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Record for a UF thesis. Title & abstract won't display until thesis is accessible after 2011-12-31.

Permanent Link: http://ufdc.ufl.edu/UFE0041122/00001

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Title: Record for a UF thesis. Title & abstract won't display until thesis is accessible after 2011-12-31.
Physical Description: Book
Language: english
Creator: Ganesan, Krishna
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2009

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Subjects / Keywords: Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
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theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
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Electronic Thesis or Dissertation

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Statement of Responsibility: by Krishna Ganesan.
Thesis: Thesis (Ph.D.)--University of Florida, 2009.
Local: Adviser: Fuchs, Gerhard E.
Electronic Access: INACCESSIBLE UNTIL 2011-12-31

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Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2009
System ID: UFE0041122:00001

Permanent Link: http://ufdc.ufl.edu/UFE0041122/00001

Material Information

Title: Record for a UF thesis. Title & abstract won't display until thesis is accessible after 2011-12-31.
Physical Description: Book
Language: english
Creator: Ganesan, Krishna
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2009

Subjects

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

Notes

Statement of Responsibility: by Krishna Ganesan.
Thesis: Thesis (Ph.D.)--University of Florida, 2009.
Local: Adviser: Fuchs, Gerhard E.
Electronic Access: INACCESSIBLE UNTIL 2011-12-31

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2009
System ID: UFE0041122:00001


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1 ON THE MICROSTRUCTURE, HIGH TEMPERATURE OXIDATION AND MECHANICAL BEHAVIOR OF RARE EARTH MODIFIED CM 247LC, A NI CKEL BASE POLYCRYSTALLINE SUPERALLOY By KRISHNA PRAKASH GANESAN 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 2009

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2 2009 Krishna Prakash Ganesan

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3 To my parents

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4 ACKNOWLEDGMENTS I take this opportunity to express my sincere gratitude to various wonderful people without whom this dissertation would not have been possible. At the outset, I would like to acknowledge my doctoral advisor, Dr. Gerhard E. Fuchs for his tutelage, support and encouragement. I derived my motivation and inspiration during this research from his unrelenting passion towards high temperature alloys. I would like to thank the professors in my supervisory committee (Dr. Paul Holloway, Dr. Martin Glicksman, Dr. Ame lia Luisa Dempere and Dr. Tim Anderson) for their valuable inputs in this research. I would like to thank the Materials Systems Development group at Siemens Power Generation, Inc., Orlando, FL for their financial assistance in this project. Special thanks to Cynthia Klein for her incredible support and patience in promptly responding to innumerable e mails and phone calls about the materials, testing, etc. Her contribution and interest made sure the project was on track all the time. Also, I would like to thank Dr. Allister James, Dr. Anand Kulkarni and Dr. Sachin Shinde for their valuable inputs during the course of the project. I would like to acknowledge Dr. Bruce Pint at Oak Ridge National Lab for his suggestions and his contributions to the field of ox idation have served as a bigger impetus to my research. During my stay in Gainesville, outside the office and apartment, I spent most of my time at Major Analytical and Instrumentation Center (MAIC) working on the characterization. I am grateful to Dr. Ame lia Dempere for introducing me to microstructural analysis and her inputs in the early stages about characterization of ppm levels; heartfelt thanks to Dr. Gerald Bourne for stimulating interest in TEM, help with FIB and valuable discussions about research; Mr. Wayne Acree for his help with MicroProbe Analysis and for making sure that I get what I asked for; Ms. Kerry Siebien for timely support with HR TEM; Mr. Eric Lambers for his help and inputs in Auger Electron Spectroscopy; Dr. Maggie Puga Lambers at Micro Fabritech for her help with

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5 SIMS. The MAIC team had been so incredibly helpful and friendly towards me, especially, Rosabel Ruiza Rosabel for her patience and understanding towards my habit of holding onto some of the facility keys indefinitely. I t ake this moment to thank the marvelous support, enthusiasm and friendship of all the fellow students in the High Temperature Materials Lab, both past and present Andrew Wasson and Brandon Wilson for their intellectual inputs and critical assessment during my research which helped me to set the bar always high; Brendan Collins for the interesting multi cultural discussions and introducing me to the amazing game of football; Alvaro Mendoza, Phillip Draa and Justin Cretty for their help. I am extremely privi leged to gain a lot of wonderful friends during my time in graduate school with whom I had the some of the best times in Gainesville Mahesh Tanniru, Rakesh Behera, Sankara Sarma Tatiparti, Shobit Omar, Abhijit Pramanick, Parag Katira, Saurabh Morarka, Pankaj Nerikar, Dilpuneet Aidhey, Priyank Shukla, and Karthik Padmanabhan to name a few. I am grateful to my brother Arul Prakash Ganesan for his encouragement and motivation throughout my life. My fiance, Saritha deserves a special mention for her love, understanding and making sure that I stayed inspired through all these four years. Lastly but more importantly, everything in life that I have attained till this time would not been possible without my parents. They have worked so hard to make sure that I ge t the best in everything. They gave me total freedom, supported all my decisions, taught me to dream big and laid a strong foundation for my life based on commitment to hard work, dedication and aspiration. To acknowledge their unconditional love and support in all walks of my life, I dedicate this dissertation to my parents.

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6 TABLE OF CONTENTS page ACKNOWLEDGMENTS ...............................................................................................................4 LIST OF TABLES .........................................................................................................................10 LIST OF FIGURES .......................................................................................................................11 ABSTRACT ...................................................................................................................................22 CHAPTER 1 INTRODUCTION ..................................................................................................................24 1.1 Hydrogen Powered Gas Turbines .....................................................................................24 1.2 Approaches towards Alloy Development .........................................................................25 1.3 Understanding the Rare Earth Effect ................................................................................25 1.4 Organization of the Dissertation .......................................................................................26 2 BACKGROUND ....................................................................................................................30 2.1 Development of Ni base Superalloys ...............................................................................30 2.1.1 Improvements in Alloy Processing ........................................................................31 2.1.2 Alloying and the Strengthening Mechanisms .........................................................32 2.1.3 Thermal Barrier Coatings .......................................................................................33 2.2 Mechanism of Oxide Scale Formation .............................................................................34 2.2.1 Scale formation in Ni Cr Al alloys ........................................................................34 2.2.2 Polymorphic Forms of Alumina .............................................................................37 2.3 Me chanistic Understanding of Rare Earth Effect .............................................................38 2.3.1 Improved Chemical Bonding at the Interface ........................................................39 2.3.2 Graded Seal Mechanism .........................................................................................40 2.3.3 Pegging Effect ........................................................................................................40 2.3.4 Scale Plasticity ........................................................................................................41 2.3.5 Change in Oxidation Kinetics ................................................................................42 2.3.6 Vacancy Sink Mechanism ......................................................................................43 2.4 Effect of Sulfur on the Oxidation Behavior ......................................................................44 2.5 Sulfur Embrittlement ........................................................................................................46 2.6 Effect of RE on the Mechanical Properties ......................................................................47 3 EXPE RIMENTAL TECHNIQUES ........................................................................................56 3.1 Materials ...........................................................................................................................56 3.1.1 Rare Earth Additions ..............................................................................................57 3.1.2 Heat Treatments ......................................................................................................58 3.1.3 Hot Isostatic Pressing (HIPing) ..............................................................................59 3.2 Oxidation ..........................................................................................................................59

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7 3.2.1 Isothermal Oxidation Testing .................................................................................61 3.2.2 Cyclic Oxidation Testing ........................................................................................62 3.3 Mechanical Testing ...........................................................................................................63 3.3.1 Tensile Testing .......................................................................................................63 3.3.2 Creep Testing ..........................................................................................................64 3.4 Microstructural C haracterization ......................................................................................65 3.4.1 Metallographic Preparation ....................................................................................65 3.4.2 Optical Microscopy ................................................................................................66 3.4.3 Scanning Electron Microscopy ...............................................................................66 3.4.4 Electron Probe Micro Analysis (EPMA) ................................................................66 3.4.5 X Ray Diffraction ...................................................................................................67 3.4.6 Auger Electron Spectroscopy (AES) ......................................................................67 3.4.7 Secondary Ion Mass Spectrometry (SIMS) ............................................................68 3.4.8 Transmission Electron Microscopy (TEM) ............................................................68 4 MICROSTRUCTURE ............................................................................................................74 4.1 Introduction .......................................................................................................................74 4.2 As Cast Microstructural Analysis .....................................................................................74 4.2.1 Inter dendritic Spacing Measurements ...................................................................75 4.2.2 Partitioning Studies Using Micro Probe Analysis .................................................76 4.3 Development of a Multi Step Solution Heat Treatment ...................................................77 4.3.1 Baseline Solution Heat Treatments ........................................................................78 4.3.2 Differential Thermal Analysis (DTA) Results .......................................................79 4.3.3 Design of Modified Solution Heat Treatments ......................................................80 4.4 Rare Earth Hafnium Phase .............................................................................................82 4.5 Summary ...........................................................................................................................85 5 OXIDATION BEHAVIOR OF UNCOATED ALLOYS ......................................................99 5.1 Isothermal Oxidation Testing .........................................................................................100 5.1.1 Oxidation Kinetics ................................................................................................100 5.1.2 Gas/Scale interface microstructure .......................................................................103 5.1.3 Scale/alloy interface microstructure .....................................................................107 5.1.4 XRD analysis of oxide phases ..............................................................................108 5.1.5 RE segregation to the oxide scale .........................................................................109 5.2 Initial Stages of Oxide Sca le Development ....................................................................109 5.2.1 Microstructure and Phase Analysis ......................................................................110 5.2.2 Analysis of Rare Earth Oxide Precipitation ......................................................110 5.3 Cyclic Oxidation Behavior .............................................................................................111 5.3.1 Oxidation Kinetics ................................................................................................112 5.3.2 Oxide Microstructure ............................................................................................114 5.3.3 Internal Oxidation Pegs .................................................................................115 5.3.4 Gamma Prime Denuded Zones (GDPZ) .......................................................116 5.4 Summary .........................................................................................................................117

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8 6 OXIDATION BEHAVIOR OF ALLOYS WITH THERMAL BARRIER COATINGS ....134 6.1 As Deposi ted Microstructure of the Coating ..................................................................135 6.2 Isothermal Spallation Testing .........................................................................................138 6.3 Microstructural Changes after Oxidation .......................................................................139 6.3.1 Bond Coat / Superalloy substrate interface ..........................................................140 6.3.2 Thermally Grown Oxide (TGO) Layer ................................................................141 6.3.2.1 Hafnium Oxide Stringers ...........................................................................142 6.3.2.2 X Ray Phase Analysis of YSZ/TGO interface ...........................................142 6.3.2.2 T hickness of TGO Layer ............................................................................143 6.3.2.3 MicroProbe Analysis on TGO BC thickness ............................................144 6.4 Segregation Analysis on the TGO Layer ........................................................................145 6.4.1 SIMS Analysis on the TGO layer .........................................................................145 6.4.2 TEM Analysis of the TGO Layer .........................................................................146 6.5 Summary .........................................................................................................................147 7 TENSILE BEHAVIOR ........................................................................................................158 7.1 Introduction .....................................................................................................................158 7.2 Phase I Un HIPed Alloys .............................................................................................159 7.2.1 Industrial HT ........................................................................................................159 7.2.2 UFHT ....................................................................................................................160 7.2.3 Fractography .........................................................................................................160 7.3 Phase I HIPed Alloys ...................................................................................................161 7.3.1 UFHT ....................................................................................................................161 7.3.2 Fractography .........................................................................................................163 7.4 Comparison between HIPed and UnHIPed Alloys .......................................................164 7.5 Phase II Alloys ................................................................................................................165 7.5.1 Industrial HT ........................................................................................................165 7.5.2 UF HT (Mod SHT 03) ..........................................................................................166 7.5.3 Fractogra phy .........................................................................................................168 7.6 Summary .........................................................................................................................169 8 CREEP BEHAVIOR ............................................................................................................183 8.1 Introduction .....................................................................................................................183 8.2 Phase I Alloys .................................................................................................................184 8.2.1 UFHT Un HIPed Alloys .......................................................................................184 8.2.2 UF HT HIPed Alloys .............................................................................................184 8.3 Phase II Alloys ................................................................................................................185 8.3.1 IHT Alloys ............................................................................................................185 8.3.2 UFHT (Mod SHT 03) Alloys ............................................................................187 8.4 Temperature Dependence of Creep Behavior .................................................................188 8.5 Summary .........................................................................................................................189

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9 9 THE RARE EARTH EFFECT .............................................................................................197 9.1 RE Effect on the Oxidation Behavior of Uncoated Alloys .............................................197 9.1.1 Thermodynamic Affinity towards Oxygen ..........................................................198 9.1.2 Ionic Size Misfit and Its Effect on Segregation ....................................................199 9.1.3 Segregation t o Oxide Grain Boundaries and Its Effect on Oxidation Mechanism .................................................................................................................200 9.1.4 Precipitation of RE Oxides at Gas/Oxide Interface ..............................................202 9.1.5 Optimum levels of RE additions ..........................................................................203 9.1.6 Advantages of CoAlloying ..................................................................................205 9.1.7 Effect of Oxidation Temperature ..........................................................................206 9.1.8 Effect of Hf on the Oxidation Behavior ...............................................................206 9.1.8.1 Internal Oxidation Stringers .......................................................................207 9.1.8.2 Role of Hf on the Oxidation Kinetics .........................................................208 9.2 RE Effect on the Performance of Thermal Barrier Coatings ..........................................209 9.2.1 Improvement in the Spallation Lifetimes .............................................................209 9.2.2 Mechanisms of Coating Failure ............................................................................209 9.2.3 Effect of RE Elements on Therm ally Grown Oxide Layers .................................212 9.2.3.1 Is RE Diffusion from the Substrate to TGO possible? ...............................213 9.2.3.2 Relating Hf Effect to RE Eff ect .................................................................214 9.3 RE Effect on the Mechanical Deformation Behavior .....................................................215 9.3.1 How are RE Different? .........................................................................................215 9.3.2 RE Effect on the Grain Boundary Sliding ............................................................216 9.3.3 Segregation to Stacking Faults .............................................................................219 9.3.4 Solute Drag on the Dislocation Movement ..........................................................220 9.3.5 Why were RE Additions Detrimental to Creep Behavior? ..................................221 10 CONCLUSIONS AND FUTURE WORK ...........................................................................232 10.1 Microstructure ...............................................................................................................232 10.2 Oxidation Behavior .......................................................................................................233 10.3 Mechanical Properties ..................................................................................................234 10.4 Future Work ..................................................................................................................234 LIST OF REFERENCES .............................................................................................................236 BIOGRAPHICAL SKETCH .......................................................................................................253

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10 LIST OF TABLES Table page 11 Comparison between Pulverized Coal Combustion (PCC) and Integrated Gasificatio n Combined Cycle (IGCC) in terms of maximum gas temperature, efficiency of power generation and emission levels. .........................................................28 31 As cast composition of baseline CM247LC alloy (Wt %). ...............................................70 32 Composition of 16 alloys in Phase I based on the RE element additions. .........................70 41 List of various heat treatment cycles used on CM247LC allo ys. ......................................89 410 Eutectic temperatures for various Ni RE systems taken from their phase diagrams. ........92 43 Differential Thermal Analy sis (DTA) results showing the various transformation temperatures. RE additions did not result in large reduction in the solidus temperatures. ......................................................................................................................93 44 List of vacuum annealing treatments conducted to understand the RE Hf phase evolution. ...........................................................................................................................97 53 Parabolic weight loss rates for regime I in cyclic oxidation testing at 1121 oC. Note the Kp values are in negative. ...........................................................................................129 91 Relative thermodynamic affinities of various elements towards oxygen at 1400K. .......224 92 Ionic size misfit ratios for various RE elements use d in this research. A ratio greater than 1.5 usually indicates that the element segregates very effectively to the grain boundary regions. .............................................................................................................224 93 Properties of various constituent in a TBC s ystem (Reproduced with permission from [50]). .................................................................................................................................229 94 Differences in the atomic radii between the RE elements and Ni matrix. Larger atomic radius of RE elements drives their segregation to th e grain boundary regions. ...231

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11 LIST OF FIGURES Figure page 11 Chart representing the relative percentages of world coal reserves as of 2005. The numbers represent the quantity of recoverable coal in million short tons (Reproduced with permission from [2]). ...........................................................................28 12 The roadmap for hydrogen powered gas turbines drafted by Siemens Power G eneration. The work performed at University of Florida is highlighted. (Reproduced with permission from [5]). ...........................................................................29 21 Improvements in the turbine inlet temperatures from the first known Whittl e W1 engine to the modern day Trent 900 engines. (Reproduced with permission from [153]). .................................................................................................................................50 22 Creep behavior of Mar M200 alloy in the conventionally cast, directionally solidified and sing le crystal conditions Superior lifetimes and higher ductility observed for single crystals (Reproduced with permission from [27]). .................................................50 23 Ni Cr Al ternary phase diagram illustrated the regio ns where Group I, II and III oxidation are dominant. (Reproduced with permission from [17]). ..................................51 24 Oxide scale development during Groups I, II and III oxidation. In the schematic, S represen ts spinel oxides of type Ni(Cr, Al)2O4 dark circ les are Cr and empty circles Al. (Reproduced with permission from [57]). ...................................................................51 25 Arrhenius plot showing the slower growth kinetics of Al2O3 co mpared to NiO and Cr2O3 scale at all the temperatures of oxidation. ...............................................................52 26 Crystal structures of Al2O3 polymorphs namely, phase and phase ............................52 27 NiAl alloy at 1150 oC showing the beneficial effects of Zr and Y additions on the spallation resistance compared to the baseline allo y. (Reproduced with permission from [9] ). ..........................................................................53 28 Al2O3 NiAl alloys at 1200 oC oxidation in unalloyed and Zr containing alloy s. (Reprod uced with permission from [9]). .........................................................................................................53 29 NiAl at 1100 oC showing the detrimental effect of S on the spallation properties and beneficial effect of Y in g ettering S (Reproduced with permission from [9]). ...........................................................................54 210 Detrimental effects of ppm levels of S in a Ni showing the ductility minimum in the temperature range of 600 1000 oC (Reproduced w ith permission from [132]). ............54

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12 211 Cavitation observed on the grain boundaries in Fe+60 ppm S tensile tested at 700 oC and the corresponding X ray map taken in SEM showing heavy concentratio n of S in the cavities. (Reproduced with permission from [154]). ...................................................55 212 Reduction in the minimum creep rate in NiCr alloys due to the addition of ppm levels of Ce. (Reproduced with permissi on from [138]). ..................................................55 31 A typical fully prepared alloy coupon for oxidation testing. The edges have been rounded to reduce the stress effects during hightemperature oxide growth. ....................70 32 The isothermal oxidation testing procedure showing the samples and the time intervals for removals for the weight measurements. Similar procedures were used for oxidation testing at 1010 oC and 1079 oC. ...................................................................71 33 A schematic illustration of a cyclic oxidation profile. A heating and cooling rate of 25 oC/min were used during the testing and the time per cycle including the heating, soaking and cooling is 24 hours. ........................................................................................71 34 An engineering drawing of a mechanical testing specimen showing the important dimensions. Pin holes were machined into the shoulders to use the screw type ex tensometers. ....................................................................................................................72 35 Larson Miller parameter plotted against the stress for Mar M247 alloy. The properties of equiaxed CM247LC are very similar to its parent alloy Mar M247. ...........72 36 Images taken in the secondary electron mode showing the various sequential steps in the preparation of TEM foils using FIB. The steps shown are deposition of a thin Pt strip followed by milling of trenc hes and final thinning. ...................................................73 41 Optical micrographs showing the dendritic stru cture in the as cast alloys for CM247LC baseline and CM247LC + 300 ppm Pr. The longer dendritic structure radiatin g across the cross section are typical of a slowly cooled structure. ......................86 42 SEM micrographs showing the structure of the as cast dendritic structure of CM247LC+860 ppm Ce. H igher magnificatio n view of the box showing inter dendritic structure showing the carbides, ....................87 43 Comparison of the average Secondary Dendritic Arm Spacing (SDAS) CM247LC alloys with various RE additions. A refinement of 34 m was observed due to RE additions. ............................................................................................................................87 44 Micro probe analysis results of RE distribution in the as cast CM247LC alloys. Note that the detected wt % values are higher than the actual alloy composition. .....................88 45 Dendritic and inter dendritic partitioning behavior of various elements in the as cast CM247LC+860 ppm Ce and CM247LC+300 ppm Pr alloys. RE elements exhibit a strong tendency to segregate to the inter dendritic regions. ..............................................88

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13 46 Micro probe analysis line scan results for heavy elements like W, Cr, Co and Ta in as cast CM247LC alloys. Large scale inhomogeneity in the chemical distribution is noticed. ...............................................................................................................................89 47 SEM mic rographs of IHT alloys showing the inter dendritic regions. Large amount ........................90 48 Micro probe analysis showing the distribution of heavier e lements like W, Co, Cr and Ta in IHT alloys. In comparison with the as cast results, improvements in the compositional homogeneity were attained. ........................................................................90 49 SEM micrograph showing t he structure o f UFHT alloys. Incipient melting and c racking along the eutectic regions due to the nonuniform deformation during creep testing are observed ...........................................................................................................91 410 A typical Differential Thermal Analysis (DTA) curve for a CM247LC baseline alloy showing the important transformation temperatures. ........................................................91 411 Ni Al phase diagram with the hatched regions showing the Ni rich regions (Reproduced from A SM Handbook on Binary Alloy Phase Diagrams). ...........................92 412 SEM micrographs showing the structure of alloys subjected to ModSHT03 solution heat treatment. Compared with the previous heat treatment, a higher degree of microstructural homogeneity was achieved through this modified heat treatment. ..........93 413 Micro probe analysis showing the distribution of heavier elements like W, Co, Cr and Ta in the alloys subjected to ModSHT03. This modified heat treatment resulted in higher degree of compositional homogeneity also. .......................................................94 414 SEM micrographs showing the fully heat treated micr ostructure illustrating the .....................................................................................................94 415 SEM micrographs showing the Hf Ce Gd phase in the alloys subjec ted to annealing heat treatment. The phase in the middle is a MC t ype carbide rich in Hf. ........................95 416 SEM micrographs and X ray mapping images showing the regions richer in Ni, Pr and Hf. Complete absence of Ni is evident in the X ray maps. .........................................95 417 SEM micrographs showing the presence of Hf Ce Dy phase in the grain boundaries of the alloy subjected to ModSHT02 solution heat treatment. .........................................96 418 FIB images showing the Hf Ce Dy p hase at the grain boundaries. A three dimensional view of the blocky phase is illustrated in this image. ....................................96 419 TEM micrograph showing the Hf Ce Dy phase in Bright field mode and Dark field mode. Interestingly, a bi crystal grain was observed and better grain contrast was obtained using tilting to different zone axis in the TEM. ..................................................97

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14 420 TEM micrograph showing the grain boundary of the Hf Ce Dy phase and the corresponding EDAX spectrum taken inside one of the grains Cu signals were from the TEM sample grid. ........................................................................................................98 51 Specific mass gain results for 1000 hours isothermal oxidation testing conducted at 1010 oC. Faster oxide growth rates associated with transient oxidation slows down to steady state oxidation after about 100 hours. ...................................................................118 52 Specific mass gain results for 1000 hours isothermal oxidation testing plotted against the square root of time. ....................................................................................................119 53 Specific mass gain results for 1079 oC isotherma l oxidation testing. Despite uncharacteristic weight loss during isothermal oxidation, such behavior is not uncommon. Refer to the text for more information. ........................................................119 54 Three prominent layers co nstituting the oxide layers namely, granular NiO, convoluted and porous spinel oxide of the type Ni (Cr, Al)2O4 and several nuclei of Al2O3 which eventually grow to generate a flat protective scale. ................................120 55 Secondary electron micrographs of CM247LC alloy subjected to 200 hours of oxidation at 1010 oC showing the s pallation pattern of the NiO scales. The spallation pattern is reminiscent of a crystallographic fracture mode. .............................................120 56 Secondary electron micrographs showing the NiO grains after 50 hours of isothermal oxidation. Porosity in the NiO layer is also very apparent. With increase in the exposure time, the faceted NiO grains assume a rounded surface. ..................................121 57 Secondary electron micrographs showing the spinel oxides at the gas/scale interface. Diffusion and oxidation mechanism through the spinels is a complex phenomenon and not fully established yet ............................................................................................121 58 Secondary electron micrograph showing decohesion observed at the interface Al2O3 scale as indicated by the arrow. ................122 59 Secondary electron micrographs showing the presence of Al2O3 whiskers growing at the gas/scale interface on CM247LC alloy exposed to isothermal oxidation at 1079 oC for 50 hours. ..............................................................................................................122 510 Secondary electron micrograph sh Al2O3 scales observed at the gas/scale interface. Wrinkling was considered to be one of the mechanisms to relax the heavy compressive growth stresses built in the oxide scale. ...123 511 Scanning electron micrographs taken from CM247LC+300 ppm Pr alloy showing the changes in the gas/scale interface with increase in the oxidation exposure at 1010 oC. .....................................................................................................................................123 512 Scanning electron micrographs taken from the cross section of CM247LC oxidized at 1010 oC for 50 hours and 200 hours. ...........................................................................124

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15 513 Scanning electron micrographs taken from the cross section of CM247LC+860 ppm Ce oxidized at 1079 oC for 50 hours and 200 hours A thicker oxide scale with layers of outer NiO, middle spinel oxide and an inner Al2O3 was observed. ............................124 514 X ray diffraction patterns taken from CM247LC alloy subjected to oxidation at 1010 oC and 1079 oC for various times. ....................................................................................125 515 X ray diffraction pattern taken from CM247LC+860 ppm Ce alloy subjecte d to oxidation at 1010 oC and 1079 oC ...................................................................................126 516 Secondary electron micrographs taken from alloys CM247LC+860 ppm Ce alloy exposed to 1079 oC oxidation for 50 hours and CM247LC+300 ppm Pr allo y oxidized at 1010 oC for 500 hours The arrows indicate the presence of higher amount of Ce and Pr in the respective samples. ............................................................127 517 Secondary electron micrographs showing the developme nt of oxide layer during the initial few hours of transient oxidation. The needle Al2O3 Al2O3. ...........................................127 518 Auger scan results conducted on CM247LC+260 ppm Ce+80 ppm Pr alloy subjected to 1010 oC oxidation for 2, 5, 10 and 25 hours. The scan result on the unoxidized sample is also given as a reference. ...............................................................................128 519 Auger scan results from the spot which emitted Cathode Luminescence on the CM247LC+260 ppm Ce+80 ppm Pr alloy subjected to 1010 oC oxidation for 10 hours. ................................................................................................................................128 520 Results of cyclic oxid ation testing conducted at 1121 oC. Heavy spallation of CM247LC+860 ppm Ce and CM247LC alloys were observed with other RE additions improving the oxide adherence under thermal cycling conditions. ...............129 521 Results of cyclic oxidation testing conducted at 1150 oC. CM247LC+860 ppm Ce and CM247LC alloys showed uncontrolled spallation behavior while the oxidation resistance of other RE modified alloys were observed to be very good. .........................130 522 Results of cyclic oxidation testing conducted on downselected alloys at 1121 oC. ........131 523 Results of cyclic oxidation testing conducted on downselected alloys at 1150 oC. ........131 524 Secondary electron micrographs taken on an alloy exposed to cyclic oxidation at 1150 oC for 176 hours showing an externa Al2O3 scale, a gamma prime denuded .......................................................................132 525 Secondary electron micrographs showing the cracked oxide/alloy interface. CM247LC alloy showed a re latively smooth interface while CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd showed oxide imprints, which represent areas of high oxide alloy adherence ......................................................................................................132

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16 526 Mean Gamma Prime Denuded Zone (GDP Z) thickness for various alloys during cyclic oxidation testing at 1150 oC for 196 and 1176 hours. ...........................................133 527 Line scan results using Electron probe microanalysis showing the depletion of Al i n the GPDZ during cyclic oxidation at 1150 oC for 196, 1176 and 5400 hours. Higher level of depletion was noticed with increasing oxidation exposures. ..............................133 61 Cross sectio nal SEM micrographs showing the various layers of a TBC system namely, Air Plasma Spray deposited Yttria Stabilized Zirconia (YSZ) and High Velocity Oxy Fuel sprayed MCrAlY bond coat ...........................................................148 62 A digital image s howing the CM247LC substrate coated with TBC system during the various stages of oxidation exposure. The failure criterion is the complete de cohesion of the YSZ ceramic top coat. ............................................................................148 63 Spa llation lifetime of TBC system on various RE modified CM247LC alloys compared with the baseline alloy. Alloys exposed to oxidation at 1010 oC did not suffer spallation even after an extended span of 10, 000 hours, as indicated the arrow marks. .............................................................................................................................149 64 Spallation lifetime of TBC system on CM247LC alloys modified with more than one RE additions. CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd showed better lifetimes. ..........................................................................................................................149 65 Cross sectional SEM micrographs of the CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd coate d with a TBC system showing as deposited structure and microstructure after exposure at 1010 oC for 3720 hours. ..............................................150 66 SEM micrographs of CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd exposed to oxidation at 1010 oC for 3720 hours, showing two layered structure of bond coat. ......150 67 SEM micrographs showing the layer I/II interface of the bond coat in CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy exposed to oxidation for 3720 hours at 1010 oC and 1079 oC .. ...............................................................................151 68 SEM micrographs showing the YSZ/BC region in CM247LC+ 200 ppm Ce+60 ppm Pr+50 ppm Gd.. ........................................................................................................151 69 SEM micrographs showing the plan view of the thermally grown oxide layer (TGO) between YSZ and BC. ....................................................................................................152 610 SEM micrographs showing the cross sectional view of the thermally grown oxide (TGO) layer between YSZ and BC. It is clear that the primary composition of TGO i Al2O3 while the regions closer to the YSZ are richer in oxides of Ni, Co, Cr and Al. .....................................................................................................................................152

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17 611 SEM micrographs showing the presence of HfO2 internal oxidation stringers encapsulated inside Al2O3 as indicated by the arrows in CM247LC baseline alloy and CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy .. ................................153 612 X ray diffraction patterns of CM247LC+860 ppm Ce alloy exposed to oxidation at 1010 oC / 3720 hours, 1079 oC / 4392 hours and 1121 oC / 1440 hours. .......................153 613 Average thickness of the TGO layer in CM247LC baseline and CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloys. RE modified alloy showed a lower thickness value and the higher temperature lead to thickening of the TGO. ...................................154 614 Line scan results of micro probe analysis conducted on CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy exposed to oxidation at 1010 oC for 3720 hours. The hatched region represents the TGO layer composition. ...................................................155 615 Line scan results of micro probe analysis conducted on CM247LC+200 pm Ce+60 ppm Pr alloy exposed to oxidation at 1079 oC for 3720 hours. .......................................155 616 Secondary Ion Mass Spectroscopy (SIMS) results obtained from CM247LC+260 ppm Ce+80 ppm P r alloy exposed to 1010 oC for 3720 hours. RE elements were not detected at the interface. ..................................................................................................156 617 A collation of cross sectional TEM micrographs taken along the length of the TGO layer show ing finer grains at the YSZ/TGO interface followed by longer columnar grain s towards the TGO/BC interface. .............................................................................156 618 Cross sectional TEM micrographs showing the presence of an oxide phase rich in Hf Y Al in a TGO grain boundary intersecting the BC layer .........................................157 71 Tensile stress st rain curves of un HIPed Industrial solution heat treated Phase I alloys tested at 760 oC. T hese alloys w ere given a Industrial solution heat treatment followed by double aging. ................................................................................................171 72 Tensile stressstrain curves of un HIP ed UF solution heat treated Phase I alloys subjected to testing at 760 oC. Only the CM247LC+160 ppm Pr alloy was observed to exhibit a better plastic deformation behavior. .............................................................172 73 SEM fractographs taken from Phase I Industrial solution heat treated alloys .. ..............172 74 Tensile stress strain curves of HIPped Phase I UFHT alloys subjected to testing at room temperature. While CM247LC+860 ppm Ce alloy failed pre maturely in the elastic region, other RE modified alloys showed a moderate str ain hardening.. .............173 75 Tensile stress strain curves of HIPped UFHT Phase I alloys subjected to testing at 650 oC. While CM247LC+120 ppm Gd alloy exhibited higher tensile strength, CM247LC+180 ppm Dy alloy showed a better tensile ductility. ....................................174

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18 76 Tensile stress strain curves of HIPped UFHT Phase I alloys subjected to testing at 760 oC. CM247LC+ 180 ppm Dy alloy exhibited a minor necking behavior during testing. ..............................................................................................................................175 77 Ultimate strength values (UTS) of all the RE modified alloys tested at RT, 650 oC and 760 oC. Generally, a reductio n in UTS with increase in testing temperature was observed. ..........................................................................................................................175 78 SEM fractographs of HIPped UFHT Phase I alloys tested at RT showing an inter dendritic fracture m ode in CM247LC+860 ppm Ce alloy. .............................................176 79 SEM fractographs of HIPped UFHT Phase I alloys tested at 650 oC .. ............................176 710 SEM fractographs of HIPped UFHT Pha se I alloys tested at 760 oC ..............................177 711 Comparison of Ultimate strength values (UTS) between un HIPed and HIPed RE modified alloys. Except CM247LC+160 ppm Pr alloy, other RE modified alloys showed a considerable increase in the UTS due to HIPing. ............................................177 712 Yield strength values of Phase II Industrial solution heat treated alloys subjected to tensile testing at RT, 760 oC and 850 oC. A general trend of decrease in yield strength with increase in temperature was noticed. .........................................................178 713 Ultimate strength values (UTS) of Phase II Siemens solution heat treated alloys subjected to tensile testing at RT, 760 oC and 850 oC. With the exception of few anomalies, a decrease in the maximum tensile strength with increase in temperature was observed. ...................................................................................................................178 714 Tensile stress st rain curves of UFHT Phase II alloys subjected to testing at RT. The strain hardening response was very moderate with all the alloys suffering failure at the maximum tensile strength. .........................................................................................179 715 Tensile stress strain curves of UFT Phase II alloys subjected to testing at 650 oC. CM247LC+300 ppm alloy exhibited a very pronounced strain hardening response. .....179 716 Tensile stress strai n curves of UFHT Phase II alloys subjected to testing at 760 oC. CM247LC baseline alloy proved to be stronger than the RE modified alloys. ...............180 717 Tensile stress strain curves of UFHT Phase II alloys subjected to testing at 850 oC. CM247LC baseline alloy showed a pronounced necking behavior while RE modified alloys suffered pre mature failure in the elastic region. ..................................................180 718 Ultimate strength values (UTS) of UFHT Phase II alloys subjected to tensile testing at RT, 650 oC, 760 oC and 850 oC. With a few exceptions, a general trend of decreasing UTS values with increasing temperature was observed. ...............................181

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19 719 SEM fractographs taken on UFHT Phase II alloys tested at 650 oC showing trans granular cleavage facets in CM247LC baseline and micro void coalescence in CM247LC+300 ppm Pr that explains the higher ductility observed during testing. .......181 720 SEM fractographs taken on UFHT Phase II alloys tested at 760 oC showing microvoids in CM247LC baseline and inter granular mode of failure in CM247LC+300 ppm Pr. .............................................................................................................................182 721 SEM fractographs taken from UFHT Phase II alloys tested at 850 oC. ..........................182 81 Creep deformation behavior of Phase I UFHT Un HIPped alloys subjected to testing at 850 oC / 240 MPa. Addition of 860 ppm Ce and 180 ppm Dy was observed to increase the creep rates compared to CM247LC baseline alloy. .....................................190 82 Creep defor mation behavior of Phase I UFHT HIPped alloys subjected to testing at 850 oC / 390 MPa and 950 oC / 204 MPa. An increase in the ppm level of RE additions was observed to have an effect in increasing the steady state creep rates. ......191 83 Creep deformation behavior of Phase II IHT alloys subjected to testing at 850 oC / 390 MPa. Addition of RE elements resulted in premature rupture during testing. .........191 84 Creep deformation behavior of Phase II IHT alloys subjected to testing at 950 oC / 204 MPa. Similar to the previous tests, RE additions had a detrimental effect on creep rupture lifetimes. ....................................................................................................192 85 Creep deformation behavior of Phase II IHT CM247LC+300 ppm Pr alloy subjected to 1000 hour creep rupture testing at 850 oC / 280 MPa and 950 oC / 130 MPa. ............192 86 SEM micrographs showing the fracture surfaces of Phase II IHT alloys subjected to creep at 850 oC / 390 MPa.. .............................................................................................193 87 SEM micrographs showing the longitudinal sections of all oys subjected to creep testing ..............................................................................................................................193 88 Creep deformation behavior of Phase II UFHT alloys subjected to creep testing at 850 oC / 390 MPa. Addition of RE elements resulted in increased creep rates but prem ature creep rupture. ..................................................................................................194 89 Creep deformation behavior of Phase II UFHT alloys subjected to creep testing at 950 oC / 204 MPa. An increase in the creep rates was observed due to RE additions but resulted in premature creep failures. ..........................................................................194 810 SEM micrographs showing the fracture surfaces of Phase II UFHT alloys subjected to creep at 850 oC / 390 MPa.. .........................................................................................195 811 SEM micrographs showing the longitudinal sections of the alloys after creep failure. ..195

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20 812 Creep deformation curves of Phase II UFHT alloys su bjected to creep testing at the conditions indicted in the graph. A clear dependence of creep rupture on temperature was observed for CM247LC+300 ppm Pr alloy. .............................................................196 91 A schematic illustratio n of the several stages involved in the diffusion of RE elements from the substrate to the gas/oxide interface. T he oxidation in the CM247LC+RE alloys occurs predominantly by the inward oxygen diffusion.. .............223 92 TEM mic Al2O3 grain boundaries in near the gas/oxide interface in the CM247LC+860 ppm Ce alloy exposed to isothermal oxidation at 1079 oC for 1000 hours ................................................................................225 93 HR TEM Al2O3 grain boundary showing the distribution pattern of Hf and Ce. A reliable signal indicating the pres ence of Ce was not obtained. ............................................................................................................225 94 Specific weight gains of CM247LC and RE modified alloys exposed to 1010 oC and 1079 oC for 25 and 200 hours. .........................................................................................226 95 Comparison of the cyclic oxidation weight loss kinetics for CM247LC and RE modifie d alloys as a function of temperature. Note that the values of Kp are negative. ..226 96 Comparison of the cyclic oxidation weight loss kinetics as a function of ppm level RE additions (* indi cates that the specific alloy contains multiple RE additions). .........227 97 A schematic illustration showing the possible build up of RE rich oxides along the grain boundaries leading to increase in the lateral stresses which are detrimental to the scale adhesion. ..........................................................................................................227 98 TEM micrographs showing the cross Al2O3 scale containing HfO2 internal oxide particles Formation of voids within the oxides was also noticed as shown. ..............................................................................................................................228 99 TBC spallation lif etimes plotted as a function of amount of RE additions for isothermal oxidation at 1079 oC. In general, a total amount of 300 ppm RE was beneficial. .........................................................................................................................228 910 TBC spallation lifetimes plotted as a function of ppm levels for isothermal oxidation at 1121 oC. Similar to the previous case, additions of 300 ppm RE was found to be beneficial. .........................................................................................................................229 911 SEM micrograph showing the defects typically observed within the YSZ coating namely, inter splat cracks, porosity and undulations of the YSZ/TGO interface. ...........230 912 SEM micrograph taken on the TBC coated alloy exposed till fai lure at 1121 oC. The failure plane traverses along the crest in the undulations as indicated by the arrows. ....230

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21 913 Schematic illustrating the effect of RE segregation to the grain boundarie s. Reduction in the grain boundary diffusivity leads to cavitation along the boundaries and at the triple junctions. ................................................................................................231

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22 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 ON THE MICROSTRUCTURE, HIGH TEMPERATURE OXIDATION AND MECHANICAL BEHAVIOR OF A RAR E EARTH MODIFIED CM 247LC A NICKEL BASE POLYCRYSTALLINE SUPERALLOY By Krishna Prakash G anesan December 2009 Chair: Gerhard E. Fuchs Major: Materials Science and Engineering Minor additions of Rare Earth (RE) elements were observed to significantly improve the oxidation resistance of model Ni Cr Al alloys. This work was performed with the objective of establishing a mechanistic understanding of four RE elements namely Ce, Pr, Dy and Gd on the high temperature oxidation and mechanical behavior of Ni base superalloys. Oxidation behavior was studied by subjecting the alloys to isothermal and cyclic oxidation exposure in the temperature range of 1000 1200 oC for a span of 1000 10, 000 hours. The RE additions were beneficial in retarding the oxidation kinetics and increasing the scale adherence under thermal cycling conditions. It was proposed that segregation of REn+ Al2O3 grain boundaries inhibited the outward diffusion of oxygen active elements like Al, Ti etc. Due to this segregation, the oxidation reaction progressed predominantly by the inward diffusion of oxygen and lead to reduction in the scale growth rate. The segregation tendency was attributed to their larger ionic radii and higher thermodynamic affinity towards oxygen. Studies conducted on CM247LC alloy with a TBC (YSZ + MCrAlY) system also revealed the beneficial effects of the RE additions in imp roving the spallation lifetimes of YSZ by 20% in the temperature range of

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23 1050 1150 oMechanical behav ior of the RE modified CM247LC alloys were also studied through tensile and creep testing in the temperature range of 650 950 C, by reducing the growth rate of thermally grown oxide layers. The segregation behavior was characterized using Auger Spectroscopy, SIMS, TEM and SEM. oC. RE elements lead to increase in the steady state creep rate with reduced creep rupture lifetimes. A change in the fracture behavior occurred from mixed mode to completely inter granular at temperatures > 800 o C, due to RE additions. The temperature dependence of the creep deformation mechanisms in RE modified was evaluated. Due to their larger atomic size and low solubility in Ni, RE preferentially segregates to the grain boundaries. The significance of RE additions on the grain boundary sliding and segregation to stacking faults and dislocation cores were analyzed.

