples failed with less than 5% creep elongation. This reduction in ductility is likely a dire ct consequence of the large strengthening effect of the Re addition to PWA 1480. The alloy may have been hardened to such a point that ductility is greatly reduce d. Though not tested during this investigation, the fracture toughness of PWA 1480+ is potentially reduced as well. Similarly to PWA 1480, the age heat treatments did not create any obvi ous differences in creep performance. The primary creep strain s produced with both age heat treatments were similar. While PWA 1480 exhibited similar primar y creep strains at all test temperatures, PWA 1480+ exhibited lower primary creep strains at higher temperat ures, Table 6-1. The primary creep strains of PWA 1480+ specimens at 704 C/862 MPa are in the same range as those produced in PWA 1480 without Re. At 815 C/621 MPa, PWA 1480+ exhibits less than half the primary creep of the lower temperature specimens. The creep behavior of PWA 1480+ still maintain ed some similarities to the behavior of PWA 1480. For instance, the minimum creep rate slowly climbed during the test and the overall creep behavior appeared to be dominated by ter tiary creep. The rise in creep rates of PWA

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117 1480+ were much less pronounced than the rise in minimum creep rates of PWA 1480. Additionally, primary creep for both alloys was similar, with the exception of 815 C. This remains true despite a significant increase in t1% and t2% times, indicating that the amount of primary creep that is attained is not necessarily li nked to the amount of time required for the completion of primary creep. For example, th e interrupted creep tests of PWA 1480 and PWA 1480+ were terminated around 0.75% creep elonga tion for TEM analysis. The PWA 1480 tests were terminated between 10 and 25 hours, Figure 6-4. The PWA 1480+ samples, however, finished between 250 and 450 hours for the same elongation, Figure 6-5. The amount of primary creep produced was the same in this case at both temperatures and for both age heat treatments. If it was only a matter of time to complete primary creep, PWA 1480+ would likely have produced much less prim ary creep than PWA 1480. In actuality, it appears that the total primary creep strain is the determining factor. Another way of stating this is to say that the onset of secondary creep te rminates primary creep. This reasoning may sound redundant; however, the onset of se condary creep is related to wo rk hardening processes in the alloys. Once significant work hardening is pr oduced to slow the primary creep deformation processes, secondary creep begins. The primar y contributor to the ons et of secondary creep, then, is not time, but strain. The amount of work hardening necessary for an alloy to enter secondary creep will vary based on microstructura l characteristics and composition. As will be shown later, substantial work hardening takes place in PWA 1480 and PWA 1480+ during the early stages of primary creep leading to rapi d work hardening and only 0.3% to 0.5% primary creep strain. PWA 1484, however, does not exhibit this behavior and instead demonstrates very little dislocation and stacking fault interaction during primary creep lead ing to reduced work hardening. Finally, because strain is the controlling factor in th e onset of secondary creep, it is

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118 not surprising that PWA 1480 and PWA 1480+ can produce the same amount of primary creep with large differences in the time required. PWA 1484 The creep behavior of the second generati on alloy PWA 1484, however, displays a brief incubation period followed by a large primary creep strain. Primary creep typically ends very soon after it starts by transitioning to secondary creep with a constant, low creep rate. Despite the large primary creep strains, most of the r upture lives of the PWA 1484 specimens occurred during secondary creep. Plotted as creep elongati on (%) vs. run time (hr.), the secondary creep stage of these alloys is very nearly linear. As th e test nears the eventual failure life of the alloy, the creep rate slowly begins to increase. The rising creep rate continues to increase until failure, and thus comprises the tertiary stage of creep for PWA 1484. Compared to PWA 1480 for the highest temper ature test, PWA 1484 has a rupture life that is roughly 10 times (1000%) the rupture life of PWA 1480, Table 6-2. As the temperature is lowered (and the stress increases), the relative difference decreases. At 760 C, PWA 1484 has a lifetime of about 4 times (400%) the lifetime of PWA 1480. At 704 C, however, PWA 1480 has the greater rupture life. In the LT age condition, for example, the rupture life of PWA 1484 is now only 1/10 (10%) of the life of PWA 1480, Figures 6-1 through 6-3. The prominence of the primary creep stage of PWA 1484 is evident for all test conditions used in this investigation. Generally, as the temperature is increased and the load decreased, the magnitude of primary creep decreases, Figure 6-6. Alternatively, as the temperature is reduced and the load is increased, th e magnitude of primary creep in PWA 1484 increases. By comparison, there does not appear to be a signif icant difference in primary creep as a result of test condition in PWA 1480 and PWA 1480+, Fi gures 6-4 and 6-5. PWA 1484 also shows a

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119 fairly strong primary creep dependence on age he at treatment. For all PWA 1484 creep tests, the LT aged specimens produced significantly more primary creep than the HT aged specimens. The effect of test temperature/stress and age heat treatment can be readily seen in Figures 6-1 to 6-3, and 6-6 to 6-7. The HT age reduced the pr imary creep strain of PWA 1484 by up to 10% of the LT primary creep strain at a ll temperatures. Alternatively, th e primary creep strains of the HT aged specimens have been reduced to as low as 50% of the primary creep strains of the LT aged counterparts. Post-test measurements revealed that th e PWA 1484 specimens deformed non-uniformly during creep testing. Inhomogeneous creep has been shown to occur frequently in second generation superalloys like CMSX-4 and PWA 1484.13, 47, 55 All of the PWA 1484 specimens were elongated from the original circular cross-se ction to an elliptical cross-section. The major and minor axes of the ellipse were parallel to <110> directions. It has been reported that the elliptical cross-section is form ed mostly during primary creep if few slip systems are active. After the initiation of secondary creep, work-hardening should result in greater uniformity of deformation due to the activation of multiple slip systems. Following this reasoning, the difference in cree p rate between the two age heat treatments of PWA 1484 at 815 C/621 MPa, Figure 6-3, may be due to changes in cro ss-sectional area during primary creep. With a smaller cross-sect ion the LT age sample will deform faster than the HT age, giving the impression of a higher creep rate. Contrasting this reasoning are the data produced at 760 C/690 MPa, Figure 6-2. Here, an even greater difference in primary creep strain yielded nearly identical secondary creep ra tes between the two age h eat treatments. More testing is needed before the cause of the vari ed secondary creep rate can be stated with confidence.

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120 Interrupted Tests Interrupted creep results are presented in Figur es 6-4 to 6-6 for all three alloys and both age heat treatments. From these results it is immediately clear that PWA 1484 exhibits unusually large primary creep strains. Additionally, the amount of time that elapses during primary creep changes significantly for each alloy. In order of increasing time to complete primary creep: PWA 1484 (5-15 hours), PWA 1480 (12-22 hours) and PWA 1480+ (230-410 hours). This order also applies to decreasing amounts of primary creep produced during interrupted testing. Comparing the age heat treatments for the th ree alloys reveals a few correlations of interest. First, the primary creep behavior of PWA 1480 may have sli ght age heat treatment dependence; however, as already mentioned the overall behavior of PWA 1480 does not appear to follow a dependence on age heat treatment te mperature. Second, the primary creep of PWA 1484 has a strong dependence on age heat trea tment temperature, Figure 6-6. At both temperatures, the HT age specimens resulted in the lowest primary creep which is consistent with full length creep testing. By comparison, at 704 C the HT age reduced the primary creep of PWA 1484 by 32% and at 815 C reduced the primary creep by 67%. Finally, PWA 1480+ shows no correlation to age heat treatment, whic h is also consistent with full length testing. The test temperature also br ought about some behavioral changes. Both PWA 1480 and PWA 1480+ exhibited less time to reach 0.7% creep at 815 C. The minimum and maximum creep rates were slightly higher at 815 C than at 704 C for both alloys, Table 6-1. PWA 1484 exhibited a much more dramatic change in prop erties with increasing temperature. While PWA 1480 and PWA 1480+ produced approximately the same amount of primary creep at both temperatures, the primary creep of PWA 1484 was greatly reduced at 815 C for both age heat treatments. For the LT age, the pr imary creep strain of PWA 1484 at 704 C was 24.44% and

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121 3.68% at 815 C. For the HT age, the results were 16.66% at 704 C and 1.20% at 815 C. For the LT age, this represents an 85% reduction in primary creep. For the HT age it is a 93% reduction. Transmission Electron Microscopy (TEM) Transmission electron microscopy was perfor med to produce a qualitative understanding of the active deformation mechanisms because the two reported deformation mechanisms are distinct in appearance. First, the more comm on deformation mechanism reported in superalloys consists of matrix dislocations bowing to fill the matrix channels and creating interfacial dislocation networks.19, 27, 56, 57 This mechanism applies to most superalloys due to the growing incoherency of the / interface during a creep test. De formation usually begins in the matrix and this leads to the formation of interfacial netw orks to one degree or anot her in all nickel base superalloys. The second deformation mechanism commonly associated with primary creep occurs by stacking fault shear of the precipitates and has been reported in several alloys.16, 17, 23, 24, 33, 58 The alloys most commonly linked to this behavior are second generation and later alloys like PWA 1484 and CMSX-4; however, one fi rst generation alloy has been reported to exhibit this method of creep deformation as well.33 Following the interrupted creep tests it would be expected that all of the specim ens would exhibit sufficient formation of / interfacial dislocation networks to cause th e onset of secondary creep, whic h was found to be true for the specimens examined here.16, 19 PWA 1480 The deformation of PWA 1480 during interrupted creep consisted primarily of matrix dislocations bowing to fill the matrix channels, Figures 611 and 6-12. Most of the deformation appears to be limited to the matrix phase; however, a low amount of stacking fault

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122 formation was also found within the specimens, Figures 6-13 to 6-15. The stacking faults that were found tended to be limited to a range of one or two precipitates only. It will be shown that the limited nature of stacking fault formation st ands in contrast to the prevalent nature of the stacking faults in PWA 1484. The limiting of dislocations to the matrix during primary creep is also consistent with low amounts of pr imary creep as reported by several sources.16, 17, 19, 20, 27, 33, 56, 57 Additionally, the presence of secondary precipitates was confirmed by TEM, Figure 616. PWA 1480+ The second generation version of PWA 1480, PWA 1480+Re, exhibited a significant change in behavior (relative to PWA 1480). Mo st notably was an increase in the number of stacking faults within the material, Figure 6-17. Additionally, stacking faults were often seen interacting with stacking fau lts lying on different slip plan es, Figures 6-17 to 6-19. While stacking fault formation is tied to large primar y creep strains, they ar e typically limited to a single slip system in alloys exhi biting large primary creep strains.16, 17, 33 PWA 1480+, however, was shown to exhibit low amounts of primary cr eep. Additionally, the fact that the stacking faults appear to be interacting with other stacki ng faults, indicating slip on more than one slip system, is worth exploring. These interactions are common in superall oys displaying stacking fault shear, but not until the final stages of creep leading to failure.33 These interactions in PWA 1480+ occurred early in the life of the specime ns during primary creep and failure is not imminent. These interactions, coupled with the active matrix dislocations, may be generating relatively large amounts of strain hardening that br ings about the onset of secondary creep earlier than alloys, such as PWA 1484, that exhibit stack ing fault shear on a single slip system during primary creep. Additionally, stacking faults in PWA 1480+ are limited to regions spanning only

