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Effect of coating and exposure on the oxidation and mechanical properties of Ti-22A1-26Nb

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Title:
Effect of coating and exposure on the oxidation and mechanical properties of Ti-22A1-26Nb
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Dobbs, James Ross, 1962-
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English
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xviii, 132 leaves : ill. ; 29 cm.

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Alloys ( jstor )
Argon ( jstor )
Cooling ( jstor )
Fatigue ( jstor )
Microhardness ( jstor )
Oxidation ( jstor )
Oxides ( jstor )
Oxygen ( jstor )
Titanium ( jstor )
Titanium alloys ( jstor )
Dissertations, Academic -- Materials Science and Engineering -- UF ( lcsh )
Materials Science and Engineering thesis, Ph.D ( lcsh )
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non-fiction ( marcgt )

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Thesis:
Thesis (Ph.D.)--University of Florida, 1997.
Bibliography:
Includes bibliographical references (leaves 121-131).
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Typescript.
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Vita.
Statement of Responsibility:
by James Ross Dobbs.

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EFFECT OF COATING AND EXPOSURE ON THE OXIDATION AND
MECHANICAL PROPERTIES OF Ti-22A1-26Nb















By

JAMES ROSS DOBBS


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


1997

































Copyright 1997

by

James Ross Dobbs

































To my loving wife, Jane.















ACKNOWLEDGMENTS


The title of Doctor of Philosophy is bestowed on the individual, but the road to

obtaining that title is truly a team effort. My road was paved with many supportive,

capable individuals whom I would like to recognize. First and foremost I thank my wife,

Jane, for the support and love during a period in our relationship when there was

something other than her that commanded much of my time and attention. I thank my

parents and brothers and sister for their support and enthusiasm and for always being

willing to listen. I would also like to thank my thesis advisors, Prof. Mike Kaufman

whose friendship I value as much as his technical interaction and Dr. Mike Gigliotti who

never let me forget the focus of my goal. I would also like to thank Prof. Teresa Pollock

whose gentle reminders and kind inquiries gave me inspiration to finish.

I would like to thank several people who helped me conduct the experimentation

and understand the results: Mr. Mike Gilhooley for his help and guidance in LCF testing

and Mr, Chris Canestraro for his assistance in tensile testing, Mr. Keith Borst for

producing the sputter coatings, Mr. Amie Henry for teaching me the LPPS system, and

Mr. Kevin Janora for CVD coatings. I would like to thank Dr. Ravi Ravikumar and Mr.

Mike Larson for their guidance and patience while I learned and practiced TEM, Ms.

Cindy Hayden for assistance in conducting and interpreting XRD, Dr. Bob Gilmore for

acoustic microscopy and acoustic modulus measurements and Dr. John Ackerman for

guidance through many experimental decisions and his wealth of CVD knowledge and









his ability to convey that knowledge. A special note of thanks to Dr Mike Henry for

lengthy and insightful discussions on fatigue, and Prof. Harry Lipsett for many hours of

discussion on titanium metallurgy and relevant experimentation.

I would like to thank Dr. Ann Ritter for support, both financial and moral.

Thanks to Dr. Jeff Graves and Dr. Tom Cox for having enough faith in me to employee

me while I completed this work. And a special note of thanks to Ms. Karen Keever for

valuable assistance in preparation of this document.















TABLE OF CONTENTS


page


ACKNOW LEDGM ENTS ................................................................ ...................... iv

L IST O F T A B L E S ............................................................ ........................................... ix

LIST O F FIG U R E S ........................................................................... ...................... x

ABSTRA CT..................................................... .......................... xvii

CHAPTERS

1 INTRODUCTION...................... ....................... .............

2 LITERATURE REVIEW.................... ........ .....................5
Overview of Ti-Al-Nb M etallurgy.................................... ............................5
H istory..................................... ............ ... .... ......................5
Description of Phases Present in Ti-Al-Nb.................... ...... ............. 7
Deformation Behavior of C,, O and B2 Phases ...........................................14
Tensile Behavior of a2 and O Alloys................... ...........................17
Fracture Toughness and Fatigue ........................................ ......... ......22
Overview of Oxidation Behavior .............................. ....................26
Oxidation of Uncoated Titanium Alloys.................................................26
Oxidation of Coated Titanium Alloys...................... ............................28
A T i ...................................................................... ............................. 2 8
M CrA lY ........................... ............................. ...........29
Pure m etals ........................ ............................. 29
Glasses, silicates and oxides............................. .................... 30
Embrittlement of Titanium Base Alloys ................................................31
Effects of Coatings on Mechanical Properties of Ti Alloys...................33
Program A pproach................................................... .........................34

3 EXPERIMENTAL PROCEDURE .......................................................35
Material Processing ................... ........ .....................35
Specim en Blanking........................... ............................ .................... 42
Test Specim ens and Test Procedures ............................... ...........................45
Oxidation............... ..................................................................... 45









S am p les .................................................................. ............................4 5
T testing .................................................. ..................................... 45
T en sile .................................................................... .................................4 9
Sam ples .................................................. .................................... 49
T testing ...................................................................... ................. 49
F atigu e.................................................................... .................................5 0
Sam ples .................................................. .................................... 50
T testing ...................................................................... ................. 53
A analysis ............................................. ........................................ 54
Coatings.............................. ...............................56
Low Pressure Plasma Spray (LPPS) Coatings.............................................56
NiCrAlY + 40 vol% A1203............................ ................................57
FeCrAY............................................. ....................................58
Sputtered Coatings. .............................................. ......................58
Sputtered A l......................................... ......................................... 59
Sputtered Si ................... .. ........ ......................60
Sputtered Pt and Cr...................................................61
C V D coatings.............................................................. ...................... 61
S iO ..................................................... ........................................ ........6 3
T a20 .......... ......................................................................................63
MgO........................................................................................... 64

4 RESULTS AND DISCUSSION..................................... ...........................65
O xidation ................................................ ............................................... 65
Cyclic Oxidation Testing ............................................. .....................65
B baseline results................................. ......................................... 65
Al and Si Coated and Reacted Results ........................................66
Oxide Coating Results.................................................67
Metallic Coating Results ........................ ............. .....................68
Microhardness and Microprobe Evaluation............... .....................70
Baseline uncoated........................... ...........................70
FeCrAlY and NiCrAIY + A1203 ........................ .....................73
Sputtered and Reacted Al and Si.................... ...........................77
Sputtered Pt and Cr coatings ................................... ....................80
CVD Oxide Coatings.......................................................................82
Effect of Forging Strain and Cooling Rate on the Tensile and LCF
B behavior ................................................................................................ 86
Fatigue Results............................ .............. ......................86
Tensile R results ............................................................ ......................87
D iscu ssion ........................................................................... ....................8 8
Effect of Exposure at 6500C on the Tensile and LCF Behavior.......................95
Tensile Results ............. ..............................................................95
Fatigue Results ......................................................................96
Discussion ................ .........................................................98



vii









Effect of Coating and Exposure on the LCF Behavior...................................105
R esults............................................................................. ..............105
Discussion ............................................ 110

5 SUMMARY AND CONCLUSIONS ...................... .......... ......... 117
O xidation ........................................................................ ................. 117
Forging Location .................................................................. .................... 118
Environm ental Exposure .............................................................. .................. 119
Coated LCF Behavior............................. ....................120

LIST OF REFERENCES............................................121

BIOGRAPHICAL SKETCH ........................................................ ... 132









































viii















LIST OF TABLES


Table page

Table 3-1. Compositions from several locations along the length of the 15.25 cm
diameter billet, surface and center, and from the forging used in this study.............36

Table 3-2. Forging parameters for the material used in this study. The first line is step
1 and the second line is step 2....................... ..... ..................36

Table 4-1. Testing parameters and results for samples used to determine effect of
location on L C F ..................................................................................................86

Table 4-2. Tensile test data. .................................................... ................................ 95

Table 4-3. Low cycle fatigue data of environmentally exposed samples........................96

Table 4-4. Coated LCF data......................................................................................106















LIST OF FIGURES


Figure page

Figure 1.1. Specific yield strength vs. temperature comparing an a2 and orthorhombic
Ti-Al-Nb alloy with a conventional Ti alloy and IN718. After [Woo93]...................2

Figure 2.1. Titanium aluminum equilibrium phase. After [Mur87]................ ............

Figure 2.2. Comparison of the specific yield strength of a, and orthorhombic alloys
as a function of temperature. After [Row91]. ....................................................... 7

Figure 2.3. Ball model of basal plane of a showing lattice sites for Ti and Al atoms.
The a phase lattice is similar, with the Ti and Al atoms arranged randomly..............8

Figure 2.4. 900C isotherm from the Ti-Al-Nb ternary phase diagram. B2, a,, and
orthorhombic phases are in equilibrium. After [Row92]....................................... 9

Figure 2.5. Ball model drawing of the B2 phase showing the relationship between the
titanium alum inum and niobium atom s. ............................. ............................... 10

Figure 2.6. Comparison of the basal planes of the a, and orthorhombic phases ..............11

Figure 2.7. Constant 25 at% Al isopleth from the Ti-Al-Nb ternary in Figure 2.4.
A after [B an95]. .......................................... ......................... ........... ...................12

Figure 2.8. SEM Backscatter micrograph ofTi-22A1-26Nb processed in the O + B2
field and directly aged at 8150C for 4 hours. After [Woo93]..................................13

Figure 2.9. SEM Backscatter micrograph of Ti-22A1-26Nb processed in the B2 field
and directly aged at 8150C for 4 hours. After [Woo93]........................................... 14

Figure 2.10. Crystal structure and burgers vectors for the a, and O phases. After
[B an 9 5 ]...................... .... ................................................ ....................... ............ 15

Figure 2.11. Slip bands in the B2 phase in Ti-24A1- 1Nb showing the inhomogeneous
nature of slip. After [Ban90]. ............ .......................................... .................... 17

Figure 2.12. Yield strength vs. cooling rate for two a, alloys. After [Gog90] ...............17

Figure 2.13. Elongation vs. cooling rate for two a2 alloys. After [Gog90]......................18









Figure 2.14. SEM backscatter micrograph of Ti-22A1-26Nb forged in the B2 region
and directly aged at 8040C. Notice the colony of O laths that have formed off of
th e large lath ..................................................... .................................................. 19

Figure 2.15. Intergranular crack observed in tensile tested Ti-22A1-27Nb solutioned
above the P(B2) transus and aged at 8700C for 50 hours.........................................20

Figure 2.16. Fractograph of forged + direct aged material. Notice the high GAR and
the initiation of failure from an internal boundary.........................................21

Figure 2.17. SEM fractograph of material solutioned above the P(B2) transus and aged
at 870'c for 50 hours. Notice that there is mixed intergranular and transgranular
fracture across the sample and the prior P grain size is approximately 500m..........22

Figure 3.1. BSE image of a typical microstructure of the Ti-22A1-26Nb material after
B2 forging and heat treatment at 8150C for 4 hours. .......................... ........... 37

Figure 3.2. Photomacrograph of one half of the forging showing flow lines....................38

Figure 3.3. Forging half showing location of 10x macrophotographs and lines for
sectioning into metallographic samples. .......................... .........................39

Figure 3.4. Calculated iso-strain lines for the forging examined in this study ................39

Figure 3.5. Photomicrograph taken from area C. Notice the equiaxed grain structure
and low grain aspect ratio. ................................. ....... .. ...........................40

Figure 3.6. Photomicrograph taken from area P. Notice the high grain aspect ratio
resulting from the higher effective forging strain.......................................40

Figure 3.7. High resolution SEM image of area S (identified in Figure 3.3) from the
slowest cooled region in the forging...................... ............................41

Figure 3.8. High resolution SEM image of area Q (identified in Figure 3.3) from the
fastest cooled region in the forging................................... .... .................... ... 42

Figure 3.9. Fatigue and oxidation specimen sectioning diagram....................................43

Figure 3.10. Tensile and oxidation sample sectioning diagram. Shaded samples are
oxidation and white samples are tensile/creep...........................................44

Figure 3.11. Details of sectioning diagram for oxidation (shaded) and tensile (open)
samples from Figure 3.10. Calculated effective strain is also depicted....................44

Figure 3.12. Drawing of oxidation sample used in this research....................................45

Figure 3.13. Photograph of furnace used to conduct cyclic oxidation tests....................46









Figure 3.14. Oxidation pins loaded on hearth plate. The samples are the darker pins,
the thermocouples are the two white rods in the foreground. The cordierite block
is resting on the furnace platen. ........................ ..... ............................47

Figure 3.15. Typical thermal cycle recorded by strip chart recorder..............................48

Figure 3.16. Drawing of tensile sample used in this research. Dimensions are in
in ch es. .......................... .............................................................................. ....4 9

Figure 3.17. Typical stress-strain curve generated from tensile testing. Line tangent to
slope of elastic portion of curve used to determine elongation at failure is shown...50

Figure 3.18. Calculated iso-strain lines from forging showing various effective strain
from forging and respective fatigue sample location.............................................51

Figure 3.19. Calculated cooling rates from 1080C to 8700C in "C/min and
corresponding fatigue specim en location........................................................... ... 51

Figure 3.20. Drawing of fatigue sample used in this research. The dimensions are in
inches ..................................... ......................................52

Figure 3.21. Photograph of LCF samples loaded for thermal cycling. Samples are
separated by the alum ina tubes. .................................................. .................... 53

Figure 3.22. Typical LCF data taken during testing showing the hysteresis loops at
various cycles. Curve for modulus determination is also shown............................55

Figure 3.23. Illustration of sectioning and metallographic mounting of the fatigue
sam ple for evaluation. ....................................... ..................... .. .............................55

Figure 3.24. Drawing of LPPS specimen fixture.................... .......................57

Figure 3.25. SEM backscatter micrograph of Al coated and reacted sample showing
the in-situ formed A13Ti coating on the right. ..................... .. ...................59

Figure 3.26. SEM backscatter micrograph of Si coated and reacted sample showing
in-situ formed TisSi3 coating on the right. ..................... ............................60

Figure 3.27. SEM backscatter micrograph of Pt coating showing the columnar
structure and porous nature of the coating. ...............................................61

Figure 3.28. Photograph of CVD reactor........................................ ......................62

Figure 3.29. Schematic diagram of CVD reactor shown in Figure 3.28.........................62









Figure 4.1. Baseline oxidation results showing no variation between three samples
from different locations run together, and a larger variation with a sample from a
different experimental run...................... ........ ......................66

Figure 4.2. Al and Si coated and reacted oxidation results showing the decrease in
oxidation rate over the baseline. .............................................. ...................... 67

Figure 4.3. Oxide coated oxidation results showing that TaO, and MgO coated
material behaves like uncoated material and SiO, coated samples begin to gain
w eight after 400 cycles. ........................................ ............. ........................68

Figure 4.4. Results showing the large reduction in oxidation rate with the application
of the MCrAIY coatings and the effect of Cr and Pt coatings.................................69

Figure 4.5. Comparison of the oxidation rate behavior of all coatings examined in this
study. ...... ........... ............................................... .............................................. 70

Figure 4.6. Uncoated baseline oxygen content and microhardness vs. depth...................71

Figure 4.7. SEM backscatter image of baseline microhardness sample cycled for 500
cy cles...................... ............................................................................................7 2

Figure 4.8. SEM backscatter image of baseline microhardness sample cycled for 1000
cy c les...................... .............................................................................................7 2

Figure 4.9. SEM backscatter image of the NiCrAlY + A1203 coated sample after 1000
cycles at 650 C ...................................................................... .............................74

Figure 4.10. Microhardness and oxygen vs. depth for the FeCrAlY coated and 1000
cycle exposed sam ple .................................................................. ......................75

Figure 4.11. Microhardness and iron vs. depth for the FeCrAlY coated and 1000 cycle
exposed sam ple. .................... ...............................................................................75

Figure 4.12. Microhardness and oxygen vs. depth for the NiCrAlY + A1203 coated
and 1000 cycle exposed sam ple. ............................................ .................... ....76

Figure 4.13. Microhardness and Ni vs. depth for the NiCrAIY + A1203 coated and
1000 cycle exposed sam ple.................................................... .......................76

Figure 4.14. Microhardness vs. oxygen content for the Al coated and reacted sample.
Note that the distance from the interface starts at -5tm..........................................77

Figure 4.15. SEM backscatter image of Al coated and reacted sample after 1000
cycles at 6500C. Notice the crack in the AITi coating .........................................78









Figure 4.16. Microhardness vs. oxygen content for the Si coated and reacted sample.
Notice that the distance from substrate starts at -5pm..............................................79

Figure 4.17. SEM backscatter image of the Si coated and reacted sample after 1000
cy c les..................... ..........................................................................................7 9

Figure 4.18. SEM backscatter micrograph of Pt sputter coated sample cycled for 1000
cycles............. ........ ...... ....... .... .......... .. .. .... .........................81

Figure 4.19. Microhardness and oxygen content vs. depth for the Pt coated and 1000
cycle exposed sam ple .................................................................. ...................81

Figure 4.20. Plot of Cr and oxygen content vs. depth in 1000 cycle exposed oxidation
sam p le ............................................... ...................................... ............................ 82

Figure 4.21. SEM backscatter image of the SiO, coated sample after 1000 cycles at
6 5 0 C .......................................................................................... .........................8 3

Figure 4.22. Microhardness vs. oxygen content for the SiO, coated sample....................84

Figure 4.23. SEM backscatter image of Ta2,O coated sample showing area analyzed
by Auger spectroscopy. Box outlines area where compositional maps were
obtained ........................................................ ........................... ............................84

Figure 4.24. Auger compositional maps showing the concentration ofTa, Ti and
oxygen in the boxed region shown in Figure 4.23.................... ....... ............ 86

Figure 4.25. Number of cycles to failure vs. strain for various locations in the forging.
The sample number is shown beneath each symbol, and the failure location is
listed above each symbol .................... ........ ....................87

Figure 4.26. Ultimate tensile strength, yield strength and percent elongation (plastic)
vs. calculated effective strain from forging.......................... ............................88

Figure 4.27. LCF data showing the calculated a, and a, values, along with failure
location and sample number for each data point...................... ...................90

Figure 4.28. Optical fractograph of specimen taken from a highly stressed region in
the forging (N o. 52). ...................................................... ................................ 92

Figure 4.29. Secondary SEM photograph of fracture initiation site from specimen #52
show n in Figure 4.28....................... .... ...................................................... 93

Figure 4.30. Secondary SEM photograph of the fracture initiation site shown in the
small white box in Figure 4.29. ......................................... 94









Figure 4.31. Optical fractograph of sample number 6 taken from a low strain region in
the forging ................ .............. .............................................. 94

Figure 4.32. Ultimate tensile strength, yield strength and percent elongation of exposed
and unexposed samples tested in air and vacuum at room temperature and 5400C...96

Figure 4.33. Plot of number of cycles to failure vs. strain. Failure location is noted
w ith each sam ple ....................................................................... ..................... 97

Figure 4.34. Number of cycles to failure vs. strain for the unexposed and exposed
samples. This construction uses logarithmic curve fits to describe the LCF
behavior of the exposed samples. ............................................. ..................... 99

Figure 4.35. SEM backscatter image of a sample exposed for 100 cycles at 650C.
The dark region at the top of the photo is an area of high oxygen concentration.
The fracture surface is oriented to the right in this photo.....................................100

Figure 4.36. Optical micrograph of cross section from exposed tensile sample
stretched to 0.7% strain. Notice cracks run the depth of the oxygen rich zone......101

Figure 4.37. Optical micrograph of cross section from exposed tensile sample stretched
to 1.2% strain. Note the crack runs beyond the oxygen rich region .....................102

Figure 4.38. Uncoated, unexposed and exposed data, showing the failure location and
the calculated a ...................................... .................................................... 104

Figure 4.39. Oxygen concentration vs. depth from surface for the exposed fatigue
sam p les................................ ...............................................................................10 5

Figure 4.40. SiO, coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material. ..................... ......................107

Figure 4.41. TaO5 coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material. ..................... ....................108

Figure 4.42. MgO coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material. ................... ......................109

Figure 4.43. Cr coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material. ............................................110

Figure 4.44. Depth of oxygen concentration in SiO, coated LCF samples exposed in
air and argon........................ .... ............................................................... 11

Figure 4.45. Depth of oxygen concentration in TaZO, coated LCF samples after
exposure in air and argon ............................................................ ..................... 112









Figure 4.46. Depth of oxygen concentration in Cr coated LCF samples exposed to air
an d argo n ....................................................................... ... .................................. 1 13

Figure 4.47. Cr concentration depth in the Cr coated LCF samples exposed in air and
arg o n .................................................................. ............. ................................. 1 14

Figure 4.48. SEM backscatter image of Cr coated LCF sample exposed in air for 100
cycles at 6 50 C .................................................... ................... ....................... 14

Figure 4.49. Coefficient of thermal expansion comparison between coatings and Ti-
22A1-26Nb substrate........................... ....... ....................... 115















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

EFFECT OF COATING AND EXPOSURE ON THE OXIDATION AND
MECHANICAL PROPERTIES OF Ti-22A1-26Nb

By

James Ross Dobbs

December 1997
Chairman: Michael J. Kaufman
Major Department: Materials Science and Engineering

High temperature titanium base alloys are being developed to increase the thrust-

to-weight ratio in aircraft turbine engines. Conventional titanium alloys, such as Ti-6AI-

2Sn-4Zr-2Mo (wt%), can only be used to 5400C. Advanced titanium aluminide alloys

based on the a2 phase, such as Ti-24Al-11Nb (at%), and the orthorhombic phase, such as

Ti-22A1-26Nb (at%), have better high temperature tensile and creep properties than the

conventional alloys but have lower ductility. In addition to having low intrinsic ductility,

the alloys are embrittled by high temperature exposure in air.

In this study, the cyclic oxidation rate of Ti-22A1-26Nb (at%) at 6500C was

determined, and several classes of coatings, including NiCrAIY+A12O,, FeCrAIY, Pt, Cr,

Al3Ti, TisSi3, SiO,, Ta2O, and MgO, were evaluated to determine their effect on the

oxidation rate. Baseline low cycle fatigue life was established, and the effect of exposure

in air at 6500C on LCF life was determined. Based on oxidation and microhardness,









coatings ofCr, SiO,, Ta2O, and MgO were selected for a detailed study of the effect of

exposure in air at 6500C on the low cycle fatigue life of coated and uncoated material.

It is shown that exposure in air at 650C reduced the LCF life of Ti-22AI-26Nb by

as much as 2 orders of magnitude for a 100 cycle exposure. A single cyclic exposure

reduces LCF by 1.5 orders of magnitude. No coating evaluated reduced this degradation,

and in the case of SiO2 and MgO, the LCF life was worse after exposure than when no

coating was applied. It is shown that the observed reduction in LCF life is associated

with an increase in oxygen content and microhardness measured in the near surface

regions of the oxidation samples after thermal cycling. The oxide coatings were reduced

by the substrate either during coating or during exposure. The Cr coating allowed oxygen

to diffuse to the substrate.

These results indicate that further work should study those coatings that would not

be chemically reduced by titanium and that would not allow oxygen to diffuse through

the coating to the substrate.














CHAPTER 1
INTRODUCTION

Titanium base materials are used in a wide variety of applications. The medical

and chemical industries use titanium because of its resistance to corrosion and its

biocompatability, the sporting goods industry uses it because it is light weight and stiff,

and the aerospace industry uses titanium alloys for their high strength-to-weight ratio at

elevated temperatures. The use of titanium alloys in gas turbine engines is the most

demanding since the material is subjected to a corrosive environment at elevated

temperatures and, in some cases, high stresses that are frequently cyclic in nature. Thus,

titanium base alloys are attractive for use in gas turbine engines because of their low

density compared to Ni-base alloys (4.5 gm/cm3 vs 8.9 gm/cm3) and their high strength at

elevated temperatures [Pos92].

Titanium alloyed with aluminum and vanadium, Ti-6wt%Al-4wt%V, has been an

industry standard for many years [Woo72]. Below 1000C, this alloy consists of a strong

hcp a-phase precipitated in a softer bcc p-phase matrix, producing a stable two-phase

microstructure of hcp a-phase and bcc P-phase [Woo72]. The P phase is stabilized by the

addition of the vanadium whereas the addition of aluminum stabilizes the a phase. In

binary Ti-Al, single phase a is stable up to aluminum levels of approximately 11 atomic

percent. At greater aluminum levels, the disordered a phase is in equilibrium with the

ordered hcp TiAI (a2) phase [Mur87]. Single phase a, is stable from 22 to 35 atomic









percent aluminum at 5000C. The substitution of niobium (also a 3 stabilizer) for titanium

in the a, Ti3Al alloys leads to a better balance of properties, i.e. improved creep

resistance, fracture toughness and oxidation behavior. Initial work on the Ti-Al-Nb

system was carried out by McAndrew and Simcoe [Ml. ,., I] who studied alloys with up

to 20wt% Al and 30wt% Nb. More recent work conducted by Blackburn [Bla78]

indicated the promise of the TiAl class of materials as engineering alloys. The specific

tensile yield strengths of the TiAI + Nb alloys as a function of temperature are superior

to that of conventional titanium alloys, and even better than cast and wrought IN718

[Pos92] as shown in Figure 1.1. Several engine components have been produced from

Ti3A1 alloys, including afterburner nozzle seals and high pressure compressor casings for

the F100 engine and exhaust seals for the F404 engine [Lip85]

900


S800


r 700 Orthorhombic Alloy


600II 718
1 \Alloy 718

.S 500 Ti-6242


400
0 200 400 600 800 1000 1200 1400
Temperature (F)

Figure 1.1. Specific yield strength vs. temperature comparing an a( and orthorhombic Ti-
Al-Nb alloy with a conventional Ti alloy and IN718. After [W...','3 I









The discovery by Banerjee et al. [Ban88] and Rowe [Row90] that higher additions

of niobium (above approximately 12 at%) to Ti3Al leads to the formation of an ordered

Ti2AlNb orthorhombic phase (O Phase) introduced another class of materials for high

temperature applications. Alloys based on this O phase tend to have higher toughness

and tensile strength than alloys based on the a2 phase [Row91, Row91a, Row93, Smi93b,

Smi94] while maintaining similar creep resistance. Extensive evaluation of this class of

materials has been conducted in recent years [Ban95]. The deformation behavior and

dislocation processes have been evaluated [Ban91, Ban92, Ban95a, Dar94, Dou93,

Pop96a, Pop96b], as well as phase transformations and relationships [Ban93, Kau88,

Mur95, Vas96, Wey89].

A major impediment to using titanium-base materials at elevated temperatures in

oxidizing environments is the propensity for oxygen to diffuse into the titanium [Unn86].

The oxygen, being a small atom, sits in the titanium lattice interstices and causes the

titanium alloy to become very hard and thus embrittled [Cha87, Unn86]. In the case of

oa-p alloys, oxygen is a strong a stabilizing element and leads to a hard surface layer

known as "a-case". In the a2 and O alloys, there is a high solubility for oxygen in the

lattice, but these a like phases are not as readily stabilized by oxygen. Therefore, there is

no discernible change in the microstructure due to oxygen ingression during exposure.

Even so, high oxygen surface regions are embrittled and lead to a reduction in ductility,

fracture toughness and fatigue resistance [Dar95, Dar96, Lip93, Sai93, Sch95]. These

reductions in properties have been a major barrier to the implementation of this class of

materials in gas turbine engines [Pos92].









Several attempts to mitigate the environmental embrittlement have been made

with limited success. For example, 02 and O alloys have been coated with a variety of

materials, including MCrAlYs [Bri92, McK93, Sch95], glasses [Dev90, Wie89, Wie91],

silicides [Coc96], aluminides [Kun90, Smi90, Smi93b, Sub88] and Pt [Nic96]. These

coatings resulted in embrittlement, presumably due to either the formation of deleterious

reaction phases, or the coating itself was brittle and caused a reduction in low cycle

fatigue (LCF). The process of applying the coating could also contribute to the reduction

in LCF life. Subsequent exposure of these coated alloys to a high temperature oxidizing

environment typically led to negligible further embrittlement over the as-coated condition

presumably because the initial deterioration had already occurred during coating.

The objectives of this research include identifying a potential coating system and

methodology that does not lead to degradation upon application and is protective during

subsequent exposure, and to identify the mechanisms that can cause changes in the

mechanical properties of Ti-22A1-26Nb due to the presence of specific protective

coatings.















CHAPTER 2
LITERATURE REVIEW

Overview of Ti-Al-Nb Metallurgy


History

The development of supersonic jet aircraft after World War II nurtured the need

for high thrust-to-weight ratio aircraft turbine engines. This initiated the search for new

structural alloys that could withstand the high temperatures and stresses induced by a

turbine engine [Jaf80]. In aerospace applications conventional aluminum alloys are only

useable up to approximately 150C. Conventional titanium alloys like Ti-6A1-4V (wt%)

can be used at temperatures exceeding 300C. Today's commercial alloys have the

strength capability to operate at temperatures approaching 600C. Their oxidation

resistance is not yet on the same level as the mechanical properties and, consequently,

coatings have been developed that mitigate the poor oxidation behavior [Eyl84]. In the

early stages of titanium alloy development, it was recognized that the addition of

aluminum would increase the tensile and creep strength as well as the elastic modulus

[Ely84]. Much work was conducted in subsequent years to develop an accurate Ti-Al

phase diagram and in 1987 Murray [Mur87] published a version of the diagram based on

much of this work. This diagram is shown in Figure 2.1 and includes all of the phases

that form with the increasing aluminum.




















1200


noD- 0T / f N







Ti Atomic Percent Alum num A]
Ti5Aii


.0 (aoTi) I





0Atomi2 P-r"2nt Aluminum Al

Figure 2.1. Titanium aluminum equilibrium phase. After [Mur87].



Based on the phase diagram and empirical studies, it was determined that the

amount of aluminum, or its equivalent, should be kept below 9 at% to remain in the a

phase region. Higher levels of aluminum led to the formation of the intermetallic Ti3Al

and a reduction in ductility. In the 1970s it was recognized that the TiAl and TiAl class

of materials could possibly be exploited as structural engineering alloys. Ti3Al, also

known as a,, was developed by the Air Force [Lip81] and at the aircraft engine

manufacturers. It was recognized that the addition of Nb to Ti3Al improved the ductility,

creep behavior and oxidation resistance [Bla78, Sas77], and a new class of TiAl alloys

was launched. At the lead of these materials was the ternary alloy Ti-24Al-I 1Nb (at%).