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24 CHAPTER 1 INTRODUCTION 1.1 Hydrogen Powered Gas Turbines Gas turbines for power and propulsion operate by the combustion of fuels derived from oil and natural gas. Dwindling resources of the conventional fuels has driven the development of alternative sources of energy to meet the ever increasing global energy dema nds. Large reserves of coal [1] (Shown in Figure 11) and relatively lower cost of energy generation has generated renewed interest in the coal based technologies [2] Present day coal combusti on methods based on Pulverized Coal Combustion (PCC) suffer from inherent disadvantages like lower energy efficiency and higher levels of harmful emissions such as CO2, SO2 and NOx. To counter these problems, Clean Coal Technologies (CCT) are being developed using gasification based systems combined with CO2 ) ( 2 2 2 y Electricit Turbine Steam Heat O H CO O Coal capture and sequestration methods. Compared to PCC, in gasification reaction, coal reacts with controlled amount of air or oxygen to produce synthetic (syn) gas. The two combustion reactions are compared below: (PCC) ) ( ) (2 2 2 2 22Steam Heat O H CO syngas H CO O CoalO (Gasification) The hydrogen derived from the syngas can be used to drive the gas turbines to generate electricity [3] Department of Energys FutureGen program proposes the development of a coal fired Integrated Gasification Combined Cycle (IGCC) which utilizes hydrogen as the primary fuel with an efficiency greater than 50 % and lower levels of emissions [4] The advantages of IGCC over PCC are evident from increased efficiency and reduced emissions as shown in Table 11.

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25 1.2 Approaches towards Alloy Development Higher temperatures close to 1400 oC in the hydrogen powered turbines combined with severe oxidation, hot corrosion conditions and ex treme mechanical loading demand the development of high performance alloys and heat resistant coatings. A roadmap for the IGCC development proposed by Siemens Power Generation [5] is illustrated in Figu re 1 2. Among the class of high temperature materials, Nibase superalloys offer the best combinati on of properties required in the H21.3 Understanding the Rare Earth Effect turbines. Polycrystalline CM247LC alloys are chosen due to their excellent castability and mechanical properties. A major modification in the alloy chemistry will be the addition of minor levels of Rare Earth (RE) elements from the lanthanide series. Four RE namely Cerium (Ce), Praseodymium (Pr), Dysprosium (Dy) and Gadolinium (Gd) will be added in ppm levels to CM247LC composition during casting. In the first phase of the project (Phase I), 15 alloys containing one and m ultiple RE elements will be conventionally cast as equiaxed structures. The composition of CM247LC and RE additions are given in Tables 31 and 32, respectively. Based on the performance during high temperature oxidation and mechanical testing, further mo difications in the alloy chemistry will be formulated. The question that may be prompted at this point is, Why RE elements? the answer to which is discussed below. Rare Earth or Reactive Elements are the 17 eleme nts in the lanthanide series of the periodic table and the definition generally includes Y, Zr and Hf also. The original Rare Earth effect was propounded by L.B. Pfeil [6] in his 1937 patent on the improvements in the properties of a heat resistant Ni Cr alloy. Since then, RE elements were used in metallurgy to eliminate S from the melt, improve the workability of alloys and more significantly, increase the oxidation resistance of Cr2O3 and Al2O3 forming alloys. Major developments were achieved in the 1970s and a breakthrough came in the form of an excellent review b y Whittle and Stringer [7] which

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26 promoted further progress in this field. Chapter 2: Background offers a detailed review of the various mechanisms proposed until this time on the beneficial effects of RE additions. Most of understanding related to the RE effect is confined to the research performed on model Ni Cr Al alloys. Though this composition served as the platform for the extensive development of Ni base superalloys chemistry, it is of interest to understand the role of minor additions in superalloys as we ll. Some of the commercial Nibase superalloys like Ren N5, CMSX 4 etc have reported improvements in the oxidation properties due to the minor additions of Y and La [8, 9] However, the role o f these minor additions on hightemperature mechanical strength remains largely inconclusive. It should be reminded that, during the earlier days of superalloys development, major emphasis was given to improving the alloy stre ngth [1012] Oxidation properti es, mostly, depended heavily on the overlay coatings. However, the coatings should not be considered as prime reliant since any local degradation of the coatings would accelerate the substrate failure. In this context, it is imperative to develop an alloy composition with optimum combination of mechanical strength and oxidation resistance. In this work, it would be of interest to apply the learning from the simple NiCr Al alloys on a commercial significant CM247LC being developed for next generation H2To obtain a mechanistic understanding of the RE effect on a commercially relevant CM247LC superalloys and Thermal Barrier Coatings thr ough hightemperature ox idation exposure and mechanical testing, microstructural characterization and optimization for better long term benefits. pow ered turbines. With these factors serving as a motivation, the objective of this research is to: 1.4 Organization of the Dissertation This section describes the general strategy for achieving the research objectives. With the sign ificance and objective of the research described in this chapter, Chapter 2 will provide a thorough review of the RE elements in context of their effect in oxidation and mechanical

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27 properties published in literature. The objectives will be achieved mainly through RE alloying, testing and characterization. To this end, Chapter 3 presents the experimental details covering the alloy compositions, oxidation and mechanical testing procedures, various microstructural characterization techniques namely Scanning El ectron Microscopy, Transmission Electron Microscopy, Focused Ion Beam milling, X ray Diffraction, Auger Electron Microscopy and Secondary Ion Mass Spectroscopy. The significance of using a specific characterization is also explained. Chapter 4 details the microstructural effects due to the addition of RE elements in CM247LC in as cast condition and heat treated condition. The details of a new RE Hf phase will also be discussed based on their structure and evolution during heat treatment. Chapter 5 will pres ent the oxidation behavior of the RE modified CM247LC alloys during the hightemperature isothermal and cyclic oxidation exposure. Results based on the reduction in oxidation kinetics and improvements in the oxide scale adhesion will be discussed. The lear ning from the RE effect on uncoated substrates will be used on CM247LC coated with Thermal Barrier Coating (TBC) system as shown in Chapter 6. The RE effect on the growth of Thermally Grown Oxide (TGO) layer and spallation lifetime of Yttria Stabilized Zir conia (YSZ) coatings will be discussed. After studying the role of minor additions of RE on the surface properties, the deformation behavior is characterized through high temperature tensile and creep testing, as will be elucidated in Chapters 7 and 8, re spectively. Chapter 9 would discuss the various results and strive to provide a mechanistic understanding on the role of the RE elements on the oxidation and the mechanical properties. In this chapter, more emphasis will be placed on what makes RE element s unique in a Ni base superalloys matrix, which lays the foundation for the RE Effect. The dissertation will conclude with Chapter 10 which outlines major conclusions, impact and future directions of this work.

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28 Figure 1 1. Chart representing the relat ive percentages of world coal reserves as of 2005. The numbers represent the quantity of recoverable coal in million short tons (Reproduced with permission from [2] ) Table 1 1. Comparison between Pulverized Coal Combustion (PCC) and Integrated Gasification Com bined Cycle (IGCC) in terms of maximum gas temperature, efficiency of power generation and emission levels.

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29 Figure 1 2. The roadmap for hydrogen powered gas turbines drafted by Siemens Power Generation. The work performed at University of Florida is hi ghlighted (Reproduced with permission from [5] )

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30 CHAPTER 2 BACKGROUND The objective of this chapter is to introduce the reader to the various effects of Rare Earth (RE) additions reported in the literature over the past several years. Due to the higher appreciation of RE elements towards improving the oxidation resistance, the emphasis of the chapter is more on oxidation compared to mechanical properties. At the outset, some of the key aspects of Ni base superalloys development and why it still remains the material of choice for high temperature applications are outlined. 2.1 Development of Ni base Superalloys Ni base superalloys are a special class of materials used primarily in the high temperature secti ons of the gas turbines for propulsion and power generation applications, generally constituting about 50% of the total engine weight [13] Their usage in such extreme high performance conditions is justified by the ir ability to bear heavy structural loads and withstand surface degradation with excellent microstructural stability even at temperatures greater than 0.6 Tm (absolute melting temperature) [1420] The development of Ni base superalloys has closely followed the improvements in jet engines since the early 1940s. Over all these years, the turbine entry temperature (TET) has been steadily rising to meet the continually growing demands for increased engine operating efficiencies. Figure 2 1 elucidates the steady increase in the TET from the first known Whittle Engine W1 to modern day Trent 900 and GEnx engines. To illustrate the significance of the material requirements, the first W1 engine delivered a thrust of approximately 1000 lb at TET of ~ 700 oC. The advanced GEnx engine developed by General Electric to power Boeing 787 Dreamliner is designed to deliver thrust levels closer t o 70,000 lbs with TET approaching 1500 oC. The increase in the high temperature capability of Ni base superalloys is primarily attributed to decades of research and development work on the various

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31 alloying strategies, improved processing conditions, advanc ed cooling channels and finally, the development of thermal barrier coatings. This section attempts to provide a short glimpse of the various significant achievements in this field which are responsible for high performance propulsion and power generation. 2.1.1 Improvements in Alloy Processing Use of Vacuum Induction Melting (VIM) as the primary melting process for superalloys has enabled a closer control of detrimental impurities and addition of highly reactive elements like Al, Ti etc which are vital for the strengthening mechanisms operating at higher temperatures [2125] Earlier methods involving wrought processing were hampered by the inability to add strengthening elements due to problems associated with hot workability. These shortcomings were overcome by the introduction of investment casting process with the capability of producing near net shape airfoils. Casting improvements also lead to the production of hollow shaped blades to reduce the centrifugal stresses associated with the rotating component weight. Grain boundaries were long considered as the weaker spots affecting the strength of Ni base superalloys leading to intergranular fracture at higher temperatures. A major breakthrough came in the form of Directional Solidification (DS) technique which eliminated the weaker grain boundaries transverse to the stress direction and produced columnar grains oriented along the low modulus crystallogr aphic direction, <001> [26, 27] Leaping one step forward, grain boundaries were completely eliminated through modifications in the DS process to make the alloy completely mono crystalline [28] As seen in the Figure 2 2, single crystals lead to massive improvements in high temperature creep resistance of superalloys and readily find application in the hottest turbine sections. On the other hand, polycrystalline alloys are still employed in the compressor, discs and cooler sections of the engines. Turbine discs are made of

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32 power metallurgy alloys, primarily manufactured through forging and extrusion techniques [2931] As evident from Figure 2 1, improvements in the engineering design of cooling schemes have also contributed majorly to the continuous increments in the TET. Modern day airfoils and vanes contain intricate cooling channels for the flow of cooling air which absorbs the heat, thus making the superalloys operate at lower temperatures. Similar to the composition of thermal insulation coatings, the cooling schemes are highly proprietary in nature. Modern day aircr aft turbine sections contain up to 100, 000 cooling holes [32] and the fabrication of the cooling chann els was made possible with the advances in the casting [33] and recently, laser machining techniques have enabled the drilling of cooling channels on blades after the deposition of overlay coatings [34, 35] 2.1.2 Alloying and the Strengthening Mechanisms The composition of Ni base superalloys originates from Ni 20% Cr alloy, a solid solution strengthened allo y. During the course of the development, several alloying additions primarily from the d block of the periodic table were added to impart strengthening through various mechanisms [36] Higher resistance to creep was achieved through solid solution strengthening effect brought about by the addition of heavier elements like W, Ta, Mo, Re and Ru. Major strengthening in the superalloys is attributed to the presence of a c 3(Al, Ti, Ta, Nb) is an ordered L12 structure which is a FCC lattice with Ni atoms in the face centers and other atoms at the face corners. D critical resolved shear stress (CRSS) with respect to increase in temperature was observed in these alloys [3739] H igh strength Ni base superalloys contain about 70 vol % of this strengthening phase [4042] Minor levels of C, B, Zr and Hf were a dded for grain boundary

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33 strengthening in polycrystalline superalloys. These elements were detrimental due to the formation of low melting eutectics and severely limited the solutioning characteristics of polycrystalline alloys. Due to the elimination of gr ain boundaries in single crystals, these elements were removed from the chemistry which increased the incipient melting temperature of the alloy leading to higher degree of microstructural homogeneity. Later, addition of these elements such as C to single crystals was found to be beneficial in enhancing the casting yields and prevent the formation of casting defects like freckles, stray grains etc [43, 44] One of the major problems associated with the microstructural stability of Nibase superalloys is the formation of detrimental Topologically C lose Packed (TCP) phases after extended exposure to higher temperatures. TCP phases are typically composed of Ni, Cr, Co, Mo, W and Re which are important for the strengthening the microstructure and this phase formation leads to the depletion of these vit al elements [45, 46] Needlelike morphology of these complex phases acts as stress raisers and contributes to crack initiation. To better predict the behavior of superalloys at higher temperatures and modify the compositions to prevent the formation of these detrimental phases, theoretical approaches using advanced methods like CALPHAD (CALculation of PHAse Diagrams) are employed in the present days for alloy design [ 4749] 2.1.3 Thermal Barrier Coatings Adding to the improvements in alloying and processing, deposition of overlay thermal insulation coatings raised the operating temperatures closer to the melting temperature of superalloys [5053] A TBC system consists of a MCrAlY (M = Ni, Co) bond coat to provide oxidation resistance and a top ceramic layer of Yttria Stabilized Zirconia (YSZ) which provides heat resistance. Originall y formulated by NASA, ZrO2 68 wt% Y2O3 was observed to have lower thermal conductivity at higher temperatures of operation. There are two major methods

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34 used for the deposition of YSZ coatings. Electron Beam Physical Vapor Deposition (EB PVD) coatings wh ich produce a columnar morphology, are used on the blades in the aero engines. Air Plasma Spray (APS) coatings which impart splat morphology are primarily employed in the longer dimension blades and vanes used in power generation turbines. The coating is beneficial in reducing the temperature at the alloy interface by about 100 200 K leading to better performance of the underlying substrate. Chapter 6: Oxidation Behavior of Alloys with TBC system gives a very detailed account of the coating properties from the perspective of the current research. 2.2 Mechanism of Oxide Scale Formation A slow growing, dense, thermodynamically stable and nonvolatile oxide layers are considered to be ideal for imparting maximum protection to the substrate against degradation from high temperature oxidation conditions. The following section will discuss the mechanism behind the establishment of a continuous protective layer in oxidation resistant Ni Cr Al type alloys. 2.2.1 Scale formation in Ni Cr Al alloys When a Ni base superalloy is exposed to high temperature oxidizing conditions, almost all of the constituent elements form their respective oxides at the alloy surface, the quantity of which is determined by their individual volume fraction. This phase of oxidation during the early stages is defined as transient stage of oxidation With increase in time, steady state o x idation is established that is characterized by the formation of stable, slow growing oxides namely Cr2O3 and Al2O3Generally, oxidation reactions progress at the surface by the outward diffusion of M n+ (M = substrate elements) and inward diffusion of O2 -. During the initial stages, a linear rate of oxide growth is observed due to the ready availability of O2 and Mn+. After a thin layer is established,

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35 furth er oxidation rate assumes a parabolic rate, indicating a diffusion controlled scale growth is prevalent. After the establishment of a continuous protective layer through the hetero diffusion mechanism, outward diffusion of Mn+ would be necessary for the ox ide layer replenishment. However, a critical amount of oxygen active solute is necessary to form a continuous oxide layer. Subcritical addition levels would result in the formation of internal oxide particles of the respective solute element. According t o earlier thermodynamic considerations, it was postulated that Cr2O3 and Al2O3 would form as the stable oxide layer even when added in ppma levels to Ni Cr and Ni Al alloys, respectively [54, 55] But, later experimental analysis reported that a critical volume fraction of Cr and Al were necessary to form and continually replenish the protective oxide layer, which depends on various factors such a s microstructure, oxidation temperature, oxygen solubility in the alloy etc. In a binary Ni Al alloy, a minimum of 17 wt% Al was found to be necessary to form a continuous protective Al2O3 scale [55] while 30 wt% Cr was required for the generation of Cr2O3 layer in NiCr system [54] As mentioned earlier, in either cases, subcritical composition of Al and Cr we re observed to nucleate faster growing, nonprotective NiO and Ni(Cr, Al)2O4 type of oxides as the external layers. Al2O3 and Cr2O3 Extensive studies were conducted on the oxidation mechanisms of model Ni Cr Al alloys, which were the base for modern Ni base superalloys [12, 17, 54, 56, 57] In Ni Cr Al alloys, oxidation mechanisms were grouped into three categories based on the amount of Cr and Al additions. With the subcritical levels of Cr and Al, Group I alloys formed a continuous NiO layer with internal oxides of Cr and Al. Addition of suf ficient Cr lead to the generation of an external Cr nucleated as internal oxides in respective cases. 2O3 layer with internal Al2O3 subscale, categorized as Group II. With higher Al levels, a stable protective Al2O3 formed with no visible internal oxidation, defined as Group III.

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36 Figure 2 3 illustrates a N i Cr Al ternary diagram indicating the various compositional ranges where these three types of oxidations are generally observed. The oxidation reaction occurring in Ni Cr Al alloys as a function of exposure time is depicted in Figure 24 and explained as following. Rapid intake of oxygen into alloy instantaneously converts the constituent elements into their respective oxides in the near surface layer. NiO and Ni(Cr, Al)2O4 spinel layers form on the surface, with a larger area of the surface covered with the spinel oxides. Faster diffusion through the NiO leads to the rapid thickening of the NiO layer and entirely covers the spinel oxide scale. At the interface between the NiO + Ni(Cr, Al)2O4 scales, Cr2O3 begins to nucleate and with sufficient concentrati on begins to a generate a thin sub scale. The formation of Cr2O3 subscale curtails the outward diffusion of other metallic ions, thus checking the further growth of NiO and spinel layers. Due to the ability of Al to oxidize even under lower oxygen partial pressure, Al2O3 subscale generates as a thin uniform layer closer to the substrate. Through the formation of an intermediate subscale of Cr2O3, Cr helps in the formation of a continuous layer of Al2O3, reducing the critical amount of Al needed for scale formation to about 5 wt% in Ni Cr Al alloys. Comparing the growth rates of NiO, Cr2O3 and Al2O3 layers, as shown in Figure 25, it is seen that Al2O3 is an ideal scale for optimal protection with reduced oxidation kinetics at higher temperatures [56] The close packed hexagonal structure and defect free nature of Al2O3 scale fu rther contributes in reducing the diffusion rate of the various species. The high temperature applicability of Cr2O3 ) ( 3 ) ( 2 ) ( 3 24 3 2g g sCrO O O Cr forming alloys is limited due to the following reaction that leads to the volatilization of oxide scale under higher oxygen partial pressur es.

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37 Al2O3 2.2.2 Polymorphic Forms of Alumina does not involve any gas phase reactions and is considered very stable at higher temperatures. Without belonging to any of the above three categories, when a mixed character of oxidation is noticed, it is termed as transition alloy oxidation. This type of oxidation is widely observed in Ni base superalloys and will be discussed in Chapter 4 with representative microstructures. Several metastable polymorphic variations of Al2O3 form on the Al containing alloys and the layer described in the previous section that generates during steady state oxidation is a Al2O3 layer, metastable polymorphic form s of Al2O3 nucleate and grow. As many as three polymorphs have been reported in literature [5860] and the transformation can generally be represented as: Among the metastable phas phase was widely observed. Below about 1000 o Al2O3 was found to be the dom inant phase for a longer period, Al2O3 begins to nucleate at higher temperatures, > 1000 oC, shortly after the exposure to oxidation [58, 61] These metastable phases were basically cubic structures with O2 in the lattice sites and Al3+ occupying the octahedral and tetrahedral vacancies. With the progression of the polymorphism from the ratio between octahedral to tetrahedral site occupancy decreases reaching a value of unity for Al2O3. Figure 2Al2O3, the metastab le phases contain many structural defects like cation vacancies, which encourages the outward diffusion of Al3+ For this reason, the metastable phase exhibits bladelike or whisker morphology, indicative of outward diffusion of

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38 cations through the lattic e. This outward diffusion mechanism was investigated using 18O tracer studies in Secondary Ion Mass Spectroscopy [62, 63] On the other hand, the stable Al2O3 grows predominantly by the inward diffusion of O2 and limited outward diffusion of Al3+It was interesting to note that due to the similar corundum structure, Cr 2O3 promotes the Al2O3 during the oxidation of Ni Cr Al alloys. At th e same time, addition of rare earth elements like Ce, La, Y etc was found to segregate to the grain boundaries, inhibiting the martensitic transformation of FCC HCP [60, 62] The arres t of the transformation reaction was useful for catalytic applications where the bladelike morphology of Al2O3 2.3 Mechanistic Understanding of Rare Earth Effect is beneficial. The transformation from is also associated with the volume reduction of about 13% leading to the cracking of the oxide s cale. This cracking was usually observed to lead to the formation of convoluted oxide morphology due to the enhanced outward diffusion of oxygen active elements. This outward diffusion was inhibited by the addition of RE elements and thus the surface scale was relatively free of ridges [60, 64] Through all these years, since the Rare Earth (RE) effect was first propounded in1935 by Pfeil [6] var ious RE elements, including Zr, Hf, Y were added to alloys to improve the oxidation resistance where it played a major role in reducing the scale growth rate and increasing the adhesion of the oxide scale to the underlying substrate. Figure 27 compares the cyclic oxidation characteristics of NiAl alloys at 1500 oC. The Whittle and Stringer review extensively discussed the role of RE dispersoids on the oxidation behavior of Cr2O3 forming alloys [7] Several authors also reported the beneficial effects of RE elements added as dispersoids on reducin g the rate of scale growth and improving the oxidation behavior [60, 6572] Addition of oxygen active RE elements as dispersoids were beneficial in serving as heterogeneous nucleation

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39 sites enabling the formation of a continuous protective scale in a short span of time and even at subcritical levels of Al or Cr. In addition, dispersoids increased the high temperature strength of the alloys by impeding the dislocation motion through the lattice. Other modes of additions were through the application of an external coating rich in RE elem ents [7378] or the ion implantation of the elements into the near surface regions [7984] These superficial methods were advantageous in the following ways higher amounts of several R E elements could be used regardless of their compatibility with the substrate. Also, it provides a direct method to understand the actual RE effect independent of the substrate chemistry. Later, it was postulated that a continuous flux of beneficial RE ele ments is important to realize long term benefits on the oxidation properties, which could be achieved only by their addition to the substrates [8587] The Whittle and St ringer review categorically described several mechanisms that were proposed in the literature prior to 1980 on the effect of RE elements. After their review, as Pint mentioned in his 2001 paper [87] advances in the highresolution characterization techniques has enabled a very detailed understanding of the RE effect. The following sections provide brief insights into various theories proposed to explain the significant improvements in the oxidation properties due to minor RE additions. Note that the order of the mechanisms is not based on any relative significance or chronological sequence. 2.3.1 I mproved Chemical Bonding at the Interface A linear dependence between the adhesive strength of a thin layer of Al2O3 scale formed on a metallic substrate on the o of formation of metallic oxide was reported by McDonald and Eberhart [88] Additions of ppm levels of oxygen active elements like Ti and Zr were proposed to improve the oxide adhesion by forming stronger M O bonds at the interface Considering the higher oxygen affinity, the same theory could be applied to RE elemental additions also. But, any experimental evidence in support of this chemical bonding theory is

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40 unavailable. Later, this proposed effect was observed to have only minor significance on the oxidation resistance. Additions of dispersoids of an element like Al with much lower stability of oxides compared to Zr, Ti, RE etc also was observed to improve remar kably th e adhesion of oxide scales [7, 87, 89] The dis persoid effect was primarily attributed to their role on the Al2O32.3.2 Graded Seal Mechanism scale. Spallation and cracking of the oxide scales during cooling are attributed primarily to the stresses imposed due to differential thermal expansion coefficient valu es between the oxide scale and the substrate. Pfeiffer [90] reported improvements in the scale adhesion in Fe Cr Al alloys due to the addition of Mischmetal (Ce La Nd Pr). The improvement was correlated to a possible formation of a graded seal rich in RE Oxides sandwiched between the oxide scale and the substrate. With a thermal expansion coefficient value intermediate between the oxide and substrate, the graded seal was believ ed to alleviate the stresses upon cooling, leading to resultant improvement in scale adhesion. Whittle and Stringer [7] reviewed Cr2O3 forming alloys that were reportedly forming an intermediate layer between host oxide and substrate, which is rich in complex perovskite oxides of the type RE Cr O or RE Al O. This layer served to act as a set of sieves blocking the inward diffusion of O2 and were associated with a reduction in oxidation kinetics. Similarly, isothermal oxidation resistance of Fe Cr Al and Ni C r Al alloys was observed to be increased due to the formation of a perovskite type YCrO3 and LaAlO32.3.3 Pegging Effect layers [91] Despite several of the cited reports, a reliable experimental evidence suggesting the formation of an intermediate layer was never reported [89] As early as 1950, additions of Zr, Si and Ca to a Ni Cr alloy were observed to form a net work of Cr2O3 stringers penetrating into the alloy which was reported to increase the scale

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41 adhesion [92] Felten [93] also observed an extensive i nternal oxide network formation due to the addition of Y to a Cr2O3Addition of rare earth elements like Hf, Y, Sc also lead to the formation of HfO forming alloy. Following these initial reports, considerable attention was given to this morphological feature, perceived as a unique effect of RE additions. Known by the names pegs or str ingers, they were believed to provide a keying in effect, anchoring the oxide scale to the substrate. 2, Y2O3 or Sc2O32.3.4 Scale Plasticity rich pegs in Fe Cr Al alloys [89, 9497] Whittle and Stringer [7] also offered extensive support to this theory, saying that uniformly spaced smaller size pegs could be helpful in improving the scale adhesion. Contradicting the proposed effect of pegs in increasing the scale adh esion, addition of Y in Ni Cr Al and Fe Cr Al alloys produced a large population of deep extending pegs, but still the oxide scales were non adherent [96] The reason was large number of pegs would develop a significant out of plane tensile stresses normal to the oxide/alloy interface, which typically promote interfacial cracking and spallation [9] With the above reports that a large number of internally protruding pegs are not a necessary condition f or scale adhesion, the effect of this morphological feature is not completely understood. Differences in the high temperature mechanical behavior of oxide scale and substrate play an important role in determining the adhesion behavior of the oxide scale. At higher temperatures, substrates such as Ni base superalloys are usually strong. A weaker oxide scale formed on superalloys may rumple due to creep to relieve the stresses and eventually spall. In this context, the greater ability o f the oxide scale to accommodate plastic strains is considered to be beneficial for longer lifetimes. Minor RE additives like Y, Ce etc were usually found segregated to the grain boundaries in Al2O3 and Cr2O3 scales. Due to the solute pinning effect, this segregation lead to the inhibition of grain growth resulting in finer grains. At high

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42 temperature, grain boundary sliding is a dominant creep deformation mechanism. Large volume fraction of grain boundaries would assist in improving the strain compatibilit y of the oxide scale through sliding process to relieve the inbuilt scale growth stresses. The mechanism is debatable since other published reports have observed an increase in flow stress due to RE segregation to grain boundaries in bulk alumina, thus inhibiting the sliding mechanism [89, 98100] 2.3.5 Change in Oxidation Kinetics The effectiveness of alloying additions depends on their ability to redu ce the oxidation kinetics that reduces the thickness of the oxide scale and the associated growth stresses. During the earlier days, Pt marker studies were used very frequently to analyze the oxidation mec hanism by spotting the location of the Pt markers at the gas/oxide interface or the oxide/substrate interface [89, 101] Whereas the former evidence indicated the inward growth of the oxide scale dominated by the inward oxygen diffusion, the latte r represented the outward growth due to metallic ions ( Mn+). Due to the arbitrary and unreliable nature of the marker studies, more effective 18O tracer analysis was employed in conjunction with Secondary Ion Mass Spectrometry (SIMS) to understand the nature of oxidation kinetics [62, 63, 102, 103] It was generally observed that the growth kinetics of Cr2O3 scales are controlled by the predominant outward diffusion of Cr3+ ions [89, 104] During the beginning of oxidation exposure, availability of oxygen near the alloy surface leads to the development of equiaxed Cr2O3 scale due to the inter diffusion be tween the two reacting species. Then, the dominant mechanism of Cr3+ diffusion takes over to result in completely outward growth of the scale with columnar grains of Cr2O3Unlike Cr growing above the equiaxed structure. 2O3, Al2O3 growth is sustained by the counter diffusion of Al3+ and O2 -. In response to the oxygen potential gradient, Al3+ and other oxygenactive elements from the substrates diffuse in the outward direction through the pre existing alumina scale. Grain

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43 boundaries of the oxide scale provide faster paths for the outward diffusion of such species, by orders of magnitude greater than the bulk diffusion. This outward diffusion of Mn+ would lead to formation of oxides within the original scale lead ing to lateral thickening o f the scale which may add to the stresses in the oxide scale [94] Also, e nhanced transport along the grain boundaries could result in the thickening of the scale near the grain boundaries leading to the scale rumpling at the gas/oxi de interface [68] RE elements were routinely observed to segregate to the oxide scal e grain boundaries. Al2O3 lead to the suppression of outward diffusion of oxygen active elements like Al, Ti etc, transforming the oxidation mechanism to predominantly by inward diffusion of O2 2.3.6 Vacancy Sink Mec hanism [9, 68, 69, 74, 77, 104] Due to this change in the oxidation mechanism, the oxide morphology changed from a convoluted to a thin uniform scale [68, 94] exhibiting columnar oxide structure as shown in Figure 2 8. Initial reports on the formation of a perovskite type oxide at the grain boundaries were later disproved by extensive TEM investigations that the segregation usually occur in the ionic form [68, 85, 105] Among the various reasons considered to be d etrimental to scale adhesion is the nucleation and growth of voids at the interface that limit the contact between the oxide and the substrate. It was discussed previously that oxidation reaction is controlled by the counter diffusion of cations and ions along the grain boundaries of the oxide scales. By the Kirkendall diffusion mechanism, large scale diffusion of a cation like Al3+ results in the counter current of vacancies as well. Microstructural observations have shown the presence of a large population of voids at the intersection of oxide scale grain boundaries with the oxide/alloy interface [106, 107] Segregation of rare earth elements to the oxide s cale grain boundaries inhibited the outward diffusion of the oxygen active elements and consequently, the internal void formation due to the

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44 vacancy coalescence was suppressed [78, 106, 107] Addition of dispersoids containing rare earth elements were also observed to inhibit the vacancy coalescence at the interfaces [7, 87, 89, 101] Extensive void formation was observed around the dispersoid interfaces within the alloy near the alloy/oxide interface [101] It was proposed that rare earth additions with large atomic radii form a vacancy com plex in an attempt to reduce the dilational strain energy [89] One of the observations reported that in spite of the improved oxi dation behavior of the Y containing Ni Cr Al alloys, pores at the alloy oxide interface were comparable in size and number to the undoped alloy [108] The larger dimension of the voids observed with the oxide scale (Refer Fig 2 8B) was explained by Pint [107] that the interfacial voids nucleated due to the scale separation at the oxide/alloy interface were incorporated into the growing scale during the inward growth of the oxide scales. Segregation of S to the free surfaces was found to be detrimental to the interfacial void nucleation and growth, as discussed in the following section. 2.4 Effect of Sulfur on the Oxidation Behavior The problems associated with S are more pronounced in industrial gas tur bines due to cheaper fuels that serve as the source of S. On the other hand, aircraft engines run on much cleaner fuels, making the problems related to S induced hot corrosion negligi ble. The other sources of S are the indigenous impurities present in the base alloy itself. Due to its small size, S even when present in ppm levels, readily segregates to free surfaces such as grain boundaries and interfaces to reduce the surface free en ergy [108110] During oxidation, the indigenous S present in the alloy was found to segregate to the oxide scale interface, weakening the interface and ultimately leading to the scale spallation. This detrimental effect of S was widely reported during the oxidation of Ni base superalloys [8, 9, 111113] Ni Cr Al alloys [96, 104, 108, 114, 115] Fe Cr Al alloys [116] and in steels [117] The subject was recently reviewed in detail in light of the various characterization techniques available to analyze the segregation behavior of S

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45 at the alloy scale interfaces [64] Theoretical calculations also have indicated that segregation of S lead to weakening of bonds across the Ni Al2O3 interfaces [118, 119] Presence of even 10 ppm levels of S lead to faster rate of scale growth and premature spallation of the protective oxide scales during cyclic oxidation [9, 115] Using Auger spectroscopic studies, it was found that the concentration of S to the alloy Al2O3 interfaces increases with time. Such segregation leading to detrimental effect on the oxidation resistance was reported for Cr2O3 forming alloys also [120] Though no reliable explanation was given until this date on the actual mechanism behind the S effe ct, it was proposed [115] that S enters the oxide lattice as singly charged Sion that leads to the formation of Al3+ vacancies. The set up of the vacancy gradient promotes the outward diffusion of Al3+ i ons along the short circuit diffusion paths available in the oxides. Thus, Al3+ becomes freely available along the oxide scale grain boundaries and scale gas interface to participate actively in the oxidation process. This increase in the oxidation rate is coupled with buildup of heavier growth stresses. In the wake of the active outward diffusion of Al3+The detrimental effect of S on the oxide scale adhesion was reduced by additions of rare earth elements like Y, Ce and La that formed stable sulphides, thus inhibiting the availability of free S to the alloy scale interfaces [9, 74, 96, 108, 110, 114, 115] Figu re 2 9 illustrates the beneficial effects of RE additions in combating the detrimental influence of S during oxidation. At the same time, addition of rare earth element as sulphides like Y large population of vacancies counter diffuse, only to coalesce at the oxide/alloy interface resulting in scale separation. 2S3 lead to the poor oxidation results due to the possible breakup o f the sulphides and by making free S available for segregation at the interfaces [96, 108] In addition to using rare earth elements to eliminate S, various processing improvements were devised to reduce the S content of the alloy. For example,

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46 CMSX 4 SLS (Super Low Sulfur) was made commercially available with about 1 ppm S produced through improvements in the Vacuum Induction Melting technique [8] Hydrogen annealing was one important desulphurization technique used to reduce the S content to negligible levels. This process was observed to deliver oxidation properties equivalent to rare earth doped alloys [112] The subject of hydrogen annealing as a technique to clean the Ni base superalloys was widely reported in literature with associated improvement in the oxidation resistance [113, 121124] It was observed that RE doped alloys always performed better than the undoped and even S free base alloys indicating that effect of RE was beyond bei ng just a gettering agent [125] 2.5 Sulfur Embrittlement Embrittlement is associated with a broad ductility minimum in the temperature range of 0.4 0.5 Tm, associated with a transition from low temperature transgranular to high temperature intergranular mode of fracture. Intergranular fracture usually occurs by the nucleation and growth of grain boundary cavities that coalesce to form cracks. S was observed to be one of the major impurities leading to embrittlement effect, even if present in the minor ppm levels. In one of the earliest observations, presence of 50 ppm S proved to be detrimental to the workability of Ni [126] at around 650 oC, possibly due to the formation of a liquid eutectic phase Ni3S2 at the grain boundaries that promoted inter granular fracture. Later reports indicated that S led to loss of ductility over a broad range of temperature, 420 760 oC [127, 128] showing that its effect on the embrittlement behavior was more complicated than the formation of a liquid phase at the grain boundaries. Figure 210 depicts the reduction in ductility as a function of increasing S levels and temperatures. The detrimental effect of S on the mechanical properties o f Ni base alloys was reviewed in detail by several authors [24, 25, 129132] Various theories were postulated to address the embrittlement phenomenon due to S additions.

