Thermomechanical Processing of a Gamma-Prime Strengthened Cobalt-Base Superalloy

DISSERTATION

Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University

Donald S. Weaver

Graduate Program in Industrial and Systems Engineering

The Ohio State University

2018

Dissertation Committee:

Professor Rajiv Shivpuri, Advisor

Professor Jerald Brevick

Professor Hamish Fraser

Copyrighted by

Donald S. Weaver

2018

Abstract

A novel class of gamma-prime strengthened cobalt-based superalloys may enable a significant temperature and efficiency capability improvement relative to nickel-base superalloys for future generation turbine engine hardware. However, little information exists regarding deformation processing of these novel Co-Al-W alloys into useable product forms with the necessary microstructure refinement at an industrially relevant scale with industrially relevant processes. To address this need, an ingot metallurgy thermomechanical processing sequence was demonstrated for a novel class of cobalt-base gamma-prime containing superalloys. From an as-cast ingot, the material was characterized and a homogenization heat treatment was developed and executed to reduce residual segregation from casting. Representative ingot conversion steps using extrusion were evaluated and performed followed by a recrystallization heat treatment to produce the desired fine-grain, wrought microstructure. Deformation processing of wrought material was completed at supersolvus hot-working temperatures using both cylindrical upset specimens to establish flow-stress behavior and custom-designed double-cone upset specimens to experimentally quantify the effect of strain, strain-rate, and temperature in microstructure evolution during hot-working, including the dynamic recrystallization and grain growth. All upset testing was completed at two supersolvus temperatures (1149 °C or 1204 °C) and one of three strain-rates (0.01/s, 0.1/s, or 1.0/s) depending on the type of

ii testing completed. Required thermophysical and thermomechanical data was determined for material property inputs to a finite element model which was used to correlate observed microstructures to location-specific thermomechanical processing history.

As part of this development, a significant effort was undertaken at each stage of processing to sufficiently characterize the microstructure through optical microscopy, electron microscopy, or other specialized testing. Large-area electron backscatter diffraction was completed to document the microstructure evolution during subsequent thermomechanical processing sequences.

The microstructure data from the characterization was used to fit Johnson-Mehl-

Avrami-Kolmogorov equations for dynamic recrystallization and a traditional grain growth equation. The results were compared to similar equations generated for a nickel- base superalloy, Waspaloy. The microstructure fits were incorporated into a commercial finite element software package and used to predict location-specific microstructure information and compared to experimental results.

This work established ingot metallurgy was a feasible processing route for this novel class of gamma-prime strengthened cobalt superalloy and determined and demonstrated thermomechanical processing sequences for successful homogenization, ingot conversion, and wrought processing at an industrially-relevant scale. The microstructure model fits, incorporated in finite element software, successfully reflected the microstructure evolution and showed similarity between a legacy nickel superalloy,

Waspaloy, and confirmed the suitability of traditional JMAK microstructure evolution models for novel gamma-prime strengthened cobalt superalloys.

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The strong similarity between nickel- and cobalt-based gamma-prime strengthened alloys was also promising from an industrial perspective. Similarity of thermomechanical processing sequences meant conventional machines, processes, and tooling could be used for manufacturing. Utilizing updated, well-accepted existing computational modeling tools enabled process and location-specific microstructure and mechanical property simulation and optimization at less effort than a full development.

The result of the demonstrated similarity was reduced barriers to transition and eventual implementation of this novel alloy class.

For my wife and my children. You are my inspiration and my joy.

I am grateful for each day we have together.

Acknowledgments

I express my sincere appreciation to my advisor, Dr. Rajiv Shivpuri, for his endless support and patience over my many, many years of study at The Ohio State

University. His input and diverse expertise have helped broaden and shape my perspective during my doctoral endeavors. I also thank my doctoral committee members,

Dr. Jerald Brevick and Dr. Hamish Fraser. I have known both for many years, and am grateful for their steadfast willingness to help me see this process finished.

I had many collaborators and supporters during this effort. I thank Dr. S. Lee

Semiatin for his continued mentorship, friendship, and enthusiasm. Dr. Semiatin’s excitement for traditional ‘heat-and-beat’ metallurgy is what initially drew me to this research area, and I will always remember to “look for the pony.” I am thankful for

(soon-to-be Dr.) Katelun Wertz for leading the in-house cobalt effort. I am grateful for her support and friendship and all she does to keep the cobalt project moving forward.

I am appreciative of Dr. Eric Payton and Dr. Adam Pilchak for providing financial support over the years to work in a novel, rapidly developing area. Special thanks go to the on-site research staff and technicians for their world-class support of this research both in the Metals Processing Laboratory (Joe, Travis, and Pat) and the

Materials Characterization Facility (Jared, Kathleen, Tommy, and Bob) in the Structural

Materials Division of the Materials and Manufacturing Directorate. I am also grateful for

Alana Parey and Ryan Hendrickson for their metallography and microscopy support.

External collaborators have also impacted the effort positively. Dr. Erin

McDevitt at ATI Specialty Materials kindly produced material for the project and was available for technical discussions on this novel alloy system. I would like to thank Dr.

Wei-Tsu Wu and Scientific Forming Technologies Corporation for use and discussions of their DEFORM software.

Lastly, I sincerely appreciate Dave M, Jay T, Eric B, TJ T, Jon M, and Dan E for their continued friendship and their various efforts to keep me “focused” on finally finishing—even after all this time—in your own unique ways.

This research was supported by the Structural Materials Division of the Materials and Manufacturing Directorate of the Air Force Research Laboratory.

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Vita

2003...... B.S. Mechanical Engineering, University of

Dayton

2007...... M.S. Mechanical Engineering, University of

Dayton

2003 to present ...... Research Engineer, Materials and

Manufacturing Directorate, Air Force

Research Laboratory, Wright-Patterson Air

Force Base, OH

Publications

[1] A.L. Pilchak, D.L. Ballard, D.S. Weaver, S.L. Semiatin, Metallurgical and Materials Transactions A, 42 (2011) 1089-1102.

[2] D.L. Ballard, D.S. Weaver, A.L. Pilchak, S.L. Semiatin, Journal of the European Ceramic Society, 30 (2010) 2305-2312.

[3] D.S. Weaver, S.L. Semiatin, Scripta Materialia, 57 (2007) 1044-1047.

[4] S.L. Semiatin, D.S. Weaver, R.L. Goetz, J.P. Thomas, T.J. Turner, Materials Science Forum, 550 (2007) 129-140.

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[5] S.L. Semiatin, D.S. Weaver, R.C. Kramb, P.N. Fagin, M.G. Glavicic, R.L. Goetz, N.D. Frey, M.M. Antony, Metallurgical and Materials Transactions A, 35 (2004) 679- 693.

Fields of Study

Major Field: Industrial and Systems Engineering

Table of Contents

Abstract ...... ii

Acknowledgments...... vi

Vita ...... viii

Publications ...... viii

Fields of Study ...... ix

Table of Contents ...... x

List of Tables ...... xiv

List of Figures ...... xv

Chapter 1: Introduction ...... 1

1.1 Motivation ...... 2

1.2 Gamma-Prime Strengthened Cobalt Superalloys ...... 3

1.2.1 Mechanical Behavior of Cobalt Superalloys ...... 4

1.2.2 Factors Relevant to Thermomechanical Processing ...... 6

1.3 Thermomechanical Processing of Superalloys ...... 7

1.4. Computational Modeling Tools ...... 9 x

1.4.1 Deformation Process Models...... 9

1.4.2 Microstructure Evolution Models ...... 10

1.4.3 Thermodynamic Modeling ...... 11

1.5 Summary of Opportunities and Objectives of this Research ...... 11

Chapter 2: As-Cast Cobalt Ingot Characterization ...... 13

2.1 Ingot Sectioning and Naming Conventions ...... 14

2.2 Macro Section Characterization ...... 16

2.2.1 Optical Macrographs ...... 18

2.2.2 Primary Solidification and Growth Direction Characterization ...... 24

2.3 As-Cast Microstructure Characterization ...... 26

2.3.1 As-Cast Texture and Grain Diameter ...... 26

2.3.2 Interdendritic Phase Identification ...... 31

2.3.3 Gamma-Prime Characterization ...... 35

2.4 Thermophysical Property Data...... 38

Chapter 3: Cobalt Ingot Homogenization ...... 43

3.1 Analytical Homogenization Time ...... 43

3.2 Experimental Homogenization Procedures ...... 49

3.2.1 Ingot Coupon Homogenization Procedures ...... 49

3.2.2 Ingot Coupon Homogenization Results ...... 51

3.3 Experimental Ingot Material Homogenization ...... 57

3.3.1 Ingot Homogenization Procedures ...... 58

3.3.2 Ingot Homogenization Results ...... 61

Chapter 4: Cobalt Ingot Conversion ...... 64

4.1 Isothermal Compression Testing of Cobalt Ingot Material ...... 64

4.1.1 Isothermal Compression Testing Materials and Procedures ...... 65

4.1.2 Isothermal Compression Testing Results and Discussion ...... 68

4.2 Cobalt Ingot Extrusion ...... 75

4.2.1 Extrusion Materials and Procedures ...... 76

4.2.2 Extrusion Results and Discussion...... 78

4.3 Post-Extrusion Recrystallization Heat Treatment ...... 82

4.3.1 Subscale Recrystallization Heat Treatment Materials and Procedures ...... 82

4.3.2 Subscale Recrystallization Heat Treatment Results and Discussion ...... 83

4.3.3 Extruded Workpiece Heat Treatment Procedures ...... 84

4.3.4 Extruded Workpiece Heat Treatment Results and Discussion ...... 84

Chapter 5: Wrought Processing of Cobalt Ingot Material ...... 88

5.1 Isothermal Cylindrical Compression Testing ...... 88

5.1.1 Wrought Cylindrical Compression Testing Materials and Procedures ...... 89

5.1.2 Wrought Cylindrical Compression Testing Results and Discussion ...... 90

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5.2 Isothermal Double-Cone Compression Testing ...... 99

5.2.1 Isothermal Double-Cone Compression Testing Materials and Procedures ...... 99

5.2.2 Isothermal Double-Cone Compression Testing Results and Discussion ...... 105

5.3 Static Grain Growth Behavior ...... 119

5.3.1 Static Grain Growth Materials and Procedures ...... 120

5.3.2 Static Grain Growth Results and Discussion ...... 121

Chapter 6: Empirical Microstructure Model for a Cobalt Superalloy ...... 123

6.1 Dynamic Recrystallization ...... 124

6.1.1 Peak Strain ...... 125

6.1.2 Strain for 50% Dynamic Recrystallization ...... 126

6.1.3 DRX as a Function of Imposed Strain ...... 128

6.1.4 DRX Grain Size ...... 129

6.2 Grain Growth/Coarsening ...... 131

6.3 Wrought JMAK Fit Model Results ...... 132

6.4 Comparison to a Wrought Nickel-Base Superalloy ...... 138

Chapter 7: Conclusions and Future Work ...... 144

7.1 Conclusions ...... 145

7.2 Future Work ...... 149

References ...... 154

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List of Tables

Table 1. Composition target and measurements at several stages of production (at %) .. 14

Table 2. Etchant recipe for Macro Krolls and Macro Cobalt...... 19

Table 3. Summary of thermophysical properties of a -prime cobalt superalloy ...... 38

Table 4. Homogenization heat treatment conditions...... 51

Table 5. HIP cycle used for extrusion and microscopy materials...... 60

Table 6. Cylindrical upset testing conditions for the coarse grain ingot material...... 67

Table 7. Area fraction recrystallized for 3:1 reduction coarse-grain ingot upset cylinders

...... 72

Table 8. Isothermal cylindrical compression testing conditions for wrought cobalt...... 91

Table 9. Predicted maximum die stress during double-cone compression tests...... 106

Table 10. Dynamically recrystallized grain diameter by thermomechanical condition. 115

Table 11. Comparison of peak strains of the JMAK fit with experiment...... 132

Table 12. Comparison of calculated JMAK fit values for strain to 50% recrystallization with experiment...... 133

Table 13. Summary of microstructure fits for cobalt- and nickel-base [41] superalloys.

...... 139

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List of Figures

Figure 1. Evolution of high-temperature capability of nickel-base superalloys [1]...... 2

Figure 2. Variation of specific fuel consumption against thrust for turbine entry temperature and overall pressure ratio [1]...... 3

Figure 3. Comparison of temperature dependence of flow stress for several ’ cobalt alloys versus a nickel alloy and a solid-solution strengthened cobalt alloy [7]...... 5

Figure 4. Assembled ingot sections, from left to right T, E, M, and B...... 15

Figure 5. Detail-view of section E with grinding marks and remaining surface defects. . 16

Figure 6. Cut up diagram for Microscopy slice with the as-cast samples marked and cut.

...... 17

Figure 7. Observed crack propagating from the cut surface during machining in the as- cast ingot material...... 18

Figure 8. Macrograph of diametral section, theta orientation, with center to left and OD on the right...... 20

Figure 9. Typical grain structures in solidified castings of ingots [33]...... 20

Figure 10. Primary dendrites and secondary arms seen in an as-cast Co-Al-W alloy...... 21

Figure 11. Transverse section of as-cast cobalt with center on the left, OD on the right. 22

Figure 12. Magnified view of dendrites of the as-cast ingot...... 23

Figure 13. Lines tracing dendrite growth direction in the ingot...... 24

Figure 14. Orthogonal sectioning lines added for each location ...... 25

Figure 15. Overlay of sectioning angle for metallographic analysis of heat treat specimens...... 26

Figure 16. As-cast IPF map of raw data as-captured (a) and after basic cleanup (b)...... 28

Figure 17. Calculated [001] PF demonstrating stack-up mis-alignment...... 29

Figure 18. Corrected IPF map of as-cast cobalt with the expected [001] fiber texture. ... 30

Figure 19. Grain ID map (left) with area fraction versus grain diameter (right)...... 31

Figure 20. EDS composition maps for as-cast cobalt superalloy...... 33

Figure 21. Solidification projection for Co-9Al-9W-2Ta...... 35

Figure 22. Backscatter electron micrograph showing high area fraction of gamma prime

...... 36

Figure 23. Backscatter electron image for determining gamma prime size and morphology...... 37

Figure 24. Specific heat traces for determination of precipitation behavior on heat up

(dashed) and cool down (solid) lines. Dotted vertical line represents Tsolvus...... 39

Figure 25. Experimental -prime approach curve (dashed line)...... 40

Figure 26. Equilibrium phase fraction predictions for Co-9Al-9W-2Ta ...... 41

Figure 27. Impact of temperature on calculated homogenization time...... 45

Figure 28. Predicted phases during homogenization...... 46

Figure 29. Estimated homogenization time including variability in calculating the effective interdiffusivity coefficient, D, at 1180 °C (top) and 1232 °C (bottom)...... 48

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Figure 30 - As-sectioned specimens for homogenization analysis...... 50

Figure 31. Encapsulated cobalt sample (left) with tantalum foil (right)...... 50

Figure 32. Cobalt ingot samples after 1232 °C heat treatment in air for times of 0, 16, 48, and 96 hours (left to right)...... 53

Figure 33. Typical minor surface oxidation of encapsulated samples as a function of time.

...... 53

Figure 34. Changes in microstructure after heat treatment time at 1232°C...... 55

Figure 35. Grain maps for (a) the as-cast ingot and (b) following 240 hours at 1232 °C. 56

Figure 36. IQ maps of as-cast cobalt (a) and cobalt after 48 hours homogenization (b). . 57

Figure 37. Typical Extrusion Section ingot pie shaped piece for heat treatment...... 59

Figure 38. Cobalt material stacked in HIP heating element assembly and carrier...... 60

Figure 39. Representative etched microstructure from the as-cast cobalt ingot...... 62

Figure 40. Representative etched microstructure of the cast plus homogenized ingot. ... 63

Figure 41. As-wire EDM bar of coarse-grain ingot compression specimens prior to final machining...... 65

Figure 42. Schematic of die-stack up and materials used during compression testing..... 67

Figure 43. Macrographs of non-uniform deformation observed in as-upset ingot material.

...... 69

Figure 44. Supersolvus true stress-true strain curves for coarse-grained cast and homogenized cobalt ingot material...... 70

Figure 45. Stress-strain curves comparing supersolvus flow behavior of coarse grain cobalt and Waspaloy deformed at 0.1/s [32]...... 71

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Figure 46. Compression direction IPF with recrystallization in black for 3:1 reduction. 74

Figure 47. Turned cobalt superalloy billet ready for canning (left), welded assembly ready for evacuation and sealing (center) and coated assembly ready for extrusion (right).