The addition of vanadium and molybdenum to a, improved properties even more, and

super a(, or Ti-25Al-10Nb-3V-1Mo, was developed. The addition of Nb at levels above









about 15 at% causes the Nb atoms order on the titanium sublattice leading to a distortion

of the ordered hexagonal cell into the orthorhombic phase [Ban88]. The orthorhombic

alloys are reported to have superior tensile strengths over the a2 alloys (Figure 2.2), better

oxidation resistance and comparable creep strengths. [Row90, Row91, Row91a, Row92,

Row93, Smi93, Smi94]. It is this phase that is the focus of this study.


3.0


2.6-


2.1
U
a 1.7
E

. 1.3-


S 0.9-


0.4-


0.0I I I I I
0 100 200 300 400 500 600 700 800
Temp (*C)

Figure 2.2. Comparison of the specific yield strength of ax and orthorhombic alloys as a
function of temperature. After [Row91].


Description of Phases Present in Ti-Al-Nb

As seen in Figure 2.1, there are five distinct phases present in the Ti-Al system.

Four of these P, a, TiAl1 or a2, and TiAI or y are interesting engineering alloys. AlTi has

shown promise as a coating for titanium alloys and will be discussed in a later section.

The 0 phase is a body centered cubic (bcc) structure and is stable at higher temperatures.


Ti-22AI-27Nb
-*-Ti-25AI-1 ONb-3V-1 Mo
--Ti-24AI-17Nb-1Mo




--..\








On cooling through the 1 transus, this phase undergoes a transformation to the a phase.

The a phase is a hexagonal close packed (hcp) phase consisting of a random arrangement

of titanium and aluminum atoms. As seen in Figure 2.1 the low temperature solubility of

aluminum in titanium is approximately 10 at%; beyond this amount a 2 phase mixture of

a and TiAl, and ordered structure known as ca, exist in equilibrium. Aluminum

additions over 22 at% will create single phase a2. The basal plane of the a, lattice is

depicted by a ball model in Figure 2.3.



Il Ti Al



Ti Ti Ti Ti



Al )l Ti ; I Al Ti -. a2



STi I Ti Ti Ti



Al Ti A


al -a3
Figure 2.3. Ball model of basal plane of a, showing lattice sites for Ti and Al atoms.
The a phase lattice is similar, with the Ti and Al atoms arranged randomly.









The addition of Nb to Ti3Al ast levels beyond 10 at% results in either a 2 phase

(a + B2 or O + B2) or 3 phase (a2 + O + B2) structure as will be discussed further

below. A partial 9000C Ti-Al-Nb isotherm (Figure 2.4) indicates that the maximum

solubility of Nb in a, is 15 at%, and above 10 at% Nb the a, is in equilibrium with B2 up

to 15 at% where the a2 and B2 are in equilibrium with the O phase. The B2 phase is an

ordered bcc phase formed from the p phase. The addition of Nb causes the p phase to

order with Ti atoms occupying one site and the Al and Nb atoms randomly occupying the

other [Ban87] as shown in Figure 2.5.

Ti

















T60Tib (6 FT--
0




B2+O+u,

Compositons o f
Interest




Ti-60Nb 9000C (16500F) Ti-60AI

Ti-Al-Nb Ternary

Figure 2.4. 900C isotherm from the Ti-Al-Nb ternary phase diagram. B2, a2, and
orthorhombic phases are in equilibrium. After [Row92].









As seen in Figure 2.4, there are three important phases of interest in the Ti-Al-Nb

system at 900C, namely, ac, B2 and 0. A ball model of the basal planes of the a2 and O

phases is shown in Figure 2.6. The orthorhombic unit cell is derived by a slight distortion

of the a2 unit cell caused by the ordering of the Nb atoms on the Ti sites [Kon86]. For

complete substitution, it is clear that the O phase is an ordered ternary phase based on the

composition TiAINb.


Figure 2.5. Ball model drawing of the B2 phase showing the relationship between the
titanium, aluminum and niobium atoms.









a3 [TT20] A1I OTi *Nb

..O ) "-o-*--a[100]



[T210] O



al [210] [010]

a2 O

Figure 2.6. Comparison of the basal planes of the a, and orthorhombic phases.


Szaruga et al. [Sza92] reported that oxygen is a strong a, stabilizer in Ti-25A1-

10Nb-3V-1Mo having an effect on both the p transus temperature and the p/B2 order-

disorder transition. Oxygen is also an important alloying element in the O phase. An

alloy of composition Ti-22A1-23Nb containing above 1000 wppm oxygen is made up of

the a2, O and B2 phases. Reducing the oxygen content below 1000 wppm produces only

two phases, O and B2 [Lip93]. This indicates that oxygen stabilizes the a, phase, much

as oxygen stabilizes the a phase in conventional titanium alloys. Rhodes et al. [Rho93]

reported the same effect in the alloy Ti-22A1-27Nb, but at oxygen concentrations of

approximately 1120 wppm the equilibrium phases were still O and B2. Therefore, at

higher Nb contents it takes more oxygen to stabilize the a, phase. Ward [War93]

concluded that interstitial atoms such as oxygen can be potent strengtheners, but have an

adverse effect on ductility by limiting the number of slip systems available.










A cut through the ternary phase diagram shown in Figure 2.4 along the constant

25 at% Al line is shown in Figure 2.7. The P to B2 ordering line is shown above 1100C

for compositions around 12 at% Nb.


B2 to ri/O
Widanstatten




82 to o2
Massive


B2toO
Massive



Tempered
martensite





B2 to (type


Nb (at %)

Figure 2.7. Constant 25 at% Al isopleth from the Ti-Al-Nb ternary in Figure 2.4. After
[Ban95].



Also shown in Figure 2.7 is the range of transformations that Ti3AI + Nb alloys

experience. The transformations associated with continuous cooling of high Nb alloys

(>16 at%) from the a, + B2 region is complex. The precipitation of a2 or O laths follows


the Burgers relationship (0001)a2, (110)1; [112 0]a2 [I1 l]p and the equivalent









relationship for the O phase (001)0 1 (110)1, [110]0 1 [111]3 [Ben91]. The

transformation from ac to O phase results in a lamellar or mosaic structure [Ban88].

Microstructures obtained during thermomechanical processing of alloys with high Nb

contents depend on the processing temperature. Processing in the B2 + a, or O phase

field will yield, on cooling, a mix of primary a, or O phase surrounded by a

Widmanstatten a2 or O + B2 mixture. Processing in the B2 field results in a 100%

Widmanstatten a2 or O laths with B2 in the interlath regions. Examples of these two

microstructures are shown for the alloy used in this study, Ti-22A1-26Nb in Figure 2.8

and Figure 2.9 respectively.




















Figure 2.8. SEM Backscatter micrograph of Ti-22A1-26Nb processed in the O + B2 field
and directly aged at 815C for 4 hours. After [Woo93].




























Figure 2.9. SEM Backscatter micrograph of Ti-22A1-26Nb processed in the B2 field and
directly aged at 8150C for 4 hours. After [Woo93].


Deformation Behavior of a,, O and B2 Phases

Dislocation arrangements in deformed a, alloys [Akk91, Ker84, Kos90, Lip 80,

Lip85] and O alloys [Ban91, Ban92, Ban95a, Dou93, Pop96b, Pop96a] have been studied

in depth. This includes deformation in the B2 phase contained in both materials.

The a, phase deforms by slip [Ban90] in two distinct systems, (1120)(0001)

and (1120){1010}, and by slip in the (1126){1121} system [Kos90]. The

crystal structure and Burgers vectors for the a2 phase are shown in Figure 2.10 along with

those for the O phase [Ban95]. slip has been observed in ac alloys but are associated

with the P to a, transformation rather than tensile deformation [Ban95]. The lack of

slip and the dependence of a, alloys on
and slip create plastic

incompatibility and thus the a, alloys are considered brittle at room temperature.









The O structure is similar to the a2 phase with a slight distortion due to ternary

ordering. Figure 2.10 compares the Burgers vectors for the O and ao phases, and it can

be seen that the three 1/6(1120) vectors in a, are not equal in the O phase in either

magnitude or generation ofAPBs. Banerjee et al. [Ban91] have shown that the

deformation associated with
slip in O is very similar to that in a,, but unlike a, there

is considerable slip in the O phase alloys and these dislocations are arranged into

well defined slip bands [Kos90]. This would suggest that the O phase alloys should

possess more plasticity at room temperature and behave better under fatigue crack growth

conditions.

1/6 [1126 Al 12102]
0 T
0 Nb
S 1/4[114]




1/6 [1120] /2[100

o 0 1/4[110]


Figure 2.10. Crystal structure and burgers vectors for the a, and O phases. After
[Ban95]


Deformation in the B2 phase occurs predominantly by (111) slip on the {110},

{121} and {123}planes [Ban90]. This deformation is extremely inhomogeneous and

localized into slip bands as shown in Figure 2.11. This localization could result in the

formation of persistent slip bands and contribute to the formation of crack nucleation sites

during fatigue [Mil97].



























a FS \ pm

Figure 2.11. Slip bands in the B2 phase in Ti-24Al-11Nb showing the inhomogeneous
nature of slip. After [Ban90].


Tensile Behavior of a, and O Alloys

The discovery that the TiAl+Nb alloys have higher tensile ductility than the

binary Ti3Al alloys [Sas77] indicate that Ti3Al base alloys, including the O phase alloys,

have to be alloyed sufficiently to stabilize the P or B2 phase with Nb as the preferential 3

stabilizing element. The B2 phase delays the cleavage cracking of the a, or O laths to

higher strains by its ability to accommodate the plastic instabilities associated with the

lack of slip systems in a,. This typically occurs at the a,/B2 interfaces with its large

number of available slip systems [Ban95].

Recall that the decomposition of the B2 phase into the a( or O phase when heat

treated above the p(B2) transus results in a lath type microstructure as shown in Figure









2.9 where the a2 or O laths have the Burgers relationship with respect to the B2 phase.

The cooling rate from the p(B2) transus strongly affects the strength, as shown in Figure

2.12, and the ductility, as shown in Figure 2.13. The yield strength increases continually

with increasing cooling rate while the ductility goes through a maximum.


1400.0 --


1200.0



-1000.0



S800.0
800.0


D 24-15
0 24-11


600.0 --



400.0 I
0.01 0.1 1 10
Cooling Rate (C/sec)

Figure 2.12. Yield strength vs. cooling rate for two a, alloys. After [Gog90].










U
O 24-15 El
O 24-11 El
4.0



3.0 -



2.0



1.0 O



0.0 -E- T .
0.01 0.1 1 10 100
Cooling Rate (C/sec)

Figure 2.13. Elongation vs. cooling rate for two a, alloys. After [Gog90].


The increasing cooling rate causes a finer size lath to be formed on cooling from

the P(B2) phase field. This finer lath arrangement yields a higher yield strength due to

the Hall-Petch relationship. The low ductility at low cooling rates results from the

formation of similarly oriented lath colonies which allow cleavage to occur across many

laths with no change in crack energy [Ban90]. The reduction in ductility at high cooling

rates is associated with the reduction in scale of the microstructure and the reduction in

volume fraction of the P phase. At high cooling rates there are also a, and O laths that

have nucleated from the grain boundaries. These grain boundary nucleated laths have a

similar orientation and thus a crack that forms in one can easily propagate into the others









with little energy loss. Therefore, the optimum lath structure is a fine basketweave of the

ca2 or O phase without any grain boundary film or grain boundary initiated lath colonies.

An example of a colony of O laths is seen in Figure 2.14 where the lath colony in the

lower right hand comer has formed sympathetically from the larger O lath running from

bottom right to top left. This is a micrograph of the Ti-22AI-26Nb alloy used in this

study.


Figure 2.14. SEM backscatter micrograph of Ti-22A1-26Nb forged in the B2 region and
directly aged at 804C. Notice the colony of O laths that have formed off of the large
lath.









The thermomechanical processing history of the material can also play a role in

the macroscopic tensile behavior. Material that has been forged and directly aged has

higher tensile strength than material that has been forged, solutioned above the p(B2)

transus and then aged [Dob94]. This is again due to the failure of the a, + P at the ca

film formed along the grain boundary interface due to strain incompatibility [C I',i.I as

seen in Figure 2.15.


Figure 2.15. Intergranular crack observed in tensile tested Ti-22A1-27Nb solutioned
above the P(B2) transus and aged at 8700C for 50 hours.


The high grain aspect ratio (GAR) left over from forging in the P forged + direct

aged material forces this failure to occur on boundaries that are parallel to the stress axis

as seen in Figure 2.16, as opposed to a few boundaries that are perpendicular to the stress

axis, as seen in Figure 2.17. The 0 forged and directly aged material also did not contain









the grain boundary a, film, and the lath size was much finer [Dob94]. This is indicative

of faster cooling rates from the P region.


Figure 2.16. Fractograph of forged + direct aged material. Notice the high GAR and the
initiation of failure from an internal boundary.









Figure 2.17. SEM fractograph of material solutioned above the P(B2) transus and aged at
870'c for 50 hours. Notice that there is mixed intergranular and transgranular fracture
across the sample and the prior p grain size is approximately 500Pm.


Fracture Toughness and Fatigue

Chan determined [Cha92] that the crack growth process in a coarse basketweave

structure of Ti-24AI-11Nb (at%), one where the laths are approximately 20 jtm long and

5 tm wide, was by decohesion of slip bands, similar to the mechanism for equiaxed a,.

The tips of the microcracks terminated in the P phase, as did the tip of the main crack. In

contrast, for a fine basketweave structure where the lath size is approximately half that of

the coarse laths and where the P phase is not as continuous, the cracks propagated around

and through the p phase. Therefore, Chan concluded that the room temperature fracture

toughness of Ti-24Al- 11Nb (at%) was imparted by the P phase inhibiting microcrack

nucleation by both relaxing the incompatibility strain in the a//p interfaces and by

blunting the crack tips.

Penton et al. reported that Super a, (Ti-25A-10-Nb-3V-lMo (at%)) containing

primary ax exhibits faster crack growth rates than p solutioned and aged material

[Pen93]. This effect appears to result from early cleavage of the brittle primary a,

[Pen92, Pen93]. Ward reported that the a2 laths appeared to cleave on the basal plane

[War93b] and that the crack is then bridged by the p phase [Gog90, Luk90]. Others have

reported that the cracks also propagate along the a,/p interfaces in transformed regions

[Luk90, Tak96]. Davidson et al. [Dav91] found that, for Super a, rolled below the p

transus and then aged, small cracks were always found in the a2 phase after deformation.

These small cracks would then grow below the AK,,, for large cracks. Ravichandran and









Larsen [Rav92] concluded that growth rates of small cracks in basketweave structures of

Ti-24Al-11Nb were consistent with large crack growth rate data. Penton et al. [Pen93]

concluded that the growth rate of solutioned and aged material is in line with the fracture

of a, laths.

Miller [Mil97] has defined three distinct conditions for initiation of fatigue:

1) Fatigue limit ofpolycrystalline material which is related to some limiting

microstructural feature such as grain size, primary precipitate size or brittle

precipitate size.

2) The mechanical stress state applied to a pre-existing crack or flaw.

3) The condition for a single crystal in which a defect, such as intrusion or

extrusion, must be induced.

Miller concludes that the fatigue crack is initiated on the first cycle by one of the

above conditions. In the case of a and O alloys with limited slip systems in the lath

structures, the first condition of some limiting microstructural feature would apply. This

feature would most likely be a primary a,/ O particle or an az/ O lath, in agreement with

the observations described above.

Fatigue is typically described as consisting of the following stages: [Sur91]

1) Nucleation of permanent damage through microstructural changes;

2) Creation of microscopic cracks;

3) Coalescence of microscopic cracks to form measurable cracks;

4) Stable growth of a preferred crack; and

5) Ultimate catastrophic failure.









In classic fracture mechanics, Stage 3 indicates the end of initiation and the

beginning of propagation of the fatigue crack. Considering this model of fatigue, linear

elastic fracture mechanics (LEFM) can be utilized to help understand the behavior of

material by calculating both the initial flaw size, a, and the size of the crack at failure, af.

The equation for the stress intensity, K,, at the notch of a sharp crack is

1.12 o'ra
K, = (2.1)


3z r a 2 7r
where + In the special case when c = a, the ( term is reduced to the
8 8 c 2

equation then becomes

K, = 0.71 acr- (2.2)

This is the case for a penny shaped flaw that is very small relative to the size of the test

bar. To determine the size of the fatigue crack at failure, af, one can use the plane strain

fracture toughness, Kic for K,, and in the case where A=l, the maximum stress is equal to

Ac. Solving equation 5.2 for a, gives


a,= ( )'2 (2.3)
0.71an_ I

da
Paris showed that the fatigue crack growth increment is related to the stress intensity
dN

factor as

da
N= C(AK)" (2.4)
dN

da
where C and m are scaling constants derived from the vs. AK curve. When equation
dN

2.4 is integrated from an assumed initial flaw size a, to the critical crack size a, the









number of cycles to failure can be calculated. When the stress intensity factor for a small

crack is defined as in equation 2.2, equation 2.4 becomes


d =C(0.71Ao-r, a)" (2.5)
dN

where a has now become Ao, defined as a,,,, a,,,,. For our case where we are assuming

that the stress intensity factor is not a function of the crack depth a, equation 2.5

integrates as


CY'(AcT),7 r ,' da
CY"(Ac)mt N "2 dN = J- (2.6)
0 .,a

The resulting fatigue life is


N,= CYAr' (Im -2 1 (2.7)
_a a, 2.

for n 2.

Comparison of the fatigue crack growth data for Ti-24Al-l 1Nb (at%) generated

by Ravichandran and Larsen [Rav92] to that for Ti-21Al-25Nb (at%) generated by

Woodfield et al. [Woo94] with a similar basketweave structure indicates that the Ti-21Al-

25Nb is slightly better. This could indicate that the O + B2 alloys would behave better in

fatigue, probably due to the increase in the number of active slip systems and the increase

in Nb content in the B2 phase. Balsone et al. [Bal93] evaluated fatigue crack growth in

Ti-25A1-25Nb (at%) with a microstructure consisting of primary a2 + O laths surrounded

by the B2 phase. Comparison of dA/dN vs. AK data for Ti-24Al- 11Nb (at%) with Ti-

25A1-25Nb with a similar microstructure [Rav92], indicates that the latter has a slightly









higher crack growth rate. This is probably due to the scale and volume fraction of the

primary ax or O and the effect of the ac2/P(B2) interface.

Overview of Oxidation Behavior


Oxidation of Uncoated Titanium Alloys

The oxidation behavior of titanium base alloys has been extensively studied since

the late 1940s, starting first with unalloyed titanium. Initially there were large

disagreements between investigators about rate laws and rates of oxidation [Kof58].

More recent investigations of commercially pure titanium by Unnam et al. have shown

that the weight gain due to oxide growth and oxygen dissolution is essentially parabolic

with respect to time [Unn86]. Unnam et al. also reported that the oxygen diffusion

coefficient in Ti(O) is independent of oxygen concentration in the 1-10 at% range, and

the effective solubility limit of oxygen in pure titanium is 20 at%. They also concluded

that the diffusion coefficient of oxygen in TiO, is about 50 times that of oxygen in the

metal. Chuanxi and Bingnan reported that the addition of Nb to pure Ti improved the

oxidation resistance by improving the surface stability of the oxides [Chu92]. Chaze and

Coddet [Cha87] conducted oxidation experiments on Ti with additions of Al, Si and Cr.

They concluded that the addition of Al up to 16.5 at% and Si up to 1.5 at% improved the

oxidation resistance of Ti, while Cr additions up to 18 at% had little or no effect. In

contrast, Kahveci and Welsch evaluated the effect of Al on oxidation ofTi [Kah87] and

concluded that at least 25 at% Al is necessary for any significant improvement in the

oxidation behavior of Ti-Al alloys and that there is no appreciable effect for the addition

of Al until about 13 at%, beyond which the oxidation behavior improves with increasing









Al content. Kahveci et al. [Kah88] showed that the kinetics for oxide growth for Ti-25A1

(at%) fall between TiO2 and AIO, but are closer to the kinetics for TiO,. Qiu et al

[Qui95] reported that the addition of 11 at% Nb or 5 at% Si can improve the oxidation

resistance of TiAl and that the effect of adding both is even greater. However, no

continuous A103 scale was formed after any additions.

In a subsequent paper, Welsch and Kahveci evaluated 3 binary Ti-Al alloys with

increasing Al content and 1 Ti-Al-Nb alloy[Wel89]. They observed that the parabolic

rate constant decreased and the thickness of the oxide scale decreased with increasing Al.

They concluded that the outer oxide scale in the binary alloys is a TiO, with A120,

channels, the intermediate layer consists of TiO2, A,103 and porosity, and the inner oxide

layer is A1203 with TiO, and porosity. The alloy surface has A1,03 lamella from internal

oxidation. On the other hand, the Ti-Al-Nb alloy had a dense outer oxide scale consisting

mostly of A1203 and TiO2, an intermediate layer of TiO, and Nb0z5; and the layer next to

the metal was an Nb2Os-TiO, layer with porosity at the interface. There was no visible

internal oxidation but there was an oxygen rich zone adjacent to the oxide scale.

Chromium has been found to help improve the ability to form a protective scale in

Ti-Al-Nb alloys and is an excellent P stabilizing element [Doy95]. The Cr tends to alloy

with the oxide scale and provide a more protective scale. This indicates that

improvements in oxidation are possible by alloying, and Cr is a good alloying addition.









Oxidation of Coated Titanium Alloys

Coatings for titanium base alloys have been evaluated based predominantly on

their ability to mitigate oxidation. The coatings for Ti alloys fall into four broad classes,

Al3Ti, MCrAlY, pure metals, and glasses or oxides.

AlTi

AI3Ti has been produced as a coating by a variety of methods including (1) laser

surface alloying [Abb93] where Al powder is either fed into the molten pool or placed on

the substrate as a slurry and melted in by laser [Gal92], (2) dipping the substrate of

interest into a molten bath of Al and allowing the reaction to take place [Abd91] (3)

deposition of Al by EBPVD and subsequently reacting with the matrix during exposure in

air [Unn85], (4) deposition by sputtering and reacting in vacuum prior to exposure in

a.i [\.- .', .1 and (5) pack cementation [Kun90, Smi90, Smi93a, Sub88] where the

substrate is placed in a mixture of alumina powder, aluminum powders and halide salt

activators and reacted at 1000C where the Al will volatile and produce a coating of Al on

the substrate, which is then reacted to form AI3Ti. In all cases the coating A13Ti was

cracked on cooling from the reaction temperature. Even so, the coatings tended to impart

improved oxidation resistance despite being cracked, although oxidation did occur down

the cracks and affect the substrate [Kun90]. It was concluded that the thicker the coating,

the more susceptible it was to cracking [Smi90]. McMordie reported that the addition of

Si to the Al3Ti coating improved oxidation even more, possibly due to the formation of a

silicide at the coating/substrate interface [McM91].









MCrAlY

MCrAIY type coatings, where M is a transition metal like Ni, Fe or Co, are

typical coatings used in the gas turbine industry. The application of these coatings to Ti

has been attempted by a number of researchers, and several patents are held on the

materials and processes [Bri92, Lut90, Tob92]. The preferred method of producing this

type of coating is plasma spray [McC90, McK93, McK93a, Sch95]. For Ti3Al + Nb

alloys, the MCrAlY is sometimes mixed with an oxide powder such as AIO3 in order to

reduce the thermal expansion of the coating to better match that of the substrate [Sch95].

In all cases, the MCrAIY coatings improve the oxidation resistance of Ti substrates, such

as Ti-6A1-2Sn-4Zr-2Mo (wt%) [McC90]. There is a slight increase in hardness beneath

the coating in the as-coated condition and this is attributed to oxygen ingress during the

pre-heating in partial vacuum prior to the plasma deposition.

Pure metals

The evaluation of pure metals as oxidation resistant coatings for Ti base materials

has mostly been limited to Pt and Cr. The preferred method for deposition of Pt is ion

plating [Eyl84, Eyl85, Fuj79] and for Cr is PVD processes such as sputtering [McK90].

Eylon et al. showed that the Pt coatings reduced the surface oxidation rate of Ti-6Al-2Sn-

4Zr-2Mo (wt%) by three orders of magnitude[Eyl85] and did not degrade the high cycle

fatigue properties [Fuj80]. PtAl2, produced by multi-layer sputtering of Pt and Al and

subsequent reaction, has been evaluated as a coating for a, type alloys by Nicholls et al

[Nic96]. They reported that a 3um thick coating protected then a2 alloy MT754 (Ti-

23A1-9Nb-2Mo-0.9Si at%) from oxygen ingress at 700 and 800C, and extended the









creep life of the alloy by a factor of two under conditions of 350MPa at 6500C. They also

reported that after exposure in air for 100 hours at 7000C, there was no hardening under

the PtAl2 coating.

Glasses, silicates and oxides

Oxides such as SiO2, Y203, MgO, A1203 and ZrO, have been deposited on Ti-base

alloys by a variety of techniques. Wiedemann et al. [Wie89, Wie91] deposited SiO,,

A1203 and B,03 by a sol-gel process, and Y203, MgO, ZrO, and HfO by sputtering onto

Ti-24Al-11Nb (at%). They reported that most coatings had poor integrity and allowed

oxygen to diffuse through. Of all the coatings they tried, the MgO based coating made

from the sol-gel process behaved the best after 1 hour at 982C in air. They concluded

that coatings applied by sputtering were not effective for oxidation protection due to their

poor integrity. In contrast, Clark et al. [Cla88] deposited SiO2 by sputtering and reported

that there was a significant reduction in weight gain at 7000C after 25 hours in air.

DeVore and Osborne [Dev90] coated TiAI with borosilicates and aluminosilicates and

concluded that all coatings showed some improvement in oxidation, but the thin sol-gel

coatings were the worst while a borosilicate coating airbrushed onto the sample exhibited

the best behavior. Bedell et al. [Bed91] utilized ion beam assisted deposition to deposit a

double coating of silicon nitride and chromium on the conventional Ti alloy IMI829.

They concluded that the coating provided protection from oxidation at 7000C for 100

hours, but the SiN4 diffusion barrier is not stable with titanium. Cockram and Rapp

[Coc96] produced silicide coatings on a2 and O phase alloys by pack cementation and

reported excellent oxidation resistance up to 10000C.









Embrittlement of Titanium Base Alloys

As mentioned previously, titanium base alloys suffer from embrittlement when

exposed to air at high temperatures. This embrittlement is due to the ingress of oxygen

and the subsequent increase in hardness in the affected layer [Sha68]. The volume

fraction of the beta phase in these near surface regions in Ti-6A1-4V (wt%) is reduced at

all temperatures when exposed in air [Kah87]. The 3 transus temperature is raised by the

presence of oxygen in Ti-6A1-4V as described by the relationship T,[C] = 937 + 242.7 x

[O wt%] [Kah87]. The depth of oxygen penetration, and thus the depth of hardening due

to interstitial ingress, is a parabolic function [Kah86, She86] thus indicating that the

hardening is due to diffusion of oxygen into the substrate. Wallace conducted tensile

tests on Ti-1100 (Ti-6A1-2.75Sn-4Zr-0.4Mo-0.070-0.02Fe wt%) foils after exposure in

air at 6000C and 800C for up to 1000 hours. He concluded that the elongation was the

most sensitive parameter to minor oxidation, but tensile strength was also affected by

long exposures. He also observed a band of brittle fracture at the metal surface due to

oxygen diffusion into the metal. This band is high in oxygen and made up of the a phase

[Wal95].

Balsone studied the tensile properties of Ti-24A1-1 1Nb (at%) after exposure in air

for 10 and 100 hours [Bal89]. He reported that an annulus of embrittled material was

formed around the surface of an exposed sample that cracked during loading. These

cracks served as notches for premature failure during testing such that both the elongation

and UTS were reduced. Embrittlement of Ti-25Al-10Nb-3V-1Mo (at%) was also studied

by Saitoh and Mino [Sai93] after exposure at 7000C for 100 hours. They also reported









that the elongation and UTS was reduced, as did Meier and Pettit [Mei92] for a Nb

modified u2 alloy after a 24 hour exposure at 9000C. Other investigators have evaluated

the effects of embrittlement on fatigue behavior. Godavarti et al. [God92] reported that

the LCF life of Ti-25Al-10Nb-3V-1Mo (at%) was reduced by 2 orders of magnitude after

short exposure times at 7200C. They attributed this reduction to the formation of a brittle

layer at the surface of the test samples. They also reported that the same pre-exposure in

air had no effect on high cycle fatigue (HCF) life due to the low strains imposed during

HCF testing. Praida and Nicholas [Pra92] found that fatigue crack growth resistance in

Ti-24Al-l Nb (at%) was strongly effected by testing in air at 6500C. Hold time fatigue

crack growth rates were the most effected by elevated temperature testing, indicating an

effect of the environment on cracking. Balsone et al. [Bal93] determined that Ti-25AI-

25Nb exhibited time-dependent crack growth behavior when tested at elevated

temperature. They reported that at 6500C and 750C, there is a contribution from

environmentally assisted crack growth.

Dary et al. evaluated the effect of exposure in air and in vacuum on the tensile

properties of Ti-22A1-23Nb [Dar96, Dar94, Dar95]. They observed a decrease in UTS

and ductility in all samples exposed in air and, although there was no loss of mechanical

properties in the vacuum exposed samples, the external surfaces contained a large number

of small cracks. Although Rhodes et al. [Rho93] observed a, in Ti-22A1-23Nb at oxygen

contents greater than 0.098wt%, Dary and Pollock did not observe any a2 precipitates in

the high oxygen regions when exposed at the same temperature. They supposed that the

760C exposure was too low to cause au precipitation or that the B2 + O phases were









compositionally stable at this temperature. Chesnutt et al. [Che93] conducted tensile tests

on Ti-22A1-27Nb after either a thermal exposure (no environment) at 6500C or an

environmental exposure at 590C in air. The results indicate no by the thermal exposure

but significant embrittlement after the environmental exposure. No evaluation was

conducted on the environmentally exposed samples to establish the cause of

embrittlement.