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123 2 to 3 precipitates (or fewer). The localized natu re of the stacking fault shear may be a result of these stacking fault interactions restricting the propagation of the shear bands as they form. In this way, the ability to produce strain is reduced and the expected primary creep strains would be lower. Despite the formation of larger numbers of stacking faults, deformation is limited and the secondary creep stage occurs early (in terms of total strain, not time). PWA 1484 The PWA 1484 specimens, when observed on the TEM, appear to share the same deformation mechanisms widely reported among second generation alloys at low temperatures and high loads as used in this investigation. Wide-spread stacking fault sh ear that appears to act on a single slip system was apparent over large re gions within the specim ens, Figures 6-20 to 621. This highly planar deformati on mechanism, as reported by others16, 17, occurs by the passage of two a/2<110> matrix dislocations (same slip plane with burgers vectors at 60 to each other) into the precipitates. These dislocations then dissociate into a/3<112> and a/6<112> partial dislocations with stacking faults in between. The two pairs of pa rtials (with stacking faults in between) are separated by an anti phase boundary (APB) due to the ordered nature of the precipitates. It is this complex system of dislocations, stacking faults, and the associated anti phase boundary that is able to sh ear large distances with relati vely little impedance, generating large primary creep strains.17 The numerous rows of stacking fault ribbons in PWA 1484 can be seen in Figures 6-20 and 6-21. Additionally, a large density of matrix dislocations has formed in PWA 1484 following primary creep, Figure 6-22. Agai n, this is consistent with the onset of secondary creep. These shear bands are also conservati ve in nature and leave no dislocation debris at the / interfaces (assuming no other dislocations are interacting with the stacking fault ribbons). As a result, the ribbons are free to expand and glide unt il they interact with other

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124 matrix dislocations. These matrix dislocations will eventually form interfacial networks; however, the amount of primary cree p that has been produced is alr eady quite large by this point. Therefore, by the time enough work hardening is pr esent to force the onset of secondary creep in PWA 1484, the amount of primary creep that has been conferred can be quite large. The matrix dislocations shown in Figure 6-22 are the likely cause of the beginning of the secondary creep stage.

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125 Table 6-1. Primary creep and creep ra tes from interrupted creep tests. 704C/862 MPa 815C/621 MPa Alloy Age HT Primary Creep (%) Max. Primary Creep Rate (%/hr.) Minimum Creep Rate (%/hr.) Primary Creep (%) Max. Primary Creep Rate (%/hr.) Minimum Creep Rate (%/hr.) LT 0.35 0.1632 0.0116 0.34 0.6001 0.0277 PWA 1480 HT 0.49 0.4015 0.0166 0.32 1.0377 0.0418 LT 0.36 0.0073 0.0007 0.22 0.0339 0.0021 PWA 1480+ HT 0.34 0.0296 0.0011 0.11 0.0304 0.0022 LT 24.44 10.704 1.8697 3.68 2.4563 0.0768 PWA 1484 HT 16.66 4.3877 0.7010 1.20 0.6751 0.0454 Figure 6-1. Creep at 704 C/862 MPa of all three alloys.

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126 Figure 6-2. Creep at 760 C/690 MPa of all three alloys. Figure 6-3. Creep at 815 C/621 MPa of all three alloys.

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127 Table 6-2. Rupture lives and tota l creep elongation from full-length creep tests. Also included is time to 1% and time to 2% creep. Note: The PWA 1480+ HT sample at 815 C/621 MPa failed before 2% creep was achieved. Alloy Age t1% (hr.) t2% (hr.) trupture (hr.) Elongation (%) LT 55.2 199 842 16.77 1480 HT 52.0 135 424 14.52 LT 63.0 427 1223 5.52 1480+ HT 619 1225 1721 3.32 LT 7.19 8.33 90.3 38.81 704 C 862 MPa 1484 HT 10.1 11.7 23.7 19.64 LT 56.6 117 311 10.78 1480 HT 47.4 124 428 13.90 LT 3.27 4.10 1465 20.08 760 C 690 MPa 1484 HT 5.96 13.9 1525 14.69 LT 11.2 28.8 99.3 15.01 1480 HT 14.4 26.8 77.2 14.93 LT 486 797 1239 8.94 1480+ HT 372 N/A 505 1.45 LT 1.15 1.18 391 13.40 815 C 621 MPa 1484 HT 4.92 49.8 914 19.06

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128 Figure 6-4. Primary creep of PWA 1480. Figure 6-5. Primary creep of PWA 1480+.

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129 Figure 6-6. Primary creep of PWA 1484. Figure 6-7. Creep at 704 C/862 MPa of PWA 1484 compared to PWA 1480 and PWA 1480+. Inset: magnified view of the creep behavior of PWA 1484.

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130 Figure 6-8. Primary creep comparison at 704 C/862 MPa for all three alloys. Note: large strain range was used better viewing of the behavior of PWA 1484. Figure 6-9. Primary creep comparison at 704 C/862 MPa for all three alloys. Note: long time range was used for better viewing of the behavior of PWA 1480 and PWA 1480+.

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131 Figure 6-10. Primary creep comparison at 815 C/621 MPa for all three alloys.

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132 Figure 6-11. Bright field(a)/D ark field(b) pair of deform ation in PWA 1480 (HT age, 704 C). Creep deformation is primarily limited to the matrix and interfacial dislocation networks have already formed. a b

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133 Figure 6-12. Bright field(a)/D ark field(b) pair of deform ation in PWA 1480 (HT age, 704 C) revealing the formation of few stacking faults. a b

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134 Figure 6-13. Bright field TEM image of disl ocation networks in PWA 1480 following primary creep. Figure 6-14. Stacking fault and di slocation shear of PWA 1480 is limited to small regions within the specimens (bright field).

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135 Figure 6-15. A stacking fau lt in PWA 1480 (dark field). Figure 6-16. Secondary precipitates (marked by arrows) in PWA 1480 following interrupted creep.

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136 Figure 6-17. Stacking fau lt interactions following pr imary creep in PWA 1480+. Figure 6-18. Stacking fault interactions and a dislo cation network in PWA 1480+.

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137 Figure 6-19. Short range stack ing fault shear of PWA 1480+. Figure 6-20. Bright field TEM image of PWA 1484 LT (704 C) following interrupted creep. Inhomogeneous deformation by stacking fault shear of the precipitates is apparent.

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138 Figure 6-21. Stacking fault shear of precipitates in PWA 1484 (bright field).

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139 Figure 6-22. Bright field(a)/Dar k field(b) pair showing the interfacial dislocation networks present in PWA 1484. Note: stacking fault shea r is also present, but out of contrast (marked by arrows). a b

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140 CHAPTER 7 RESULTS: ADDITIONAL CHARACTERIZATION In addition to common metallographic ch aracterization, mechanical testing, and transmission electron microscopy, two additional appr oaches were used to characterize the three alloys in this investig ation. The first is a re cently developed method called the Local Electrode Atom Probe (LEAP). This technique improves upon ideas established with the Scanning Atom Probe (SAP) and the 3 Dimensional Atom Probe (3DAP) that were developed in the early 1990s. The LEAP allowed for high resoluti on compositional characterization in three dimensions and was used to observe the secondary in PWA 1484 as well as the segregation behavior of Re to the matrix. The results presented in the LEAP section of this chapter were conducted at the University of North Texas by Anantha Puthicode and Mike Kaufman in conjunction with the University of Florida. The second characterization method included in this chapter is X-ray diffraction (XRD). X-ray diffraction was used to study the lattice misf it of all three alloys an d to observe how misfit changes with heat treatment. Additionally, the data collected by XRD needed to be processed to separate the contributions of the and phases that otherwise overlap. This extra step involved deconvolution of the intensity data produced by th e (002) and (004) planes within the alloys. Both of these techniques were used in an attemp t to gain an understand ing of the fundamental differences between the alloys as a result of chemistry and processing. Local Electrode Atom Probe (LEAP) The Local Electrode Atom Probe is capable of near atomic reso lution in both three dimensional space and mass number. As a result these instruments are capable of rendering three dimensional representations of the distribut ion of atoms within a specimen. Due to the relatively recent development of the LEAP, a brie f discussion of its operation will be included

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141 below. A full review of the development and capability of the LEAP can be found in (Kelly and Larson, 2000).52 At its most basic, the LEAP functions by extrac ting ions (atoms) from a specimen and accelerating the ions towards a detect or. With the system highly calibrated, the time of flight of an ion is used to determine the mass and the lo cation of the ion on the detector determines the location in two dimensions of the point of origin for the ion at the tip of the specimen. The third dimension (along the axis of the specimen) is controlled by careful adjustment of the specimen and rate of extrac tion. A two dimensional side-view schematic of the device is given in Figure 7-1. As shown in the figure, the specimen is posi tioned beneath the extraction electrode. The extraction electrode resembles a hollow cone with th e tip removed to create an opening at the tip. It is through this opening that the extracted ions are accelerat ed towards the detector. The secondary electrode is a disk (with a concentric hole in the cente r) of larger diameter than the base of the extraction electrode. The presence of the secondary electrode is required because the LEAP requires a relatively low extraction potential, thus the ions must be accelerated towards the detector. Additionally, the extraction elect rode can be pulsed with a high frequency to control the rate of extraction events.52 Specimens for the LEAP take the form of tall na rrow cones resembling spikes or needles. These spikes need to be very small in diameter with a tip radius less than 10 to 50 nm. Smaller tip radii aid in focusing the applie d field at the tip, improving the ab ility to extract ions from the specimen. The specimen size utilized for this inve stigation consisted of a diameter of less than 50 nm and a length (of analyzed volume) of at least 150 nm. The complete method of sample preparation is given in Chapter 3. Creating the specimens, however, involved two stages of electropolishing in order to achieve the dimensions required for the LEAP system, Figure 7-2.

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142 The first step was a bulk electrothinning opera tion that resulted in tip radii between 250 nm and 500 nm. The second step used a small platin um loop to refine the tip to a radius below 50 nm. All LEAP specimens in this investigation are oriented parallel to the [001] direction (also parallel to the applied stress axis). Add itionally, the compositions of the individual LEAP specimens may vary slightly based on location. Due to incomplete homogenization, a specimen near a dendrite will be more enriched in Re, Mo, and W and depleted of Al and Ta, while a specimen near the interdendritic region will express the opposite. PWA 1484, however, exhibits relatively little segregation and a large solution heat treatment window so these effects should be small. Four sets of LEAP specimens were created from a single bar of PWA 1484. The four sections of PWA 1484 received heat treatments to induce changes in th e microstructure for observation of secondary and the segregation behavior of rh enium. Two sections received the LT age and two received the HT age. One sectio n from each of the two age heat treated groups was then subjected to another br ief solution heat treatment. This extra solution heat treatment was conducted with the goal of quickly dissolving the precipitates but not allowing enough time for the enriched Re layer around the prior interface boundaries time to homogenize. The samples were then quenched rapidly. The goal wa s to create a condition that would allow the Re shells (or clusters) to be examined by the LEA P method. The results of the LEAP analysis are presented below and separated by heat treatment. It should be noted that a successful analysis of PWA 1484 LT (age heat treatment only) could not be accomplished due to losses during specimen preparation.