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47 Segregation of S to the grain boundaries were reported to reduce the boundary diffusivity leading to higher stress concentration and cavity nucleation during creep [133] By reducing the interfacial energy at the cavity surface, S segregation promoted the formation of more small sized cavities The presence and coalescence of these cavities lead to nucleation and propagati on of cracks, ending in intergranular failure. Cavities richer in S were observed in a Fe base alloy which failed inter gra nularly, as shown in Figure 211. Development of advanced melting techniques like ElectroSlag Remelting (ESR), Electron Beam Remel ting (EBR) and Vacuum Arc Remelting (VAR) enabled the production of cleaner alloys with closer control over the impurity levels [23] In addition, Vacuum Induction Melting (VIM) techniques were also helpful in the addition of elements with higher vapor pressures at casting temperatures. In this context, surface active RE elements like Ce, La etc which form high ly stable sul phides were included to counter the detrimental effects of S [25, 134, 135] 2.6 Effect of RE on the Mechanical Properties In addition to the gettering ability, RE elements were also intentionally added to alloys to influence the mechanical behavior. This section reviews the various publications that have reported the significance of RE additions on the mechanical behavior of Ni base systems. Addition of 60 180 ppm of Ce inhibited the change in the fracture mechanism from low temperature transgranular to high temperature intergranular fracture, otherwise observed in Ni Cr alloys w ithout RE additions [136] Similar results were also obtained during the tensile testing of a Udimet901 alloy containing Ce in the range of 90 300 ppm [137] These improvements were associated with major changes in the tensile ductility with Ce additions in creasing the reduction in the cross sectional area and elongation. However, RE additions did not produce any significant changes in maximum tensile strength values in both cases. Addition of Ce in ppm levels were found to decrease the minimum creep rate in Ni Cr alloys at 700 oC [138] as shown

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48 in Fig 2 12. These additions also played a role in improving the creep rupture ductility compared to the alloys without Ce additions [137] Observations of fracture surfaces also indicated a change in the fracture mechanism to transgranular mode due to the presence of RE elements. General ly, among the most important factors that determine an increase in creep resistance of the alloys are an increase in activation energy for self diffusion and a decrease in the stacking fault energy [139] Venkataraman [140] observed an increase in the activation energy for lattice diffusion from 295 kJ/mol to 355 kJ/mol with the addition of 63 ppm Ce in Ni Cr alloys. It was also reported that 225 ppm Y increased the activation energy from 264 kJ/mol to 310 kJ/mol for Ni Cr alloys [141] Detailed TEM investigations conducted to calculate the stacking fault energy by measuring the distance between the partial dislocations showed large variations in the energy values and results were largely inconclusive [138] Stress rupture tests conducted at 1100 oC and 80 MPa on Ni Al alloys containing Y in the range of 0.04 0.20 wt % showed improvement in the properties over the baseline alloy [142, 143] Y additions were reported to improve the stress rupture behavior in superalloys also [144, 145] A refinement in the inter dendritic spacing leading to better fatigue properties was reported in Ni Al intermetallics due to Y additions [142, 143] A Fe base superalloy of Incoloy 600 series cont aining trace additions of Y was also observed to exhibit longer stress rupture life [144, 146] The rupture life at the testing conditions, 540 oC and 825 MPa of the conventional alloy was found to be 50 hours, whereas Y containing alloys show ed a stress rupture life of 200 hours. The improvement in the rupture lif etimes was attributed to changes in the micr ostructure due to Y additions. These series of superalloys have been found to contain platelet precipitates known as phase and represented stoichiometrically as (Fe, Ni, Co)3(Nb, Ti, Si). With the addition of Y, a transformation in the structure of phase from FCC to an hexagonal phase called, H phase

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49 represented as (Fe, Ni, Co)5In addition to their effect on creep and tensile properties, RE elements like Dy Nd etc. were shown to increase compressive ductility values [149, 150] Segregation of these elements to grain boundaries was believed to make the boundaries stronger by increasing electron density, disordering the Ni Al bonds across the boundary resulti ng in an overall increase of compressive ductility. (Nb, Ti, Si) was observed [144, 146148] Compositional analysis illustrated segregation of Y to the H phase The orientation relationship between H matrix indicated that the closed packed planes and directions are parallel to each other. High resolution TEM investigations showed that these phases share a semi coherent interface that made the phase fairly stable due to the lower interfaci al energy. Significant improvement in the workabi lity of Nibase superalloys was also observed due to the presence of trace levels of Y [151, 152] Alloys with > 20 ppm wt% Y containing alloys exhibi ted the highest RA of about 80% during high temperature, high strain Gleeble tests. Optimal additions of Ce and La to the melt were also reported to induce favorable hot ductility to the al loys [134] While addressing the beneficial effects, higher levels of additions lead to the formation of brittle intermetallics phases like Ni5Hf, Ni5Ce etc leading to reductio n in the ductility values and generally the entire mechanical properties [25, 130, 142, 144, 145]

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50 Figure 2 1. Improvements in the turbine inlet temperatures from the first known Whittle W1 engine to the modern day Trent 900 engines. Higher temperature capability is attributed to advancements in the alloy design, processing improvements, intricate cooling channels and thermal barrier coatings (Reproduced with permission from [153] ). Figure 2 2. Creep behavior of Mar M200 alloy in the conventionally cast, directionally solidi fied and single crystal conditions Superior lifetimes and higher ductility observed for single crystals ( Reproduced with permission from [27] ).

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51 Figure 2 3. Ni Cr Al ternary phase diagram illustrated the regions where Group I, II and III oxidation are dominant ( Reproduced with permission from [17] ). Figure 2 4. Oxide scale development during Groups I, II and III oxidation. In the schematic, S represents spinel oxides of type Ni(Cr, Al)2O4 dark circles are Cr and empty circles Al ( Reproduced with permission from [57] ).

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52 Figure 2 5. Arrhenius plot showing the slower growth kinetics of Al2O3 compared to NiO and Cr2O3 scale at all the temperatures of oxidation. O Al A B Figure 2 6. Crystal structures of Al2O3 pol ymorphs (A) phase and (B) phase

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53 Figure 2 NiAl alloy at 1150 o C showing the beneficial effects of Zr and Y additions on the spallation resistance compared to the baseline alloy. Note that Zr and Y as an alloying addition and dispersoid within the alloy improved the properties better that implanted Y (Reproduced with permission from [9] ). 4 m A B 4 m A B Figure 2 8. SEM micrographs showing the development of Al2O3 NiAl alloys at 1200 oC oxidation, (A) unalloyed showing voids and whiskers (B) Zr containing alloy showing a flat uniform scale with a columnar morphology (Reproduced with permission from [9] ).

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54 Figure 2 NiAl at 1100 o C showing the detrimental effect of S on the spallation properties and beneficial effect of Y in gettering S (Reproduced with permission from [9] ). Figure 2 10. Detrimental effects of ppm levels of S in a Ni showing the ductility minimum in the temperature range of 600 1000 oC (Reproduced with permission from [132 ] ).

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55 A B Figure 2 11. Cavitation observed on the grain boundaries in Fe+60 ppm S tensile tested at 700 o C (A) and the corresponding X ray map taken in SEM showing heavy concentration of S in the cavities (B) (Reproduced with permission from [154] ). Figure 2 12. Reduction in the minimum creep rate in Ni Cr alloys due to the addition of ppm levels of Ce (Reproduced with permission from [138] ).

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56 CHAPTER 3 EXPERIMENTAL TECHNIQ UES The primary objectives of the research were to evaluate the role of minor rare earth alloying additions on the microstructure, oxidation and mechanical deformation behavior. The oxidation properties were studied by conducting isothermal and cyclic oxidation testing at high temperatures. Similarly, the mechanical deformation behavior was analyzed using tensile and creep testing at elevated temperatures. This chapter deals with the testing procedures and the characterization techniques employed to gain a detailed understanding of the Ra re Earth effect. 3.1 Materials CM247LC (Canon Muskegon 247Low Carbon) is a Ni base superalloy developed from Mar M247 by Canon Muskegon Inc, in the later 1970s primarily for use with equiaxed and directionally solidified structures [155] By introducing Hf and closely controlling the levels of Ti, Zr, Si and S, CM247LC alloys achieved remarkable resistance to hot tearing [156] thus imparting the alloys with excellent castability and mechanical properties. Blades used in the hot section of the industrial gas turbines are typically large and achieving higher casting yields of these components during the directional solidification processing has proven to be difficult. However, there is a need for these types of Ni base superalloys for hydrogen powered industrial gas turbines. Keeping this reason in mind, in this proj ect, CM247LC alloy chemistry was adopted to cast an equiaxed structure. The alloys used in this research were cast at Howmet Corporation, Inc., Whitehall, MI. Cylindrical bars approximately 15 cm long with 1.5 cm diameter were conventionally cast using vacuum induction melting using Howmet supplied master alloy. A typical baseline composition of the as cast CM247LC bars is given in Table 3 1. Compositional details of the major constituents were obtained using X Ray Fluorescence (XRF), whereas the minor addi tives were analyzed using Inductively Coupled Plasma (ICP) analysis.

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57 Howmet carried out nondestructive X ray technique to analyze the casting defects on the surface of the as cast bars. 3.1.1 Rare Earth Additions In the literature, Y, La and Ce were the most popular rare earth elements used in the oxidation studies. Though all these elements in the lanthanide series along with Y were recognized as rare earth elements, some researchers prefer to use the term Reactive Elements. In the current work, the author prefers to use the terminology Rare Earth Elements to address these special additions. For clarity, an abbreviated form RE will be used to address these additions. Four elements belonging to the Lanthanide series namely Cerium (Ce), Praseodymium ( Pr), Dysprosium (Dy) and Gadolinium (Gd) were chosen for this study. Though the primary objective was to establish the effect of single additions on the oxidation and mechanical properties, it was also of interest to identify the effect of multiple additio ns, more commonly referred to as co doping effect. So, alloys with two and three additions were also cast. Using a design of experiments factorial method, a total of 15 different compositions were selected. Including the baseline alloy for properties compa rison, a total of 16 different compositions was cast. Higher level additions of RE introduce low melting eutectics leading to incipient melting and associated reduction in the mechanical properties. Due to this constraint, the addition levels were maintain ed in the parts per million (ppm) levels. These additions were added in the melt during the final stages of casting in the form of powders wrapped in aluminum foils. The composition of the RE additives are given in Table 32. Due to the higher surface acti vities of these elements at the casting temperatures, retention levels were expected to be in the range of 4050% or in some cases less [8, 157] It was also found that all the detrimental impurities like S, P etc were maintained at less than 2 ppm levels in the as cast bars. In all of these modifications,

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58 the total additions were in the range of 100900 ppm. In the literature, the term dopants is used in some cases for such low level additions. In this work, though, the terms dopants and additives will be used interchangeably. Examination of the as cast microstructure to analyze the inter dendritic spacing, par titioning coefficient etc would be discussed in detail with the methods and results in the following chapter. 3.1.2 Heat Treatments Like most Nibase superalloys, CM247LC has been reported to exhibit severe solidification segregation and solution heat tre atment is required to avoid nonuniform compositions leading to local differences in the oxidation kinetics [158] or mechanical properties. Upon the re ceipt, the bars were subjected to a general precipitation hardening heat treatment comprising three stages, namely solution heat treatment at a higher temperature, quenching and aging at intermediate temperatures. Two types of solution heat treatments were used namely, Industrial Heat Treatment (IHT) and a baseline UF Solution Heat Treatment (UFHT) as given below: 1185 oC/10 minutes 1232 o1221 C/2h + rapid fan quench in Ar (IHT) oC/2h 1232 oC/2h 1244 oC/2h 1260 oThese solution heat treatments were followed by the double aging treatment: C/2h + rapid fan quench in Ar (UFHT) 1079 oC/4h air quench + 871 oTo eliminate the possibility of oxidation at such higher solutioning temperatures, both the IHT and UFHT were performed in a vacuum f urnace maintained at a pressure of 10 C/20h air quench 3 Torr. The furnace was preprogrammed for the desired ramping rates and soaking times at the appropriate temperatures. The ramp rates of 25 oC/minute from the room temperature to the first stage, 5 oC/minute between the consecutive stages and a 100 oC/minute quench rate from the maximum soak temperature during the cooling stage was used. It has to be mentioned that a slower heat up

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59 rate during the successive steps was used to provide higher level of solutioning and minimize the possibility of incipient melting. The above heat treatments were used only for the first part of the study. Based on the solidus values, several modi fied heat treatments were developed to maximize the microstructural homogenization of the alloy. The development of new modified heat treatments and comparison of the microstructural and compositional properties will be explained in detail in Chapter 4: Mi crostructure. 3.1.3 Hot Isostatic Pressing (HIPing) It is typical for the as cast bars to contain a small volume fraction of pores due to solidification shrinkage that can be detrimental to the mechanical properties. In this work, the examination of the as cast bars along longitudinal section showed the presence of about 2 vol % porosity. Hot Isostatic Pressing (HIPing) treatment was used to reduce the volume fraction of the porosities. Prior to HIPing, the as cast bars were given the following partial sol ution heat treatment in vacuum to decrease the chances of incipient melting: 1185 oC/10 minutes 1232 oThese partially solutioned bars were HIPed at PCC Airfoils, Minerva, OH. The bars were HIPed at 118515 C/30 minutes + rapid fan quench in Ar oC at a pressure of 175 7 MPa for 4 hours in an Argon atmosphere. A maximum heatup rate of 10 o3.2 Oxidation C/minute was used during HIPing. The HIPed samples were then given the remainder of the heat treatment and then machined into test specimens and their mechanical properti es were compared with the Un HIPed alloys. Alloy bars subjected to an IHT solution heat treatment followed by double aging treatment were electric d ischarge machined at Advanced Manufacturing Techniques (AMT), Clifton Park,

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60 NY to obtain cylin drical oxidation coupons ( 15mm diameter and 3m m thickness). Machining was preceded by heat treatment since the latter could possibly introduce a thin oxide film on the surface. Sharper edges in the oxidation coupons could lead to pronounced stress raiser e ffects leading to decohesion of the oxide layer at the edges [9, 94, 159] To avoid such geometrical variations having an effect on the oxidation kinetics, the edges were rounded prior to oxidation exposure. T he machined samples were ground using SiC grit papers and fine polished using ypical fully prepared sample is shown in Figure 31. Prior to oxidation exposure, the samples were degreased in acetone and cleaned ultrasonically using methanol. The sample preparation was similar for both isothermal and cyclic oxidation testing. The oxi dation testing was carried out in a hot zone of a furnace. Oxidation coupons were placed inside an alumina crucible and introduced into the furnace. The oxidation reaction was allowed to progress in static air at 1 atm pressure. At appropriate time interva ls, the samples were removed, cooled to room temperature and weighed in a microbalance. The coupons were re introduced into the furnace and the same procedure was followed for the entire span of exposure. The temperature and testing durations will be discussed in the respective sections. Due to the difference in the thermal expansion coefficient between the oxide scale and the metallic substrate, cooling stresses would lead to cracking and spallation of the oxide scales in some cases. Due to this reason, m ost oxidation testing cited in the literature use Thermo Gravimetric Analysis (TGA) set up to continuously monitor the weight gain/loss behavior at the test temperatures. Due to the difficulties associated with operating a TGA for a large number of samples and for long term exposure, a semi gravimetric analysis involving weight measurement after cooling to room temperature was employed. Most of the long term testing discussed in this work was conducted at Siemens Laboratories at Casselberry, FL.

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61 3.2.1 Isoth ermal Oxidation Testing The objectives of the isothermal oxidation testing were to study the oxidation kinetics during both the transient stage and steady state, oxide morphology at the gas scale interface, microstructure and composition of the oxide scale s evolving during various stages of exposure. The coupons were oxidized at two temperatures namely, 1010 oC and 1079 oOnly the alloys with single RE addition were subjected to isothermal oxidation and the baseline alloy was included for comparative purposes. This makes the total number of compositions exposed for the study to be 6. For each type of alloy, a total of 6 oxidation coupons were introduced into the preheated furnaces at the start of exposure at one temperature. Again referring to Figure 3 2, three samples were taken ou t after 50, 200 and 500 hours for microstructural analysis while the remaining three samples were used for weight gain measurements. These samples were measured after 25, 50, 100, 200, 300, 500, 700 and 1000 hours of oxidation. Due to the more transient na ture of the oxidation reaction during the initial few hours of exposure, it was necessary to conduct weight measurement at close intervals. C for a total span of 1000 hours. A schematic illustration of the isothermal oxidation testing procedure is shown in Figure 3 2. The figur e specifies the time intervals after which the coupons were removed for weight measurements. In addition to the above tests, short time oxidation exposures were carried out for 2, 5, 10 and 25 hours to study the oxide phase formation during the transient stage of oxidation. These shorter exposures were conducted using box furnaces at University of Florida. Oxidation coupons were placed in an alumina crucible and maintained at 1010 oC for the desired time perio ds. These samples were used for Auger Electron Spectroscopy (AES) studies also to study the segregation of RE elements to the gas scale interface during the initial few hours of oxidation.

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62 3.2.2 Cyclic Oxidation Testing While isothermal conditions would i nduce growth stresses into the oxide scale, cycling the coupons between high and room temperatures would introduce an additional factor of thermal stresses. This type of testing is more reflective of the conditions existing in the gas turbines during servi ce. Thus, the objec tives of this testing method were to understand the response of the alloys to thermal cycling conditions and effect of RE on the oxide scale spallation behavior Testing was carried out at three temperatures namely, 1079 oC, 1121 oC and 1150 oC with a cycle time of 24 hours. One cycle comprised of heating the samples from room temperature to the testing temperature at the rate of about 25 oC/minute, soaking at the test temperature and then cooling down to the room temperature at the rate of about 25 oCyclic oxidation tests were conducted in two batches. In batch I, all the 16 alloys were tested, wher eas batch II was confined to only six alloys downselected from the original group of the alloys based on the performance from the batch I tests. Performance of the alloys was judged on the basis of spallation properties, weight gain/loss and compatibility with an overl ay thermal barrier coating that was determined from a separate testing process. Oxidation performance of alloys with thermal barrier coatings will be discussed in Chapter 6: Oxidation Behavior of Coated Alloys. Batch I testing was conducted fo r 5400 hours, whereas in Batch II the exposure was extended to 9600 hours. This testing durations contributed to gain an understanding of the long term stability of the alloys with different RE elemental additions. C/minute which made the entire cycle time to be 24 hours. The cycle was then repeated again for the entire testing duration. Figure 3 3 schematically illustrates the thermal cycling. In this testing also, the samples were removed after pre determined time intervals for the weight measurement, which would be apparent from the weight gain/loss graphs in Chapter 5: Oxidation on Uncoated Alloys.

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63 3.3 Mechanical Testing High temperature mechanical testing namely, tensile and creep were conducted to evaluate the effect of ppm level RE additions on the deformation behavior of CM247LC. Including IHT and UFHT, modified solution heat treated alloys were subjected to mechanical testing and would be discussed below. Threaded mechanical testing specimens with a cylindrical gage section of diameter ~ 4.52 mm and gage length of ~ 26.035 mm were machined from fully heat treated alloy bars at Joliet Metallurgical Labs, Joliet, IL. The dimensions of bot h the tensile and creep specimens were similar, except that creep specimens had pin holes for extensometers, machined into the shoulders as shown in Figure 34. 3.3.1 Tensile Testing Tensile testing was conducted using a Satec servo hydraulic mechanical t esting frame with Instron 8800 instrumentation. The testing was conducted at an initial strain rate of 8 x 105 s1, which is equivalent to a machine cross head speed of 0.0144 inches/minute. The specimens were tensile tested at room temperature, 650 oC, 760 oC and 850 oC in air. The testing equipment is equipped with a clamshell type furnace with a maximum temperature capability of 1000 oC. Temperatures were monitored using two K type thermocouples wired onto the gage section and were maintained within 3 oC. A high precision extensometer was attached on to the specimen directly at room temperature or an external assembly at high temperatures to record the changes in the specimen dimension during tensile testing. The instrument was interfaced to a computer running Instron Merlin Software that monitored various tensile testing parameters continuously during testing. After the tensile testing, all the required information was retrieved from the software in the form of a spreadsheet and the tensile curves were plotted. The final gage length and gage diameter were measured using a Vernier Caliper to calculate the tensile ductility. Other

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64 important outputs including yield strength and Ultimate Tensile Strength (UTS) were obtained directly from the software. 3.3.2 Creep Testing Constant load creep tests were conducted in a Satec Model M3 frames with a Nuvision Mentor control system and a lever ratio of 16:1. The machine was equipped with a clamshell furnace with a temperature capability of more than 1000 oC. Three K type thermocouples were attached to the gage section of the specimen to record the temperature and maintain it within 3 oThe creep specimens were tested using four conditions namely, 750 C. Linear Variable Differential Transformer (LVDT) based extensometers were used to record the extension of the samples during cree p and the data was continuously fed into a computer interfaced with the testing equipment. The software was programmed to store the desired outputs such as time to 0.1%, 0.2%, 0.3%, 0.4%, 0.5%, 1%, 2% and 5% creep strain. Hot modulus values were also provi ded by the software directly. Similar to the Merlin program used in tens ile testing, the output data were obtained in the form of a spreadsheet useful for further analysis and presentation. Creep ductility values were measured from the sample dimensions us ing a Vernier caliper. oC / 650 MPa, 800 oC / MPa, 850 oC / 390 MPa and 950 o C / 204 MPa. These conditions were derived from LarsonMiller plots for Mar M247 with a projected lifetime of 30 0 hours (Figure 35). With the input of test temperature and desired life time, P value was obtained from the following equation: 310 20 log 460 t T P (Eqn. 3 1) where P is the Larson Miller parameter, T is temperature in oF and t is time in hours. With the calculated P value, the corresponding load level from the ordinate axis was found.

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65 3.4 Microstructural Characterization After studying the RE related properties during oxidation and mechanical test ing, it was important to relate the observed properties to the microstructure. This section discussed the analytical characterization tools used for studying the microstructure. 3.4.1 Metallographic Preparation Surface topography of the oxide scales at th e scalegas interface of oxidation samples was studied by depositing a thin layer of carbon to avoid any charging effect under the electron beam due to the nonconductivity of the oxide scales. To analyze the scale metal interface, the specimens were secti oned during a diamond blade. These samples were mounted in a thermosetting epoxy resin and ground flat using SiC grit papers of grades 240, 320, 400, 600, 800 and 1200. The samples were then polished to a mirror finish using alumina suspensions of sizes 5, 3, 1 and 0.3 m. To analyze the microstructure of the as cast, heat treated and mechanical tested specimens, the samples were sectioned using a cut off wheel. The mounting and polishing procedures were similar to those discussed above. Samples for fract ography were sectioned using a diamond blade and mounted on an aluminum stub. To study the grain sizes, the specimens were etched using Kallings Reagent (50 ml HCl, 50 ml C2H5OH, and 2.5 g CuCl2 2O, 100 ml HCl, 100 ml HNO3 and 3 g of MoO3) and AG21 etchant (25 ml Lactic Acid, 15 ml HNO3 and 1 ml HCl). Any special sample requirements for specific characterization equipment would be discussed in the corresponding section.

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66 3.4.2 Optical Microscopy As cast microstructural analysis to study the inter dendritic spacing and grain sizes were conducted using a Leica light microscope with an inverted mounted platform. The microscope was equipped with a digital camera and only low magnification images like 50X a nd 100 X were captured using this microscope with the help of imaging software. 3.4.3 Scanning Electron Microscopy A JEOL 6400 Scanning Microscope equipped with Energy Dispersive Spectroscopy (EDS) was used for studying the microstructure of oxidation and mechanical tested specimens. A working distance of 15 mm and an accelerating voltage of 15 kV were used for secondary electron imaging and compositional analysis of the microstructure while fractography was conducted at a higher working distance of 25 mm for improved depth of focus. The features of interest during this analysis were; As occurrence of incipient melting and RE Hf phases were studied. Oxidation: Gas scale and scalealloy interfaces were studied to characterize various oxide phases and their composition. Additionally, Gamma Prime Denuded Zones (GPDZ) and Thermally Grown Oxide (TGO) layer thickness were also measured from the cross sectional images. Mechan ical Testing: Fracture surfaces were studied to understand the initiation and mode of fracture. Longitudinal sections were analyzed for the presence of secondary cracks, rafting and cracking of RE Hf phases. 3.4.4 Electron Probe Micro Analysis (EPMA) Oxida tion has been recognized as a very dynamic process with the inter diffusion of several ionic species across the gas, scale and alloy interfaces. It becomes imperative to analyze the compositional variations along the scale alloy cross section. Since the composition studies

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67 using SEM EDS are semi quantitative, JEOL Super Probe EPMA with a probe size of 1 m with a step size of 2 or 5 m was used for the line scan analysis. Samples prepared in a similar way to SEM were used in the un etched condition. EPMA wa s also employed extensively to compare several heat treatment modifications for elemental homogeneity. 3.4.5 X Ray Diffraction Due to the complexity of the alloy chemistry under investigation, it was expected that various oxide types and complex spinels w ould nucleate and grow during the high temperature exposure. A Philips APD Powder Diffractometer was used for the oxide phase analysis under the working conditions of 40 kV and 20 mA. As oxidized samples were mounted on glass slides, clamped inside the dif fractometer set o to 70o3.4.6 Auger Electron Spectroscopy (AES) PCPDF database. Valuable insight into the se gregation and precipitation of oxygen active RE elements on the gas scale interface could be provided by Auger Electron Spectroscopy. AES PHI 660 Scanning Auger Multiprobe was used with an accelerating voltage of 10 kV and a total of 10 sweeps on a sample. A RBD analysis suite software was used to plot the Auger scan results in the form of dN (E) vs Kinetic energy (eV) plots. Conductive samples were required for the analysis. Since the depth of Auger analysis is confined to the Angstrom levels from the surf aces, any carbon coating to make the oxide scale conductive would potentially impact the outcome of the analysis. So, only the samples that were exposed to 2, 5, 10 and 25 hours of oxidation which were expected to develop only a very thin oxide film were s ubjected to this study. When charging effects due to nonconductivity became an issue during analysis, Ar gas was used to sputter away an ultra fine layer of the gas scale interface.

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68 3.4.7 Secondary Ion Mass Spectrometry (SIMS) To analyze the segregation of RE elements to the Thermally Grown Oxide (TGO) layers grown between the ceramic top coating and the metallic bond coat, a Perkin Elmer/Physical Electronics model 6600 Quadrupole based Dynamic SIMS system was used. The ceramic coating was polished off to leave a very fine layer of TGO on the metallic bond coat. 56 keV Oxygen or c esium primary ion beams were used to sputter the sample surface. The ions formed during the sputtering process were analyzed using a mass spectrometer. 3.4.8 Transmission Electron Microscopy (TEM) Segregation analysis in the oxide scale grain boundaries and the alloy scale inter face were investigated using TEM. A JEOL TEM200CX was used at an accelerating voltage of 200 kV in the bright field and dark field conditions for studying the oxide grains, grain boundaries and interfaces. Crystal structure of the oxide layers was determined by studying the diffraction pattern taken during TEM observation. Segregation of RE to the thermally grown oxide layers were also studied in a similar way. RE Hf phases were also subjected to extensive TEM analysis structure near the grain boundaries in the creep tested samples were also studied in TEM. When necessary, compositional data was gathered using a high resolution TEM2010 equipped with an Energy Dispersive Analysis using X rays (EDAX). TEM thin foils were prepared using Focused Ion Beam (FIB) technique s As oxidized samples were given a carbon coating and mounted in FIB specific stubs. A beam of Ga ions was used for the milling process. A layer of Pt was deposited before FIB milling to protect the sample from beam damage. An AUTOFIB program was used to mill the trenches into the specimen and mill it to the sample thickness of about 200 nm. Following that, manual thinning using a lower ion beam current was performed to thin the samples down to the electron transparent levels (~

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69 100 nm). The various stages in the milling process are shown in Figur e 36. For oxide samples, transverse sections were prepared by milling on the plan section of gas scale interface, whereas polished cross sections were used for obtaining TEM foils from RE Hf and mechanical tested samples. Thinned down samples were cut fre e using the ion beam and lifted ex situ using a micro manipulator provided with glass needles with micron sized tips. Lifted out foils were placed on a C grid for TEM analysis.

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70 Table 3 1. As cast composition of baseline CM247LC alloy (Wt %). Table 3 2. Composition of 16 alloys in Phase I based on the RE element additions. Figure 3 1. A typical fully prepared alloy coupon for oxidation testing. The edges have been rounded to reduce the stress effects during high temperature oxide growth.

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71 1 2 6 3 5 4SamplesTo be taken out after 200 hours To be taken out after 500 hours Samples to be weighed at specific intervals for weight gain data which will later be averaged out. 6 5 4 6 5 4Weigh 50 h 25 h 1Take out 100 h 6 5 4 6 5 4Weigh 6 5 4 6 5 4Weigh 2Take out 200 h 300 h 6 5 4 6 5 4Weigh 6 5 4 6 5 4Weigh 3Take out 500 h 700 h 6 5 4 6 5 4Weigh 6 5 4 6 5 4Weigh 1000 h Figure 32. The isothermal oxidation testing procedure showing the samples and the time intervals for removals for the weight measurements. Similar procedures were used for oxidation testing at 1010 oC and 1079 o C. Figure 3 3. A schematic illustration of a cyc lic oxidation profile. A heating and cooling rate of 25 oC/min were used during the testing and the time per cycle including the heating, soaking and cooling is 24 hours.

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72 Figure 3 4. An engineering drawing of a mechanical testing specimen showing the im portant dimensions. Pin holes were machined into the shoulders to use the screw type extensometers. Figure 3 5. LarsonMiller parameter plotted against the stress for Mar M247 alloy. The properties of equiaxed CM247LC are very similar to its parent allo y Mar M247.

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73 Figure 3 6. Images taken in the secondary electron mode showing the various sequential steps in the preparation of TEM foils using FIB. The steps shown are deposition of a thin Pt strip for protection against beam damage followed by milling o f trenches and final thinning.

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74 CHAPTER 4 MICROSTRUCTURE 4.1 Introduction The excellent high temperature mechanical properties of Ni base superalloys are attributed ffect of minor RE additions on the high temperature oxidation and mechanical properties of a polycrystalline Ni base superalloy. While trying to establish the RE effect on these major properties, it is important to understand the influence of these minor a dditions on the microstructure. As cast microstructures are highly segregated and direct exposure of the alloy to high temperature would lead to significant local differences in the oxidation and deformation behavior. So, a high temperature heat treatment is required to homogenize the elemental precipitates and attain better high temperature properties [160, 161] In this research, for the first time, a RE Hf phase was observed during solution heat treatment. With only a few reports concerning the microstructura l effects of RE elements in Ni base superalloys [144146, 148] this finding assumes greater importance. The objectives of the work described in the chapter are: To establish the role of RE elements on the as cast structural parameters such as inter dendritic spacing and partitioning coefficient. To design a multi ste p solution heat treatment cycle for obtaining the homogeneity of chemical composition and microstructure. To conduct detailed microanalytical studies on the RE Hf phase to obtain the information on the structure, composition, growth kinetics and orientati on relationship with the matrix. 4.2 As Cast Microstructural Analysis To understand the effect of RE additions on the as cast structures, preliminary analysis was conducted using a light optical microscope. Low magnification microstructures obtained from t sections of CM247LC and CM247LC+300 ppm Pr alloys

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75 are illustrated in Figure 4 1. The microstructures were very typical of a conventionally cast and slow cooled dendritic structure with long dendritic arms radiating along the cross section. The darker contrast of the inter dendritic regions between the secondary dendrite arms represented the heavier nature of segregation during solidification. Low magnification images did not show any appreciable differences between the b aseline and RE modified alloys. The finer details of the dendritic and inter dendritic regions were studied using SEM. Typical features of the microstructure closer to the inter dendritic regions are presented in Figure 4 2. The dendritic regions showed the presence of prim whereas the inter dendritic g solidification, the eutectics, which were the last to solidify, were very coarse in nature. Carbides phase were MC carbides type rich in Hf, Ti and Ta and the morphologies were of blocky and script type. Despite the extensive analysis of the inter dendritic regions in various alloys, there were no significant differences in the overall m icrostructure due to the additions of RE elements. 4.2.1 Inter dendritic Spacing Measurements RE elements have a larger atomic radii compared to the base metal Ni and can be expected to segregate preferentially to the inter dendritic regions during the so lidification process. This segregation to the inter dendritic regions may potentially limit the growth of dendrites and leading to their refinement. A very fine dendritic structure is favorable for better fatigue properties [142] To evaluate the effect of various RE on the dendritic structure, the inter dendritic spacings were measured using the low magnification images taken using an optical microscope. The linear measurements spanning several sec ondary dendrites parallel to the primary dendrite was performed using SPOT image analysis software. For the purpose of consistency, all the measurements were taken over 11 secondary dendritic arms (n) and the total length obtained was divided by (n1) (i.e 10) which gave the required Secondary Dendritic Arm

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76 Spacing (SDAS) [162] The results obtained from CM247LC baseline and CM247LC alloys containing onl y one RE additions are represented in Figure 43. As inferred from the figure, CM247LC baseline alloy showed an average SDAS value of about 50 m. Due to the addition of RE elements, a refinement of about 34 m was obtained. Measurements conducted on allo ys with multiple RE elements also confirmed the limited dendritic refinement due to the RE additions. 4.2.2 Partitioning Studies Using MicroProbe Analysis Using a probe size of 1 m, an Electron Probe Micro Analysis (EPMA) was employed in the line scan mo de with a step size of 5 m to analyze the distribution of RE elemental additions over a distance of 150 m. The line scanning was performed approximately in the center of the cross section in the as cast alloy bars, only for CM247LC alloys containing one RE addition. The results of the distribution of Ce, Pr, Dy and Gd in different alloys are complied in Figure 44. Irrespective of the higher level of Ce present in the alloy, Pr exhibited the highest levels during the analysis. Also, the evaluated weight percent values were about 17.5 times higher than the original composition for Pr and about 24 times higher for the other three RE elements. This observation implied a higher degree of compositional inhomogeneity for the RE elements. However, due to the und efined location of the line used for scanning, a thorough understanding of the distribution pattern for the various regions was not obtained. To further understand the distribution behavior of these RE elements, micro probe point scans were conducted on tw o different locations on the alloy cross sections over a length of 10 secondary dendritic and the adjoining inter dendritic regions. In addition to RE elements, major constituent elements like Ni, Cr, Co, W, Ta and Al were also included in the analysis. The results were used to calculate the partitioning coefficient of the various elements, k which is defined by the ratio of the composition of the element in the dendritic region divided by the composition of the same

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77 element in the inter dendritic region [163] This analysis was conducted on two alloys namely, CM247LC+860 ppm Ce and CM247LC+300 ppm Pr and the k values are shown in Figure 45. The segregation behavior of refractory elements like W and Co to the dendritic regions were repre sented by their k values higher than unity. With the Cr showing a k value of approximately unity, other elements exhibited their preference to segregate to inter dendritic regions with their value being less than unity. In line with our previous observat ion of the role of RE additions on the dendritic refinement, Ce and Pr exhibited kvalues much, less than unity, indicating their stronger tendency to segregate in the inter dendritic regions. 4.3 Development of a Multi Step Solution Heat Treatment As seen in the last section, some of the characteristic features of an as cast structure in superalloys are: non in the inter dendritic regions and nonuniform dendritic structure. All these irregularities could be represented by the inhomogeneities associated with the elemental distribution in the as cast alloys as shown in the Figure 46. The extent of segregation was illustrated by the results of line scan performed over a length of 150 m for the heavier elements like W, Ta, Co and Cr. On other hand, Ni, be ing the base, does not usually show larger differences in the distribution as shown by these elements. The nonuniform microstructure resulting directly from the solidification conditions and the elemental distribution leads to detrimental mechanical prope rties. All the superalloys are subjected to a solution heat treatment at a higher temperature segregation from the baseline solution heat treatment of CM247L C and the development of modified multistep solution heat treatment cycles to obtain a more uniform microstructure. The effectiveness of a specific heat treatment modification was determined by means of SEM

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78 characterization for microstructural uniformity and micro probe analysis for compositional homogeneity. 4.3.1 Baseline Solution Heat Treatments During the initial stages, two different types of baseline solution heat treatment were conducted on the alloys as mentioned below: Industrial Heat Treatment ( IHT) which was a single step heat treatment with a maximum soaking temperature of 1232 oC. The alloy bars were heated at the rate of 20 oC / min with a hold for 5 minutes at 1185 oC. From this point, the alloys were heated at a slow rate of 2 oC / min and soaked at 1232 oC for 2 hours. Finally, the alloys were rapidly cooled at the rate of 200 o Multistep Heat Treatment (MHT) which consisted of four soaking stages at 1121 C / min using a gas forced cool. oC, 1232 oC, 1245 oC and 1260 oC for 2 hours each. T he rate of heating from the room temperature to the first step was 25 oC / min and from one temperature to the next was 5 o C / min. The cooling rate was similar to the previous case. Refer to Table 4.1 for the detailed outline of the various baseline and m odified heat treatments used in this research. Typically, in the industry, the elaborate multi step heat treatments are not frequently used for the larger blades for power generation applications. They follow a single step heat treatment like the one ment ioned above followed by double aging at 1079 oC for 4 hours and 871 oAs shown in Figure 4 7, the microstructural observations conducted on alloys given the IHT solution heat treatment showed the presence of a significant fraction of residua dendritic regions, the structure was not very different from the as cast alloy. The non uniformity in the compositional distribution of heavier elements lik e W, Ta, Co and Cr was also evident from the micro probe results shown in Figure 48. Although, the micro probe results were significantly different from that of the as cast structure, still the desirable level of homogeneity C for 20 hours. Industry resort s to such simple heat treatme nt to achieve cost reduction, increased productivity and avoid difficulties in achieving uniformity in the large dimension blad es.