Effects of Coatings on Mechanical Properties of Ti Alloys

Coatings to prevent the degradation of mechanical properties in Ti alloys have not

received as much attention as coatings to mitigate the poor oxidation behavior. As

described above, Ti base alloys are embrittled after exposure in air at elevated

temperatures. To prevent or at least reduce this effect, a series of evaluations have been

conducted to determine the best coatings and the effect of these coatings on the

mechanical properties in both the as-coated and exposed conditions. Fujishiro and Eylon

demonstrated that the HCF life of Ti-6A1-2Sn-4Zr-2Mo (wt%) was improved at 455C

with the application of a platinum coating [Fuj80]. For the cx alloy Ti-24.5A1-12.5Nb-

1.5Mo (at%), Schaeffer and McCarron evaluated NiCrAl + 40 vol% AlO0 deposited by

plasma spray [Sch95]. They reported that the LCF of as coated samples was

approximately the same as for the uncoated and exposed samples where the LCF life was

reduced by three orders of magnitude. Subsequent exposure of the coated sample did not

degrade the LCF life any further. McKee reported that for the same alloy, an underlayer

of chromium or tungsten prior to deposition of the NiCrAl would eliminate this

embrittlement [McK93]. He based this conclusion on microhardness profiles beneath the









coating, rather than on mechanical property data. Chesnutt et al. evaluated the LCF

response of the O alloy Ti-22A1-27Nb (at%) with and without a coating [Che93] and

concluded that the LCF life was degraded by the coating in a manner similar to that seen

by Scaeffer and McCarron [Sch95].

Program Approach

While there have been various studies addressing the oxidation of coated O phase

alloys and the effect of coatings on mechanical properties, to date there has not been a

systematic study addressing both issues. Such an study will be critical to the use of this

class of alloys in high temperature applications in air, particularly in static applications

with large thermal gradients where the stresses developed by the thermal gradients can be

large and the number of thermal cycles small, therefore subjecting a static part to low

cycle fatigue.

The approach of this program is to evaluate the oxidation, tensile and LCF

behavior of O phase alloy Ti-22A1-26Nb in both the uncoated and coated conditions. The

material will be taken from a forged disk, and the baseline properties of the material will

be determined, including the effect of location in the forging. Coatings will include

oxides and pure metals, and are intended to be barriers for oxygen diffusion and diffusion

of any topcoat that may be identified.














CHAPTER 3
EXPERIMENTAL PROCEDURE

Material Processing


A 6100 pound ingot of target composition Ti-22A1-26Nb (atom percent) was

produced at Teledyne Allvac/Vasco. The material was produced by an initial plasma arc

melting (PAM) into a 43 cm (17 inch) diameter ingot, followed by vacuum arc remelting

(VAR) into a 53 cm (21 inch) diameter ingot. This 53 cm (21 inch) diameter ingot was

sectioned into two pieces. The top section was cogged down from 53 cm (21 inch) to 41

cm (16 inch) diameter billet by heating to 12320C (2250F), and using 2.5 cm (1 inch)

drafts. This was followed by hot grinding and cooling to room temperature. This top

half was converted to 20 cm (8 inch) diameter billet and sheet bar for use in other

programs. The bottom half was cogged down from 53 cm (21 inch) to 40 cm (16 inch)

billet using a similar approach as the top half, except there was a 40% reduction initially

and fewer reheats were used. This 40 cm (16 inch) diameter billet was then GFM

converted to a 15 cm (6 inch) diameter billet. This conversion took place at 1010C

(18500F), based on beta transus measurements of a slice from the 41 cm (16 inch)

diameter billet. Four billets approximately 2.4 m (94 inches) long and 16 cm (6.3 inches)

in diameter weighing a total of approximately 2000 pounds were produced. Chemistries

were measured from several locations along the length of the ingot and from the surface

and center of these locations. The compositions are listed in Table 3-1 [Woo84].









Table 3-1. Compositions from several locations along the length of the 15.25 cm
diameter billet, surface and center, and from the forging used in this study.
Billet Forging Aim
Al (at%) 21.35 21.29 22.0+1.0
Nb (at%) 25.75 26.15 26.0+1.0
Ti Balance Balance Balance
Fe (wt%) 0.036 0.06 <0.07
C (wt%) 0.013 0.0063 <0.05
H (wt%) 0.020 0.0014 <0.0125
O (wt%) 0.071 0.075 <0.08
N (wt%) 0.012 0.012 <0.05


A 30.5 cm long section from one of these billets was forged at Wyman-Gordon in

Houston. The two step forging consisted of an O + P forging followed by a P forging.

The forging parameters are listed in Table 3-2.


Table 3-2. Forging parameters for the material used in this study. The first line is step 1
and the second line is step 2.
Initial Final Reduction Billet Temp Die Temp Strain Cooling
Height Height ratio C (F) o C (F) Rate (offdies)
cm (in) cm (in) (/min)
30.48(12) 15.88 (6.25) 1.92/1 996 (1825) 870 (1600) 0.95 Air
15.88 (6.25) 6.35 (2.50) 2.5/1 1100(2010) 870 (1600) 0.92 Fan


After the billet was forged, it was removed from the forging press and allowed to

cool to room temperature. It was then given an aging heat treatment at 8150C (15000F)

for 4 hours followed by air cooling without being subjected to a solution heat treatment.

This procedure is termed direct aging. A typical microstructure is shown in Figure 3.1. A

diametrical slice was then taken across the center of the forging for evaluation of the

material flow. Unfortunately a carbide cut-off wheel without cooling was used to section

this slice and ,as a result, the slice and both cut faces of the forging were severely









cracked. This caused some concern as to the integrity of the forging, so scanning

acoustic microscopy (SAM) was conducted to evaluate the depth of the cracking.


Figure 3.1. BSE image of a typical microstructure of the Ti-22Al-26Nb material after B2
forging and heat treatment at 815C for 4 hours.


A 1.27 cm (0.5 inch) thick section was cut off of each forging face prior to the

SAM analysis. The SAM analysis revealed no cracking on the remaining faces indicating

that all of the cracks were confined to the near surface region. The 1.27 cm (0.5 inch)

thick section was cut in half and polished and macro etched to reveal the flow lines in the

forging. The macro flow lines are shown in Figure 3.2.
























Figure 3.2. Photomacrograph of one half of the forging showing flow lines.


After the forging was pronounced sound, an extensive analysis of the variation in

grain size and grain aspect ratio versus location in the forging was conducted. The

effective strain from forging was calculated and iso-strain lines, shown in Figure 3.4,

were constructed. The calculated iso-strain lines were superimposed onto the forging

macrograph, and the location of various forging strains were identified. Figure 3.3 shows

the half diametrical slice from the forging center marked for photomacrographs at 10x

magnification. After the photographs were taken, the diametrical slice was sectioned

along the line shown into metallographic samples for microstructural analysis. The

forging was assumed to be symmetrical about the centerline and thus was sectioned to

reveal the grain structure in one quadrant of the forging. The diametrical slice was

sectioned with a carbide cut-off wheel and the samples were mounted in bakelite and

polished and etched. Microphotographs from representative locations in each samples

were taken at 50x, 500x and 1000x magnification.























*S STEL



Figure 3.3. Forging half showing location of 1Ox macrophotographs and lines for
sectioning into metallographic samples.



3.00
Eff. Strain
S. A 0.131
.' : e~%"" -" -;~, N A= 0.300
2.00 B 0600
N\ C 0.800
.2- -- -.7. N D=1.200
a aN /G F IE D E= 1.500
1.00- -- / F 1.800
G=2.100
H= 2.400
I= 2.700
.00 I I- =2.957
.00 1.00 2.00 3.00 4.00 5,00 6.00 7.00

Radius

Figure 3.4. Calculated iso-strain lines for the forging examined in this study.



Representative micrographs from the dead zone in the forging, location C, and a

more highly strained area, location P, are shown in Figure 3.5 and Figure 3.6,

respectively.



























Figure 3.5. Photomicrograph taken from area C. Notice the equiaxed grain structure and
low grain aspect ratio.


2 L-I 0 .-,

Figure 3.6. Photomicrograph taken from area P. Notice the high grain aspect ratio
resulting from the higher effective forging strain.









Notice the equiaxed grain structure from area C and the high grain aspect ratio in

area P. These two areas encompass the range of microstructures (grain sizes and aspect

ratios) that resulted from the forging process.

The rate at which the forging cooled from the P transus temperature would effect

the size and distribution of the O phase lath. Thus, the cooling rate of the forging from

the p transus temperature through 870C was calculated. There was little variation in the

cooling rate across the forging, and thus little variation in the O phase lath size would be

expected. This was confirmed by performing high resolution SEM analyses, shown in

Figure 3.7 and Figure 3.8, and little or no variation in the orthorhombic lath was

observed.


.4 t'" l .- ,. u. -. 1 : i i l r : ... : r ,:. j L" .
Figure 3.7. High resolution SEM image of area S (identified in Figure 3.3) from the
slowest cooled region in the forging.





































Figure 3.8. High resolution SEM image of area Q (identified in Figure 3.3) from the
fastest cooled region in the forging.


Specimen Blanking


After determining that the forging was sound, a plan for sectioning the oxidation

samples from the near surface dead zone and the mechanical property samples from

inside the forging was devised. This was based on the premise that the microstructure

found in this dead zone should not effect oxidation. The sectioning plan along with the

two ingot sections were submitted to the wire electro-discharge machining (EDM) shop

for sectioning as shown schematically in Figure 3.9.



































0.250" 117


Figure 3.9. Fatigue and oxidation specimen sectioning diagram


The remaining oxidation samples and the tensile samples were sectioned from the

other forging half as shown in Figure 3.10 which also shows the scrap blocks, numbered

1 through 6, that were left after sectioning. The oxidation samples were cut from near the

surface and from a block in the middle of the forging, some in low strain and some in

high strain regions, as shown in Figure 3.10 surface and the tensile samples were taken

from the more highly worked material. The details of this sectioning is shown in Figure

3.11. The open circles again represent the 5.08 cm (2 inch) long tensile/creep sample and

the shaded circles represent the oxidation samples.

































876 B70 L BI 247 260
iA
BOTTOM

B121 B105 1.20 134 135 148
B131 163 1 167








Figure 3.10. Tensile and oxidation sample sectioning diagram. Shaded samples are
oxidation and white samples are tensile/creep.


Figure 3.11. Details of sectioning diagram for oxidation (shaded) and tensile (open)
samples from Figure 3.10. Calculated effective strain is also depicted









Test Specimens and Test Procedures


Oxidation

Samples

The oxidation samples were 8.9 cm (3.5 inches) long and 0.635 cm (0.25 inch) in

diameter with a 0.3175 cm (0.125 inch) radius machined on each end (Figure 3.12). The

total surface area of each sample was 17.72 cm2. The specimen blanks were EDMed

oversize from the forging, then centerless ground to the final diameter. The radius was

machined by turning on a lathe; this made complete coverage of the sample easier and

allowed for a simpler calculation of surface area as needed for the oxidation

measurements. More importantly, the rounded ends eliminated any sharp comers which

might act as stress risers for the coating leading to premature failure of the coating during

thermal cycling.

0.250| 3.250
\ 0.250



R.125


Figure 3.12. Drawing of oxidation sample used in this research


Testing

The oxidation testing was conducted in a bottom-loading thermal cycle furnace

built by Ted Kominsky Ovens and Furnaces, Model number KS16000B. A photograph

of the furnace is shown in Figure 3.13. The temperature was controlled with a Honeywell

UDC 3000 controller using input from a type K (NiCr-NiAl) thermocouple. This









thermocouple was calibrated against a NIST standard calibrated thermocouple. The

samples were loaded vertically into a sample fixture made from a cordierite open cell

foam to facilitate rapid cooling as shown in Figure 3.14. Note the thermocouples that

protrude through the hearth plate at the same height as the oxidation pins. One

thermocouple is attached to a the controller and isthe other is attached to a strip chart

recorder that records temperature and number of cycles.


Figure 3.13. Photograph of furnace used to conduct cyclic oxidation tests.









The samples are raised into the furnace and the furnace is automatically started.

The samples remain up in the furnace for 55 minutes, then are lowered and fan cooled for

5 minutes. A typical thermal cycle taken from the strip chart recorder is shown in Figure

3.15. As seen in Figure 3.15, it takes approximately 12 minutes to reach 6500C, and 5

minutes to cool down to approximately 100C. The samples were weighed at 1, 3, 5 and

10 cycles, then every 10 cycles to 100 cycles and then every 100 cycles to 1000 cycles

where the test was terminated and the samples analyzed. The samples were handled

wearing cotton gloves during weighing to prevent contamination from the examiner and

after each weighing the samples were flipped over to ensure that one end was not always

in contact with the fixture. The samples were weighed on a Mettler 200XT digital scale

with an accuracy of 0.1 mg.











Platen




Figure 3.14. Oxidation pins loaded on hearth plate. The samples are the darker pins, the
thermocouples are the two white rods in the foreground. The cordierite block is resting
on the furnace platen.


The samples were then sectioned transverse and longitudinally and mounted in

Conductomet for analysis. Electron microprobe analysis (EMPA) was conducted at









RPI by Dr. David Wark. Microhardness measurements were conducted on a LECO DM-

400FT at a O1gfload and a 20 second load time. The tester was calibrated prior to taking

each set of measurements using the LECO calibration blocks. The data presented is from

one data point at each location; therefore, there will be some experimental error due to

variations in the microstructure or operator error in reading the microhardness

indentation. The results from the microhardness were compared to the oxygen levels

detected by the EMPA, and correlations were made concerning the effectiveness of the

various coatings to prevent oxygen ingress. The EMPA results were also used to

correlate other diffusion phenomena, such as that of the Ni and Fe into the alloy substrate,

to the microhardness results.


700 I --



600


500--


400-


S300-


200-

100 -

0 10 20 30 40 50 60
Time (mins)

Figure 3.15. Typical thermal cycle recorded by strip chart recorder.









Tensile

Samples

The tensile samples are 5.08 cm (2 inches) long and have a gauge length of 1.9 cm

(0.75 inches). The gauge diameter is 0.40 cm (0.16 inches). The sample is gripped by /4-

20 threads with a thread root radius of 0.25 mm (0.01 inches). This sample is shown

schematically in Figure 3.16.

2.000
S1.000
.375




.250-20 threads
2 places

S.160
Figure 3.16. Drawing of tensile sample used in this research. Dimensions are in inches.


Testing

The tensile tests were conducted on an Instron Model 1125 screw driven test stand

using a 5,000 pound load cell. The cross-head speed was 0.05 cm/min (0.02 in/min) and

the sample were held in threaded grips. The strain was measured from the cross-head

displacement, and load vs. time was recorded electronically and on a strip chart recorder.

Percent strain was calculated from the cross-head displacement divided by the initial

gauge length multiplied by 100. The stress was calculated by dividing the load by the

initial cross sectional area. A typical stress-strain curve is shown in Figure 3.17 for

sample B75.















S150










50




0 2 2 .2
0 2 4 6 8
Engineering Strain (%)

Figure 3.17. Typical stress-strain curve generated from tensile testing. Line tangent to
slope of elastic portion of curve used to determine elongation at failure is shown.



Fatigue

Samples

Fatigue samples were sectioned from many different locations in the as-received

forging, as discussed above. One facet of this study was to determine is the location of

the sample from the forging had any effect on the fatigue properties. To do this, the

sample location was recorded and compared to the calculated effective strain from

forging, as shown in Figure 3.18. The samples that were taken from various locations

have been divided into three groups: low (,ff <1.0), medium (1.0< 6,, < 2.0) or high (Ef

> 2.0) strain locations.


















Low Med High
0=.13160 D=1.200 G-2.100
A=.300 E= 1.500 H 2.400
B= .600 F= 1.800 1= 2.700
C= .800 2.957

Figure 3.18. Calculated iso-strain lines from forging showing various effective strain
from forging and respective fatigue sample location.



The effect of cooling rate on microstructure was also discussed above, and the

calculated cooling rate from forging was also superimposed over the diagram of fatigue

samples, as shown in Figure 3.19. The effect of location on the fatigue results was

determined and will be discussed in Chapters 4 and 5.


Low Med High
1=55 E 78 -= 115
J152 F 72 B =107
K=48 G=66 C=96
L=45 H=60 D 87
S=43

Figure 3.19. Calculated cooling rates from 1080C to 8700C in C/min and corresponding
fatigue specimen location.



The fatigue samples were 8.89 cm (3.5 inches) long with a uniform gauge length

of 2.22 cm (0.875 inches). The gauge diameter was 5.08 mm (0.200 inches) with a

surface finish of 8 microinches, and the sample was gripped by /2-20 threads. This









sample is shown schematically in Figure 3.20. In order to avoid stress concentrations at

the surface, the samples were longitudinally polished in the uniform gauge to eliminate

all circumferential grinding marks and to produce all polishing marks parallel to the

loading axis of the sample.

3.500.010
8 --- 1.750.005-






.500-20 UNF 3A
2 PLACES _- .2001.001 DIA

Figure 3.20. Drawing of fatigue sample used in this research. The dimensions are in
inches


Thermal cycling of all fatigue samples was conducted in the Kominsky thermal

cycling furnace shown in Figure 3.13. The samples were arranged horizontally on a

cordierite platen as shown in Figure 3.21. Samples were thermally cycled for 1, 10 and

100 cycles, as well as exposed in static air for 92 hours to compare with the 100 cycle

samples. Samples were also encapsulated in quartz filled with argon. These samples

were placed inside titanium tubes and inserted along with titanium sponge which was

added as an oxygen getter in quartz tubes. The quartz ampoules were evacuated prior to

backfilling and then filled to 200 torr with argon before sealing. On heating to 6500C the

pressure inside the ampoule reached approximately 760 torr.






















Figure 3.21. Photograph of LCF samples loaded for thermal cycling. Samples are
separated by the alumina tubes.


Testing

The fatigue testing was conducted on an MTS 810 servohydraulic test machine

with a 20,000 pound load cell. The testing was conducted under strain control using

resistive strain gauge extensometers (model number MTS 632.53E-14) with high

temperature alumina rods. The A ratio was 1 (,,,, = 0), calculated as

A = o/o, (3.1)




where o, is the stress amplitude defined as

O max O mn
2 (3.2)


and o,,, the mean stress defined as

O max+ O mml
n,=- 2 (3.3)




The cycling rate was 20 cycles per minute. The testing was computer controlled

using Testware SX software that incorporated data collection. Peak and valley load data









were collected electronically for every cycle. Load-strain information was recorded at

various cycles on a strip chart recorder, as shown in Figure 3.22. Hydraulic grips were

used to eliminate the need for a backing nut on the threads thus allowing for tension and

compression loading. The elastic modulus of each sample was determined by loading the

sample in the elastic region and recording the resultant strain on a strip chart recorder, as

shown in Figure 3.22. The applied load was divided by the sample cross sectional area in

the uniform gauge to determine the applied stress. The calculated stress was then divided

by the measured elastic strain in order to estimate the room temperature elastic modulus.

This room temperature modulus was multiplied by the applied strain during testing in

order to give the average, or pseudo-stress applied during the fatigue testing. This

pseudo-stress is typically reported as the alternating pseudo-stress, which is the pseudo-

stress divided by two.

Analysis

After fatigue testing, two cross-sections were cut from the sample, one containing

the fracture surface and one behind the fracture surface, as shown in Figure 3.23. The

cross section containing the fracture surface was sectioned longitudinally to observe the

crack path, the initiation site and opposite the initiation site on the longitudinal section.

These samples were mounted in Conductomet such that the initiation site in the

longitudinal section was on the same side as the cross section as shown in Figure 3.23.

This allowed the initiation site to be identified during metallography.










Specimen #52, Strain Control, 0.95%E, 20CPM, R.T.


3000 .


S 2000

1000


/ i 0.004


Strain (in/in)

Figure 3.22. Typical LCF data taken during testing showing the hysteresis loops at
various cycles. Curve for modulus determination is also shown.


Figure 3.23. Illustration of sectioning and metallographic mounting of the fatigue sample
for evaluation.


The samples were then examined in the SEM using backscatter analysis or on the


optical metallograph after etching.









Coatings


All of the samples to be coated were ultrasonically cleaned in a warm Alconox

bath for 30 minutes, followed by a de-ionized water rinse and a 2-propanol rinse. The

samples were dried with a K ,.i- .pi;' i:. 'cl and placed in sterile plastic bags, and only

handled with gloves prior to coating.

Low Pressure Plasma Spray (LPPS) Coatings

For LPPS, the oxidation pins were secured in a custom designed fixture that holds

the pin at each end and thus allows the complete pin to be coated instead of just half of

the pin. This cuts the processing time in half and eliminates any overspray in the center

of the pin. The fixture is shown in Figure 3.24. After the sample was loaded into the

fixture, the fixture was loaded into the LPPS chamber. The chamber was then pumped

down to 50 torr and backfilled with argon. The sample oscillation and rotation were

initiated and the plasma guns were turned on to pre-heat the sample to approximately

650'C. The powder feeders were then started and the sample was coated to the proper

thickness. The plasma guns and feeders were then shut off and the sample was allowed

to cool in the chamber. The sample was then removed from the single pin fixture and

placed in a fixture with several other samples where only the end radii are exposed.

These radii were also coated to insure that the weight gain versus unit area measured

during oxidation was as accurate as possible.

















































Figure 3.24. Drawing ofLPPS specimen fixture.


NiCrAlY + 40 vol% A1203.

NiCrAIY (Ni-21.7Cr-10AI-1.13Y (wt%)) powder (purchased from Praxair) was

low pressure plasma sprayed (LLPS) onto a scrap oxidation pin to assess the LPPS









process. The NiCrAIY broke away from the pin in three large sections after processing.

Since it was initially suspected that the fixture might be the cause of the lack of adherence

another trial was attempted, this time holding the pin on one half only. The result was the

same; the coating broke away from the pin after processing. The pin was sectioned in the

location of the coating failure and examined using SEM backscatter techniques. It was

determined that there was residual Ni left on the surface of the pin, indicating that the

coating was initially adherent but subsequently spalled. This spallation was probably a

result of the large thermal expansion differences between the substrate and the coating.

Consequently, a mixture of 60 vol% NiCrAlY and 40 vol% A1203 was produced by

blending, and this "cermet" was deposited on the substrate. This mixture provided a

suitable thermal expansion match between the coating and substrate and, as a

consequence, the coating was adherent.

FeCrAlY.

FeCrAIY (Fe-29.9Cr-4.9A1-0.6Y (wt%)) powder was also applied to the oxidation

pins via LPPS. Since this coating was successfully applied and did not spall, no changes

were made to the deposition process.



Sputtered Coatings.

Several elemental coatings were deposited onto oxidation pins via vacuum

sputtering. The samples were held on one end and rotated in the sputtering chamber via a

magnetic feed through. The elements were either reacted with the substrate at elevated









temperature to create a stable compound at the surface(e.g., A13Ti) or were used in their

pure elemental form.

Sputtered Al.

25 pm of Al was sputtered onto the oxidation pins in a vacuum sputtering unit.

These pins were then reacted in dry, pure argon using the following heat treatment

schedule:

Ramp to 6000C at 3000C/hour,

Hold at 6000C for 2 hours,

Ramp to 6300C at 100C/hour,

Hold at 6300C for 10 hours,

Furnace cool.

X-ray diffraction confirmed that a mixture of A13Ti and A13Nb were formed at the

surface, as shown in Figure 3.25.


Figure 3.25. SEM backscatter micrograph of Al coated and reacted sample showing the
in-situ formed A13Ti coating on the right.









Sputtered Si

Elemental Si was also sputtered onto the orthorhombic Ti-aluminide substrate in a

vacuum sputtering unit. The sputtering rate was very low and, as a consequence, only

12.7 pm of Si was deposited. The Si coated pins were reacted as follows:

Ramp to 12000C,

Hold at 12000C for 16 hours,

Furnace cool.

This cycle was designed to produce the TisSi, phase on the surface and XRD

indicated that this phase was indeed produced, as shown in Figure 3.26. The reaction

temperature was above the p transus for this alloy, and the microstructure indicated that

the prior 3 grains had indeed grown and the as forged microstructure was eliminated.




















Figure 3.26. SEM backscatter micrograph of Si coated and reacted sample showing in-
situ formed TisSi3 coating on the right.









Sputtered Pt and Cr

Pt and Cr was sputtered onto the oxidation pins in order to provide protection

based on the low oxidation rates of pure Pt, and the stability with titanium [Fuj79, Eyl85],

and because Cr has been show to provide good oxidation resistance as a coating for

titanium [McK93, McK90]. Both coatings were adherent, although the Pt coating was

porous and columnar, as shown in Figure 3.27.










P. o acting .. ...
*1; i. .... ^ .






Figure 3.27. SEM backscatter micrograph of Pt coating showing the columnar structure
and porous nature of the coating.


CVD coatings

Several oxide coatings were produced by metal organic chemical vapor deposition

(MOCVD). The idea was to evaluate the stability of the oxide in contact with the

substrate and determine if the Ti would reduce the oxide. MOCVD was selected because

of the capability of producing adherent coatings of oxides on any and every surface in the

reactor. The CVD process, unlike PVD or LPPS, is not line-of-sight, and therefore every










hot surface in the reactor will be coated. The MOCVD process consists of a metal

organic precursor that is evaporated and flowed over the substrate held in a reaction

chamber at elevated temperature, as shown in Figure 3.28 and Figure 3.29. The selection

of the precursors and the deposition temperatures are described below.



















Figure 3.28. Photograph of CVD reactor.



/ Pressure
Gauge

Furnace

Samples



Furnace -






To Vacuum Matl
Pump Liquid Mantle
Nitrnn-n Heater


Figure 3.29. Schematic diagram of CVD reactor shown in Figure 3.28.









SiO2

Oxide and silicide coatings have been reported to provide oxidation protection for

ca-p titanium [Sol85, Cla88, Bed91], r2 alloys [Wei89, Dev90] and orthorhombic alloys

[Coc96]. The mechanism is through the production of an oxide comprised mainly of

SiO,. For this reason, SiO2 was chosen as a diffusion layer for an outer coating, and

possibly as a primary oxidation coating. The reagent used to produce the SiO, was

silicon tetraethoxide [Si(OC2H,)4]. The reactor temperature was 6700C at a pressure of

500pim. The Effusion cell temperature was 900C and the flow rate of reagent was 2x10-5

gs 'cm-2. The deposition rate of the SiO, was approximately O.lnm/sec with a reagent

utilization of approximately 0.5%. The reaction taking place was

SI II i I, i SiO2+ CHsOH+C2H4

Tao,

Tantala was chosen as a barrier layer due to its higher CTE over SiO, and greater

oxidation resistance relative to TiO,. The reagent used to produce the Ta20s was tantalum

ethoxide dimer [Ta,(OC,H,) ,,]. The reactor temperature was 4250C at a pressure of

100m. The Effusion cell temperature was 1250C and the reagent flow rate was 3x10'-

gs 'cm-2. These conditions led to a deposition rate of the Ta,O, of approximately

0.lnm/sec with a reagent utilization of approximately 0.5%. A coating thickness of 0.45

tm was typically deposited on the oxidation pins and fatigue samples. The reaction

taking place was

Ta2(OC.H,),, = TaO, + 5 CZHsOH + 5 CH4









MgO

MgO has a high oxygen diffusivity in air at 6500C and therefore it probably would

not provide adequate protection as a stand alone oxidation coating. But MgO has been

evaluated as a barrier layer for fiber-reinforced titanium aluminide composites [McG95].

It was shown to dissolve oxygen into Ti-24A1-11Nb after 100 hours at 1000C, but at a

reasonably low rate. MgO has also been formed in-situ in Ti-Mg alloys produced by high

rate evaporation and quenching. After oxidation at 850C for only 10 minutes, MgO

particles were formed in the titanium matrix [War95]. For this reason and because MgO

has a higher thermal expansion coefficient than the Ti-22A1-26Nb alloy, it was chosen as

a candidate coating system.

The reagent used to form MgO was magnesium 2,4 pentane dionate and the

chemical reaction was

Mg (C,HO,), HO-, MgO + 4 CHsOH + CH4

The reaction temperature was 540C at a partial pressure of 0.5 torr. Oxygen was

the initial carrier gas and was chosen to reduce the level of carbon in the deposited film.

After deposition, the film thickness was 0.85 tim. The samples were baked out in an air

furnace for 1/2 hour at 575C, and the film thickness decreased to 0.3 pm. This is

probably due to evolution of carbon into carbon dioxide in the furnace. The source of

carbon is from the reaction of the reagent in the furnace and the incomplete formation of

the reaction products. This process is still in the lab stages and is not yet fully reduced to

practice.















CHAPTER 4
RESULTS AND DISCUSSION

Oxidation


Cyclic Oxidation Testing

Baseline results

Four baseline samples were cyclically oxidized at 6500C. Three of the samples

(136, 187 and 191) were chosen to have different levels of strain from forging, and thus

different grain size and grain aspect ratios, as discussed in Chapter 3. These samples

were chosen to determine if the strain history, which results in a slight variation in

microstructure, has any effect on the oxidation behavior. The results from these baseline

studies are shown in Figure 4.1. The difference in oxidation rates between these three

samples is small as is the weight change over time. Therefore, it can be stated that there

is little or no effect on oxidation from the strain history. The fourth sample, 174, was run

in a separate oxidation test with another series of samples as a control. The differences

observed between the samples indicates typical experimental scatter.























--&- 136
-- 187
-- 191
- 174
+-Average


0 200 400 600 800 1000
Time (hrs)

Figure 4.1. Baseline oxidation results showing no variation between three samples from
different locations run together, and a larger variation with a sample from a different
experimental run.



Al and Si Coated and Reacted Results


The results from the oxidation testing on Al and Si coated and reacted samples are

shown in Figure 4.2. The average of the baseline samples is also plotted as a reference.