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143 PWA 1484 LT Reconstruction (solution HT) The first of the LEAP reconstructions is gi ven in Figure 7-3. Fi gure 7-3a is an 18% aluminum iso-surface and Figure 7-3b is an 18% chromium iso-surface from PWA 1484 LT with the extra solution heat treatment following aging. An iso-surface is a concept used to develop an understanding of the distribution of specific elements within the sp ecimen. In the case of Figure 7-3a, any region with a concentration of at leas t 18% Al will create a surface that encompasses the region. In this case, the areas that are colored in green are most likely precipitates, while the areas with no color are most likely the matrix. The precipitate with its thickness fully contained within the LEAP speci men is potentially a secondary precipitate based on the measured thickness of 75 nm. Primary precipitates, by comparison, in PWA 1484 measure between 300 and 500 nm in thickness. The sm all, Al-rich regions in Figure 7-3 in the matix between the larger secondary precipitates measure less than 5 nm in diameter. These locations may mark clustering of Al atoms leading to the development of secondary precipitates. Without further characterization, no specific conclusion can be made as to the structure of these regions. Figure 7-3b is an 18% Cr iso-surface. Because Cr pa rtitions more to the matrix, the the phase becomes apparent. It sh ould be noted, however, that the phase still contains a significant Cr concentration and, as a result, blue surfaces appear in the phase as well. Composition profile (solution HT) The image in Figure 7-4 is an SEM image of the specimen tip used to create the reconstruction in Figure 7-3. The line along the sp ecimen axis is the direction and length of the analyzed volume for both the reconstruction in Figure 7-3 and for the composition profile in Figure 7-5. The composition profile in Figure 7-5 is very similar in appearance to a composition

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144 line scan that can be produced on an Electron Probe Micro Analysis (EPMA) microscope except with greater accuracy. In actuality, the profile produced by the l eap is created from a cylinder of user defined length, diameter, and orientation within the collected data. This cylinder can be oriented parallel to the axis of the specimen or perpendicular to the specimen or at any other angle that is deemed useful. In the case of the profile in Figu re 7-5, the cylinder of analyzed volume was oriented within the center of the specim en (concentric) and parallel to the long axis. From the compositional profile, it can be seen that fine precipitates are present between the distances of 20 and 95 nm and 125 and 170 nm. Studying the profile reveals several distinct partitioning behaviors among the elements added to PWA 1484. First, aluminum and tantalum partition strongly to the phase. As a result, the composition of these elements is much greater in the phase than in the matrix. Additionally, it can be seen that Al partitions so strongly to the phase that a significant depl etion of Al exists for the matrix in the vicinity of the precipitate. The second behavior of interest results from those elements that partition to the matrix. These elements are rhenium, molybdenum, chromium, and cobalt. Rhenium and chromium appear to partition the most the phase with a slight enrichment layer on the side of the / interface. Molybdenum only slightly partitions to the phase as it expresses significant solubility in both the and phases. The third behavior is exhibited by tungsten. Tungsten does not partition particularly str ong to either phase and, as a re sult, maintains relatively uniform composition throughout the analyzed volume. This be havior is interesting as previous research has indicated that W additions enable increased Re solubility in the phase.35 If W does not partition strongly to the phase like the other solid solution strengtheners, then a significant portion of the added W is present in the phase. This allows for the possibility that W additions

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145 increase the average lattice parameter in the resulting in a synergistic effect between W and Re. PWA 1484 HT Reconstruction (age HT) The next reconstruction is presented in Fi gure 7-6 for a PWA 1484 HT specimen in the age heat treated condition. This recons truction is an example with all of the recorded ions present. Each ion species is recorded as a different colo r dot placed within the volume at its place of origin. Figure 7-7 is another iso-surface constructi on to illustrate the part itioning behavior of Cr (Figure 7-7a) and Al (Figure 7-7b). From these figures it is apparent that the lower phase is most likely due to the high Al concentration. The upper phase is most likely No fine precipitates were observed for this specimen, however, it should be noted that only 30 nm in the length of the specimen were analyzed due to excessive specimen thickness. Composition profile (age HT) The composition profile given in Figure 7-8 was produced from the analyzeld volume presented in the reconstruction in Figure 7-6. This profile is si milar in nature to the one in Figure 7-5 except for a higher scale on the compositi on axis to allow for th e Ni profile. Based on the composition profiles it is apparent that the phase is present over the distances 0 to 20 nm and the phase is present over the distances 20 to 30 nm. All of the partitioning behaviors discussed above are present with the addition Ni partitioning to the phase (due to the Ni3Al formula). Unfortunately, few conclusions can be drawn from this specimen due to the small volume that was analyzed. Additionally, significant enrichment and/or depletion were not observed with this specimen.

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146 Reconstruction (solution HT) The final reconstruction was produced with a PWA 1484 HT specimen exposed to the extra solution heat treatment following aging. The two images in Figur e 7-9 are two different views of the same reconstruction. Figure 7-9 is an 18% Al iso-surface, again coloring regions of green. The precipitates all appe ar to be secondary with thicknesses of 20 nm or less (except for the relatively large precipitate in the bottom corner). Additionally, as with Figure 73, there are many small aluminum rich regions of less than 5 nm in thickness. Again, it is difficult to conclude if these are ultra-fine precipitates or just Al rich clusters that are precursors to precipitate formation. Another represen tation of the same reconstruction is provided in Figure 7-10. Here onl y the Al, Ta, Cr, and Mo ions are displayed. Because Al and Ta partition to the phase and Cr and Mo partition to the phase, the contrast between the two phases is still evident. From both Figure 7-3 a nd Figure 7-9, it is appa rent that PWA 1484 is capable of producing secondary precipitates. It has alread y been shown that PWA 1480 and PWA 1480+ are capable as well via SEM techniques, Chapter 4. Composition profile (solution HT) Finally, Figure 7-11 is a composition profile generated from the reconstruction in Figure 79. Figure 7-12 illustrates the location and orientation of the cylinder of material selected to produce the composition profile. Th is region spans two secondary precipitates between the distances 8 and 15 nm and 36 and 47 nm. The meas ured thickness of these two precipitates is then 7 and 11 nm, respectively. Slight local enrichment by Re and depletion of Al in the matrix in the vicinity of the 7 nm precipitate can be seen on the profile. Additionally, the partitioning behaviors already discussed were also present for the PWA 1484 HT (solution HT) specimen as expected.

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147 Secondary Concentrations One benefit of the LEAP system is the ability to determine the concentrations of the secondary precipitates. These concentrations were determined from the line scan data generated with cylinders of data selected from the reconstructi ons. Of the most interest to the current investigation are the differences in secondary concentration as a f unction of precipitate size (thickness). These precipitate compositi on data were collected from the composition profiles in Figures 7-5, 7-8, and 711. Contained within these profiles are the entire thicknesses of 3 precipitates and partial thicknesses of 2 precipitates. The compositions, listed in order of thickness, are given in Table 7-1. The alum inum concentration is similar in all five precipitates (roughly 18 wt%). Th e elements chromium, molybdenum, and tungsten were also maintained at the same composition regardless of precipitate size (approximately 2wt% for Cr, 3wt% for Mo, and 3wt% for W). The elements nickel, cobalt, tantalum, and rhenium, however, all varied with precipitate size. As the precipitates increased in size, the concentrations of Ni and Ta increased while the concentrations of Co and Re decreased. The relative changes in composition for these four elements are given in Table 7-2. Compared to the smallest precipitate, the largest precipitate (in this set of measurements) exhibite d an increase in Ni concentration of 10wt%, an incr ease in Ta concentration of n early 5wt%, a decrease in Co concentration of 7wt%, and a decrease in Re conc entration of nearly 2wt% These results appear to be consistent with normal precipitate coar sening behavior. As the precipitate grows, the concentration of the formers would be expected to grow (Ni and Ta) while the elements that are rejected from the phase would be expected to decrease in concentration (Co and Re). The elements Cr, Mo, and W, however, do not displa y this behavior as they maintain uniform composition with precipitate size.

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148 X-Ray Diffraction (XRD) The X-ray diffraction study was initiated to dete rmine the lattice misfit of the three alloys in this investigation: PWA 1480, PWA 1480+, and PWA 1484. Sections were taken from the single crystal bars of each alloy and heat treated to 4 hr., 10 hr., 100 hr., and 1000 hr. at 1080 C. This temperature was selected as a continuation of the coating heat treatment cycle that these alloys were subjected to prior to service whic h is consistent with e xperimental procedures utilizing over-aging heat treatments between 950 C and 1100 C used in other sources.27, 59-62 Additionally, this temperature was selected, rather than the LT or HT age temperature, because the higher temperature allows faster coarsening and a better approximation of the equilibrium structure after 1000 hours. Following the over-aging heat treatments, th e specimens were thinned and polished to produce the best results. Upon recording the intensity data, p eak deconvolution was employed to differentiate the contributions of the and phases in agreement with other sources working with nickel base superalloys.46, 59 Figure 7-13 illustrates the de convolution process. Two peaks are assigned to both phases creating a total of 4 peaks. The two peaks for each phase represent the Cu k 1 and Cu k 2 wavelengths of X-ray radiation to which the specimens were exposed. Due to the slight difference in wavelength betw een the two, the diffracted peak produced by each is slightly displaced. This e ffect necessitates the separation of both wavelengths for each phase accounted for in a given peak. The MDI Jade (ver 7) software is programmed to account for the difference in Cu k 1 and Cu k 2 radiation and is calibrated to the diffractometer. Creating four peaks from the original is a mathematic operation and takes place utilizing an iterative algorithm. Upon completion of a deconvolution routine, a report is generated detailing critical information about the peaks and the fit of the model. Thes e reports have been incl uded in full in Appendix

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149 B. Repeated deconvolution of the same data se t does not guarantee the same result every time. Depending on the selected starti ng conditions and the choices made among seven peak fit options, these results can vary. As a result, it is critical that each report is observed to verify if the model is useful or if it has strayed too far from the expected result. An example of this is Figure B-57 in which the contribution of the phase has been almost entirely eliminated due to the ability of the phase to fit the original intensity da ta without any furthe r contribution by the phase. In this case, the deconvolution was repeated producing Figure B-59. Due to these concerns, repeated deconvolution was necessary to generated accurate predictions that could be averaged to produce a valid la ttice parameter measurement. The results of the deconvolution procedures were then processed using a spreadsheet to calculate the lattice parameters and lattice mi sfit values from each of the alloys. These calculations are also presented in Appendix B. The results presen ted here are the most reliable data from the deconvolution proce ss. The lattice misf it of the three alloys was found for the 4 hr., 10 hr., and 1000 hr. specimens. The 100 hour specimens produced low peak intensities and were not useful for deconvolution and subsequent misfit calculation. The calculated lattice parameters for the (002) plane are provided in Tabl e 7-3. These results are also plotted in Figure 7-14 along with the limited results from the (004) plane. From observation of Figure 7-14, it is apparent that PWA 1484 has the largest, negativ e misfit value of the three alloys. This is expected as PWA 1484 contains more solid solution strengtheners th at partition strongly to the matrix, including Re which is known to partition to the phase particularly potently. The lattice misfit of PWA 1480, however, varies on the slightly negative side of zero. The magnitude of the misfit of PWA 1480 is low (less than 0.15%). The addition of Re to PWA 1480 resulted in a

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150 continually decreasing (from positive to negative, but increasing in magnitude) lattice misfit from an initially positive value of +0.034% to a final value of -0.114%. The results from the (004) plane were more difficult to obtain due to the lower intensity produced through diffraction. As a result, fewe r specimens yielded strong enough intensity measurements over the 2 range 117 to 120 For this reason, fewer results are reported for the (004) plane in Figure 7-14. Additionally, the gr eater spacing achieved at this high angle proved more difficult for the deconvolution process. In creased difficulty in data acquisition coupled with decreased analysis accur acy resulted in fewer reportable results. The results shown in Figure 7-14 consistently indicated more positive values for lattice misfit for all three alloys. PWA 1480, for instance, is calculated to have a positive misfit of roughly the same magnitude that was calculated from the (002) data. PWA 1480+ is also shown to have positive misfit of the same range in magnitude. Finally, the (004) result for PWA 1484 after 1000 hours is over 0.1% more positive representing greater than 50% difference in calculated misfit values from the same specimen. The result of these inconsistencies is an apparent need to continue examining the lattice misfit of these alloys. Alternative methods (TEM based) al so likely need to be employed to verify the results from th e XRD deconvolution process.