...... 77

Figure 48. Macrographs of 6:1 extrusions with nose and tail detail...... 79

Figure 49. Macrograph of 4.3:1 cobalt extrusion with nose and tail detail...... 80

Figure 50. IPF map of as-extruded cobalt material for 6:1 reduction at 1204 °C...... 81

Figure 51. Extrusion direction IPF maps showing relative grain size in the center (left) and outer diameter (right) for a 4.3:1 extrusion at 1204 °C with slow cool...... 82

Figure 52. IPF map of 6:1 reduction extruded material after recrystallization heat treatment at 1204 °C for 60 minutes with center at left...... 83

Figure 53. Typical as-sectioned extruded workpieces prior to recrystallization heat treatment...... 85

Figure 54. Canned, extruded bars following recrystallization heat treatment with instrumented dummy sample...... 86

Figure 55. Close-up of ends of the as-heat treated bar exposed to oxygen for 1 hour. .... 87

Figure 56. Extruded and heat treated bar with cylindrical compression specimens...... 89

Figure 57. Typical as-upset wrought cobalt superalloy cylindrical specimen from top view (left) and side view (right)...... 92

Figure 58. True stress-true strain curves for a wrought Co-Al-W superalloy...... 93

Figure 59. Comparison of ingot and wrought flow stress curves of a ' cobalt superalloy.

...... 94

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Figure 60. Determining slope of 1/T versus log peak strain to determine temperature sensitivity of the flow stress, S, in K...... 96

Figure 61. Changes to calculated strain and strain rate profiles with friction coefficient, mu, for a 2.75:1 upset cylinder at 1149 °C and 0.1/s...... 97

Figure 62. Load stroke comparison of experiment versus simulation for different friction conditions for cylindrical upset tests conducted at 1149 °C 0.1/s...... 98

Figure 63. Double-cone sample orientation relative to the extruded bar...... 100

Figure 64. Dimensions in millimeters for double-cone geometry used in present work. 100

Figure 65. Size comparison of double-cone geometry for coarse grain ingot work (left) versus present fine-grain wrought work (right)...... 101

Figure 66. Double-cone experimental configuration...... 102

Figure 67. Finite element representation of cobalt double cone hot compression test setup.

...... 103

Figure 68. Compression Axis IPF maps as-scanned (left) and after threshholding to highlight recrystallized areas in black (right)...... 104

Figure 69. Predicted maximum die stresses for 1204 C 0.1/s upset...... 106

Figure 70. Press load versus stroke for 1149 °C double cone upset test at 0.01/s, 0.1/s, or

1.0/s measured experimentally (dashed lines) or predicted by FEM (solid lines)...... 108

Figure 71. Wrought double-cone compressed at 1204 °C 0.01/s to 3:1 reduction...... 109

Figure 72. Simulated strain (left) and strain rate (right) profiles for 3:1 double cone upset at 1149 °C at 0.1/s for different friction coefficients, mu...... 111

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Figure 73. Finite element representation of as-deformed double-cone geometry (left).

Effective strain profile for the mid-plane computed from the finite element model (right).

...... 113

Figure 74. Compression-axis (vertical) IPF map of diametral mid-plane for all test conditions...... 114

Figure 75. Imposed strain versus area fraction recrystallized in wrought cobalt superalloy.

...... 116

Figure 76. Comparison of wrought cobalt data to wrought Waspaloy model [41]...... 117

Figure 77. Recrystallization data plotted to determine, n, or the slope of the line...... 129

Figure 78. Plot of ln (Z) versus ln (diameter) in microns from the double cone upset testing. The slope provides the value of m for Equation 19...... 130

Figure 79. JMAK fit versus experimental data for fraction recrystallized versus strain. 134

Figure 80. Improved JMAK fit for DRX by removing constraints to apparent activation energy...... 136

Figure 81. Half-symmetry double-cone upset simulation results for Von Mises strain, recrystallized fraction, and average grain size...... 137

Figure 82. JMAK fits for wrought ' cobalt and Waspaloy for similar test conditions. . 142

Chapter 1: Introduction

Metallic materials used in turbine engines are highly-engineered, multi- component alloy systems that have been optimized and improved over many decades to provide the required balance of properties at elevated temperatures. Due to the demanding service environment, significant effort has been expended developing metallic materials, more specifically superalloys, to meet strength and toughness requirements.

Traditional nickel-base superalloys have been incrementally modified both through compositional adjustments and via alternate processing methods over time to maintain the required balance of properties while increasing operating temperatures

(Figure 1) [1]. Transitioning from polycrystalline wrought alloys to polycrystalline cast alloys, then advancing to directionally solidified and eventually single crystal casting methods has improved performance by reducing grain boundaries to improve creep properties for blade materials. Similarly, incremental alloying changes to existing superalloys has increased the performance of each generation by improving targeted material properties, but room for significant improvements to nickel-base superalloys with small chemistry adjustments or elemental additions has rapidly decreased with time.

A similar improvement has been slower for polycrystalline disk materials, but is also asymptotic.

Figure 1. Evolution of high-temperature capability of nickel-base superalloys [1].

1.1 Motivation

Modern turbine engines are designed using the Brayton thermodynamic cycle, and as such, efficiency is improved through higher compression ratios and turbine inlet temperatures (Figure 2). For commercial aero-platforms and ground-based power generation, increasing efficiency is largely driven financially as fuel savings. While fuel costs are concerning to military platforms, performance-based capability and strategic advantages are also drivers. Increased thrust performance or aircraft availability/loiter time would be enabled by an increase to the overall temperature capability of a turbine engine. Thus, any increase in high temperature capability is desirable. Nickel-base alloys are relatively mature, featuring multi-component systems and specialty processing, 2

Figure 2. Variation of specific fuel consumption against thrust for turbine entry temperature and overall pressure ratio [1].

yet there continues to be significant investment in this area every year. Successive generations of single-crystal nickel alloys historically have only required a 25-30 °C increase in temperature capability to justify investment in scale-up and would represent a significant advancement in capability [2-4].

1.2 Gamma-Prime Strengthened Cobalt Superalloys

Solid solution strengthened cobalt superalloys have been used successfully in elevated temperature applications for many years but not temperatures approaching 90% of the melting point like nickel-base superalloys due to the loss of strength at elevated temperatures [1, 5]. Recent development efforts by Sato [6] have outlined an alloy 3 composition space demonstrating a metastable L12 Co3(Al-W) ’ phase structurally similar to those found in traditional, precipitation strengthened nickel-base superalloys.

The ’ precipitates, coupled with their size and distribution, are what allow the unique combination of high temperature properties required for turbine engine application. For early Co-Al-W alloys, investigations on the solidus temperature of this class of cobalt superalloys showed a potential increase of 100-150 °C relative to nickel-base superalloys

[4]. This large difference was expected to translate to a commensurate improvement in ’ solvus, but so has not yet been realized [7]. Rather, peak strength of cobalt ’,

Co3(Al,W), was found to occur at temperatures approaching 40 °C higher than those observed in nickel ’, Ni3Al [8]. This improvement is still significant and merits further investigation. Directly-replacing a polycrystalline wrought nickel-base superalloy disk with a polycrystalline wrought cobalt disk would support increasing the overall temperature capability of the turbine engine.

Prior literature in the areas of cobalt superalloys was reviewed to comprehensively understand the opportunity and determine the research approach. This evaluation focused on thermomechanical processing and robust microstructure evolution models that could be demonstrated and transitioned to a cobalt-base, ’ containing superalloy and compared to more well-documented nickel-base superalloys.

1.2.1 Mechanical Behavior of Cobalt Superalloys

The service environment of a turbine engine disk is challenging, requiring strength at temperatures of 650-800 °C and cyclic loading often exceeding 600 MPa for long durations [1]. One of the main properties that makes ’ containing cobalt 4 superalloys attractive to meet that requirement is the anomalous yield stress bump at elevated temperatures. Suzuki, incorporating data from several sources [5, 8-10], compared the 0.2% flow stress, or yield, versus temperature for a variety of alloys

(Figure 3) [7]. The ’ containing cobalt alloys were stronger across all temperature ranges where the ’ is expected to be stable than the traditional solid-solution

Figure 3. Comparison of temperature dependence of flow stress for several ’ cobalt alloys versus a nickel alloy and a solid-solution strengthened cobalt alloy [7].

strengthened cobalt alloy, Haynes 188. The strength of the ’-containing cobalt alloys is lower than a traditional nickel superalloy, MarM247, until reaching elevated temperatures. For a simple composition relative to the nickel alloy, the polycrystalline cobalt alloys, specifically those containing tantalum, were encouraging due to the

5 strength of the yield curve above 900 °C relative to the nickel-based MarM247.

Compositional changes to the cobalt alloy should improve this behavior.

Other work has shown that different elements partition to each of the phases differently. It was shown that tantalum was a ’ stabilizer to higher temperatures while chromium had the opposite effect [11-13], but chromium was needed for oxidation resistance. Similar to nickel-base superalloys, boron has been shown to be a grain boundary strengthener [14] and improve ductility during hot working. These functional composition relationships are critical to the future development of these alloys.

1.2.2 Factors Relevant to Thermomechanical Processing

Determining which of the Co-Al-W alloy compositions meet thermal and mechanical demands is only part of the requirement – the alloy still has to be manufactured into a useful shape with properties suitable for the chosen application.

Hence, thermomechanical processing, microstructure, and the relevant properties must be considered.

When considering processability, a few key properties provide insight for workability. Cobalt alloys are said to exhibit a narrow solidification range, meaning the material transitions from all liquid to all solid within a small temperature window. This would indicate large amount of segregation is less likely because solidification is occurring quickly. A large supersolvus hot working window equates to a large range between when the ’ solvus temperature and solidus temperature. A large range is useful because most initial hot working is conducted supersolvus, and maintaining temperature in the workpiece often requires careful monitoring and several reheats during open-die

6 forming operations to reduce risk of cracking due to thermal gradients across large sections of material. Lastly, slow ’ precipitation kinetics suggests the strengthening phase is slow to form which increases workability at intermediate temperatures.

Numerous studies have shown that Co-Al-W alloys may be easier to manufacture than traditional nickel-base superalloys [4, 6, 7, 10, 15-17]. Traditional alloying additions for

’ phase stability and improved mechanical properties in nickel-base superalloys also tend to increase the solidification window and limit the useable processing window [18], but few formal studies on industrially relevant scales have been conducted.

A single effort to hot work a similar composition of polycrystalline Co-Al-W was conducted [15, 17]. Three pilot-scale ingots were melted, a ternary (Co-Al-W), and two quaternaries (Co-Al-W-Ti), one containing boron for grain boundary strengthening and one without. Severe cracking was observed in the ternary alloy on deformation, but the quaternary alloys were successfully hot worked. Only qualitative results were presented without quantitative assessments of homogenization or microstructure evolution as a function of hot-working sequence.

1.3 Thermomechanical Processing of Superalloys

The metallic components in traditional nickel-base superalloy disks are made via several methods including both ingot metallurgy (IM) and powder metallurgy (PM) [1,

5]. For this effort, traditional ingot metallurgy techniques will be investigated as the most affordable manufacturing method for structural disk materials. This process begins by melting a known weight mix of metals or alloys matching the desired chemistry in a

7 vacuum arc furnace and casting the alloy into an electrode for further vacuum arc remelting or electroslag remelting (VIM/VAR/ESR). The electrode is melted and then re-cast into a large chilled copper mold and allowed to cool.

Once cooled, the ingot will then be reheated in a large furnace to stress relieve and homogenize the material. The homogenization heat treatment is done above the gamma-prime solvus in nickel superalloys so all secondary strengthening phases that may have precipitated during slow cooling of the ingot are dissolved and the constituent elements are free to diffuse within the material. The homogenization process is limited by the diffusivity of the alloying elements in the base elements. Faster homogenization is accomplished either by increasing the temperature, and thus diffusivity of the mobile elements, or by increasing homogenization time at the expense of potential grain growth and possible surface reaction. Once the segregation has been reduced sufficiently, the ingot is allowed to cool and the surfaces are cleaned and prepared for ingot-to-billet conversion.

Ingot conversion consists of several processes with the goal of converting the coarse, columnar ingot microstructure with anisotropic properties to a nominally uniform, finer equiax recrystallized microstructure, resulting in isotropic properties and finer grain size to aid in subsequent forming options. To accomplish this, several iterations of upsetting the material, or making it shorter along the length by end-pressing, then drawing the material, or making it longer along the length by side-pressing, are accomplished with intermittent reheating performed as required to maintain sufficient temperature for workability and promote recrystallization. This can also be completed

8 through extrusion. Ingot conversion is performed supersolvus to dissolve the strengthening gamma-prime and encourage recrystallization. Further heat treatment of the converted billet may be necessary to further promote recrystallization and grain growth.

After being converted to wrought billet, forging mults will be cut for blocker and final shape forgings. Final shape forgings are generally conducted sub-solvus as the strengthening gamma-prime phase will limit grain growth during processing. Machining, inspection, and subsequent heat treatments aimed at developing a specific grain size and precipitate strengthening phase size and distribution are performed.

Due to the similarities in the strengthening gamma-prime phase and the wide temperature window for hot-working, thermomechanical processing of a similarly strengthened, similarly structured cobalt-base superalloy should be achievable using a conventional ingot metallurgy processing methodology.

1.4. Computational Modeling Tools

Computational modeling tools are an essential way to investigate key components of the research area and include a range of tools. Deformation process models, microstructure evolution models, and thermodynamic models all played a significant role in the completion of the contained work, where appropriate.

1.4.1 Deformation Process Models

Continuum-level, finite element process models provide component-level information on the process used and how the workpiece responds to thermomechanical

9 inputs. Models can be used in a variety of ways including calculation of strain or strain rate profiles for different geometries of specimens, or to predict die stresses and press loads to make decisions on how material should be processed. Models of this type are usually material agnostic, but require basic material property information to function.

1.4.2 Microstructure Evolution Models

Two types of microstructure models were considered for this work, those that explicitly model the microstructure and its evolution versus those that capture bulk statistics. The empirical Johnson-Mehl-Avrami-Kolmogorov (JMAK) method is a fit to many transformation reactions and can be described phenomenologically in terms of nucleation and growth processes. Early work by Kolmogorov [19], Johnson and Mehl

[20], and Avrami [21] formed the basis for the JMAK model. Books are available which speak more specifically to the kinetics of phase transformations in metals which utilize the JMAK methodology [22]. When applied to processes like recrystallization, JMAK fits are generally very quick to tune and compute due to only tracking the average properties at a specific location. A similar statistical microstructure model would be the so-called “meso-scale” model that quantifies different meso-scale units and tracks the evolution of the different grain categories statistically rather than explicitly modeling each grain and its evolution [23]. Models that explicitly model the microstructure include

Cellular Automata (CA) [24-26] and Monte Carlo (MC) [27, 28] methods. These models work better at capturing specific microstructure and neighborhood effects, but are more computationally intensive and have their own limitations, including translating from simulation steps to real-time.

1.4.3 Thermodynamic Modeling

Calphad-type thermodynamic models are used for prediction of several material properties based on using the concept of equilibrium simulations of Gibbs free energy

[29]. These tools can predict things like phase fractions, precipitate formation, and metallurgically relevant temperatures such as solvus, solidus, liquidus temperatures that are necessary for making assessments on the suitability of the material for thermomechanical processing or determining processing route. Thermodynamic models are extremely material data sensitive and require knowledge of data pedigree and material system and components and extrapolation beyond known ranges is difficult.

1.5 Summary of Opportunities and Objectives of this Research

While a significant body of work had been completed since 2008, there remain opportunities for research in the area of ’ strengthened cobalt-base superalloys. A strong need exists for demonstrating viability of typical ingot metallurgy thermomechanical processing techniques and routes on an industrially relevant scale, including homogenization, ingot conversion, and wrought processing. Demonstration of applicability of industrially relevant microstructure evolution models for this class of alloys is necessary for maturation and eventual transition. Comparison to traditional nickel-base superalloys produced using ingot metallurgy techniques will also be useful in identifying the differences between the alloys and assessing the difficulty expected of industry to transition from a nickel-base to a cobalt-base system production.