As can be seen, both the Al and Si coated and reacted samples behaved well in cyclic

oxidation tests. This behavior has been observed previously for Ti coated with AIlTi

[81Str, 85Unn, 91Abd, 91McM, 92Gal, 93Abb, 93Wie] and a, coated with Al3Ti [88Sub,

90Kun, 90Smi, 93Smi]. Silicide coated Ti-22A1-27Nb has also been evaluated and


~t

.-fi


X,


- -








shown to have similar oxidation behavior as the Si coated and reacted Ti-21A1-26Nb

[96Coc].


SAl
-E- Si
--- Baseline


0 200 400 600 800
Cycles (hrs)
55 mms @ 650C, 5 mms @ R T
Figure 4.2. Al and Si coated and reacted oxidation results showing the decrease in
oxidation rate over the baseline.


Oxide Coating Results

The effect of cyclic oxidation on the oxide coated samples are shown in Figure

4.3. Both Ta20, and MgO behave similarly to the baseline alloy while the SiO, coated

sample exhibited better oxidation behavior than the baseline at short times, but began to

gain weight in a linear fashion with time after 400 cycles. This indicates that the coating

was not protective and that oxygen was diffusing through after 400 hours causing the


1000


o-
'Ir


I










sample to gain weight. These results indicate that the oxide coatings by themselves do

not provide protection from an oxidizing environment.


1.2-~~--..--------
-0- SiO2
a205
S + MgO
Baseline


S0.8 v/















0 200 400 600 800 1000
+

I 0.6'












0 200 400 600 800 1000
Cycles (hrs)
55 mins @ 650'C, 5 mins @ R.T.

Figure 4.3. Oxide coated oxidation results showing that TaO, and MgO coated material
behaves like uncoated material and SiO2 coated samples begin to gain weight after 400
cycles.



Metallic Coating Results


The results from the cyclic oxidation testing of the metallic coated samples are

shown in Figure 4.4. The results show that the MCrAlY coatings provide excellent

oxidation resistance, that Cr provides intermediate protection, and that Pt has little effect

over baseline behavior. The protection by the MCrAlY coatings has been shown in

previous work on Ct alloys [92Bri, 93McK, 95Sch] and titanium alloys [92Tob, 90Lut,









90McC, 93McK]. On the other hand, platinum, in the form of PtAl2, has been shown to

reduce the cyclic oxidation rates in conventional titanium alloys and in Ti3AI [93Nic].

Pure platinum coatings have also been shown to improve creep resistance [85Eyl] and

decrease oxidation rate [79Fuj] of Ti-6242. Finally, pure Cr coatings have also been

shown to reduce the weight gain over time of Ti-64 and Ti-6242 [90McK].

1.2 -
-NiCrAIY+AI203
-A- FeCrAIY
1- -Pt
--C--- Cr Bs
---- Baseline

le0.8 -V

& \

en~
0.6
U


0.4


0.2-- -- -


0 200 400 600 800 1000
Cycles (hrs)
55 mins @ 650aC, 5 mins @ R T

Figure 4.4. Results showing the large reduction in oxidation rate with the application of
the MCrAlY coatings and the effect of Cr and Pt coatings.


The results from all of the cyclic oxidation tests are compared in Figure 4.5. As

can be seen, the metallic coatings, excluding Pt, provide excellent oxidation resistance










after 1000 cycles, whereas the Ta20O and MgO coatings behave similarly to the baseline,

and the SiO, coated samples begin to degrade after 400 hours.


1.2



1



- 0.8



S
. 0.6



0.4


- Al
-- Si
-o S10
X- Ta20
+ MgO
N-- NiCrAIY+AlO0
-A- FeCrAIY
-- Pt
- Cr
-- Baseline


0 200 400 600 800 1000
Cycles (hrs)
55 mms 650C, 5 mins @ R.T

Figure 4.5. Comparison of the oxidation rate behavior of all coatings examined in this
study.


Microhardness and Microprobe Evaluation


Baseline uncoated


As shown in Figure 4.1, the three samples with varying forging strains behaved

similarly, so after 500 cycles samples 191 was removed for analysis. Both microhardness

and oxygen content vs. depth are plotted in Figure 4.6. Notice that in both curves, the

oxygen content and microhardness follow the same trend. Further work is needed to









determine if the high value of oxygen in the 1000 cycle sample at 25gm is scatter or an

indication of localized oxidation, possibly an oxide.

25 --- 1100

SOxygen (at%) 1000 cycles
SOxygen (at%) 500 cycles 1000
20-A- O p hardness (KHN) 1000 cycles
'- hardness (KHN) 500 cycles 900


800
15 800

nt -700

O ,10 \
St 600

500
5-
400

0 300
0 10 20 30 40 50 60
Distance from Coating/Substrate Interface (gim)

Figure 4.6. Uncoated baseline oxygen content and microhardness vs. depth.



The backscatter SEM images of the two microhardness samples (Figure 4.7 and

Figure 4.8) indicate that the dark zone near the surface of the sample is high in oxygen.

The microhardness indentations visible in these figures clearly show that these regions

are much harder that the base material. The depths of oxygen rich zone in the 500 and

1000 cycle samples are approximately 20pm and 35pm respectively. Dividing the square

of the depth (in cm) by the time in seconds, it is possible to estimate the diffusivity of

oxygen into the O + B2 lattice at 6500C as:

x = (cm) [4.1]































Figure 4.7. SEM backscatter image of baseline microhardness sample cycled for 500
cycles.
























Figure 4.8. SEM backscatter image of baseline microhardness sample cycled for 1000
cycles.
cycles.









or

X2
D =- (cm2/sec) [4.2]
t

This yields a D of approximately 3.4 x 10-12 cm2/sec. This estimate ignores the oxide

present on the surface of the sample as well as any interaction that oxide has with the

substrate, and the fact that these are no isothermal experiments. These are a valid

assumptions given the fact that the oxide has a diffusivity of oxygen that is 50x that of the

metal [Unn86]. The diffusion coefficient of oxygen for CP cL-Ti at 6500C is

approximately 2.5 x 10-12 cm2/sec [Kub83] and for Ti-24Al-15 Nb at 10270C is

approximately 2.6 x 10-12 cm2/sec [Roy96]. These values correlate well to the 3.4 x 10-

12 cm2/sec value calculated from the baseline samples.

FeCrAIY and NiCrAlY + A1203

Brindley et al. [Bri92] have evaluated the oxidation behavior of both of these

coatings, and found that while they provide excellent resistance to oxygen ingress into

titanium aluminides, the elements in the coatings themselves diffuse into the substrate

and cause embrittlement. Figure 4.9 is an SEM backscatter image of the NiCrAlY +

A1l03 coated specimen after 1000 cycles. The NiCrAlY + A1203 plasma sprayed coating

can be seen in the lower right hand comer. Notice that the microstructure of the matrix

next to the coating has changed from the lath morphology to a fine, equiaxed structure.

Work conducted by Schaeffer et al. [Sch95] on U2 coated with NiCrAIY + A1203

demonstrated that the coating alone, without exposure, degraded the LCF life by two


orders of magnitude.

























Figure 4.9. SEM backscatter image of the NiCrAIY + A1203 coated sample after 1000
cycles at 6500C.


The results from this study support Schaeffer's results, as can be seen from the

microhardness and oxygen profiles in Figure 4.10 through Figure 4.12. Figure 4.10 and

Figure 4.12 compare the oxygen level in the substrate to the measured microhardness.

Notice that there is simply one point under the coating that contains oxygen, and beyond

that point the level is below the detectability limit of the EMPA. Figure 4.11 and Figure

4.13 show the Fe and Ni concentration vs. microhardness. The Fe and Ni content seems

to correlate with microhardness, and the Ni and Fe levels both taper off with distance

from the interface, as does the microhardness. This could contribute to a loss in ductility

and thus fatigue life. Recall that LPPS is conducted in a partial pressure of argon and air,

and the sample is preheated to 650C for several minutes prior to coating during which

oxygen could ingress. A diffusion barrier between the substrate and coating might

possibly control the ingress of Ni or Fe from the coating into the substrate, but the pre-

heat might still cause embrittlement.














10 Oxygen (at%) 650
hardness (KHN)
600
8

o :550
S6
4500
4
0 450 ,

2
400

0 350


-2 -' 300
0 10 20 30 40 50 60 70
Distance from Coating/Substrate Interface (pm)

Figure 4.10. Microhardness and oxygen vs. depth for the FeCrAlY coated and 1000
cycle exposed sample.


Fe (at%)
hardness (KHN)


S02


S015


550


500


450


0 10 20 30 40 50 60 70
Distance from C (,.,Iri Su.,hiraji Interface (pm)

Figure 4.11. Microhardness and iron vs. depth for the FeCrAlY coated and 1000 cycle
exposed sample.













Oxygen (at%)

hardness (KHN)


2.5


2


S 1.5


0


500

C-
450


400


0 10 20 30 40 50 60 70
Distance from Coating/Substrate Interface (pm)

Figure 4.12. Microhardness and oxygen vs. depth for the NiCrAIY + Al203 coated and
1000 cycle exposed sample.


NI (at%)

hardness (KHN)


02


500
C-

450


400


-0,1 .1 .. 300
0 10 20 30 40 50 60 70
Distance from Coating/Substrate Interface (pm)

Figure 4.13. Microhardness and Ni vs. depth for the NiCrAIY + A1203 coated and 1000
cycle exposed sample.










Sputtered and Reacted Al and Si


Elemental Al and Si were sputtered onto the orthorhombic under vacuum. The Al

coated samples were then ramped to 6000C in two hours and held for two hours followed

by ramping to 6300C in 3 hours and holding for 10 hours. This produced A13(Ti,Nb) as a

surface reacted coating, as verified by XRD. The Si coated samples were ramped to

12000C as quickly as possible and held for 16 hours. The intent was to produce the TiSi3

phase and miss the TiSi phase. This was accomplished as can be seen in Figure 4.17,

and also verified by XRD. The microhardness profiles vs. the oxygen content are shown

in Figure 4.14 for the Al coated sample and in Figure 4.16 for the Si coated sample.

Figure 4.15 and Figure 4.17 are SEM backscatter images of the Al and Si coated and

reacted samples, respectively, obtained after cycling at 650'C for 1000 cycles.

025 -- 800

Oxygen (at%)
0.2
hardness (KHN) 700

0.15


0 5

O 500
0.05


0 400



-0.05 -- -- 300
-5 5 15 25 35 45 55 65
Distance from Coating/Substrate Interface (pm)

Figure 4.14. Microhardness vs. oxygen content for the Al coated and reacted sample.
Note that the distance from the interface starts at -5tm.





















i I E




Figure 4.15. SEM backscatter image of Al coated and reacted sample after 1000 cycles at
6500C. Notice the crack in the AlTi coating.


Notice that the A13(Ti,Nb) coating has cracks that run to the substrate, yet the

coating behaved very well in oxidation. The SEM backscatter images do not show any

indications that oxygen diffused preferentially down the cracks and into the substrate.

This could be due to the fact that on heating, the coating expanded more than the

substrate and thus the cracks were not open during exposure. Microhardness

measurements were taken in the reacted layer because in optical microscopy the layer

looked like the substrate. Therefore, the oxygen analysis and the microhardness start at

minus 5tm in both the Al and Si coated cases. The zero point is the location of the

coating/substrate interface after the reaction. This indicates that the coating was the

major source of oxidation and hardness. Directly beneath the reacted coating, the oxygen

content was negligible and the microhardness had returned to that of the substrate. But

since the coating was atomistically attached to the substrate, any cracks in the coating

would propagate into the substrate and cause limited LCF life.











0.14-


0.12-


0.1 -


0.08-
--

0.06-
S0.04
0 0.04-


Oxygen (at%)

- hardness (KHN)


900


800


700
(5
600


500


002 400


0 300


-0.02 -. 200
-5 5 15 25 35 45 55 65
Distance from Coating/Substrate Interface (im)

Figure 4.16. Microhardness vs. oxygen content for the Si coated and reacted sample.
Notice that the distance from substrate starts at -5pm.


, i '- "-'r*''i.




1 1- 1 f '.

*


Figure 4.17. SEM backscatter image of the Si coated and reacted sample after 1000
cycles.









Sputtered Pt and Cr coatings

Pt was sputtered onto the orthorhombic substrate to a thickness of about 1 pm.

The Pt coating (Figure 4.18) was columnar, porous and discontinuous. As a result, it

allowed oxygen to diffuse to a similar depth as in the uncoated material. The

microhardness and oxygen profiles (Figure 4.19) are consistent with the poor behavior of

the coating in oxidation. Had the coating been dense and adherent, it should not have

allowed oxygen to pass at 6500C according to previous work conducted on Pt coated Ti-

6242 [Fuj79] where the rate of weight gain was reduced from 3x10-1 to 2x104 at 5930C.

Their coating, deposited by ion plating, was approximately 1 Im thick and was dense and

continuous. Sputtered Cr exhibited better oxidation behavior than Pt during cyclic

oxidation, as shown in Figure 4.4. The level of oxygen beneath the coating after

exposure is shown in Figure 4.20, along with the level of Cr. Oxygen content for the

unexposed sample was zero as determined by EMPA. There is essentially no Cr diffused

into the matrix after 1000 cycles at 6500C, and very little oxygen, compared to the

baseline of 20 at%.





















Pt coating

*hA.^ fc:--`--; .-<, ..j '


Figure 4.18. SEM backscatter micrograph of Pt sputter coated sample cycled for 1000
cycles.


25


Oxygen (at%)
hardness (KHN)


700 _

0-
600


500


0 10 20 30 40 50 60
Distance from Coating/Substrate Interface (I'm)

Figure 4.19. Microhardness and oxygen content vs. depth for the Pt coated and 1000
cycle exposed sample.










2-
S-- Oxygen
S- Cr


1.5-





1-





0.5-



0 Im n m 1


0 5 10 15 20 25
Depth (pim)
Figure 4.20. Plot of Cr and oxygen content vs. depth in 1000 cycle exposed oxidation
sample.



CVD Oxide Coatings


The CVD SiO2 coating was dense, adherent, and appeared iridescent green in

color. The Ta20, and MgO coatings were also dense and adherent. All coatings were

approximately 0.5 im thick.

As can be seen from Figure 4.3, the initial oxidation behavior of the SiO2 coating

was excellent. This data is an average of three separate tests conducted with the SiO,

coating. However, after 400 cycles, the oxidation behavior of the SiO2 coated samples

began to behave in a more linear relationship with time instead of parabolic. This is

typically an indication that there is no protective oxide present and that oxygen is simply




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FILES


EFFECT OF COATING AND EXPOSURE ON THE OXIDATION AND
MECHANICAL PROPERTIES OF Ti-22Al-26Nb
By
JAMES ROSS DOBBS
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
1997

Copyright 1997
by
James Ross Dobbs

To my loving wife, Jane.

ACKNOWLEDGMENTS
The title of Doctor of Philosophy is bestowed on the individual, but the road to
obtaining that title is truly a team effort. My road was paved with many supportive,
capable individuals whom I would like to recognize. First and foremost I thank my wife,
Jane, for the support and love during a period in our relationship when there was
something other than her that commanded much of my time and attention. I thank my
parents and brothers and sister for their support and enthusiasm and for always being
willing to listen. I would also like to thank my thesis advisors, Prof. Mike Kaufman
whose friendship I value as much as his technical interaction and Dr. Mike Gigliotti who
never let me forget the focus of my goal. I would also like to thank Prof. Teresa Pollock
whose gentle reminders and kind inquiries gave me inspiration to finish.
I would like to thank several people who helped me conduct the experimentation
and understand the results: Mr. Mike Gilhooley for his help and guidance in LCF testing
and Mr, Chris Canestraro for his assistance in tensile testing, Mr. Keith Borst for
producing the sputter coatings, Mr. Amie Henry for teaching me the LPPS system, and
Mr. Kevin Janora for CVD coatings. I would like to thank Dr. Ravi Ravikumar and Mr.
Mike Larson for their guidance and patience while I learned and practiced TEM, Ms.
Cindy Hayden for assistance in conducting and interpreting XRD, Dr. Bob Gilmore for
acoustic microscopy and acoustic modulus measurements and Dr. John Ackerman for
guidance through many experimental decisions and his wealth of CVD knowledge and
IV

his ability to convey that knowledge. A special note of thanks to Dr Mike Henry for
lengthy and insightful discussions on fatigue, and Prof. Harry Lipsett for many hours of
discussion on titanium metallurgy and relevant experimentation.
I would like to thank Dr. Ann Ritter for support, both financial and moral.
Thanks to Dr. Jeff Graves and Dr. Tom Cox for having enough faith in me to employee
me while I completed this work. And a special note of thanks to Ms. Karen Keevem for
valuable assistance in preparation of this document.
v

TABLE OF CONTENTS
page
ACKNOWLEDGMENTS iv
LIST OF TABLES ix
LIST OF FIGURES x
ABSTRACT xvii
CHAPTERS
1 INTRODUCTION 1
2 LITERATURE REVIEW 5
Overview of Ti-Al-Nb Metallurgy 5
History 5
Description of Phases Present in Ti-Al-Nb 7
Deformation Behavior of a,, O and B2 Phases 14
Tensile Behavior of a2 andO Alloys 17
Fracture Toughness and Fatigue 22
Overview of Oxidation Behavior 26
Oxidation of Uncoated Titanium Alloys 26
Oxidation of Coated Titanium Alloys 28
Al3Ti 28
MCrAlY 29
Pure metals 29
Glasses, silicates and oxides 30
Embrittlement of Titanium Base Alloys 31
Effects of Coatings on Mechanical Properties of Ti Alloys 33
Program Approach 34
3 EXPERIMENTAL PROCEDURE 35
Material Processing 35
Specimen Blanking 42
Test Specimens and Test Procedures 45
Oxidation 45
vi

Samples 45
Testing 45
Tensile 49
Samples 49
Testing 49
Fatigue 50
Samples 50
Testing 53
Analysis 54
Coatings 56
Low Pressure Plasma Spray (LPPS) Coatings 56
NiCrAlY + 40 vol% AI2O3 57
FeCrAlY 58
Sputtered Coatings 58
Sputtered A1 59
Sputtered Si 60
Sputtered Pt and Cr 61
CVD coatings 61
SiO, 63
Ta,65 63
MgO 64
4 RESULTS AND DISCUSSION 65
Oxidation 65
Cyclic Oxidation Testing 65
Baseline results 65
A1 and Si Coated and Reacted Results 66
Oxide Coating Results 67
Metallic Coating Results 68
Microhardness and Microprobe Evaluation 70
Baseline uncoated 70
FeCrAlY and NiCrAlY + AI2O3 73
Sputtered and Reacted A1 and Si 77
Sputtered Pt and Cr coatings 80
CVD Oxide Coatings 82
Effect of Forging Strain and Cooling Rate on the Tensile and LCF
Behavior 86
Fatigue Results 86
Tensile Results 87
Discussion 88
Effect of Exposure at 650°C on the Tensile and LCF Behavior 95
Tensile Results 95
Fatigue Results 96
Discussion 98
vii

Effect of Coating and Exposure on the LCF Behavior 105
Results 105
Discussion 110
5 SUMMARY AND CONCLUSIONS 117
Oxidation 117
Forging Location 118
Environmental Exposure 119
Coated LCF Behavior 120
LIST OF REFERENCES 121
BIOGRAPHICAL SKETCH 132
viii

LIST OF TABLES
Table page
Table 3-1. Compositions from several locations along the length of the 15.25 cm
diameter billet, surface and center, and from the forging used in this study 36
Table 3-2. Forging parameters for the material used in this study. The first line is step
1 and the second line is step 2 36
Table 4-1. Testing parameters and results for samples used to determine effect of
location on LCF 86
Table 4-2. Tensile test data 95
Table 4-3. Low cycle fatigue data of environmentally exposed samples 96
Table 4-4. Coated LCF data 106
ix

LIST OF FIGURES
Figure page
Figure 1.1. Specific yield strength vs. temperature comparing an a, and orthorhombic
Ti-Al-Nb alloy with a conventional Ti alloy and IN718. After [Woo93] 2
Figure 2.1. Titanium - aluminum equilibrium phase. After [Mur87] 6
Figure 2.2. Comparison of the specific yield strength of a2 and orthorhombic alloys
as a function of temperature. After [Row91] 7
Figure 2.3. Ball model of basal plane of a, showing lattice sites for Ti and A1 atoms.
The a phase lattice is similar, with the Ti and A1 atoms arranged randomly 8
Figure 2.4. 900°C isotherm from the Ti-Al-Nb ternary phase diagram. B2, a2, and
orthorhombic phases are in equilibrium. After [Row92] 9
Figure 2.5. Ball model drawing of the B2 phase showing the relationship between the
titanium, aluminum and niobium atoms 10
Figure 2.6. Comparison of the basal planes of the a2 and orthorhombic phases 11
Figure 2.7. Constant 25 at% A1 isopleth from the Ti-Al-Nb ternary in Figure 2.4.
After [Ban95] 12
Figure 2.8. SEM Backscatter micrograph of Ti-22Al-26Nb processed in the O + B2
field and directly aged at 815°C for 4 hours. After [Woo93] 13
Figure 2.9. SEM Backscatter micrograph of Ti-22Al-26Nb processed in the B2 field
and directly aged at 815°C for 4 hours. After [Woo93] 14
Figure 2.10. Crystal structure and burgers vectors for the a2 and O phases. After
[Ban95] 15
Figure 2.11. Slip bands in the B2 phase in TÍ-24A1-1 INb showing the inhomogeneous
nature of slip. After [Ban90] 17
Figure 2.12. Yield strength vs. cooling rate for two a2 alloys. After [Gog90] 17
Figure 2.13. Elongation vs. cooling rate for two a2 alloys. After [Gog90] 18
x

Figure 2.14. SEM backscatter micrograph of Ti-22Al-26Nb forged in the B2 region
and directly aged at 804°C. Notice the colony of O laths that have formed off of
the large lath 19
Figure 2.15. Intergranular crack observed in tensile tested Ti-22Al-27Nb solutioned
above the p(B2) transus and aged at 870°C for 50 hours 20
Figure 2.16. Fractograph of forged + direct aged material. Notice the high GAR and
the initiation of failure from an internal boundary 21
Figure 2.17. SEM fractograph of material solutioned above the p(B2) transus and aged
at 870°c for 50 hours. Notice that there is mixed intergranular and transgranular
fracture across the sample and the prior p grain size is approximately 500pm 22
Figure 3.1. BSE image of a typical microstructure of the Ti-22Al-26Nb material after
B2 forging and heat treatment at 815°C for 4 hours 37
Figure 3.2. Photomacrograph of one half of the forging showing flow lines 38
Figure 3.3. Forging half showing location of lOx macrophotographs and lines for
sectioning into metallographic samples 39
Figure 3.4. Calculated iso-strain lines for the forging examined in this study 39
Figure 3.5. Photomicrograph taken from area C. Notice the equiaxed grain structure
and low grain aspect ratio 40
Figure 3.6. Photomicrograph taken from area P. Notice the high grain aspect ratio
resulting from the higher effective forging strain 40
Figure 3.7. High resolution SEM image of area S (identified in Figure 3.3) from the
slowest cooled region in the forging 41
Figure 3.8. High resolution SEM image of area Q (identified in Figure 3.3) from the
fastest cooled region in the forging 42
Figure 3.9. Fatigue and oxidation specimen sectioning diagram 43
Figure 3.10. Tensile and oxidation sample sectioning diagram. Shaded samples are
oxidation and white samples are tensile/creep 44
Figure 3.11. Details of sectioning diagram for oxidation (shaded) and tensile (open)
samples from Figure 3.10. Calculated effective strain is also depicted 44
Figure 3.12. Drawing of oxidation sample used in this research 45
Figure 3.13. Photograph of furnace used to conduct cyclic oxidation tests 46
xi

Figure 3.14. Oxidation pins loaded on hearth plate. The samples are the darker pins,
the thermocouples are the two white rods in the foreground. The cordierite block
is resting on the furnace platen 47
Figure 3.15. Typical thermal cycle recorded by strip chart recorder 48
Figure 3.16. Drawing of tensile sample used in this research. Dimensions are in
inches 49
Figure 3.17. Typical stress-strain curve generated from tensile testing. Line tangent to
slope of elastic portion of curve used to determine elongation at failure is shown. ..50
Figure 3.18. Calculated iso-strain lines from forging showing various effective strain
from forging and respective fatigue sample location 51
Figure 3.19. Calculated cooling rates from 1080°C to 870°C in °C/min and
corresponding fatigue specimen location 51
Figure 3.20. Drawing of fatigue sample used in this research. The dimensions are in
inches 52
Figure 3.21. Photograph of LCF samples loaded for thermal cycling. Samples are
separated by the alumina tubes 53
Figure 3.22. Typical LCF data taken during testing showing the hysteresis loops at
various cycles. Curve for modulus determination is also shown 55
Figure 3.23. Illustration of sectioning and metallographic mounting of the fatigue
sample for evaluation 55
Figure 3.24. Drawing of LPPS specimen fixture 57
Figure 3.25. SEM backscatter micrograph of A1 coated and reacted sample showing
the in-situ formed A13TÍ coating on the right 59
Figure 3.26. SEM backscatter micrograph of Si coated and reacted sample showing
in-situ formed Ti5Si3 coating on the right 60
Figure 3.27. SEM backscatter micrograph of Pt coating showing the columnar
structure and porous nature of the coating 61
Figure 3.28. Photograph of CVD reactor 62
Figure 3.29. Schematic diagram of CVD reactor shown in Figure 3.28 62
xii

Figure 4.1. Baseline oxidation results showing no variation between three samples
from different locations run together, and a larger variation with a sample from a
different experimental run 66
Figure 4.2. A1 and Si coated and reacted oxidation results showing the decrease in
oxidation rate over the baseline 67
Figure 4.3. Oxide coated oxidation results showing that Ta^ and MgO coated
material behaves like uncoated material and Si02 coated samples begin to gain
weight after 400 cycles 68
Figure 4.4. Results showing the large reduction in oxidation rate with the application
of the MCrAlY coatings and the effect of Cr and Pt coatings 69
Figure 4.5. Comparison of the oxidation rate behavior of all coatings examined in this
study 70
Figure 4.6. Uncoated baseline oxygen content and microhardness vs. depth 71
Figure 4.7. SEM backscatter image of baseline microhardness sample cycled for 500
cycles 72
Figure 4.8. SEM backscatter image of baseline microhardness sample cycled for 1000
cycles 72
Figure 4.9. SEM backscatter image of the NiCrAlY + AI2O3 coated sample after 1000
cycles at 650°C 74
Figure 4.10. Microhardness and oxygen vs. depth for the FeCrAlY coated and 1000
cycle exposed sample 75
Figure 4.11. Microhardness and iron vs. depth for the FeCrAlY coated and 1000 cycle
exposed sample 75
Figure 4.12. Microhardness and oxygen vs. depth for the NiCrAlY + AI2O3 coated
and 1000 cycle exposed sample 76
Figure 4.13. Microhardness and Ni vs. depth for the NiCrAlY + AI2O3 coated and
1000 cycle exposed sample 76
Figure 4.14. Microhardness vs. oxygen content for the A1 coated and reacted sample.
Note that the distance from the interface starts at -5pm 77
Figure 4.15. SEM backscatter image of A1 coated and reacted sample after 1000
cycles at 650°C. Notice the crack in the Al3Ti coating 78
xiii

Figure 4.16. Microhardness vs. oxygen content for the Si coated and reacted sample.
Notice that the distance from substrate starts at -5pm 79
Figure 4.17. SEM backscatter image of the Si coated and reacted sample after 1000
cycles 79
Figure 4.18. SEM backscatter micrograph of Pt sputter coated sample cycled for 1000
cycles 81
Figure 4.19. Microhardness and oxygen content vs. depth for the Pt coated and 1000
cycle exposed sample 81
Figure 4.20. Plot of Cr and oxygen content vs. depth in 1000 cycle exposed oxidation
sample 82
Figure 4.21. SEM backscatter image of the Si02 coated sample after 1000 cycles at
650°C 83
Figure 4.22. Microhardness vs. oxygen content for the Si02 coated sample 84
Figure 4.23. SEM backscatter image of Ta205 coated sample showing area analyzed
by Auger spectroscopy. Box outlines area where compositional maps were
obtained 84
Figure 4.24. Auger compositional maps showing the concentration of Ta, Ti and
oxygen in the boxed region shown in Figure 4.23 86
Figure 4.25. Number of cycles to failure vs. strain for various locations in the forging.
The sample number is shown beneath each symbol, and the failure location is
listed above each symbol 87
Figure 4.26. Ultimate tensile strength, yield strength and percent elongation (plastic)
vs. calculated effective strain from forging 88
Figure 4.27. LCF data showing the calculated a„ and af values, along with failure
location and sample number for each data point 90
Figure 4.28. Optical fractograph of specimen taken from a highly stressed region in
the forging (No. 52) 92
Figure 4.29. Secondary SEM photograph of fracture initiation site from specimen #52
shown in Figure 4.28 93
Figure 4.30. Secondary SEM photograph of the fracture initiation site shown in the
small white box in Figure 4.29 94
xiv

Figure 4.31. Optical fractograph of sample number 6 taken from a low strain region in
the forging 94
Figure 4.32. Ultimate tensile strength, yield strength and percent elongation of exposed
and unexposed samples tested in air and vacuum at room temperature and 540°C. ..96
Figure 4.33. Plot of number of cycles to failure vs. strain. Failure location is noted
with each sample 97
Figure 4.34. Number of cycles to failure vs. strain for the unexposed and exposed
samples. This construction uses logarithmic curve fits to describe the LCF
behavior of the exposed samples 99
Figure 4.35. SEM backscatter image of a sample exposed for 100 cycles at 650°C.
The dark region at the top of the photo is an area of high oxygen concentration.
The fracture surface is oriented to the right in this photo 100
Figure 4.36. Optical micrograph of cross section from exposed tensile sample
stretched to 0.7% strain. Notice cracks run the depth of the oxygen rich zone 101
Figure 4.37. Optical micrograph of cross section from exposed tensile sample stretched
to 1.2% strain. Note the crack runs beyond the oxygen rich region 102
Figure 4.38. Uncoated, unexposed and exposed data, showing the failure location and
the calculated a„ 104
Figure 4.39. Oxygen concentration vs. depth from surface for the exposed fatigue
samples 105
Figure 4.40. Si02 coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material 107
Figure 4.41. Ta,Os coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material 108
Figure 4.42. MgO coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material 109
Figure 4.43. Cr coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material 110
Figure 4.44. Depth of oxygen concentration in Si02 coated LCF samples exposed in
air and argon 111
Figure 4.45. Depth of oxygen concentration in Ta205 coated LCF samples after
exposure in air and argon 112
xv