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151 Figure 7-1. Schematic illustrating the basic f unction of the LEAP system. The specimen is positioned beneath the cone of the extractor anode and ions are extracted using a pulsing voltage. The ions are then accelerated towards the detector by a secondary electrode.52

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152 Figure 7-2. LEAP specimens before, (a), and after, (b), the final polishing step with the Imago Electropointer. a b

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153 Figure 7-3. Iso surfaces created with the LEAP system (PWA 1484 LT with solution HT). (a) 18% Aluminum surface, (b) 18% Chromium surface. a b

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154 Figure 7-4. Magnified (SEM) view of the tip analyzed in Figure 7-3. Figure 7-5. Composition profile from the specime n in Figure 7-3 (PWA 1484 LT with solution HT). Fine precipitates are present at distan ces between 20 and 95 nm and between 125 and 170 nm (Corresponds with Figure 7-3a). 0 5 10 15 20 25 020406080100120140160 Ni % Co % Cr % Al % Mo % Ta % W % Re %Composition Distance

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155 Figure 7-6. The distribution of all recorded ions for PWA 1484 HT with no additional solution heat treatment. Figure 7-7. Iso surfaces created with the LEA P system (PWA 1484 HT no solution HT). (a) 18% Chromium surface, (b) 18% Aluminum su rface. These results are from the same specimen shown in Figure 7-6. a b

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156 Figure 7-8. Composition profile from the speci men in Figure 7-7 (PWA 1484 HT no solution HT). A primary precipitate is present at distances between 20 nm and 30 nm (corresponds with Figure 7-7a)Composition % Distance (nm) 0 10 20 30 40 50 60 70 0 5 10 15 20 25 30 Ni % Co % Cr % Al % Mo % Ta % W % Re %

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157 Figure 7-9. Iso surface (18% Aluminum) created with the LEAP system (PWA 1484 HT with solution HT). Secondary precipitates are apparent in the matrix.

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158 Figure 7-10. LEAP results with only Al, Ta, Cr and Mo ions represented (PWA 1484 HT with solution HT). The results shown here are from the same specimen reported in Figure 7-9.

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159 Figure 7-11. Composition profile from the specimen in Figures 7-9 and 7-10 (PWA 1484 HT with solution HT). An illustration of the volume of material used for this composition profile is given in Figure 7-12. Figure 7-12. Illustration of the data selected for the compositional profile shown in Figure 7-11. The cylinder is user selected and only the data contained within this region is counted in the profile. Composition % Distance (nm) 0 10 20 30 40 50 60 05101520253035404550 Ni % Co % Cr % Al % Mo % Ta % W % Re %

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160 Table 7-1. Composition (wt%) of secondary precipitates of varyi ng diameter in PWA 1484. Diameter: 6nm Ni Co Cr Al Mo Ta W Re 53 17 3 18 3 <1 3 2 Diameter: 11nm Ni Co Cr Al Mo Ta W Re 54 16 2 19 4 1 3 1 Diameter: 74nm Ni Co Cr Al Mo Ta W Re 58 14 2 18 3 2 3 <1 Diameter: 46+nm Ni Co Cr Al Mo Ta W Re 62 9 2 18 3 3 3 <1 Diameter: 10+nm Ni Co Cr Al Mo Ta W Re 63 10 2 18 2 5 2 <1 Table 7-2. Net changes in secondary concentration with increasing precipitate size. Particle size (nm): 6 11 74 46+ 10+ Net Change (wt%) % change Nickel 53 54 58 62 63 +10 +19 Cobalt 17 16 14 9 10 -7 -41 Tantalum <1 1 2 3 5 +5 +1000 Composition (wt%) Rhenium 2 1 <1 <1 <1 -2 -75

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161 Figure 7-13. An example of the deconvolution pro cess used to separate th e contributions of the and phases. Notice, both phases have two peaks associated with them from the Cu k 1 and Cu k 2 wavelengths. Table 7-3. Lattice misfit (%) for the (002) plane following heat treatments at 1080 C. Heat treatment time at 1080 C 4 hr. 10 hr. 1000 hr. PWA 1480 -0.075 +0.002 -0.138 PWA 1480+ +0.034 -0.052 -0.114 PWA 1484 -0.228 -0.251 -0.201

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162 -0.30 -0.20 -0.10 0.00 0.10 0.20 0.30123Heat TreatmentLattice Misfit (%) 0911-002 0911-004 0922-002 0922-004 1212-002 1212-004 Figure 7-14. Lattice misfit vs. heat treatment fr om measurements of both the (002) and (004) peaks. Note: The heat treatments for each point are given by nu mbers (1: 4 hr. at 1080 C, 2: 10 hr. at 1080 C, and 3: 1000 hr. at 1080 C). The 100 hr. at 1080 C heat treatment is not shown due to the unreliable nature of the data produced from samples given the 100 hour heat treatment. Also the sample names are as follows: 0911 is PWA 1480, 0922 is PWA 1484, and 1212 is PWA 1480+.

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163 CHAPTER 8 DISCUSSION Through the course of the investigation into the primary creep behavior of PWA 1480 and PWA 1484 several relationships between processi ng and creep behavior were observed. The composition and processing an alloy has received pl ays a vital role in determining the nature of the creep behavior that should be expected. Numerous studies have been performed to incorporate the many inherent differences in mi crostructure into a material based modeling scheme. The following list includes several mate rial conditions that are impacted by chemistry and/or processing often in cluded in creep models: Stacking fault energy (both and phases)13, 63, 64 Anti-phase boundary energy ( phase only)13, 63, 64 / morphology o Precipitate dimensions and distributions21, 23, 63 o The absence/presence of secondary precipitates16, 20 o volume fraction63, 65 o channel thickness66, 67 o composition (relates to the strength of the phase)11, 68, 69 o composition (relates to the strength of the phase)10, 11, 68, 69 o Microstructural stability ( coarsening, rafting)38, 70-73 Chemistry effects o Segregation (eg. Re segregation to the matrix)10, 11 o Short range order (eg. presence of DO22 ordered clusters of W and/or Cr)69 o Formation of secondary phases (carbides and TCP phases)35, 38 Active deformation mechanism o Dislocation shear of the matrix21, 74 o Stacking fault shear of the phase12, 13, 16, 17, 23, 64 o Mixed-mode deformation63 Each of the above points can impact the behavior of a single crystal superalloy during creep. Consequently, it is necessa ry to investigate any relationships that may exist between these material attributes in order to simplify any mode ls that are derived. Th e current investigation represents an attempt to unders tand the effect of secondary precipitates and the element rhenium on the primary creep behavior of firs t and second generation superalloys PWA 1480 and

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164 PWA 1484, respectively. A thir d alloy, named PWA 1480+ for th is investigation, was also produced by adding 3wt% Re to PWA 1480 to make a second generation version of PWA 1480. The following discussion will focus on the differe nt behaviors among the three alloys produced by two different aging heat treatment schemes a nd alloy chemistry. Special attention will be given to several of the material at tributes listed above and their e ffect on the creep behavior of all three alloys at low temperatures (700 C-815 C) and high stresses ( >500 MPa) where primary creep dominates the creep behavior of PWA 1484 Finally, methods of applying the knowledge gained from this investigation towards creep behavior modeling and future alloy development will be discussed. Microstructure The three alloys used for this investigation a ll share several microstructural similarities. First, the typical cuboidal / microstructure for high volume fraction superalloys extends to all three alloys. Second, all th ree contain a low volume fraction of carbide phases due to the addition of between 0.02wt% and 0.04wt% C. Additionally, secondary precipitates were found in all three alloys with both age heat treatments. All of thes e similarities were impacted by alloy chemistry and processing. / Morphology The cuboidal microstructure common among modern single crystal nick el base superalloys is produced due to a combination of the high volume fraction of and a negative lattice misfit.40 In the case of PWA 1480, the la ttice misfit is small and only slightly negative at room temperature. The addition of Re to create PWA 1480+ did not signifi cantly reduce the lattice misfit as expected. For both PWA 1480 and PW A 1480+, the measured values of lattice misfit at room temperature would be expected to beco me more negative as the temperature is raised

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165 due to thermal expansion effects. The low misfit values have led to slightly rounded precipitates and slightly larger channels (than PWA 1484). A be nefit of low lattice misfit, and rounded precipitates, is increased stability agai nst coarsening and coherency loss. As the misfit is increased, the internal stress near the / interfaces is increased. Under an applied stress, dislocations will be attracted to the interfaces to alleviate the high stresses. Additionally, diffusion-based processes will oc cur to reduce the misfit stress es still further. The second generation PWA 1484, however, has a larger misfit resulting in sharper edges and corners on the cuboidal precipitates. Rhenium additions have also been tied to changes in the / morphology.8, 75 In particular, the addition of Re has been linked to smaller precipitate size and an increase in rounded edges (to the point of spherical shapes in so me lower volume fraction superalloys).8 Within the present investigation, the e ffect of Re on the / microstructure is small. Following the HT3 solution heat treatment, there is littl e difference between PWA 1480 and PW A 1480+ in size and shape of the phase. The effect of Re, as reported in the l iterature, is caused due to the low diffusivity of the element in both and During cooling from the solution heat treatment, the supersaturated nature of the phase brings about nucle ation of fine primary precipitates. The low diffusivity of Re prevents these precipitates fr om growing significantly during cooling. Upon subsequent aging heat treatments, the growth of the precipitates is slower due, again, to the rejection of the slow diffusing Re from the precipitates into the matrix. Additionally, the rounded corners of precipitates were caused by Re due to reduced growth kinetics. These results were produced in alloys with lower volume fractions than are present in PWA 1480 and PWA 1484.8, 75

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166 Precipitate coarsening is also driven by the Gibbs-Thompson e ffect by which the reduction in interfacial energy drives th e growth of larger precipitate s at the expense of smaller precipitates.76 The Oswald coarsening of precipitates in a matrix has been shown to be impacted by a number of factors including co mposition, coherency, precipitate si ze, radii of curvature near corners (morphology), and applied stresses.76 In the case of single crystal superalloys, both primary and secondary precipitates are subject to the Gibbs -Thompson effect. This effect can usually be seen in the matrix near the / interface where the population of secondary precipitates is reduced. The grow th of the much larger primary precipitate depletes the matrix near the / of formers and causes the elimination of the ultra-fine secondary precipitates in this region. These growth processes occur during isothermal aging heat treatments in order to obtain optimal mechanical properties s hown to be maximized with a primary size between 0.30 m and 0.45 m for high volume fraction nickel-base superalloys.36 Modelling the compositions of precipitates with varied sizes and shapes can be performed based on this effect. Another application of the Gibbs -Thompson effect can be found with the application of an external stress at elevated temperature on the primary precipitates. The Gibbs-Thompson equation can be modified to account for applie d stress on coherent, or partially coherent, precipitates in a matrix. The application of st ress can create directional growth of normally spherical or cuboidal precipitates do to the supe rposition of interfacial stresses that vary by location around the precipitates. As these stresses are relaxed, the microstructure can become directionally oriented. This has been shown experimentally as we ll as theoretically and has been called rafting or topo logical inversion.70, 71, 76 For the alloys in this inves tigation, these effects are limited by the high equilibrium volume fraction of The predicted equilibrium volume fractions for PWA 1480, PWA 1480+, and