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79 was not achieved. In the cas e of MHT, the local differences in the composition were further Figure 4 9 A. Eutectic regions, being the last to solidify, had the lowest melting point in the mic rostructure and failing to exercise care during heat treatment design would lead to local melting. Under mechanical load, these regions would act as the initiation site of failure. ad to the local variation in the stress fields during mechanical deformation. As shown in the Figure 49 B, such differences result in the nonuniform rates of deformation and cracking at the interfaces of eutectics and matrix. The issues of residual eutectics and incipient melting necessitated the development of a modified solution heat treatment, as discussed below. 4.3.2 Differential Thermal Analysis (DTA) Results Despite the widely recognized advantages of RE additions on the high temperature propertie s, two factors impose limitations on inclusion of RE elements in Ni base superalloys. Firstly, they have lower solubility in the Ni matrix and additions beyond some limit would lead to the formation of detrimental brittle inter metallic Nix( RE) phases. Sec ondly, RE elements form low melting eutectics with Ni that lead to the incipient melting during high temperature exposure. The eutectic temperatures taken from the respective Ni RE phase diagrams are tabulated in Table 4 2 [164] The results shows that any local RE enrichment could lead to the formation of a very low melting eutectic. In this research, the RE elements have been added to CM247LC in the range of 100 900 ppm. Thus, to design an optimum solution heat treatment, knowledge on the ef fect of RE additions on the solidus temperatures of the alloys was considered to be important. To this end, Differential Thermal Analysis was conducted on the alloy bars as explained below.

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80 CM247LC baseline alloy and 15 alloy bars with different RE additions were sectioned using a cut off wheel to obtain cylindrical samples, 15 mm diameter and 15 mm thick. These samples were sent to Dirats Laboratories in Westfield, MA for DTA to evaluate the effect of RE samples were heated at a constant r ate of 20 oC / min with a reference sample of pure Ni and the difference in the temperature between the two samples were noted. The temperature difference and the derivative of temperature difference between the samples were plotted as a function of temper ature. Figure 4 10 shows a typical DTA curve obtained in CM247LC+860 ppm Ce alloy showing the inflection point which corresponds to the various transformation temperatures, as indicated. The details of the solvus, solidus and liquidus temperatures obtained from this analysis are given in Table 4 3. At the levels examined in this study, the addition of various ppm level RE additions did not significantly lower the solidus temperatures of the alloys in the as cast condition with an average reduction of only 10 o temperature was about 1210 o4.3.3 Design of Modified Solution Heat Treatments C. The objective of this part of the work was to design a successful heat treatment with a higher degree of compositional homogeneity of the various elements and microstructural time of heat trea tment. Micro probe analysis on the as cast alloys revealed the nature of segregation. So, there should have been local differences in the solvus and solidus temperatures values in the dendritic and inter dendritic regions depending on the composition. Thus the DTA results did not necessarily reflect the true solvus and solidus values. Inter dendritic regions, being the last to solidify, would naturally have the lowest solidus temperatures. Since the local

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81 composition determines the solidus, it may be more appropriate to term solidus as local incipient melting temperature. locally melts inci pient melting temperature define the region within which the heat treatment has to be carried out. In an alloy, a near homogeneous composition would take the incipient melting temperature close to the original solidus temperature of the alloy. A considerable amount of time has to be provided to enable the diffusion of elements in the alloy to achieve homogenization. This can be achieved by introducing several steps with successively higher temperatures with a specific soaking time at the temperatures. Preliminary studies showed 1260 oC to be the appropriate highest holding temperature. Despite a lengthier heat treatment and multiple steps to ensure homogenization, heating at temperatures higher than 1260 oC, say 1265 oC, proved to be counter productive with the appearance of severe incipient melting. Thus, fixing 1260 oC as the highest temperature and 1210 o step heat treatments were designed within this heat treatment window of 50 oBased on the above expl anation, four modified heat treatments ranging from an elaborate 5step 11 hours heat treatment to 3step 5 hours heat treatment were conducted. The list of the modified heat treatments is given in Table 4 1. The heat treatments were conducted in a vacuum furnace. Generally, in all the heat treatments, the alloy samples were heated from ambient temperature at the rate of 20 C. For reference, the heat treatment window has been schematically illustrated using a NiAl phase diagram in Figure 4 11. oC / min up to a temperature slig temperature, 1190 ooC / min up to the first soaking stage. The same lower degree of heating was followed between the

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82 soaking steps also. As mentioned previously, soaking at successively higher temperatures and heating slowly into and through the solvus region are important to allow sufficient time for the heavier elements to diffuse and avoid incipient melting. All the heat treatments proved to be very beneficial i n hom ogenizing the microstructure, with negligible incipient melting and only a very 12. Micro probe results obtained from the alloy subjected to ModSHT03 heat treatment is shown in Figure 4 13. The comparison of the compositional homogeneity with the previous IHT clearly proved the effectiveness of the modified solution heat treatment. This was attributed to the present design homogenization. Taking into consi deration the compositional homogeneity, microstructural uniformity and time required for the heat treatment, Mod SHT03 was found to be the effective solution heat treatment. In the Phase II and Phase III mechanical testing in this work, all the alloys woul d be subjected to only ModSHT03 and would be compared against IHT. Following the solution heat treatment, the alloys were subjected to a primary aging / coating cycle heat treatment at 1080 oC for 4 hours followed by the secondary aging at 870 o4.4 Rare Earth Hafnium Phase C for 20 hours. The aging treatments were conducted in a box furnace and the microstructure of a 14. equiaxed alloys as a grain boundary strengthener. In the directionally solidified alloys, addition of Hf was found to enhance the hot tearing resistance by preventing cracking along the grain boundaries [156] Hf is also known as an effective carbide former and primary MC type carbides containing Hf were observed in the as cast alloys in this study. The hi gher oxygen affinity of Hf was also found beneficial in increasing the high temperature oxidation resistance [87, 165] In this research,

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83 Al2O3As presented before, the as cast microstructure was composed primarily of the usual sectional analysis of the as cast alloys using SEM did not show the presence of any RE rich phases within the grains or the grain boundaries where RE elements tend to segregate. While characterizing the alloys to determine the microstructural homogenization after modif ied solution heat treatments, phases rich in RE and Hf were detected. These phases were of lighter contrast compared to the carbides and the morphology of the phases were observed to be either blocky or a long plate like structure, situated mostly along th e grain boundaries. SEM/EDS analysis showed the presence of all the RE elements added to the specific alloy and Hf in the phases. The compositional analysis performed using EDS showed the presence of the respective RE element/elements added to the specific alloy and Hf. Figure 415 shows an Hf Ce Gd phase formed during a multi step solution heat treatment cycle. A Hf carbide particle is also shown in the same micrograph to provide a comparison for the phase contrast. A semi quantitative analysis using EDS g ave the composition of the phase to be 54.18 % Hf, 35.58 % Ce and 10.24 % Gd (given in wt %). To obtain the relative distribution of the different elements within the phase and phase/matrix interfaces, X ray mapping was also performed and the results illus trated in Figure 4 16. A complete absence of major elements like Ni, Al, Cr etc was evident from the maps. The reason why these phases were not observed during the baseline solution heat treatment is still uncertain. One of the advantages offered by the mo dified solution heat treatments was the almost co mplete absence of residual which allowed for a better analysis of the microstructural features. After all the modified solution heat treatment cycles, this phase was noticed. A quantitative analysis to gather the in both the uncoated and alloys with TBC were observed. A detailed account of this Hf effect in the current research is given in Chapter 9.

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84 volume fraction of the phases was not attempted due to the limited amount of the phase observed. strongly to the grain boundaries to reduce the misfit strain energy. Hf is also a stronger se gregant to the grain boundaries, but the interaction phenomenon between these two elements was not documented before. A SEM micrograph showing the phase along the grain boundaries is given in Figure 4 17. The common m orphologies of the phase were blocky, about 5 m in diameter and rod like structure about 10 m long. To obtain a three dimensional perspective of the blocky structure, a trench was milled into the alloy cross section using Ga ion beam in Focused Ion Beam (FIB) instrument as shown in Figure 418. It is clear that the phase formation was taking place during the solution heat treatment stage. To understand the phase evolution and to identify at what stage of solutioning, the phases are nucleating, various mul tistep annealing treatments were designed to be conducted in vacuum as listed in Table 4 2. Due to the higher affinity of Hf and RE towards oxygen, performing heat treatments in box furnace atmospheres would lead to the contamination of the phase due to oxidation damage. Microstructural homogenization is important to reduce the volume fraction of times were required during the annealing treatments. As the ma ximum temperature was lowered to study the phase evolution at various stages of the actual solution heat treatments, successively greater amounts of time were required. SEM analysis of the RE Hf01 and RE Hf02 treatments showed the presence of RE Hf phases, eutectics could not be reduced in the remaining annealing treatments and after careful observation, the phases were not present. Next step, the modified solution heat treatments were

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85 conducted and after every solutioning stage, the samples were rapidly quenched in He to obtain the structural information pertaining to the respective stages. To gain a detailed understanding of the RE Hf phase, TEM analysis was conducted on thin foils prepared using F IB technique. A typical TEM micrograph taken on a Ce Dy Hf phase is shown in Figure 419. Interestingly the sample characterized contained a bi crystal of the phase as noticed from the grain contrast obtained by tilting the foil with respect to the beam direction. To complement the compositional details obtained in SEM using EDS analysis, a 3 nm size probe EDAX facility in the High Resolution TEM (HR TEM) instrument was employed to obtain an accurate composition of the phase. Figure 420 shows phase analyze d and the corresponding EDAX spectrum that gave the composition of the phase to be 58.23 % Hf, 36.34 % Ce and 5.43 % Dy (given in % wt). Lattice imaging to obtain the dspacing of the structure did not give any conclusive results due to the thickness varia tions in the foil. 4.5 Summary The results of the study conducted on as cast and solution heat treated alloys can be summarized as follows: Larger atomic radii and lower solubility in the Nimatrix resulted in the pronounced segregation of RE elements to the inter dendritic regions. This segregation lead to dendritic refinement by about 34 m. with negligible incipient melting. During the solution heat treatment, a RE Hf phase was observed to form as blocky or platelet morphologies along the grain boundaries. A detailed TEM investigation was conducted to analyze the phase composition and structure.

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86 0.25 mm A B Figure 4 1. Optical micrographs showing the dendritic structure in the as cast alloys, (A) CM247LC baseline and (B) CM247LC + 300 ppm Pr. The longer dendritic structure radiating across the cross section are typical of a slowly cooled structure. The darker regions in the inter dendritic re gions represent the higher level of segregation.

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87 100 m 5 m A B Figure 4 2. SEM micrographs showing the structure of the as cast CM247LC+860 ppm Ce, (A) dendritic structure and (B) higher magnification view of the box showing inter dendritic structure showing the pre eutectics. The carbides are of MC type rich in Hf, Ti and Ta. 45 4647 48 49 50 51 CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr 45 4647 48 49 50 51 CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr Figure 4 3. Comparison of the average Secondary Dendritic Arm Spacing (SDAS) CM247LC alloys with various RE additions. A refinement of 34 m was observed due to RE additions.

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88 0.00 0.05 0.10 0.15 0.20 0.25 0.30 0.35 0 20 40 60 80 100 120 140 160 % Wt 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd Figure 4 4. Microprobe analysis results of RE distribution in the as cast CM247LC alloys. Note that the detected wt % values are higher than the actual alloy composition. 0 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 1.8 Ni Cr Co Al Ta W Ce/Pr Partitioning Coefficient, k' CM247LC+860 ppm Ce CM247LC+300 ppm Pr Dendritic Inter Dendritic Figure 4 5. Dendritic and inter dendritic partitioning behavior of various elements in the as cast CM247LC+860 ppm Ce and CM247LC+300 ppm Pr alloys. RE elements exhibit a strong tendency to segregate to the inter dendritic regions.

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89 0 2 4 6 8 10 12 14 0 20 40 60 80 100 120 140 Wt % Cr Co W Ta Figure 4 6. Microprobe analysis line scan results for heavy elements like W, Cr, Co and Ta in as cast CM247LC alloys. Large scale inhomogeneity in the chemical distribution is noticed. Table 4 1. List of various heat treatment cycles used on CM247LC alloys. Refer to the text for the rate of heating between consecutive cycles and rate of cooling after soaking at the maximum temperature 1079 C / 4 hr Aging 1185 C / 10 min + 1232 C / 2 hr 1121 C / 2 hr + 1232 C / 2 hr + 1245 C / 2 hr + 1260 C / 2hr 1190 C / 5 min + 1225 C / 2 hr + 1235 C / 2 hr + 1244 C / 2 hr + 1252 C / 2 hr + 1260 C / 2 hr 1190 C / 5 min + 1235 C / 2 hr + 1244 C / 2 hr + 1252 C / 2 hr + 1260 C / 2 hr 1190 C / 5 min + 1244 C / 2 hr + 1252 C / 2 hr + 1260 C / 2 hr 1190 C / 5 min + 1244 C / 1 hr + 1252 C / 1 hr + 1260 C / 2 hr Mod-SHT01 Mod-SHT02 Mod-SHT03 Mod-SHT04 Type Heat Treatment Cycle* Baseline UFHT

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90 200 m 50 m A B Figure 4 7. SEM micrographs showing IHT alloys, (A) Low magnification ima ge and (B) High magnification image showing the inter residual eutectics were present even after solution heat treatment. 0 2 4 6 8 10 12 0 20 40 60 80 100 120 140 Wt% Co Cr W Ta Figure 4 8. Microprobe analysis showing the distribution of heavier elements like W, Co, Cr and Ta in IHT alloys. In comparison with the as cast results, improvements in the compositional homogeneity were attained.

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91 10 m A B Figure 4 9. SEM micrograph showing the structure of UFHT alloys, (A) Incipient melting and (B) Cracking along the eutectic regions due to the nonuniform deformation during creep testing. Figure 4 10. A typical Differential Thermal Analysis (DTA) curve for a CM247LC baseline alloy showing the important transformation temperatures.

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92 Table 4 10. Eutectic temperatures for various N i RE systems taken from their phase diagrams. Binary System Ni rich side RE rich side Composition (wt%) Eutectic Temp., o C Composition (wt%) Eutectic Temp., o C Ni Y 92 Ni 8 Y 1285 25 Ni 75 Y 805 Ni Ce 81 Ni 19 Ce 1210 11 Ni 89 Ce 477 Ni Pr 77 Ni 23 Pr 1280 9 Ni 91 Pr 460 Ni Dy 83 Ni 17 Dy 1279 14 Ni 86 Dy 693 Ni Gd 87 Ni 13 Gd 1275 14 Ni 86 Gd 635 Figure 4 11. Ni Al phase diagram with the hatched regions showing the Ni rich regions (Reproduced from ASM Handbook on Phase Diagra ms)

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93 Table 4 3. Differential Thermal Analysis (DTA) results showing the various transformation temperatures. RE additions did not result in large reduction in the solidus temperatures. Ce Pr Dy Gd Solidus Liquidus 1 1204 1299 1381 2 860 1204 1289 1377 3 160 1201 1289 1379 4 180 1199 1291 1379 5 120 1197 1288 1378 6 260 80 1206 1289 1377 7 250 70 1207 1287 1378 8 350 70 1228 1294 1376 9 70 80 1235 1289 1378 10 70 47 1201 1289 1378 11 80 70 1211 1288 1379 12 150 70 50 1206 1290 1378 13 200 60 50 1212 1289 1378 14 200 70 50 1206 1279 1378 15 60 50 50 1230 1288 1377 16 300 1207 1290 1379 Baseline RE # Temperature (C) 20 m 2 m A B Figure 4 12. SEM micrographs showing the structure of alloys subjected to ModSHT03 solution heat treatment. Compared with the previous heat treatment, a higher degree of microstructural homogeneity was achieved through this modified heat treatment.

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94 0 2 4 6 8 10 12 0 20 40 60 80 100 120 140 Wt % Co Cr W Ta Figure 4 13. Microprobe analysis showing the distribution of heavi er elements like W, Co, Cr and Ta in the alloys subjected to ModSHT03. This modified heat treatment resulted in higher degree of compositional homogeneity also. 5 m 1 m A B Figure 4 14. SEM micrographs showing the fully heat treated microstructure in (A) low magnif ication and (B) higher magnification micrograph precipitates.

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95 2 m A B Figure 4 15. SEM micrographs showing the Hf Ce Gd phase in the alloys subjected to annealing heat treatment in (A) secondary electron mode and (B) back scattered mode. The phase in the middle is a MC type carbide rich in Hf. Hf Pr Ni 1 m Figure 4 16. SEM micrographs and X ray mapping images showing the regions richer in Ni, Pr and Hf. Complete absence of Ni is evident in the X ray maps.

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96 A B 10 m 2 m Figure 4 17. SEM micrographs showing the presence of Hf Ce Dy phase in the grain boundari es of the alloy subjected to ModSHT02 solution heat treatment. 5 m A B Figure 4 18. FIB images showing the Hf Ce Dy phase at the grain boundaries in (A) Ion beam mode and (B) Electron image mode. A three dimensional view of the blocky phase is illustrated in this image.

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97 Table 4 4. List of vacuum annealing treatments conducted to understand the RE Hf phase evolution. Refer to the text for the rate of heating between consecutive cycles and rate of cooling after soaking at the maximum temperature 1190 C / 5 min + 1225 C / 1 hr + 1232 C / 1 hr + 1245 C / 1 hr + 1252 C / 10 hr RE-Hf 01 RE-HF 02 Mod-SHT03 Mod-SHT04 1190 C / 5 min + 1225 C / 30 min + 1232 C / 30 min + 1245 C / 30 min + 1260 C / 8 hr 1190 C / 5 min + 1225 C / 2 hr + 1232 C / 2 hr + 1245 C / 15 hr 1190 C / 5 min + 1225 C / 4 hr + 1232 C / 20 hr Type Vacuum Annealing Treatments 0.5 m / RE Hf A B Figure 4 19. TEM micrograph showing the Hf Ce Dy phase in (A) Bright field mode and (B) Dark field mode. Interestingly, a bi crystal gra in was observed and better grain contrast was obtained using tilting to different zone axis in the TEM.

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98 0.2 m A B Figure 4 20. TEM micrograph showing the grain boundary of the Hf Ce Dy phase (A) and the corresponding EDAX spectrum taken inside one of the grains (B) showing the presence of various elements. Cu signals were from the TEM sample grid.

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99 CHAPTER 5 OXIDATION BEHAVIOR OF UNCOATED ALLOYS In the quest to improve the high temperature mechanical strength, oxidation properties were largely an after thought during the development of Ni base superalloys [12] Nevertheless, among the class of high temperature materials, Ni base superalloys possess satisfactor y oxidation resistance due to the presence of Al and Cr. This property was considered to be one of the barriers preventing the development of alternative high temperature alloys to move beyond Ni base superalloys [10] Largely attributed to the significant developments in the ceramic based thermal barrier coatings, modern day gas turbines were able to achieve an operating gas temperature of ~1600 oDue to the following reasons, the presented work is unique compared to the previously published works on the similar subject; C [51] But, overlay coatings should not be considered to be prime reliant and any local spallation of the coatings would accelerate the failure of the superalloy substrate. In the light of this possibility, improving oxidation resistance of the underlying substrate also should be given due importance. Minor additions of rare earth elements were found to reduce the scale growth rate and increase the oxide scale adherence, considered to be two basic requirements for optimal oxidation resistance. Refer Chapter 2 for a detailed review on the significant impact of such minor additions on the high temperature oxidation properties. The objective of the chapter is to evaluate the oxidation properties of CM247LC with minor additions of Ce, Pr, Dy and Gd. A commercially viable high temperature alloy was investigated compared to m ost of the previous work which focused exclusively on simplified model alloys of the type Ni Al, Ni Cr or Ni Cr Al which typically formed either Al2O3 or Cr2O3 Owing to the desired longevity of the components used in the gas turbines, it was important to understand the long term effects of any alloying addition meant to improve layers for oxidation protection.

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100 the oxidation resistance. This chapter includes the results of the 5000 to 10000 hours oxidation testing conducted to evaluate the RE effect. The i sothermal and cyclic oxidation properties of rare earth modified alloys are compared to a baseline alloy. Results of importance in this chapter are oxidation kinetics, oxide microstructure evolution as a function of time and the segregation of rare earth e lements to the various interfaces namely gas/scale, scale/alloy and the oxide grain boundaries. 5.1 Isothermal Oxidation Testing 5.1.1 Oxidation Kinetics The weight gain profiles of the alloys subjected to 1010 oC isothermal oxidation are shown in Figure 5 1. As indicated in the figure, during the first 100 hours of exposure, all the alloys were observed to exhibit a linear increase in the weight, usually termed as the transient oxidation phase followed by steady state oxidation with a reduced scale growth rate. Of the alloys exposed at 1010 oThe rapid weight gains observed during the transient stage were attributed by the faster access of oxygen t o the substrate surface leading to the nucleation and growth of the oxides of all the constituent elements. In addition, the transient phase is underlined by the observation of spinel oxides of the type AB C, CM247LC baseline and 860 ppm Ce containing modification experienced rapid increase in specific mass gain compared to Pr, Dy and Gd containing variants which exhibited similar but reduced oxidation kinetics. Among the CM247LC alloys containing Pr, the alloy with 300 ppm Pr showed higher scale growth rate during the later stage of oxidation matching up with baseline and Ce containing alloy. The non uniform weight gain behavior in some of the alloys characterized by wavi ness of the weight vs. time graphs was attributed to spallation and re growth of the oxide scales, more typical of a dynamic process like oxidation. 2O4 composed of the constituent elements and various metastable phases of Al2O3. The extent of the transient stage depends on various factors such as alloy

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101 chemistry and exposure temperature. In general, the higher the temperature of the exposure, the shorter the term of transient stage. The steady state t erm is characterized by the growth and establishment of a slow growing and stable oxide scale. The diffusion of oxygen and alloying elements through these scale during this steady state condition determine the oxidation kinetics. Thus, the alloys exposed t o oxidation environment exhibit a typical parabolic behavior of weight gain during isothermal oxidation. From Figure 5 1, the oxidation rates can be represented by the following equation; n pt K A m (Eqn. 5 1) where p is the parabolic rate constant, t is time of oxidation and n is the rate exponent. The value of n was calculated from the slope of log ( during oxidation would result in a n value of 0.5 The values of the rate exponent, n tabulated in Table 51 were less than 0.5 which showed that all the alloys followed a subparabolic oxidation behavior. Generally, it is not uncommon to assume the oxidati on behavior to be parabolic while trying to determine the oxidation rates. In order to obtain the value of Kp, as shown in Figure 5 2. The slopes were calculated from these profiles and were tabulated along with rate exponent values in Table 51. Note that the values of n and KpComparing the scale growth rates of the alloys oxidized using K were calculated using the method of l east squares fitting with a correlation coefficient greater than 0.9. p values, the addition of 860 ppm Ce to CM247LC resulted in an increase in the oxidation kinetics by nearly 50% while the other RE alloying additions were beneficial and reduced the growth rate by about 25% on average.

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102 Plotted in Figure 5.3 are the specific mass gain data for isothermal oxidation conducted at 1079 oFollowing an initial shorter span of accelerated oxidation kinetics, all the alloys experienced weight losses. As included in the figure, there were three regimes that were apparent namely Regime I with the rapid weight gain during the first 25 hours follow ed by II with severe spallation, and finally, III with steady state characteristics. Similar to the 1010 C. The data are more representative of a cyclic oxidation testing that reflects the oxidation kinetics as a function of both the growth stresses and the thermal cycling stress, the latter more responsible for the spallation of the oxide scale. But, a weight gain followed by immediate weight loss characteristics dur ing isothermal oxidation at higher temperatures is not uncommon and is alloy dependent [115, 166] oC results, the CM247LC+860 ppm Ce containing alloy exhibited a more rapid growth rate and suffered significant weight losses compared to th e other alloys. Surprisingly, the baseline alloy, in spite of showing slower oxidation kinetics in regime I experienced heavy spallation in Regime II. Similar to the 1010 oC testing, the other RE modified alloys showed similar behavior with smaller weight losses in regime II and were observed to generate more stable oxide scales for oxidation protection. Using the oxidation rate kinetics method again, the oxidation behavior between all the alloys were compared for the weight gain rate in regime I and weight loss rate in regime II. The calculated values are shown in Table 5.2. On an average, the RE modified CM247LC alloys showed an increase in the oxidation rate over baseline by about 50% in the regime I. CM247LC+860 ppm Ce alloy exhibited nearly 50 % increas e in the weight loss rate compared to baseline in regime II, whereas other alloys had a weight rate loss decreased by 30% during the same period.

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103 Commenting on the cyclic oxidation like behavior, at higher temperatures the coefficient of thermal expansion between the oxide scale and substrate plays a role in the spallation properties during cooling. After the first 25 hours of exposure, the sample s were unloaded for weighing. The scale might have cracked due to the cooling stresses and subsequently spalled off. The differences in spallation behavior due to test temperatures could be attributed to the severity of cooling stresses and difference in the coefficient of thermal expansion, both of which depend strong on temperature. This is one of the reasons why Thermo Gravimetric Analysis is employed on small scale tests to study the oxidation kinetics at the testing temperature directly instead of having to cool down for weighing and face the effect of cooling stresses. In spite of such a shortcoming, similar p rocedure employed for all the samples still provided a reliable means of comparison to evaluate the RE effect. 5.1.2 Gas/Scale interface microstructure Due to the large number of samples involved in this oxidation analysis, only a selected group of samples were chosen to be characterized for th e microstructural features that could provide an understanding of the RE effect on the oxidation behavior. As mentioned earlier, the multiple constituents of the superalloy chemistry would lead to the formation of several spinel oxide phases. Before going into the details of the microstructure, the three most prominent features observed on the gas/scale interface, irrespective of the alloying additions and time of exposure, are presented in Figure 5.4. Seen at the gas/ scale interface is a NiO layer and at the scale/alloy interface is an Al2O3 with a spinel oxide of the general type (Ni, Co) (Cr, Al)2O4 sandwiched between them. Using the traditional explanation for oxide growth owing to the faster diffusion of Ni2+, NiO forms a complete l ayer at the gas/scale interface, offering little protection. Slow growing compact oxides such as Al2O3, form at the oxide/alloy interface, providing a barrier for the inward diffusion of O2 and outward diffusion of other oxygen ac tive

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104 elements. An exclusive Cr2O3 layer formation was not noticed, and Cr was found mostly incorporated as a part of these three layers. Mostly, Cr was found incorporated in either NiO or Al2O3 Examination of the gas/scale interface after 50 hours of exposure in the baseline CM247LC alloy showed a NiO scale with well defined multi faceted grains with an average grain size of 3 5 m Generally, the layer was non uniform with protrusive grains, which was found to be extremely porous in agreement with a previously reported work [158] During oxidat ion, NiO is the first oxide layer to be established as a continuous scale. The grain boundaries of the alloy play a role during oxidation in providing faster diffusion paths for the species at higher temperature. But, the much fi ner grain size of the NiO s cale, compared to the alloy substrate indicated that there is no dependence of the oxide scale growth on the texture of the alloy. Along similar lines, the finer grain size of NiO was made possible by the abundance of Ni at the gas/scale interface that he lped in the nucleation of large number of NiO immediately upon exposure. During the same period of oxidation, Pr and Ce containing alloys exhibited a discontinuous NiO scale that often spalled, only to be replaced by Al scales. The salient features of the above discussed oxide mo rphologies will be detailed in discussion below. 2O3 scales. The spallation pattern of the NiO scale after 200 hours of exposure is shown in Figure 5 5. The multi facets of the grain s were replaced by uneven brick like layer with almost flat surfaces. With cracks running along the flat edges, it was assumed that the observed feature was an intermediate step in the spallation process of the oxide. With increase in oxidation times, the NiO layers were found to be discontinuous, with the edges of the grain more rounded and the grain size coarsened, as shown in Figure 5.6. The arrows in the figu re indic ate steps on the NiO grain that were representative of their crystallographic growth mode.

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105 The role of spinel oxides on the oxidation mechanism is not yet fully understood. It is generally believed that due to the defect structure, the diffusion th rough spinels are more rapid compared to Al2O3 and Cr2O3 scales [12] In the current analysis, the spinel oxide scales were not observed to exhibit any consistent morphology. One of the observations was the presence of a ridge like structure rich in Ni and Cr observed in baseline CM247LC, presumably a NiCr2O4 spinel, as shown in Figure 5.7 A. Due to the convoluted shape of the structure, it was assumed that they had formed along the grain boundaries of the substrate during the initial stage of oxidation. Grain boundaries provide faster paths for the diffusion of Mn+, and unless blocked by the presence of RE elements, their growth continues unabated [68, 94] Due to the increased thickness and associated stress buildup at the interse ction of oxide scales with the substrate grain boundaries, local detachment and spallation of the scale could occur. Spinels with such ridge like morphology were not observed on the RE modified alloys. In Figure 5.7 B, one observes a very porous spinel oxi de, which was found to be NiAl2O4The above discussed NiO and spinel oxide scales are fast growing, largely in part to the loosely packed struc ture t hat facilitates the faster diffusion of the ionic species through them. Eventually, the faster growing scales spall, and the oxidation reaction is primarily controlled by the diffusion through a dense, stable, relatively defect free and nonvolatile Al based on the EDS results. Several elemental peaks were observed in the EDS results which indicates that these other elements were incorporated into the spinel oxide through diffusion from the substrate or these scales were in fact, transition alloy oxides. XRD results that will follow shortly would be instrumental in confirming the presence of spinel oxides. 2O3. This signifies the onset of steady state oxidation as already shown in the Figure 51. The cracking and spallation of the transient oxides after 200 hours of oxidation is depicted by the microstructure

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106 given in Figure 5.8. The Al2O3 was darker i n contrast with colonies of alumina nucleating and growing to form a flat scale. Al2O3 scale observed here was very different compared to the scale observed on exclusiv ely alumina forming alloys that had similarities with the NiO scales observed in this study [68] The gas/scale interface of CM247LC baseline alloys oxidized at 1079 oC exhibited A l2O3 rich whiskers as shown in Figure 5.9, specifically at locations wher e the transient oxide scale has spalled. On closer observation, these whiskers were found to be growing from the previously established Al2O3 layer. This type of morphology is typical of an Al2O3 that grows predominantly by the outward diffusion of Al3+ through the Al2O3 grain boundaries [9] Al2O3 phases grow by the outward diffusion of Al3+ but Al2O3 growth progresses through inward diffusion with very minimal outward diffusion. At such higher temperatures and longer periods of exposure, it is unlikely that the metastable phases would exist which leaves only the outward diffusion of Al3+ as the only possible mechanism of whisker growth in this case. The addition of RE elements suppressed the outward diffusion of Al3+Furthermore, the Al and the whisker morphology was not observed at the gas/scale interface in the RE contain ing alloys. 2O3 also exhibited an unusual morphology due to the presence of a very pronounced wrinkling, as shown in Figure 5 10. As discussed earlier in previous section, CM247LC+860 ppm Ce alloys displayed more rapid oxidation kinetics during the 1010 oC oxidati on exposures. The higher oxidation rate is a result of the thicker oxide which leads to greater compressive stresses in the scale. Cracking, wrinkling, rumpling leading to detachment are one of the many ways through which the oxide scale can relax their in ternal stresses [94, 167, 168]

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107 To summarize the changes occurring at the gas/scale interface as a function of exposure time, refer to Figure 5.11 which shows the microstructure during the initial 200 hours where oxide transformed to a stable Al2O3 at 1010 o5.1.3 Scale/alloy interface microstructure C. At 1010 oC, CM247LC+860 ppm alloys exposed for 50 hours showed an extensive vein like network of Al2O3 shown in Figure 5.12 A. On further oxidation, due to the continuous inward diffusion of O2 -, the internal oxide veins transformed into a complete Al2O3 scale as seen in Figure 5.12 B. Comparing the thickness difference between the previously established Al2O3 network and the newly formed Al2O3 scale showed that the spallation of the scale had taken place. This internal oxide network creates detrimental out of plane tensile stresses that lead to spallation of the oxide scale. This spallation and possible re growth of the oxide scale combined with heavy internal diffusion of O2 explains the faster oxidation kinetics of Ce modified alloy. This type of network was not observed in the CM247LC baseline alloy oxidized under similar conditions. Oxide scales formed during 1079 oIn the transition oxides, one of the interesting features noticed in 1079 C expos ure were observed to be very irregular and were primarily transition oxides of elements Ni, Cr, Al, W and Hf in both the baseline and RE modified alloys. oC case which was not very pronounced in 1010 oC was the presence of W rich oxide phases found between NiO and NiAl2O4 spinel phase as shown in Figure 5.13. A continuous layer of NiO was not observed but it was observed beneath the spinel oxide scale. Also seen in the figure is the thickening of the Al2O3 during longer exposure time. It has to be mentioned that W forms a volatile WO3 oxide phase which is detrimental since it leads to the depletion of an important solid solution strengthening element.