Figure 4.46. Depth of oxygen concentration in Cr coated LCF samples exposed to air
and argon 113
Figure 4.47. Cr concentration depth in the Cr coated LCF samples exposed in air and
argon 114
Figure 4.48. SEM backscatter image of Cr coated LCF sample exposed in air for 100
cycles at 650°C 114
Figure 4.49. Coefficient of thermal expansion comparison between coatings and Ti-
22Al-26Nb substrate 115
xvi

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
EFFECT OF COATING AND EXPOSURE ON THE OXIDATION AND
MECHANICAL PROPERTIES OF Ti-22Al-26Nb
By
James Ross Dobbs
December 1997
Chairman: Michael J. Kaufman
Major Department: Materials Science and Engineering
High temperature titanium base alloys are being developed to increase the thrust-
to-weight ratio in aircraft turbine engines. Conventional titanium alloys, such as TÍ-6A1-
2Sn-4Zr-2Mo (wt%), can only be used to 540°C. Advanced titanium aluminide alloys
based on the a2 phase, such as TÍ-24A1-1 INb (at%), and the orthorhombic phase, such as
Ti-22Al-26Nb (at%), have better high temperature tensile and creep properties than the
conventional alloys but have lower ductility. In addition to having low intrinsic ductility,
the alloys are embrittled by high temperature exposure in air.
In this study, the cyclic oxidation rate of Ti-22Al-26Nb (at%) at 650°C was
determined, and several classes of coatings, including NiCrAlY+Al203, FeCrAlY, Pt, Cr,
Al3Ti, Ti5Si3, Si02, Taj05 and MgO, were evaluated to determine their effect on the
oxidation rate. Baseline low cycle fatigue life was established, and the effect of exposure
in air at 650°C on LCF life was determined. Based on oxidation and microhardness,
xvn

coatings of Cr, Si02, Taj05 and MgO were selected for a detailed study of the effect of
exposure in air at 650°C on the low cycle fatigue life of coated and uncoated material.
It is shown that exposure in air at 650°C reduced the LCF life of Ti-22Al-26Nb by
as much as 2 orders of magnitude for a 100 cycle exposure. A single cyclic exposure
reduces LCF by 1.5 orders of magnitude. No coating evaluated reduced this degradation,
and in the case of Si02 and MgO, the LCF life was worse after exposure than when no
coating was applied. It is shown that the observed reduction in LCF life is associated
with an increase in oxygen content and microhardness measured in the near surface
regions of the oxidation samples after thermal cycling. The oxide coatings were reduced
by the substrate either during coating or during exposure. The Cr coating allowed oxygen
to diffuse to the substrate.
These results indicate that further work should study those coatings that would not
be chemically reduced by titanium and that would not allow oxygen to diffuse through
the coating to the substrate.
18

CHAPTER 1
INTRODUCTION
Titanium base materials are used in a wide variety of applications. The medical
and chemical industries use titanium because of its resistance to corrosion and its
biocompatability, the sporting goods industry uses it because it is light weight and stiff,
and the aerospace industry uses titanium alloys for their high strength-to-weight ratio at
elevated temperatures. The use of titanium alloys in gas turbine engines is the most
demanding since the material is subjected to a corrosive environment at elevated
temperatures and, in some cases, high stresses that are frequently cyclic in nature. Thus,
titanium base alloys are attractive for use in gas turbine engines because of their low
density compared to Ni-base alloys (4.5 gm/cm3 vs 8.9 gm/cm3) and their high strength at
elevated temperatures [Pos92],
Titanium alloyed with aluminum and vanadium, Ti-6wt%Al-4wt%V, has been an
industry standard for many years [Woo72], Below 1000°C, this alloy consists of a strong
hep a-phase precipitated in a softer bcc (3-phase matrix, producing a stable two-phase
microstructure of hep a-phase and bcc (3-phase [Woo72]. The (3 phase is stabilized by the
addition of the vanadium whereas the addition of aluminum stabilizes the a phase. In
binary Ti-Al, single phase a is stable up to aluminum levels of approximately 11 atomic
percent. At greater aluminum levels, the disordered a phase is in equilibrium with the
ordered hep Ti3Al (a2) phase [Mur87], Single phase a2 is stable from 22 to 35 atomic
1

2
percent aluminum at 500°C. The substitution of niobium (also a p stabilizer) for titanium
in the a2 Ti3Al alloys leads to a better balance of properties, i.e. improved creep
resistance, fracture toughness and oxidation behavior. Initial work on the Ti-Al-Nb
system was carried out by McAndrew and Simcoe [McA60] who studied alloys with up
to 20wt% A1 and 30wt% Nb. More recent work conducted by Blackburn [Bla78]
indicated the promise of the Ti3Al class of materials as engineering alloys. The specific
tensile yield strengths of the Ti3Al + Nb alloys as a function of temperature are superior
to that of conventional titanium alloys, and even better than cast and wrought IN718
[Pos92] as shown in Figure 1.1. Several engine components have been produced from
Ti3Al alloys, including afterburner nozzle seals and high pressure compressor casings for
the F100 engine and exhaust seals for the F404 engine [Lip85]
Figure 1.1. Specific yield strength vs. temperature comparing an a, and orthorhombic Ti-
Al-Nb alloy with a conventional Ti alloy and IN718. After [Woo93].

3
The discovery by Baneijee et al. [Ban88] and Rowe [Row90] that higher additions
of niobium (above approximately 12 at%) to Ti3Al leads to the formation of an ordered
Ti2AlNb orthorhombic phase (O Phase) introduced another class of materials for high
temperature applications. Alloys based on this O phase tend to have higher toughness
and tensile strength than alloys based on the a2 phase [Row91, Row91a, Row93, Smi93b,
Smi94] while maintaining similar creep resistance. Extensive evaluation of this class of
materials has been conducted in recent years [Ban95], The deformation behavior and
dislocation processes have been evaluated [Ban91, Ban92, Ban95a, Dar94, Dou93,
Pop96a, Pop96b], as well as phase transformations and relationships [Ban93, Kau88,
Mur95, Vas96, Wey89],
A major impediment to using titanium-base materials at elevated temperatures in
oxidizing environments is the propensity for oxygen to diffuse into the titanium [Unn86].
The oxygen, being a small atom, sits in the titanium lattice interstices and causes the
titanium alloy to become very hard and thus embrittled [Cha87, Unn86], In the case of
a-p alloys, oxygen is a strong a stabilizing element and leads to a hard surface layer
known as “a-case”. In the 0.2 and O alloys, there is a high solubility for oxygen in the
lattice, but these a like phases are not as readily stabilized by oxygen. Therefore, there is
no discernible change in the microstructure due to oxygen ingression during exposure.
Even so, high oxygen surface regions are embrittled and lead to a reduction in ductility,
fracture toughness and fatigue resistance [Dar95, Dar96, Lip93, Sai93, Sch95]. These
reductions in properties have been a major barrier to the implementation of this class of
materials in gas turbine engines [Pos92],

4
Several attempts to mitigate the environmental embrittlement have been made
with limited success. For example, 012 and O alloys have been coated with a variety of
materials, including MCrAlYs [Bri92, McK93, Sch95], glasses [Dev90, Wie89, Wie91],
silicides [Coc96], aluminides [Kun90, Smi90, Smi93b, Sub88] and Pt [Nic96], These
coatings resulted in embrittlement, presumably due to either the formation of deleterious
reaction phases, or the coating itself was brittle and caused a reduction in low cycle
fatigue (LCF). The process of applying the coating could also contribute to the reduction
in LCF life. Subsequent exposure of these coated alloys to a high temperature oxidizing
environment typically led to negligible further embrittlement over the as-coated condition
presumably because the initial deterioration had already occurred during coating.
The objectives of this research include identifying a potential coating system and
methodology that does not lead to degradation upon application and is protective during
subsequent exposure, and to identify the mechanisms that can cause changes in the
mechanical properties of Ti-22Al-26Nb due to the presence of specific protective
coatings.

CHAPTER 2
LITERATURE REVIEW
Overview of Ti-Al-Nb Metallurgy
History
The development of supersonic jet aircraft after World War II nurtured the need
for high thrust-to-weight ratio aircraft turbine engines. This initiated the search for new
structural alloys that could withstand the high temperatures and stresses induced by a
turbine engine [JafBO], In aerospace applications conventional aluminum alloys are only
useable up to approximately 150°C. Conventional titanium alloys like TÍ-6A1-4V (wt%)
can be used at temperatures exceeding 300"C. Today’s commercial alloys have the
strength capability to operate at temperatures approaching 600°C. Their oxidation
resistance is not yet on the same level as the mechanical properties and, consequently,
coatings have been developed that mitigate the poor oxidation behavior [Eyl84], In the
early stages of titanium alloy development, it was recognized that the addition of
aluminum would increase the tensile and creep strength as well as the elastic modulus
[Ely84], Much work was conducted in subsequent years to develop an accurate Ti-Al
phase diagram and in 1987 Murray [Mur87] published a version of the diagram based on
much of this work. This diagram is shown in Figure 2.1 and includes all of the phases
that form with the increasing aluminum.
5

6
Weight Percent Aluminum
Figure 2.1. Titanium - aluminum equilibrium phase. After [Mur87].
Based on the phase diagram and empirical studies, it was determined that the
amount of aluminum, or its equivalent, should be kept below 9 at% to remain in the a
phase region. Higher levels of aluminum led to the formation of the intermetallic Ti3Al
and a reduction in ductility. In the 1970s it was recognized that the Ti3Al and TiAl class
of materials could possibly be exploited as structural engineering alloys. Ti3Al, also
known as a2, was developed by the Air Force [Lip81] and at the aircraft engine
manufacturers. It was recognized that the addition of Nb to Ti3Al improved the ductility,
creep behavior and oxidation resistance [Bla78, Sas77], and a new class of Ti,Al alloys
was launched. At the lead of these materials was the ternary alloy TÍ-24A1-1 INb (at%).
The addition of vanadium and molybdenum to a2 improved properties even more, and
super a2, or Ti-25Al-10Nb-3V-lMo, was developed. The addition of Nb at levels above

7
about 15 at% causes the Nb atoms order on the titanium sublattice leading to a distortion
of the ordered hexagonal cell into the orthorhombic phase [Ban88], The orthorhombic
alloys are reported to have superior tensile strengths over the a2 alloys (Figure 2.2), better
oxidation resistance and comparable creep strengths. [Row90, Row91, Row91a, Row92,
Row93, Smi93, Smi94], It is this phase that is the focus of this study.
Figure 2.2. Comparison of the specific yield strength of a2 and orthorhombic alloys as a
function of temperature. After [Row91 ].
Description of Phases Present in Ti-Al-Nb
As seen in Figure 2.1, there are five distinct phases present in the Ti-Al system.
Four of these p, a, Ti3Al or a2, and TiAl or y are interesting engineering alloys. Al3Ti has
shown promise as a coating for titanium alloys and will be discussed in a later section.
The P phase is a body centered cubic (bee) structure and is stable at higher temperatures.

On cooling through the p transus, this phase undergoes a transformation to the a phase.
The a phase is a hexagonal close packed (hep) phase consisting of a random arrangement
of titanium and aluminum atoms. As seen in Figure 2.1 the low temperature solubility of
aluminum in titanium is approximately 10 at%; beyond this amount a 2 phase mixture of
a and Ti3Al, and ordered structure known as a2, exist in equilibrium. Aluminum
additions over 22 at% will create single phase a2. The basal plane of the a2 lattice is
depicted by a ball model in Figure 2.3.
Figure 2.3. Ball model of basal plane of a2 showing lattice sites for Ti and A1 atoms.
The a phase lattice is similar, with the Ti and A1 atoms arranged randomly.

9
The addition of Nb to Ti3Al ast levels beyond 10 at% results in either a 2 phase
(a2 + B2 or O + B2) or 3 phase (a2 + O + B2) structure as will be discussed further
below. A partial 900°C Ti-Al-Nb isotherm (Figure 2.4) indicates that the maximum
solubility of Nb in a2 is 15 at%, and above 10 at% Nb the a2 is in equilibrium with B2 up
to 15 at% where the a2 and B2 are in equilibrium with the O phase. The B2 phase is an
ordered bcc phase formed from the (5 phase. The addition of Nb causes the p phase to
order with Ti atoms occupying one site and the A1 and Nb atoms randomly occupying the
other [Ban87] as shown in Figure 2.5.
Ti
Ti-Al-Nb Ternary
Figure 2.4. 900°C isotherm from the Ti-Al-Nb ternary phase diagram. B2, a2, and
orthorhombic phases are in equilibrium. After [Row92],

10
As seen in Figure 2.4, there are three important phases of interest in the Ti-Al-Nb
system at 900°C, namely, a2, B2 and O. A ball model of the basal planes of the a2 and O
phases is shown in Figure 2.6. The orthorhombic unit cell is derived by a slight distortion
of the a2 unit cell caused by the ordering of the Nb atoms on the Ti sites [Kon86], For
complete substitution, it is clear that the O phase is an ordered ternary phase based on the
composition Ti2AlNb.
Figure 2.5. Ball model drawing of the B2 phase showing the relationship between the
titanium, aluminum and niobium atoms.

11
Figure 2.6. Comparison of the basal planes of the a2 and orthorhombic phases.
Szaruga et al. [Sza92] reported that oxygen is a strong a2 stabilizer in TÍ-25A1-
10Nb-3V-lMo having an effect on both the (3 transus temperature and the p/B2 order-
disorder transition. Oxygen is also an important alloying element in the O phase. An
alloy of composition Ti-22Al-23Nb containing above 1000 wppm oxygen is made up of
the two phases, O and B2 [Lip93]. This indicates that oxygen stabilizes the a, phase, much
as oxygen stabilizes the a phase in conventional titanium alloys. Rhodes et al. [Rho93]
reported the same effect in the alloy Ti-22Al-27Nb, but at oxygen concentrations of
approximately 1120 wppm the equilibrium phases were still O and B2. Therefore, at
higher Nb contents it takes more oxygen to stabilize the a2 phase. Ward [War93]
concluded that interstitial atoms such as oxygen can be potent strengtheners, but have an
adverse effect on ductility by limiting the number of slip systems available.

12
A cut through the ternary phase diagram shown in Figure 2.4 along the constant
25 at% A1 line is shown in Figure 2.7. The (5 to B2 ordering line is shown above 1100°C
for compositions around 12 at% Nb.
Nb (at %)
Figure 2.7. Constant 25 at% A1 isopleth from the Ti-Al-Nb ternary in Figure 2.4. After
[Ban95],
Also shown in Figure 2.7 is the range of transformations that Ti3Al + Nb alloys
experience. The transformations associated with continuous cooling of high Nb alloys
(>16 at%) from the oc2 + B2 region is complex. The precipitation of az or O laths follows
the Burgers relationship (0001)a21| (110)(3 ; [112 0]a2 || [111](3 and the equivalent

13
relationship for the O phase (001)0 I (110)P, [110]O || [11 l]p [Ben91], The
transformation from a2 to O phase results in a lamellar or mosaic structure [Ban88],
Microstructures obtained during thermomechanical processing of alloys with high Nb
contents depend on the processing temperature. Processing in the B2 + a2 or O phase
field will yield, on cooling, a mix of primary a, or 0 phase surrounded by a
Widmanstatten a2 or O + B2 mixture. Processing in the B2 field results in a 100%
Widmanstatten a2 or O laths with B2 in the interlath regions. Examples of these two
microstructures are shown for the alloy used in this study, Ti-22Al-26Nb in Figure 2.8
and Figure 2.9 respectively.
Figure 2.8. SEM Backscatter micrograph of Ti-22Al-26Nb processed in the O + B2 field
and directly aged at 815°C for 4 hours. After [Woo93],

14
Figure 2.9. SEM Backscatter micrograph of Ti-22Al-26Nb processed in the B2 field and
directly aged at 815°C for 4 hours. After [Woo93],
Deformation Behavior of ct?, O and B2 Phases
Dislocation arrangements in deformed a2 alloys [Akk91, Ker84, Kos90, Lip 80,
Lip85] and O alloys [Ban91, Ban92, Ban95a, Dou93, Pop96b, Pop96a] have been studied
in depth. This includes deformation in the B2 phase contained in both materials.
The a, phase deforms by slip [Ban90] in two distinct systems, (l 120)(0001)
and (l 120){1010}, and by slip in the (l 126){1121} system [Kos90], The
crystal structure and Burgers vectors for the a, phase are shown in Figure 2.10 along with
those for the O phase [Ban95], slip has been observed in a2 alloys but are associated
with the p to a2 transformation rather than tensile deformation [Ban95]. The lack of
slip and the dependence of a, alloys on
and slip create plastic
incompatibility and thus the a2 alloys are considered brittle at room temperature.

15
The O structure is similar to the a2 phase with a slight distortion due to ternary
ordering. Figure 2.10 compares the Burgers vectors for the O and a2 phases, and it can
be seen that the three 1/6 (l 120/ vectors in a2 are not equal in the O phase in either
magnitude or generation of APBs. Baneijee et al. [Ban91] have shown that the
deformation associated with
slip in O is very similar to that in a2, but unlike a2 there
is considerable slip in the O phase alloys and these dislocations are arranged into
well defined slip bands [Kos90]. This would suggest that the O phase alloys should
possess more plasticity at room temperature and behave better under fatigue crack growth
conditions.
Figure 2.10. Crystal structure and burgers vectors for the a, and O phases. After
[Ban95]
Deformation in the B2 phase occurs predominantly by (l 11) slip on the {110},
{121} and {123}planes [Ban90]. This deformation is extremely inhomogeneous and
localized into slip bands as shown in Figure 2.11. This localization could result in the
formation of persistent slip bands and contribute to the formation of crack nucleation sites
during fatigue [MÍ197].

16
Figure 2.11. Slip bands in the B2 phase in TÍ-24A1-1 INb showing the inhomogeneous
nature of slip. After [Ban90].
Tensile Behavior of a, and 0 Alloys
The discovery that the Ti3Al+Nb alloys have higher tensile ductility than the
binary Ti3Al alloys [Sas77] indicate that Ti3Al base alloys, including the O phase alloys,
have to be alloyed sufficiently to stabilize the (3 or B2 phase with Nb as the preferential p
stabilizing element. The B2 phase delays the cleavage cracking of the a, or O laths to
higher strains by its ability to accommodate the plastic instabilities associated with the
lack of slip systems in a2. This typically occurs at the a,/B2 interfaces with its large
number of available slip systems [Ban95],
Recall that the decomposition of the B2 phase into the a2 or O phase when heat
treated above the P(B2) transus results in a lath type microstructure as shown in Figure

Yield Strength (MPa)
17
2.9 where the a2 or O laths have the Burgers relationship with respect to the B2 phase.
The cooling rate from the p(B2) transus strongly affects the strength, as shown in Figure
2.12, and the ductility, as shown in Figure 2.13. The yield strength increases continually
with increasing cooling rate while the ductility goes through a maximum.
Figure 2.12. Yield strength vs. cooling rate for two oc2 alloys. After [Gog90],

18
5.0
4.0
g 3.0
c
_o
00
a
o
“ 2.0
1.0
0.0
0.01 0.1 1 10 100
Cooling Rate (C/sec)
Figure 2.13. Elongation vs. cooling rate for two a2 alloys. After [Gog90].
The increasing cooling rate causes a finer size lath to be formed on cooling from
the P(B2) phase field. This finer lath arrangement yields a higher yield strength due to
the Hall-Petch relationship. The low ductility at low cooling rates results from the
formation of similarly oriented lath colonies which allow cleavage to occur across many
laths with no change in crack energy [Ban90], The reduction in ductility at high cooling
rates is associated with the reduction in scale of the microstructure and the reduction in
volume fraction of the p phase. At high cooling rates there are also a2 and O laths that
have nucleated from the grain boundaries. These grain boundary nucleated laths have a
similar orientation and thus a crack that forms in one can easily propagate into the others
â–¡ 24-15 El
O 24-11 El
o
o
-s-
O
â–¡
o
â–¡
o
â–¡

19
with little energy loss. Therefore, the optimum lath structure is a fine basketweave of the
a2 or O phase without any grain boundary film or grain boundary initiated lath colonies.
An example of a colony of O laths is seen in Figure 2.14 where the lath colony in the
lower right hand comer has formed sympathetically from the larger O lath running from
bottom right to top left. This is a micrograph of the Ti-22Al-26Nb alloy used in this
study.
S,3kV,SE@17mm,10kX,4844S14.ras 2 microns/div.
Figure 2.14. SEM backscatter micrograph of Ti-22Al-26Nb forged in the B2 region and
directly aged at 804°C. Notice the colony of O laths that have formed off of the large
lath.

20
The thermomechanical processing history of the material can also play a role in
the macroscopic tensile behavior. Material that has been forged and directly aged has
higher tensile strength than material that has been forged, solutioned above the (3(B2)
transus and then aged [Dob94], This is again due to the failure of the a2 + (3 at the a2
film formed along the grain boundary interface due to strain incompatibility [Cha90], as
seen in Figure 2.15.
Figure 2.15. Intergranular crack observed in tensile tested Ti-22Al-27Nb solutioned
above the P(B2) transus and aged at 870°C for 50 hours.
The high grain aspect ratio (GAR) left over from forging in the p forged + direct
aged material forces this failure to occur on boundaries that are parallel to the stress axis
as seen in Figure 2.16, as opposed to a few boundaries that are perpendicular to the stress
axis, as seen in Figure 2.17. The p forged and directly aged material also did not contain

21
the grain boundary a2 film, and the lath size was much finer [Dob94], This is indicative
of faster cooling rates from the p region.
Figure 2.16. Fractograph of forged + direct aged material. Notice the high GAR and the
initiation of failure from an internal boundary.

22
Figure 2.17. SEM fractograph of material solutioned above the P(B2) transus and aged at
870°c for 50 hours. Notice that there is mixed intergranular and transgranular fracture
across the sample and the prior P grain size is approximately 500pm.
Fracture Toughness and Fatigue
Chan determined [Cha92] that the crack growth process in a coarse basketweave
structure of TÍ-24A1-1 INb (at%), one where the laths are approximately 20 pm long and
5 pm wide, was by decohesion of slip bands, similar to the mechanism for equiaxed a2.
The tips of the microcracks terminated in the p phase, as did the tip of the main crack. In
contrast, for a fine basketweave structure where the lath size is approximately half that of
the coarse laths and where the p phase is not as continuous, the cracks propagated around
and through the P phase. Therefore, Chan concluded that the room temperature fracture
toughness of TÍ-24A1-1 INb (at%) was imparted by the p phase inhibiting microcrack
nucleation by both relaxing the incompatibility strain in the a2/p interfaces and by
blunting the crack tips.
Penton et al. reported that Super a2 (Ti-25Al-10Nb-3V-lMo (at%)) containing
primary a2 exhibits faster crack growth rates than p solutioned and aged material
[Pen93], This effect appears to result from early cleavage of the brittle primary [Pen92, Pen93], Ward reported that the a2 laths appeared to cleave on the basal plane
[War93b] and that the crack is then bridged by the p phase [Gog90, Luk90], Others have
reported that the cracks also propagate along the a2/p interfaces in transformed regions
[Luk90, Tak96], Davidson et al. [Dav91] found that, for Super a2 rolled below the p
transus and then aged, small cracks were always found in the a2 phase after deformation.
These small cracks would then grow below the AK,,, for large cracks. Ravichandran and

23
Larsen [Rav92] concluded that growth rates of small cracks in basketweave structures of
TÍ-24A1-1 INb were consistent with large crack growth rate data. Penton et al. [Pen93]
concluded that the growth rate of solutioned and aged material is in line with the fracture
of a2 laths.
Miller [MÍ197] has defined three distinct conditions for initiation of fatigue:
1) Fatigue limit of polycrystalline material which is related to some limiting
microstructural feature such as grain size, primary precipitate size or brittle
precipitate size.
2) The mechanical stress state applied to a pre-existing crack or flaw.
3) The condition for a single crystal in which a defect, such as intrusion or
extrusion, must be induced.
Miller concludes that the fatigue crack is initiated on the first cycle by one of the
above conditions. In the case of a2 and O alloys with limited slip systems in the lath
structures, the first condition of some limiting microstructural feature would apply. This
feature would most likely be a primary a2l O particle or an a2l O lath, in agreement with
the observations described above.
Fatigue is typically described as consisting of the following stages: [Sur91]
1) Nucleation of permanent damage through microstructural changes;
2) Creation of microscopic cracks;
3) Coalescence of microscopic cracks to form measurable cracks;
4) Stable growth of a preferred crack; and
5) Ultimate catastrophic failure.

24
In classic fracture mechanics, Stage 3 indicates the end of initiation and the
beginning of propagation of the fatigue crack. Considering this model of fatigue, linear
elastic fracture mechanics (LEFM) can be utilized to help understand the behavior of
material by calculating both the initial flaw size, a„ and the size of the crack at failure, a,.
The equation for the stress intensity, K„ at the notch of a sharp crack is
1.12 crV/ra
K,= (2.1)
9
hK k a2 . , , n
where 6 = 1 • —r. In the special case when c = a, the m term is reduced to — the
Y 8 8 cz 2
equation then becomes
K, = 0.71 crVia (2.2)
This is the case for a penny shaped flaw that is very small relative to the size of the test
bar. To determine the size of the fatigue crack at failure, af, one can use the plane strain
fracture toughness, KIC for K„ and in the case where A=l, the maximum stress is equal to
A a. Solving equation 5.2 for af gives
Klr , 1
af= ( lc )2.~ (2.3)
0.71er,,,,. k
da
Paris showed that the fatigue crack growth increment — is related to the stress intensity
dN
factor as
da
dN
= C(AK)ra (2.4)
da
where C and m are scaling constants derived from the — vs. AK curve. When equation
dN
2.4 is integrated from an assumed initial flaw size a„ to the critical crack size af, the

25
number of cycles to failure can be calculated. When the stress intensity factor for a small
crack is defined as in equation 2.2, equation 2.4 becomes
= C(0.71AcrVia y1 (2.5)
where cr has now become Act, defined as crmax - omill. For our case where we are assuming
that the stress intensity factor is not a function of the crack depth a, equation 2.5
integrates as
0
The resulting fatigue life is
forn * 2.
Comparison of the fatigue crack growth data for TÍ-24A1-1 INb (at%) generated
by Ravichandran and Larsen [Rav92] to that for Ti-21Al-25Nb (at%) generated by
Woodfield et al. [Woo94] with a similar basketweave structure indicates that the TÍ-21A1-
25Nb is slightly better. This could indicate that the O + B2 alloys would behave better in
fatigue, probably due to the increase in the number of active slip systems and the increase
in Nb content in the B2 phase. Balsone et al. [Bal93] evaluated fatigue crack growth in
Ti-25Al-25Nb (at%) with a microstructure consisting of primary a2 + O laths surrounded
by the B2 phase. Comparison of dA/dN vs. AK data for TÍ-24A1-1 INb (at%) with Ti-
25Al-25Nb with a similar microstructure [Rav92], indicates that the latter has a slightly

26
higher crack growth rate. This is probably due to the scale and volume fraction of the
primary a2 or O and the effect of the ct2/p(B2) interface.
Overview of Oxidation Behavior
Oxidation of Uncoated Titanium Alloys
The oxidation behavior of titanium base alloys has been extensively studied since
the late 1940s, starting first with unalloyed titanium. Initially there were large
disagreements between investigators about rate laws and rates of oxidation [Kof58],
More recent investigations of commercially pure titanium by Unnam et al. have shown
that the weight gain due to oxide growth and oxygen dissolution is essentially parabolic
with respect to time [Unn86], Unnam et al. also reported that the oxygen diffusion
coefficient in Ti(O) is independent of oxygen concentration in the 1-10 at% range, and
the effective solubility limit of oxygen in pure titanium is 20 at%. They also concluded
that the diffusion coefficient of oxygen in Ti02 is about 50 times that of oxygen in the
metal. Chuanxi and Bingnan reported that the addition of Nb to pure Ti improved the
oxidation resistance by improving the surface stability of the oxides [Chu92]. Chaze and
Coddet [Cha87] conducted oxidation experiments on Ti with additions of Al, Si and Cr.
They concluded that the addition of Al up to 16.5 at% and Si up to 1.5 at% improved the
oxidation resistance of Ti, while Cr additions up to 18 at% had little or no effect. In
contrast, Kahveci and Welsch evaluated the effect of Al on oxidation of Ti [Kah87] and
concluded that at least 25 at% Al is necessary for any significant improvement in the
oxidation behavior of Ti-Al alloys and that there is no appreciable effect for the addition
of Al until about 13 at%, beyond which the oxidation behavior improves with increasing

27
Al content. Kahveci et al. [Kah88] showed that the kinetics for oxide growth for TÍ-25A1
(at%) fall between Ti02 and A1203 but are closer to the kinetics for Ti02. Qiu et al
[Qui95] reported that the addition of 11 at% Nb or 5 at% Si can improve the oxidation
resistance of Ti2Al and that the effect of adding both is even greater. However, no
continuous A1203 scale was formed after any additions.
In a subsequent paper, Welsch and Kahveci evaluated 3 binary Ti-Al alloys with
increasing Al content and 1 Ti-Al-Nb alloy[Wel89]. They observed that the parabolic
rate constant decreased and the thickness of the oxide scale decreased with increasing Al.
They concluded that the outer oxide scale in the binary alloys is a Ti02 with A1203
channels, the intermediate layer consists of Ti02, A1203 and porosity, and the inner oxide
layer is A1203 with Ti02 and porosity. The alloy surface has A1203 lamella from internal
oxidation. On the other hand, the Ti-Al-Nb alloy had a dense outer oxide scale consisting
mostly of A1203 and Ti02, an intermediate layer of Ti03 and Nb205; and the layer next to
the metal was an Nb205-Ti02 layer with porosity at the interface. There was no visible
internal oxidation but there was an oxygen rich zone adjacent to the oxide scale.
Chromium has been found to help improve the ability to form a protective scale in
Ti-Al-Nb alloys and is an excellent ft stabilizing element [Doy95], The Cr tends to alloy
with the oxide scale and provide a more protective scale. This indicates that
improvements in oxidation are possible by alloying, and Cr is a good alloying addition.