This research has several objectives. The first objective is to demonstrate feasibility of typical ingot conversion and representative wrought thermomechanical processing of a novel ’ strengthened cobalt-base superalloy on an industrially relevant scale. The second objective is to develop the data necessary to enable finite element thermomechanical process simulation techniques. The final objective of this research is to demonstrate feasibility of using existing microstructure evolution models for nickel- base superalloys by tailoring them for cobalt-base superalloys to accurately predict the resultant microstructure from a representative thermomechanical processing sequence.

Developing the necessary background information and model data for the thermomechanical processing of cobalt superalloys and subsequently demonstrating the ability to manufacture an industrially relevant product form with a predictable microstructure would be a significant step forward in utilizing ’ cobalt-base superalloys in turbine engine disk hardware and the increased capability it could provide.

To address these objectives, this document is structured such that each discrete thermomechanical processing step—homogenization, ingot conversion, and wrought processing—is discussed separately in-sequence as each depends on the results of the prior processing step. This is followed by the demonstration of an empirical fit JMAK microstructure model using data generated with material converted and processed during the course of this investigation. Conclusions and recommendations finish the document.

Within the document, please note gamma-prime is interchangeable with ’ and gamma prime.

Chapter 2: As-Cast Cobalt Ingot Characterization

The material used in the present research was a 125 kg, 20 cm diameter, 40 cm long, solid as-cast Co-Al-W ingot produced specifically for this effort by a primary mill supplier using industrially-relevant intermediate-scale melt and casting facilities. The composition used in this work was chosen to closely mirror compositions investigated in the open literature [7, 18, 30, 31]. This alloy was used as a model system to demonstrate the feasibility and suitability of typical thermomechanical processing of an L12 ’ cobalt superalloy and the applicability of traditional microstructure evolution models for nickel- based superalloys.

The desired material composition was weighed and melted in an intermediate- scale furnace. The constituents were vacuum induction melted (VIM) and cast into a 10 cm round electrode then vacuum arc remelted (VAR) and cast into a 20 cm round, water- cooled copper mold and allowed to cool. Ladle composition was checked during melting but corrections were not made in-process. As such, some deviation from target chemistry was expected. Bulk chemical composition was checked independently upon receipt of the material from the supplier (Table 1). Compositions during and after casting were within the expected composition range to show the desired meta-stable ’ L12 precipitate phase. During the later stages of melting and subsequent casting after the ladle

13 composition was checked, some lighter alloying element losses were observed for Al, Ta, and B.

At% Al W Ta B Co Target 9.0 9.0 2.0 trace bal Ladle 9.3 8.6 2.1 0.02 bal Final 8.4 8.9 1.9 - bal Table 1. Composition target and measurements at several stages of production (at %)

The ingot material was provided as-cast from the supplier with none of the traditional subsequent thermomechanical processing having been performed so they could be performed as part of this investigation . This included homogenization and ingot conversion steps for chemistry and microstructure refinement and control that would typically be conducted at the supplier. Prior to undergoing subsequent thermomechanical processing, a thorough characterization of the as-cast material was performed to baseline the microstructure for comparative analysis.

2.1 Ingot Sectioning and Naming Conventions

The cobalt ingot was cut using wire electrical discharge machining (wire-EDM) into four major sections (Figure 4). These sections: T (Top), B (Bottom), M

(Microscopy), and E (Extrusion), were tailored to provide material for subsequent characterization and thermomechanical processing evaluations. A longitudinal reference notch was wire-EDM on the ingot as a rotational reference and to provide stability for subsequent machining operations.

Figure 4. Assembled ingot sections, from left to right T, E, M, and B.

As is standard industry practice, the surface of the ingot was mechanically ground after casting to remove surface defects, but some surface-connected porosity from the casting remained (Figure 5). The grinding resulted in a slight concavity or hour-glass shape to the final delivered ingot. The remaining surface porosity, generally less than 5 mm in depth, was not a concern in the areas it was still present for the purposes of this effort. Locations for thermomechanical testing and extrusion specimens were chosen to minimize overlap with those areas to reduce risk of contacting the surface and capturing surface-connected defects.

Figure 5. Detail-view of section E with grinding marks and remaining surface defects.

2.2 Macro Section Characterization

The Microscopy section of the cobalt ingot was wire-EDM longitudinally to make several smaller pieces suitable for characterization. A half-ingot D-shaped section of the

M-slice was used for the as-cast characterization. The remaining half was saved for later homogenization and subsequent microscopy (Figure 6).

Figure 6. Cut up diagram for Microscopy slice with the as-cast samples marked and cut.

Considerable residual stress remained in the ingot material from the differential thermal profiles during ingot casting. This was expected as the material was not stress- relieved or homogenized while being produced. However, the residual stress state resulted in challenges specifically related to cracking and distortion while machining and subsequent heat treatment. Precautions were taken to mitigate cracking during both operations but cracking remained a concern due to the potential loss of usable material for the experiments. Significant cracking was observed during wire-EDM of the initial ingot sections for homogenization (Figure 7). During heat treatment, conservative

17 heating rates were used to limit thermal cracking by reducing thermal gradient between the surface and interior of the coupons or sections being heat treated.

Figure 7. Observed crack propagating from the cut surface during machining in the as- cast ingot material.

2.2.1 Optical Macrographs

Samples were excised from the as-cast Microscopy section to produce radial sections both longitudinal and transverse to the ingot axis. These samples were prepared for optical microscopy using standard metallographic preparation techniques using progressively finer grit papers, diamond pastes, and vibratory polishers with polishing solutions. To provide contrast, the prepared samples were chemically etched by 18 submersion for approximately 15 seconds in a modified Krolls (Macro Krolls) or modified Cobalt (Macro Cobalt) etchant (Table 2). The ‘macro’ recipes contained extra acid relative to the original recipes swapping water for acid while keeping the total volume constant for increased potency due to the known corrosion resistance of cobalt.

The prepared specimens were imaged using an inverted light microscope equipped with digital camera and automated, mechanized sample stage for sequential image capture.

The images were assembled using built-in software to produce a macro-section of the entire cross section piece.

Macro Krolls Macro Cobalt Water 200 mL 50 mL Hydrofluoric Acid 16 mL - Nitric Acid 48 mL 50 mL Hydrochloric Acid - 100 mL Table 2. Etchant recipe for Macro Krolls and Macro Cobalt.

The radial section, displaying the longitudinal axis of the ingot (Figure 8), showed the typical columnar dendritic ingot solidification texture (Figure 9) with the strongest thermal gradient occurring from outer diameter of the ingot section (right) to the center

(left) and from bottom to top. The thermal gradient observed during solidification caused a gradual and local shift in dendrite orientation as the bottom and sides of the ingot cool while the center and top remained hot during casting. While the orientation angles differ radially by becoming more vertical approaching the center, similar inclination angles were observed along the length of the ingot in the steady state region for a given radius.

10 mm

Figure 8. Macrograph of diametral section, theta orientation, with center to left and OD on the right.

Similar variation in grain orientation and structure was observed in cast nickel- base superalloy ingots [32]. Samples for thermomechanical processing, including extrusion and cylindrical upset specimens, were excised from the steady state region of the cast ingot, Section E, to reduce the variability in the mechanical properties due to significant microstructural changes from test coupon to test coupon.

Figure 9. Typical grain structures in solidified castings of ingots [33].

For etched and optically imaged specimens, lighter areas indicated the dendrite cores while the interdendritic regions appeared much darker. If sectioned properly, the

20 dendrites (lighter areas) would be seen as long, continuous bands. The visible hatching was a result slight mis-alignment during sectioning along the primary solidification direction. A higher magnification image (Figure 10) shows the longer, primary solidification direction of the dendrites and the much shorter secondary arm spacing orthogonal to the primary growth direction. At higher magnification, a secondary, discontinuous particulate phase appears in the interdendritic regions as a bright white dispersion and was characterized further in §2.3.2.

200 m

Figure 10. Primary dendrites and secondary arms seen in an as-cast Co-Al-W alloy.

Similarly, the transverse section showed strong segregation and expected dendritic structure from ingot casting (Figure 11). These samples exhibited some surface cracking due to high residual stresses remaining in the ingot after casting without stress 21 relief. The roughly round ends of grains can be seen from about the three-quarters radius to the outer diameter despite the strong etch used for dendrites. The as-cast grain size was smaller than expected, but was likely due to the interdendritic particles pinning grain boundaries and will be quantified in §2.3.1.

20 mm

Figure 11. Transverse section of as-cast cobalt with center on the left, OD on the right.

Groups of similarly oriented dendrites in the transverse sections appeared as plusses and defined a single grain. The dendrites were more easily seen at higher magnification (Figure 12), while grains were more easily identified using orientation mapping tools such as electron back-scattered diffraction (EBSD). As expected, and demonstrated in the longitudinal section, the orientation of the dendrites followed the same shift from center to outer diameter. Again, this changed the relative appearance of

22 the dendrites to more plus-like nearer the center of the ingot where the section was more orthogonal to the dendrite growth direction than at the outer diameter.

Figure 12. Magnified view of dendrites of the as-cast ingot.

2.2.2 Primary Solidification and Growth Direction Characterization

Quantification of the primary growth direction was necessary for accurately conducting several measurements of both the as-cast and the eventual cast-and- homogenized material. These measurements included grain diameter, primary dendrite arm spacing, and composition gradients between dendrite core and interdendritic regions.

More specifically, the angle of growth as a function of radial position in the ingot was needed to compensate for the varying angle of the dendrites. This compensation allowed proper measurement of the desired characteristics independent of location in the ingot and for the results to be interpreted and used meaningfully.

An optical macrograph (Figure 8) was used as a representative cross-section of the ingot. Due to radial symmetry, samples taken from a radial location should have a similar microstructure as the corresponding area in the macrograph. Lines were made tracing the primary dendrite growth direction in several areas on the macrograph of the ingot (Figure 13). The angle indicated in the figure is the deviation from horizontal

Figure 13. Lines tracing dendrite growth direction in the ingot.

24 measured counter-clockwise from each line. The necessary measurements will need to be made on sectioning planes orthogonal to the growth direction, so orange dashed lines were drawn orthogonal to the solid blue dendrite growth directions (Figure 14). The angles for the required sectioning were measured clockwise to horizontal for each line.

The determined angle is valid for any section along a continuous dendrite or fiber. This method was useful for determining the proper sectioning angle of the homogenization trial specimens (Figure 15) to help ensure the required end-on view of the dendrites was achieved and measurements were conducted in the correct plane.

Figure 14. Orthogonal sectioning lines added for each location

Figure 15. Overlay of sectioning angle for metallographic analysis of heat treat specimens.

2.3 As-Cast Microstructure Characterization

Several features were of interest when quantifying the as-cast microstructure of the cobalt ingot material, including grain diameter, volume fraction and morphology of - prime, primary dendrite arm spacing, and particle phase identification. Each of these microstructure features impacted the required thermomechanical processing during this effort and would impact subsequent service properties of the resultant material.

2.3.1 As-Cast Texture and Grain Diameter

As-cast grain diameter, or grain size, was determined using Electron Back Scatter

Diffraction (EBSD) methods. Using the location and corresponding sectioning angle

26 information from §2.2.2, a sample was excised, sectioned, mounted in conductive mounting material, and polished using standard metallographic techniques for electron microscopy. The sample was imaged using a standard scanning electron microscope

(SEM) equipped with an EBSD detector and acquisition software [34]. There was no material file for cobalt in the TSL software so nickel, which is also face centered cubic, was the material selected for analysis. Due to the large area of interest, custom batch scanning software was used to couple beam and stage controls to acquire approximately

2000 scans of a standard field of view at 350x magnification, using a 5 micron step size.

Data files for the individual scans were combined using a semi-automated custom software [35] script. A representative EBSD Inverse Pole Figure (IPF) was calculated

(Figure 16a). Unindexed points, or those the software was unable to index properly, showed up as small clumps of either rainbow colors or black in otherwise solid-colored regions. Additionally, slight vertical color gradients were observed due to small changes in calculated orientations from incident beam deflection due to the low magnification used for the scans. Basic clean-up algorithms within the EBSD analysis software were used to reduce noise in the data from unindexed points due to grain boundaries, non-FCC interdendritic particles, or other defects in the material such as cracks or surface contaminants. Results were compared before and after cleanup (Figure 16b) to determine minimally invasive cleanup procedures. Care was taken to remove mis-indexed points but not substantively alter the overall microstructure. The sample was sectioned such that an [001] IPF should show red grains, meaning a surface of the FCC crystal was pointed normal to the sample surface in the primary growth direction.

Figure 16. As-cast IPF map of raw data as-captured (a) and after basic cleanup (b).

An [001] Pole Figure (PF) was calculated using a standard harmonic series expansion and inversion symmetry and plotted (Figure 17). A strong [001] pole was 28 observed which indicated a majority of the crystals were aligned in that direction, accompanied by a circle around the periphery illustrating rotational crystal symmetry and was consistent for a columnar fiber texture of an as-cast ingot. The strength of the pole also confirmed the strength of the fiber texture present in the columnar region of the ingot. Together these indicated the crystals were aligned properly in the growth direction from the thermal gradient but allowed to rotate freely around that axis. The slight misalignment of the central crystal axis was reasonable when considering error from marking, sectioning, mounting and polishing, and alignment when mounting in the electron microscope. Correcting the IPF to account for the observed alignment errors resulted in the expected red [001] IPF map (Figure 18).

Figure 17. Calculated [001] PF demonstrating stack-up mis-alignment.

Figure 18. Corrected IPF map of as-cast cobalt with the expected [001] fiber texture.

For materials with large grains, an optical micrograph of a lapped or polished and etched sample surface is sufficient for determining grain size using standard methods such as ASTM E112 [36]. While Macro Krolls and Macro Cobalt were suited to revealing the dendritic structure, they were not well-suited to highlight grain boundaries.

Other attempts to determine a grain boundary etch were unsuccessful and resulted in severe over-etching of the dendritic structure and would not have been suitable for grain size measurements of the cast ingot material. Instead, EBSD IPF and grain identification maps were used to determine the grain size. For the purposes of automating the image analysis, point-to-point orientation changes in the material greater than 3 degrees were considered grain boundaries. This approach was deemed a reasonable compromise between capturing boundaries and not artificially creating new ones. The grain

30 boundaries present in the as-cast cobalt were pinned and highly tortuous (Figure 18) making traditional measurement of grain size difficult. An equivalent diameter circle method was used to calculate average grain size using ASTM E1382 [37]. The mean as- cast grain size was determined to be 500 m with a significant standard deviation. Grain sizes ranging from tens of microns to several millimeters in diameter were observed, with the largest area fraction occurring in the range of 1-2 mm diameter (Figure 19). The mean grain length was not able to be determined from this section transverse to the growth direction. Grains were substantially longer in the solidification direction due to the columnar solidification texture, and extended through and beyond the 25 mm height of the Microscopy section of the ingot used for characterization. Similar elongations of grains were documented in as-cast Waspaloy [38].

Figure 19. Grain ID map (left) with area fraction versus grain diameter (right).

2.3.2 Interdendritic Phase Identification

The chosen composition of Co-Al-W -prime superalloy resulted in a hard, intermetallic phase from the melting, casting, and solidification of the original ingot. The

31 distribution of these particles was not randomized and were clearly concentrated within the interdendritic regions of the as-cast ingot (Figure 10, Figure 12). The overall area fraction of particles from top-to-bottom and center-to-edge was consistently between 1-3 percent depending on the field of view. The exact size and morphology of the particles varied but the average as-cast diameter was approximately 25 microns.

Chemical mapping using Electron Dispersive Spectroscopy (EDS) was conducted to determine the local segregation of alloying elements between the matrix and the interspersed particles (Figure 20). The back-scattered electron (BSE) image showed the overall structure and distribution of the particles, which display as light grey or white due to their density. Each map below the BSE image showed the relative distribution of the named element within the field of view, with brighter intensity colors representing more of the element. The elemental maps showed the particles were lower in both cobalt and aluminum than the bulk matrix. Tungsten was similar, if not slightly higher, in the particles than the bulk, but tantalum was strongly segregated to the interdendritic region and the particles.

Figure 20. EDS composition maps for as-cast cobalt superalloy. 33

Calphad-type thermodynamic modeling [39] was completed on the nominal alloy composition to gain insight to the phase formation within the material during solidification. The equilibrium solidification projection indicated no additional phases were precipitated other than the FCC matrix. By using a non-equilibrium Scheil approach, the solidification projection predicted the formation of a deleterious intermetallic and laves phase that were among the last to solidify at approximately 1275

°C (Figure 21) and resulted from non-equilibrium solidification conditions. The calculation predicted a similar fraction these undesirable phases as was observed in the experimental material at approximately 5% of volume. Given this, the particles were assumed thermally stable eutectic-type solidification remnants and unresponsive to typical heat treatments used for homogenization of compositional gradients, similar to those seen in other investigations of similar compositions [17]. Later homogenization heat treatments confirmed that the particles remained stable even after a homogenization of 240 hours at 1232 °C.