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108 It is widely understood that an y residual nonuniformity in the chemical composition would lead to differences in the oxidation kinetics at the alloy surface [158] All the alloys we re subjected to Industrial solution heat treatment and as discussed in Chapter 4, this solutioning was lso have an impact on the oxidation microstructure. A thicker oxide nodule was formed in the vicinity of the residual eutectic due to the presence of heavier elements like Co, W etc. compared to a uniform scale in locations t hat were relatively homogeneous in chemical composition. Very deep internal oxidation stringers were also observed after exposures at 1079 o5.1.4 XRD analysis of oxide phases C. This feature is covered in greater detail in the section describing the cyclic oxidation behavior. The previ ous section outlined the observation of several complex spinel and transition oxides that were difficult to be studied using SEM EDS techniques and thus, demanded the use of X r ay diffraction for identifying the various oxide phases. XRD patterns obtained from CM247LC and CM247LC+860 ppm Ce alloys exposed at both 1010 oC and 1079 oC for various exposure i ntervals are shown in Figures 5 14 and 515 respectively. The peaks identified belong Al2O3, NiCr2O4, NiAl2O4, HfO2 Al2O3 were the major peaks with the peak intensity of the latter increasing with the exposure time. In addition, other peaks were also increasing in intensi ty due to their progressive growth and thickening as a function of oxidation time. Due to this increasing thickness of the oxide, the alloy substrate peak decrease correspondingly. The presence of HfO2 was attributed to the higher af finity of Hf2+ towards O2 and HfO2 was observed during the microstructural characterization to be dispersed within the oxide or as internal oxidation stringers. Comparing the peaks between the two temperatures, no visible difference was readily apparent. In

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109 Al2O3 peaks were very prominent due to the extensive spallation of the NiO scales during 1079 oC oxidation. In spite of the higher oxygen affinity of Ce, CeO25.1.5 RE segregation to the oxide scale peaks were not observed, probably due to the lower fraction of suc h oxides which could be well below the detectable limit of the instrument. The segregation behavior of RE elements to the oxide scale grain boundaries and interface is one of the most popular mechanisms used to expla in behavior of RE containing alloys [9, 68, 69, 77, 104, 106] Detailed microstructural analysis using SEM was conducted on various oxide phases at the gas/oxide interface for identifying the presence of RE elements. P resence of Ce and Pr was detected on the oxide scales of CM247LC+860 ppm Ce and CM247LC+300 ppm Pr al loys respectively at both 1010 oC and 1079 oC. See Figure 5.16. The presence of RE was observed, but only on the transition oxide scales and probably in solution with other constituent elements. Since SEM is not capable of detecting ionic segregation, it would be more appropriate to consider the presence of RE elements as either CeO2 or PrO25.2 Ini tial Stages of Oxide Scale Development phases precipitated at the gas/scale interface. In spite of extensive analysis on alumina scales for the presence of similar precipitated RE oxides, it was not found at the interface examined in this work. Even on the transition oxide scales, RE oxides were not uniformly precipitated, primarily due to the ppm le vel additions of these elements, which indicate that the levels examined in this work were too low. The transient stage of oxidation involving the faster rate of oxide growth extended for ~100 hours during 1010 o To understand the oxide phase and mor phological evolution during the initial few hours of transient stage of oxidation C isothermal exposure. The objectives of this part of the work are:

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110 To study the segregation and precipitation of RE oxides to the gas/scale interface during the same period. The weight gain s of the samples were not measured and the exposur es were performed on CM247LC+260 ppm Ce+80 ppm Pr alloy at 1010 o5.2.1 Microstructure and Phase Analysis C in a box furnace. Figure 5.17 shows the microstructure at the gas/scale interface for the alloys exposed to 2, 5, 10 and 25 hours of isothermal oxida tion. After 2 hours, the surface of the alloy was converted into a uniform layer with the oxide nuclei of the various substrate constituents. Nodules of spinel oxides rich in Ni, Co, Cr and Al were also noticed randomly distributed on the surface. Acicular Al2O3 Al2O3 phase [58, 60, 61] The metastable phases were reported to be stable for a longer time at lower exposure temperatures [60] On closer obs Al2O3 needlelike oxides were aligned in a specific direction and it is possible that the alignment was along the polishing marks as noted previously [94] After 25 hours of oxidation, the entire surf ace was found to be covered with a uniform NiO scale, whose morphology was explained in detail in the previous section on NiO peak appearing after 25 hours oxidati on and also had weaker peaks of Al2O3 5.2.2 Analysis of Rare Earth Oxide Precipitation for all the conditions. SEM examinations of isothermally oxidized samples showed the precipitation of RE oxides or the presence of RE elements in solution within the transition oxides at the gas/scale interface. The previously discussed results were obtained from SEM and XRD techniques which have a depth of resolution about 12 m from the surface of the oxide layer. To understand the composition of the surface layer a nd also to gain more insights into the presence of RE elements

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111 at the gas/scale interface, Auger Electron Spectroscopy was employed. Auger Electron Spectroscopy (AES) typically collects signals only from a few Angstroms deep into the sample and would be ve ry helpful in determining the surface composition. Due to the higher sensitivity of the AES techniques, ppm level additions could also be reliably detected. The Auger scan results for CM247L+260 ppm Ce+80 ppm Pr alloy exposed for 2, 5, 10 and 25 hours are shown in Figure 5.18. To obtain the unoxidized surface composition, unexposed alloy was also subjected to the Auger scan and the results included in the figure. As oxidation progresses, a predominantly Ni rich surface were transformed to Al rich surface wi th oxides containing Cr and Ni also growing on the surface. In the alloys oxidized for 10 hours, a Cathode Luminescence (CL) effect was observed. This interesting observation, which was signified by a small circular green light emitted from the surface, wa s linked to the presence of Ce on the surface. Figure 5.19 shows the Auger results gathered from the gas/scale interface where the Luminescence effect was noticed. Due to the lower level of Pr in the oxide, the appropriate peak was not detected. This obser vation is supportive of SEM results of the possible RE oxide precipitation at the gas/oxide interface. Limited analysis was conducted using an SEM equipped with a CL detector to study the distribution of RE oxides on the surface using CL effect. But, the r esults were inconclusive and will therefore not be discussed further. 5.3 Cyclic Oxidation Behavior Two types of stresses need to be considered while subjecting the alloys to cyclic oxidation, namely, (1) growth stresses that arise due to the thickening of the oxide scales and (2) thermal stresses that result from cycling the alloys between room temperature and test temperatures. Other factors that influence the oxidation performance of the alloys are the difference in thermal expansion coefficient between the alloy and oxide s cale, high temperature mechanical strength of the base alloy and the rate of oxide scale growth. To a major extent, the isothermal oxidation

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112 behavior of the alloys discussed in the previous section was influenced by the growth stresses The operating conditions existing in the gas turbine engines impose severe thermal cycling stresses on the alloys, which make it important to understand the oxidation behavior of the alloys under the combined influence of growth and thermal cycling stres ses. With this objective in mind, the alloys were exposed to cyclic oxidation conditions at 1121 oC and 1150 oThis section outlines the oxidation kinetics and microstructure with emphasis on internal oxidation pegs and elemental distribution across the gamma prime denuded zones C 5.3.1 Oxidation Kinetics In the first place, it is important to answer why a 24 hour cyclic oxidation testing procedure was chosen instead of widely used and reported 1 hour cyclic oxidation procedure. The reasons are as follows ; 1 hour thermal cycling conditions are more represe ntative of the aircraft engines, whereas the alloys being investigated in this research are meant for industrial gas turbines where longer thermal cycling conditions exist. This makes the choice of 24 hours testing procedure more appropriate. Another advantage that was mentioned earlier is the objective of this analysis to study the long term effect of alloying additions on the high temperature properties. Since we are dealing with exposure time of about 5000 to 10000 hours, 1 hour cycling would be more demanding for the oxidation testing equipment. Due to the large scale of the alloy samples involved in this study, scheduling the weighing time table for all the alloys is another difficult process to be con sidered. For small scale laboratory level testing for shorter exposure periods, 1 hour cycling would be more relevant and easier to process. The results of the cyclic oxidation testing conducted at 1121 oC on the alloys with single RE additions are shown i n Figure 5.20. From the beginning of oxidation exposure, all the alloys experienced spallation at a high rate until about 500 hours, with the heaviest rates of spallation observed in the CM247LC and CM247LC+860 ppm Ce alloys. On closer examination, Ce and Pr containing alloys showed an earlier weight gain, immediately followed by spallation of the

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113 oxide scale, as shown in the inset in the Figure 5.20. The weight loss behavior could be divided in three regimes namely Regime I with a higher rate of weight los s, Regime II with a moderate degree of oxide regrowth occurring in the alloy and finally Regime III which was more similar to the steady state oxidation phase observed in 1010 oC isothermal oxidation exposure. Although the baseline alloy sustained its we ight loss during Regimes II and III, CM247LC+860 ppm was stable during II phase followed by uncontrolled spallation all through the III regime. On the other hand, the other RE modified alloys showing excellent oxidation behavior in Regimes II and III with CM247LC+180 ppm Dy alloy showing a higher degree of oxide re growth for high temperature protection. To compare the differing rates of spallation of alloys during Regime I, KpAt 1150 was calculated from specific mass gain vs. square root of time plot as presented in Table 5.3. CM247LC+860 ppm Ce showed ~70 % increase in the weight loss rate compared to baseline whereas other RE modified alloys showed similar weight loss rates as the baseline alloy itself. oC cyclic oxidation exposure, CM247LC and CM247LC+ 860 ppm Ce alloys exhibited weight loss and spallation very similar to the previous case as shown in Figure 5.21. All the alloys suffered weight loss at similar rates in Regime I, whereas Regimes II and III were not similar to the previously observed behavior for 1121 oIt has to be mentioned that the oxidation testing was conducted at 1079 C case. Other RE modified alloys showed better oxidation behavior with mostly steady state oxidation during Regimes II and III. Comparing the weight loss rate in the Regime I for all the alloys, CM247LC+860 ppm Ce alloy exhibited a weight los s rate ~40 % higher than the baseline whereas other RE modified alloys showed a decrease in the weight loss rate by ~ 30 %. The values are given in Table 5 4. oC, 1121 oC and 1150 oC for all the 16 different alloys. For the sake of simplicity, only the results of the alloys with single RE elements are discussed in this section. Since it was observed that the kinetics data

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114 obtained from 1079 oC was not conclusive, the data are incorpora ted into the Appendix A for reference. The 1121 oC and 1150 oAt 1121 C cyclic oxidation data for the alloys with more than one RE additions are included in the same section. It should be noted that alloys with multiple RE elements exhibited cyclic oxidation behavi or very similar to the single RE elements in the current discussion. Based on the oxidation performance during the cyclic testing combined with adhesion behavior of the TBC coated alloys, 5 alloys were downselected and subjected to long term 10, 000 hour c yclic oxidation testing to re evaluate the oxidation kinetics. The following section describes the test results of the second phase studies. oC, CM247LC with one RE addition (300 ppm Pr), two additions (260 ppm Ce + 80 ppm Pr) and three additions ( 60 ppm Pr + 50 ppm Dy + 50 ppm Gd) performed better with steady state oxidation resistance following an initial weight loss as shown in Figure 5 22. The baseline CM247LC and other RE containing alloys with two additions (250 ppm Ce + 70 ppm Dy) and three a dditions (200 ppm Ce + 60 ppm Pr + 50 ppm Dy) continue to suffer spallation through the entire oxidation testing exposure. It should be noted that all the tested alloys exhibited weight loss during 1150 oC long term testing with only CM247LC+60 ppm Pr+50 ppm Dy+50 ppm Gd showing a moderate arrest in the weight loss. Refer Figure 5.23. Expectedly, the weight loss rates were very high at 1150 oC compared to the results from the cyclic oxidation at 1121 o5.3.2 Oxide Microstructure C. The evolution of oxide mi crostructure and phases during the transient and steady state oxidation was covered in the section explaining isothermal oxidation behavior. The difference here is the inclusion of thermal stresses promotes the cracking at the scale/alloy interface, lead ing to spallation. Only the scale/alloy cross section was characterized and some of the salient features observed in the alloys are explained here as under.

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115 A typical cross section of an alloy exposed to cyclic oxidation is shown in Figure 5.24. Despite t Al2O3 scale, a complete decohesion between the alloy and oxide was observed after exposure for 196 hours at 1150 oC. Though the oxidation temperature during isothermal exposure were lower, such widespr ead decohesion was not observed, which clearly indicates the effect of thermal stresses on the scale spallation. Transition oxides, similar to the ones noticed during isothermal oxidation, were also noticed but to a limited degree. The difference in the scale/alloy interface micr ostructure between baseline and RE modified alloys can be seen in Figure 5 25. The cracked open interface of the baseline alloy showed the exposed alloy surface with relatively smoother areas. On the other hand, CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd allo y exposed to 196 hours at 1150 o5.3.3 Internal Oxidation Pegs C showed oxide imprints that were representative of the areas with better oxide scale adhesion [89, 94] Proposed by several authors as one of the structural features mechanically anchoring the scale to the substrate [7, 92, 93, 9597] a large population of oxide pegs were observed in the cross section extending from the oxide scale into the alloy substrate, with the length increasing as a function of exposure time. Looking back at Figure 5.25, the internal oxidat ion pegs were in fact related to the locations where the oxide scale had connected to the alloy. The stringe rs consist of HfO2 encapsulated inside an Al2O3 layer as shown in same figure The mechanism of the peg formation and their relevance in the oxidati on resistance will be presented in Chapter 9: Discussion. The difference between CM247LC and the RE modified alloys in terms of the peg depth and density, was not apparent. Isothermal oxidation at 1079 oC resulted in relatively few pegs present compared to the cyclic oxidation case, meaning the mechanism of peg formation was largely temperature dependent.

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116 5.3.4 Gamma Prime Denuded Zones (GDPZ) One of the characteristic features of oxidation microstructure in Ni base superalloys is the Gamma Prime Denuded Z ones (GDPZ ). Ni precipitates which have a L12 structure, a FCC like structure with Ni at the face centers and Al at the face corners. During the oxidation process, due to the higher affinity of Al3+ ions towards O2 -, outward diffusion of Al3+ is observed towards the oxide/alloy interface and then onto the oxide itself. This leads to the depletion of s in the regions immediate ly beneath the oxide scale that could lead to weakening of the alloy at the surface. Thus, it is important to study the role of RE elements in containing the outward diffusion of Al3+The extent of depletion was measured on the cross sectional microstructures using a software application, I mageTool The distance from the oxide al loy interface to the GDPZ structure interface was measured by constructing a straight line and the average values measured for the alloys exposed to 1150 ions by evaluating the GPDZ thickness. oC for 196 and 1176 hours are shown in Figure 5.26. The measured values seemed to correlate fai rly well with the extent of weight loss at 1150 oTo study the depletion of Al in the GDPZ, Electron Probe Micro Analysis was conducted e/alloy interface. As shown in the Figure 5.28, there was a depletion of Al in the region close to the oxide/alloy interface and it was increasing with the time of oxidation. The peaks showing increased Al presence was due to the presence of s C at specified times. CM247LC, despite showing very lower depletion, suffered extensive spallation after about 2000 hours. CM247LC+60 ppm Pr+50 ppm Dy+50 ppm Gd suffered a smaller amount of depletion and als o exhibited better oxidation resistance. Advanced metal loss calculations involve oxide scale thickn ess and weight loss considerations. The measurement

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117 along t he line of the microprobe scan, which were smaller than the regular size. Noticeable changes were not observed in the distribution of other elements like Ni, Cr, Co and W while minor elements like Hf, Zr and Pr were not detected during the analysi s. The distribution pattern of these elements is included in the Appendix section. 5.4 Summary Improving the oxidation resistance is integrally coupled with reducing the oxidation kinetics and increasing the oxide adherence under thermal cycling conditions Isothermal and cyclic oxidation testing were conducted to study the effect of RE elemental additions on these two important requirements. With the exception of CM247LC+860 ppm Ce containing alloy, other RE modified alloys exhibited better oxidation beha vior with reduced oxidation kinetics during isothermal oxidation testing. The addition of multiple RE elements was also found to improve the oxidation resistance during cyclic oxidation testing. Due to the complexity of the alloy chemistry, transition oxidation was predominantly Al2O3 was generated to provide a barrier for the counter diffusion of Mn+ (M Metal) and O2 One of the important observations was the precipit ation of RE oxides at t he gas/oxide interface, as confirmed using SEM and Auger Spectroscopy.

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118 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 0 200 400 600 800 1000 Time, h CM247LC 120 ppm Gd 300 ppm Pr 860 ppm Ce 180 ppm Dy 160 ppm Pr Specific Mass Gain, mg/cm2 Steady state oxidation Transient stage 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 0 200 400 600 800 1000 Time, h CM247LC 120 ppm Gd 300 ppm Pr 860 ppm Ce 180 ppm Dy 160 ppm Pr Specific Mass Gain, mg/cm2 Steady state oxidation Transient stage Figure 5 1. Specific mass gain results for 1000 hours isothermal oxidation testing conducted at 1010 o C. Faster oxide growth rates associated with transient o xidation slows down to steady state oxidation after about 100 hours. Table 5 1. Rate exponent, n and parabolic reaction rate constant, Kp values for isothermal oxidation testing at 1010 oC. The value of n indicates a subparabolic growth behavior while a ddition of RE elements was found to reduce the oxidation kinetics as given by Kp 160 ppm Pr 180 ppm Dy Alloy 120 ppm Gd 300 ppm Pr 0.332 0.227 0.155 0.189 0.239 0.311 CM247LC 860 ppm Ce 5.889 5.278 Rate exponent, n 7.389 11.139 5.694 4.722 Kp, 10-12 g2/cm4 s

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119 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 0.0 5.0 10.0 15.0 20.0 25.0 30.0 35.0 Sq. Root of Time (hr0.5 ) CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr Specific Mass Gain, mg/cm2 Figure 5 2. Specific mass gain results for 1000 hours isothermal oxidation testing plotted against the square root of time. -1.75 -1.25 -0.75 -0.25 0.25 0.75 0 200 400 600 800 1000 Time, h CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr Specific Mass Gain, mg/cm2 II I III Figure 5 3. Specific mass gain results fo r 1079 oC isothermal oxidation testing. Despite uncharacteristic weight loss during isothermal oxidation, such behavior is not uncommon. Refer to the text for more information.

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120 10 ma d c b 10 m Figure 5 4. Three prominent layers constituting the oxide layers, (a) illus trated using a spalled region at the gas/scale interface, (b) granular appearance of NiO grains with well defined multi faceted grains, (c) convoluted and porous structure of spinel oxide of the type Ni (Cr, Al)2O4 Al2O3 5 m 20 m which e ventually grow to generate a flat protective scale. Figure 5 5. Secondary electron micrographs of CM247LC alloy subjected to 200 hours of oxidation at 1010 o C. (A) Spallation pattern of the NiO scales and (B) Cracks, as indicated by the arrows, running along the smooth plates of spalling NiO layer. The spallation pattern is reminiscent of a crystallographic fracture mode. A B

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121 5 m A 1010 oC B 1079 oC Figure 5 6. Secondary electron micrographs showing the NiO grains after 50 hours of isothermal oxidation (A) Finer grains observed at 1010 oC compared to coarser grains at 1079 o C (B) Porosity in the NiO layer is also very apparent. With increase in the exposure time, the faceted NiO grains assume a rounded surface. 10 m A B Figure 5 7. Secondary electron micrographs showing the spinel oxides at the gas/scale interface, (A) Convoluted morphology of the Ni (Cr, Al)2O4 spinel and (B) Highly porous NiAl2O4 spinel. Diffusion and oxidation mechanism through the spinels is a complex phenomenon and not fully established yet

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122 10 m Figure 5 8. Secondary electron micrograph showing decohesion observed at the interface Al2O3 scale as indicated by the arrow. The different coefficient of thermal expansion between the different oxide layers was cited as one of the prim ary reasons for stress build up in the scale that leads to spallation. 5 m Figure 5 9. Secondary electron micrographs showing the presence of Al2O3 whiskers growing at the gas/scale interface on CM247LC alloy exposed to isothermal oxidation at 1079 oC for 5 0 hours. Formation of whisker morphologies indicates an outward diffusion of Al2O3 which can be retarded by the addition of minor RE elements.

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123 10 m Figure 5 Al2O3 scales observed at the gas/scale interface. Wrinkling was considered to be one of the mechanisms to relax the heavy compressive growth stresses built in the oxide scale. 50 h 200 h 500 h 1000 h 50 m Figure 5 11. Scanning electron micrographs taken from CM247LC+300 ppm Pr alloy showing the changes in the gas/scale interface with increase in the oxidation exposure at 1010 oC. Al2O3 grows as the most stable layer after the spallation of NiO and spinel ox ide.

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124 10 m A B Figure 5 12. Scanning electron micrographs taken from the cross section of CM247LC oxidized at 1010 oC for 50 hours (A) and 200 hours (B). An internal veinlike network of Al2O3 develops at shorter times which transforms to a continuous scale at longer times. Also noticeable in the microstructures are a zone below the oxide scale which diffusion zone. 10 m Al2O3 Al2O3 A B Figure 5 13. Scanning electron micrographs taken from the cross section of CM247LC+860 ppm Ce oxi dized at 1079 oC for 50 hours (A) and 200 hours (B). A thicker oxide scale with layers of outer NiO, middle spinel oxide and an inner Al2O3 was observed. Oxide rich in W was noticed within the NiO layer as indicated by the arrows.

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125 NiO Al2O3 NiCr2O4 NiAl2O4 HfO2 20 30 40 50 60 70 2 Theta 50 h 200 h 500 h 1000 h NiO Al2O3 NiCr2O4 NiAl2O4 HfO2A 20 30 40 50 60 70 2 Theta 50 h 200 h 500 h 1000 h B Figure 5 14. X ray d iffraction patterns taken from CM247LC alloy subjected to oxidation at 1010 oC (A) and 1079 oC (B) for various times.

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126 20 30 40 50 60 70 2 Theta 50 h 200 h 500 h 1000 h NiO Al2O3 NiCr2O4 NiAl2O4 HfO2A 20 30 40 50 60 70 2 Theta 50 h 200 h 500 h 1000 h B Figure 5 15. X ray diffraction pattern taken from CM247LC+860 ppm Ce alloy subjected to oxidation at 1010 oC (A) and 1079 oC (B).

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127 A B 5 m Figure 5 16. Secondary electron micrographs taken from alloys CM247LC+860 ppm Ce alloy exposed to 1079 oC oxidation for 50 hours (A) and CM247LC+300 ppm Pr alloy oxidized at 1010 o 50 m 2 m 10 h 2 h 5 h 25 h 50 m 2 m 10 h 2 h 5 h 25 h C for 500 hours (B). The arrows indicate the presence of higher amount of Ce and Pr in the respective samples. Both the oxide scales were transition oxides and the detection represents the possible precipitation of RE oxides. Figure 5 17. Secondary electron micrographs showing the development of oxide layer during the initial fe w hours of transient oxidation. The needle Al2O3 Al2O3. Nodules of transition oxides were also observed as shown in the micrographs.

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128 0 200 400 600 800 1000 1200 1400 Kinetic Energy (eV) dN(E), (A.U.) 0 h 2 h 5 h 10 h 25 h O Ni Cr Al Figure 5 18. Auger scan results conducted on CM247LC+260 ppm Ce+80 ppm Pr alloy subjected to 1010 oC oxidation for 2, 5, 10 and 25 hours. The scan result on the unoxidized sample is also given as a reference. Metastable Al2O3 50 250 450 650 850 1050 1250 1450 Kinetic Energy (eV) dN(E) (A.U) Ce C Ni Al O phases were found to be very dominant in the transient oxidation stage. Figure 5 19. Auger scan results from the spot which emitted Cathode Luminescence on the CM247LC+260 ppm Ce+80 ppm Pr alloy subjected to 1010 oC oxidation for 10 hours.

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129 -5 -4 -3 -2 -1 0 1 0 1000 2000 3000 4000 5000 6000 Time (hrs) CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 300 ppm Pr 120 ppm Gd -1.5 -1 -0.5 0 0.5 0 50 100 II I III Specific Mass Gain, mg/cm2 Figure 5 20. Results of cyclic oxidation testing conducted at 1121 o C. Heavy spallation of CM247LC+860 ppm Ce and CM247LC alloys were observed with other RE additions improving the oxide adherence under thermal cycling conditions. As shown in the inset, only 860 ppm Ce and 300 ppm Pr containing alloys showed an initial weight gain. Table 53. Parabolic weight loss rates for regime I in cyclic oxidation testing at 1121 oC. Note the Kp Kp (I), 10-12 g2 /cm4s 32.167 54.500 24.056 28.806 29.361 24.944 120 ppm Gd 300 ppm Pr Alloy CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy values are in negative.

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130 -7 -5 -3 -1 1 0 1000 2000 3000 4000 5000 6000 Time (hrs) CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr II I III Specific Mass Gain, mg/cm2 Figure 5 21. Results of cyclic oxidation testing conducted at 1150 o C. CM247LC+860 ppm Ce and CM247LC alloys showed uncontrolled s pallation behavior while the oxidation resistance of other RE modified alloys were observed to be very good. Table 5 4. Parabolic weight loss rates for cyclic oxidation testing for regime I in 1150 oC cyclic oxidation testing. Note that the Kp Kp (I), 10-12 g /cm4s 44.528 63.028 31.694 32.472 34.722 27.333 120 ppm Gd 300 ppm Pr Alloy CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy values are negative.

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131 -6 -4 -2 0 0 2000 4000 6000 8000 10000 Time, h CM247LC 260 ppm Ce + 80 ppm Pr 250 ppm Ce + 70 ppm Dy 200 ppm Ce + 60 ppm Pr + 50 ppm Gd 60 ppm Pr + 50 ppm Dy+ 50 ppm Gd 300 ppm Pr Specific Mass Gain, mg/cm2 Figure 5 22. Results of cyclic oxidation testing conducted on downselected alloys at 1121 o -12 -7 -2 0 2000 4000 6000 8000 10000 Time, h CM247LC 260 ppm Ce + 80 ppm Pr 250 ppm Ce + 70 ppm Dy 200 ppm Ce + 60 ppm Pr + 50 ppm Gd 60 ppm Pr + 50 ppm Dy + 50 ppm Gd 300 ppm Pr Specific Mass Gain, mg/cm2 C. Figure 5 23. Results of cyclic oxidation testing conducted on downselected alloys at 1150 oC.

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132 50 m Figure 5 24. Secondary electron micrographs taken on an alloy exposed to cyclic oxidation at 1150 o Al2O3 scale, a gamma prime denuded oxidation suffered from decohesion at the oxide/alloy interface. 10 m A B Figure 5 25. Secondary electron micrographs showing the cracked oxide/alloy interface. CM247LC alloy showed a relatively smooth interface (A) while CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd showed oxide imprints, which represent areas of high oxide alloy adherence (B).

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133 0 50 100 150 200 250 CM247LC Ce+Pr Ce+Dy Ce+Pr+Gd Pr+Dy+Gd Pr 196 h 1176 h Figure 5 26. Mean Gamma Prime Denuded Zone (GDPZ) thickness for various alloys during cyclic oxidation testing at 1150 o 0 2 4 6 8 0 50 100 150 200 250 300 350 Al (wt%) 196 h 1176 h 5400 h C for 196 and 1176 hours. Figure 5 27. Line scan results using Electron probe microanalysis show ing the depletion of Al in the GPDZ during cyclic oxidation at 1150 oC for 196, 1176 and 5400 hours. Higher level of depletion was noticed with increasing oxidation exposures.

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134 CHAPTER 6 OXIDATION BEHAVIOR OF ALLOYS WITH THERMA L BARRIER COATINGS To satisfy the demands for higher operating efficiencies, the service temperatures of the gas turbines have continually increased since their inception in the 1940s. During the initial phase, the hightemperature operating capability was realized through the design of high strength alloys, advanced processing techniques and efficient cooling systems. With the improvements in these fields potentially reaching the saturating point, but the demands for higher efficiency growing unabated, focus was shifted towards the de velopment of t hermal insulation coatings that could reduce the surface temperature of the alloys. Due to the dramatic improvements in the Thermal Barrier Coatings (TBC) [5053, 169173] modern day gas turbines operate at extreme gas temperatures in the range of 1300 1500 oC, which is higher than the incipient melting temperatures of the superalloys. Chapter 2: Background provided a brief account on the development of these heat resistant coatings. The TBC syst em comprises of four components, namely, the ceramic top coat (TBC itself) for thermal protection, the Bond Coat (BC) that serves as a reservoir for Al, the thermal ly grown oxide (TGO) laye r that forms due to the oxidation of the BC and the superalloy substrate that bears the structural loads. The TGO which is basically Al2O3 layer provides a barrier for the diffusion of O2 and Mn+The results obtained from the oxidation behavior of bare alloy substrates clearly illustrated the beneficial effects of minor RE additions on reducing the growth rate and increasing the adhesion of the oxi de layers under thermal cycling conditions. With the efforts taken to increase the durability and develop TBC into a prime reliant coating system, it becomes imperative to address the effect of substrate additions on the TGO layer that determines the perfo rmance of the ions. The performance and durability of the coatings have been primarily attributed to the growth and spallation of the TGO layer, which has been labeled as a weak link in the TBC syst em [86, 170, 173178]

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135 TBC system. Although there have been several published reports on the effect of alloy substrate on the performance of TBC system [86, 159, 165, 179184] there were only minimal accounts on the effect of minor RE elements incorporated into the substrate. Thus, to address the effect of RE additions on the oxid ation behavior of coated alloys, this specific study was taken up with the following objectives: To study the spallation behavior of the TBC system as a function of RE additions and the exposure temperature. To analyze the changes in the microstructure of YSZ, TGO and BC as a result of high temperature oxidation. Al2O3 To establish the role of other oxygen active elements like Hf and Y on the properties of TGO layer. ) and to the interface between the TGO/BC layer. To realize these outlined objectives, isothermal oxidation experiments were conducted on CM247LC samples coated with TBC. High Resolution Transmission Electron Microscopy (HR TEM) were conducted to study the segregation behavi or of RE elements to the TGO layer, and Secondary Ion Mass Spectroscopy (SIMS) was employed to qualitatively characterize the segregation of RE and Y to the TBC/TGO interface. Comparisons were made between the microstructures of as deposited and oxidized a lloys. 6.1 As Deposited Microstructure of the Coating The cross sectional microstructure of the as deposited coating consisted of a three layered structure with a ceramic coating at the top, metallic bond coat in the middle and a Nibase superalloy substra te, as shown in the Figure 6 1. The ceramic coating deposited for the purposing of imparting thermal insulation to the underlying substrate was composed of a ZrO2 partially stabilized by 8 wt% Y2O3. This coating more commonly referred to as Yttria Stabiliz ed Zirconia (YSZ) was a tetragonal t structure. Deposited using Air Plasma Spray (APS), the 200 300 m

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136 thick ceramic layer exhibited a splat morphology which was typical of a particle by particle deposition during the spraying process. As clearly evident in the figure, the cross section was populated with a significant fraction of porosity and cracks oriented in different directions. One of the reasons attributed to the formation of cracks was the oxidation of the particles during their flight to the sub strate during the air plasma spraying process. The oxide layers around the particle would lead to the poor bonding between the splats resulting in cracks during the thermal cycling process. These structural defects were deliberately engineered into the co ating for the following reasons: (a) Lowering the thermal conductivity of the ceramic structure by deflecting and providing a discontinuous path for the heat radiations or phonons. This property is further enhanced by the presence of the oxygen vacancies and presence of Y2O3D ue to the structural defects incorporated into the YSZ, the ingress of oxygen through the entire thickness of the coating would continue making it oxygen transparent. At such higher temperatures of operation, availability of oxygen at the superalloy subs trate would lead to degradation of the alloy. In a less critical application like diesel engines, thermal protection in solution. (b) Imparting strain tolerance with the residual stresses during thermal expansion and contraction accommodated through the presence of internal cracks and pores. The durability issues of TBC system are primarily related to the different thermo mechanical characteristics of the various layers involved. Higher degree of strain tolerance would be beneficial in improving the lifetime and reliability of the coatings for protection of superalloy substrates. Prior to the deposition of the YSZ, the bond coat surface was shot peened, which le d to the undulations or roughening as seen at the YSZ/BC interface. This roughening was believed to impart a keying in effect of the YZS to the BC, thereby increasing its spallation lifetime.

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137 coatings can be directly deposited on to the alloy substrates. But, in a higher temperature gas turbine application, direct deposition of Y SZ onto the superalloy would lead to strain incompatibility due to the different thermo mechanical characteristics of the two entities. In order to make the superalloy oxidation resistant and at the same time utilize the excellent thermal insulation charac teristics of YSZ, a metallic bond coat layer of the type MCrAlY (M = Ni and Co) was deposited on the superalloy substrate. The composition of the BC layer is Ni ( 2426 %) Co ( 1618 %) Cr ( 9.5 11 % Al ) (0.3 0.5 % Y ) ( 11.8 % Re ) (Note: All the compositions are given in wt %). MCrAlY composition was d Al2O3Two separate layers of MCrAlY coating were dep osited using High Velocity Oxy Fuel (HVOF) process on the superalloy substrates. The layers are not readily apparent in the 100 150 m thick MCrAlY layer as shown in the Figure 61 C. Interestingly, the unheat treated and as deposited BC also was compos ed of a splat structure. The first layer close to the substrate was deposited with coarse powders and the second layer with finer powders. Although the objective of depositing coarser powders was to impart optimal roughness to increase the adherence with s ubstrate, finer powders were intended to provide higher coating density. as the stable protective oxide layer and thus acts as an Al reservoir for high temperature oxidation resistance. The primary objective behind the addition of Y was to getter the S present in the substrate from segregating to the coating interfaces. The detrimental effects of S and the RE effect on S gettering was discussed in detail in Chapter 2. Re was added for the purpose of enhancing the hightemperature strength of the BC layer. During the APS coating of YSZ layer, the substrate + BC were subjected to a coating heat treatment at 1079 oC for 4 hours to improve the bonding quality at the YSZ/BC interface. As observed in several cases, this heating would lead to the oxidation of the bond coat, creating a

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138 thin oxide layer called the TGO at the YSZ/BC interface. In this case, the TGO layer might have been formed during the coating heat treatment, but was not present in the observed cross section. The growth of the TGO layer during the exposure to isothermal oxidation conditions would be discussed in more detail in the following sections. 6.2 Isothermal Spallation Testing The spallation lifetimes of the TBC system c an be evaluated through both the isothermal and cyclic oxidation testing. Ni base superalloys used in the industrial power generation applications are used for longer periods of time under isothermal conditions with relatively few cycles. Thus, it was more appropriate to select isothermal testing for studying the spallation resistance of the coatings. After the deposition of coating layers and secondary aging heat treatment at 870 oC, alloy coupons were placed inside an alumina crucible and exposed to isoth ermal oxidation at three higher temperatures namely, 1010 oC, 1079 oC and 1121 oThe spallation lifetimes of alloys with single RE elements exposed to isothermal spallation testing are plotted in Figure 63. Since the alloys exposed to 1010 C. From Phase I, two groups of alloys were exposed to test conditions alloys with single RE additions and with multiple RE additions. At the three temperatures of oxidation, three samples per each alloy were subjected to oxidation to obtain an average spallation lifetime. With the failure criterion set as the complete decohesion of the YSZ top coat, the testing was interrupted at pre determined intervals and samples taken out of the furnace to check for the adhesion of the coating. Figure 6 2 illustrates the typical as deposited, as exposed and spalled stages of the alloy coupons. Studying the oxidation kinetics was not one of the objectives in this analysis. So, the weight me asurements were not taken into consideration. oC did not suffer spallation even after prolongin g the exposure to more than 10, 000 hours, only the results from other two exposure temperatures will be discussed. The figure also shows the dashed and dotted lines

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139 corresponding to the spallation lifetime of CM247LC baseline alloy at two temperatures. As seen from the figure, all the RE modified alloys exhibited better average coating lifetimes compared to the CM247LC baseline. Contrary to the oxidation results reported in the previous chapter, CM247LC+860 ppm alloy showed the largest gain in the spallati on resistance at 1079 oFigure 6 4 shows the spallation lifetimes of alloys with two or three RE additions subjected to isothermal oxidation. Similar to the pre vious case, alloys exposed to 1010 C but CM247LC+300 ppm Pr clearly showed better properties at both the temperatures of exposure. oC did not suffer failure even after extended span of exposure. Here as well, all the RE modified CM247LC alloys exhibited better TBC adhesi on compared to baseline. A ll the alloys with RE additions performed well at 1079 oC, but CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd showed remarkable spallation resistance at 1121 oSome of the observations from the testing results are summarized below: C with about a 27% improvement in the lifetime over baseline alloy. Based on the reported oxidation performance, alloys were downselected from Phase I for extended cyclic oxidation testing on uncoated substrates. Refer to Chapter 5 for the details of the testing results. Isothermal oxidation exposure at 1121 oC, in spite of being only 42 oC higher than the 1079 o In general, the group with multiple RE additions exhibited better lifetimes compared to alloys with single RE element. C exposure temperature, reduced the spallation life time by about ~ 2500 hours. On the whole, CM247LC+300 ppm Pr and CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloys were observed to exhibit better TBC adhesion properties under high temperature oxidation exposure. 6.3 Microstructural Changes after Oxidation Microstructural analysis on alloys exposed to isothermal spalla tion testing at 1010 oC and 1070 oC were conducted on samples that were pulled out of the testing after 3720 hours. Also, it should be noted that most of the samples oxidized at the higher temperature of 1121 oC

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140 experienced spallation failure after an aver age time of ~1500 hours. The changes in the microstructure at the BC the YZS/BC interface and general insights on the failure of the coatings will be discussed in this section. 6.3.1 Bond Coat / Superalloy s ubstrate interface The microstructural differences between the as deposited and exposed samples, with reference to the NiCoCrAlY bond coat is presented in Figure 6 5. The BC serves as an Al Al2O3 T GO at the YSZ/BC interface. Due to this reason, a slight reduction in the thickness of the BC was observed after the exposure to oxidation. On closer observation of the alloys exposed to 1010 oC, two regions of BC were apparent and are illustrated in Figur e 6 6. Similar to the closer to the substrate. On exposure to oxidation condition and outward diffusion of Al to aid in the formation of a continuous rich needle shaped phases were noticed at the interface between BC and substrate, a thin layer referred to as Inter phase, one of the categories of Topologically Controlled Phases(TCP) that act as stress raisers, detrimental to phases was very low and not considered to be influential on the mechanical properties. Internal oxide particles rich in Al were also noticed at various locations along the length of the BC. Alloys exposed to the higher temperatures of 1079 oC and 1121 oC exhibited exte nsive regions.

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141 One of the very interesting observations was the presence of HfO2 particles and Hf Ce Gd rich phase at the interface between regions I and II. As shown in the Figure 67, these features were often observed associated with internal oxidation stringers of Al2O36.3.2 Thermally Grown Oxide (TGO) Layer Detailed compositional analysis did not divulge the presence of Hf, Ce and Gd as a part of the BC composition. This gives rise to the possibility of diffusion of these elements from the substrate into the bond coat regions. It was shown before that YSZ was composed of a large populat ion of porosity and cracks that were instrumental in providing thermal insulation and strain tolerance. On the other hand, these defects mak e the YSZ transparent to oxygen at higher temperatures leading to the oxidation of the MCrAlY coatings. The composition of MCrAlY coatings were selected to preferentially Al2O3 scale at the inter face Al2O3 layer slows down the outward diffusion of Mn+ ions from the BC and the growth depends primarily on the inward diffusion of O2 Al2O3 layer further limits the degradation of t he bond coat elements to oxidation. Figure 68 illustrates the difference between YSZ/BC interface in the as deposi ted and oxidized state, clearly Al2O3 Al between the YSZ and the BC in the oxidized alloy. The undula tions in the oxide scale were due to the roughening of the BC surface using shot peening prior to the YSZ deposition. 2O3 scales formed at temperatures greater than 1000 oStudying in detail the structure and composition of TGO layer, Figures 69 and 610 shows the plan view and cross sectional microstructures of YSZ/TGO interface. The contrast in C are generally referred to as thermally grown oxide layers (TGO) but the terminology is more commonly used to identify the oxide layers formed in TBC systems.