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167 PWA 1484 are shown in Figure 414. The super-saturated phase, upon cooling, nucleated a fine dispersion of coherent precipitates. Because the equilibrium volume fraction of was so high, it is expected that the this thermodynamic driving force caused growth of the precipitates, possibly with slightly rounded corners, until the narrow channels became depleted with former elements. As two precipitates grew with parallel {001} faces approaching each other, the depletion of the channel adjacent to the centers of the precipitate surfaces in formers slowed the local growth of precipitates. The channel regions near what would become the corners of the precipitates would have c ontained a greater amount of formers and precipitate growth would have continued until these regions also became depleted. The result would be cuboidal precipitates with fairly well defined corn ers and edges even if Re additions would dictate slightly rounded co rners in lower volume fracti on alloys. Comparing the / microstructure of PWA 1480 and PW A 1480+ revealed nearly identical shapes and sizes despite the addition of Re. The channel widths also shared th e same size and shape within both alloys. PWA 1484, however, exhibited sharper corners and slightly narrower channels. This effect can be described by comparing the all oy chemistry with that of PWA 1480. Aside from the addition of Re which would be expect ed to increase misfit and create rounded cubes, an increase in Co and reduction in Ta were made in the newer alloy.77 The presence of both of these elements has been shown to reduce the la ttice misfit which would lead to further rounding of the precipitate edges.8, 32, 75 Countering these effects, though, is the increased diffusion coefficient produced by increasing Co and reduc ing Ta. Reported results have shown that increased cobalt concentrations can reduce or ev en eliminate any adverse effects caused by the addition of Re.75 The increase in diffusion caused by these alloy modifications are possibly

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168 responsible for the slight improvement in shape and the reduction in channel thickness in PWA 1484 when compared to PWA 1480. Carbides Carbides were found to be present for al l three alloys. PWA 1480 and PWA 1480+ both contained about 0.04wt% C and PWA 1484 contained about 0.02wt% C. These small C additions were sufficient to bri ng about precipitation of carbide phases in the form of localized carbide networks. The level of carbon in the three alloys was insufficient to cause dendritic carbide formations.78, 79 All three alloys contained script carbide networks with a few small blocky carbides associated with the local networks. Carbon additi ons have been shown to reduce casting defect formation and reduce / eutectic formation.79 Because the carbide content for all three alloys was low, the carbides did not imp act the primary creep process significantly. Research has shown that carbide interfaces are often the site of void and crack formation leading to failure.39 It was not found that the carbides present in PWA 1480, PWA 1480+, and PWA 1484 played a significant ro le during primary creep. Topologically Close Packed Phases The addition of rhenium to PWA 1480+, while improving creep properties, brought about precipitation of plate-like TCP phases produced during high temperature exposure consistent with the formation of phase.35, 38 The TCP phases present in PWA 1480+ grew throughout the heat treated microstructure with the length parallel to <110> dire ctions. The presence of these phases during primary creep indicates that the all oy is very unstable with regard to TCP phase formation. A direct consequen ce of this is that PWA 4180+ is not a suitable alloy for commercial use. These phases are expected to be the cause of the reduced ductility exhibited during creep and tensile testing. It is also expected that the fa tigue life of PWA 1480+ would be

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169 reduced as a result of the formation of the TCP pr ecipitates. As a conseq uence, the applicability of PWA 1480+ is primarily limited to experimental testing and invest igation of the rhenium effect. The formation of TCP phases during prim ary creep occurred over a much longer time period than the time for completion of prim ary creep in PWA 1480 and PWA 1484 due to the reduction of creep rate caused by the addition of rhenium. Increased rhenium content has been tied to increased TCP phase content as a resu lt of the increase in overall refractory element content.35, 38, 50 While TCP phases have received a great amount of attention due to the acicular morphology that is often associat ed with their presence, PWA 1480+ demonstrated the greatest creep lifetime at every test condition. Th e PWA 1480+ specimens failed with low ductility which might be tied to the formation of cracks and voids associated with TCP precipitates. Interesting future work could pursue the link be tween TCP precipitation a nd the early failure of PWA 1480+. Additionally, slight modifications in alloy chemistry might be made that preserve the excellent creep behavior of PWA 1480+ while reducing the TCP content to increase lifetime and/or ductility. While all three alloys share seve ral similarities in microstructure, differences in performance are apparent as a result of the impact the microstructures of the three alloys in this investigation had on the active deformation me chanisms during tensile and creep testing. Tensile Behavior The tensile strength properties of the three a lloys of interest are n ecessary for a complete understanding of their respective creep behavior. The strain controlled natu re of a tensile test reveals a different material response than a load controlled creep test. For example, superalloys most often exhibit <110> type deformation in the phase.80, 81 Depending on a variety of factors, including lattice misfit, dislocations will either shear or bypass the precipitates during tensile testing. Shear of the phase often occurs with pairs of a/2<110> dislocations. Similarly

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170 to creep testing, the a/2<110> super-dislocati on pair can further di ssociate into partial dislocations separated by a complex fault.80 Deformation during creep, however, has been shown to occur through a variety of methods: dislocation shear of the matrix by <110> and <112> dislocations, shear by <112> super-dislocation pa irs and stacking fault pairs, and <010> cube slip.12, 13, 16, 17, 24, 56, 63, 74, 82, 83 Both test methods bring about complex deformation mechanisms. Tensile testing, therefore, is complementary to creep testing and can aid in generating a more complete understanding of the uniaxial properties of these alloys. As with creep deformation, tensile defo rmation primarily begins within the matrix through the generation of a/2<110>{ 111} dislocations. Initially, the microstructure is relatively free of dislocations. Upon yiel ding, these dislocations begin m oving through the microstructure as more dislocations are created by the availabl e sources. During the early stages of plastic deformation, dislocations begin bowing between precipitates, lining the / interfaces and filling the channels rather than entering and shearing the relatively hard precipitates. It would be expected that this proce ss will be affected by the pres ence (or absence) of secondary precipitates within the narrow channels. The ultra-fine secondary may play a role in tensile deformation by a number of possible methods. Secondary First, because secondary precipitates are very small in size (spherical and 10-50 nm in diameter), they would be highly cohe rent, increasing the probability of shear. In this case, the <112> dislocations would be likely to shear the without dissociating into partials due to the small size of the precipitates. Anti phase boundaries would be created and so additional dislocations would be drawn into these precipitates to relieve this energy. While shear of the secondary precipitate would be more difficult for th e first dislocation, the second dislocation

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171 would experience softening while eliminating APB.20 The second potential interaction between secondary precipitates and matrix dislocations is due to the possibility of dislocation bypass mechanisms. At high temperatures dislocation climb would be exp ected to play a vital role in bypass operations. Within the lower temperature range of 700 C to 815 C, however, climb is kinetically slow. As a result, cross-slip is more likely to account for precipitate bypass by dislocations. Both mechanisms would require added stress to accomplish. The composition of the secondary precipitates is also useful while considering either method of secondary -dislocation interaction. The composition of the secondary precipitates (determined by LEAP analysis, Tables 7-1 and 72) revealed that the precipitates contained lowered levels of Ta than is present in the base alloy (1-5 wt% vs. 9 wt% nominal composition). Additionally, in the smaller secondary the Mo content was nearly double the nominal alloy composition (3-4 wt% vs. 2 wt%), the Co cont ent was higher (14-17 wt% vs. 10 wt%), the W content was reduced to half (3 wt% vs. 6 wt%), and the Cr content was lowe r (2 wt% vs. 5 wt%). With reduced strengthener content, namely Ta, and lower solid solution strengthener content these precipitates may be lower in strength than the larger primary precipitates. As a result, the strength benefit due to the presence of secondary would be expected to be relatively small and interactions would be frequent due to the dense, fine dispersion within the matrix. The serrations that appeared in the plastic deformation regime duri ng tensile testing of PWA 1480, PWA 1480+, PWA 1484 are likely due to the presence of secondary Of the possible interactions described above, shear is the most likely due to the coherent nature (and potentially lower strength) of the secondary Similar interactions would be expected during creep with some exceptions. The differences that are observed during creep are due to

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172 differences in deformation mechanism and will be discussed further in the Creep Mechanism discussion. Channel Thickness Another microstructural feature that impacts tensile properties is the thickness of the channels. This feature is primarily important for the early stages of tensile deformation because as precipitate shear becomes more common in th e later stages, the channel widths become less important as dislocation bowing becomes less preferred. The importance of the channel width can be seen in Equation 8-1 that describes the stress increment to cause bowing of a straight dislocation between two precipitates: Equation 8-1: )2( r Gb Where the stress increment due to bowing is G is the elastic modulus, b is the burgers vector for the dislocation, is the center-to-cente r particle spacing, and r is the particle radius (assuming the particles are the same size). In th is case, it can be seen that the strengthening effect is increased if the inter-particle distance is reduced.84 Strengthening due to dislocation bowing is likely to play a role in the tensile deformation of PWA 1480, PWA 1480+ and PWA1484; however, there are severa l interactions that occur that complicate the problem. For example, the already mentioned secondary would serve to drastically reduce the inter-particle spacing (if no shear was assumed and bypass was not active). It is likely that a combination of these eff ects takes place during tensile deformation. For example, when a dislocation is stretched across a channel (pinned by the primary precipitates), the first strength increment to be accounted for should be the bowing stress from the primary precipitates. As the stress is incr eased, the dislocation will come against many fine secondary precipitates. The effective inter-particle spacing becomes significantly reduced and

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173 the bowing strength effect is increased. As the applied stress continues to increase, either the stress to shear the secondary or the stress to cross-sl ip and bypass the secondary is reached and the dislocation begins to glide. These inte ractions continue until the stress to shear the primary precipitates is reached and large scale shear becomes dom inant leading to eventual failure. Lattice Misfit Additional considerations are related to th e effect of alloy composition and, more specifically, the equilibrium concentration of the phase. Three aspects related to composition and heat treatment are lattice misfit, stacking fault energy, and anti-phase boundary energy. The magnitude of the lattice misfit can impact the resolved shear stresses within the channels. High misfit alloys experience large misfit stresses that can create compressive stresses in the channels parallel to the applied stress direction. These misfit stress es can be quite large. During high temperature, low stress creep testing, for in stance, these compressive stresses can exceed the applied external stress crea ting a very large driving force for microstructural change and, subsequently, drives rafting. During tensile testing, these comp ressive stresses can relieve the applied stress from a dislocati on and reduce the actual shear stre ss felt by the dislocation creating a strengthening effect. In order to relieve misfit stresses, dislocations can be attracted to the / interfaces to create dislocation networks. The formation of these networks increases the stress in the matrix leading to shear.65, 80 Stacking Fault Energy Shear of the phase usually involves the formati on of stacking faults as matrix dislocations enter the Stacking faults are regions of at oms shifted from normal lattice sites by the passage of a partial disl ocation. Stacking faults are contained between two partial