Figure 21. Solidification projection for Co-9Al-9W-2Ta.

EBSD of the intermetallic particles indicated a nominally hexagonal crystal structure and as such was able to be indexed using a hexagonal material pattern. Actively partitioning EBSD by indexing both FCC-cobalt and the hexagonal interdendritic material while performing the data collection allowed the phase to be easily differentiated from the standard gamma/gamma-prime matrix during subsequent analyses.

2.3.3 Gamma-Prime Characterization

The structure and area fraction of the L12 gamma-prime phase of this Co-Al-W alloy were determined using backscatter imaging. Samples were aged at 850 °C and quenched to preserve the gamma-prime microstructure and prepared using standard 35 metallographic preparation techniques and imaged on a scanning electron microscope equipped with backscatter electron imaging detector. Image processing and thresholding techniques were used on the image to determine the overall area fraction of gamma prime and agreed with manual point counting methods. The area fraction of gamma-prime was high when compared to nickel superalloys at 80% (Figure 22) but agreed with literature for similar cobalt compositions [4, 7, 10, 11, 40]. This was also consistent given the content of typical gamma prime formers in the alloy.

Figure 22. Backscatter electron micrograph showing high area fraction of gamma prime

Higher magnification imaging (Figure 23) confirmed the gamma-prime phase morphology was cuboidal. The cuboidal precipitates were an indicator of the high area 36 fraction and are prevalent for materials containing gamma-prime formers such as titanium and aluminum. The cuboidal precipitates also indicated a relative positive lattice-particle misfit of 0.5-1%. The average size of the gamma-prime phase particles was 125 nm. The alignment of the gamma-prime particles was regular and linear with small 15-20 nm gaps between the individual gamma-prime within a row and 50-60 nm thick gamma channels separating the rows and should exhibit good rafting behavior..

Figure 23. Backscatter electron image for determining gamma prime size and morphology.

2.4 Thermophysical Property Data

Thermophysical property data for the cobalt alloy was necessary for a variety of reasons from defining processing windows to implementing the alloy in a material database for finite element modeling simulations. As-cast and homogenized material was sent out to a vendor to obtain the necessary thermophysical property data. Properties reported included density, specific heat, thermal diffusivity, and thermal conductivity

(Table 3).

Specific Thermal Temperature Density Heat Diffusivity Conductivity (°C) (g/cm3) (Wsec/gK) (cm2/sec) (W/cmK) 23 10.2416 0.3668 0.03733 0.14024 50 10.2416 0.3731 0.03839 0.1467 100 10.2416 0.3838 0.03997 0.15712 200 10.2416 0.4022 0.04199 0.17299 300 10.2416 0.4177 0.04346 0.18596 400 10.2416 0.4314 0.04466 0.19731 500 10.2416 0.4441 0.04584 0.2085 600 10.2416 0.4552 0.04657 0.21714 700 10.2416 0.4625 0.04724 0.22379 800 10.2416 0.4729 0.04808 0.23287 900 10.2416 0.4806 0.04854 0.23896 1000 10.2416 0.488 0.04649 0.23235 1100 10.2416 0.492 0.04326 0.21798 1200 10.2416 0.494 0.05143 0.2602 1240 10.2416 0.495 0.05184 0.2628 Table 3. Summary of thermophysical properties of a -prime cobalt superalloy

Gamma-prime solvus temperature was determined experimentally by interrogating experimental data for specific heat capacity as a function of temperature.

The experimental data for specific heat on heating and cooling was plotted to show evolution with temperature (Figure 24). During heating, the alloy requires temperature 38 beyond the gamma-prime solvus (over-heating) for the gamma prime to fully dissolve.

Similarly, during cooling, some amount of under-cooling was needed to thermodynamically drive the onset of precipitation of the -prime particles. The effective gamma-prime solvus temperature sits between the temperature where precipitation starts on cooling and when dissolution finishes on heating. Using this method, the gamma prime solvus temperature was estimated at Tsolvus = 1080 °C.

1.2

1.1 C) ° 1.0 Cooling 0.9 Heating

0.8

0.7 Dissolution Finished

0.6 Specific Heat (Wsec/g Heat Specific

0.5 Precipitation Started 0.4 600 700 800 900 1000 1100 1200 1300 Temperature (°C)

Figure 24. Specific heat traces for determination of precipitation behavior on heat up (dashed) and cool down (solid) lines. Dotted vertical line represents Tsolvus.

A traditional ’ solvus approach curve was completed for this alloy by heating coupons to a range of temperatures near the expected solvus temperature and allowing samples to soak to equilibrate the ’-volume fraction prior to water quenching to preserve the microstructure. After metallography, the results from the solvus approach curve compared favorably with the specific heat data showing complete dissolution of the gamma prime above 1080 °C but not below 1040 °C (Figure 25). An additional intermediate temperature at or near 1070 °C would have been illustrative but was not completed.

Figure 25. Experimental -prime approach curve (dashed line).

Calphad-based modeling tools [39] were also used to estimate the -prime solvus temperature for the composition used in the present work. The -prime solvus temperature was estimated at 1104 °C (Figure 26) using this method. This temperature was higher than the experimental -prime approach curve and specific heat data. The software also under-predicted the volume fraction of gamma-prime as a function of temperature. The equilibrium phase map did not predict the interdendritic phase that remained after casting. Several improvements were recently made to the cobalt material database used in these calculations and results were much more aligned with experimental measurements.

Figure 26. Equilibrium phase fraction predictions for Co-9Al-9W-2Ta

Given agreement between several experimental methods of determination and the proximity of the computational methods, Tsolvus = 1080 °C was used as the effective gamma-prime solvus temperature for this composition and research work. By comparison, this was approximately 60 °C higher than the estimated gamma-prime solvus temperature of Waspaloy [41] and indicates the ability of the material to retain strength at higher temperatures than a traditional gamma-prime strengthened nickel superalloy. This temperature was the basis for identifying and conducting subsequent thermomechanical processes.

Chapter 3: Cobalt Ingot Homogenization

Homogenization is an elevated temperature heat treatment of the material intended to reduce the residual compositional segregation between major constitutive elements of the alloy that naturally occur during ingot melting, casting, and solidification.

The heat treatment necessary is dependent on a number of factors including the composition and diffusivities of those elements in the base material, the required diffusion distance, amount of segregation observed and degree of homogenization desired, and the solvus temperature for any precipitated strengthening phase. The necessary time to sufficiently reduce the residual segregation from casting was determined using a combination of both analytical and experimental methods.

3.1 Analytical Homogenization Time

Bulk interdiffusion rates vary by element in cobalt, and correlates to the atomic diameter of the diffusing atom relative to the bulk cobalt. Diffusion of tungsten in cobalt is slow [42-44] and was the rate-limiting element in this alloy. The homogenization time, t, is determined by the interdiffusion rate of tungsten in cobalt, D, and the required distance, which is related to the Primary Dendrite Arm Spacing (PDAS) (Equation 1).

1 푃퐷퐴푆 2 푡 = ( ) (1) 2퐷 2 43

The inter-diffusion coefficient is dependent on temperature and compositional gradient of the diffusion couple used in the measurements.

Samples from smaller ingots were initially homogenization heat treated to help baseline the homogenization required as part of another effort. The PDAS of prior-cast 8 cm round ingots was 100 m. Using Equation 1, at 1232 °C the smaller ingots were predicted to take 90 hours to homogenize, but 48 hours was demonstrated sufficient time experimentally for the heavier elements to diffuse the necessary distance in those smaller experimental ingots [45]. However, for the larger 20 cm diameter ingot used in this work, the PDAS was larger at 200 m due to changes in thermal conditions relative to the smaller ingots. The larger PDAS required a much longer homogenization treatment.

Bulk interdiffusion of the tungsten over that distance would require over 720 hours at

1180 °C (Figure 27). During homogenization, the compositional differences between the regions decreases, resulting in a decrease of the effective interdiffusion coefficient/driving force for diffusion, resulting in significantly longer homogenization times. While 720 hours is high, it represents the time required for complete homogenization, or no compositional gradient from dendrite core to interdendritic region.

This level of homogenization is rarely the case in practice as some residual compositional gradient is expected and acceptable as long as the degree of the compositional differences do not significantly impact the microstructure such as precipitated phases or mechanical response. Even considering the approximate factor of 2x observed with the smaller ingots to reduce the time from 720 to 360 hours, any homogenization treatment over a few days, much less 15 days, is an impracticable number for commercial production.

800 700 600 500 1180 °C 400 300

Time (hours) Time 200 1232 °C 100 0 0 50 100 150 200 250 PDAS (m)

Figure 27. Impact of temperature on calculated homogenization time.

The thermophysical data and -prime approach curve data (§ 2.4) indicated the - prime solvus was approximately 1080 °C. A decision was made to increase the homogenization temperature further over the -prime solvus to 1232 °C to increase the rate of interdiffusion of the elements. Thermodynamic modeling [39] indicated this temperature would still produce the desired single FCC-phase within the material for homogenization (Figure 28).

Figure 28. Predicted phases during homogenization.

Increasing temperature decreased the time required for homogenization. A modest 50 °C increase in homogenization temperature effectively doubled the interdiffusion coefficient but would still result in an unacceptably long homogenization time of almost 350 hours. Even halving this estimate, as was needed for the smaller ingots, still suggested 175 hours would be required. The possible ranges for homogenization times were determined by examining how the interdiffusion coefficients,

D, were calculated using the Arhennius relationship (Equation 2) and the expected

−푄 퐷̅ = 퐷 푒푥푝 ( ) (2) 0 푅푇

46 variability of the input data used [43], where D0 is the frequency factor, Q is the activation energy, R is the gas constant, and T is the diffusion temperature. The range of times from variability in calculated interdiffusion coefficient was plotted (Figure 29).

While the averages followed the expected trend of decreasing time with increasing temperature, the variability showed significant room for interpretation. The dashed lines to the left represented the time predicted using the using the lower bound of the frequency factor and higher bound of activation energy resulting in the lowest possible interdiffusivity coefficient. Conversely, the dashed lines to the right represented the highest interdiffusivity coefficients, obtained by using the upper bound of the frequency factor and the lower bound of the activation energy. The solid lines in the center were the times calculated from the stated averages.

800 700 Low D 600 500 Average D 400 300 High D

Time (hours) Time 200 100 0 0 50 100 150 200 250 1180 °C PDAS (m)

800 700 600 Low D 500 400 Average D 300

Time (hours) Time 200 High D 100 0 0 50 100 150 200 250 1232 °C PDAS (m)

Figure 29. Estimated homogenization time including variability in calculating the effective interdiffusivity coefficient, D, at 1180 °C (top) and 1232 °C (bottom).

While the literature data was useful for estimating purposes, it was the large variability in expected times, and the significant overlap of the temperatures under consideration, that necessitated an experimental investigation specific to the alloy being

48 considered. Prior to homogenizing the entire ingot, initial exploratory heat treatments were completed on smaller coupons of material over a wide range of times to experimentally determine an appropriate homogenization time at the higher temperature,

1232 °C. Once determined, the required ingot material was heat treated using the same sequence to sufficiently homogenize the structure and prepare for ingot conversion.

3.2 Experimental Homogenization Procedures

A multi-staged approach was employed to determine the homogenization heat treatment necessary for the cast cobalt ingot. Initial homogenization quantification work was completed using smaller, subscale coupon specimens cut from the Microscopy section of the ingot. After determining the homogenization heat treatment necessary from the coupon specimens, homogenization was completed on larger sections of the full-size ingot and verified using optical microscopy prior to undertaking any additional thermomechanical processing.

3.2.1 Ingot Coupon Homogenization Procedures

An initial ingot material trial homogenization was performed in air at 1232°C for

48 hours on a roughly 1 cm cube specimen to determine the extent of homogenization and the amount of surface oxide formation at homogenization temperature for this alloy.

Additional samples were cut from a radial section of the ingot for temperature- and environment-controlled and instrumented heat treatments (Figure 30). Due to the expected reactivity of the material with oxygen at homogenization temperature, steps were taken to protect the roughly 1 cm cubes used for the more detailed analysis.

Figure 30 - As-sectioned specimens for homogenization analysis.

Each piece was encapsulated separately in a tubular quartz glass vessel with a piece of tantalum as an oxygen-getter, then evacuated and backfilled with argon prior to being sealed (Figure 31). Using the ideal gas law at the homogenization temperature, the backpressure needed to avoid bursting the vessel while at temperature was determined at

-24psi. The sealed quartz tubes containing the specimens were cleaned to remove any surface contamination from handling and reduce the risk of cracking of the quartz vessel

Figure 31. Encapsulated cobalt sample (left) with tantalum foil (right).

50 during heat treatment. Six samples were homogenization heat treated in an instrumented, at-temperature box furnace for times varying from 48 hours to 10 days and water quenched to freeze the resultant microstructure (Table 4).

Sample Number Temp (°C ) Time (hours) Quench M1397-M-A-5 1232 48 Water Quench M1397-M-A-7 1232 96 Water Quench M1397-M-A-8 1232 192 Water Quench M1397-M-A-6 1232 240 Water Quench Table 4. Homogenization heat treatment conditions.

The samples were sectioned orthogonal to the dendrite growth direction for each sample per the angles measured in §2.2.2. Both adjacent faces from each specimen were mounted and polished for optical microscopy and EBSD. The samples for optical microscopy were etched as before using MacroKrolls while the samples for electron microscopy including EBSD were left as-polished, cleaned in an ultrasonic cleaner to remove surface contaminants, and then demagnetized to reduce distortion from the magnetic tendency of the material.

Optical microscopy was conducted on an inverted light metallograph equipped with an automated stage and batch-stitching software. EBSD was conducted on a scanning electron microscope equipped with an EBSD detector and acquisition software for large scans.

3.2.2 Ingot Coupon Homogenization Results

The ingot coupon experiments were successful in demonstrating the reactivity with air and homogenization behavior of the gamma-prime containing cobalt alloy. The results are split by test and analysis method. 51

3.2.2.1 In-Air Homogenization Results

The heat treatments conducted with the sample exposed to air showed significant reactivity and oxidation. After 48 hours at 1232°C, the test sample had ballooned significantly due to surface oxidation. EDS showed this to be aluminum oxide, or alumina, formed during heat treatment. Removal of the surface oxide showed the specimen lost approximately 50% of its mass to non-adherent oxide formation (Figure

32). While the high percentage can be largely attributed to a surface-area to volume effect that would not be the case when homogenizing a large ingot section, it was surprising to see the extent of the oxide formation given the short time at temperature.

Early in production of nickel-base superalloy ingots, it is common for vendors to over- size their raw materials so that surface oxides or other contaminants can be removed and the ingot will still meet dimensional requirements for delivery. The amount of oxidation after a relatively short exposure time factors negatively against the use of cobalt due to the special processing and handling precautions required to limit exposure to oxygen at temperature. As a result, subsequent testing and heat treatments were conducted in inert environments either by encapsulation, canning, or through utilization of low-reactivity environmental chambers.

Figure 32. Cobalt ingot samples after 1232 °C heat treatment in air for times of 0, 16, 48, and 96 hours (left to right).

3.2.2.2 Optical Microscopy Results

Unlike the exposed-air homogenization sample, the homogenization samples for optical and post processing were encapsulated and backfilled with argon to prevent oxidation. Efforts to reduce oxidation of the cobalt were successful. The encapsulated samples experienced minor, consistent surface oxidation (Figure 33) instead of the significant bulk material oxidation and loss observed during heat treatment in air.

Figure 33. Typical minor surface oxidation of encapsulated samples as a function of time.

The as-cast dendrites and segregation were easily observed using optical microscopy after standard metallographic sample preparation techniques and chemical etching (Figure 34). As expected, the optical micrographs show decreasing contrast from the residual segregation as homogenization time was increased. The dendrites were clearly visible in both the as-cast material and after homogenization of 48 hours at 1232

°C. After 48 hours, however, the dendrites were much more diffuse in appearance. After

96 hours, the prior dendritic structure was no longer visible optically. This homogenization time was one quarter of the 360-720 hours originally estimated and was selected as the experimentally confirmed time to homogenize the as-cast material.