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142 backscattered imaging mode in SEM helped to identify the three regions namely : (1) the ZrO2 which were remenants of YSZ layer, (2) a continuous but wavy Al2O3 layer and (3) isolated colonies of transition oxides rich in Ni, Co and Cr at the YSZ/TGO interface. Observation of the underside of a spalled YSZ top coat also revealed similar contrast, but the fraction of ZrO26.3.2.1 Haf nium Oxide Stringers phase was much larger. Microstructural analysis on uncoated alloys exposed to cyclic oxidation which revealed the presence of HfO2 particles encapsulated in Al2O3 was reported in Chapter 5. Similar HfO2 Hf is not a constituent element of the MCrAlY bond coat composition. Only the substrate contained 1.5 wt % Hf. internal oxidation stringers were found as a part of TG O at locations where it was extended into the underlying BC layer. This was observed both in the CM247LC baseline and RE modified alloys at all the temperatures of oxidation as shown in Figure 6 11. The observation is very significant for the following rea sons: Formation of HfO2 In the present case, the substrate/BC interface is about ~100 m away from the TGO/BC interface and for the formation of HfO stringers in uncoated alloys was relatively a simple phenomenon due to the availability of Hf in the immediate vicinity o f the oxide scales. 2 pegs, Hf2+The im plications of this observation upon the effect of RE elements and their contribution to the spallation lifetime will be discussed in detail in Chapter 9. had to diffuse through the entire thickness of bond coat at the given temperature and time. 6.3.2.2 X Ray Phase Analysis of YSZ/TGO interface SEM observations showed the presence of three major phases at the YSZ/TGO interface. To analyze the phase formation in detail at various temperatures of oxidation, X ray diffraction was conducted on CM247LC+860 ppm Ce alloy exposed to the following oxidation temperatures and times: 1010 oC / 3720 hours, 1079 oC / 4392 hours and 1121 oC / 1440 hours

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143 and the results are illustrated in Figure 6 12. Alloy subjected to 1010 oC did not suffer spallation and the YSZ layer had to be carefully polished away to expose the TGO layer for XRD analysis. It would have been ideal to have the samples exposed for similar time periods at all the temperatures. The alloys subjected to 1121 oC failed earlier compared to the other two conditions. But, the time required to form similar oxide phases at higher temperatures is decreased compared to that observed after exposure at lower temperatures. In this context, 1010 oC oxidation would require considerably longer extent of exposure to produce intensity similar to 1079 oC and 1121 oC. By the same reasoning, the dominant phase at 1010 oC w Al2O3 phase while several oxides of MCrAlY composition namely NiO, CoO combined with spinels like NiCr2O4 and NiAl2O4 were observed. The ZrO2 and Y2O3 peaks were primarily from the remenants of YSZ top coat on the TGO layer. There was also a poss ibility of Y in the MCrAlY Al2O36.3.2.2 Thickness of TGO Layer several complex spinel oxides also contributed in determining the growth stresses and spallation properties of TBC system. The differences between the oxidati on on bare substrates and under a TBC will discussed in Chapter 9. Higher growth rates of the TGO lead to the buildup of growth stresses resulting in cracking and spallation failure. For this reason, the TGO was long considered to the weaker link in achieving higher lifetimes of TBC systems. The addition of elements to contain the growth rates of the TGO would be highly beneficial in improving the adherence and durability of TBC layer. The isothermal spallation lifetimes we re improved by the addition of one or a combination of RE elements to the base alloy. One of the primary aspects of RE additions discussed in the previous chapter was its ability to reduce the growth rate of the oxide layer. One of the questions

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144 in this study was the role of RE additions on the TGO layer growth that can be inferred from its thickness. For this purpose, TBC coated CM247LC baseline and CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy which were subjected to isothermal oxidation at 1010 oC and 1079 oC for 3720 hours were selected. These two alloys were picked on the basis of their spallation lifetime with the baseline showing the lowest and the selected RE modified alloy showing the highest lifetime after all the RE modified alloys. The results of the thickness measurements are shown in Figure 613. As seen in the Figure, an average reduction in the thickness by ~ 21 % and 12% at 1010 oC and 1079 oC respectively were observed. Due to the earlier spallation of alloys exposed to 1121 o6.3.2.3 MicroProbe Analysis on TGO BC thickness C, the thicknes s analysis was not conducted. From this analysis, t he higher lifetime of RE modified alloys ca n be correlated to the reduced TGO growth rate. To analyze the variations in the composition of Ni, Co, Cr, Al, Y, Hf and Re across the thickness of the TGO and BC layer, line scan compositional analysis was conducted using Electron Probe Micro Analysis (EPMA). On the TGO and in the regions closer to the TGO/BC interface, a step size of 2 m and in the locations far ther from the interface, a step size of 5 m was used. The results are for the compositional analysis for alloys exposed to 1010 oC and 1079 oC are shown in Figures 614 and 615, respectively. At 1010 oC, the results showed that the regions closer to the YSZ/TGO interface were richer in Ni, Co and Cr, indicative of the spinel oxide phases as observed during SEM studies. As seen from the scan results, the TGO was primarily composed of Al with only minor levels of Ni, Co and Cr. In the BC regions, the values were more representative of its actual composition. Similar results were observed on alloys oxidized at 1079 oC also. Comparing the line scan profiles for both the temperatures, the length of Al enrichment was higher at 1010 oC than at 1079 oC. This lengt h is not to be correlated with the overall thickness of the TGO layer. During the characterization process, the optimal location

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145 to perform the line scan on 1079 oC alloy had a relatively thin Al2O3 scale. Minor levels of Re and Y were also detected as a part of the BC composition, whereas Hf was detected at lower levels through the length of the scanning. Due to the presence of the RE elements in the ppm levels, they were not detected during the microprobe analysis. Reliable line scan profiles could not be obtained from the alloys exposed to 1121 o6.4 Segregation Analysis on the TGO Layer C due to the extensive cracking in the regions close to the YSZ/TGO interface. The previous section addressed the structural details in the TBC system with respect to the microstructural changes at the TGO/BC interface and near the BC/Superalloy interface. Interestingly, even in their absence as a part of the MCrAlY chemistry, phases consisting of Hf crostructure of MCrAlY. Further more, Hf was found as an oxide encapsulated within the TGO layer, about 150 m ahead of the source, the superalloy substrate. The next question to be answered is what would be the effect of this long distance diffusion of Hf and possibly the RE elements on the microstructure of Al2O3. In this section, attempts were made to address this question using the results obtained Al2O36.4.1 SIMS Analysis on the TGO layer grain boundaries. Due to the high sensitivity of SIMS, it was employed to quantitatively identify the presence of minor RE additions within the TGO and at the TGO/BC interface. For this reason, CM247LC+200 ppm Ce+80 ppm Pr alloy exposed to isothermal oxidation at 1079 oC for 3720 hours was analyzed. To avoid the difficulties associated with sputtering a ceramic, YSZ layer was polished down to leave only a thin layer of TGO on the BC and Ni base superalloy substrate. The sample surface was sputtered using a 6 keV o xygen primary ion beam and the qualitative results were obtained as shown in the Figure 6 16. A high intensity of Al was

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146 detected along with the constituent elements of the BC, but the presence of RE elements were not observed using SIMS. The presence of only a total of 280 ppm RE could have limited the possibility of their detection during SIMS survey. The encouraging observation was the presence of Y in the TGO layer which might have diffused from the BC into the TGO during the high temperature oxi dation. Alternatively, a 5 keV cesium primary ion beam was also employed but produced little change in the outcome of the qualitative analysis. 6.4.2 TEM Analysis of the TGO Layer Focused Ion Beam (FIB) milling technique was used to produce a 80 nm thick TEM foi l from a CM247LC+300 ppm Pr alloy exposed to oxidation at 1079 oC for 3720 hours. The grain Al2O3, YSZ/TGO and TGO/BC interface regions were analyzed in the TEM. The composition of various elements within the grains, at the grain boundaries and at the interfaces was also obtained using EDAX facility equipped in the TE M. A combination of three TEM micrographs that were taken at different locations along the length of the TGO was collated into one single TEM micrograph as shown in Figure 617. A mixed zone composed of the oxides of various constituent elements of the BC was also seen close to the YSZ/TGO interface. These mixed type oxides were also observed during the SEM analysis on thermally grown oxide layers. Al2O3The grain boundaries of the TGO layer intersecting the BC were analyzed using EDAX for the segregation of RE, Y and Hf. An oxide phase rich in Hf, Al and Y with the approximate composition of 64.09 % Hf, 9.82 % Y, 7.50 % Al and 18.60 % O (all given in % wt) w as noticed at one of the grain boundaries intersecting the BC. Refer to Figure 6 18 for the TEM micrograph along with the EDAX spectrum for the above mentioned phase. Using a probe size of 3 nm, was finer in nature close to the mixed zone followed by coarser columnar grains through the remaining length of the TGO layer. The columnar grains were not observed to be oriented perpendicular to the YSZ/TGO interfaces.

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147 extensive scanning was conducted at various grain boundaries to detect the presence of RE elements. Hf was observed to be present but RE element presence was not detected. 6.5 Summary The observations of the oxidation behavior of RE modified alloys coated with TBC are summarized as follows; The spallation lifetimes of RE modified CM247LC alloys were significantly better than the baseline alloys. Also, the alloys with multiple RE elements exhibited better lifetimes than alloys with single RE additions. The spallation of the YSZ layer was observed to occur in the regi on within the ceramic near to the YSZ/TGO interface. SIMS and TEM EDAX analysis did not reveal the RE elements at the grain bounda ries and TGO/BC interfaces. On the other hand, presence of an oxide phase rich in Hf, Y and Al was found near one of the TGO grain boundaries intersecting the BC layer.

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148 200 m 50 m A B C TBC BC / Figure 6 1. Cross sectional SEM micrographs showing (A) the va rious layers of a TBC system, (B ) AirPlasma Spray deposited Yttria Stabilized Zirconia (YSZ) layer with porosity and cracks running in different directions and (C) High Velocity Oxy Fuel sprayed MCrAlY bond coat with a splat morphology. The defective nature of the YSZ coating is important for deflecting the heat waves to impart maximum thermal protection. A coating heat treatment would sinter the splat like morphology of the bond coat. Figure 6 2. A digital image showing the CM247LC substrate coated with TBC system during the various stages of oxidation exposure. The failure criterion is the complete de cohesion of the YSZ ceramic top coat.

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149 0 1000 2000 3000 4000 CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr Spallation lifetime (hrs) 1010 C 1079 C 1121 C 0 1000 2000 3000 4000 CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr Spallation lifetime (hrs) 1010 C 1079 C 1121 C Figure 6 3. Spallation lifetime of TBC system on various RE modified CM247LC alloys compared with the baseline all oy. Alloys exposed to oxidation at 1010 o 0 1000 2000 3000 4000 CM247LC Ce+Pr Ce+Dy Ce+Pr+Dy Ce+Pr+Gd Pr+Dy+Gd Spallation lifetime (hrs) 1010 C 1079 C 1121 C 0 1000 2000 3000 4000 CM247LC Ce+Pr Ce+Dy Ce+Pr+Dy Ce+Pr+Gd Pr+Dy+Gd Spallation lifetime (hrs) 1010 C 1079 C 1121 C C did not suffer spallation even after an extended span of 10, 000 hours, as indicated the arrow marks. Overall, CM247LC+300 ppm Pr alloy exhibited better spallation behavior. Figure 6 4. Spallation lifetime of TB C system on CM247LC alloys modified with more than one RE additions. CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd showed better lifetimes. Refer to the text for the actual ppm level compositions of RE additions.

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150 100 m BC / BC / I II A B Figure 6 5. Cross sectional SEM micrographs of t he CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd coated with a TBC system showing, (A) as deposited structure and (B) microstructure after exposure at 1010 o 5 m 20 m A B C HfC 5 m C for 3720 hours. Two layers are visible in the BC and are explained in the following figure. Figure 6 6. SEM micrographs of CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd exposed to oxidation at 1010 oC for 3720 hours, showing (A) two layered structure of bond coat. The magnified micrographs of rectangular box with continuous lines (B) and dotted lines (C), showing shaped phase present in inter diffusion zone.

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151 1 m A B Figure 6 7. SEM micrographs showing the layer I/II interface of the bond coat in CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy exposed to oxidation for 3720 hours at (A) 1010 oC and (B) 1079 oC. The dotted lines tentatively demarcate the layers I and II. The oxide inclusions are darker in contrast. The dark arrows indicate the HfO2 phases while the white arrows indicate the oxide phase richer in Hf Ce Pr Gd. 100 m A B Figure 6 8. SEM micrographs showing the YSZ/BC region in CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd. (A) As deposited structure and (B) Presence of a thermally grown oxide layer between YSZ and BC as indicated by the arrow, after exposure for 3720 hours at 1010 oC.

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152 20 m A B C 5 m A B Figure 6 9. SEM micrographs showing the plan view of the thermally grown oxide layer (TGO) between YSZ and BC. (A) Back scattered electron image showing the YSZ, spinel Al2O3 layer represented respectively by A, B and C (B) A high Al2O3 5 m 10 m A B C A B Figure 6 10. SEM micrographs showing the cross sectional view of the thermally grown oxide (TGO) layer between YSZ and BC. (A) Secondary electron mode image and higher magnific ation view of the rectangular box shown in the backscattered electron mode Al2O3 while the regions closer to the YSZ are richer in oxides of Ni, Co, Cr and Al.

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153 5 m A B Figure 6 11. SEM micrographs showing the presence of HfO2 internal oxidation stringers Al2O3 as indicated by the arrows in (A) CM247LC baseline alloy subjected to oxidation at 1010 oC for 3720 hours and (B) CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy oxidized a t 1079 o 20 30 40 50 60 70 2 Theta Intensity (A.U) 1010 C 1079 C 1121 C NiO Al2O3 NiCr2O4 NiAl2O4 NiAl10O16ZrO2 CoO Y2O3 C for 3720 hours. Figure 6 12. X ray diffraction patterns of CM247LC+860 ppm Ce alloy exposed to oxidation at 1010 oC / 3720 hours, 1079 oC / 4392 hours and 1121 oC / 1440 hours. The peaks of ZrO2 and Y2O3 originated from the isolated colonies o f unspalled YSZ while spinel oxides corresponding to MCrAlY bond coat were also noticed.

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154 Figure 6 13. Average thickness of the TGO layer in CM247LC baseline and CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloys. RE modified alloy showed a lower thickness value and the higher temperature lead to thickening of the TGO.

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155 0 10 20 30 40 50 60 70 80 90 100 0 20 40 60 80 100 120 Wt % Al Ni Co Cr TGO Figure 6 14. Line scan results of microprobe analysis conducted on CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy exposed to oxidation at 1010 o 0 10 20 30 40 50 60 70 80 90 100 0 20 40 60 80 100 120 Wt % Al Ni Co Cr TGO C for 3720 hours. The hatched region represen ts the TGO layer composition. Figure 6 15. Line scan results of microprobe analysis conducted on CM247LC+200 pm Ce+60 ppm Pr alloy exposed to oxidation at 1079 oC for 3720 hours.

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156 0 1 23 4 5 6 0 15 3045 60 75 90 105 120 135 150 Mass (AMU) Log Cps O 16 Al 27 Cr 52 Ni 58 Co 59 Y 89 Zr 90 Figure 6 16. Secondary Ion Mass Spectroscopy (SIMS) results obtained fro m CM247LC+260 ppm Ce+80 ppm Pr alloy exposed to 1010 o 1 m Pt deposit Bond Coat YSZ Coat Mixed Zone 1 m Pt deposit Bond Coat YSZ Coat Mixed Zone C for 3720 hours. The survey was conducted on the YSZ/TGO interface to obtain the qualitative information about the composition. RE elements were not detected at the interface. Figure 6 17. A collation of cross sectional TEM micrographs taken along the length of the TGO layer showing finer grains at the YSZ/TGO interface followed by longer columnar grains towards the TGO/BC interface. A mixed zone consisting of oxides of the constituent elements in BC were also observed closer to the YSZ/TGO interface.

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157 0.5 m 0.2 m BC Al2O3 A B D C 0.5 m 0.2 m BC Al2O3 A B D C Figure 6 18. Cross sectional TEM micrographs showing (A) the presence of an oxide phase rich in Hf Y Al in a TGO grain boundary intersecting the BC layer, (B) the corresponding EDAX spectrum of the phas e, (C) higher magnification image of the phase in bright field and (D) dark field.

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158 CHAPTER 7 TENSILE BEHAVIOR 7.1 Introduction The previous two chapters dealt in greater detail with the significant improvements in the high temperature oxidation resistanc e imparted by the minor additions of RE elements in both the uncoated and coated alloys. This chapter will try to answer the question of what would be the effect of minor RE alloying additions on the deformation behavior of superalloys under the influence of an applied tensile stress? Mechanical deformation behavior, unlike oxidation, is influenced not only by the alloy constituents but also by the microstructure. RE additions, despite their strong influence in improving the oxidation resistance, were disre garded as a potential alloying addition for increasing mechanical strength due to the formation of low melting eutectics with transition group elements [12] To further weaken the case for RE additions, the elements of interest in this research have very low solubility in Ni and tend to form brittle NixRE additions are becoming increasingly popular in the commercial alloys due to their well documented benefits in improving the oxidation properties [8, 9, 157, 185] Due to the very limited amount of research conducted to understand the role of these elements on mechanical properties and deformation [136, 137,.142, 186189] the RE effe ct on deformation behavior is still not fully characterized. In this context, the objectives of the presented work here are: ( RE ) based intermetallics, which have a detrimental impact on the mechanical properties [137, 142] Even in this research, RE Hf rich phases were observed in the solution heat treated structures as discussed in Chapter 4 To establish the effect of RE elements on the tensile behavior of Ni base superalloys as a function of temperature, heat treatment and various ppm levels of the RE additions. To understand the role of RE additions on the fracture mechanisms of the alloys at various temperatures.

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159 Alloys from Phase I and II were subjected to tensile testing in the temperature range of 650 C to 850 C. An additional test at room temperature was also conducted to study the deformation without the influence of temperature. In Phase I, the alloys tested were exposed to solution heat treatment conditions using Industrial Heat Treatment (IHT) and UF Heat Trea tment (UFHT) cycles. Comparison between unHIPed and HIPed alloys is also discussed. In Phase II, the downselected alloys were given a Mod SHT 03 heat treatment. The major interests in this chapter are yielding behavior and the fracture modes for the allo y as a function of rare earth additions and temperature. The main objectives were addressed using the 0.2% yield strength, Ultimate Tensile Strength (UTS), % reduction in cross sectional area, % elongation together with detailed fractographic analysis usin g SEM. 7.2 Phase I Un HIPed Alloys 7.2.1 Industrial HT As discussed in the Chapters 5 and 6, alloys were downselected for further analysis from the initial batch of 16 alloys based on the adhesion behavior of TBC coated alloys as a function of RE additi ons. Only those 6 downselected alloys were subjected to a Industrial solution heat treatment with a maximum temperature of 1232 oC followed by the conventional double aging at 760 C and the tensile flow curves are compiled in Figure 7 1. The testing was carried out at a cross head velocity of 0.1 in/min. As seen from the figure, CM247LC+300 ppm alloy exhibited the maximum tensile strength while CM247LC+60 ppm Ce+50 ppm Dy+50 ppm Gd showed higher tensile ductility. With the exception of these two alloys, all the other alloys exhibited little or no plastic behavior. There was no significant difference in the modulus of elasticity of the alloys tested and all of the samples failed immediately after reaching the maximum stress levels.

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160 All the alloys exhibited very negligible strains to fa ilure and due to this reason, ductility was not taken into consideration. 7.2.2 UFHT Alloys containing only one RE addition were subjected to a multi step UF solution heat treatment with a maximum temperature of 1260 oC. Similar to the previous case, the testing on these alloys was conducted at 760 o7.2.3 Fractography C and the tensile stress vs. strain graphs are plotted in Figure 2. CM247LC baseline alloy exhibited the highest modulus of elasticity but fractured only after a very short plastic region. Only CM247LC+160 ppm Pr alloy showed a visible strain hardening behavior and possessed the maximum strength among all the alloys tested. On the other hand, other RE containing alloys exhibited premature failure with very low modulus of elasticity values. All the alloys tested from both the heat treatments s howed very low ductility values, with no evidence of necking, which was apparent from their failure at the maximum stress. Limited analysis on the fracture surfaces using SEM was conducted on the CM247L C+250 ppm Ce+70 ppm Dy and CM247LC+300 ppm Pr alloys from SHT to understand why all alloys exhibited such poor tensile behavior. Observation of the fracture surface gave no indication of any specific fracture mechanism. A clear distinction of any inter gra nular or trans granular fracture was not readily apparent. As seen in the Figure 73A, the fracture surface of the alloys were populated with a large volume fraction of porosities. The spheroidal nature of the pores which were 3 5 m in diameter was repres entative of solidification porosity. The presence of larger sized pores reduced the effective load bearing cross section of the samples at the gage area, which was responsible for the pre mature failure of the alloys. On the other hand, the 300 ppm Pr modi fied alloy contained relatively fewer pores and exhibited an inter dendritic mode of fracture. On closer observation, several features like tear ridges and micro voids were observed

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161 which were indicative of ductile mode of fracture. This explained why this alloy exhibited better tensile behavior compared to rest of the alloys subjected to testing. 7.3 Phase I HIPed Alloys 7.3.1 UFHT The tensile behavior of unHIPed alloys was observed to be strongly influenced by the presence of large sized pores. To eli minate the effect of porosities on the tensile results, the alloys were Hot Isosta tically Pressed (HIPed) at 1185 15 oC with a pressure of 175 MPa 7 MPa for 4 hours. Only the alloys containing single RE additions were subjected to tensile testing in t his phase and unfortunately, due to the shortage of alloy bars, the baseline alloy was not included in this study. All these alloy bars were given a multistep UFHT solution heat treatment with a maximum temperature of 1260 oC followed by double aging. A s train rate of 8 x 105 s1 was used at three temperatures namely, room temperature, 650 oC and 760 o The results of tensile testing at room temperature are compiled in Figure 7 4. Reminiscent of unHIPped alloy tensile results, CM247LC+860 ppm Ce alloy showed the lowest modulus of elasticity and failed without any sign of plastic deformation. But, the other three RE modified alloys showed a very pronounced plasticity, deforming under tension in a manner similar to each other. After the start of yielding, these alloys showed only moderate strain hardening, with their yielding relatively flat over a longer range of straining. All the alloys failed at the maximum stress level. CM247LC+160 ppm Pr and CM247LC+180 ppm Dy alloys s howed the highest strains to fracture with the former being the strongest alloy with an UTS value 10 20 % greater than the other alloys tested at room temperature. C. The objective of testing at various temperatures was to gain an understanding of the flow character and fracture mechanism as a function of test temperat ure. The following is the discussion of the results obtained from this testing and the fractography results on a few selected samples.

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162 Figure 7 5 shows the tensile testing results of the alloys at 650 oC. The CM247LC+860 ppm Ce alloy was stronger in the elastic region but failed immediately after the onset of yielding. The CM247LC+120 ppm Gd alloy exhibited a comparatively better strain hardening character with a maximum strength among the alloys tested at 650 oThe tensile behavior of the alloys at 760 C. Pr and Dy modi fied alloys deformed in a similar way with the 160 ppm Pr containing alloy showing the highest strain to fracture. oSummarizing the results of HIPed alloys: C is shown in Figure 76. The CM247LC+860 ppm did not show any change in the deformation characte r compared to the previous two reports. CM247LC+160 ppm Pr and CM247LC+120 ppm Gd modified alloys were weaker in the elastic region but showed a moderate strain hardening and higher strain to fracture. CM247LC+180 ppm Dy alloy proved to be the strongest wi th a UTS value nearly 40% greater than Ce and 10 % greater than Pr and Dy modified alloys. Contrary to the previous observations of the alloys failing at the maximum stress, the Dy containing alloy failed at a stress about 60 MPa lesser than the maximum stress, indicating some necking had occurred. The CM247LC+860 ppm Ce alloys failed in the elastic region without showing any signs of plastic deformation. This behavior was very similar to results obtained from the unHIPed samples and thus, it raised questions about the HIPing response of this type of alloy. The other three RE modified versions showed plastic behavior and significant tensile ductility. With the exception of CM247LC+180 ppm Dy alloy at 760 oC, all the other alloys at three test temperatures were observed to fail at the maximum stress levels. The UTS values of the alloys at three test temperatures are shown in Figure 7 7. While CM247LC+860 ppm Ce showed a very strong dependence of temperature on UTS values, CM247LC+180 ppm Dy alloy showing almost similar UTS values at all the three test temperatures. The temperature dependence of modulus of elasticity could not be established with only the CM247LC+860 ppm alloy showed the normal behavior of decrease in modulus with increase in temperature of testing.

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163 With variations only in the strain to failure and UTS, Pr and Gd modified alloys exhibited similar flow behavior under all the conditions. 7.3.2 Fractography It was important to understand the reason behind the poor performance of Ce modified alloy. So, the fractographic analysis of CM247LC+860 ppm Ce alloy was conducted using SEM and compared to the fracture behavior of a CM247LC+180 ppm Dy sample which was observed to exhibit better tensile properties. The sa lient features of the fracture surfaces at various temperatures are discussed below and correlated with their tensile deformation characteristics. Figure 7 8 shows the fractography of the two alloys fractured at room temperature. CM247LC+860 ppm Ce alloy s howed an inter dendritic fracture pattern but did not show any sign of porosity. The inter dendritic pattern was similar to a conventional cast structure with finer structure near the periphery followed by columnar dendritic structure radiating in towards the center of the alloy. The separation having taken place completely along the inter dendritic regions, the cleavage planes were not readily apparent. On closer observation, nodules were seen protruding out from a relatively smoother dendrite surface as m arked by arrows in Figure 7 8 B. In one of the earlier reports, this type of appearance was attributed to the possible presence of a thin liquid film remaining at the inter dendritic regions after solidification [156] CM247LC+180 ppm Dy alloy showed a dominant trans granular fracture with tear ridges representative of a ductile mode of failure. As seen in t he Figure 7 8 C, these equiaxed tear ridges were possibly aligned in the direction of the grain orientation of the alloy. The fractographic details of alloys exposed to tensile testing at 650 oC are given in Figure 79. CM247LC+860 ppm Ce alloy showed pri marily but not completely an inter dendritic fracture mode. As indicated by arrows in Figure 7 9 A, secondary cracks were observed along the inter de ndritic regions, which explained the lack of strength and ductility seen in this type of

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164 alloy. Mixed mode of fracture with tear ridges characteristic of trans granular mode and grain boundary separation typical of inter granular mode were observed for CM247LC+180 ppm Dy alloy. Secondary cracking was also observed along the grain boundaries with microporositie s populated on the grain facets as seen from Figure 79 B. The fracture surface o f CM247LC+860 ppm Ce alloy that failed at 760 oSummarily, with increase in temperature; C was similar to the previously observed i nter dendritic fracture pattern, with the presence of secondary cracking along the inter dendritic regions. Refer Figure 710A. The dominant failure mode in CM247LC+180 ppm Dy alloy was through microvoid coalescence as seen in Figure 7 10 B. The micro voids were nucleated at the interface between carbides and the matrix, growing in size under the influence of an externally applied tensile stress and finally coalescing to result in fracture. The higher strength and ductility observed in this type of alloy could be correlated to the fracture mode. Other observable features of minor significa nce were the presence of transgranular cleavage facets and tear ridges. The Ce modified alloy did not exhibit major changes in the fracture mode with cracking appearing only in the inter dendritic region at higher temperature. A change in the dominant mode of fracture from transgranular to a mixed mode (trans and intergranular) to primarily failure by micro void coalescence was observed in the Dy modified alloys. 7.4 Comparison between HIPed and UnHIPed Alloys Ba sed on the tensile behavior discussed in the previous two sections, it was possible to compare the deformation characteristics of un HIPed and HIPed alloys subjected to testing at 760 oC. Similar to the previous results, the CM247LC+860 ppm Ce alloy did not show changes in the fracture feature, whereas major changes were seen in CM247LC+180 ppm Dy and CM247LC+120 ppm Gd alloys. Compared to the total absence of plastic region in the un-

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165 HIPped conditions, both these alloys showed a significant yielding behavi or. Unlike the other alloys, the CM247LC+160 ppm Pr alloy exhibited plastic deformation in the unHIPed condition and the characteristics did not change much in the HIPed condition. The differences in the UTS values could be inferred from Figure 7 11. Dy a nd Gd containing alloys achieved significant gains due to HIPing, whereas the Pr containing alloy exhibited similar properties in both the conditions. Regarding the fracture mechanism, porosities did not play a significant role in the deformation behavior of the HIPed alloy and deformation through micro void coalescence was the primary mode of failure. The question of why Ce modified alloys exhibited a poor response to the HIPping process remains to be addressed. 7.5 Phase II Alloys The development of the a lloys in different phases was explained in detail in the Chapter 3, Materials and Experimental Methods. The cyclic oxidation behavior of the Phase I downselected alloys was discussed in Chapter 5. Based on the oxidation behavior, two alloys namely, CM247LC +300 ppm Pr and CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloys were chosen for the Phase II studies. All of these alloys were received in the HIPed condition. The tensile deformation characteristics of these alloys under various conditions of solution heat treatment and testing temperature will be discussed in the following section. 7.5.1 Industrial HT One set of Phase II alloys were given a Siemens solution heat treatment with a maximum temperature of 1232 oC followed by double aging. Testing on these speci mens was carried out at Joliet Metallurgical Laboratories, Joliet, IL at a strain rate of 8 x 105 s1 and at three temperatures namely, room temperature, 760 oC and 850 oC. Due to the nature of the testing equipment used, continuous monitoring of the test ing data was not feasible and only important parameters like 0.2 % yield strength and UTS values were recorded as shown in Figures 712

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166 and 713 respectively. From the yield strength values of the three alloys under consideration, the CM247LC baseline allo y proved to be slightly stronger than the two RE modified alloys at room temperature and 760 oC. On the other hand, all the three alloys were observed to exhibit similar yield strength levels at 850 oConsidering UTS values, the CM247LC baseline alloy proved to be stronger than the RE modified alloys at both the room temperature and 760 C. Also, all the alloys followed the trend of yielding a t lower at lower stresses with an increase in temperature. Considering the differences in the yield strength and UTS values of the alloys at three test temperatures, all of the alloys exhibited an approximately 20 % increase in the stress level indicating a good strain hardening behavior. oC testing. Similar to the yield strength data, all of the three alloys showed similar yield stress values at a higher temperature of 850 oC. The CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloys showed a higher UTS value at 760 oC than at room temperature and was considered to be an anomaly in the testing. There was only a minor difference in the UTS values between room temperature and 760 oC, whereas the UTS value observed at 850 oC decreased by 14% compared to 760 o7.5.2 UF HT (Mod SHT 03) C. A nother set of alloy bars were subjected to a modified solution heat treatment cycle, ModSHT03, t he development of which was explained in greater detail in the Chapter 4 This solution cycle was comparable to previous heat treatment employed on other samples but was more effective in reducing residual eutectics with negligible incipient melting. The machined samples were subjected to tensile testing at the strain rate of 8 x 105 s1 at four different temperatures namely, room temperature, 650 oC, 760 oC and 850 oC. The inclusion of 650 oC testing provided an additional range in analyzing the tensile b ehavior and fracture mechanisms as a function of temperature.

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167 The tensile deformation behavior at room temperature for the three alloys tested is shown in Figure 7 14. The CM247LC baseline alloy exhibited a lower modulus of elasticity but showed higher te nsile strength and higher tensile ductility compared to the RE modified alloys. Similar to the previously discussed Phase I tensile results, strain hardening was limited with all of the alloys showing a flow stress plateau and failing immediately after rea ching the maximum tensile strength. Comparing the UTS values, the baseline proved to be about 7% stronger than the RE modified alloys. Figure 7 15 shows the tensile results from the testing conducted at 650 oC. With all the alloys showing similar modulus o f elasticity values, the CM247LC+300 ppm Pr alloy exhibited a very pronounced strain hardening character until failure at the maximum tensile strength. Pr modified alloys also showed a higher tensile ductility compared to the other two alloys, also recordi ng an increase in the UTS value by 20 %. Due to the work hardening, the behavior of Pr modified alloy in this case was similar to the CM247LC baseline alloy while testing at 760 oAt 850 C as shown in Figure 716. Despite showing a good strain hardening character, RE modified alloys suffered only a limited tensile strain to failure. oSummarizing the above behavior of the alloys at various temperatures; C testing, the baseline alloy was stronger in the elastic region and showed an extensive flow stress plateau of tensile straining with negligible strain hardening as shown in th e Figure 7 17. A pronounced necking behavior was observed for the CM247LC baseline alloy as also indicated by the deformation behavior. Interestingly, RE modified alloys failed in the elastic region with no signs of plastic deformation. This behavior was s imilar to the tensile behavior of unHIPed alloys tested in Phase I.

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168 All the alloys showed a greater amount of strain hardening at 650 oC and 760 oC but showing only a flow stress plate au of extensive tensile straining at room temperature and 850 o RE modified alloys exhibited a decrease in the yield strength and modulus of elasticity as function of increasing temperature. But, CM247LC did not show any appreciable decr ease in the elastic modulus but followed the general trend of decreasing yield strength at higher temperatures. C (only baseline). With the exceptions of the CM247LC baseline showing higher strength at 760 oC and the CM247LC+300 ppm Pr at 650 o There was a general increase in the strain to fracture with increase in temperature, except the case of RE modified alloys at 850 C, the relation of UTS and temperature was in line with the expected trend. The UTS values at different temperatures are graphed in Figure 7 18. o7.5.3 F ractography C. Analysis of the fracture surfaces using SEM was conducted for CM247LC baseline and the CM247LC+300 ppm Pr alloys subjected to tensile testing at four test temperatures. At room temperature, the failure initiated by transgranular cleavage with f lat facets showing river mark like patterns indicating the direction of crack propagation. Tear ridges were also predominantly noticed in the alloy explaining the ductility observed during testing. On the other hand, the CM247LC+300 ppm Pr alloy failed thr ough a mixed mode of inter dendritic and transgranular mechanism with the latter appearing to the dominant one. Nodules, similar to the ones observed in the Phase I Ce modified alloys, were noticed on the fracture surface. The fracture behavior of the CM2 47LC baseline alloy at 650 oC was very similar to the room temperature case with a transgranular cleavage facet which initiated the failure and the alloy failed by tearing due to overload. Refer Figure 7 19 A for the respective fractograph. The CM247LC+300 ppm Pr alloy exhibited a mixed mode of transgranular cleavage and a significant amount of microvoid coalescence. Similar to the previously mentioned cases, the voids nucleated around carbide particles as shown in Figure 719 B. This explains the increase d strength and ductile character of Pr modified alloy at 650 oC testing.

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169 At 760 oUpon initiation by transgranular cleavage, the CM247LC baseline alloy failed pr edominantly through the microvoid coalescence as shown in Figure 7 21 A. Explaining the pre mature failure of the CM247LC+300 ppm Pr alloy, the alloy fractured by total inter granular separation as shown in the Figure 721 B. On closer observation, deep s econdary cracks were visible along the grain boundaries as shown in Figure 721 C. One of the interesting observations was that th at grain facets were not smooth, which indicated the occurrence of sliding along grain boundaries prior to graingrain separat ion (Refer to Figures 7 20 and 721). C, the CM247LC baseline alloy exhibited a very complicated mode of fracture with several features similar a general transgranular appearance, secondary cracking along the gra in boundaries and the presence of micro voids. On closer observation, microvoids were present uniformly on all the transgranular fracture surfaces as shown in Figure 720 A. Except the relative amounts, the appearance was very much similar in the CM247LC+300 ppm Pr alloy which also showed lesser incidence of microvoids. Secondary cracking was not observed on the inter granular fracture surfaces as shown in the Figure 7 20 B. The CM247LC+300 ppm Pr alloy exhibited several features, basically failing through a pred ominant transgranular mode that transformed into total inter granular fracture at the highest test temperature. This observation points to a definitive role of Pr in the grain boundary deformation mechanisms that depended strongly on temperature. 7.6 Summary The observations of the tensile deformation characteristics and fractographic analysis can be generally summarized as follow s; Minor additions of RE did not produce the expected results of alloy strengthening. The CM247LC+860 ppm Ce alloy failed without showing any signs of plastic deformation under both unHIPed and HIPed conditions. Fractographic analysis showed the presence of nodules protruding from the dendritic surface, which raised questions about the quality of

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170 the specific alloy bar casting. But, similar observation of nodules in the CM247L C+300 ppm Pr alloys in Phase II, which exhibited good tensile ductility contradi cts the failure hypothesis for Ce modified alloys. So, the specific effect of 860 ppm Ce additions needs to be addressed. Most of the alloys showed a mixed mode of fracture and the samples with a larger fraction of dimples on the fracture surface exhibiti ng better tensile ductility. Complete inter granular cleavage of the CM247LC+300 ppm Pr alloy at 850 o C points to the increased influence of RE on the grain boundary deformation at higher test temperatures.

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171 Figure 7 1. Tensile stress strain curves of unHIPped Siemens solution heat treated Phase I alloys tested at 760 oC. With the exception of CM247LC+300 ppm Pr and CM247LC+60 ppm Pr+50 ppm Dy+50 ppm Gd, all the other alloys showed little signs of plastic deformation. These alloys were given a Sie mens solution heat treatment followed by double aging.

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172 Figure 7 2. Tensile stress strain curves of unHIPped UF solution heat treated Phase I alloys subjected to testing at 760 o 10 m A B C. Only the CM247LC+160 ppm Pr alloy was observed to exhibit a better plastic deformation behavior. Figure 7 3. SEM fractog raphs taken from Phase I Industrial solution heat treated alloys, (A) CM247LC+250 ppm Ce+70 ppm Pr showing large sized solidification porosities that reduced the effective load bearing cross section leadin g to pre mature failure during testing (B) CM 247LC+300 ppm Pr showing the presence of micro voids which indic a ted the ductile mode of fracture.