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174 dislocations that can be comb ined, with enough stress, to form a perfect dislocation. The stacking fault energy, SFE, is composition dependent and is related to the energy required to create the fault. The larger the SFE the smaller the spacing between partial dislocations. With large SFE alloys recombination a nd cross-slip is easier, but w ith low SFE alloys, the partial dislocations can spread out furt her and recombination becomes more difficult and interactions with other stacking faults and/or disloca tions becomes more likely. The SFE of the phase can also impede deformation of the phase. If the SFE of the primary precipitates is high, more resolved shear stress will be needed to force a di slocation to enter the precipitate. Shear of the precipitates is slightly easier with lower stacking fault energy alloys.65, 80 Anti-Phase Boundary Energy Related to SFE is the anti-phase boundary en ergy. While stacking faults displace atoms a partial atomic spacing, anti-phas e boundaries are created by displacing atoms in an ordered precipitate by a whole atomic spaci ng. The result is a change in nearest neighbor species type. For the highly ordered phase, an APB will position Al atom s next to Al atoms and Ni atoms next to Ni atoms across the boundary. The re sult is increased energy and resistance to deformation. To relieve the energy produced by the creation of an APB, a second dislocation must shear along the same plane to shift the atom s into proper atomic positions. Like the SFE discussed above, the magnitude of the APB relates to the energy re quired to create it. Alloys with large APB energies will have short disloc ation spacing and vise-versa for low APB energy alloys. There is a strength increment required to produce an anti-phase boundary that prevents dislocations from entering the phase. Once a single dislocati on has entered (either in whole form or dissociated into partials with a st acking fault), though, a second dislocation will be attracted into the phase to relieve the APB. The c onsequence of this behavior is the

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175 observation of stacking fau lt pairs that shear the phase. As a result of each of these strengthening effects, the tensile deformati on of single crystal s uperalloys can become complex.65, 80 Tensile Results The results of the tensile testing co nducted on PWA 1480, PWA 1480+, and PWA 1484 revealed several differences between the al loys. The first gene ration superalloy PWA 1480 provides a basis of comparison for both the Re modified PWA 1480+ and the second generation PWA 1484. Both second generation alloys are de scendant from the original PWA 1480 alloy. As discussed earlier, the transition from PW A 1480 to PWA 1484 during the development cycle was brought about by several key changes in com position including Cr, Co, Ta, Ti, and Re. The creation of PWA 1480+ was performed by adding 3 wt% Re (the same amount added to PWA 1484) to PWA 1480 to create a second generation version of the alloy. The second generation superalloy, PWA 1484, has demonstrated significant improvements in high temperature creep capability compared to PWA 1480.77 A consequence of the changes in composition that led to improved creep strength was a decrease in tensile strength for PWA 1484 relative to the original PWA 1480. The yield strength of PWA 1484 is over 400 MPa lower at 700 C and over 300 MPa lower at 815 C. PWA 1480 exhibited high yield strength, UTS, and failure strength values for both aging heat treatm ents and test temperatures used in this study, Table 5-1. PWA 1484 exhibited significantly lower yield strengt h, as mentioned earlier, but demonstrated a greater ability to work ha rden following yielding. In fact, at 815 C the failure strength of PWA 1484 LT exceeds the failure strength of PWA 1480 LT. Additionally, the strength of the HT version of both alloys at 815 C is only separated by 16 MPa at failure despite

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176 a 214 MPa disadvantage in yield strength. Clear ly the alloy design approach taken for PWA 1484 significantly changed the mechanical beha vior of the alloy relative to PWA 1480. While the behavior of PWA 1484 is very diffe rent from the origin al PWA 1480 alloy, the performance of PWA 1480+ was quite similar to PWA 1480. The modified alloy is marked by a greater yield strength (in the HT condition) and less ductility. This response is not unexpected due to the addition of rhenium. Rhenium has long been known to be a potent solid solution strengthener as well as a retardant of coarsening.77 Simply adding Re raised the yield strength of PWA 1480+ HT by 42 MPa at 700 C and162 MPa at 815 C. The LT condition, however, exhibited lower yield strength than HT condition at both test temperatures. Another effect of the age heat treatments that were given to the three alloys can be seen in PWA 1480 with which the LT age produced greater than 10% ductility during tensile testing. The HT age, however, significantly reduced the du ctility, yield strength, and UTS values. The HT age does not, however, produce this same effect in PWA 1484. The HT aged specimens of PWA 1484 exhibited a more rapid increase in flow stress during work hardening, but the ductility and UTS values were similar to the LT aged specimens. Additionally, it is noteworthy that while the tensile behavior of PWA 1480 was significantly altered by age heat treatment and PWA 1484 was not, creep testing, as discussed below, produced creep behavior of PWA 1484 that was significantly altered by age heat treatment while PWA 1480 was not. Creep Behavior Primary creep at low temperatures and high stresses has beco me an important topic as models of creep behavior are being developed for modern superalloys. The early superalloys required relatively simple to design creep models These first generation superalloys, such as PWA 1480, exhibit tertiary creep behavior across a broad range of temperatures and stresses.

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177 Modern superalloys, beginning with the second generation superall oys, however, exhibit a range of creep behaviors depending on the stress and te mperature regime to which they are subjected. For CMSX-4, for example, the three creep regimes that are displayed are primary creep (T<850 C, >500 MPa), tertiary creep (700 C950 C, <200 MPa).63 Each of these creep regimes are related to the microstructure and, consequently, the processing that the alloy was subjected to prior to servi ce. Additionally the behaviors at these three temperature/stress regi mes have been shown to be highly dependent on the chemistry of the alloy. Tertiary Creep The most common creep regime that is encounte red in testing of single crystal nickel-base superalloys is tertiary creep. For example, the alloy PWA 1480, a first generation superalloy, exhibits increasing creep rate as creep strain accumulates (tertiar y creep), over a wide range of temperatures and stresses including those used for this investigation. Tert iary creep behavior is has been linked to a continuously increasing di slocation density implying a creep softening process.63 Often, the tertiary creep rate can be modeled with the relatively simple equation given below: Equation 8-2: ) 1(0 C Where is the creep strain rate, 0 is the initial creep rate, C is a fitting constant and is creep strain.43 The dislocations responsible for ter tiary creep are primarily a/2<110>{111} type dislocations that are mostly contained within the matrix. These dislocations, under the applied stress, are forced to bend between the primary precipitates and stretch such that they line both sides of the channels. As dislocations from di fferent sources begin to meet along the /

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178 interfaces, nodal networks are formed. As the di slocation density continues to rise, the spacing between dislocations at the in terfacial networks is reduced. For alloys that display tertiary cr eep in the absence of rafting, the / microstructure is relatively stable for the duration of the creep test. This is largely due to the fact that shearing of primary precipitates is low and migration of vertical channels to horizontal channels is slow. Failure during creep of si ngle crystal superalloys typically involves the condensation and coalescing of vacancies to create voids. Ofte n, cracks and voids are nucleated in conjunction with casting porosity. In alloys containing hard, br ittle carbides and/or TCP phases, voids typically form at the interface be tween the hard particle and the more ductile matrix. For cases where these hard interf aces are not present (o r present in low amounts) it is thought that void formation may also occur through the devel opment of dense dislocation cells near / interfaces where dislocation annihilation can lead to a local increase in vacancy concentration.63 Changing the temperature and/or stress such that another creep regime becomes active reduces the accuracy of the basic creep rate model shown above. As rafting or primary creep become dominant, more complex models are required. Rafting Rafting in single crystal nick el base superalloys has become commonplace at low stresses and high temperatures. Rafting, also known as t opological inversion, is a process by which the original cuboidal precipitate morphology evolves into lo ng plate-like raf ts. The final orientation of the rafts depends on the lattice misfit between the and phases. Most often this misfit value is negative, leading to rafts oriented perpendicular to the applied stress. The rafting response occurs through diffusion of matter from thinner rafts through the channels to the thicker rafts. The result is a widening of the perpendicular channels. This morphological

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179 shape change results in the elimination of channels parallel to the ap plied stress direction. The <110>{111} matrix dislocations begin gathering along the / interfaces creating dense dislocation networks. Because matter is primarily diffusing across the channels during rafting, the channel widening can be described by a parabolic rate law: Equation 8-3: tCw1 Where w is the change in thickness, C1 is a material and temperat ure dependent constant (for isothermal creep), and t is time. Notice, equation 8-2 describes widening of the channels using a t1/2 time dependence. Coarsening of the precipitates, or Ostwald ripening, typically follows a t1/3 time dependence.32 As matter diffuses from the channels that are parallel to the applied stress into channels perpendicular to the applied stre ss, matter from the cube surfaces of the precipitates parallel to the applied stress is simultaneously diffusing into the parallel channels, effectively closing the channels oriented parallel to th e applied stress. The change in channel width due to this morphology change eff ect has been described by the following: Equation 8-4: 2 22)(2 ) 2( 2oo oooo MCsw swsw w Where wo is the initial channel thickness and so is the initial edge length. The morphology effect is used to calculate the amount of change in channel width as a result, simply, of the formation of rafts (ie. the change in shape necessary to create rafts). Fina lly, the parabolic rate constant from equation 8-3 can be found by: Equation 8-5: MC owwtww ) (

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180 Plotting w versus tallows for determination of the rate c onstant (slope of the line). Modeling the channel widening kinetics of superalloys experiencing rafting is necessary for the development of constitutive equations to pr edict the creep behavior at these conditions.66, 67 Primary Creep While numerous studies have focused on developing accurate descriptions of tertiary creep behavior and / rafting, the origin of primary creep has relatively little atte ntion until recently. Large primary creep strains produced at low temp eratures and high loads are typically found in second generation and later single crystal superall oys. While creep deformation in the tertiary creep regime is governed primarily by <110>{111} dislocation shear in the matrix, deformation during primary creep under these conditions is produced by pairs of stacking faults that cooperatively shear both the and phases. Specifically, stacking faults in the are formed by reactions of dislocations of the type a/2<110> to form a/3<112> and a/6<113> dislocations. An example reaction is given below: Equation 8-6: ]211[6/]211[3/]101[2/]110[2/a a a a 17 Between the a/3<112> and s/6<112> dislocations lies a stacking fault that exist in both the and phases. As the pair of dislocations (wit h the stacking fault between) enter an ordered precipitate an anti-phase boundary is created in the phase, but not in the phase. To reduce the energy required to create the APB, a sec ond pair of <112> dislocations enters the in the opposite configuration. As they progress, the APB is eliminated. An example of this configuration is given below: Equation 8-7: a/3<112>+SISF+a/6<112>+APB+a/6<112>+SESF+a/3<112>

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181 Where SISF and SESF are intrinsic and extrinsic stacking faults, respectively, and APB is antiphase boundary. It is this confi guration that is refered to as a dislocation or stacking fault ribbon.14 These dislocation ribbons, two stacking faults separated by an APB cooperatively shearing both phases, are capable of travel ing relatively long distances within th e material without leaving dislocation segments behind at the / interfaces. It should also be noted that while the first pair of dislocations is impeded by the formation of the APB, the second pair is aided by the elimination of the APB. Large creep strains can be produced by this mechanism in the absence of forest dislocations at the / interfaces. As common a/2<110> matrix dislocations expand to fill the channels and create in terfacial networks the difficulty of coope rative shear increases. Eventually, the networks that are formed are su fficient to reduce the rate of shear of the / microstructure and a steady-state condition is reached where the dislocation density would remain constant under constant stress creep conditions.16, 17, 63 Adding primary creep behavior to existing creep behavior models is a challenge due to the complexity of the behavior. For example, as shown in Figure 8-5, the dislocation pairs required to form the stacking fault pairs form from a/2<110> type dislocations. This implies that before cooperative shear can take place, a population of a/2<110> dislocations must already be present in the microstructure.16 For this reason, it is believed that this population is produced during the incubation period that often precedes primary cr eep. In this way, alloys that deform by shear exhibit the same deformation mechanisms prior to the start of primary creep that are found in alloys that display tertiary cr eep behavior only. Developing a m odel that can predict when the transition to / shear will take place has proven difficult.63