Homogenization time beyond 96 hours appeared unnecessary as there was little change in the observable microstructure and what remained of the dendrites past that time. After 240 hours, there was a visible denuded zone extending inward approximately

750 m from the surfaces of the sample resulting in a lighter appearance when etched than the bulk material. This lighter region caused difficulties optical imaging due to brightness and lack of contrast. To compensate, the exposure was automatically adjusted and resulted in the center of the 240 hour sample imaging significantly darker than the rest of the heat treated samples despite appearing similar overall by-eye to the other samples. A homogenization time of 96 hours, or 4 days, was a significant amount of time for heat treatment but would not be dismissed if the advantages warranted the cost.

Figure 34. Changes in microstructure after heat treatment time at 1232°C 55

3.2.2.3 EBSD Results

Grain size measurements were calculated from EBSD data collected from each of the heat-treated specimens. Grain growth was not significant during the homogenization heat treatment (Figure 35). Using the EBSD specification for grain size [36] on large area scans, the average as-cast grain size was 505 m. After 240 hours at 1232 °C the grain size was only slightly larger at 610 m per specification.

Figure 35. Grain maps for (a) the as-cast ingot and (b) following 240 hours at 1232 °C.

Such minimal growth over a long time indicated the as-cast grains were nearing a stable, Zener pinned grain size, DZener, where dp is the diameter of the pinning particles and Fv is the area fraction of the pinning phase. For the nominal area fraction of particles of 1-3% observed in the ingot material and the average size of 25 m, the effective

Zener-pinned stable grain size was given by Equation 333:

2푑푝 퐷푍푒푛푒푟 = (3) 3퐹푣 This produced a pinned grain size that ranged from ranged from 550-1600 m but was highly sensitive to the characteristics of the pinning phase. The inhomogeneous

56 distribution from casting and solidification of the intermetallic pinning phase also contributed to variability in what would be expected from a stable pinned grain size. As such, excessive growth for shorter even multi-day homogenization heat treatments was not a concern.

Segregation was reflected in the EBSD Image Quality (IQ) maps as the absorption and reflectivity of the incident beam changed with composition traversing from dendrite core through the interdendritic region to adjacent dendrite core. Clusters of the cores representing grains were visible in the as-cast IQ map (Figure 36a). After 48 hours the remnant dendritic-scale segregation was significantly reduced such that the IQ map no longer showed an observable difference between the dendrite cores and interdendritic regions (Figure 36b). The tortuous grain boundary character for this alloy stayed similar even after homogenization heat treatment.

Figure 36. IQ maps of as-cast cobalt (a) and cobalt after 48 hours homogenization (b).

3.3 Experimental Ingot Material Homogenization

Having determined the necessary amount of time for the required homogenization to occur analytically and experimentally on small coupons, full sections of ingot material

57 were to be homogenized and stress-relieved in one operation to generate material for subsequent thermomechanical processing, characterization, and analysis. Unfortunately, in-house vacuum heat treatment for larger – and very heavy – metallic samples was non- operational and local vacuum furnace capabilities were prohibitively expensive considering the multiple days of heat treatment. Alternatives to these approaches were investigated. The decision was made to avoid heat treating them in air due to the undesirable severe oxidation behavior (Figure 32) coupled with the lengthy 96-hour exposure. Considering the large diameter and the significant weight of the cobalt ingot, the best solution of those available was to perform the homogenization using an on-site

Hot Isostatic Press (HIP) unit using a non-traditional HIP cycle.

3.3.1 Ingot Homogenization Procedures

A traditional HIP treatment involves highly-pressurized inert gas for a short time at temperature and is intended to force interior pores and cracks to close in a material.

This was unnecessary for this alloy because there was no evidence of internal porosity, and surface-connected porosity is unchanged by high pressure HIP. Instead, a low- pressure HIP cycle was used to maintain inert atmosphere around the material with the intent to purge the oxygen in the chamber, effectively reducing oxidation behavior over a much longer duration than would typically be used in a HIP. The HIP pressure valve control settings were modified to limit the pressure venting from the coarse purge valve in favor of the fine pressure venting valve. This change reduced the cycling of the argon venting system which in-turn reduced argon usage and limited oxygen reintroduction

58 through the coarse venting valve preventing additional oxidation of the cobalt material at such long times at temperature.

The limited size of the heated zone of the HIP necessitated sectioning the ingot into four, 90-degree pie-shaped pieces (Figure 37). One extrusion quarter-pie section and one microscopy section were stacked vertically using alumina spacers between the

Figure 37. Typical Extrusion Section ingot pie shaped piece for heat treatment.

graphite base-plate and cobalt pieces and heat treated together per run (Figure 38).

Temperatures were monitored real-time using three thermocouples in the low-, mid-, and upper-zone of the furnace with the mid-height used for control. The furnace was cleaned and refurbished after each run to remove any residual contaminants and ensure viability

59 of the required seals, heating elements, and thermocouples. No reactive foils or getters were used for the runs either surrounding or in crucibles within the vessel during HIP.

Figure 38. Cobalt material stacked in HIP heating element assembly and carrier.

The entire E-section of the ingot was homogenized along with half of the M- section in four batches using the same HIP recipe (Table 5) following intermediate temperature vacuum purge to remove contaminants from the chamber.

Step # Operation Details 1 Heat Up from 20°C at 2°C/min 2 Hold at 1232°C for 96 hours 3 Cool down at 2°C/min Table 5. HIP cycle used for extrusion and microscopy materials.

Heating or cooling the ingot too quickly could result in undesirable cracking due to differences in temperature between the outside and inside of the ingot section from thermal expansion. Having not received a stress relief at the mill will make propensity for cracking worse. A material database was generated for this cobalt composition using the acquired thermophysical property data (§2.4). A finite element heat treatment simulation was performed [46] to predict the maximum temperature deviation between center and outer surface of the ingot during heat-up and cool-down that could lead to cracking. The 2°C per minute was chosen as a reasonable compromise between time required and risk of cracking the cobalt pieces in the furnace.

Following the HIP homogenization cycle, metallographic specimens in both the as-cast and cast plus homogenized conditions were excised and prepared identically using standard metallographic procedures and etched using Macro Krolls. After etching, these samples were imaged using an optical microscope and compared to visually confirm the expected reduction in segregation.

3.3.2 Ingot Homogenization Results

The entire Extrusion-section of the ingot was successfully homogenized in four pieces along with half of the Microscopy-section. A light surface layer of oxidation was observed with some surface greying and chalkiness that was not present prior to homogenization. To evaluate whether the ingot material using the HIP cycle had effectively seen the same homogenization heat treatment as the experimental coupons utilizing a traditional furnace setup, representative areas of the as-cast microstructure

(Figure 39) and the cast-plus-homogenized microstructure (Figure 40) were imaged and

61 compared. While the as-cast microstructure exhibited dendritic and interdendritic structures and severe gradients in etching from compositional segregation, the cast plus homogenized sample lacked those features and was much more even in etching at the grain level. The interdendritic phase remained stable at the 1232°C homogenization temperature as the relative amount, dispersion, and morphology appeared similar between the as-cast and homogenized material even within the ingot. With these results, there was strong confidence the segregation had been reduced sufficiently to proceed with ingot conversion.

Figure 39. Representative etched microstructure from the as-cast cobalt ingot.

Figure 40. Representative etched microstructure of the cast plus homogenized ingot.

Concurrent with the homogenization, stress relief of the ingot sections was also accomplished. Wire EDM of the as-cast material resulted in several instances of cracking of the ingot material during cutting and numerous broken or pinched wires due to the distortion of the workpiece due to residual stress from differential thermal profiles resulting from casting and solidification. Efforts to wire EDM the cobalt material after the homogenization/stress relief treatment were more easily accomplished.

Chapter 4: Cobalt Ingot Conversion

Ingot conversion was accomplished to refine the microstructure of the entire ingot cross section, converting the coarse, columnar microstructure resulted from casting and homogenization to a much finer equiax microstructure. Conversion aids in producing material with isotropic mechanical properties for manufacturing and service through microstructure control and facilitates ease of subsequent manufacturing. The goal of this

Chapter was to demonstrate traditional ingot conversion thermomechanical processes for nickel-base alloys were also successful in refining the microstructure in the new class of gamma-prime strengthened cobalt superalloys, as well as to produce material suitable for wrought processing.

4.1 Isothermal Compression Testing of Cobalt Ingot Material

Isothermal, hot compression testing was conducted to determine a number of necessary parameters for thermomechanical processing and subsequent process modeling for the coarse grain ingot material. Documenting the coarse grain ingot material response to deformation and temperature well help define and demonstrate relevant hot working conditions. Coarse grain material will exhibit much stronger anisotropy of properties than would be expected for fine-grain recrystallized materials. As such, much of this

64 testing will be repeated on fine-grained converted material and results compared in §5.1.

Primary results included stress-strain data, or characteristic flow curves, and observations of microstructure evolution.

4.1.1 Isothermal Compression Testing Materials and Procedures

Material from the cast and homogenized E-section of the Co-Al-W ingot was wire-EDM to produce 3 bars 9 cm radially from the center of the ingot. Each bar was approximately 2.1 cm diameter by 15 cm long (Figure 41). After machining, it was

Figure 41. As-wire EDM bar of coarse-grain ingot compression specimens prior to final machining.

discovered one bar exhibited a surface-crack along a majority of its length. This crack aligned with an existing crack that opened during sectioning of the ingot into quarters prior to homogenization and stress relief. All three bars were sectioned to produce 4 cylindrical upset specimens per bar, labeled A to L. Also cut were center and end pieces for metallography, if required later to confirm starting texture. Following separation to individual specimens, the blanks were low-stress ground to ensure contact surfaces were parallel and centerless edge ground remove the recast layer from EDM, leaving them 2 65 cm diameter by 3 cm length. The top and bottom corners were chamfered to reduce friction between the compression dies and the specimen. While the cracked specimens were not mechanically tested, they were used for machine setup and thermal testing purposes to calibrate thermal profiles of the specimen during testing.

Hot isothermal compression testing was completed in a servohydraulic test frame with a load capacity of 900 kN with an environmental enclosure. This system has a pair of 83-mm-diameter TZM dies which are induction/susceptor heated in a nitrogen atmosphere to prevent oxidation of the dies and cobalt specimens. A spare upset sample was instrumented at several locations on the surface and the sample center then heated to establish a surface-to-bulk temperature correction factor. The upset specimens were coated with lubricant after being instrumented using an external thermocouple mounted at mid-height. The sample and die stack-up consisted of layers of grafoil, nickel foil, alumina insulation, and tungsten wire to facilitate separation of the specimen from the dies at the elevated temperature being used for these tests (Figure 42). The sample and die stack were all heated to test temperature over 45 minutes and then soaked at temperature, correcting for the natural thermal gradient center-to-edge from the induction setup, for 15 minutes. The samples were then compressed to a known reduction at a constant strain rate. Following upset, the bottom die, hopefully with the compressed specimen, was lowered to expose the sample where a pneumatically operated rod was used to push the tested specimen into a water-filled quench tank in the environmental chamber within 2-3 seconds. The test conditions (Table 6) were chosen to include supersolvus temperature ranges that would be feasible for ingot conversion of this alloy

66 but were limited on the upper end due to the temperature capability of the test equipment.

The load-stroke data was recorded and corrected for compliance of the load frame prior to being converted to true stress-true strain or flow stress data and reported.

Figure 42. Schematic of die-stack up and materials used during compression testing.

Temp (°C) Strain Rate (s-1) Reduction Water Quench 1149 0.01 3:1 Yes 1149 0.1 3:1 No 1149 1.0 3:1 Yes 1204 0.01 3:1 No 1204 0.1 3:1 Yes 1204 1.0 3:1 Yes Table 6. Cylindrical upset testing conditions for the coarse grain ingot material.

Following mechanical testing, the specimens were sectioned orthogonally along the major (long) and minor (short) axes of the as-deformed specimen and prepared using standard metallographic techniques for EBSD and optical microscopy for measurements of area fraction recrystallized and grain size. EBSD scans were completed using large- area batch-scanning software, capturing points in a square-grid at a 3 m step size to cover the entire height of the as-compressed sample from the center of the sample outward by approximately 5-8 mm depending on the scan time available. Analysis of the

EBSD scan data included observations of changes to microstructure.

4.1.2 Isothermal Compression Testing Results and Discussion

Compression testing of the cast and homogenized coarse-grain ingot material was completed on the six upset samples as prescribed. Unfortunately, two of the six specimens failed to water quench to preserve the as-compressed microstructure but were still successful in the primary objective of establishing the flow behavior of coarse grain ingot material during hot working conditions at various strain rates and supersolvus temperatures. The as-upset ingot specimens exhibited expected strong flow anisotropy resulting in a non-round as-compressed specimen shape (Figure 43). The remnant coarse, columnar casting structure of the ingot after homogenization results in a different material flow behavior depending on how the imposed upset was aligned with the grain structure. Similar macroscopic anisotropy of flow behavior was documented in coarse- grained, columnar Waspaloy ingot material upset in the 45° orientation relative to the compression axis and was found strongly dependent on starting columnar texture relative to the compression axis [32]. Also observed was the imprint from the tungsten wire loop

68 used to encourage separation of the specimen from the top die so the sample could be quenched after deformation.

Figure 43. Macrographs of non-uniform deformation observed in as-upset ingot material.

The stress-strain behavior of the coarse-grained, cast-and-homogenized cobalt ingot material showed a strong dependence on temperature and strain rate (Figure 44).

The higher temperature produced a 25 percent reduction in flow stress at all strain rates relative to the lower temperature. The higher strain rate samples (1.0/s) had different response than the lower strain rates (0.01/s and 0.1/s). The lower strain rate samples exhibited initial strain hardening, a peak stress, then flow softening with continued deformation. In contrast, the higher strain rate samples at both temperatures exhibited a two-part yield behavior where the rate of strain hardening changed prior to a more diffuse

69 peak stress and flow softening. Across all samples, the rate of softening increased with strain rate.

400 1149 °C 350 1204 °C 300 1.0/s 250

200 0.1/s 150

True Stress (MPa) Stress True 100 0.01/s

0 0.0 0.2 0.4 0.6 0.8 1.0 True Strain

Figure 44. Supersolvus true stress-true strain curves for coarse-grained cast and homogenized cobalt ingot material.

Comparisons with supersolvus flow behavior of coarse-grained, cast and homogenized ingot Waspaloy with the similarly processed cobalt alloy were favorable

(Figure 45). Despite differences in orientation dependence - transverse in the Waspaloy versus 45° in the cobalt relative to the compression direction - similar softening behavior was exhibited between both materials indicating microstructure evolution mechanisms may also be common. Also observed was a significant bump, 60 MPa, in flow stress of

70 the cobalt alloy relative to Waspaloy when ingot materials were compared at their respective supersolvus processing temperatures.

300

250 Cobalt - ’ + 70 °C

200 Waspaloy - ’ + 50 °C

150 Cobalt - ’ + 125 °C

100 True Stress (MPa) Stress True

Waspaloy - ’ + 160 °C 50

0 0.0 0.2 0.4 0.6 0.8 1.0 1.2

True Strain

Figure 45. Stress-strain curves comparing supersolvus flow behavior of coarse grain cobalt and Waspaloy deformed at 0.1/s [32].

Microstructure observations were preliminarily conducted using EBSD due to difficulty of etching grain boundaries without adversely impacting the rest of the sample surface. The orthogonal sections along the longest and shortest axes of the compression specimens were scanned for each test condition and several observations were noted.

First, the imposed deformation, a 3:1 reduction, during monotonic hot compression testing was sufficient to induce dynamic recrystallization but not fully recrystallize the prior ingot microstructure. The area fraction recrystallized for those compression samples that were water quenched exhibited behavior similar to nickel-base superalloys

(Table 7). Increasing the temperature resulted in an increase in area fraction recrystallized for a given imposed strain. Increasing strain rate also increased the fraction recrystallized. The difference between how much recrystallization was present for each strain rate after deformation was larger at higher temperatures, even despite the lowest strain rate specimen not quenching for the higher temperature.

Temp Strain- Water Area °C Rate Quench Fraction RX (mm/mm) % 1149 0.01 Y 41 1149 0.1 N 43 1149 1.0 Y 45 1204 0.01 N 96 1204 0.1 Y 59 1204 1.0 Y 65 Table 7. Area fraction recrystallized for 3:1 reduction coarse-grain ingot upset cylinders

The unquenched samples displayed divergent behavior, despite having nominally similar cool-down times after deformation. The higher temperature unquenched sample was nearly fully recrystallized while the lower temperature unquenched sample exhibited significantly less recrystallization. In fact, the amount of recrystallization for the lower temperature unquenched sample was found to be consistent with the quenched ones. This was explained by the proximity of the test temperature to the -prime solvus temperature.