28
Oxidation of Coated Titanium Alloys
Coatings for titanium base alloys have been evaluated based predominantly on
their ability to mitigate oxidation. The coatings for Ti alloys fall into four broad classes,
AljTi, MCrAlY, pure metals, and glasses or oxides.
Aljli
Al3Ti has been produced as a coating by a variety of methods including (1) laser
surface alloying [Abb93] where A1 powder is either fed into the molten pool or placed on
the substrate as a slurry and melted in by laser [Gal92], (2) dipping the substrate of
interest into a molten bath of A1 and allowing the reaction to take place [Abd91] (3)
deposition of A1 by EBPVD and subsequently reacting with the matrix during exposure in
air [Unn85], (4) deposition by sputtering and reacting in vacuum prior to exposure in
air[Wie93], and (5) pack cementation [Kun90, Smi90, Smi93a, Sub88] where the
substrate is placed in a mixture of alumina powder, aluminum powders and halide salt
activators and reacted at 1000°C where the A1 will volatile and produce a coating of A1 on
the substrate, which is then reacted to form Al3Ti. In all cases the coating Al3Ti was
cracked on cooling from the reaction temperature. Even so, the coatings tended to impart
improved oxidation resistance despite being cracked, although oxidation did occur down
the cracks and affect the substrate [Kun90]. It was concluded that the thicker the coating,
the more susceptible it was to cracking [Smi90]. McMordie reported that the addition of
Si to the Al3Ti coating improved oxidation even more, possibly due to the formation of a
silicide at the coating/substrate interface [McM91],

29
MCrAlY
MCrAlY type coatings, where M is a transition metal like Ni, Fe or Co, are
typical coatings used in the gas turbine industry. The application of these coatings to Ti
has been attempted by a number of researchers, and several patents are held on the
materials and processes [Bri92, Lut90, Tob92], The preferred method of producing this
type of coating is plasma spray [McC90, McK93, McK93a, Sch95], For Ti3Al + Nb
alloys, the MCrAlY is sometimes mixed with an oxide powder such as A1203 in order to
reduce the thermal expansion of the coating to better match that of the substrate [Sch95].
In all cases, the MCrAlY coatings improve the oxidation resistance of Ti substrates, such
as Ti-6Al-2Sn-4Zr-2Mo (wt%) [McC90]. There is a slight increase in hardness beneath
the coating in the as-coated condition and this is attributed to oxygen ingress during the
pre-heating in partial vacuum prior to the plasma deposition.
Pure metals
The evaluation of pure metals as oxidation resistant coatings for Ti base materials
has mostly been limited to Pt and Cr. The preferred method for deposition of Pt is ion
plating [Eyl84, Eyl85, Fuj79] and for Cr is PVD processes such as sputtering [McK90],
Eylon et al. showed that the Pt coatings reduced the surface oxidation rate of Ti-6Al-2Sn-
4Zr-2Mo (wt%) by three orders of magnitude[Eyl85] and did not degrade the high cycle
fatigue properties [Fuj80], PtAl2, produced by multi-layer sputtering of Pt and Al and
subsequent reaction, has been evaluated as a coating for a, type alloys by Nicholls et al
[Nic96]. They reported that a 3pm thick coating protected then a2 alloy MT754 (Ti-
23Al-9Nb-2Mo-0.9Si at%) from oxygen ingress at 700 and 800°C, and extended the

30
creep life of the alloy by a factor of two under conditions of 350MPa at 650°C. They also
reported that after exposure in air for 100 hours at 700°C, there was no hardening under
the PtAl, coating.
Glasses, silicates and oxides
Oxides such as Si02, Y203, MgO, A1203 and Zr02 have been deposited on Ti-base
alloys by a variety of techniques. Wiedemann et al. [Wie89, Wie91] deposited Si02,
A1203 and B203 by a sol-gel process, and Y203, MgO, Zr02 and HfO by sputtering onto
TÍ-24A1-1 INb (at%). They reported that most coatings had poor integrity and allowed
oxygen to diffuse through. Of all the coatings they tried, the MgO based coating made
from the sol-gel process behaved the best after 1 hour at 982°C in air. They concluded
that coatings applied by sputtering were not effective for oxidation protection due to their
poor integrity. In contrast, Clark et al. [Cla88] deposited Si02 by sputtering and reported
that there was a significant reduction in weight gain at 700°C after 25 hours in air.
DeVore and Osbome [Dev90] coated Ti3Al with borosilicates and aluminosilicates and
concluded that all coatings showed some improvement in oxidation, but the thin sol-gel
coatings were the worst while a borosilicate coating airbrushed onto the sample exhibited
the best behavior. Bedell et al. [Bed91] utilized ion beam assisted deposition to deposit a
double coating of silicon nitride and chromium on the conventional Ti alloy IMI829.
They concluded that the coating provided protection from oxidation at 700°C for 100
hours, but the Si3N4 diffusion barrier is not stable with titanium. Cockram and Rapp
[Coc96] produced silicide coatings on a2 and O phase alloys by pack cementation and
reported excellent oxidation resistance up to 1000°C.

31
Embrittlement of Titanium Base Alloys
As mentioned previously, titanium base alloys suffer from embrittlement when
exposed to air at high temperatures. This embrittlement is due to the ingress of oxygen
and the subsequent increase in hardness in the affected layer [Sha68], The volume
fraction of the beta phase in these near surface regions in TÍ-6A1-4V (wt%) is reduced at
all temperatures when exposed in air [Kah87]. The P transus temperature is raised by the
presence of oxygen in TÍ-6A1-4V as described by the relationship T„[°C] = 937 + 242.7 x
[O wt%] [Kah87], The depth of oxygen penetration, and thus the depth of hardening due
to interstitial ingress, is a parabolic function [Kah86, She86] thus indicating that the
hardening is due to diffusion of oxygen into the substrate. Wallace conducted tensile
tests on Ti-1100 (Ti-6Al-2.75Sn-4Zr-0.4Mo-0.070-0.02Fe wt%) foils after exposure in
air at 600°C and 800°C for up to 1000 hours. He concluded that the elongation was the
most sensitive parameter to minor oxidation, but tensile strength was also affected by
long exposures. He also observed a band of brittle fracture at the metal surface due to
oxygen diffusion into the metal. This band is high in oxygen and made up of the a phase
[Wal95],
Balsone studied the tensile properties of TÍ-24A1-1 INb (at%) after exposure in air
for 10 and 100 hours [Bal89], He reported that an annulus of embrittled material was
formed around the surface of an exposed sample that cracked during loading. These
cracks served as notches for premature failure during testing such that both the elongation
and UTS were reduced. Embrittlement of Ti-25Al-10Nb-3V-lMo (at%) was also studied
by Saitoh and Mino [Sai93] after exposure at 700°C for 100 hours. They also reported

32
that the elongation and UTS was reduced, as did Meier and Pettit [Mei92] for a Nb
modified a2 alloy after a 24 hour exposure at 900°C. Other investigators have evaluated
the effects of embrittlement on fatigue behavior. Godavarti et al. [God92] reported that
the LCF life of Ti-25Al-10Nb-3V-lMo (at%) was reduced by 2 orders of magnitude after
short exposure times at 720°C. They attributed this reduction to the formation of a brittle
layer at the surface of the test samples. They also reported that the same pre-exposure in
air had no effect on high cycle fatigue (HCF) life due to the low strains imposed during
F1CF testing. Praida and Nicholas [Pra92] found that fatigue crack growth resistance in
TÍ-24A1-1 INb (at%) was strongly effected by testing in air at 650°C. Hold time fatigue
crack growth rates were the most effected by elevated temperature testing, indicating an
effect of the environment on cracking. Balsone et al. [Bal93] determined that TÍ-25A1-
25Nb exhibited time-dependent crack growth behavior when tested at elevated
temperature. They reported that at 650°C and 750°C, there is a contribution from
environmentally assisted crack growth.
Dary et al. evaluated the effect of exposure in air and in vacuum on the tensile
properties of Ti-22Al-23Nb [Dar96, Dar94, Dar95], They observed a decrease in UTS
and ductility in all samples exposed in air and, although there was no loss of mechanical
properties in the vacuum exposed samples, the external surfaces contained a large number
of small cracks. Although Rhodes et al. [Rho93] observed a2 in Ti-22Al-23Nb at oxygen
contents greater than 0.098wt%, Dary and Pollock did not observe any a2 precipitates in
the high oxygen regions when exposed at the same temperature. They supposed that the
760°C exposure was too low to cause a2 precipitation or that the B2 + O phases were

33
compositionally stable at this temperature. Chesnutt et al. [Che93] conducted tensile tests
on Ti-22Al-27Nb after either a thermal exposure (no environment) at 650°C or an
environmental exposure at 590°C in air. The results indicate no by the thermal exposure
but significant embrittlement after the environmental exposure. No evaluation was
conducted on the environmentally exposed samples to establish the cause of
embrittlement.
Effects of Coatings on Mechanical Properties of Ti Alloys
Coatings to prevent the degradation of mechanical properties in Ti alloys have not
received as much attention as coatings to mitigate the poor oxidation behavior. As
described above, Ti base alloys are embrittled after exposure in air at elevated
temperatures. To prevent or at least reduce this effect, a series of evaluations have been
conducted to determine the best coatings and the effect of these coatings on the
mechanical properties in both the as-coated and exposed conditions. Fujishiro and Eylon
demonstrated that the HCF life of Ti-6Al-2Sn-4Zr-2Mo (wt%) was improved at 455°C
with the application of a platinum coating [Fuj80]. For the a2 alloy Ti-24.5Al-12.5Nb-
1.5Mo (at%), Schaeffer and McCarron evaluated NiCrAl + 40 vol% A1,03 deposited by
plasma spray [Sch95], They reported that the LCF of as coated samples was
approximately the same as for the uncoated and exposed samples where the LCF life was
reduced by three orders of magnitude. Subsequent exposure of the coated sample did not
degrade the LCF life any further. McKee reported that for the same alloy, an underlayer
of chromium or tungsten prior to deposition of the NiCrAl would eliminate this
embrittlement [McK93], Fie based this conclusion on microhardness profiles beneath the

34
coating, rather than on mechanical property data. Chesnutt et al. evaluated the LCF
response of the O alloy Ti-22Al-27Nb (at%) with and without a coating [Che93] and
concluded that the LCF life was degraded by the coating in a manner similar to that seen
by Scaeffer and McCarron [Sch95],
Program Approach
While there have been various studies addressing the oxidation of coated O phase
alloys and the effect of coatings on mechanical properties, to date there has not been a
systematic study addressing both issues. Such an study will be critical to the use of this
class of alloys in high temperature applications in air, particularly in static applications
with large thermal gradients where the stresses developed by the thermal gradients can be
large and the number of thermal cycles small, therefore subjecting a static part to low
cycle fatigue.
The approach of this program is to evaluate the oxidation, tensile and LCF
behavior of O phase alloy Ti-22Al-26Nb in both the uncoated and coated conditions. The
material will be taken from a forged disk, and the baseline properties of the material will
be determined, including the effect of location in the forging. Coatings will include
oxides and pure metals, and are intended to be barriers for oxygen diffusion and diffusion
of any topcoat that may be identified.

CHAPTER 3
EXPERIMENTAL PROCEDURE
Material Processing
A 6100 pound ingot of target composition Ti-22Al-26Nb (atom percent) was
produced at Teledyne Allvac/Vasco. The material was produced by an initial plasma arc
melting (PAM) into a 43 cm (17 inch) diameter ingot, followed by vacuum arc remelting
(VAR) into a 53 cm (21 inch) diameter ingot. This 53 cm (21 inch) diameter ingot was
sectioned into two pieces. The top section was cogged down from 53 cm (21 inch) to 41
cm (16 inch) diameter billet by heating to 1232°C (2250°F), and using 2.5 cm (1 inch)
drafts. This was followed by hot grinding and cooling to room temperature. This top
half was converted to 20 cm (8 inch) diameter billet and sheet bar for use in other
programs. The bottom half was cogged down from 53 cm (21 inch) to 40 cm (16 inch)
billet using a similar approach as the top half, except there was a 40% reduction initially
and fewer reheats were used. This 40 cm (16 inch) diameter billet was then GFM
converted to a 15 cm (6 inch) diameter billet. This conversion took place at 1010°C
(1850°F), based on beta transus measurements of a slice from the 41 cm (16 inch)
diameter billet. Four billets approximately 2.4 m (94 inches) long and 16 cm (6.3 inches)
in diameter weighing a total of approximately 2000 pounds were produced. Chemistries
were measured from several locations along the length of the ingot and from the surface
and center of these locations. The compositions are listed in Table 3-1 [Woo84],
35

36
Table 3-1. Compositions from several locations along the length of the 15.25 cm
diameter billet, surface and center, and from the forging used in this study.
Billet
Forging
Aim
A1 (at%)
21.35
21.29
22.0±1.0
Nb (at%)
25.75
26.15
26.0±1.0
Ti
Balance
Balance
Balance
Fe (wt%)
0.036
0.06
<0.07
C (wt%)
0.013
0.0063
<0.05
H (wt%)
0.020
0.0014
<0.0125
0 (wt%)
0.071
0.075
<0.08
N (wt%)
0.012
0.012
<0.05
A 30.5 cm long section from one of these billets was forged at Wyman-Gordon in
Houston. The two step forging consisted of an O + p forging followed by a p forging.
The forging parameters are listed in Table 3-2.
Table 3-2. Forging parameters for the material used in this study. The first line is step 1
and the second line is step 2.
Initial
Height
cm (in)
Final
Height
cm (in)
Reduction
ratio
Billet Temp
°C (°F)
Die Temp
°C (°F)
Strain
Rate
(/min)
Cooling
(off dies)
30.48 (12)
15.88 (6.25)
1.92/1
996(1825)
870(1600)
0.95
Air
15.88 (6.25)
6.35 (2.50)
2.5/1
1100 (2010)
870 (1600)
0.92
Fan
After the billet was forged, it was removed from the forging press and allowed to
cool to room temperature. It was then given an aging heat treatment at 815°C (1500°F)
for 4 hours followed by air cooling without being subjected to a solution heat treatment.
This procedure is termed direct aging. A typical microstructure is shown in Figure 3.1. A
diametrical slice was then taken across the center of the forging for evaluation of the
material flow. Unfortunately a carbide cut-off wheel without cooling was used to section
this slice and ,as a result, the slice and both cut faces of the forging were severely

37
cracked. This caused some concern as to the integrity of the forging, so scanning
acoustic microscopy (SAM) was conducted to evaluate the depth of the cracking.
Figure 3.1. BSE image of a typical microstructure of the Ti-22Al-26Nb material after B2
forging and heat treatment at 815°C for 4 hours.
A 1.27 cm (0.5 inch) thick section was cut off of each forging face prior to the
SAM analysis. The SAM analysis revealed no cracking on the remaining faces indicating
that all of the cracks were confined to the near surface region. The 1.27 cm (0.5 inch)
thick section was cut in half and polished and macro etched to reveal the flow lines in the
forging. The macro flow lines are shown in Figure 3.2.

38
Figure 3.2. Photomacrograph of one half of the forging showing flow lines.
After the forging was pronounced sound, an extensive analysis of the variation in
grain size and grain aspect ratio versus location in the forging was conducted. The
effective strain from forging was calculated and iso-strain lines, shown in Figure 3.4,
were constructed. The calculated iso-strain lines were superimposed onto the forging
macrograph, and the location of various forging strains were identified. Figure 3.3 shows
the half diametrical slice from the forging center marked for photomacrographs at lOx
magnification. After the photographs were taken, the diametrical slice was sectioned
along the line shown into metallographic samples for microstructural analysis. The
forging was assumed to be symmetrical about the centerline and thus was sectioned to
reveal the grain structure in one quadrant of the forging. The diametrical slice was
sectioned with a carbide cut-off wheel and the samples were mounted in bakelite and
polished and etched. Microphotographs from representative locations in each samples
were taken at 50x, 500x and lOOOx magnification.

39
Figure 3.3. Forging half showing location of lOx macrophotographs and lines for
sectioning into metallographic samples.
Eff.
Strain
A=
0.131
A=
0.300
B=
0.600
C=
0.800
D=
1.200
E=
1.500
F=
1.800
G=
2.100
H=
2.400
1=
2.700
â–¡ =
=2.957
Representative micrographs from the dead zone in the forging, location C, and a
more highly strained area, location P, are shown in Figure 3.5 and Figure 3.6,
respectively.

40
Figure 3.5. Photomicrograph taken from area C. Notice the equiaxed grain structure and
low grain aspect ratio.
Figure 3.6. Photomicrograph taken from area P. Notice the high grain aspect ratio
resulting from the higher effective forging strain.

41
Notice the equiaxed grain structure from area C and the high grain aspect ratio in
area P. These two areas encompass the range of microstructures (grain sizes and aspect
ratios) that resulted from the forging process.
The rate at which the forging cooled from the (5 transus temperature would effect
the size and distribution of the O phase lath. Thus, the cooling rate of the forging from
the p transus temperature through 870°C was calculated. There was little variation in the
cooling rate across the forging, and thus little variation in the O phase lath size would be
expected. This was confirmed by performing high resolution SEM analyses, shown in
Figure 3.7 and Figure 3.8, and little or no variation in the orthorhombic lath was
observed.
Figure 3.7. High resolution SEM image of area S (identified in Figure 3.3) from the
slowest cooled region in the forging.

42
Figure 3.8. High resolution SEM image of area Q (identified in Figure 3.3) from the
fastest cooled region in the forging.
Specimen Blanking
After determining that the forging was sound, a plan for sectioning the oxidation
samples from the near surface dead zone and the mechanical property samples from
inside the forging was devised. This was based on the premise that the microstructure
found in this dead zone should not effect oxidation. The sectioning plan along with the
two ingot sections were submitted to the wire electro-discharge machining (EDM) shop
for sectioning as shown schematically in Figure 3.9.

43
Figure 3.9. Fatigue and oxidation specimen sectioning diagram
The remaining oxidation samples and the tensile samples were sectioned from the
other forging half as shown in Figure 3.10 which also shows the scrap blocks, numbered
1 through 6, that were left after sectioning. The oxidation samples were cut from near the
surface and from a block in the middle of the forging, some in low strain and some in
high strain regions, as shown in Figure 3.10 surface and the tensile samples were taken
from the more highly worked material. The details of this sectioning is shown in Figure
3.11. The open circles again represent the 5.08 cm (2 inch) long tensile/creep sample and
the shaded circles represent the oxidation samples.

44
T71 TI
Figure 3.10. Tensile and oxidation sample sectioning diagram. Shaded samples are
oxidation and white samples are tensile/creep.
Eff. Strain
A= .300
B= .600
C= .800
D= 1.200
E= 1.500
F= 1.800
G=2.100
H= 2.400
1= 2.700
0=2.957
á=.13160
Figure 3.11. Details of sectioning diagram for oxidation (shaded) and tensile (open)
samples from Figure 3.10. Calculated effective strain is also depicted

45
Test Specimens and Test Procedures
Oxidation
Samples
The oxidation samples were 8.9 cm (3.5 inches) long and 0.635 cm (0.25 inch) in
diameter with a 0.3175 cm (0.125 inch) radius machined on each end (Figure 3.12). The
total surface area of each sample was 17.72 cm2. The specimen blanks were EDMed
oversize from the forging, then centerless ground to the final diameter. The radius was
machined by turning on a lathe; this made complete coverage of the sample easier and
allowed for a simpler calculation of surface area as needed for the oxidation
measurements. More importantly, the rounded ends eliminated any sharp comers which
might act as stress risers for the coating leading to premature failure of the coating during
thermal cycling.
Figure 3.12. Drawing of oxidation sample used in this research
Testing
The oxidation testing was conducted in a bottom-loading thermal cycle furnace
built by Ted Kominsky Ovens and Furnaces, Model number KS16000B. A photograph
of the furnace is shown in Figure 3.13. The temperature was controlled with a Floneywell
UDC 3000 controller using input from a type K (NiCr-NiAl) thermocouple. This

46
thermocouple was calibrated against a NIST standard calibrated thermocouple. The
samples were loaded vertically into a sample fixture made from a cordierite open cell
foam to facilitate rapid cooling as shown in Figure 3.14. Note the thermocouples that
protrude through the hearth plate at the same height as the oxidation pins. One
thermocouple is attached to a the controller and isthe other is attached to a strip chart
recorder that records temperature and number of cycles.
Figure 3.13. Photograph of furnace used to conduct cyclic oxidation tests.

47
The samples are raised into the furnace and the furnace is automatically started.
The samples remain up in the furnace for 55 minutes, then are lowered and fan cooled for
5 minutes. A typical thermal cycle taken from the strip chart recorder is shown in Figure
3.15. As seen in Figure 3.15, it takes approximately 12 minutes to reach 650°C, and 5
minutes to cool down to approximately 100°C. The samples were weighed at 1, 3, 5 and
10 cycles, then every 10 cycles to 100 cycles and then every 100 cycles to 1000 cycles
where the test was terminated and the samples analyzed. The samples were handled
wearing cotton gloves during weighing to prevent contamination from the examiner and
after each weighing the samples were flipped over to ensure that one end was not always
in contact with the fixture. The samples were weighed on a Mettler 200XT digital scale
with an accuracy of ± 0.1 mg.
Figure 3.14. Oxidation pins loaded on hearth plate. The samples are the darker pins, the
thermocouples are the two white rods in the foreground. The cordierite block is resting
on the furnace platen.
The samples were then sectioned transverse and longitudinally and mounted in
Conductomet© for analysis. Electron microprobe analysis (EMPA) was conducted at

48
RPI by Dr. David Wark. Microhardness measurements were conducted on a LECO DM-
400FT at a 1 Ogf load and a 20 second load time. The tester was calibrated prior to taking
each set of measurements using the LECO calibration blocks. The data presented is from
one data point at each location; therefore, there will be some experimental error due to
variations in the microstructure or operator error in reading the microhardness
indentation. The results from the microhardness were compared to the oxygen levels
detected by the EMPA, and correlations were made concerning the effectiveness of the
various coatings to prevent oxygen ingress. The EMPA results were also used to
correlate other diffusion phenomena, such as that of the Ni and Fe into the alloy substrate,
to the microhardness results.
too 1 1 | 1 1 1 1 I 11 11 I 11 1 1 I 1 1 1 1 i 11 1 —
0 10 20 30 40 50 60
Time (mins)
Figure 3.15. Typical thermal cycle recorded by strip chart recorder.

49
Tensile
Samples
The tensile samples are 5.08 cm (2 inches) long and have a gauge length of 1.9 cm
(0.75 inches). The gauge diameter is 0.40 cm (0.16 inches). The sample is gripped by Vt-
20 threads with a thread root radius of 0.25 mm (0.01 inches). This sample is shown
schematically in Figure 3.16.
Figure 3.16. Drawing of tensile sample used in this research. Dimensions are in inches.
Testing
The tensile tests were conducted on an Instron Model 1125 screw driven test stand
using a 5,000 pound load cell. The cross-head speed was 0.05 cm/min (0.02 in/min) and
the sample were held in threaded grips. The strain was measured from the cross-head
displacement, and load vs. time was recorded electronically and on a strip chart recorder.
Percent strain was calculated from the cross-head displacement divided by the initial
gauge length multiplied by 100. The stress was calculated by dividing the load by the
initial cross sectional area. A typical stress-strain curve is shown in Figure 3.17 for
sample B75.

50
Figure 3.17. Typical stress-strain curve generated from tensile testing. Line tangent to
slope of elastic portion of curve used to determine elongation at failure is shown.
Fatigue
Samples
Fatigue samples were sectioned from many different locations in the as-received
forging, as discussed above. One facet of this study was to determine is the location of
the sample from the forging had any effect on the fatigue properties. To do this, the
sample location was recorded and compared to the calculated effective strain from
forging, as shown in Figure 3.18. The samples that were taken from various locations
have been divided into three groups: low (eeff <1.0), medium (1,0< seff < 2.0) or high (seff
> 2.0) strain locations.

51
Low
Med
High
❖=.13160
D= 1.200
G=2.100
A= .300
E= 1.500
H= 2.400
B= .600
F= 1.800
1= 2.700
C= .800
â–¡ =2.957
Figure 3.18. Calculated iso-strain lines from forging showing various effective strain
from forging and respective fatigue sample location.
The effect of cooling rate on microstructure was also discussed above, and the
calculated cooling rate from forging was also superimposed over the diagram of fatigue
samples, as shown in Figure 3.19. The effect of location on the fatigue results was
determined and will be discussed in Chapters 4 and 5.
Figure 3.19. Calculated cooling rates from 1080°C to 870°C in °C/min and corresponding
fatigue specimen location.
The fatigue samples were 8.89 cm (3.5 inches) long with a uniform gauge length
of 2.22 cm (0.875 inches). The gauge diameter was 5.08 mm (0.200 inches) with a
surface finish of 8 microinches, and the sample was gripped by ‘/2-20 threads. This

52
sample is shown schematically in Figure 3.20. In order to avoid stress concentrations at
the surface, the samples were longitudinally polished in the uniform gauge to eliminate
all circumferential grinding marks and to produce all polishing marks parallel to the
loading axis of the sample.
Figure 3.20. Drawing of fatigue sample used in this research. The dimensions are in
inches
Thermal cycling of all fatigue samples was conducted in the Kominsky thermal
cycling furnace shown in Figure 3.13. The samples were arranged horizontally on a
cordierite platen as shown in Figure 3.21. Samples were thermally cycled for 1, 10 and
100 cycles, as well as exposed in static air for 92 hours to compare with the 100 cycle
samples. Samples were also encapsulated in quartz filled with argon. These samples
were placed inside titanium tubes and inserted along with titanium sponge which was
added as an oxygen getter in quartz tubes. The quartz ampoules were evacuated prior to
backfilling and then filled to 200 torr with argon before sealing. On heating to 650°C the
pressure inside the ampoule reached approximately 760 torr.

53
Figure 3.21. Photograph of LCF samples loaded for thermal cycling. Samples are
separated by the alumina tubes.
Testing
The fatigue testing was conducted on an MTS 810 servohydraulic test machine
with a 20,000 pound load cell. The testing was conducted under strain control using
resistive strain gauge extensometers (model number MTS 632.53E-14) with high
temperature alumina rods. The A ratio was 1 (orain = 0), calculated as
A = cra/CTra (3.1)
where (7 max- (7 min
<*.= T (3-2)
and am the mean stress defined as
(7 max+ <7 min
°m= T (3-3)
The cycling rate was 20 cycles per minute. The testing was computer controlled
using Testware SX software that incorporated data collection. Peak and valley load data

54
were collected electronically for every cycle. Load-strain information was recorded at
various cycles on a strip chart recorder, as shown in Figure 3.22. Hydraulic grips were
used to eliminate the need for a backing nut on the threads thus allowing for tension and
compression loading. The elastic modulus of each sample was determined by loading the
sample in the elastic region and recording the resultant strain on a strip chart recorder, as
shown in Figure 3.22. The applied load was divided by the sample cross sectional area in
the uniform gauge to determine the applied stress. The calculated stress was then divided
by the measured elastic strain in order to estimate the room temperature elastic modulus.
This room temperature modulus was multiplied by the applied strain during testing in
order to give the average, or pseudo-stress applied during the fatigue testing. This
pseudo-stress is typically reported as the alternating pseudo-stress, which is the pseudo¬
stress divided by two.
Analysis
After fatigue testing, two cross-sections were cut from the sample, one containing
the fracture surface and one behind the fracture surface, as shown in Figure 3.23. The
cross section containing the fracture surface was sectioned longitudinally to observe the
crack path, the initiation site and opposite the initiation site on the longitudinal section.
These samples were mounted in Conductomet® such that the initiation site in the
longitudinal section was on the same side as the cross section as shown in Figure 3.23.
This allowed the initiation site to be identified during metallography.

55
Figure 3.22. Typical LCF data taken during testing showing the hysteresis loops at
various cycles. Curve for modulus determination is also shown.
Figure 3.23. Illustration of sectioning and metallographic mounting of the fatigue sample
for evaluation.
The samples were then examined in the SEM using backscatter analysis or on the
optical metallograph after etching.

56
Coatings
All of the samples to be coated were ultrasonically cleaned in a warm Alconox®
bath for 30 minutes, followed by a de-ionized water rinse and a 2-propanol rinse. The
samples were dried with a Kimwipes® towel and placed in sterile plastic bags, and only
handled with gloves prior to coating.
Low Pressure Plasma Spray (LPPS) Coatings
For LPPS, the oxidation pins were secured in a custom designed fixture that holds
the pin at each end and thus allows the complete pin to be coated instead of just half of
the pin. This cuts the processing time in half and eliminates any overspray in the center
of the pin. The fixture is shown in Figure 3.24. After the sample was loaded into the
fixture, the fixture was loaded into the LPPS chamber. The chamber was then pumped
down to 50 torr and backfilled with argon. The sample oscillation and rotation were
initiated and the plasma guns were turned on to pre-heat the sample to approximately
650°C. The powder feeders were then started and the sample was coated to the proper
thickness. The plasma guns and feeders were then shut off and the sample was allowed
to cool in the chamber. The sample was then removed from the single pin fixture and
placed in a fixture with several other samples where only the end radii are exposed.
These radii were also coated to insure that the weight gain versus unit area measured
during oxidation was as accurate as possible.