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173 Figure 7 4. Tensile stress strain curves of HIPped Phase I UFHT alloys subjected to testing at room temp erature. While CM247LC+860 ppm Ce alloy failed pre maturely in the elastic region, other RE modified alloys showed a moderate strain hardening response followed by failure at the maximum stress.

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174 Figure 7 5. Tensile stress strain curves of HIPp ed UFHT Phase I alloys subjected to testing at 650 o C. While CM247LC+120 ppm Gd alloy exhibited higher tensile strength, CM247LC+180 ppm Dy alloy showed a better tensile ductility.

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175 Figure 7 6. Tensile stress strain curves of HIPped UFHT Phase I alloys s ubjected to testing at 760 o 500 600 700 800 900 1000 1100 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd UTS (MPa) 760 C 650 C RT C. CM247LC+180 ppm Dy alloy exhibited a minor necking behavior during testing. Figure 7 7. Ultimate strength values (UTS) of all the RE modified alloys tested at RT, 650 oC and 760 oC. Generally, a reduction in UTS with increa se in testing temperature was observed.

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176 1 mm 100 m 50 m A B C 1 mm 100 m 50 m A B C Figure 7 8. SEM fractographs of HIPped UFHT Phase I alloys tested at RT showing (A) an inter dendritic fracture mode in CM247LC+860 ppm Ce alloy, (B) the same surface at higher magnification showing nodules protruding from the inter dendritic regions and (C) fracture surface of CM247LC+180 ppm Dy alloy showing tear ridges which were equiaxed and oriented according to the alloy grain structure. 200 m A B Figure 7 9. SEM fractographs of HIPped UFHT Phase I alloys tested at 650 oC showing (A) inter dendritic cracking in CM247LC+860 ppm Ce alloy and (B) secondary cracking along the grain boundaries. Also noticed on the grain facets is the presence of microporosities.

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177 50 m 200 m A B Figure 7 10. SEM fractographs of HIPped UFHT Phase I alloys tested at 760 o 200 400 600 800 1000 Ce Pr Dy Gd UTS (MPa) unHIP HIP C showing (A) inter dendritic cracking in CM247LC+860 ppm Ce and (B) micro void coalescence in CM247LC+180 ppm Dy which explained its higher ductility. Figure 7 11. Comparison of Ultimate stren gth values (UTS) between unHIPed and HIP ed R E modified alloys. Except CM247LC+160 ppm Pr alloy, other RE modified alloys showed a considerable increase in the UTS due to HIP ing.

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178 0 200 400 600 800 1000 Baseline Ce+Pr+Gd Pr 0.2 % Yield Strength, MPa RT 760 C 850 C Figure 7 12. Yield strength values of Phase II Industrial solution heat treated alloys subjected to tensile testing at R T, 760 oC and 850 o 0 200 400 600 800 1000 1200 Baseline Ce+Pr+Gd Pr Ultimate Tensile Strength, MPa RT 760 C 850 C C. A general trend of decrease in yield strength with increase in temperature was noticed. Figure 7 13. Ultimate strength values (UTS) of Phase II Siemens solution heat treated alloys subjected to tensile testing at RT, 760 oC and 850 oC. With the exception of few anomalies, a decrease in the maximum tensile strength with increase in temperature was observed.

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179 Figure 7 14. Tensile stress strain curves of UFHT Phase II alloys subjected to testing at RT. The strain hardening response was very moderate with all the alloys suffering failure at the maximum tensile strength. Figure 7 15. Tensile stress strain curves of UFT Phase II alloys subjected to testing at 650 oC. CM247LC+300 ppm alloy exhibited a very pronounced strain hardening response.

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180 Figure 7 16. Tensile stress strain curves of UFHT Phase II alloys subjected to testing at 760 o C. CM247LC baseline alloy proved to be stronger than the RE modified alloys. Figure 7 17. Tensile stress strain curves of UFHT Phase II alloys subjected to testing at 850 oC. CM247LC baseline alloy showed a pronounced necking behavior while RE modified alloys suffered pre mature failure in the elastic region.

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181 Figure 7 18. Ultimate strength values (UTS) of UFHT Phase II alloys subjected to tensile testing at RT, 650 oC, 760 oC and 850 o 50 m A B C. With a few exceptions, a general trend of decreasing UTS values with increasing temperature was observed. Figure 7 19. SEM fractographs taken on UFHT Phase II alloys tested at 650 oC showing (A) trans granular cleavag e facets in CM247LC baseline and (B) micro void coalescence in CM247LC+300 ppm Pr that explains the higher ductility observed during testing.

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182 20 m 50 m A B Figure 7 20. SEM fractographs taken on UFHT Phase II alloys tested at 760 o 50 m 1 mm A B C C showing (A) micro voids in CM247LC baseline and (B) inter granular mode of failure in CM247LC+300 ppm Pr. Figure 7 21. SEM fractographs taken from UFHT Phase II alloys tested at 850 oC showing (A) fracture through the coalescence of micro voids in CM247LC baseline alloy, (B) low magnifica tion micrograph of CM247LC+300 ppm Pr showing a total inter granular fracture mode and (C) secondary cracking observed along the grain boundaries.

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183 CHAPTER 8 CREEP BEHAVIOR 8.1 Introduction elements have a stronger tendency to segregate to the alloy grain boundaries. This segregation is beneficial in two ways: Since they have the ability to form stable sulphides, RE can getter S, if any, present at the grain boundaries to inhibit their detri mental effect on mechanical behavior (Refer to Chapter 2 for the role of S on the embrittlement behavior during mechanical deformation). The alloys used in this study were relatively free of S and thus, the consumption of RE in the gettering process is red uced and can be considered to be deliberate additions to influence the high temperature properties. Secondly, the presence of RE as free atoms or as a part of RE Hf phase at the grain boundaries can increase the activation en ergy for grain boundary sliding which is an important creep deformation mechanism. When there are 10 alloy constituents in the Ni base superalloys which play a role in a vari ety of strengthening mechanisms, such as precipitation hardening, solid solution strengthening, grain boundary strengthening etc, addressing the role of a ppm level RE additions can be di fficult. Unfortunately, to date, very little work has been published in this area and the RE effect on the creep defo rmation behavior has not been fully addressed. With the object ive of answering the questions about the RE effect, creep tests were conducted at 750 oC, 800 oC, 850 oC and 950 oC with the load levels designed for rupture after 300 hours. Alloys from Phase I and II that were exposed to both the baseline and modified he at treatments were subjected to creep testing. Analysis of the creep deformation behavior was carried out basically through the use of creep curves and fractographic analysis using SEM. Due to the very low ductility values obtained during creep testing, which were typical of

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184 polycrystalline materials, the ductility values were not used as a criterion for the comparison of alloy performance. 8.2 Phase I Alloys In this section, creep behavior of Phase I CM247LC alloys containing only single RE additions were studied in the unHIPed and HIPed condition. Only the results of the alloys subjected to baseline multi step UFHT solution heat treatment are discussed below. 8.2.1 UFHT Un HIPed Alloys The creep deformation of the CM247LC baseline and the RE modified all oys at 850 o8.2.2 UFHT HIPed Alloys C / 240 MPa is illustrated in Figure 8 1. The stress levels were reduced in the Un HIPed condition from 390 MPa in order to reach the target of 300 hour rupture life due to the inherent porosity in the structure that le d to lower lifetimes. The CM247LC baseline and the CM247LC+860 ppm Ce alloys did not e xhibit any primary creep regime, with the latter exhib iting higher minimum creep rate, but the former exhibited a relatively long tertiary creep regime. On the other hand, CM247LC+180 ppm Dy allo y did exhibit primary creep behavior and also showed the highest minimum creep rate. It is interesting to note that the alloy showing higher steady state creep rate exhibited better lifetime. The rupture lifetime of Dy modified alloy was ~ 19 % greater tha n the baseline and about ~ 40 % higher than the CM247LC+860 ppm Ce alloy. The fractographic analysis of the crept alloys showed features very similar to that of the alloys that failed during tensile testing. Refer Figure 7 3. To in hibit the prominent role of casting porosities on the creep deformation characteristics, the alloys were HIPed at 1185 15 oC and a pressure of 175 7 MPa. Unfortunately, due to the shortage of the alloy bars, CM247LC baseline alloy could not be included in this study. The results of the creep study conducted at 850 oC / 390 MPa and 950 oC / 204 MPa are compiled in

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185 Figure 8 2. At 850 oC / 390 MPa, the CM247LC+120 ppm Gd, the CM247LC+160 ppm Pr and the CM247LC+180 ppm Dy alloys showed a higher minimum cre ep rate (in that order). This made the creep rate seemingly a function of the ppm level of the RE addition at this creep testing condition. In the context of rupture lifetimes, CM247LC+120 ppm Gd with a lower creep rate exhibited the highest lifetime among the three alloys. The trend of rupture lifetime reversed at 950 oC / 204 MPa with the CM247LC+180 ppm Dy alloy showing the longest lifetime but both alloys exhibited similar creep behavior. HIPed alloys were observed to show exhibit greater amounts of cre ep ductility compared to the Un HIPed samples. SEM characterization of the fracture surface of CM247LC+180 ppm Dy alloy showed a mixed mode of fracture with microvoid coalescence, intergranular and transgranular cleavage facets at 850 oC. On other hand, f ractographs at 950 o8.3 Phase II Alloys C showed predominantly transgranular characteristics with some intergranular features. The Phase II alloys tested included the CM247LC baseline and two RE modified alloys namely, CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd and CM247LC+300 ppm Pr. All the alloys were HIPped following the casting. In this section, the creep behavior of these alloys that were subjected to a single step Industrial Solution Heat Treatment (IHT) and multistep modified solution heat treatment ( ModSHT 03) are discussed. 8.3.1 IHT Alloys Figure 8 3 shows the creep curves of the three alloys subjected to deformation under constant load at 850 oC / 390 MPa. All the alloys were found to have similar primary creep rates but the secondary creep rate o f the CM247LC baseline was lower than that of the two RE modified alloys. T he baseline alloy underwent an extensive secondary creep with a total strain of 4 % and showed lifetimes higher than 300 hours. RE modified alloys suffered premature failure

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186 in less than 100 hours with creep strains around 1 %. Similar rupture lifetimes were obtained for the creep testing conducted at 950 oTo understand the long term deformation behavior due to RE additions, creep tests were conducted at 850 C / 204 MPa also as graphed in Figure 84. The differences in this case were the primary creep rate of the CM247LC baseline was h igher and so was the steady state creep rate. The CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy exhibited lower creep rates and shorter lifetimes compared to the CM247LC+300 ppm Pr alloy. oC / 280 MPa and 950 oC / 130 MPa with a designed rupture lifetime of 1000 hours. The results are plotted in Figure 85 for CM247LC+300 ppm Pr alloy. Comparing the previous results for the same alloy at 850 oFractographic analysis on the CM247LC baseline alloy subjected to creep at 850 C, lower load levels were observed to decrease the steady state creep rate and at the same time delivered designed rupture lifetimes. oC / 390 MPa showed the presence of extensive dimples c ombined with tear ridges as sh own in Figure 86 A. Steps that resembled fatigue striations, but were found to be slip steps, were also observed on large portions of the fracture surface. Dimpling on the fracture surface was accompanied by the presence of m icro voids coalesced around the carbide particles. These features supported the observation of higher creep ductility in the baseline alloy. On the other hand, analysis on the CM247LC+300 ppm Pr alloy showed a mixed mode of fracture with transgranular tear ridges but the lower creep strain to rupture was supported by the presence of secondary cracking along the inter granular regions and a small fraction of porosity as shown in Figure 8 6 B. The fracture surfaces in the samples tested at 950 oC / 204 MPa di d not exhibit any distinct features and the information obtained was considered inconclusive. Microstructural analysis along the longitudinal sections did not show any appreciable differences between the alloys. The commonly observed features were the pres ence of long network of cracks running a long the

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187 uniform deformation leading to cracking at the interfaces. Refer to Figure 4 9 B for the SEM micrograph illustrating the above mentioned cracking. The presence of cracks was also noticed at distances farther from the fracture sur face. Extensive, wide, grain boundary cracking was observed in both al loys as shown in Figure 8precipitates, was observed in all the samples subjected to creep at 950 o8.3.2 UFHT (Mod SHT 03) Alloys C / 204 MPa. The creep deformation character of the three alloys at 850 oC / 390 MPa is illustrated in Figure 8 8. Similar to the creep behavior exhibited by IHT alloys, both the RE modified alloys suffered premature rupture within about 80 hours of exposure to test conditions. As evident from the figure, RE modified allo ys underwent higher rate of secondary creep before failure, whereas the CM247LC baseline alloy exhibited a higher lifetime of ~ 350 hours. Correlating the steady state creep rate with the amount of RE additions, the CM247LC baseline alloy exhibited the low est rate while the CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd alloy underwent the highest creep rate. Figure 8 9 illustrates the creep behavior of the alloys subjected to deformation at 950 oC / 204 MPa. Both the RE modified alloys showe d increased primary cre ep rates compared to the CM247LC baseline samples, where the primary creep regime was virtually non existe nt. The rupture time for the RE modified alloys decreased significantly with lifetimes below than 50 hours. At both the test temperatures, the CM247LC baseline alloy exhibited higher creep ductility compared to the modified alloys. SEM analysis on the baseline alloy crept at 850 oC / 390 MPa revealed the presence of a large population of microvoids and transgranular cleavage facets, indicative of the h igher ductility observed during the creep testing. See Figure 8 10 A. On the other hand, The CM247LC+300 ppm Pr alloy showed primarily a transgranular

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188 fracture mode with occasional intergranular characteristics and cracking in the grain boundaries as shown in Figure 8 10 B. Information obtained from the fractography of 950 oAlthough the ModSHT03 type of heat treatment did not exhibit any significant differences in the creep behavior, microstructural observations showed that the volume of cracking in the longitudinal section was very low c ompared to the samples given the IHT, as seen in Figure 8deformation behavior. Thus the deformation characteristics definitely changed due to the elimination of residu created such a large debit in the creep rupture lifetimes. Note that a similar detrimental effect of RE additions was observed during tensile testing at temperatures above 850 C / 204 MPa alloys were largely inconclusive. o8.4 Te mperature Dependence of Creep Behavior C. Addition of RE elements to CM247LC was found to be detrimental to the creep rupture lifetimes at 850 oC and 950 oC. Even during the tensile testing, the RE modified alloys failed without showing any signs of plasticit y but it was not the case for temperatures less than 850 oC. Thus to understand the significance of testing temperature, creep deformation studies were conducted at 800 oC / 475 MPa and 750 oC / 625 MPa. In line with the previous test parameters, these con ditions of temperature and stress were chosen to achieve a lifetime of 300 hours. The creep results obtained from this intermediate temperature testing are plotted in Figure 8 12 and compared with previously obtained results from 850 oC and 950 oC. In all the respective categories, the CM247LC+300 ppm Pr alloy exhibited higher creep rates compared to the baseline alloy. Interestingly, the CM247LC+300 ppm Pr alloy exhibited better rupture lifetimes at 800 oC and 750 oC as seen from the figure and also to be noted is the higher minimum creep rate of this alloy at 750 oC. Recall that the CM247LC baseline alloy, when tested at 750 oC,

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189 exhibited an increased creep rate initially and then the rate of deformation slowed down to a minimum. The baseline alloy exhibited similar creep rates at 850 oC and 950 oIn the case of RE modified alloys, a pronounced temperature dependence was observed with higher temperatures having a detrimental effect on the creep properties. Fractography on these alloys exposed to creep at 850 C with the testing at the higher temperature show ing an extensive tertiary creep, whereas the former had better rupture lifetime. oC or 950 oC did not show a complete intergranular behavior as seen on the alloys subjected to tensile testing at 850 o8.5 Summary C. Si nce RE elements have a propensity to segregate to free surfaces, which include grain boundaries, further investigation into the effect of RE segregation to grain boundaries on the creep results have to be conducted. A detailed explanation on the possible r ole of RE elements on the strengthening mechanism at higher temperatures and the reasons why RE modified alloys showed an unexpectedly poor creep behavior are analyzed in Chapter 9. In all the results, the creep character was reflective of both the temper ature and stress. It would have been interesting to study the behavior at constant stress levels and at different temperatures to gain an insight into the effect of the parameters individually. The results obtained from studying the creep beha vior of RE modified CM247LC alloys can be summarized as: Addition of RE elements resulted in premature creep rupture compared to CM247LC baseline alloys. Significant differences in the creep behavior between IHT and UFHT samples were not observed. Longitudinal sections showed a large population of cracking along the primary not lead to any changes in the creep properties.

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190 Interestingly, at lower creep testing temperatures RE modified alloys were observed to exhibit better creep rupture lifetimes. Figure 8 1. Creep deformation behavior of Phase I UFHT Un HIPped alloys subjected to testing at 850 o C / 240 MPa. Addition of 860 ppm Ce and 180 ppm Dy was observed to incr ease the creep rates compared to CM247LC baseline alloy.

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191 Figure 8 2. Creep deformation behavior of Phase I UFHT HIPped alloys subjected to testing at 850 oC / 390 MPa and 950 o C / 204 MPa. An increase in the ppm level of RE additions was observed to have an effect in increasing the steady state creep rates. Figure 8 3. Creep deformation behavior of Phase II IHT alloys subjected to testing at 850 oC / 390 MPa. Addition of RE elements resulted in premature rupture during testing.

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192 Figure 8 4. Creep defo rmation behavior of Phase II IHT alloys subjected to testing at 950 o C / 204 MPa. Similar to the previous tests, RE additions had a detrimental effect on creep rupture lifetimes. Figure 8 5. Creep deformation behavior of Phase II IHT CM247LC+300 ppm Pr a lloy subjected to 1000 hour creep rupture testing at 850 oC / 280 MPa and 950 o C / 130 MPa.

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193 100 m A B Figure 8 6. SEM micrographs showing the fracture surfaces of Phase II IHT alloys subjected to creep at 850 o C / 390 MPa. (A) CM247LC alloy showing tear ridges a nd (B) CM247LC+300 ppm Pr alloy showing the presence of pores and cracking along inter granular regions. A B 10 m Figure 8 7. SEM micrographs showing the longitudinal sections of alloys subjected to creep testing. (A) CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd allo y after failure at 850 oC / 390 MPa showing grain boundary cracking and (B) CM247LC baseline alloy after failure at 950 o

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194 Figure 8 8. Creep deformation behavior of Phase II UFHT alloys subje cted to creep testing at 850 o C / 390 MPa. Addition of RE elements resulted in increased creep rates but premature creep rupture. Figure 8 9. Creep deformation behavior of Phase II UFHT alloys subjected to creep testing at 950 oC / 204 MPa. An increase i n the creep rates was observed due to RE additions but resulted in premature creep failures.

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195 A B 20 m Figure 8 10. SEM micrographs showing the fracture surfaces of Phase II UFHT alloys subjected to creep at 850 o C / 390 MPa. (A) CM247LC baseline alloy showing tr ansgranular cleavage facets and microvoids coalesced around carbide particles (B) CM247LC+300 ppm Pr showing very severe cracking in the intergranular region. A B 10 m Figure 8 11. SEM micrographs showing the longitudinal sections of the alloys after creep fail ure. ( precipitates in the regions closer to the fracture surface.

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196 Figure 8 12. Creep deformation curves of Phase II UFHT alloys subjected to creep testing at the conditions indicted in the graph. A clear dependence of creep rupture on temperature was observed for CM247LC+300 ppm Pr alloy.

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197 CHAPTER 9 THE RARE EARTH EFFEC T The objective of the ch apter is to use the results obtained in this work and apply some of the previously established theories in literature to build a mechanistic understanding of the ef fect of RE elements on the hightemperature oxidation and mechanical deformation behavior of Ni base polycrystalline superalloys. The inherent properties of RE elements, which make them s o unique in a superalloy matrix, are used extensively in this chapter to address several questions. 9.1 RE Effect on the Oxidation Behavior of Uncoated Alloys D ue to the extensive and nonuniform transition oxidation observed during the high temperature exposure of CM247LC alloys, thickness differences in the oxide scale as a function of RE element additions could not be reliably quantified. But, it was very obvi ous from the weight gain data that RE additions have a definite influence on the oxide growth and thereby, oxidation kinetics. During exposure to cyclic oxidation, two types of stresses, namely, growth stre sses due to the scale formation, and thermal stres ses due to the cycling between low and high temperatures are observed. During the cooling stage, differences in thermal expansion coefficient between the oxide and substrate lead to cracking at the oxide/alloy interface and subsequently to the spallation o f the oxides. These stresses are a function of the scale thickness with the thinner scales generated in RE modified alloys helping to mitigate the cooling stresses leading to greater spallation resistance. With the absence of convincing information on the effect of RE elements on the scale thickness and microstructure, the only evidence that can be further disseminated was the observation of Ce and Pr incorporated into the NiCr2O4 spinel oxide scales at the gas/oxide interface at 1010 oC and 1079 oC. Auger spectroscopy studies performed on the CM247LC+860 ppm Ce alloys exposed to isothermal oxidation at 1010 oC also supported the presence of Ce at the gas/oxide interface and showed that they were precipitates of CeO2. Now, considering the

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198 observation of (RE) Ox at the gas/oxide interface as the end point, the following sections will discuss the possible chain of events that lead to the diffusion of RE elements from its alloy source to the gas/oxide interface. Figure 9 1 schematically illustrates the possible c hain of events Al2O3 formed as the continuous, thermodynamically stable and protective oxide scale closer to the alloy interface. A l2O3, Cr did not form as continuous oxide layer and was observed to be a part of the spinel oxides. Oxidation behavior due to the formation of spinel oxides are not well understood at this time. In this context, the subsequent discussions would be centered Al2O39.1.1 Thermodynamic Affinity towards Oxygen The primary reason that justified the use of RE elements as gettering agents during the casting process was their favorably higher free energy of formation of sulphides and oxides [25] CM247LC alloys employed in this research contain S levels only in the range of 12 ppm. Thus, the property that is being considered in this disc ussion is their higher affinity towards O2 -. Havi Al2O3 is the thermodynamically stable and protective oxide scale, Table 9 1 co mpares the various cations in terms of their affinity towards oxygen at 1400 K. [68, 190] It is usu al to compare the relative affinities of various elements in terms of the free energy per mole of O. Since the following discussion would be based on the segregation of cations to the oxide scales, it would be more appropriate to compare the oxygen affinit ies based on per mole of the respective cations [68] Considering Al as the basel ine for comparison, RE elements, namely, Ce, Pr, Dy and Gd are observed to possess much higher affinity towards the formation Al2O3 is the only oxide scale present on the alloy, at an oxidation temperature of say, 1100 oC, the corresponding oxygen partial pressure for the Al/Al2O3 equilibrium is PO2 = 1029 atm compared to a pressure of 1 atm existing at the gas/alloy

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199 interface. Thus, an oxygen potential gradient is set up along the entire cross section of the oxidized al loy extending all along from the gas interface. It is clear that the major factor that lead to the diffusion of RE from the alloy substrate to the gas interface is the combined effect of oxygen affinity and a potential gradient to drive the diffusion of REn+9.1.2 Ionic Size Misfit and Its Effect on Segregation towards the gas/oxide interface [9, 68, 69, 85] I n response to the oxygen potential gradient, REn+ diffuses outwards and momentarily segregates at the oxide/alloy interface. Their segregation to this interface is governed by their larger ionic radii compared to the Ni base matrix. Following the advantage s of RE elements explained before, one other major factor that influences the oxidation and will be described below is their ionic size misfit with the other constituent elements in the oxide scale. Table 9 2 illustrates the ionic size misfit of the REn+ c ompared to three major ions in the oxide scales, Ni2+, Cr3+ and Al3+. As seen from the results, REn+ exhibit ionic radii which is 1.7 1.9 times greater than that of Al3+. Further outward diffusion from the oxide/alloy interface is determined by the above mentioned size misfit. Accommodation of larger REn + in the bulk of the oxide lattice would lead to increase in the elastic strain energy which indicates that the easiest path of the diffusion for REn+ are through oxide scale grain b oundaries and line defe cts such as dislocations. Al2O3 is close packed and relatively free from line defects, which leaves grain boundaries as the most probable paths for the diffusion of REn+A factor called critical ionic radii (R) defined by the ratio of ionic radii of segregant and ionic radii of the host metal was considered important in determining the effectiveness of a specific segregan t, with R values in the range of 1.4 1.6 considered to be beneficial to effect Another obvious reason is the rapidity of the diffusion along gra in boundaries compared to through the bulk lattice.

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200 the interfacial segregation [191] For oxygenactive elements to be effective segregants, Pint [68] suggested a minimum ionic radii of ~ 0.7 In fact as listed in Table 9 2, the R value s and ionic radii are higher than the required levels to be effective segregants. Thus, with ionic radii greater than 0.9 along with their higher oxygen affinity would allow for the effective segregation of REn+ (C along the oxide scale grain boundaries. Al so, an expression that relates the grain boundary segregation to the ionic size misfit is given as, gb / Clat) x rAl/rAl)2 where C ( Eqn. 9 1 ) gb and ClatAlso to be mentioned is the beneficial role played by RE elements in eliminating the harmful effects of any residual indigenous S present in the material, which has a higher propensity to segregate to the oxide/alloy interfaces. RE elem ents increase the activation energy for void growth, which is otherwise encouraged by the presence of S [10, 64, 107, 108, 112, 116, 121, 123] are the concentration of the segreg ant (x) in the grain boundaries and lattice respectively. Thus, the greater the misfit, the greater is the level of ionic segregation along the grain boundaries. The mechanism of how this segregation assumes greater significance will be the subject of the following section. 9.1.3 Segregation to Oxide Grain Boundaries and Its Effect on Oxidation Mechanism The segregation of REn+ to the oxide scale grain boundaries leads to significant changes in the oxidation mechanism that directly affects the oxidation kinetics. In the alloys without RE additions, the oxidation reaction progr esses by a mixed mode of inward oxygen diffusion and an Al2O3 scale grain boundaries and to a minimal extent through the bulk scale. This heterodiffusion mechanism results in the formation of an equiaxed ox ide scale. For the reasons emphasized in the previous sections, REn+ segregate to the oxide scale grain boundaries and block the outward diffusion of other oxygen active elements

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201 like Al3+, Ti4 + etc. Due to this effective blocking mechanism, the oxidation reaction depends primarily on the inward diffusion of O2 with the oxidation reaction occurring at the oxide/alloy interface. Thus, a preliminary short lived hetero diffusion creates an equiaxed oxide layer followed by inward O2 diffusion creating a colum nar oxide scale. Question may be asked why Ti4+, despite its near equal affinity towards O2 would be ineffective in acting as a blocking agent similar to REn+. However Ti4+ has a smaller ionic radius with an R value Al2O3 lattice. Ti4+ similar to Al3 + has a higher diffusivity, whereas a slower diffusivity through the grain boundaries would be more effective in retarding the outward diffusion of other oxygen active elements. Summarily, larger ionic radii, higher oxygen affinity and sl ower diffusivity are the primary requirements for an ion to be effective in reducing the oxidation kinetics. This change in oxidation behavior with the addition of RE elements has been routinely observed in several Ni Cr Al alloys where the primary oxide s Al2O3Unlike the previously cited reports, SEM observations did not reveal any differences in the scale morphology in the alloys modified with RE additions. To analyze the microstructural changes due to RE additions, thin foils w ere prepared using FIB techniques from the CM247LC+860 ppm Ce alloy exposed to isothermal oxidation at 1079 oC for 1000 hours. A typical microstructure taken using TEM near the grain boundary region is given in Figure 92. Long columnar grains were observe d but the presence of an equiaxed zone near the gas/oxide inter face reported in literature was not noticed. It should be noted that the published data pertains to the model Al2O3, as oxidation studies reported on Ni base superalloys discussed only the oxidation kinetics without mentioning such microstructural c hanges, attributed to the large scale transition oxidation. To obtain qualitative information about the segregation of REn+ to the oxide grain boundaries, extensive analysis was conducted using EDAX i n HR TEM through line

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202 scanning and mapping analysis methods. The compositional analysis of Ce and Hf taken on the grain boundary are shown in Figure 93. Though Hf was observed to be segregated to the grain boundary region, a reliable signal for the prese nce of Ce was not obtained. Even after lengthy compositional analysis covering several grain boundaries and oxide/alloy interface in the TEM foil, it was not possible to get the required information about the REn+9.1.4 Precipit ation of RE Oxides at Gas/Oxide Interface segregation. The reason why 860 ppm Ce was chosen is the thought that 860 ppm would logically have higher probability of appreciable grain boundary segregation. At this point, more samples exposed to different oxidation span of oxidation and even higher ppm levels may be necessary. REn+ segregated to the oxide scale grain boundaries are not static segregants and continue to diffuse outwards towards the higher oxygen potential situated at the gas interface as proposed by Pint [85] in his widely accepted and popular Dynamic Segregation Theory. As a final point in the diffusion process, REn+ reach the gas interface and oxidize forming the (RE)Ox nuclei. As the REn+ continues to diffuse, they reach the gas int erface and combine wi th the previously formed nuclei, which, upon supersaturation, lead to the formation of (RE)Ox precipitates. At this point two things remain unclear. Firstly, due to the inflow of O2 and outflow of REn+ through the same channel of grai n boundaries, why would not the oxidation reaction taking place at the gas interface occur at the grain boundary itself? Secondly, are the REn+ oxidizing to form (RE)Ox or are they incorporated int o the spinel oxide scales? Although segregation and diffus ion through Al2O3Due to the above complications, to simplify the mechanistic understanding, the transition oxida tion has to be neglected while confining the RE effect to only Al scale grain boundaries is fairly well established, similar phenomena through transition oxide and spinel oxide scales are not yet understood. 2O3 scales and grain boundaries. Thus, it is assumed that observed precipitation during SEM observation were

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203 (RE)Ox forming at the gas interface. Due to the transition oxidation, a major fraction of gas/oxide interface is constituted by spinel oxides and that is the reason why RE elements were observed in them. Similar to HfO2, these RE oxides are also typically fast oxygen conductors and excessive formation of RE oxide precipitates would lead to faster inward conduction of O2 and enhanced oxidation kinetics. One major inference from the observation of (RE)Ox at the gas/alloy interface is REn+ from the alloy substrate definitely influences the oxidation behavior. The only available paths for the diffusion of REn+9.1.5 Optimum levels of RE additions through the oxide scales are the grain boundaries and despite the lack of TEM evidence, it is proven indirectly through this precipitation. In case of CM247LC alloys with only one RE additions, the overall oxidation results indicated a very beneficial effect of Pr, Dy and Gd in improving high temperature oxidation behavior. On the other hand, CM247LC+860 ppm Ce containing alloy exhibited reduced oxidation resistance under both isothermal and cyclic oxidation conditions. All of these four elements possess all the required characteristics for better oxidation behavior, as outlined previously. The poor performance of Ce modified alloys during oxidation raises one important question Under similar conditions, does oxidation behavior depend on the quantity of RE e lements added? To illustrate the weight gains during isothermal oxidation as a function of ppm levels of RE additions, Figure 94 shows the weight gains after 25 and 200 hours during isothermal exposure at 1010 oC and 1079 oC. At 1010 oC, though not very significant after 25 hours, a pronounced reduction in weight gain kinetics was observed after 200 hours in alloys containing RE additions in the range of 120300 ppm. The relationship is stronger at 1079 oC with 860 ppm resulting in rapid weight gains afte r 25 hours and subsequently suffer higher weight losses after

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204 200 hours. A similar comparison was done for cyclic oxidation testing using the spallation kinetics during the initial few hundred hours of exposure for three temperatures, 1079 oC, 1121 oC and 1150 oC as shown in Figure 95. Conclusions similar to the previous discussion can be drawn here as well regarding the detrimental effects of 860 ppm RE addition. One interesting observation was that the magnitude of detrimental behavior due to 860 ppm RE addition was found to be increasing with the temperature of exposure, being less detrimental at 1079 oC and becoming to significant difference at 1150 oC. Attempting to include the CM247LC alloys which contain more than one RE elements into this discussion, Figure 96 illustrates the oxidation kinetics as a function of temperature for alloys containing RE elements in the range of 0 860 ppm. In case of alloys with multiple RE additions, the total ppm levels were taken into consideration. A distinct function relating the ppm levels to Kp could not be drawn. But, it is very clear that at higher temperatures, the differences are very significant. Based on the comp arisons, the results are indicative of the fact that the optimal levels for better oxidation perform ance are in the range of 120 340 ppm RE. The reason why 860 ppm Ce additions proved to be detrimental needs to be addressed. Along the same fundamentals of segregation, diffusion along the grain boundaries and finally precipitation of (RE)Ox at the gas/o xide interface, increasing the levels of RE additions would lead to enhanced diffusion and higher volume fraction of (RE)Ox. Due to the higher oxygen conductivity through these oxides, higher volume fraction would naturally lead to higher ingress of O2 -, t hus increasing the oxide scale thickness. This explains the increased oxidation kinetics at 1010 oC isothermal oxidation for CM247LC+860 ppm Ce alloys. Thicker scales would lead to higher buildup of stresses leading to cracking at the oxide/alloy interface and consequently, spallation of the oxide exposing the base metal to further oxidation degradation.