The lower upset test temperature was performed approximately 70 °C over the ’-prime 72 solvus temperature and quickly cooled prior to sufficient time for metadynamic or static recrystallization or grain growth. The other unquenched sample, upset at the higher temperature or 125 °C above solvus, likely had enough thermal energy to promote metadynamic or static recrystallization and/or grain growth prior to cooling below the gamma-prime solvus temperature and effectively freezing the microstructure.

For the Co-Al-W alloy, the remnant intermetallic particles from the original ingot casting served as nucleation sites for recrystallized grains through a common nucleation process in superalloys, particle stimulated nucleation (PSN). The remaining dendritic- shaped distribution of the intermetallic phase, stuck in what was the interdendritic region, was plainly evident by highlighting the recrystallized regions in the as-deformed ingot sample (Figure 46). Grain boundary nucleation and PSN accounted for most of the recrystallized area, particularly at low temperatures and low amounts of imposed deformation. This was consistent with observations of PSN at carbides and grain boundary nucleation in coarse-grain, cast and homogenized Waspaloy during ingot conversion [32].

Figure 46. Compression direction IPF with recrystallization in black for 3:1 reduction.

Comparisons between the short- and long-sections revealed unusual macroscopic flow behavior. For traditional nickel-base superalloys, after deformation the longer dimension is often associated with the softer cube orientation of the coarse grains, generally with the elongated grains, and the shorter direction reflects a harder cube orientation for dislocation slip activity across the elongated grains. Coarse grain

Waspaloy exhibits flow anisotropy in this way, where the short axis of a sample will display the circular ends of grains. Conversely, the long axis would display grain

74 boundaries across the length of the sample. The cobalt alloy exhibited behavior opposite traditional alloys and permitted more deformation across the grain boundaries than along them. It was unclear if this is due to the slip systems activated during hot deformation and the inherent stacking fault of the material and differences to nickel superalloys, or whether a larger macroscopic flow localization issue occurred due to a lack of grain boundary strengthening elements that would be present in more typical alloys. The larger non-uniform flow seems to account for the gross deformation observed, and was supported by the double-yield behavior exhibited by the higher strain rate tests (Figure

44) and as seen in the macrographs of those samples (Figure 43). No other coarse grain alloys were available for comparison to see if this effect was unique to this alloy.

4.2 Cobalt Ingot Extrusion

Once the flow stress behavior was established, three extrusions were completed to convert the textured, coarse-grain, homogenized cobalt ingot material into fine-grained, randomly oriented recrystallized material for subsequent thermomechanical testing and microstructure evolution quantification. Conversion using extrusion is common among nickel superalloys and was successfully used to produce sufficient material for subsequent testing and analysis.

After an extensive search for suitable lubricants for use during high-temperature extrusion, none were found to sufficiently adhere and protect the cobalt alloy at hot- working temperatures for the times necessary for during heating and extrusion. The current alloy, by choice, did not contain any oxidation or corrosion inhibitors such as

75 chromium that would be typical of a nickel-base superalloy. Hence, the decision was made to can and co-extrude the material to show feasibility of thermomechanical processing the cobalt alloy while reducing the impact of the strong oxidation behavior.

4.2.1 Extrusion Materials and Procedures

Material from the cast and homogenized E-section of the cobalt ingot were wire-

EDM to produce preforms for extrusions roughly 6 cm diameter by 12.5 cm length. The extrusion preforms were turned to remove the recast layer from wire-EDM and to fit snugly within a section of 304 stainless steel pipe while including a barrier of SILTEMP

® insulation for heat retention and as a diffusion barrier between the can and the cobalt preform. Stainless nose and tail pieces were manufactured on a lathe and welded to the pipe to fully enclose the billet, resulting in a total stack length of approximately 18 cm.

The nose piece included an evacuation pipe to remove air from the assembly prior to sealing. Once sufficiently evacuated at temperature, the pipe was pinched and sealed shut to maintain vacuum. The entire preform assembly was then coated in glass lubricant and allowed to dry (Figure 47).

Figure 47. Turned cobalt superalloy billet ready for canning (left), welded assembly ready for evacuation and sealing (center) and coated assembly ready for extrusion (right).

The extrusions were performed locally at the Air Force Research Laboratory

Material and Manufacturing Directorate’s Material Processing Laboratory using a 635 metric tonnes extrusion press and round-to-round extrusion dies using an extrusion ratio of either 6:1 or 4.3:1. The canned stack was placed in a pre-heated furnace and soaked for 1.5 hours at 1204°C. The extrusion die was preheated to 260°C to reduce die chill and encourage uniform flow. Just prior to extrusion, the die stack received a coating of lubricant. For the extrusion, the hot can assembly was transferred within 5-7 seconds from the hot soak furnace to the extrusion die. A carbon block was placed behind the billet and a steel follower completed the stack-up in the die. The extrusions were conducted at a constant die speed of 1.5 meters per minute and the material was extruded

77 completely through the die. After extrusion, the as-extruded bars were slowly-cooled overnight in vermiculite insulation to reduce the cooling rate and accompanying risk of cracking from a stronger thermal gradient.

Following extrusion, the bars were cut to remove the extra 304 stainless steel from the nose and tail sections. Metallographic sections were taken from the steady-state region in the nose both longitudinal and transverse to the extrusion direction and prepared similarly to prior specimens for optical imaging and EBSD analysis.

4.2.2 Extrusion Results and Discussion

Three extrusions were completed successfully, with the both 6:1 extrusions resulting in a extruded length of approximately 89 cm and the 4.3:1 measuring 66 cm.

Press loads were consistent among similar reductions and exhibited a small increasing stress region where the press fills the cavity followed by a large steady-state region during the extrusion. Peak loads were well within the capability of the extrusion press at approximately 358 metric tonnes for 6:1 and 313 metric tonnes for 4.3:1.

Some tearing was exhibited in the tail welds of both 6:1 extrusions (Figure 48) and the nose of the 4.3:1 extrusion (Figure 49). This type of tearing is common for the canned co-extrusion with this type of can construction and the cobalt material inside was not impacted. After removing the nose and tail and excising check samples for metallography with an abrasive cut-off saw, there was approximately 55 cm from each of usable material for sample generation from the smaller diameter 6:1 bars and 40 cm of the larger diameter 4.3:1 extrusion.

Figure 48. Macrographs of 6:1 extrusions with nose and tail detail.

Figure 49. Macrograph of 4.3:1 cobalt extrusion with nose and tail detail.

After de-canning, the bars were inspected for defects and were determined free of surface and centerline cracking and porosity but did exhibit slight ovaling. The minor anisotropic macroscopic flow behavior was due to the strongly-textured, coarse-grain ingot starting material used as the extrusion preform. The softer stainless steel can was able to accommodate the anisotropic deformation and the cobalt remained encapsulated resulting in an overall round, canned extrusion. The anisotropic flow behavior mirrors the behavior exhibited in the coarse-grained ingot upset specimens (Figure 43).

The as-extruded microstructure was documented using EBSD methods. Both reductions featured a strongly recrystallized structure in the center with regions of unrecrystallized areas near the outer diameter (Figure 50).

Figure 50. IPF map of as-extruded cobalt material for 6:1 reduction at 1204 °C.

Also exhibited in both reductions was a gradient of grain size from center to outer diameter with the center being much larger as a result of the thermal gradient during slow cooling after extrusion (Figure 51). The center contained a high fraction of larger annealing twins, whereas the outer diameter contained a larger fraction of smaller twins.

The recrystallized and grown grain size variations were found much larger when compared to the purely dynamically recrystallized grain sizes from the coarse-grain ingot specimens. This was due to the grain growth that occurred during the slow cool of the extruded bar whereas the upset specimens were water quenched to preserve the microstructure.

Figure 51. Extrusion direction IPF maps showing relative grain size in the center (left) and outer diameter (right) for a 4.3:1 extrusion at 1204 °C with slow cool.

4.3 Post-Extrusion Recrystallization Heat Treatment

While the structure following extrusion and slow-cool from temperature was largely recrystallized for both extrusion ratios, a small area fraction of unrecrystallized material near the outer diameter of the extrudate at the can/extrudate interface remained.

The observed gradient of grain size from center to outer diameter was also undesirable

(Figure 51). A post-extrusion recrystallization heat treatment was performed to fully recrystallize the microstructure and enlarge the overall grain size to increase the difference between the starting material and subsequently recrystallized grains after further deformation processing.

4.3.1 Subscale Recrystallization Heat Treatment Materials and Procedures

Coupon heat treatments were conducted on three as-extruded samples removed from the nose-end of the steady-state region of the extrusion. These samples were 82 encapsulated in a quartz tube and backfilled with Argon and then placed with an instrumented extra sample in a pre-heated furnace at 1204 °C. The samples were allowed to reach temperature and then soaked for 5, 30, or 60 minutes before being water quenched to preserve the as-heat treated microstructure. Following heat treatment, the samples were sectioned and metallographically prepared for EBSD.

4.3.2 Subscale Recrystallization Heat Treatment Results and Discussion

The subscale recrystallization heat treatments were completed successfully. All samples, even the short 5 minute soak, displayed nearly full recrystallization across the entire diameter of the extrusion (Figure 52). Compared to the as-extruded grain sizes, the recrystallization heat treatment also successfully increased the overall grain size, with the longer heat treatment resulting in largest grain sizes. With all the grains having grown during the heat treatment, the difference in size between the smallest and largest was decreased.

Figure 52. IPF map of 6:1 reduction extruded material after recrystallization heat treatment at 1204 °C for 60 minutes with center at left.

As each heat treatment resulted in a similarly recrystallized structure, the decision was made to use the heat treatment producing the largest grain size, or 60 minutes. While the grain size obtained was not as large as the stable pinned grain size, it was sufficiently differentiated from the recrystallized grains for subsequent analysis and was chosen as an industrially relevant time for heat treatment.

4.3.3 Extruded Workpiece Heat Treatment Procedures

The as-extruded workpieces required recrystallization heat treatment. The bar was sectioned into smaller segments for evaluation, ease of handling, and so they fit within the surveyed and certified zone of the furnace used for the heat treatments. The stainless can material from extrusion was left on to protect the majority of the bar, but the ends were left as-cut and exposed. The segments were loaded onto a purpose-built tray to allow several to be heat treated concurrently with an instrumented dummy sample. The heat treatment sequence per §4.3.2 was used, and the tray, with bars and instrumented dummy sample, were loaded into a pre-heated furnace at 1204°C, allowed to reach temperature then soaked for one hour. Following the hour, the entire tray was removed from the furnace and allowed to cool in air to room temperature. Once cooled, the pieces were turned on a lathe to remove the stainless steel can and machined into specimens for subsequent mechanical testing. This process was repeated for both extrusion diameters.

4.3.4 Extruded Workpiece Heat Treatment Results and Discussion

Prior to heat treatment, the as-sectioned extruded workpieces were examined

(Figure 53). No cracks were revealed in any of the surfaces, indicating a sound extruded workpiece. Also, consistent can thickness was observed between the pieces, front-to-

84 back in the extrusion, which indicated a long, steady-state flow region. The SILTEMP ® successfully maintained a barrier between the can and the billet material during extrusion and reduced diffusion of alloying elements from the stainless steel into the cobalt billet material inside.

Figure 53. Typical as-sectioned extruded workpieces prior to recrystallization heat treatment.

The extruded workpiece segments were heat treated successfully in two batches

(Figure 54). The can material from extrusion was left intact to protect the bulk of the material from oxidation in the air of the furnace. As the ends were not protected, they showed an expected blue oxide layer (Figure 55). Some of the brittle oxide material fractured and spalled off during heat treatment. The blue coloring extended approximately 5 mm in on the surface of each end of the workpiece from oxygen ingress at the gap in the can-billet interface left by the SILTEMP ® barrier. Volumetric oxygen

85 ingress was much less. Following heat treatment, the material on the exposed ends was removed during subsequent machining operations to produce mechanical test samples.

Figure 54. Canned, extruded bars following recrystallization heat treatment with instrumented dummy sample.

Figure 55. Close-up of ends of the as-heat treated bar exposed to oxygen for 1 hour.

The recrystallization heat treatment was completed successfully. The resultant microstructure in the heat-treated extrusions was equiaxed with very low area fraction of unrecrystallized remnant prior grains and exhibited a typical annealed structure. By eye, some regions of macro-texturing were visible but EBSD measurements confirmed a nominally randomly oriented recrystallized structure.

Chapter 5: Wrought Processing of Cobalt Ingot Material

Supersolvus wrought processing was completed to establish the mechanical response of fine-grain wrought cobalt superalloy materials and quantify how the wrought microstructure evolved as a function of additional thermomechanical processing treatments. While useful empirically, this data was also used as inputs for microstructure evolution models in Chapter 6. Using the material resulting from ingot conversion, necessary information on flow stress and microstructure evolution during processing were obtained through a series of specifically chosen compression tests and pre- and post-deformation thermal treatments. These treatments were chosen to reflect conditions expected during typical industrial practice for supersolvus wrought forging operations.

5.1 Isothermal Cylindrical Compression Testing

Isothermal, cylindrical compression testing was conducted on the fine-grain billet material to determine flow behavior during hot working conditions for input to finite element deformation models and to gain insight into the microstructure evolution mechanisms dominant during supersolvus wrought processing.

5.1.1 Wrought Cylindrical Compression Testing Materials and Procedures

Wrought material from the heat treated, extruded bars was used as starting material for the cylindrical compression testing specimens. Once removed from the can, the 6:1 diameter extrusions were used to machine cylindrical upset specimens to minimize waste. In total, 16 fully recrystallized specimens, 2 cm diameter by 3 cm in length, were produced along with 4 specimens to confirm starting microstructure (Figure

56)

Figure 56. Extruded and heat treated bar with cylindrical compression specimens.

Mechanical testing was completed identically to §4.1.1 using the same servohydraulic test frame with a load capacity of 900 kN with an environmental enclosure filled with nitrogen to reduce oxidation of the cobalt and the dies. The samples were coated, instrumented, and inserted in the die stack assembly (Figure 42). The 89 samples were preheated, allowed to equilibrate at test temperature, then deformed at a constant, specific strain-rate to a known reduction prior to being quenched in water to preserve the microstructure. The conditions for testing were chosen to bracket those typical of industrial practice for supersolvus thermomechanical processing and will be sufficient for populating the required material flow stress data for process models. These conditions also mirrored the conditions for the ingot conversion testing so that flow behavior between the coarse grain and fine grain materials could be directly compared.

The load-stroke data was corrected for machine stiffness and translated to stress-strain curves.

Following mechanical testing, the as-deformed specimens were sectioned diametrically and prepared using standard metallographic techniques for EBSD and optical microscopy similar to prior sections. EBSD and optical microscopy were performed to conduct measurements of area fraction recrystallized and grain size. EBSD scans were conducted using a traditional SEM with EBSD detector and software, coupled with large-area batch-scanning software. The scans captured points in a square-grid at a

3 m step size to cover the entire 11.3 mm height of the as-compressed sample from the center of the sample outward by approximately 5-8 mm depending on the scan time available. Of that, approximately the middle 1/3 height was used to make measurements of recrystallization fraction and grain size.

5.1.2 Wrought Cylindrical Compression Testing Results and Discussion

Hot isothermal cylindrical compression testing was completed successfully for the

6 test conditions defined (Table 8). For this set of experiments, one of the specimens

90 failed to water quench to lock in the microstructure but was re-run with successful water quench. The coatings used for lubrication remained intact and the nitrogen-filled environmental enclosure prevented oxidation of the sample and the dies as intended.

Temp (°C) Strain Rate (s-1) Reduction Water Quench 1149 0.01 3:1 Yes 1149 0.1 3:1 Yes 1149 1.0 3:1 Yes 1204 0.01 3:1 Yes 1204 0.1 3:1 Yes 1204 1.0 3:1 Yes Table 8. Isothermal cylindrical compression testing conditions for wrought cobalt.

The extrusion and recrystallization heat treatment successfully randomized the starting ingot microstructure and texture, which resulted in round, axisymmetric compression samples after deformation. A slight barreling was consistently observed on all upset specimens which indicated friction between the upset specimen and the forming dies during the testing was reasonably low (Figure 57). The exterior surface of the specimens was smooth, without tearing or cracking, and the specimens did not exhibit any of the bulk deformation flow irregularities observed in the ingot upset test specimens

(Figure 43).