Figure 3.24. Drawing of LPPS specimen fixture.
NiCrAlY + 40 vol% Al?Ch.
NiCrAlY (Ni-21.7Cr-10Al-l.13Y (wt%)) powder (purchased from Praxair) was
low pressure plasma sprayed (LLPS) onto a scrap oxidation pin to assess the LPPS

58
process. The NiCrAlY broke away from the pin in three large sections after processing.
Since it was initially suspected that the fixture might be the cause of the lack of adherence
another trial was attempted, this time holding the pin on one half only. The result was the
same; the coating broke away from the pin after processing. The pin was sectioned in the
location of the coating failure and examined using SEM backscatter techniques. It was
determined that there was residual Ni left on the surface of the pin, indicating that the
coating was initially adherent but subsequently spalled. This spallation was probably a
result of the large thermal expansion differences between the substrate and the coating.
Consequently, a mixture of 60 vol% NiCrAlY and 40 vol% AI2O3 was produced by
blending, and this “cermet” was deposited on the substrate. This mixture provided a
suitable thermal expansion match between the coating and substrate and, as a
consequence, the coating was adherent.
FeCrAlY.
FeCrAlY (Fe-29.9Cr-4.9Al-0.6Y (wt%)) powder was also applied to the oxidation
pins via LPPS. Since this coating was successfully applied and did not spall, no changes
were made to the deposition process.
Sputtered Coatings.
Several elemental coatings were deposited onto oxidation pins via vacuum
sputtering. The samples were held on one end and rotated in the sputtering chamber via a
magnetic feed through. The elements were either reacted with the substrate at elevated

59
temperature to create a stable compound at the surface(e.g., Al3Ti) or were used in their
pure elemental form.
Sputtered Al.
25 pm of Al was sputtered onto the oxidation pins in a vacuum sputtering unit.
These pins were then reacted in dry, pure argon using the following heat treatment
schedule:
Ramp to 600°C at 300°C/hour,
Hold at 600°C for 2 hours,
Ramp to 630°C at 10°C/hour,
Hold at 630°C for 10 hours,
Furnace cool.
X-ray diffraction confirmed that a mixture of Al3Ti and Al3Nb were formed at the
surface, as shown in Figure 3.25.
Figure 3.25. SEM backscatter micrograph of Al coated and reacted sample showing the
in-situ formed A13TÍ coating on the right.

60
Sputtered Si
Elemental Si was also sputtered onto the orthorhombic Ti-aluminide substrate in a
vacuum sputtering unit. The sputtering rate was very low and, as a consequence, only
12.7 pm of Si was deposited. The Si coated pins were reacted as follows:
Ramp to 1200°C,
Hold at 1200°C for 16 hours,
Furnace cool.
This cycle was designed to produce the Ti5Si3 phase on the surface and XRD
indicated that this phase was indeed produced, as shown in Figure 3.26. The reaction
temperature was above the (3 transus for this alloy, and the microstructure indicated that
the prior (3 grains had indeed grown and the as forged microstructure was eliminated.
Figure 3.26. SEM backscatter micrograph of Si coated and reacted sample showing in-
situ formed TisSi3 coating on the right.

61
Sputtered Pt and Cr
Pt and Cr was sputtered onto the oxidation pins in order to provide protection
based on the low oxidation rates of pure Pt, and the stability with titanium [Fuj79, Eyl85],
and because Cr has been show to provide good oxidation resistance as a coating for
titanium [McK93, McK90]. Both coatings were adherent, although the Pt coating was
porous and columnar, as shown in Figure 3.27.
3,50KX 28KU WD•9MM S-00000 P:00000
i0 ij n
PT COATED, 1000HRS 6580
Figure 3.27. SEM backscatter micrograph of Pt coating showing the columnar structure
and porous nature of the coating.
CVD coatings
Several oxide coatings were produced by metal organic chemical vapor deposition
(MOCVD). The idea was to evaluate the stability of the oxide in contact with the
substrate and determine if the Ti would reduce the oxide. MOCVD was selected because
of the capability of producing adherent coatings of oxides on any and every surface in the
reactor. The CVD process, unlike PVD or LPPS, is not line-of-sight, and therefore every

62
hot surface in the reactor will be coated. The MOCVD process consists of a metal
organic precursor that is evaporated and flowed over the substrate held in a reaction
chamber at elevated temperature, as shown in Figure 3.28 and Figure 3.29. The selection
of the precursors and the deposition temperatures are described below.
Figure 3.28. Photograph of CVD reactor.
0
7) Pressure
| Gauge
I
Furnace |
Samples
To Vacuum
Pump
Liquid
Nitrogen
Condenser
Furnace
Figure 3.29. Schematic diagram of CVD reactor shown in Figure 3.28.

63
SiO,
Oxide and silicide coatings have been reported to provide oxidation protection for
a-p titanium [Sol85, Cla88, Bed91], 0.2 alloys [Wei89, Dev90] and orthorhombic alloys
[Coc96]. The mechanism is through the production of an oxide comprised mainly of
Si02. For this reason, Si02 was chosen as a diffusion layer for an outer coating, and
possibly as a primary oxidation coating. The reagent used to produce the Si02 was
silicon tetraethoxide [Si(OC2H5)4]. The reactor temperature was 670°C at a pressure of
500pm. The Effusion cell temperature was 90°C and the flow rate of reagent was 2x10‘5
gs''cm'2. The deposition rate of the Si02 was approximately O.lnm/sec with a reagent
utilization of approximately 0.5%. The reaction taking place was
Si(OC,H5)4 O Si02 + C2H5OH + C2H4
TaA
Tantala was chosen as a barrier layer due to its higher CTE over Si02 and greater
oxidation resistance relative to Ti02. The reagent used to produce the Ta^ was tantalum
ethoxide dimer [Ta^OQFy The reactor temperature was 425°C at a pressure of
100pm. The Effusion cell temperature was 125°C and the reagent flow rate was 3xl0'5
gs 'cm'2. These conditions led to a deposition rate of the Ta2Os of approximately
O.lnm/sec with a reagent utilization of approximately 0.5%. A coating thickness of 0.45
pm was typically deposited on the oxidation pins and fatigue samples. The reaction
taking place was
Ta^OC,!!,),» ■=> TaA + 5 C2H5OH + 5 C2H4

MgO has a high oxygen diffusivity in air at 650°C and therefore it probably would
not provide adequate protection as a stand alone oxidation coating. But MgO has been
evaluated as a barrier layer for fiber-reinforced titanium aluminide composites [McG95].
It was shown to dissolve oxygen into TÍ-24A1-1 INb after 100 hours at 1000°C, but at a
reasonably low rate. MgO has also been formed in-situ in Ti-Mg alloys produced by high
rate evaporation and quenching. After oxidation at 850°C for only 10 minutes, MgO
particles were formed in the titanium matrix [War95]. For this reason and because MgO
has a higher thermal expansion coefficient than the Ti-22Al-26Nb alloy, it was chosen as
a candidate coating system.
The reagent used to form MgO was magnesium 2,4 pentane dionate and the
chemical reaction was
Mg (C5H70,)2 H20O MgO + 4 C2H5OH + C,H4
The reaction temperature was 540°C at a partial pressure of 0.5 torr. Oxygen was
the initial carrier gas and was chosen to reduce the level of carbon in the deposited film.
After deposition, the film thickness was 0.85 pm. The samples were baked out in an air
furnace for 1/2 hour at 575°C, and the film thickness decreased to 0.3 pm. This is
probably due to evolution of carbon into carbon dioxide in the furnace. The source of
carbon is from the reaction of the reagent in the furnace and the incomplete formation of
the reaction products. This process is still in the lab stages and is not yet fully reduced to
practice.

CHAPTER 4
RESULTS AND DISCUSSION
Oxidation
Cyclic Oxidation Testing
Baseline results
Four baseline samples were cyclically oxidized at 650°C. Three of the samples
(136, 187 and 191) were chosen to have different levels of strain from forging, and thus
different grain size and grain aspect ratios, as discussed in Chapter 3. These samples
were chosen to determine if the strain history, which results in a slight variation in
microstructure, has any effect on the oxidation behavior. The results from these baseline
studies are shown in Figure 4.1. The difference in oxidation rates between these three
samples is small as is the weight change over time. Therefore, it can be stated that there
is little or no effect on oxidation from the strain history. The fourth sample, 174, was run
in a separate oxidation test with another series of samples as a control. The differences
observed between the samples indicates typical experimental scatter.
65

66
Figure 4.1. Baseline oxidation results showing no variation between three samples from
different locations run together, and a larger variation with a sample from a different
experimental run.
A1 and Si Coated and Reacted Results
The results from the oxidation testing on A1 and Si coated and reacted samples are
shown in Figure 4.2. The average of the baseline samples is also plotted as a reference.
As can be seen, both the A1 and Si coated and reacted samples behaved well in cyclic
oxidation tests. This behavior has been observed previously for Ti coated with Al3Ti
[81Str, 85Unn, 91Abd, 91McM, 92Gal, 93Abb, 93Wie] and a2 coated with Al3Ti [88Sub,
90Kun, 90Smi, 93Smi]. Silicide coated Ti-22Al-27Nb has also been evaluated and

67
shown to have similar oxidation behavior as the Si coated and reacted Ti-21 Al-26Nb
[96Coc],
Figure 4.2. A1 and Si coated and reacted oxidation results showing the decrease in
oxidation rate over the baseline.
Oxide Coating Results
The effect of cyclic oxidation on the oxide coated samples are shown in Figure
4.3. Both TajOj and MgO behave similarly to the baseline alloy while the Si02 coated
sample exhibited better oxidation behavior than the baseline at short times, but began to
gain weight in a linear fashion with time after 400 cycles. This indicates that the coating
was not protective and that oxygen was diffusing through after 400 hours causing the

68
sample to gain weight. These results indicate that the oxide coatings by themselves do
not provide protection from an oxidizing environment.
Figure 4.3. Oxide coated oxidation results showing that Ta,05 and MgO coated material
behaves like uncoated material and Si02 coated samples begin to gain weight after 400
cycles.
Metallic Coating Results
The results from the cyclic oxidation testing of the metallic coated samples are
shown in Figure 4.4. The results show that the MCrAlY coatings provide excellent
oxidation resistance, that Cr provides intermediate protection, and that Pt has little effect
over baseline behavior. The protection by the MCrAlY coatings has been shown in
previous work on a2 alloys [92Bri, 93McK, 95Sch] and titanium alloys [92Tob, 90Lut,

69
90McC, 93McK], On the other hand, platinum, in the form of PtAl2, has been shown to
reduce the cyclic oxidation rates in conventional titanium alloys and in Ti3Al [93Nic].
Pure platinum coatings have also been shown to improve creep resistance [85Eyl] and
decrease oxidation rate [79Fuj] of Ti-6242. Finally, pure Cr coatings have also been
shown to reduce the weight gain over time of Ti-64 and Ti-6242 [90McK],
Figure 4.4. Results showing the large reduction in oxidation rate with the application of
the MCrAlY coatings and the effect of Cr and Pt coatings.
The results from all of the cyclic oxidation tests are compared in Figure 4.5. As
can be seen, the metallic coatings, excluding Pt, provide excellent oxidation resistance

70
after 1000 cycles, whereas the Ta,Os and MgO coatings behave similarly to the baseline,
and the Si02 coated samples begin to degrade after 400 hours.
Figure 4.5. Comparison of the oxidation rate behavior of all coatings examined in this
study.
Microhardness and Microprobe Evaluation
Baseline uncoated
As shown in Figure 4.1, the three samples with varying forging strains behaved
similarly, so after 500 cycles samples 191 was removed for analysis. Both microhardness
and oxygen content vs. depth are plotted in Figure 4.6. Notice that in both curves, the
oxygen content and microhardness follow the same trend. Further work is needed to

71
determine if the high value of oxygen in the 1000 cycle sample at 25 pm is scatter or an
indication of localized oxidation, possibly an oxide.
- 1000
600
500
300
Figure 4.6. Uncoated baseline oxygen content and microhardness vs. depth.
The backscatter SEM images of the two microhardness samples (Figure 4.7 and
Figure 4.8) indicate that the dark zone near the surface of the sample is high in oxygen.
The microhardness indentations visible in these figures clearly show that these regions
are much harder that the base material. The depths of oxygen rich zone in the 500 and
1000 cycle samples are approximately 20pm and 35pm respectively. Dividing the square
of the depth (in cm) by the time in seconds, it is possible to estimate the diffusivity of
oxygen into the O + B2 lattice at 650°C as:
x =VDt (cm) [4.1]

72
Figure 4.7. SEM backscatter image of baseline microhardness sample cycled for 500
cycles.
Figure 4.8. SEM backscatter image of baseline microhardness sample cycled for 1000
cycles.

73
or
D = (cmVsec) [4.2]
This yields a D of approximately 3.4 x 10-12 cm2/sec. This estimate ignores the oxide
present on the surface of the sample as well as any interaction that oxide has with the
substrate, and the fact that these are no isothermal experiments. These are a valid
assumptions given the fact that the oxide has a diffusivity of oxygen that is 50x that of the
metal [Unn86], The diffusion coefficient of oxygen for CP a-Ti at 650°C is
approximately 2.5 x 10-12 cm2/sec [Kub83] and for TÍ-24A1-15 Nb at 1027°C is
approximately 2.6 x 10'12 cm2/sec [Roy96]. These values correlate well to the 3.4 x 10_
12 cm2/sec value calculated from the baseline samples.
FeCrAlY and NiCrAlY + AI2O3
Brindley et al. [Bri92] have evaluated the oxidation behavior of both of these
coatings, and found that while they provide excellent resistance to oxygen ingress into
titanium aluminides, the elements in the coatings themselves diffuse into the substrate
and cause embrittlement. Figure 4.9 is an SEM backscatter image of the NiCrAlY +
A1203 coated specimen after 1000 cycles. The NiCrAlY + A1203 plasma sprayed coating
can be seen in the lower right hand comer. Notice that the microstructure of the matrix
next to the coating has changed from the lath morphology to a fine, equiaxed structure.
Work conducted by Schaeffer et al. [Sch95] on 012 coated with NiCrAlY + A1203
demonstrated that the coating alone, without exposure, degraded the LCF life by two
orders of magnitude.

74
Figure 4.9. SEM backscatter image of the NiCrAlY + AI2O3 coated sample after 1000
cycles at 650°C.
The results from this study support Schaeffer’s results, as can be seen from the
microhardness and oxygen profiles in Figure 4.10 through Figure 4.12. Figure 4.10 and
Figure 4.12 compare the oxygen level in the substrate to the measured microhardness.
Notice that there is simply one point under the coating that contains oxygen, and beyond
that point the level is below the detectability limit of the EMPA. Figure 4.11 and Figure
4.13 show the Fe and Ni concentration vs. microhardness. The Fe and Ni content seems
to correlate with microhardness, and the Ni and Fe levels both taper off with distance
from the interface, as does the microhardness. This could contribute to a loss in ductility
and thus fatigue life. Recall that LPPS is conducted in a partial pressure of argon and air,
and the sample is preheated to 650°C for several minutes prior to coating during which
oxygen could ingress. A diffusion barrier between the substrate and coating might
possibly control the ingress of Ni or Fe from the coating into the substrate, but the pre¬
heat might still cause embrittlement.

75
Figure 4.10. Microhardness and oxygen vs. depth for the FeCrAlY coated and 1000
cycle exposed sample.
550 I"
c_

500 %
450 ^
400
Figure 4.11. Microhardness and iron vs. depth for the FeCrAlY coated and 1000 cycle
exposed sample.

76
r.
I
z
Figure 4.12. Microhardness and oxygen vs. depth for the NiCrAlY + AI2O3 coated and
1000 cycle exposed sample.
Figure 4.13. Microhardness and Ni vs. depth for the NiCrAlY + AI2O3 coated and 1000
cycle exposed sample.

77
Sputtered and Reacted A1 and Si
Elemental A1 and Si were sputtered onto the orthorhombic under vacuum. The A1
coated samples were then ramped to 600°C in two hours and held for two hours followed
by ramping to 630°C in 3 hours and holding for 10 hours. This produced Al3(Ti,Nb) as a
surface reacted coating, as verified by XRD. The Si coated samples were ramped to
1200°C as quickly as possible and held for 16 hours. The intent was to produce the Ti5Si3
phase and miss the Ti3Si phase. This was accomplished as can be seen in Figure 4.17,
and also verified by XRD. The microhardness profiles vs. the oxygen content are shown
in Figure 4.14 for the A1 coated sample and in Figure 4.16 for the Si coated sample.
Figure 4.15 and Figure 4.17 are SEM backscatter images of the A1 and Si coated and
reacted samples, respectively, obtained after cycling at 650°C for 1000 cycles.
800
700
"C
p*
600 g_
3
CD
co
CO
500 £h
3
400
300
-5 5 15 25 35 45 55 65
Distance from Coating/Substrate Interface (pm)
Figure 4.14. Microhardness vs. oxygen content for the A1 coated and reacted sample.
Note that the distance from the interface starts at -5 pm.
tí
S o.i
l—■ ■ ■ I
Oxygen (at%)
H hardness (KHN)

78
Figure 4.15. SEM backscatter image of Al coated and reacted sample after 1000 cycles at
650°C. Notice the crack in the Al3Ti coating.
Notice that the Al3(Ti,Nb) coating has cracks that run to the substrate, yet the
coating behaved very well in oxidation. The SEM backscatter images do not show any
indications that oxygen diffused preferentially down the cracks and into the substrate.
This could be due to the fact that on heating, the coating expanded more than the
substrate and thus the cracks were not open during exposure. Microhardness
measurements were taken in the reacted layer because in optical microscopy the layer
looked like the substrate. Therefore, the oxygen analysis and the microhardness start at
minus 5 pm in both the A1 and Si coated cases. The zero point is the location of the
coating/substrate interface after the reaction. This indicates that the coating was the
major source of oxidation and hardness. Directly beneath the reacted coating, the oxygen
content was negligible and the microhardness had returned to that of the substrate. But
since the coating was atomistically attached to the substrate, any cracks in the coating
would propagate into the substrate and cause limited LCF life.

79
1000
900
800
700
600
500
400
300
200
-5 5 15 25 35 45 55 65
Distance from Coating/Substrate Interface (pm)
Figure 4.16. Microhardness vs. oxygen content for the Si coated and reacted sample.
Notice that the distance from substrate starts at -5pm.
o.i --
^ 0.08
ca
§ 0.06
OX)
X
O 0.04
Figure 4.17. SEM backscatter image of the Si coated and reacted sample after 1000
cycles.
^hardness (KHN)

80
Sputtered Pt and Cr coatings
Pt was sputtered onto the orthorhombic substrate to a thickness of about 1 pm.
The Pt coating (Figure 4.18) was columnar, porous and discontinuous. As a result, it
allowed oxygen to diffuse to a similar depth as in the uncoated material. The
microhardness and oxygen profiles (Figure 4.19) are consistent with the poor behavior of
the coating in oxidation. Had the coating been dense and adherent, it should not have
allowed oxygen to pass at 650°C according to previous work conducted on Pt coated Ti-
6242 [Fuj79] where the rate of weight gain was reduced from 3xl0'‘ to 2x1o*4 at 593°C.
Their coating, deposited by ion plating, was approximately 1 pm thick and was dense and
continuous. Sputtered Cr exhibited better oxidation behavior than Pt during cyclic
oxidation, as shown in Figure 4.4. The level of oxygen beneath the coating after
exposure is shown in Figure 4.20, along with the level of Cr. Oxygen content for the
unexposed sample was zero as determined by EMPA. There is essentially no Cr diffused
into the matrix after 1000 cycles at 650°C, and very little oxygen, compared to the
baseline of 20 at%.

81
3¡50KX 20KU HO = 9MM S• 00000 p:00000
i Bun •
PT COATED, 1008HRS 650C
Figure 4.18. SEM backscatter micrograph of Pt sputter coated sample cycled for 1000
cycles.
Figure 4.19. Microhardness and oxygen content vs. depth for the Pt coated and 1000
cycle exposed sample.
phardness (KHN)

82
Figure 4.20. Plot of Cr and oxygen content vs. depth in 1000 cycle exposed oxidation
sample.
CVD Oxide Coatings
The CVD SÍO2 coating was dense, adherent, and appeared iridescent green in
color. The Ta^ and MgO coatings were also dense and adherent. All coatings were
approximately 0.5 pm thick.
As can be seen from Figure 4.3, the initial oxidation behavior of the SiO, coating
was excellent. This data is an average of three separate tests conducted with the SiO,
coating. However, after 400 cycles, the oxidation behavior of the Si02 coated samples
began to behave in a more linear relationship with time instead of parabolic. This is
typically an indication that there is no protective oxide present and that oxygen is simply

83
diffusing into the material [66Kof]. Figure 4.21 is a backscatter SEM image of the Si02
coated sample. The microhardness vs. oxygen content is shown in Figure 4.22. No Si
was detected beneath the coating by microprobe, and work done on coating TÍ6242 with
SiO, [Cla88] also concluded that the SiO, did not react with the Ti.
The MgO and Ta205 coated samples behaved similarly to the baseline samples
tested during the same experiment, i.e. oxidation was parabolic. SEM, Auger and
microprobe data support this by showing a region of high oxygen and Ti below the Ta,05
coating (Figure 4.24), which has been reduced to a sub-oxide coating during exposure.
Figure 4.21. SEM backscatter image of the SiO, coated sample after 1000 cycles at
650°C

84
Figure 4.22. Microhardness vs. oxygen content for the Si02 coated sample.
Figure 4.23. SEM backscatter image of Ta^ coated sample showing area analyzed by
Auger spectroscopy. Box outlines area where compositional maps were obtained.
phardness (KHN)

85
Figure 4.24. Auger compositional maps showing the concentration of Ta, Ti and oxygen
in the boxed region shown in Figure 4.23.
Effect of Forging Strain and Cooling Rate on the Tensile and LCF Behavior
Fatigue Results
Twelve samples representing low (6-9), medium (67-69) and high ( 46, 50, 52 and
53) effective strains were taken from a variety of locations in the forging. These samples
were tested at a variety of strains and corresponding pseudo-stresses that spanned a range
from 7000 cycles to runout. The fatigue results for the twelve samples are listed in

86
Table 4-1 and shown graphically in Figure 4.25. The effective strain is listed as
well as the measured modulus, the percent strain used to control the test, the alternating
pseudo-stress and the number of cycles to failure. Also listed is the location of failure in
the samples. A thread failure indicates that the sample failed prematurely in the threads
and may have run for more cycles until failure in the gauge.
Table 4-1. Testing parameters and results for samples used to determine effect of
location on LCF.
Specimen #
Effective
Strain
E
Ksi
%£
Alt. P.S.
Ksi
Cycles
#
Failure Location
6
Low
18111
0.90
81.50
38403
Gauge
7
Low
17785
1.00
88.93
7483
Gauge
8
Low
18073
1.00
90.37
13273
Gauge
9
Low
18243
0.95
86.65
14287
Thread
10
Low
18571
1.10
102.14
6598
Gauge
67
Med.
17989
0.85
76.45
114000
Runout
68
Med.
18737
0.95
89.00
37971
Thread
69
Med.
18571
1.00
92.86
9763
Shoulder
46
High
18054
0.85
76.73
57806
Thread
50
High
17291
1.00
86.46
30223
Gauge
52
High
17129
0.95
81.36
19802
Gauge
53
High
17518
1.10
96.35
12491
Gauge
In Figure 4.25, the sample number is shown beneath each symbol, and the failure
location is listed above each symbol. As seen in Figure 4.25, there is no strong correlation
between the location of a sample in the forging and the LCF behavior. Since the scatter
in fatigue data is typically large [Her83] and since each data point represents a single
LCF test, the statistical probability of the data point representing the fatigue life of the
material at the given stress is much less than 100% [Die86]. The scatter of the data does
follow a logarithmic function typical of fatigue behavior in most metals [Die86]. Given

87
the fact that the data is derived from only one test per data point, an average curve was
not constructed. Instead, a polygon was drawn that encompasses all of the relevant data
points, excluding the thread failures, as shown in Figure 4.25. It is this polygon that will
be used as the baseline for further comparisons.
Figure 4.25. Number of cycles to failure vs. strain for various locations in the forging.
The sample number is shown beneath each symbol, and the failure location is listed
above each symbol
Tensile Results
Samples for tensile testing were also taken from various regions in the forging to
corresponding low, medium and high calculated effective strain. One sample
representing each strain level was tested at both room temperature and 540°C to
determine if the variation in grain structure effected the room temperature or elevated

88
temperature properties. The samples were tested on an Instron screw driven machine and
the percent elongation was taken from the stroke-displacement information and is shown
as plastic deformation. The tensile testing results are shown graphically in Figure 4.26..
All of the data comes from a single tensile test so no standard deviation is available.
o
3
Figure 4.26. Ultimate tensile strength, yield strength and percent elongation (plastic) vs.
calculated effective strain from forging.
Discussion
The tensile results shown in Figure 4.26 indicate that there is little or no effect of
plastic strain from forging on the tensile properties. There appears to be a slight effect on
the ultimate elongation, where the more highly strained material with a larger grain aspect
ratio has high elongation and lower yield strength, but the same ultimate strength
compared to material with a more equiaxed grain structure. This same effect has been
observed when comparing forged material that was directly aged, and thus having grains

89
with a high aspect ratio, to material with large equiaxed grains obtained by heating above
the beta transus and then aged [Dob94], The beta solutioned and aged material failed at
lower yield strengths and elongation than the directly aged material.
The low cycle fatigue testing results also do not indicate any variation with
respect to effective strain, as shown in Figure 4.25. Assuming that the initial
microstructural change that could occur in this material is limited to local deformation of
the p phase or cracking of individual orthorhombic lath, it is not surprising that there is
no distinct variation in the LCF behavior for different locations in the forging. As shown
in Figures 3.7 and 3.8, there is little or no variation in the orthorhombic lath size or
morphology, so following the above argument for initiation from an orthorhombic lath or
interlath p phase, there is no reason to suspect a variation in fatigue behavior. To better
understand the effect of lath size on fatigue behavior, the initial flaw size of each sample
was calculated using the methods described in Chapter 2. Woodfield generated fatigue
crack growth data [Woo94] for Ti-21Al-25Nb forged above the beta transus and aged at
da
860°C for 4 hours and reported it as vs. AK. Using this curve, and recalling the Paris
law [Sur91]of
^=C(AKr (4.1)
one can calculate the Paris law constants of C and m as 5.5xl0'13 in/cycle and 6.9,
da
respectively. These constants are the Y intercept on the -— vs. AK curve (C) and the
dN
slope of the fatigue crack growth curve (m). Solving equation 2.7 for a,, yields

90
a» =
m m m_2
NfCY A 2
m-2
(4.2)
Using equations 2.2 and 4.2 and plugging in the Paris constants and the data from
all of the LCF tests, af and a„ for each test condition was calculated, respectively. This
data is shown in Figure 4.27, along with a best curve fit for all of the data and the
equation of this curve. Notice that the values of a„ are very close, from 7 to 12 pm, or an
average of 9.5±2.5 pm. The values of a, decrease with decreasing stress, as one would
expect. The slope of the best fit curve of the strain vs. number of cycles to failure curve
is inversely proportional to the slope of the fatigue crack growth curve, m.
Figure 4.27. LCF data showing the calculated a„ and a,- values, along with failure location
and sample number for each data point.

91
The slope of -0.178, which to the would indicate an m in the Paris equation of 5.6,
is very close to the calculated m of 6.9 and thus indicates that this analysis is valid. An
attempt was made to relate these values of a„ to some microstructural feature in order to
yield some insight concerning the mechanism of initiation in this material. Miller has
stated that cracks start to grow, or initiate, on the first cycle with the growth rate
depending on the microstructural scale and size of defects [MÍ197], The samples in this
study have had the surface defects removed by longitudinal polishing and there are no
intrinsic defects such as porosity or inclusions. Therefore, the initiation, or primary
growth depends on some microstructural feature. In Ti-25Al-10Nb-3V-lMo alloys
(Super a2) processed below the p transus to contain primary a2, it has been shown that
the fatigue cracks tend to propagate through the primary a2 and probably initiate there
[Dav91, Gog90]. Takashimi et al.[Tak96] showed that fracture during cyclic loading
occurs by cleaving the a2 and the cracks bridging the P phase. Cracks also propagated
along the a2/p interface in the P-transformed regions. Penton et al. [Pen93] has shown
that Super a2 containing primary a2 exhibits faster fatigue crack growth rates at room
temperature than p solutioned and aged material. This effect seems to be consistent with
the presence of the primary a2 and the fact that it cleaves readily. It was also suggested
that the growth rates in the p solutioned and aged material is consistent with the scale of
fracture in the a2 laths on each cycle. Ward et al. [War93] showed that cracking in the
transformed p regions propagates through the a, laths on what appears to be the basal
plane and the p phase between the laths experiences bridging and pull-out. These
observations could explain why the initial flaw size, a„, is larger than the lath size shown

92
in Figure 3.7 and 3.8. The initial crack may have been across the basal plane of several
lath that have grown sympathetically from a larger lath or from a prior |3 boundary. This
material may also contain orthorhombic lath colonies, which would allow crack initiation
and propagation across several laths with little change in direction or energy for arresting
the crack.
Understanding that the initial flaw size and crack growth rate is related to some
microstructural feature, possibly the orthorhombic lath size, one would not expect a
variation in the fatigue behavior versus location in the forging. The fracture surfaces bear
this out, i.e. there is little or no difference in the fracture surface morphologies as shown
in Figure 4.28 and Figure 4.31. Figure 4.28 shows the fracture surface of a sample (No.
52) taken from a high strain region of the forging. The initiation point is at the top of the
photo and the crack propagated downward until catastrophic failure. The grain aspect
orientation is from left to right in the photo as can be seen from the crack path. The
initiation site and crack path did not appear to be influenced by the grain aspect ratio.
Figure 4.28. Optical fractograph of specimen taken from a highly stressed region in the
forging (No. 52).

93
Figure 4.29 shows a secondary SEM photo at a higher magnification than Figure
4.28 of the same sample. Notice the bright region in the center of the photo. This region
is approximately 200pm in length, which is reasonably close to the calculated % of
220pm. A higher magnification secondary SEM photograph of the region in the white
box in Figure 4.29 is shown in Figure 4.30. It appears that there is evidence of fatigue
crack propagation along the surface of the fracture, and the steps of the stria are
approximately 1pm in length. Sample No. 52 lasted 19,800 cycles, so this would indicate
that the fatigue crack would be on the order of 20mm if each cycle generated a striation.
This crack morphology is most likely caused by microstructural features, such as
orthorhombic laths, failing during the crack growth period of testing.
Figure 4.29. Secondary SEM photograph of fracture initiation site from specimen #52
shown in Figure 4.28.
Figure 4.31 shows the fracture surface of a sample (No. 6)taken from a low strain
region in the forging. Again, the initiation site is at the top of the photo and the crack
path is downward. Notice the larger grains compared to sample number 53. Again there
appears to be no influence of the grain size or morphology on crack growth. This is

94
strong microstructural evidence that there is no influence of location in the forging to
LCF behavior and, therefore, the LCF data is not dependent on location.
Figure 4.30. Secondary SEM photograph of the fracture initiation site shown in the small
white box in Figure 4.29.
Figure 4.31. Optical ffactograph of sample number 6 taken from a low strain region in the
forging.