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182 Additionally, primary creep of supe ralloys does not always occur by shear due to stacking fault ribbons. Alloys e xhibiting tertiary creep do experi ence primary creep in the form of a higher initial creep rate that is reduced during the first stages of exposure to creep conditions. For conditions governed by tertiary creep, primary creep typically occurs by <110> dislocations bowing to fill the channels. This process is si milar to the deformation process exhibited during the later stages of tertiary creep; however, prior to the start of a creep test the native dislocation density is low. Due to this low initial density, the creep rate is more rapid than in later stages when dislocatio n-dislocation interactions are common. Primary creep under these conditions occurs as the small dislocation population is increased until the dislocation interactions bring about the transi tion to tertiary creep (ie. th e creep rate is reduced to the minimum creep rate immediately following primary creep).74 Creep Results The primary creep behavior of single crysta l nickel-base superalloys has proven to be controlled by multiple factors. Previous research has shown a dependence of primary creep on orientation, magnitude of th e load applied, secondary precipitates, lattice misfit, and possibly rhenium and/or ruthenium content, all of which may lead to non-uniform deformation (only 1 or 2 slip systems) and stacking fault shear (of the phase).16, 17, 19, 20, 23, 24, 27, 33, 47, 56-58 Comparing yield strength to the applied initial creep stress, Table 5-3, it is apparent that the large primary creep strains of PWA 1484 at 704C were produced at greater than 90% of the yield strength. This high load condition produced primary creep strains of 17% for the LT age and 24% for the HT age. Such a high stress level, however, was not necessary to produce large primary creep strains in PWA 1484. Also at 704C and a reduced load of 758 MP a, primary creep strains of 5% for the LT age and 10-14% for the HT age we re observed. This clearly indicates that the

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183 magnitude of the applied initial stress contributes to the amount of primary creep that results, as reported by (Rae and Reed, 2006) and (Shah et al., 2004).16, 47 Under creep testing at conditions that yield large primary creep stra ins, raising the initial stress results in larger primary creep strains and shorter times to complete primary creep. This lower stress ratio for PWA 1484 at 815C is similar to the stress ratios for creep testing of PWA 1480 and PWA 1480+ at 704C. Even still, the primary creep of PWA 1484 is significantly larger than the other two alloys. Additionally, a PWA1480 HT specimen was tested at 704C/1200 MPa which represents a creep stress to yield stress ratio of 0.9 (90%). The PWA 1480 HT specimen failed in 0.9 hours with no discernable primary creep stage because the creep rate exhibited continuous accelera tion beginning with the earliest measurements. These results demonstrate two related points. First, the ratio of creep stress to yield stress does not cause primary creep nor do high ratios increase the primary creep seen in PWA 1480. The second point is that the magnitude of the applied st ress, while increasing the primary creep strain produced in alloys prone to large primary creep strains, does not cause large primary creep strains in alloys resistant to primary creep. Related to the previous discussion are reports that there may be a stress threshold that, above which, stacking faults in the phase may form due to the entry of pairs of dislocations that dissociate into pairs of stacking faults separated by an APB, that then shear the phase. It is thought that at stresses below this threshold matrix dislocati ons do not shear precipitates but, instead, exhibit cross-slip or climb to bypass barriers such as the fine secondary in the channels. During this process work hardening is more likely due to the distribution of slip on multiple planes. If the stress exceeds this thresh old, however, climb and cross-slip do not occur and deformation is able to proceed without generating significant work hardening and secondary

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184 creep is delayed.16, 17, 20 Incorporating the findi ngs of the current inves tigation, these results are likely to only apply to alloys prone to primary cr eep. The distinguishing f actors that differentiate alloys exhibiting primary creep are not low yield strengths or the magnitudes of the applied loads, but microstructural differences as a resu lt of alloy chemistry and processing that resist primary creep strain production. When dislocations are confined to the matrix, primary creep is low due to the interaction of dislocations within the narrow spaces of the channels. During shear, dislocations can move longer distances without interacting with an obstacle in the precipitates resulting in large creep strains and few dislocation interactions.16 Creep behavior following the HT age heat treatments supports these findings. The longer time and higher temperature of the HT age allows the to coarsen, decreasing coherency. The greater degree of incoherency of the precipitates after the HT age may be responsible for the decrease in primary creep at all test conditions for PWA 1484 while no significant ch ange in primary creep can be found in PWA 1480 and PWA 1480+. Increasing incoherency in PWA 1484 following heat treatmen t may reduce the ability of dislocations to enter the phase to form stacking faults. PW A 1480 shows little effect with heat treatment because the formation of stacking faul ts is already difficult, Figures 6-11 to 6-15. PWA 1480+, while able to produce stacking faults, deforms by multiple deformation systems so that work hardening is rapid and secondary creep starts soon after the initiation of primary creep. Because both PWA 1480 and PWA 1480+ are prone to wide-spread dislocation and stacking fault interactions, leading to hardening, they show little change in primary creep strain with age heat treatment. Research into primary creep mechanisms has yielded interesting findings into the efficiency of these two shear mechanisms. Calc ulating dislocation dens ities and predicting the

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185 expected amount of strain produced for each m echanism, differences in the ability of the different processes to confer shear can be s een. For alloys that deform primarily by the movement of matrix dislocations, the matrix channels become filled with dislocations and dislocation interactions become common following 0.3% to 0.5% creep deformation. The large amount of dislocation interactions at this point is rela ted to the onset of secondary creep, where a balance exists between deformation processes (slip) and recovery processes. As a result, it is predicted that deformation provided primarily th e movement of matrix dislocations will only yield about 0.5% primary creep at the maximum. If the abilit y to form stacking faults is included, the dislocation density in the matrix increases more slowly. The result is a deformation process that can operate much longe r before work hardening in the matrix causes secondary creep. For this scenario, primary cr eep strains greater than 5% can be expected.16, 17 These results are consistent with the current investigation and those of other researchers.19, 20, 57 In the narrow sense of comparing overall rupture lives only, PWA 1484 out-performed PWA 1480 at all but highest load at the lowest temperature (704C/862 MPa). Simply adding Re to PWA 1480, however, remarkably improved th e rupture life of PWA 1480. Creep ductility and toughness were reduced, but the minimum creep rate was decreased by over an order of magnitude and the same low primary creep behavior was maintained. The time to 1% creep, however, shows the difference in primary creep behaviors very clearly, Table 6-2. While the lifetime of PWA 1484 is the longest at several conditions, the time to 1% creep is the shortest, indicating rapid deformation early in the life of the specimens. Continuing the discussion to time from 1% to 2% creep strain, it can be seen that PWA 1480 and PWA 1480+ experience a much longer time from 1 to 2% than from 0 to 1%. For several PWA 1484 samples, however, the time

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186 from 1% to 2% creep was achieved faster than the first 1% of creep strain because the creep rate was still increasing through the first several percent primary creep. If the usual benchmarks of rupture life, tim e to 1% creep, and minimum creep rate are utilized in turbine engine design then an alloy might be selected that could exhibit rapid creep during the early stages of its service life. Thus it is necessary to incorp orate an understanding of the nature of the active primar y creep behavior into the usual design schemes. Within the context of current investigation, a simple change in the aging heat treatment temperature of a turbine component may be enough to manage the primary creep that will be produced. Another method of avoiding large primary creep strains is to pre-creep the alloy prio r to service. Some reports indicate that specimens that usually yiel d large primary creep strains were crept to small strains (<0.1%) at 950C before being subjected to low temp erature, high stress creep testing. The primary creep produced during the second test was reduced to low values similar to those produced by PWA 1480 and PWA 1480+. It is thought that this reduction in the expected primary creep strain is related to the formation of matrix dislocations that then interfere with the passage of stacking faults and bring about the onse t of secondary creep earl ier than in specimens that were not pre-crept. These changes in pr ocessing are both examples of simple techniques that can dramatically improve the usefulness of a lloys that are prone to excessive primary creep by stacking fault shear. Modeling Primary Creep There have been several models developed to describe the creep proc ess for single crystal nickel-base superalloys since the introduction of PWA 1480 in the early 1980s. Most of these were made to describe the tertiary creep regime that is most common among superalloys. PWA 1480, for example, exhibits tertiary creep across a very wide range of stress and temperature

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187 including loads as high as 90% of the yield strength of PWA 1480 at 704C. The second generation superalloy PWA 1484, however, exhib its a more complex combination of creep behaviors. During high temperature, low stress creep PWA 1484 displa ys tertiary creep behavior. At low temperatures and low stress es, tertiary creep still persists. If the stress is raised above a critical threshold, however, PW A 1484 (and other second generation superalloys) produces large primary creep strains followed by steady-state creep.16 It is this transi tion that has proven difficult to predict with conven tional models often leading to large inconsistencies between models and experimental data.63 The alloy that has received the most atte ntion for primary creep modeling is CMSX-4.16, 17, 63 CMSX-4 and PWA 1484 share fairly similar compositions, Table 2-1, and the reported creep behavior and deformation substructures are simila r in nature to those found for PWA 1484 in the current investigation so modeli ng approaches for both alloys are expected to be similar.3, 16, 47 Deformation modeling of creep of superalloys usually begins by sepa rating the deformation gradient into constituents representing the matrix and the precipitates using an equation similar to the one given below: Equation 8-8: 12 1 12 1 ') ~ ~ () ~ ~ ( nd nd L f fyP where is the shear strain rate,d ~ is the unit vector representing the slip direction,n ~ is the slip plane normal unit vector, designates a slip system (12 systems included in this calculation), and f is volume fraction of each phase denoted by and subscripts. Because the model ignores any contribution by TCP phases or carbides, the sum of the volume fractions of and are set equal to one.

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188 The next step is to define the shear strain rate ( ). This step typical ly begins with the Orowan equation: Equation 8-9: vb where is the mobile dislocation density, b is the burgers vector, and v is the velocity of mobile dislocations. Note that Equation 8-8 is specific to a given slip sy stem. Also, shear strain rate calculations are performed for each phase indepe ndently. Substituting an expression for the mobile dislocation velocity gives: Equation 8-10: c b oro pass mis b slip attack FCC FCC FCCV Tk Tk Q Fb 110exp Equation 8-11: c b pass b slip attack L L LV TkTk Q Fb 112 1 1 1exp2 2 2 Where the subscripts FCC and L12 denote the and phase, respectively. Also, denotes dislocation jump distan ce, F denotes dislocati on jump frequency, kb is Boltzmanns constant, is the resolve shear stress. Also note, the misfit stress, mis, passing stress, pass, and Orowan stress, oro, are all accounted for in the phase, but just the passing stress is involved in the calculation for the phase. The Orowan stress as de fined in the model is based on the channel width and does not account for secondary within the channels. The strengthening contribution from secondary was not included as a separate contribution within any descri ption of the creep behavior so its contribution was most likely accoun ted for in the more general terms like passing stress. Each of these terms are described in detail as is the rest of the model in the work by A. Ma et al.63 As the model was developed, additional terms were added to account for additional behaviors to be described.