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205 Assuming that REn+ elements diffusing outward along the grain boundary do react with the inward diffusion O2 -9.1.6 Advantages of CoAlloying this would lead to the formation of oxides near the grain boundaries within the existing oxide scale. The volume expansion would lead to the build up of lateral stresses increasing with time and oxidation along the grain boundaries. With addition of higher ppm levels, this reaction leading to later al stress build up would probably continue until the oxide scale buckles to relieve the stresses. This buckling would further lead to decohesion of the oxide scale, resulting in spallation. Both the above discussed effects are illustrated in Figure 9 7. While considering RE elements for improving the oxidation resistance, it is important to carefully control their ppm levels within the optimal range. Co doping or coalloying was proposed as one alternate strategy to improve the oxidation resistance of the alloys. Two commercial Ni base superalloys namely Rene N5 (34 ppm Y, 33 ppm Zr, 540 ppm Hf) and Haynes 214 superalloy (25 ppm Y and 120 ppm Zr) were found to exhibit reduced oxidation kinetics and better scale adhesion under thermal cycling conditions at 1100 oC compared to the respective alloys without RE additions [192, 193] In this research, CM247LC alloy contains about 1.5 wt % Hf. Due to the large differences in the addition levels, having Hf in addition to RE element can not be considered coalloying. To quantify the effect of co alloying in this research, refer back to Figure 9 6 which graphically represents the lifetimes as a function of ppm levels in various alloys containing single or multiple additions. As discussed previously, RE additions in the optimal ranges contributed in a beneficial towards oxidation resistance, but the effect of co alloying is not readily apparent. Through a combination of any specific element with another RE element, it was possible to observe to any synergistic effect on oxidation resistance. Coalloying could potentially redu ce the amount of individual RE element

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206 levels but impacts the oxidation kinetics in a similar way for segregation to the oxide scale grain boundaries and blocking the outward diffusion of oxygen active cations. 9.1.7 Effect of Oxidation Temperature The ef fectiveness of RE in improving the oxidation resistance is considered to strongly depend on the temperature of oxidation also for the following reasons. REn+ are heavy and usually require sufficient thermal activation to diffuse outwards to appreciable lengths along the oxide scale. Secondly, at lower temperatures (< 1000 oC), fast growing and less dense metastable forms of Al2O3 namely Al2O3 [58, 60, 61] To see more clearly the effect of temperature on the RE effect on the oxidation behavior, refer to Figures 94 and 9 5. Using the case of the CM247LC+860 ppm Ce alloy as a reference for this discussion, at higher temperatures, it becomes more clear that the 860 ppm Ce addition increases the rate of oxidation that lead to a higher degree of spallation. At 1150 oC, Ce4 + diffused outwards and by the mechanisms previously established, lead to higher degree of internal oxidation within the oxide scale. But, the effect was less significant at either 1121 oC or 1079 oC due to lower driving force available for diffusion outwards. Also in other RE additions, oxidation at 1150 o9.1.8 Effect of Hf on the Oxidation Behavior C resulted in significant reduction in oxidation rate over the baseline alloy compared to other two temperatures. Similar to RE elements, Hf also was traditionally employed to getter S from the melt. Hf also found use for improving grain boundary strength, increasing the resistance to hot tearing in directionally solidified alloys. During oxidation, the higher thermodynamic affinity of Hf for oxygen drives the outward diffus ion towards the gas/oxide interface. Segregation of Hf to the oxide/alloy interface eliminates the detrimental effects of S and, hence, improves the adhesion of

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207 the oxide scale [64, 194] In addition, Hf was observed to segregate strongly to Al2O39.1.8.1 Internal Oxidation Stringers grain boundaries [68, 69, 165] CM247LC alloys used in this study contained about 1.5 wt% Hf and it is the subject of this section to analyze the effects of Hf on the oxidation characteristics of uncoated alloys. As disc ussed in the Chapter 2, the formation of inwardly protruding stringers rich in the oxides of reactive elements were widely believed to provide a keying in effect, anchoring the oxide scale to the underlying alloy substrate [7, 92, 93] Due to ppm levels of RE additions, stringers rich in RE were not observed but HfO2 stringers were widespread in the alloys exposed to both isothermal and cyclic oxidation conditions. In the alloys oxidized isothermally at 1010 oC and 1079 oC, HfO2 was frequently observed as small particles incorporated into the oxide scales. Diffusion of O2 through HfO2 was reported to be many orders of magnitude higher than the Al2O3. Continued diffusion of O2 through its lattice leads to thickening or elongation of the HfO2 particles, protruding into the substrate giving the appearance of a peg. Also, the incoherent boundary between HfO2 Al2O3 provides a channel for the faster transport of O2 along the HfO2/ Al2O3 interface and lead to enhanced oxidation of Al. Consequently, this oxidation reaction result in the complete encapsulation of HfO2 Al2O3. A Similar concept can be applied to other reactive or RE elements like Y, Ce etc [78, 96] The primary reason why Hf and for that matter, any RE element fails to form a protective continuous layer is the rapidit y of O2 diffusion through their lattices [97] Time and temperature also seemed to play a role in their formation, with the cyclic oxidation exposure at 1150 oC exhibited a large population of pronounced HfO2 string ers encapsulated within Al2O3. Refer back to Figure 5 25, an interesting feature of the HfO2 Al2O3 pegs can be appreciated. A three dimensional perspective showing that oxide scales were connected to the underlying

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208 substrate through the internally prot ruding pegs is readily apparent. 9.1.8.2 Role of Hf on the Oxidation Kinetics Pegs were considered as a crack deflecting obstacle locat ed at the oxide/alloy interface, but the cracks were planar and almost parallel to the interface. Notable differences in the population of pegs between CM247LC and R E modified alloys were not present. This indicates that having a large number of internally protruding pegs did not necessarily provide any anchoring effect. Large numbers of pegs lead to out of plane tensile stresses that are detrimental to the scale adhesion. Thus, the adhesion of the oxide scale to the underlying substrate is not a mechanical effect but purely a chemical phenomenon involving the curtailed outward diffusion of metallic ions. The effect of pegs, if any, is only secondary [87, 89] Analysis of oxide scales in TEM revealed the presence of blocky internal oxides of Hf formed within Al2O3 scales as shown in Figure 9 8. These oxide particles might have been the core of the internal oxidation stringers around which the Al2O3 encapsulation usually develops as discussed in the previous section. Previous publications reported a stronger segregation of Hf Al2O3 in various Fe Cr Al alloys containing only 0.2 at% Hf [68] oxide dispersion strengthened NiAl with 1 vol% of HfO2 [69] Hf rich oxide particles were Al2O3 formed on a Ni 34.4 Al 0.056 Hf alloy after 2 hours of oxidation at 1200 oC [86] .Mapping studies conducted using HR TEM/EDAX failed to reveal any strong segregation tendency of Hf to the grain boundaries. The Hf was rather found randomly Al2O3 scale as illustrated in Figure 9 3. Formation of internal oxides could have possibly prevented the availability of free Hf2+ and hence its segregation to the grain boundaries. The purpose of Hf addition to CM247LC alloys in this research was not for oxidation resistance. Though the 1.5 wt % addition of Hf might be beneficial for mechanical properties, from the view point of oxidation resistance, it may be too

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209 high. But, it remains to be answered if 1.5 wt % Hf was beyond the optimal levels required for imparting oxidation resistance? It may also be interesting to study the synergistic effects of co alloying of Hf with the concerned RE elements, provided that Hf was added in ppm levels. 9.2 RE Effect on the Performance of Thermal Barrier Coatings 9.2.1 Improvement in the Spallation Lifetimes In the uncoated alloys, the effectiveness of RE additions was compared based on the rate of oxide growth during isothermal oxidation and scale adhesion during cyc lic oxidation. In the present case, a more straightforward approach for comparison was employed based on the spallation lifetimes of an Yttria Stabilized Zirconia (YSZ) coating. It was discussed in the Chapter 6 t hat the CM247LC+200 ppm Ce+60 ppm Pr+50 ppm Gd, CM247LC+300 ppm Pr and CM247LC+250 ppm Ce+70 ppm Dy alloys exhibited better overall spallation properties compared to other alloys exposed to isothermal oxidation at 1079 oC and 1121 oC. Figure 9 9 and 910 illustrate the effects of co alloying and op timal ppm additions on the spallation lifetimes at 1079 oC and 1121 o9.2.2 Mechanisms of Coating Failu re C respectively. It is clear that addition of RE elements in the range of 300 320 ppm was found to be beneficial in increasing the spallation resistance of the TBC coatings. When it comes to co alloying, three RE additions with their total a mount ~ 300 ppm was found to impart maximum benefit. As in the uncoated case, any synergistic effect between two or three RE additions, if present, is not apparent. Understanding the failure mechanisms of the TBC system is crucial to predict its durability under the service conditions. TBC system comprises of three important constituents namely top ceramic coat, alumina layer and a metallic bond coat. Table 9 3 lis ts the various important properties that are instrumental in predicting the lifetimes of the coating system. It is evident from these complex values that mechanisms underlying the failure phenomenon are very difficult

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210 to understand. Various theoretical models based on experimental observations have been proposed to deliver guidelines on the durability of the coatings [169, 170, 173, 195, 196] In the present context, discussing the spallation process of TBC system is beyond the scope of this discussion. Thus, a brief account of the failure process witnessed in this research is provided in the following discussion. Failure is usually defined as the localized or a complete spallation of the top YSZ coating from the rest of the structure. In EB PVD YSZ coatings, failure usual ly occurs at the TGO/BC interface whereas in APS YSZ coatings, the failure plane is mostly noticed within the YSZ layer, lying closer to the YSZ/TGO interface. Since the alloys used in the current research employed APS method to deposit YSZ, subsequent dis cuss would be limited to only this type of coating. Some of the most important features which govern the failure mechanisms are structural imperfections produced by the virtue of processing conditions like inter splat cracks, porosities and undulations / r umpling of the YSZ/TGO/BC interface [50, 170, 173, 195] These defects are schematically illustrated in Figure 9 11. The inter splat cracking, which are aligned in the direction normal to the YSZ/TGO interface, originates from the cooling and cracking of the layers when they are successively deposited during APS process. With exposure to high temperature, the severity of th e cracks also increases. Undulations were deliberately engineered into the structure to improve the anchoring of the top YSZ coating to the underlying MCrAlY BC. Only the alloys exposed to 1150 oC isothermal oxidation suffered complete delamination of the YSZ coating. On careful observation, the failure in all the cases traversed within the YSZ coating and the failure plane was coincident with the top/crest of the undulations as schematically shown in Figure 9 12. Observation on alloys exposed to 1079 oC al so showed cracking at similar locations which have not coalesced yet to articulate the complete failure.

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211 On the other hand, microstructural features, which are also very significant in determin ing the lifetimes of TBC system, are the TGO layer and transit ion/spinel oxide colonies located at the YSZ/TGO interface as shown in the Figure 610. Transition/spinel oxide colonies are brittle but in this observations so far, they did not seem to be particularly detrimental. Almost all of the reports about TBC fail ure unanimously highlight the TGO as the most damaging constituent. In uncoated substrates, the oxide layer is subjected to growth stresses but the scales are free to relieve their residual stresses by various relaxation mechanisms like rumpling and interf acial cracking. The scenario is very complicated for the constrained oxide growth like in the case of the TGO growing at the YSZ/BC interface. During exposure to high temperatures, due to the ingress of O2 Al2O3 scale forms at the undulated YSZ/BC interface. During the growth at high temperatures, the stresses are compressive in nature at the crest of the undulations. Remember that the stresses at the inter splat cracks present in the YSZ are tensile. When the TB C system is cooled down to ambient temperature, the differences in the thermal expansion coefficient leads to the resi dual stresses with the TGO that transform from previously compressive to residual tensile stresses. As well known, these tensile stresses are detrimental since they encourage the nucleation of cracks. Upon subsequent thermal cycling and with the growth of the TGO layer, the magnitude of the residual stresses continues to increase. With the lateral expansion of the stress and inter linking wi th the neighboring inter splat cracks, delamination or spallation of the YSZ coating occurs. It was reported that a direct correlation between the differences in thermal expansion coefficient of various constituents in the TBC system and the spallation lifetimes was difficult to establish [197] The case i s different for EB PVD coatings, where YSZ/BC interface is usually planar and the TGO lengthen s itself by rumpling to relieve the stresses [195] In addition to the cracking and rumpling, creep of the

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212 TGO layers can also be considered to be one of the stress relaxation mechanisms at high temperatures [178, 198] 9.2.3 Effect of RE Elements on Thermally Grown Oxide Layers During the YSZ coating deposition, the MCrAlY BC oxidizes to form a thin layer of TGO due to the heat involved in the processing conditions. Following up with the discussion in the previous section about the detrimental tensile stresses upon cooling that are held responsible for the YSZ coating, it is implicit that the TGO layer thickness contributes strongly to the magnitude of the stress. An important parameter called critical thickness of the TGO layer, hcrit is defined beyond failure usually occurs in a TBC system. In an aircraft engine application where the thermal cycling is very frequent, a 1 5 m thickness is identified as the maximum thicknes s for YSZ adhesion, whereas in the industrial gas turbines, where the isothermal cycles are usually longer, hcrit values are usually in the range of 5 15 m. An expression relating the time at a specific temperature, tt required to achieve hcrit takes the fo llowing simplified form: p crit tk h t 22 (Eqn. 9 2) From the above expression, the greater the value of oxidation rate constant, kp, shorter the time that is required to achieve the critical thickness leading to TBC failure. Now, the effect of RE elements on the spallation behavior takes perspective. From the results obtained in Chapters 5 and 6, addition of ppm levels of RE were found to slow down the oxidation kinetics w hich indicates a reduction in the scale thickness. In the alloys coated with TBC, addition of 200 ppm Ce+60 ppm Pr+50 ppm Gd reduced the thickness of the TGO layer by about 2 m at both 1010 oC and 1079 oC. It is very straightforward to attribute this red uction in thickness to the increase in spallation lifetimes due to RE additions. The background mechanism involved with how RE elements contribute to reduction in the scale growth kinetics have already been established in the

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213 section discussing RE effect o n uncoated alloys. Though the mechanistic understanding is very similar, there are important differences TGO grows under constrained conditions in the presence of a top YSZ layer. More significantly, a 150 m MCrAlY separates the substrate which is the s ource of RE and the TGO on which we are proposing it is having an effect. These questions are addressed in the following section. 9.2.3.1 Is RE Diffusion from the Substrate to TGO possible? At the YSZ/TGO interface, isolated colonies of transition oxides r ich in the oxides of BC elemen ts were observed. Previous work has reported that the formation of a mixed zone at this interface composed of t ZrO2 and Al2O3 in EB PVD deposited YSZ coatings [86, 174, 175, 177] This zone was formed during the initial stages of YSZ deposition process when the heat involved in the process lead to intermixing between the BC and the YSZ elements. In this research, TEM and SEM observations did not indicate a formation of any mixed zone similar to the published reports. Using cross sectional TEM analysis, it was observed that the Al2O3 were finer in the regions closer to the YSZ followed by a columnar layer. See Figure 617. As emphasized in Chapter 2, initial inter diffusion lead to the format ion an equiaxed zone followed by the columnar zone indicating predominant inward anion diffusion. The grains were about 0.5 m wide and about 23 m long. Though proposed by Pint [86] experimental evidence of the diffusion of RE from the substrate across the MCrAlY to segregate to the TGO was not reported. Recently, diffusion of S from the substrate to the TGO/BC interface, where it was playing a detrimental role, was reported [194] With the addition of RE elements, the segregation of S was inhibited, leading to improvement in the adhesion of TGO. In this work, despite extensive compositional analysis HR TEM, RE segregation to the TGO grain boundaries could not be observed. But, in the alloys exposed for isothermal oxidation at 1010 oC and 1079 oC, Hf C e Pr NiAl deple

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214 layer). This could probably be an intermediate step in the RE diffusion towards TGO. With the lack of experimental evidence, the possibility of RE diffusion towards TGO and segregation can be indirectly addressed using the observed Hf effect as discussed next. 9.2.3.2 Relating Hf Effect to RE Effect Similar to the internal oxidation stringers observed on uncoated alloys, HfO2 particles were encapsulated with Al2O3 in the parts of the TGO layer which was extending into bond coat. After careful analysis, it was found that Hf was not a constituent of MCrAlY bond coat which leaves the substrate as the only possible source. TEM observations of the TGO layer revealed the presence of Hf Y Al rich particles in the grain bou ndaries intersecting the BC layer near the TGO/BC interface. Refer to Figure 6 18. Similar observation was reported in a PWA 1484 sample but the NiCoCrAlY BC contained 0.25 wt% Hf [199] The ZrO2 TBC is a known source for the reactive elements like Hf, Zr, Y etc. But, the inward diffusion of these elements from YSZ ceramic top coat is not possible due to the very high oxygen potential gr adient extending from the gas interface to throughout the TBC system till the substrate. This potential gradient prevents the inward diffusion of oxygen active elements [85, 86] Thus, the only source of Hf in this case was the substrate and it was remarkab le that Hf was able to diffuse under i sothermal conditions at 1010 oC, 1079 oC and 1121 oC outwards through the bond coat and segregate to the TGO before nucleating particles rich in Hf and Y. It is now easy to relate the influence of Hf observation on our hypothesis that RE elements to do segregate to the TGO layers. Ce, Pr, Dy and Gd have higher affinity towards oxygen when compared to Hf. Thus, diffusion of Hf from the substrate to the TGO layer indirectly proves that RE elements can also behave similarly. The compositional difference betwee n Hf (1.5 wt %) and RE (ppm levels) is the only factor that lead to a very pronounced Hf observation. The significant improvements in

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215 the spallation lifetimes are testament to the fact that RE elements are indeed playing a beneficial role in the TGO layers 9.3 RE Effect on the Mechanical Deformation Behavior The mechanisms for RE elements improving the oxidation behavior even when added in ppm levels were analyzed in the previous sections. It is very important that any alloying addition meant to increase t he oxidation resistance should not play a detrimental role in the mechanical deformation characteristics and the vice versa. For instance, during the development of superalloys, Cr was reduced to give way for the addition of other solid solution strengthen ers and the RE elements have become an increasingly popular addition to Ni bas e superalloys, it becomes imperative to address their effect on high temperature defo rmation also. Some of the well documented problems associated with RE elements a re their propensity to form low melting eutectics with Ni and beyond the solubility limit, nucleate brittle intermetallic phases of the form NixCe, NixIn the light of the few published reports, the effect of REs on the deformation behavior is discussed against the backdrop of a few already established mechanisms. Although it may be easier to consider these ppm additions to be relatively trivial in the mechanical propert ies, the pronounced differences in the actual tensile and creep results are suggestive that these elements do play a marked role. Pr etc. 9.3.1 How are RE Different? When present in a Ni matrix, RE elements are unique in two ways atomic size and solid solubility. Table 94 compares the atomic radii of the RE additions with that of the Ni [200] which shows that RE elements are larger than Ni by an average of ~35 %. Other RE elements of commercial importance such as La (~ 45 % larger) and Y (~ 33 % larger) are also

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216 listed in the table. Similarly, r eactive elements like Hf and Zr, which are added for oxidation resistance and grain boundary strengthening, are larger by about ~ 15%. The difference in the size is one of the prominent reasons why RE exhibit lower solubilities in Ni. For example, the room temperature solubility of Ce in Ni is only 0.04 at % [201] Presence of a larger RE atom increases the misfit strain energy in the matrix, thus promoting the segregation to surfaces like grain boundaries, dislocations etc. Even during solidification, they exhibit stronger tendency to segregate to inter dendriti c regions as shown in Chapter 4, where Ce and Pr had partitioning coefficient values less than unity. Having identified their most probable locations in the alloys, the next step is to address the how this segregation behavior impacts the properties during high temperature loading. 9.3.2 RE Effect on the Grain Boundary Sliding At higher temperature and lower stress, grain boundary sliding becomes a pronounced deformation mechanism. This is attributed to the irregularity of atomic arrangement along the grain boundaries and availability of free spaces. Sliding along the grain boundaries leads to stress concentration at locations like grain boundary junctions, second phase p articles like carbides etc. These stresses are usually relieved by creep in the adjoining grains or cavitation in the grain boundaries and boundary intersections. This section describes the effects of RE segregation at the grain boundaries. Due to the factors defined previously, RE elements enrich the grain boundaries. This was explained by Seah and Hondros [202] b given as o c bx k (Eqn. 9 3 )

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217 where o cx is the solubility of the solvent in the matrix, k = exp ( H / RT) where H is the enthalpy, R is gas constant and T is the temperature. Preferential segregation of RE to the grain boundaries is related to a reduction in the grain boundary energy as given by Gibbs adsorption theory: b o s xRT xc 0 (Eqn. 9 4 ) o s is the solute concentration at the grain boundary and cx is the solute concentration. The solute enrichment leading to reduction in the grain boundary energy further leads to reduction in the boundary diffusivity as shown by the following expression [203] : kT D Dgb gb lat gb exp (Eqn. 9 5 ) where gbD is grain boundary diffusivity, latDis lattice diffusivity, gb is grain boundary diffusivity, is atomic volume and gb is interfacial width of the grain bounda ry. According to the brief account given above on the effect of insoluble solutes, S exhibits a very pronounced segregation along grain boundaries and free surfaces like freshly formed cavities along grain boundaries [25, 129132] By the ability to reduce surface S tends result i n smaller cavities being stable, which would otherwise need to grow to a critical size before getting s intered under the influence of applied stress ( / 2 r is interfacial energy and is the external applied stress). Under the influence of S, a large number of smaller but stable cavities would eventually linkup to produce cracking along the grain boundary leading to unexpected inter granular failure. The addition of RE elements counteracts the detrimental effect of S by locking up the S atoms as stable sulphide phases. In addition to the strong er sulfur gettering ability, the presence of RE in a relatively S free microstructure like the CM247LC alloys used in

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218 this research could impact the grain boundary related mechanisms in the following way: The addition of a larger solute atom, which has a hb would tend to improve the fit in the grain boundary regions, closer to the bulk arrangement. This increased regularity in the grain boundary regions contributes to reduction in the grain boundary diffusion as per the following expressio n [204] : c b v b bx a a m b D D2 2 2 1 *2 1 (Eqn. 9 6 ) where bDand bD are respectively, grain boundary diffusivities in the baseline and solute containing alloys, 1a and 2a are respectively, the at omic radii of the solvent and solute, m and vb are constants. This equation actually confers the combined explanation of the previous two equations. This reduction in the boundary diffusion by virtue of RE additions was also related to the decrease in the cavity growth rate [203] Further, the creep rate was found to depend on the grain boundary diffusivity as given in the following modified expression for Coble creep [133] for the minimum creep rate, : kT d D Bb 3 2/ (Eqn. 9 7 ) where 2B is a constant, d is grain size and othe r entities in this expression have already been defined. Contradicting the above description based on the reduction in interfacial energy due to RE segregation, Koval et al. [205] reported an increase in energy. There is no clear evidence in the literature pertaining to the effect of RE on the grain boundaries, but the effect of Zr on reducing the creep rate and improving the creep ductility have been reported [206, 207]

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219 9.3.3 Segregation to Stacking Faul ts It was observed using microprobe analysis that as cast CM247LC alloy was heavily segregated. During heat treatment, RE elements in the vicinity of grain boundaries can diffuse and segregate to them. But, the RE elements present in the interior of the c oarser grains would have to find locations similar to grain boundaries in order to reduce the misfit strain energy. The following expression indicates a dissociation of a/2<101> type partial dislocation into a/6<112> type Shockley partials which produce a Complex Stacking Fault (CSF): ] 1 1 2 [ 6 ] 2 11 [ 6 ] 1 10 [ 2 a a a the description of which is not within the scope of the current discussion. Assuming a single phase FCC structure, the following effect of RE elements can be proposed. In response to the applied stress, the partial dislocation s move through the matrix as pairs with the inter partial spacing governed by the stacking fault energy, SFE. SFE Larger inter pa rtial spacing indicates a lower energy. The movement of partial dislocations is confined to one specific plane. When they encounter any obstacle to the gliding motion in the lattice, they have to cross slip to continue their movement. But, for the cross sl ip, a recombination of the partials is required which is dependent on Lowering the energy would lead to difficult in the recombination and consequently a reduction in the creep rate. The stronger dependence of minimum creep rate on SFE is illustrated by the following equation: n RT Q SFEe b kT b A 3 (Eqn. 9 8 )

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220 where A is a constant dependent on dislocation substructure and mobility, is shear modulus, b is burgers vector, Q is activation energy for creep, n is stress exponent and k is Boltzmanns constant. A stronger dependence of SFE on the minimum creep rate is given by the above expression. The influence of alloying additions on the SFE SFE is well known [208, 209] In this case, RE elements present in the grain interiors could segregate to the region between the partial dislocations and increase the width of the partials, reducing the energy and increasing the activation required for recombination. The segregation of RE elements to the partial dislocations was observed in Ni Cr Ce alloys by Cosandey and Kandra [138] But the results on change with Ce additions were not conclusive due to the inherent complexities associated with obtaining a reliable value of partial separation using HR TEM. 9.3.4 Solute Drag on the Dislocation Movement In Ni rix channels are achieved by bowing only at stresses greater than the Orowan stress, l bor where l is the channel width. During the bowing process, solute atoms are attracted to the dislocations by means of strain fi elds. Cottrell and Jaswon [210] through their theoretical work, postulated that larger solute atoms have a tendency to segregate to the dilated regions in a dislocation and impose a dragging effect on their motion through the lattice. According to Orowan [211] this increase in the resistance to dislocation motion reduces the dislocation velo city v thereby reducing the minimum creep rate, v b where is dislocation density. To free the dislocation from the solute drag, higher stresses are required [212] For the case of interstitial solutes like C, N etc, this phenomenon leads to dynamic strain aging during the plastic yielding.

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221 Nabarro proposed that in addition to glide through the lattice, climb i s also equally affected by the solute atmospheres in the dislocations [213] Several theoretical calculations have identified the dragging effect of the solutes at low stresses [214217] However, at higher applied stresses, dislocations liberate themselves from the solute atmospheres due to the sufficient stress activation to glide and climb. Due to only ppm level presence, the solute drag imposed by the RE elements can be expected to be only localized. 9.3.5 Why were RE Additions Detrimental to Creep Behavior? The main criteria for improving the creep resistance was increase in the modulus and activation energy for self diffusion [36] Despite the various above proposed effect of RE elements, at the time of writing, the understanding in this research was limited to only testing and fractography. High resolution characterization to understand the deformation behavior along the grain boundaries is currently underway. Against all the proposed effect and observed results in literature, RE elements were very injurious to the creep behavior. This section analyzes the possible reasons behind the detrimental effect. Effect of RE additions in altering the grain sizes in CM247LC was not readily apparent. According to MonkmanGrant relationship, .. const tf which indicates that the steady state creep rate, is inversely proportional to the time to creep failure, ft Compared to the baseline, RE modified alloys were observed to deform at higher steady state creep rates. The next logical question is why are the RE additions increasing the minimum creep rate at all the temperatures. The fractographic analysis revealed the presen ce of secondary cracking along the grain boundaries in the CM247LC+300 ppm Pr alloys whereas the CM247LC baseline exhibited a predominantly transgranular failure. Taking help from tensile results, with increase in temperature of testing, at 850 oC, the CM247LC+300 ppm alloy suffered complete intergranular

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222 fracture. However, creep testing at low temperature high stress showed better lifetimes (Refer to Figure 8 12). With the results from the characterization of tensile fracture that intergranular failure occurred only at higher temperature and the results from creep testing that indicated that low temperature/high stress conditions performed better, the grain boundary related mechanisms assume importance. As discussed previously, segregation of RE to the gr ain boundaries reduces the diffusivity along the grain boundaries. Inhibition of boundary sliding would lead to increase in the stress intensification which can be relieved only by creep in the vicinity of the boundaries and cavitation at grain boundary tr iple junctions etc. The latter possibility, as schematically illustrated in Figure 9 13, is detrimental since the coalescence of cavities along grain boundaries would result in crack formation. It is possible that secondary cracks observed along the grain boundaries could have originated from curtailed boundary diffusivity. As per the established mechanisms, reduction in diffusivity should limit the creep rate but the results obtained illustrate an increased creep rate. So, there could be two related but competing mechanisms namely reduced diffusivity and stress buildup along grain boundary. The intensity of the latter was possibly higher that it promoted excessive cavitation. Along similar lines, over alloying of RE could have also lead to higher stress build up. At this point with the available results, correlating this relationship is not easy. Comparison of creep rate as a function of ppm levels of same addition is a part of an undergoing study in Phase III.

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223 Figure 9 1. A schematic illustration of th e several stages involved in the diffusion of RE elements from the substrate to the gas/oxide interface. Due to the larger ionic radii and thermodynamic affinity towards oxygen, REn+ Al2O3 grain boundaries. They are dynamic segregant s diffusing along the grain boundaries in response to the oxygen potential gradient between the oxide/alloy and gas/oxide interfaces. The segregation inhibits the outward diffusion of other oxygen active elements and leads to reduction in oxidation kinetic s. The oxidation in the CM247LC+RE alloys occurs predominantly by the inward oxygen diffusion leading to a columnar structure compared to an equiaxed structure in baseline CM247LC alloys.

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224 Table 9 1. Relative thermodynamic affinities of various elements towards oxygen at 1400K. Table 9 2. Ionic size misfit ratios for various RE elements used in this research. A ratio greater than 1.5 usually indicates that the element segregates very effectively to the grain boundary regions.

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225 A B 50 nm Figure 9 2. TEM microg Al2O3 grain boundaries in near the gas/oxide interface in the CM247LC+860 ppm Ce alloy exposed to isothermal oxidation at 1079 o 0 1 2 3 4 5 6 7 8 0 50 100 150 200 250 300 350 Scan distance (nm) Wt % Ce Hf C for 1000 hours in (A) Bright Field and (B) Dark Field. Figure 9 3. HR TEM EDAX line scan analys Al2O3 grain boundary showing the distribution pattern of Hf and Ce. A reliable signal indicating the presence of Ce was not obtained. The inset shows the TEM micrograph with a white line along which the line scan was conducted.

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226 -1.50 -1.00 -0.50 0.00 0.50 1.00 0 1 2 3 4 5 6 Specific Weight Gain (mg/cm2) 1010 C 25 h 1079 C 25 h 1010 C 200 h 1079 C 200 h CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr Figure 9 4. Specific weight gains of CM247LC and RE modified alloys exposed to 1010 oC and 1079 o 0 10 20 30 40 50 60 0.7 0.71 0.72 0.73 0.74 103/T (K-1) Kp, g / cm4.s CM247LC 860 ppm Ce 160 ppm Pr 180 ppm Dy 120 ppm Gd 300 ppm Pr C for 25 and 200 hours. Figure 9 5. Comparison of the cyclic oxidation weight loss kinetics for CM247LC and RE modified alloys as a function of temperature. Note that the values of Kp are negative.

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227 0 10 20 30 40 50 60 0.7 0.71 0.72 0.73 0.74 103/T (K-1) Kp, g2/cm4.s CM247LC 120 ppm 150 ppm (*) 160 ppm 180 ppm 300 ppm 310 ppm (*) 320 ppm (*) 340 ppm (*) 860 ppm Figure 9 6. Comparison of the cyclic oxidation weight loss kinetics as a function of ppm level RE additions ( indicates that the specific alloy contains multiple RE additions). Figure 9 7. A schematic illustration sho wing the possible buildup of RE rich oxides along the grain boundaries leading to increase in the lateral stresses which are detrimental to the scale adhesion. Formation of (RE)Ox at the gas/oxide interface leads to faster ingress of oxygen leading to inc rease in oxidation along grain boundaries.

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228 0.1 m A B Voids Figure 9 8. TEM micrographs showing the cross Al2O3 scale containing HfO2 3600 3800 4000 4200 0 100 200 300 400 RE Additions (ppma) TBC Spallation Lifetimes (hrs) CM247LC Single RE Two REs Three REs 860 ppm Ce 3600 3800 4000 4200 0 100 200 300 400 RE Additions (ppma) TBC Spallation Lifetimes (hrs) CM247LC Single RE Two REs Three REs 860 ppm Ce internal oxide particles as indicated by the arrows in (A) Bright field and (B) Dark field image. Formation of voids within the oxides was also noticed as shown. Figure 9 9. TBC spall ation lifetimes plotted as a function of amount of RE additions for isothermal oxidation at 1079 oC. In general, a total amount of 300 ppm RE was beneficial.

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229 1200 1400 1600 0 100 200 300 400 RE Additions (ppma) TBC Spallation Lifetimes (hrs) CM247LC Single RE Two REs Three REs 860 ppm Ce 1200 1400 1600 0 100 200 300 400 RE Additions (ppma) TBC Spallation Lifetimes (hrs) CM247LC Single RE Two REs Three REs 860 ppm Ce Figure 9 10. TBC spallation lifetimes plotted as a function of ppm levels for isothermal oxidat ion at 1121 oTable 9 3. Properties of various constituent in a TBC system (Reproduced with permission from [50] ) C. Similar to the previous case, additions of 300 ppm RE was found to be beneficial.

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230 Pores Inter splat cracks Undulation 50 m Pores Inter splat cracks Undulation 50 m Figure 9 11. SEM micrograph showing the defects typically observed within the YSZ coating namely, inter splat cracks, porosity and undulations of the YSZ/TGO interface. 200 m 200 m 200 m Figure 9 12. SEM micrograph taken on the TBC coated alloy exposed till failure at 1121 o C. The failure plane traverses along the crest in the undulations as indicated by the a rrows.

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231 Table 9 4. Differences in the atomic radii between the RE elements and Ni matrix. Larger atomic radius of RE elements drives their segregation to the grain boundary regions. Figure 9 13. Schematic illustrating the effect of RE segregation to the grain boundaries Reduction in the grain boundary diffusivity leads to cavitation along the boundaries and at the triple junctions.

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232 CHAPTER 10 CONCLUSIONS AND FUTURE WORK Four Rare Earth (RE) elements from the lanthanide series namely Ce, Pr, Dy and Gd were added to CM247LC, a Ni base polycrystalline superalloys in the range of 100900 ppm levels. Thes e alloys were subjected to high temperature oxidation exposure and mechanical testing to gain a mechanistic understanding of the RE effect on Ni base s uperalloys. The important observations and conclusions from this work are summarized in the following sections. 10.1 Microstructure Due to the larger size and low solubility in Ni, RE exhibited a strong tendency to segregate to the inter dendritic regions during solidification. This segregation behavior was illustrated using micro probe analysis on the as cast alloys which showed the partitioning coefficient (k) values less than unity. Similar to Hf, Zr, B, C etc, RE elements were considered detrimental due to their propensity to suppress the melting point and affect the solutioning behavior. Differential Thermal Analysis (DTA) showed that addition of the RE elements suppressed the solidus temperature only by an average of 10 oC over the baseline CM247LC al loy. By designing multistep solution heat treatments with a maximum temperature of 1260 oC, an optimum microstructural uniformity a nd compositional homogeneity were achieved. Although significant differences in the mechanical properties between the alloys with industrial heat treatment and modified heat treatment were not observed, microstructural observations along the longitudinal section revealed extensive cracking in the former al the latter showed uniform deform ation. In this work, for the first time, a RE Hf phase was observed to form along the grain boundaries during the solution heat treatment. Despite extensive annealing treatment conducted to characterize the evolution of this phase, the mechanism behind the formation remains still unclear. High resolution characterization is currently underway to

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233 obtain indepth information about the structure and orientation relationship of this phase with the matrix. These phases were observed to crack during the mechanica l testing but their effect on deformation remains to be studied in detail. 10.2 Oxidation Behavior Compared to the baseline CM247LC alloys, RE additions were observed to reduce the oxidation kinetics during isothermal exposure at 1010 oC and 1079 oC. Similarly, improvements in the scale adhesion were also obtained during cyclic oxidation at 1079 oC, 1121 oC and 1150 oC. Due to their larger ionic radii and thermodynamic affinity towards oxygen, RE elements segregate to oxide/alloy interface an Al2O3 scale grain boundaries. This segregation was believed to inhibit the outward diffusion of other oxygen active elements, leading to retarded oxide scale growth rate and increased adherence under thermal cycling conditions. These improvement s were noticed only within a specific range of RE additions namely 100 300 ppm. Add ing more than 300 ppm levels le d to enhanced oxidation kinetics possibly due to the faster inward oxygen conductivity of (RE)OxTo understand the effect of RE additions on the growth kinetics of therma lly grown oxide layer, and, therefore, on the spallation behavior of YSZ coatings, the RE modified CM247LC alloys coated with a TBC system were subjected to isothermal exposure at 1010 forming at the gas/oxide interface and buil dup of oxide scales along the grain boundaries resulting in increase in the lateral stresses. Both these cases le d to increase in the growth rate and spallation due to the inbuilt stresses. oC, 1079 oC and 1121 oC. Remarkably, by diffusing through the bondcoat layer which is 150 m thick, RE elements added to the substrate reduced the growth rate of TGO layer and resulted in the improvement in the spallation lifetimes by 20% over the CM247LC baseline alloy. In this case also, RE additions in the range of 100 300 ppm were found to be beneficial.

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234 10.3 Mechanical Properties During tensile testing at room temperature, 650 oC, 760 oC and 850 oC, the CM247LC baseline alloy was found to be relatively stronger and exhibited higher ductility compared to the RE modified CM247LC alloys. In the latter case, a transformation from mixed mode to complete intergranular fracture was observed at 850 oC. Creep testing at 850 oC/390 MPa and 950 oC/204 MPa also illustrated the detrimental effect of the RE additions which increased the steady state creep rate and exhibited reduced lifetimes compared to the CM247LC baseline alloy. Fractographic analysis revealed the presence of secondary cracking along the grain boundary regions in the case of t he RE containing CM247LC alloys which indicated the effect of these additions on the grain boundary related mechanisms during creep. It is believed that RE additions, on segregation to the grain boundaries reduced the boundary diffusivity resulting in the stress intensification. Cavita tion was one of the possible ways to relax the stresses at the grain boundarie s. Inter linking of cavities le d to crack formation and propagation under the influence of applied stress, ultimately leading to inter granular failure. Since grain boundary slid ing is a dominant creep deformation mechanism at higher temperatures, creep tests were conducted in the intermediate temperature, 750 o10.4 Future Work C/625 MPa. The RE modified alloys achieved longer creep rupture lifetimes but the steady state creep rate was still high er than the CM247LC baseline alloy. The grain boundary effect due to RE additions needs to be addressed with additional testing. Based on the effect of RE elements obtained from this work, the following suggestions are made pertaining to the future work on this subject: In this work, adding Pr was found to be beneficial in improving the oxidation resistance of the CM247LC alloys. Thus, in Phase III, the effect of adding 100 ppm Pr and 170 ppm Pr to the CM247LC alloys on the high temperature properties will be evaluated. To establish the

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235 oxidation kinetics, spallation properties and mechanical strength as a function of ppm level of one single RE element, it would be interesting to study the effect of Pr additions in the range of 100 500 ppm levels. Due to the semi gravimetric method of obtaining the oxidation kinetics, the cooling stresses inevitably influenced the isothermal weight gain data in this work. Using Thermo Gravimetric Analysis (TGA) to study the weight gains under isothermal ex posure would be helpful in eliminating the effect of these stresses on the outcome of the kinetics measurements. Also, comparison between the results of these two procedures would reflect the contribution of cooling stress on the isothermal weight gain dat a. To gain a qualitative understanding of the distribution of RE elements in the grain boundaries, using Local Electrode Atom Probe (LEAP) technique is recommended. The resultant information would be very valuable in addressing the very important segrega tion behavior of RE in the as cast and solution heat treated structures. Similarly, the segregation tendency of the other elements in CM247LC can also be characterized. At higher levels of RE additions, Secondary Ion Mass Spectroscopy (SIMS) also would be useful in obtaining a qualitative representation of the segregation pattern. In order to further investigate the effect of RE elements on grain boundaries during mechanical deformation, it would be a very useful approach to employ insitu mechanical fractu re inside an Auger chamber. At temperatures greater than 800 o If the segreg ation of RE elements to the grain boundaries influenced the grain boundary sliding mechanisms, what if the grain boundaries are eliminated in the transverse directions or altogether? To this end, studies would be conducted on Directionally Solidified and S ingle Crystal superalloys to determine the RE effect on the high temperature mechanical behavior. In these alloys, the possibility of the RE elements exhibiting higher tendency towards segregation in stacking faults and dislocation cores are higher. HR TEM analysis would be instrumental in validating these proposed effects. C where the inter granular fracture is dominant, Auger spectroscopy can provide a reliable compositional data of ppm level RE additions from the freshly exposed fracture surfaces. Studies to understand the evolution of the RE Hf phase during solution heat treatment and TEM analysis to establish the crystal structure and orientation relationship with the matrix ar e currently in progress and will be reported in the future publications.

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253 BIOGRAPHICAL SKETCH Krishna Prakash Ganesan was born in September 12th 1983 in Virudhunagar in the southern state of Tamil Nadu, India to Kasi Ganesan and Chitra Ganesan. His father is a project director in a nonprofit organization while his mothe r is a home maker. He grew up with a strong interest in space science, athletics, cricket and global affairs In a nation where cricket enjoys a larger than life status, he played for his high school team as a pace bowler and had been fanatic follower of the Indian national c ricket team since childhood. He graduated high school from KVS Matriculation Highe r Secondary School where he was the captain of the school quiz club and represented the school in numerous inter school competitions Aspiri ng to become a cardiologist, he fell short of the medical college requirements by one point in his public examination Then, he decided to take up Metallurgical Engineering in one of the premier engineering colleges in the country, National Institute of Te chnology, Trichy. While at NITT, he was the chairman of Mettle 2004, a national level technical symposium for students on metals and materials. After graduating with distinction, he was offered a j ob at TVS Motor Company, Hosur where he worked as a process engineer for one year His strong interest in high temperature materials inspired him to join Materials Science and Engineering program at the University of Florida during fall 2005. He started working for Dr. Gerhard Fuchs during the summer 2006 towards his dissertation. In addition to the privilege of working in the area he was interested in, he thoroughly enjoyed his days at Gainesville with his wonderful friends. He was fortunate to attend UF during the days when the gator football was reigning the co llege football with 2 NCAA championships in 3 years. Apart from rooting for the gators, he loves watching sports (cricket, football and baseball), reading fiction listening to heavy metal and traveling He graduated with a PhD degree in Materials Science & Engineering in December 2009.