Figure 57. Typical as-upset wrought cobalt superalloy cylindrical specimen from top view (left) and side view (right).

Using this series of tests, the supersolvus deformation behavior was captured.

The load-displacement data was corrected for machine deflection and true stress-true strain was plotted for each temperature and strain rate condition (Figure 58). Similar to

§4.1.1, the stress-strain curves showed typical dependence on temperature and strain rate with decreasing temperature and increasing strain rates resulting in increased flow stresses. The relative magnitudes between conditions remained consistent. Both the wrought and the ingot material displayed similar peak and yield behaviors. Behavioral consistency was demonstrated after similar flow behavior was observed for both the quenched and unquenched samples at the same test condition on being recompleted.

400 1149 °C 350 1204 °C 300 1.0/s

250

200 0.1/s 150 0.01/s

True Stress (MPa) Stress True 100

0 0.0 0.2 0.4 0.6 0.8 1.0 True Strain

Figure 58. True stress-true strain curves for a wrought Co-Al-W superalloy.

The flow behaviors of the wrought material and the ingot material were more similar than anticipated (Figure 59). From a bulk flow perspective, this was explained by comparing the impact of starting crystallographic texture on the compression behavior.

While the ingot specimens were very strongly textured with coarse, columnar grains, the grains were oriented relative to the compression direction such that the strength of that orientation (Taylor factor, M~3) was near that of a nominally randomly oriented recrystallization texture (M~3) [47]. As such, both starting textures would accommodate upset loading similarly. If the columnar texture of the coarse grains in the ingot

93 specimens were more strongly aligned with the compression direction the differences in the stress-strain curves would be more significant, as demonstrated on coarse-grained

Waspaloy ingot material [32].

400 Wrought 350 Ingot 300

250 1149 °C - 1.0/s

200 1204 °C - 1.0/s 1149 °C - 0.1/s 150 1204 °C - 0.1/s 1149 °C - 0.01/s True Stress (MPa) Stress True 100 1204 °C - 0.01/s 50

0 0.0 0.2 0.4 0.6 0.8 1.0 1.2 True Strain

Figure 59. Comparison of ingot and wrought flow stress curves of a ' cobalt superalloy.

While the flow behavior was similar between the ingot and wrought materials, some differences were observed. The rates of softening for the elevated strain rates, 0.1/s and 1.0/s, was higher during deformation of the wrought material (dashed lines).

Assuming similar rates of deformation heating for both the wrought and ingot compression testing for a given strain rate, the increased softening was attributed to

94 increased dynamic recrystallization (DRX) because of the higher fraction of grain boundaries present for nucleating recrystallization in the wrought material versus the ingot material. Area fraction measurements from EBSD in the center of each specimen for all conditions showed a significant increase in fraction recrystallization relative to the that observed in the ingot specimens.

The strain rate sensitivity, m, was calculated from the wrought flow stress data using Equation 4 and calculating the log of the stress ratio divided by the log of the strain-rate ratio. For the wrought specimens, the strain rate sensitivity varied slightly, but m = 0.2. 휎 log( 2⁄휎 ) 푚 = 1 (4) 휀̇ log ( 2⁄ ) 휀1̇ Also calculated from the stress strain curves of the wrought upset cylinders was the apparent activation energy for deformation [48], Qapp, using Equation 5, where R was the gas constant, S was the temperature sensitivity of the flow stress, and m was the strain rate sensitivity from Equation 4. The temperature sensitivity, S, was calculated by

푅푆 푄 = 2.303 (5) 푎푝푝 푚 determining the slope of the lines made by plotting (1/Temperature) versus the logarithm of the peak stress from the flow stress data for each temperature and strain rate (Figure

60). All three lines had a similar slope and averaged S = 4825 K. Using S = 4825 K and the strain rate sensitivity, m=0.2 resulted in a deformation Qapp = 462 kJ/mol for the fine grain, wrought cobalt. This number was effectively identical to the value determined for

95 a fine grain nickel superalloy, Waspaloy, which claims Qapp = 468 kJ/mol [41], but higher than the coarse grain Waspaloy calculated the same way with Qapp = 367 kJ/mol [32].

2.7 2.6 1.0/s 2.5 0.1/s

2.4 p

e 2.3 log 2.2 0.01/s 2.1 2.0 1.9 6.8 6.8 6.9 6.9 7.0 7.0 7.1 10^4 * 1/T (1/K)

Figure 60. Determining slope of 1/T versus log peak strain to determine temperature sensitivity of the flow stress, S, in K.

Impact of friction coefficient on the calculated strains was evaluated for the cylindrical upset tests. Typical idealized hot forging friction coefficients for the finite element model vary between 0.3 for lubricated and 0.7 for non-lubricated and are not generally considered temperature sensitive. Experimentally, borosilicate glass lubricants should have maintained their lubricity better than boron-nitride lubricants because boron- nitride spray lubricants were not well suited to the elevated test temperatures used for the cobalt. However, the boron nitride lubricants were stable and did not react with the

96 cobalt as the borosilicate glass lubricants had during earlier testing and exposure in air.

Neither accounted for the graphite foil and the nickel foil stack-up to reduce sticking of the sample to the dies. Instead, a single composite friction coefficient between the cylindrical workpiece and the dies was used. The friction coefficient in the model was found to have an impact on the final calculated imposed strain (Figure 61) and was also found to vary with strain rate and temperature. These findings were consistent with experimental ring tests conducted on a similar cobalt alloy [49]. Higher friction values resulted in predictions of increased barreling and a higher calculated variation of imposed strain between center and outside of the as-deformed sample.

Figure 61. Changes to calculated strain and strain rate profiles with friction coefficient, mu, for a 2.75:1 upset cylinder at 1149 °C and 0.1/s.

To facilitate determination of the appropriate friction coefficient, a cylindrical upset test was conducted using only boron nitride spray coatings for lubrication. Load- stroke data for both the typical- and high-friction experiments were compared to finite element simulation results for cylindrical upset tests using different coefficients of friction. A representative example was shown for upset at 1149 °C upset at 0.1/s strain 97 rate (Figure 62). Increasing friction did result in a predicted increase in load and when compared to the analytical results confirmed the increase of stress with increase in friction. The simulation behavior best reflected the experimental behavior when using a friction coefficient of  = 0.3 to represent the combination of boron nitride spray coatings and foil stacks, and would be used for subsequent finite element simulations.

180000 160000 140000 120000 100000

80000 Load(N) 60000 FEM Simulation, mu = 0.3 40000 Experiment, with foil stack 20000 Experiment, no foils 0 0 5 10 15 20

Stroke (mm)

Figure 62. Load stroke comparison of experiment versus simulation for different friction conditions for cylindrical upset tests conducted at 1149 °C 0.1/s.

While the friction fit was successful, this exercise highlighted an issue that the pre-load on the foil stack may not be sufficient to compress them completely prior to testing. A higher pre-load may have reduced the apparent drift initial loading behavior that was not observed in the higher friction, non-lubricated condition. 98

5.2 Isothermal Double-Cone Compression Testing

Additional isothermal compression testing was completed at supersolvus hot working temperatures to establish microstructure evolution behavior for a wrought cobalt superalloy under a variety of temperature, strain rate, and total imposed strain conditions.

Double-cone geometry specimens were used to obtain a large imposed strain gradient across the diameter of each specimen, effectively reducing the number of tests required for the desired microstructure data versus traditional cylindrical upset testing. The results of the testing were used as inputs to the supersolvus microstructure evolution models being developed for cobalt superalloys in Chapter 6. The results were also compared with a typical nickel-base superalloy.

5.2.1 Isothermal Double-Cone Compression Testing Materials and Procedures

Wrought material from the extruded and recrystallization heat treated bars were used as starting material for the cylindrical double cone compression testing specimens.

The 4.3:1 diameter extrusion material was utilized as it had large enough as-extruded diameter to produce the larger diameter double cone geometry. Once decanned, 13 wrought double-cone specimens were machined along with several extra smaller coupons for confirming starting microstructure and subsequent heat treatment experiments (Figure

63). The dimensions of the double cones (Figure 64) were scaled and tailored relative to those used in previous work on Waspaloy [32] (Figure 65) to fit within diameter of the

4.3:1 extrusion while maintaining the overall aspect ratio. Some measurements were modified to aid in ease of machining, but the flat ends were made larger to make the sample more stable on the dies during deformation and reduce stress with a larger contact

99 area. With a significantly smaller starting grain size, the number of grains along the diameter was not a concern but ensuring the sample remained large enough to quench successfully in the test equipment was. Due to the smaller as-compressed height, the maximum reduction attempted was 3:1 for the safety of the mechanical test equipment.

Figure 63. Double-cone sample orientation relative to the extruded bar.

Figure 64. Dimensions in millimeters for double-cone geometry used in present work.

100

Figure 65. Size comparison of double-cone geometry for coarse grain ingot work (left) versus present fine-grain wrought work (right).

Similar to §4.1.1 and §5.1.1, this mechanical testing was completed using the same 900 kN servohydraulic test frame utilizing an environmental enclosure filled with nitrogen to reduce oxidation of the cobalt test specimens and protect the dies. The samples were instrumented, coated with boron nitride, and inserted in the die stack assembly (Figure 66). The samples and die stack were preheated and allowed to equilibrate at test temperature then deformed at a constant, specific strain-rate to a known reduction relative to the original sample height and quenched in water to preserve the resultant microstructure.

101

Figure 66. Double-cone experimental configuration.

The conditions for testing were chosen to be broadly representative of industrial practice for supersolvus thermomechanical processing and largely mirror the wrought and ingot upset test specimens. A die stress analysis was conducted using DEFORM [46] version 11.2 to determine the expected die stresses as a result of the smaller initial contact area of the double-cone geometry and the higher strength at temperature of the cobalt superalloy. An axisymmetric and isothermal 2D simulation of the dies and double-cone workpiece was constructed using standard lubricated hot-forming parameters like friction

(Figure 67). The double-cone geometry was directly imported into DEFORM and the cobalt material flow stress data was a result from the wrought hot isothermal cylindrical

102 upset testing in §5.1.2. The TZM dies were sized to match the experimental equipment and were considered purely elastic for the purposes of the simulation and subsequent analysis using temperature-sensitive elastic modulus data [50]. Predicted die stresses above 275MPa resulted in further evaluation while stresses below that were considered lower risk.

Figure 67. Finite element representation of cobalt double cone hot compression test setup.

After the prescribed thermomechanical processing, the compressed double cone specimens were sectioned longitudinally along the diameter for metallography and additional heat treatments. Once completed, the samples were prepared for metallography using standard techniques. Specimens were then mapped using EBSD.

Using a combination of grain size and misorientation-based methods, segmentation was

103 completed to differentiate recrystallized and unrecrystallized areas (Figure 68) and measurements of recrystallized fraction were recorded as a function of radial location.

Within the EBSD software, the average recrystallized grain sizes were determined, after accounting for twins, by measuring only the recrystallized grains in a field of view. Due to the double-cone geometry, the effective strain varied significantly across the diameter of the deformed upset specimen. The imposed strain at each radial location was determined using an axisymmetric, isothermal FEM simulation [46] using the material flow stress data from the wrought cylindrical compression testing completed in §5.1.2 and the thermophysical property data from §2.4. This methodology for determining local imposed strain was similar to approaches used previously [38, 51, 52].

Figure 68. Compression Axis IPF maps as-scanned (left) and after threshholding to highlight recrystallized areas in black (right).

104

5.2.2 Isothermal Double-Cone Compression Testing Results and Discussion

A series of supersolvus isothermal compression tests on wrought cobalt double- cone specimens was completed successfully. The primary results were observations and analysis of the microstructure evolution behavior including dependence of recrystallization behavior on imposed temperature, strain, and strain-rate.

5.2.2.1 Die Stress Analysis by FEM Simulation

A die stress analysis was completed for all six planned processing conditions, including two temperatures (1149 °C and 1204 °C) and three strain rates (0.01/s, 0.1/s, and 1/s). The maximum effective stress in the dies followed a similar pattern for each condition and consisted of a peak stress early in the deformation stroke followed by a gradual decrease in the stress from both the additional contact area gained by deforming the double-cone geometry and the flow softening of the alloy due to recovery and recrystallization. The maximum predicted stress occurred just below the die surface near where the double cone made contact (Figure 69). This number may be over-predicted slightly due to the non-chamfered corner of the double-cone geometry simulated and varied with friction coefficient utilized.

Increasing the strain-rate and lowering the temperature both resulted in prediction of higher die stresses. Specifically, the intermediate strain rate (0.1/sec) simulations resulted in die stresses approaching the expected capability of the TZM dies while the lowest strain rate (0.01/sec) simulations were safely within the assumed capability of the dies (Table 9). The highest strain-rate conditions (1.0/sec) for both temperatures predicted die stresses significantly over the expected yield stress capability of the TZM

105 dies and were removed from the double-cone test matrix for the safety of the test equipment.

Figure 69. Predicted maximum die stresses for 1204 C 0.1/s upset.

Temp ( °C) Strain Rate (1/s) Max Die Stress (MPa) 1149 0.01 222 1149 0.1 350 1149 1 809 1204 0.01 176 1204 0.1 273 1204 1 625 Table 9. Predicted maximum die stress during double-cone compression tests. 106

Experimental load-stroke curves were corrected for system compliance and compared to the predicted press loads from the FEM under identical simulated test conditions to corroborate the results. The loads were compared rather than stress in this case because the changing geometry and contact area of the double-cone specimens during compression made an analytical solution for stress from the experimental data more difficult. Loads, on the other hand, were available from both experimental and finite element results. On a larger scale, the FEM reflected the experimental results well, and captured the slope and offset increase in load due to increasing strain rate from 0.01/s to 0.1/s (Figure 62). For very small displacements, the FEM underpredicted the loads, but not significantly, and was attributed to the input stress-strain data and the impact of the foil stack-up. If the FEM was underpredicting the loads of the 1.0/s strain rate, then given the very small diameter contact area of the double cones yield of the TZM would likely be exceeded. After a small deformation stroke the contact area would have increased sufficiently, and continued to do so with increasing deformation, that the stress would decrease to safe levels despite increasing loads. The reasonable agreement with experimental results observed for the 0.01/s and 0.1/s simulations provided confidence against completing the 1.0/s testing for the safety of the test equipment. Machine inserts of higher capability die materials were considered but due to time, cost, and complexity constraints were not procured for the testing and the 1.0/s strain rate double cone testing remained uncompleted.

107

200000 180000 Experiment 160000 Simulation 140000 1.0/s 120000 0.1/s 100000

80000 Load(N) 60000 40000 0.01/s 20000 0 0.0 2.0 4.0 6.0 8.0 10.0 12.0 Stroke (mm)

Figure 70. Press load versus stroke for 1149 °C double cone upset test at 0.01/s, 0.1/s, or 1.0/s measured experimentally (dashed lines) or predicted by FEM (solid lines).

5.2.2.2 Macroscopic Deformation Behavior

Similar to the wrought cylindrical upset specimens, the double-cone specimens remained round after deformation due to the alignment of the compression direction with the prior extrusion axis and the nominally randomized wrought recrystallization texture of the starting material. Final reductions were slightly less than expected due to the height and deformation of the graphite foils used for lubrication and system compliance.

Overall the testing was successful and samples did not exhibit indicators of issues with the test setup for bulk macroscopic flow like sticking or laps/folds (Figure 71). The tungsten wire fastening the specimen to the bottom die remained intact and resulted in the

108 samples being quenched as desired. A small depression was observed at the initial contact point on both the top and bottom surfaces similar to the upset tests where the combination nickel foil and graphite foil stack was initially pinned between the sample and both top and bottom dies.

Figure 71. Wrought double-cone compressed at 1204 °C 0.01/s to 3:1 reduction.

The impact from friction and not sticking to the compression dies made the slight differences in as-compressed heights a worth-while trade. All of the double cone specimens separated from the dies for water quenching. The lubricity provided was determined acceptable after only slight barreling of the sides was observed with no surface cracking or tearing. The geometry of the double-cones still provided a gradient in imposed strain, despite not reaching a full 3:1 reduction. Had a specific reduction been required, additional upset compensation would have been needed to achieve the desired reduction.