95
Effect of Exposure at 650°C on the Tensile and LCF Behavior
Tensile Results
The tensile test results are listed in Table 4-2 and shown graphically in Figure
4.32 through. The tensile samples tested were both unexposed and exposed for 100
cycles at 650°C. The samples were tested in air at room temperature and in air and
vacuum at 540°C. The elevated temperature testing in vacuum was to determine if there
was an effect of the environment on the tensile properties at 540°C.
Table 4-2. Tensile test data.
Sample
No.
Forging
Location
Exposure
Temp
°C
Exposure
Cycles
Test
Temp.
°C
Test
Envir.
UTS
psi
YS
psi
%E1
Ult.
%E1
Failure
B72
High
NONE
NONE
RT
Air
170488
152284
4.4
4.5
B75
Med
NONE
NONE
RT
Air
171822
160218
2.3
2.3
B33
Low
NONE
NONE
RT
Air
169388
157947
2.7
2.7
AVG
170566
156816
3.1
3.2
B66
High
NONE
NONE
540
Air
143158
121203
5.3
11.2
B118
Med
NONE
NONE
540
Air
154056
133815
5.4
11.2
B69
Med
NONE
NONE
540
Air
156444
137782
4.3
8.7
B27
Low
NONE
NONE
540
Air
152209
132779
4.8
12.1
AVG
151467
131395
5.0
10.8
B117
Med
NONE
NONE
540
Vacuum
156120
134868
5.4
12.5
B94
650
100
RT
Air
144436
144436
0
0
B103
650
100
RT
Air
163912
160531
0.655
0.813
AVG
154174
152484
0.33
0.41
B110
650
100
540
Vacuum
155373
133905
5.7
13.7
Bill
650
100
540
Air
156370
135912
4.5
9.9
Figure 4.32 compares the ultimate tensile strength, yield strength and elongation
to failure of both exposed and unexposed tensile samples tested at room temperature and
540°C in air and vacuum.

96
Exposed in Air Unexposed Exposed in Air
Tested in Air Tested in VacuumTested in Vacuum
at 540C at 540C at 540C
Figure 4.32. Ultimate tensile strength, yield strength and percent elongation of exposed
and unexposed samples tested in air and vacuum at room temperature and 540°C.
Fatigue Results
The results from the exposed samples are listed in Table 4-3.
Table 4-3. Low cycle fatigue data of environmentally exposed samples.
Sample
Forging
Exposure
Exposure
Modulus
%8
Alt.
Cycles to
Failure
22
Med.
650
10
18247
0.70
63.86
45241
Gauge
32
Med.
650
100
18443
0.70
64.55
44292
Shoulder
74
Low
650
92hrs t
17948
0.70
62.82
16840
Shoulder
14
Low
650
1
17738
0.75
66.52
37428
Shoulder
C
Med.
650
10
18172
0.80
72.69
5639
Shoulder
31
Med.
650
100
18425
0.85
78.31
1347
Shoulder
30
Med.
650
100
19047
0.90
85.71
287
Shoulder
11
Low
650
10
18515
0.95
87.95
664
Shoulder
Í61
Med.
650
1
18274
0.95
86.80
1364
Gauge
29
Med.
650
100
18275
0.95
86.81
559
Gauge
61
Med.
650
92hrs t
18148
0.95
86.20
265
Gauge
17
Med.
650
92hrs J
18550
0.85
78.84
3033
Shoulder
39
Med.
650
92hrs J
17361
0.85
73.78
1948
Shoulder
t Isothermal exposures for the times listed, £ Isothermal exposures in argon environment for times listed

97
These results are also shown graphically in Figure 4.33 along with the unexposed
results for reference. Each data point is from a single LCF test and so the scatter may be
large [Die86]. The location of failure (G = Gauge, S = Shoulder and T = Thread) is again
shown with the respective data point for each sample. Again, due to the scatter
associated with fatigue testing, a polygon was constructed around the data points to
represent the LCF data as opposed to the construction of a single representative curve.
Lines for each exposure condition were fit to the individual data sets and will be
presented in the discussion.
Figure 4.33. Plot of number of cycles to failure vs. strain. Failure location is noted with
each sample.

98
Discussion
Figure 4.32 shows that the tensile strength of the Ti-21 Al-26Nb is slightly
effected by the environmental exposure. The ultimate tensile strength (UTS) has
decreased by approximately 10% and the yield strength has decreased by approximately
3%. The parameter that is the most effected is the room temperature ductility which,
after exposure, dropped to effectively zero. This effect on ductility is the most
pronounced tensile effect from embrittlement and is the reason for the decrease in UTS
and yield strength. Dary and Pollock [Dar96] evaluated the effects of exposure in air and
vacuum on the tensile properties of Ti-22Al-23Nb. They observed a decrease in UTS and
ductility in all samples exposed in air and, although there was no loss of mechanical
properties in the vacuum-exposed samples, the external surfaces contained a large
number of small cracks. Although Rhodes et al. [Rho93] observed a2 in Ti-22Al-23Nb at
oxygen contents greater than 0.098wt%, Dary and Pollock did not observe any a2
precipitates in the high oxygen regions. They supposed that the 760°C temperature of
exposure was too low to cause a2 precipitation or that the B2 + O phases were
compositionally stable at that temperature. Chesnutt et al. [Che93] conducted tensile
testing on Ti-22Al-27Nb after both a thermal exposure (test samples machined from
exposed bulk material) at 650°C and an environmental exposure (test samples machined
then exposed) at 590°C. After the thermal exposure there was no degradation in the
tensile properties, but after the environmental exposure the samples failed in the threads.
No evaluation was conducted on the environmentally-exposed samples.

99
As seen in Figure 4.33 and again in Figure 4.34, there is an order of magnitude
degradation in LCF life for the exposed samples compared to the unexposed samples.
The data has been plotted in Figure 4.34 using a logarithmic curve fit for each data set.
Each data point in the set was constructed from a single test, so the confidence in the
actual numbers is low. Even so, the trends indicate that as the exposure times increase,
LCF life decreases. For the 92 hour isothermal exposure and the 100 cycle exposure, the
LCF degradation is 2 orders of magnitude. For the single cycle exposure, the degradation
is only 1.5 orders of magnitude.
Figure 4.34. Number of cycles to failure vs. strain for the unexposed and exposed
samples. This construction uses logarithmic curve fits to describe the LCF behavior of
the exposed samples.

100
An SEM backscatter image of the longitudinal section of a 100 cycle exposed
LCF sample after failure is shown in Figure 4.35. This photo was taken at the initiation
side of the fatigue sample. Notice that at the top of the photo the sample contains a dark
region which indicates a high oxygen content(lower atomic number). Areas of high
oxygen have been associated with embrittlement in titanium [Sha68, She86, Kah86,
Wal95], a2 [Bal89, God92, Mei92, Par92, Rak93, Sai93, Smi95, Liit96] and
orthorhombic [Ban93, Che93, Dar95, Dar96] alloys. Godavarti et al. [God92] evaluated
the LCF behavior of Ti-25Al-10Nb-3V-lMo after exposure at 650°C for 95 hours in air
and concluded that this exposure resulted in a reduction in LCF life of up to 2 orders of
magnitude in agreement with the degradation observed in this study.
Figure 4.35. SEM backscatter image of a sample exposed for 100 cycles at 650°C. The
dark region at the top of the photo is an area of high oxygen concentration. The fracture
surface is oriented to the right in this photo.

101
Godavarti et al. also exposed their material at 650°C for 1 hour and observed a
lesser degradation in LCF life. Their data was highly scattered, but the trends agree well
with the trends observed in this study.
It was speculated that on the first LCF cycle cracks formed to full depth of the
embrittled region, thus forming the initiation crack for growth on the first cycle. To
determine if this was indeed the case, two tensile samples were prepared in the same
manner as the fatigue samples by longitudinally polishing to remove circumferential
machining marks. These samples were then exposed in air at 650°C for 100 thermal
cycles in order to duplicate the exposure seen by the fatigue samples. The tensile samples
were then pulled for one cycle at 0.7% strain and 1.2% strain without failing. They were
sectioned longitudinally to look for cracks initiated on the first cycle. For the sample
strained to 0.7% cracks running the full depth of the high oxygen region are clearly
visible (Figure 4.36).
Region
. 10'jjtn
Figure 4.36. Optical micrograph of cross section from exposed tensile sample stretched
to 0.7% strain. Notice cracks run the depth of the oxygen rich zone.
This indicates that the LCF samples cracked on the first cycle during testing, and
that the life was now dependent on fatigue crack growth and no longer on the initiation of
the crack from some microstructural feature. For the sample strained to 1.1% the crack
that is visible not only runs through the high oxygen zone but also into the substrate

102
(Figure 4.37). This indicates that at higher strains, the cracks are even longer and thus the
LCF life may be shortened even more.
Figure 4.37. Optical micrograph of cross section from exposed tensile sample stretched to
1.2% strain. Note the crack runs beyond the oxygen rich region.
Relating the number of cycles to failure to the critical flaw size would help to
better define the mechanism for the decrease in LCF life. The use of the defect tolerant
method (where a defect of known size is expected to propagate at a given rate with failure
determined by fracture mechanics) of lineal' elastic fracture mechanics (LEFM) is valid
for small scale yielding, when the crack tip plastic zone is small relative to the cracked
component and where mostly elastic loading conditions dominate [Sur91], Analyzing the
LCF behavior of Ti-21Al-26Nb would involve generating fatigue crack growth data to
da
obtain -— and an understanding of the fracture toughness, KIC. Both of these values
dN
have been determined for Ti-21Al-26Nb by Woodfield [94Woo] as discussed in Chapter
5. Recalling the equation for the number of cycles to failure, Nf, from Chapter 5
for n * 2 (6.1)
and the solution for a„

103
a» =
m-2
NfCY Act —-— + af
(6.2)
and applying this equation to the uncoated, exposed data set, a„ was calculated for each
data point. Figure 4.38 shows the fatigue data and the calculated a„ for each data point
along with the data from the unexposed material presented in Chapter 5. The value for a„
can typically be related back to some microstructural feature in the material [MÍ197] and
in this case the critical flaw size values for the exposed material fall between 14 and 50
pm. The lower value of a„ at the lower strains may be related to the size of the initial
flaw size introduced on the first fatigue cycle. As related above, when a tensile sample
was strained to a value of only 0.7%, a crack formed in the full depth of the oxygen rich
zone, approximately 10pm. When a tensile sample was strained to a value of 1.2%, the
crack extended below the obvious oxygen rich zone to approximately 12pm. While these
values do not agree with the absolute numbers in the calculation, they do agree with the
trend of shorter lives at higher strains due to the formation of a larger crack on initial
loading. Recalling that the oxygen rich layer is a not a discrete layer, but is more diffuse
due to the nature of oxygen diffusion into the matrix, there may be varying levels of
embrittlement under this layer. Electron microprobe analysis was conducted on the
exposed samples to determine the oxygen profile (Figure 4.39). The curves shown on the
plot are a power law fit to the data points and show how the oxygen profile is a parabolic
function relative to time.

104
Figure 4.38. Uncoated, unexposed and exposed data, showing the failure location and the
calculated a„.
Notice that the 92 hour isothermal exposure is higher in oxygen than the 100
cycle material. Recall that during thermal cycling the furnace took almost 12 minutes to
get to temperature. Over 100 cycles that could account for as much as 20 hours below
650°C which would create a shallower oxygen rich region than a 92 hour isothermal
exposure. The LCF life for the 92 hour isothermal exposure is shorter than that of the
100 cycle exposure, and the calculated a„ is larger. This is in line with a deeper oxygen
rich region (embrittled zone) in the 92 hour sample, as indicated by the microprobe
results.

105
Figure 4.39. Oxygen concentration vs. depth from surface for the exposed fatigue
samples.
Effect of Coating and Exposure on the LCF Behavior
Results
The results of the LCF testing are listed in Table 4-4 and shown graphically in
Figure 4.40 through Figure 4.43. All of the data for the coated samples are included in
Table 4-4, including coating, exposure time and temperature, strain used for testing,
calculated pseudo-stress and number of cycles to failure, Nf. The graphs show the percent
strain used for testing plotted against the number of cycles to failure on a log-log plot.

106
Table 4-4. Coated LCF data.
Exposure
Low Cycle Fatigue Test Results
Spec.
Effective
Coating
Temp.
Cycles
E
%E
Alt.P.
Cycles
Sample
#
Strain
*
S.
Nf
Failure
Forging
°C
#
ksi
ksi
#
Location
55
High
Cr
650
100
16853
0.80
67.41
3638
Gauge
76
Med.
Cr
650
100
17967
0.85
76.36
720
Extenso.
34
Med.
Cr
650
100
17939
0.95
85.21
937
Gauge
B
Med.
Cr
650
92hrs Í
18077
0.70
63.27
4492
Shoulder
73
Low
Cr
650
92hrs J
18399
0.85
78.20
435
Shoulder
33
Med.
Cr
18319
0.80
73.28
110020
Stopped
75
Med.
Cr
18239
0.95
86.64
9356
Gauge
54
High
Cr
17700
1.05
92.93
9961
Shoulder
82
Med.
MgO
650
100
18267
0.95
86.77
47
Gauge
A
Med.
MgO
650
92hrs !
18076
0.70
63.27
3233
Shoulder
40
Med.
MgO
650
100
18027
0.70
63.09
13005
Shoulder
64
Med.
MgO
650
92hrs J
18181
0.85
77.27
990
Shoulder
47
High
Si02
650
100
18520
0.70
64.82
3087
Extenso.
49
High
Si02
650
100
17718
0.80
70.87
1008
Gauge
27
Med.
SiO,
650
100
17687
0.95
84.01
393
Gauge
4
Med.
Si02
650
92hrs t
17757
0.85
75.47
1
Gauge
71
Low
Si02
650
92hrs X
18077
0.70
63.27
691
Shoulder
48
High
Si02
17903
0.70
62.66
15385
Shoulder
28
Med.
Si02
18372
0.85
78.08
3874
Shoulder
26
Med.
SiO,
18295
0.95
86.90
2235
Gauge
56
High
Si02 + Cr
650
100
17635
0.70
61.72
3292
Shoulder
37
Med.
Si02 + Cr
650
100
18091
0.80
72.36
494
Shoulder
79
Med.
Si02 + Cr
650
100
18270
0.95
86.78
30
Shoulder
77
Med.
SiO, + Cr
17941
0.70
62.79
13058
Shoulder
35
Med.
SiO, + Cr
18165
0.85
77.20
3498
Shoulder
78
Med.
Si02 + Cr
18510
0.95
87.92
2288
Gauge
57
High
Si02 + Cr
36
Med.
Si02 removed
18197
0.70
63.69
22288
Shoulder
58
High
Si02 removed
18558
0.95
88.15
173
Gauge
24
High
Ta205
650
100
18189
0.70
63.66
112236
Stopped
38
Med.
Ta,Os
650
100
18083
0.85
76.85
835
Shoulder
59
High
Ta,05
650
100
17363
0.95
82.47
558
Shoulder
3
Med.
Ta,Os
650
92hrs X
17732
0.85
75.36
1
Gauge
72
Low
Ta,Os
650
92hrs X
18240
0.70
63.84
3686
Shoulder
80
Med.
Ta,05
18261
0.95
86.74
4615
Gauge
45
High
Ta205
18271
0.80
73.08
192000
Stopped
66
Med.
Ta,Os
18356
1.05
96.37
3450
Gauge
5
Low
Ta,05
17686
0.80
70.74
142367
Stopped
1
Med.
Ta205
650
100
44
High
Ta,Os+Cr
650
100
17976
0.70
62.92
2920
Gauge
23
High
Ta205+Cr
650
100
18079
0.80
72.32
648
Gauge
65
Med.
Ta205+Cr
650
100
18194
0.95
86.42
233
Shoulder

107
I— 1 '—
— 1—■ ■ ' "
TTT ' |
—O— Si02 coated, unexposed
♦ Si02 coated, exposed
• Si02 coated, 92hrs, Argon
—B— Si02+Cr coated, unexposed
â–  Si02+Cr coated, exposed
—— Si02 coating removed
â– 
■ ♦
• m
10 100 1000 10* 10s
LCF Cycles to Failure
Figure 4.40. Si02 coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material.
The Si02 data includes the Cr overcoat samples. These samples were first coated
with SiO, by CVD then removed and coated with Cr in the vacuum sputter unit. Notice
that the SiO, and the Si02 + Cr as coated data correlate very well. Recall that the Cr is
deposited at low temperature (200°C) relative to the Si02 (700°C), so the data is really for
Si02 as coated. Exposure in air increases the degradation of the Si02 coated samples
while the argon environment exposure degrades the LCF life the most.

108
Figure 4.41. Ta^Oj coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material.
In the case of the Ta205 coating, there is a slight degradation in the LCF life after
coating, but not as large as in the case of Si02. On exposure to air the degradation is
similar to that of uncoated material, but in argon again the degradation is much worse. A
Cr overlay coating, similar to that used in the Si02 case, degrades the LCF life more than
just the Ta^ by itself. Again, there is a large amount of scatter in the data, and the
Ta^ coated sample exposed to air and tested at 0.7% strain was a runout test, where the
uncoated samples failed before 1x10s cycles.

109
Figure 4.42. MgO coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material.
MgO was applied at the same temperature as Ta^Oj during CVD, therefore, the as-
coated behavior is expected to be the same and no tests of as-coated samples were
conducted. The MgO samples were cyclically exposed to air and statically exposed to
argon and, in both instances, there was a degradation over uncoated. In the case of MgO,
the argon exposure did not degrade the LCF life more than the air exposure as it did for
material coated with Si02 and TajOs.

110
Figure 4.43. Cr coated LCF data. Plot includes the polygons drawn for the baseline
uncoated unexposed and exposed material.
Since Cr is sputter deposited on the LCF samples in a vacuum at temperatures
below 200°C, the as-coated material experiences no degradation in the LCF life. After
cyclic exposure in air, the degradation is in the scatter for the uncoated material. After
static exposure in argon, the degradation is more than that from cyclic exposure in air, but
is similar to than of cyclically exposed Si02 + Cr and Taj05 + Cr coated material.
Discussion
The degradation from the coated samples is consistent with the results from the
uncoated exposed samples. The lack of degradation in LCF life for the Ta^ and Cr
samples in the as coated condition is due to the low temperature of deposition for both

Ill
coatings. The degradation seen in the as coated Si02 samples, including the Cr
overcoated samples, is due to the high temperature of deposition of the Si02, (700°C) in a
partial vacuum. It is evident from the single cycle exposured samples that not much
oxygen is needed for embrittlement to occur. The reasons for the increase in degradation
in LCF of the Si02 and MgO coated samples over the uncoated exposed samples is not
clear. The Si02 coated samples have a maximum of approximately 5 at% oxygen
diffused in below the surface (Figure 4.44) which is obviously enough to cause
embrittlement when compared to the value of the single thermal cycled sample of 1.5
at%, but less that that of the 100 cycle exposed samples at 20 at%.
Figure 4.44. Depth of oxygen concentration in Si02 coated LCF samples exposed in air
and argon.

112
The coatings could possibly provide detrimental alloying elements like Si or Mg
to the substrate, although EMPA conducted on the exposed LCF samples did not indicate
an increased level of these elements under the coatings.
The decrease in LCF life for Ta^ and Cr coated and exposed samples is within
the scatter of the uncoated samples. It is apparent from the SEM and Auger evaluations
discussed in Chapter 4 that the titanium substrate is reducing the oxide coating and a
subsequent increase in oxygen is occurring in the substrate underneath the coating. The
Ta205 LCF sample exposed in an inert environment contained almost as much oxygen
beneath the surface as the sample exposed in air (Figure 4.45), even though the TajOj
coated material gained weight like the uncoated in cyclic oxidation. This indicates that
the substrate is reducing oxygen from the coating.
Figure 4.45. Depth of oxygen concentration in Taj05 coated LCF samples after exposure
in air and argon.

113
The Cr coated LCF samples also showed an increase in oxygen under the surface
in the air exposed sample, but still well below the uncoated, exposed for 100 cycles
(Figure 4.46) which had a longer LCF life. The argon exposed sample showed no
increase in oxygen. Both the argon and air exposed samples showed a slight increase in
Cr under the coating (Figure 4.47), but to a depth less than that of the oxygen in the
uncoated, exposed samples. Even if Cr caused embrittlement similar to oxygen, the
shallow depth of Cr diffusion does not support the concept of a crack forming on the first
cycle in the brittle region being the cause of the reduction in LCF life. Examination of
the failed LCF sample (Figure 4.48) shows that secondary cracks, those that formed but
did not fail, ran into regions of less than 0.1 at% Cr concentration. This concentration of
Cr may be enough to cause embrittlement to occur and create a crack to this depth on the
first loading cycle.
Figure 4.46. Depth of oxygen concentration in Cr coated LCF samples exposed to air and
argon.

114
Figure 4.47. Cr concentration depth in the Cr coated LCF samples exposed in air and
argon.
Secondary
Cracks
Figure 4.48. SEM backscatter image of Cr coated LCF sample exposed in air for 100
cycles at 650°C.

115
The differences in coefficient of thermal expansion (CTE) between the coatings
and the matrix could impose residual stresses in both materials. Figure 4.49 shows the
variation in CTE with respect to temperature. The CTE of Si02 is approximately 20
times less than the substrate, Ta,05 is approximately half and MgO is 1.5 times greater.
All of the oxide coatings exhibited the same LCF behavior after exposure.
Figure 4.49. Coefficient of thermal expansion comparison between coatings and Ti-
22Al-26Nb substrate.
Cr has a CTE similar to the substrate, yet after cyclic exposure the Cr coated
samples had lower LCF lives than most coatings. This lack of correlation between the
CTE and effect of coatings on LCF life indicate that the residual stress from the CTE

116
mismatch that may exist after coating is overwhelmed by the effect of the embrittled layer
under the coatings.
The differences in elastic modulus between the coatings and substrate may
contribute to the LCF behavior. The Cr coating has a higher modulus than that of the
substrate (40 Msi vs. 18 Msi) and thus when the coating and substrate are loaded in
tension, the coating experiences a much higher stress (over two times) than that of the
substrate. This high stress in the coating could lead to failure of coating on the first LCF
cycle. Even if the coating were cracked on the first LCF cycle, it is only 1pm in
thickness, much thinner than the embrittled region for the 1 thermal cycle sample, yet the
LCF life was closer to that of the 100 thermal cycled samples with a much thicker
embrittled region.

CHAPTER 5
SUMMARY AND CONCLUSIONS
The orthorhombic titanium aluminide alloy Ti-22Al-26Nb was evaluated for its
response of oxidation, tensile and low cycle fatigue to exposure to an ambient air
environment at 650°C both uncoated and with a variety of coatings.
Oxidation
The cyclic oxidation behavior of Ti-22Al-26Nb (at%) at 650°C is improved by the
application of FeCrAlY, NiCrAlY+Al203> Al3Ti, Ti5Si3 and Cr. Si02 provides initial
protection from oxidation, but after 400 cycles the oxidation behavior departs from
parabolic and becomes more linear. Ta^ and MgO coatings provide no oxidation
protection.
Microhardness and oxygen levels measured by electron microprobe analysis
correlated well in samples that exhibited poor oxidation behavior.
Al3(Ti,Nb) and TÍ5SÍ3 coatings behaved well in oxidation, but are brittle and
contained cracks after thermal cycling.
FeCrAlY and NiCrAlY + AI2O3 coatings also behaved well in oxidation, but Fe
and Ni diffused into the substrate and are associated with increased hardness
(embrittlement) at the surface.
117

118
Pt behaved poorly in oxidation, probably due to the poor coating quality. Pt also
diffused into the substrate, but any effect on hardness was masked by the presence of
oxygen.
Cr reduced the oxidation rate of the substrate but allowed oxygen to diffuse into
the matrix thus creating an embrittled zone at the surface.
Si02 allowed oxygen to penetrate, and after only 400 hours. No Si diffused into
the substrate, but it appeared that the Si02 was reduced by the titanium and thus oxygen
was allowed to diffuse into the substrate.
Ta^ and MgO coatings provided minimal protection against oxygen ingress
during thermal cycling.
Forging Location
The tensile and low cycle fatigue (LCF) behavior of the Ti-22Al-26Nb alloy
forged above the B2 transus and heat treated at 815°C for four hours is independent of the
location in the pancake forging (i.e. the grain aspect ratio).
Variations in the grain aspect ratio in the forging are due to the variations in the
effective strain induced during forging. This effective strain is a function of friction,
material flow and temperature variation throughout the forging.
There was little or no variation in the orthorhombic lath size or distribution in the
forging due to any radial variations in cooling.
LCF behavior occurs by initiation of cracks followed by fatigue crack growth.
Since fatigue crack growth appears to be strongly dependent on orthorhombic lath size,

119
and lath size does not vary with location in the forging, neither does the LCF behavior.
LCF behavior is also independent of grain aspect ratio in this forging.
All samples from the forging can be used to compare to one another with out
regard to previous forging history (location).
Environmental Exposure
Cyclic exposure in air at 650°C for 100 thermal cycles (1 cycle being 55 minutes
at 650°C, 5 minutes at room temperature) degrades the LCF life by approximately 2
orders of magnitude and results in an essential elimination of tensile ductility. A single
cycle degrades the LCF life by 1.5 orders of magnitude. Therefore, there is not a
parabolic relationship with LCF behavior and exposure time as diffusion of oxygen
would suggest.
Linear elastic fracture mechanics (LEFM) can be used to explain the LCF
behavior of exposed Ti-22Al-26Nb by predicting an increased initial flaw size for the
exposed material relative to the unexposed material.
The reduction in tensile ductility results from an embrittled region in the surface
of the exposed test sample; This embrittled region is caused by the ingress of oxygen
during high temperature exposure to air. This embrittled region produces cracks to the
full depth of oxygen penetration on the first LCF cycle thus forming a crack of length a,
to be formed. Shorter exposure times create shallower embrittled zones thus forming
shorter cracks on the first LCF cycle. These shorter cracks yield longer LCF life because
the crack must grow to a longer length for final failure.

120
Coated LCF Behavior
LCF life after cyclic exposure is not improved by coating with Si02, TaADs, MgO
or Cr. Si02 coated samples exhibited a reduction in LCF life after coating similar to
uncoated samples after cyclic exposure.
The oxide coatings evaluated in this study do not provide protection from
embrittlement for orthorhombic titanium aluminide due to the high solubility of oxygen
in these alloys and the subsequent reduction of oxides in contact with the substrate at
elevated temperatures.
Oxide coatings deposited by MOCVD are chemically reduced by the TÍ-22A1-
26Nb substrate at 650°C. This reduction creates a layer of surface embrittlement due to
the high oxygen content in the titanium alloy. The reduction of the oxides is due to the
large reduction in free energy associated with the solubility of oxygen in titanium.
Sputtered Cr does not cause a degradation in LCF properties in the as-coated state.
After exposure in air for 100 cycles, the degradation is similar to that of uncoated
exposed material. After an isothermal exposure in argon for 92 hours the reduction in
LCF life is greater than that from the cyclic exposure in air.
No coating was identified in this study that would prevent surface embrittlement
after exposure in air at 650°C. This embrittlement leads to a reduction in LCF life and
tensile ductility.

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BIOGRAPHICAL SKETCH
James Ross Dobbs was bom on 6 September 1962 in Birmingham, Alabama,
USA. He attended Shades Valley High School in Birmingham, and graduated in 1980.
He then enrolled at the University of Alabama at Birmingham in the mechanical
engineering department. After several courses in mechanical engineering, he was
introduced to the field of materials engineering by Professor Ray Thompson and quickly
changed majors. He received a Bachelor of Science degree in materials engineering in
1985 and pursued his master’s degree until he obtained a job at GE Aircraft engines in
1987. He married Jane Gray Richardson in 1988 and subsequently completed his
master’s thesis in early 1991 and received his Master of Science degree in materials
science and engineering. He took a leave of absence from GE in late 1991 and enrolled at
the University of Florida as a doctoral student under the direction of Prof. Mike Kaufman.
He completed the course work and passed the qualifying exams for doctoral candidacy
while at U.F., then returned to GE Aircraft Engines to complete his research. He took a
position at the GE Corporate Research and Development Center in 1994 to complete his
research under the direction of Dr. Mike Gigliotti.
132

I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and quality,
as a dissertation for the degree of Doctor of Philosophy.
Professor of Materials Science and
Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and quality,
as a dissertation for the degree of Doctor of Philosophy.
Robert T. DeHoff //
Professor of Materials Science and
Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and quality,
as a dissertation for the degree of Doctor of Philosophy.
Fereshteh Ebrahimi
Associate Professor of Materials
Science and Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and quality,
as a dissertation for the degree of Doctor of Philosophy.
Metallurgist
GE Corp. Research and Development

I certify that I have read this study and that in ray opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and quality,
as a dissertation for the degree of Doctor of Philosophy.
Paul H Holloway
Professor of Materials Science and
Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and quality,
as a dissertation for the degree of Doctor of Philosophy.
Ellis D. Verink, Jr.
Distinguished Service Professor
Emeritus of Materials Science and
Engineering
I certify that I have read this study and that in my opinion it conforms to
acceptable standards of scholarly presentation and is fully adequate, in scope and quality,
as a dissertation for the degree of Doctor of Philosophy.
Ashok KumSr^
Assistant Professor of Mechanical
Engineering
This dissertation was submitted to the Graduate Faculty of the College of
Engineering and to the Graduate School and was accepted as partial fulfillment of the
requirements for the degree of Doctor of Philosophy.
December, 1997
Winfred M. Phillips
Dean, College of Engineering
Karen A. Holbrook
Dean, Graduate School

LD
1780
1991/
UNIVERSITY OF FLORIDA
3 1262 08556 6015




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