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189 The above described model goes on to account for <110> dislocation shear within the matrix as well as the generation of <112> stacki ng fault ribbons. To account for tertiary creep, the model ignores dislocation ribbons and shear because tertiary creep occurs almost exclusively by deformati on contained within the matrix. The distribu tion of slip between channels was found to be critical to the success of the model. The simplest microstructural constituents are a cube shaped phase, and three rectangular channels (two sides equal to the edge length of the precipitate and one side equal to the specified channel thickness) oriented in three directions with the normal to the plate parallel to the cube di rections [100], [010], and [001].63 With the applied stress axis pa rallel to the [001] direction, one channel is perpendicular to the stress [001] and two are parallel [100] a nd [010]. Misfit stresses were found to interact quite strongly with the active deformation mech anisms. Initially, deformation takes place almost exclusively in the perpendicula r, [001], channels due to the superposition of the misfit stress and applied stress. As long as the / interfaces remain coherent, the [001] channels are preferred. As the misfit stresses are relieved, however, the preference is reduced and deformation increases within the [100] and [010] channels.63 Lattice Misfit Lattice misfit plays a significant role during bot h tertiary creep and primary creep. Alloys with lower values for lattice misfit are reportedly more prone to the formation of the dislocation ribbons responsible for shear and, therefore, la rge primary creep strains.59-61, 71, 85 Large, negative lattice misfit values have been show n to produce better creep resistance at high temperatures and are thought to be beneficial to the prevention of large primary creep strains. Alloys with large lattice misf it have been shown to form dislocation networks along the / interfaces more quickly due the large misfit stresses that result. Dislocations then build up at the

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190 / interfaces to relieve the misfit. As the inte rface becomes lined with di slocations forming the interfacial network, shear of the phase becomes more difficult, the creep rate is reduced, and secondary creep begins. Alloys w ith low misfit are more coherent and it has been suggested that low misfit alloys are prone to shear of the phase. In addition, the di slocation ribbons that form glide relatively easily and leave little, if any, dislocation segments at / interfaces. These two factors combine to delay the formation of the interfacial networks that he rald the start of the secondary creep stage. As a re sult, the primary creep stage ( shear) occurs over a longer time allowing for the production of greater primary creep strains.16 In the current investigation, PWA 1484 was shown to have the largest, negative lattice misfit (-0.20% to -0.25%) of the three alloys tested in this investigation. This finding would tend to contradict conventional wisdom regarding the effect of lattice misfit due to the large degree of stacking fault shear in PWA 1484 during primary creep. PWA 1480 exhibited low lattice misfit values ranging from 0.00% to -0.14% and the Re addition to PWA 1480 increased the lattice misfit values to a slightly more positive ra nge (-0.05% to +0.034%) for PWA 1480+. These results suggest that the emphasi s placed on lattice misfit in cr eep modeling may need to be reevaluated. While it is clear, that lattice misfit plays a large role in tensile and creep deformation of superalloys, the more recent seco nd and third generation superalloys have larger magnitudes for lattice misfit than the older first generation alloys. The first generation alloys, however, do not deform by large-scale shear unlike the second generation alloys PWA 1484 and CMSX-4. Additional work is clearly necessary to clarify the effect of lattice misfit while modeling primary creep be havior these alloys.

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191 Secondary In addition to accounting for the nature of the / interface during creep deformation, the channels play a vital role in creep (both primary and tertiary). The channels are the location where <110> dislocations nucleat e and propagate leading to the nucleation of di slocation ribbons in the phase or / interfacial networks. In addition to channel thickness, discussed earlier, the secondary that forms within the matrix interact wit h dislocations as they expand to fill the matrix. Despite these in teractions, the secondary precipitates are omitted from most models even though channel thickness is nearly unanimously utilized for creep and tensile modeling. There is some disagreement as to which st age of heat treatment is responsible for precipitating secondary The presence of secondary has been shown to occur following solution heat treatment with coarsening occuri ng during the ensuing aging heat treatment69 and secondary has been reported to appear during ra pid cooling from the final aging heat treatment.20 In either case, the secondary are likely to always be present in the channels prior to service for most high volume fracti on superalloys. For PWA 1480, PWA 1480+, and PWA 1484, secondary was shown to be present in all th ree alloys following both age heat treatments. These precipitates are expected to be highly coherent with the matrix due to their small size resulting in an increased probability of shear by matrix dislocations bowing within the channels. Using the LEAP technique, compositions for a small sample of secondary precipitates in PWA 1484 were determined. While no statements can be made with high statistical certainty, four composition variation trends could be seen with increasing size. The larger secondary precipitates contained higher levels of Ni and Ta and lower levels of Co and Re. This indicates that as these precipitates, or potentially clusters, grow in size, strengthening is likely to occur due

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192 to the increased former content. The larger secondary precipitates reside towards the centers of the channels with smaller secondary precip itates near the rela tively large primary precipitates. As a result, the interaction between secondary and matrix dislocations is complex. Despite the differences in strength and size distribution across the channel, it is likely that matrix dislocations shear these ultrafine precipitates during tensile testing. Creep testing, however, allows for thermally activated processes to occur that mi ght alter the paths of dislocations when they meet these precipitates. For alloys that exhibit tertiary creep, such as PWA 1480 and PWA 1480+, matrix dislocations re main in the <110> form and likely shear through or cross-slip past the secondary Alloys that produce large primary creep strains, such as PWA 1484, deform largely by <112> dislocations separated by stacking faults. While these dislocations and stacking faults typically nucleate / interfaces prior to shearing the phase, this configuration is relatively stable in the matrix as well. As a result, these dislocation ribbons likely shear the secondary in a similar manner to shear of the primary It has even been suggested that the secondary may act to stabilize the di slocation ribbons through the channels between primary precipitates. Additio nally, it should be noted that there is no APB formed in the matrix so the dislocation pairs (and a ssociated stacking faults) tend to separate further apart in the channels.16, 17, 63 Clearly the secondary present in the channels plays a role in creep deformation during both primary a nd tertiary creep. The ability of PWA 1484 to produce large primary creep strains while PWA 1480 does not, though, must be linked to another material attribute as the uniform presence of secondary in all three alloys did not yield similar behaviors.

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193 Composition Alloy composition clearly controls several aspects of mechanical behavior. From controlling the strengths of phases to the lattice misfit, composition plays a direct role in deformation. Several key compositional cha nges were made from PWA 1480 to PWA 1484. While the effect of composition has already been discussed, a few points will be readdressed. Rhenium additions to superalloys have been praised for their potency in solid solution strengthening of the phase. This effect was so significan t that the first three generations of superalloys are defined by their Re content (0 wt%, 3 wt%, and 6 wt%, respectively). Despite the clear benefit in strength, the addition of Re has been identified as potential contributor to large primary creep strains due to the large numbe r of second and third generation alloys that show this behavior. While it is possible that Re has some sort of contributory effect on primary creep in second generation and late r alloys, this investigation s howed an improvement in creep properties and a slight decrease in primary creep when Re was added to PWA 1480. It was found, however, that PWA 1480+ appeared to produce more stacking faults than PWA 1480 indicating a possible decrease in the stacking fault energy of the phase. A lower stacking fault energy would allow dislocations shearing the phase to spread further leading to an increased potential of dislocation-disloca tion interactions. In fact, the stacking faults observed in PWA 1480+ exhibited many interactions similar to the deformation found in specimens of other alloys interrupted near failure. As a result of this a nd other investigations, th e effect of Re on the microstructure and properties of single crystal nickel base superalloys has been regarded as positive. Specifically, Re additions increase lattice misfit, improve creep life, improve tensile strength, reduce the size, and reduce the rate of coarsening.2-4, 8, 9, 77 One disadvantage with

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194 the use of large Re additions, though, is decreas ed stability against TCP phase formation as shown by the precipitation of TCP phas es during primary creep of PWA 1480+. There is also agreement within the superall oys community that Re segregates to the matrix during nucleation and growth of precipitates. This effect is expected to cause an increase in lattice misfit near the / interface as Re is rejected from the precipitates result in an enriched region in the matrix. Research has also shown, however, that in the presence of W, up to 20% of the Re that has been added will segregate to the precipitates.35 As rhenium additions modify the compositions of the and phases, changes in deformation mechanisms might be expected as a consequence of the potent strengthening effect. Aside from the well known solid solution streng thening effect are reports of the formation of Re-rich clusters in the phase near the / interface. Understand ing the nature of the distribution (in the matrix) of the Re that has been rejected from the phase could lead to improved model accuracy and potentially aid in th e search for a Re replacement. Reported research is still contradictory in this area. Th e two leading theories, currently, indicate that Re may either form clusters around 10 nm in diam eter or may form hardened shells in the matrix around the precipitates. Recent studies, including th e present one, have attempted to utilize the Local Electrode Atom Probe (LEAP) to examin e the local concentration distribution of Re in the matrix with limited results. While publicat ions have yet to be produced, a recently presented collaboration between Cambridge Un iversity (Cambridge UK) and Oak Ridge National Lab (Oak Ridge, TN) pointed to possible ev idence of the formation of Re clusters. The formation of clusters by solid solution stre ngtheners is not a new phenomenon and has been shown to occur with Cr and W as well.8, 35, 50 Chromium clusters have been found to have the DO22 crystal structure and are also fo rmed due to rejection from the phase. It is expected that

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195 these clusters would make dislocation shear of the phase in the vicinity of the / interface more difficult. Dislocations that encounter these regions are likely to be disrupted by changes in crystal structure and strength leading to greater buildup of dislocations at the / interface. How Re clusters or shells interact with dislocations and stacking faults during primary creep is certain to be the subjec t of many future investigations as, clearly, additional work is necessary. It still remains unclear what the in teractions are between th e rejected Re in the phase and the matrix dislocations. Additionally the high price of Re has forced both industry and academia to begin researching viable alternative strengthene rs to replace or reduce Re in future superalloys. Continued research into the rhen ium effect as it relates to these interactions is likely to grow in the coming years. The difference between PWA 1480 and PWA 1 484, however, is more than just the addition of rhenium. PWA 1484 has increased solid solution strengthener/r efractory content and decreased hardening. Based on the present investig ation, Re alone does not appear to cause increased primary creep. In fact, the opposite effect was realized in PWA 1480 when Re was added. With the many changes in alloy com position between the first generation PWA 1480 and the second generation PWA 1484, it is likely that the stacking fault energy and anti-phase boundary energy were also modified. These two pr operties very strongly in fluence shear of the phase during creep leading to the possibility of reduced resistance to large primary creep strains caused by shear. Concluding Remarks The large primary creep strains exhibite d by PWA 1484 were shown to occur through deformation processes consistent with shear reported among some second generation superalloys. The first generation alloy PWA 1480, however, exhibited tertiary creep behavior

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196 that is commonly reported among s uperalloys. The goal of this investigation was to study the effect of Re and secondary precipitates on the primary cree p behavior of PWA 1480 and PWA 1484. In order to see the effect of Re addi tions on PWA 1480, a third alloy was created by adding 3 wt% Re to PWA 1480 (the new all oy was named PWA 1480+ for this study). Through this investigation, several aspects of