109

5.2.2.3 Strain Determination by Finite Element Modeling

The final strain calculation was made by correlating the radial location of each scanned area in the specimen to corresponding locations within an axisymmetric finite element simulation [46] of the same geometry, upset using the same temperature and strain rate parameters and compressed to the same final specimen height as each individual upset test. Prior to the final determinations, two separate decision points in the model were addressed to better align the experimental and simulation results: friction and final sample height. Both the friction and final sample height were concerns because the model did not explicitly reflect the entire foil stack-up, similar to the cylindrical upset testing in §5.1.2.

As mentioned previously, the foil stack-up had implications to the final reduction imposed on the test specimen in the load frame. The final as-deformed height of the double cone specimens was taller – had less reduction – than a true 3:1 reduction. This was attributed to the finite thickness and compliance of the foil stack-up during testing and was resolved by limiting final imposed deformation for the simulations to match the resultant experimental sample height.

The impact of friction coefficient was also determined for the double-cone upset test geometry. Similar to the cylindrical upset test geometry, finite element modeling was used to show the variation of effective strain and strain rate along the mid-plane of the sample from center to outer diameter (Figure 72). The relatively constant slope of the effective strain curve for each of the friction conditions was promising for the subsequent analysis. While the difference in the center of the specimen were significant, ranging

110 from e~1.6 to e~2.3, once past mid-radius the differences were negligible for all three friction conditions.

Figure 72. Simulated strain (left) and strain rate (right) profiles for 3:1 double cone upset at 1149 °C at 0.1/s for different friction coefficients, mu.

The same was true for the calculated strain rates from mid-radius outward.

Again, higher friction coefficients resulted in higher predicted effective strain rates at the center of the compressed double-cones resulting. The overall model prediction showed strain rates over double the intended target strain rate in the center, but half the intended strain rate at the outer diameter for those conditions. This illustrated a key weakness in the mechanical testing approach being employed for this work. Specifically, while the bulk macroscopic strain rate of the test is set and controlled accurately by changing the cross-head speed of the test frame, the unique geometry used caused local heterogeneities in strain rate that were sensitive to friction coefficient. This may confound some of the analysis and results on materials that are determined strain rate sensitive by exhibiting

111 local material responses to other strain rate regimes despite using a defined macroscopic strain rate.

A better friction model could be determined experimentally and implemented for this specific test setup and material to more accurately define the strain profile as a function of temperature but was outside the scope and time available. For the purposes of exercising a microstructure model on the double-cone geometry upset test, using a constant value for the friction coefficient was deemed sufficient, while understanding the limitations. The friction coefficient was easily updated within the finite element framework and processing path information subsequently transferred to the microstructure evolution models.

Using the determined friction,  = 0.3 and the experimentally measured sample heights, the resultant strain profiles were calculated for each test condition (Figure 73).

The minor differences in sample height did not change the imposed strain profile significantly, but were accounted captured. Plots correlating the calculated imposed strain to the area fraction recrystallized are shown in later sections.

112

Figure 73. Finite element representation of as-deformed double-cone geometry (left). Effective strain profile for the mid-plane computed from the finite element model (right).

5.2.2.5 EBSD Microscopy Results and Recrystallized Grain Sizes

EBSD was completed for each test condition on the as-compressed double cone specimens and area fraction dynamically recrystallized and grain sizes were recorded as a function of radial location. Compression-axis IPF maps of each test condition were produced with grain boundaries greater than 15 degrees highlighted in black (Figure 74).

Darker regions toward the left of the images correlated to the center and high-strain regions of the specimens. Recrystallized regions with smaller grains appeared darker than regions with remnant larger, unrecrystallized grains. All exhibited the expected gradient of recrystallized fraction with higher amounts in the higher strain regions at the center transitioning to lower amounts in the low strain region at the outer diameter.

Similar to the upset testing results, increasing temperature resulted in increased fraction recrystallized for a given strain rate, while increasing the strain rate decreased the overall fraction recrystallized for a given temperature.

113

Figure 74. Compression-axis (vertical) IPF map of diametral mid-plane for all test conditions.

The compression-axis IPF maps also showed the evolution of texture from a nominally random recrystallization texture at the outer diameter to a much finer and still random recrystallization texture nearer the center. The larger remnant unrecrystallized grains within otherwise recrystallized areas showed a copper deformation texture [53] that increased in intensity from outer diameter inward and scaled with the amount of imposed deformation. This was evidenced by the strong green [101] coloration in the

IPF maps, particularly of the remnant, unrecrystallized grains in the higher deformation areas. The remnant unrecrystallized areas in the ingot upset specimens also exhibited this behavior (Figure 46) as did other FCC alloys like Waspaloy [38].

The average grain sizes in the recrystallized areas were determined using EBSD by filtering the EBSD data to remove the unrecrystallized areas. The as-dynamically 114 recrystallized average grain size was sensitive to temperature. The lower deformation temperature produced a finer grain size than was observed for the higher temperature testing. Neither the strain-rate nor the imposed strain had a significant effect on the as- recrystallized grain size (Table 10). These grain sizes were strongly dependent on the step size of the scan (2m) and the cleaning of the data performed. However, the trends held regardless of cleaning process performed assuming all data were treated similarly.

Test Conditions e ~ 1.3 e ~ 0.7 e ~ 0.3 1149 °C – 0.01/s 5.3 m 5.0 m 5.3 m 1149 °C – 0.1/s 5.3 m 5.2 m 4.4 m 1204 °C – 0.01/s 6.9 m 6.9 m 7.7 m 1204 °C – 0.1/s 7.1 m 7.3 m 7.2 m Table 10. Dynamically recrystallized grain diameter by thermomechanical condition.

5.2.2.6 Dynamic Recrystallization Behavior

The dynamic recrystallization behavior was established as a function of temperature, strain rate, and imposed strain. All of the samples showed a sigmoidal dependence of fraction recrystallized versus strain, similar to behavior expected of nickel-base superalloys (Figure 75). As shown, increasing the temperature increases the fraction DRX observed in the microstructure. Increasing the strain-rate decreased the fraction DRX. Full DRX was observed for the center of all of the double cones at an imposed strain e~1.5 except for the lower temperature, higher-strain-rate. At the lower temperature, there was a higher sensitivity to changes in strain rate than was observed at higher temperatures.

115

1.0 0.9 0.8 0.7 0.6 0.5 0.4 0.3 1149C 0.01/s 1149C 0.1/s 0.2 X, X, Fraction Recyrstalilized 1204C 0.01/s 0.1 1204C 0.1/s 0.0 0 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 Von Mises Strain (mm/mm)

Figure 75. Imposed strain versus area fraction recrystallized in wrought cobalt superalloy.

The double cone geometry was successful in imparting the imposed strain gradient and capturing the corresponding recrystallized fraction. Per the FEM, the 3:1 reduction selected resulted in an imposed strain from e~0.25 at the outer diameter to e ~

1.9 in the center but varied slightly depending on test setup. The fraction recrystallized, especially for the lower strains, was much higher than expected. The slope, or rate of increased DRX fraction with strain, was rapid for lower strain values. This rate was higher than expected for traditional grain boundary nucleation and was attributed to the contribution of PSN at the interdendritic phase at low imposed strains. 116

When compared to the model developed for wrought Waspaloy [41] using similar grain sizes, strain-rates, and supersolvus temperatures, the cobalt compared favorably

(Figure 76). As mentioned previously, the rate of recrystallization in the cobalt was more rapid than predicted for Waspaloy for the higher temperature, but comparable or perhaps slower for the lower temperature tests as strain was increased. At lower temperature, the

1.0

0.9

0.8

0.7

0.6

0.5

0.4

0.3

X, X, Fraction Recrystallized 0.2

0.1

0.0 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 Von Mises Strain (mm/mm)

WaspMod Low T Low edot WaspMod Low T Mid edot WaspMod Hi T Low edot WaspMod Hi T Mid edot Cobalt Low T Low edot Cobalt Low T Mid edot Cobalt Hi T Low edot Cobalt Hi T Mid edot

Figure 76. Comparison of wrought cobalt data to wrought Waspaloy model [41]. 117 imposed strain required for full recrystallization was much higher than needed for

Waspaloy. The cobalt showed a higher sensitivity to temperature than was predicted for

Waspaloy. The sensitivity of Waspaloy to strain-rate was much lower overall than the cobalt for lower test temperatures but appeared similar for the higher temperature. Due to the missing low strain, low area fraction experimental measurements, it was not clear if the cobalt exhibited a critical strain requirement for recrystallization which would dictate a delay in the recrystallizated fraction to a critical strain value before a sharp increase or if it predicted recrystallization from zero strain as observed in Waspaloy [41].

These results were encouraging as they were structurally consistent with the

Waspaloy and was evidence a similar model could be developed to predict the dynamic recrystallization in a cobalt superalloy of this type. Model development occurred in

Chapter 6.

5.2.2.8 Summary of Recrystallization kinetics

Dynamic recrystallization behavior was exhibited at much lower strains than expected. It was unclear from the evaluation if this behavior was the result of a difference in stacking fault energy and the accumulation of strain that is inherent to the

Co-Al-W alloy family or an influence on nucleation rate from the intermetallic particles remaining from casting and solidification, or some combination of both. Unpublished research from Wertz [49] on a similar cobalt alloy lacking the interdendritic particles showed significantly lower recrystallized fraction for the same imposed strain than was observed in the alloy with particles. Comparing the recrystallized fraction of the cobalt alloys to a wrought model of a typical nickel-based superalloy, the particle containing

118 alloy was moderately higher than the nickel alloy while the non-particle containing cobalt alloy was significantly lower than the fraction predicted from the wrought nickel alloy model for a known temperature, strain, and strain rate. This finding indicated potential stacking fault energy differences between the cobalt and nickel superalloys was not a significant cause of the differences in the dynamic recrystallization behavior. Rather,

PSN or starting microstructure prior to deformation appeared to be driving the difference in recrystallization behavior. The lesser amount of recrystallization for the same strain may be due to the larger starting grain size of the non-particle containing alloy. A larger starting grain size would have inherently less sites for traditional site-saturated nucleation that would typically occur on grain boundaries and recrystallize more slowly with deformation.

5.3 Static Grain Growth Behavior

The microstructure was not stable at supersolvus hot working temperatures.

Quantitative information regarding the change in the microstructure, and more specifically grain size, during representative supersolvus processing was desirable for microstructure control. Quantification of the evolution of grain size during simple supersolvus heat treatments such as during hot soak prior to and between deformation at temperature was needed. A series of supersolvus heat treatments were completed to address that requirement and demonstrate feasibility of using a traditional approach to quantifying and predicting grain growth.

119

5.3.1 Static Grain Growth Materials and Procedures

After recrystallization heat treatment, the larger 4.3:1 extrusion was sectioned and quartered to produce coupons for heat treatment and subsequent grain growth analysis.

The coupons were then encapsulated in quartz tubing with an oxygen getter and backfilled with argon to an appropriate pressure for subsequent heat treatment before being sealed similar to the homogenization samples in §3.2.

The sealed quartz vessels with the coupons were loaded onto a custom fabricated rack for heat treatment in a certified and instrumented furnace. A similar size coupon of the same material was instrumented using a platinum thermocouple and placed on the rack with the other encapsulated coupons in the furnace to approximate the temperature of the coupons during initial heat-up. After reaching the desired temperature, 1149 °C or

1204 °C, the coupons were allowed to soak for 2, 6, or 24 hours prior to being removed from the furnace and water quenched.

Once cooled, the heat-treated specimens were sectioned and prepared for electron imaging and EBSD using standard metallographic methods. EBSD was conducted on a 2 mm by 2 mm area of each sample using a 3 m step-size due to the large expected grain size. Measurements of average grain size after removing twins and without including edge grains were calculated from the EBSD grain files for each temperature and time condition consistent with the specification used prior [37]. The interdendritic phase was also segmented and statistics such as size and aspect ratio were recorded for each condition.

120

5.3.2 Static Grain Growth Results and Discussion

The grain growth heat treatments were successfully completed with all samples successfully quenched. Grain growth was surprisingly rapid and reached steady state at d

= 48-50 m after the initial 120 minutes of heat treatment, and remained similar through

24 hours. At times beyond the 120 minutes, the grain size was limited by the interdendritic eutectic phase particles in the material, similar to carbides pinning boundaries in nickel-base superalloys. The particle limited grain size, otherwise known as the Zener-limited grain size, was provided by Equation 6:

2푑푝 퐷푍푒푛푒푟 = (6) 3퐹푣

With DZener being the Zener-pinned grain size, dp the diameter of the pinning phase, and

Fv the area fraction of the pinning phase. After extrusion, the interdendritic particles had an average diameter of 9 m and maintained a similar area fraction at 1-2 %, for a DZener between 300 and 600 m. This was much larger than the observed pinned grain size of

48-50 m. This order of magnitude difference was accounted for when reviewing the relevant assumptions for this derivation requiring spherical particles that were even dispersed such that a single grain boundary would be considered planar among several particles. The Zener size also assumed the geometric constant, , in the original formulations of grain growth kinetics by Burke and Turnbull [54, 55] could be set such that  = 1 and could be neglected. The version with  (Equation 7) was generally found

2훼푑푝 퐷푧 = (7) 3퐹푣

121 more accurate with alpha values between 0.25 <  < 0.5 [56-60] reducing the predicted limited grain size for cobalt to Dz = 75-300 m. When the interparticle spacing was found similar to the grain size, a separate limiting formulation was required (Equation 8)

푑푝 8훼 퐷푍푙 = √ (8) 2 3퐹푣 which relates the pinning pressure of the particles to the driving pressure for growth [60].

The resulting pinning-limited grain size ranged from DZl = 25 – 50 m was consistent with experimental observations, or slightly smaller. The best agreement when using scan data from EBSD for average size, dp = 8.7 m, and area fractions, Fv = 0.01, were achieved using the upper-end of the typical geometric fitting parameter,  = 0.5, resulting in the 50 m pinned size.

122

Chapter 6: Empirical Microstructure Model for a Cobalt Superalloy

The ability to predict resultant microstructures from sequences of supersolvus thermomechanical processing is necessary to reduce the trial-and-error approach for the desired microstructure and augmenting processing-microstructure-property relationships for eventual material maturation, qualification, and insertion. Dominant microstructure evolution mechanisms, including Dynamic Recrystallization (DRX) and grain growth, were modeled by fitting experimental data to the empirical Johnson-Mehl-Avrami-

Kolmogorov (JMAK) equation for DRX and a traditional power law for grain growth.

Both approaches are well known and would demonstrate compatibility of traditional physical metallurgy-based empirical models on novel cobalt superalloys.

This approach did not explicitly model specific microstructures or grains and microstructure evolution with deformation and temperature but rather determined bulk average statistics that were none-the-less useful in representing the microstructure and the response to thermomechanical processes However, an additional step would be required to work backwards and create a synthetic structure that fit the average statistics. Lastly, this method was limited as the fits are employed as a post-process after thermomechanical simulation had already been completed. While sufficient for a first order approximation and generally capturing trends, the longer a simulation conducted

123 the less accurate the results were likely to be without intervention. Methods to address these concerns were available but generally increased complexity, so the standard JMAK formulation was utilized for this investigation.

6.1 Dynamic Recrystallization

Dynamic recrystallization is a softening mechanism during which the accumulation of dislocations at prior grain boundaries and particles cause the nucleation and growth of a new, strain-free grain. During hot deformation, the cycle of strain hardening and subsequent softening through recrystallization and recovery occurs in a continuous cycle called continuous DRX and was the focus of this investigation. The fraction of dynamic recrystallization was observed by quenching samples immediately after deformation and analyzing resultant microstructures. Utilizing the JMAK formulation, the amount of dynamic recrystallization was related to the as-heated grain size prior to upset, imposed strain, temperature, and strain-rate during deformation. Input data for the JMAK-style DRX model required the peak strain, strain for 50% DRX, and fraction DRX as a function of strain. The DRX grain size was also needed for the model as a function of strain rate and temperature. The earlier wrought testing was structured with the intent to provide the required data for model formulation using the least testing possible. Mechanical test conditions and analysis data was used as input to the Materials

Suite Post-Processor Module of DEFORM [46], which uses a formulation by Sellars [61] previously detailed on Waspaloy [41, 52] to fit the data to the JMAK governing equations. The platform used a parametric design study to fit the constants. As such, it

124 was necessary to be mindful of the values determined for fitting because the solution presented may not be unique as adjustments to some values could be offset by others.

Experimental results and measures were used where possible to limit this occurrence, but did not always produce the best fit.

6.1.1 Peak Strain

The strain corresponding to the peak of a typical flow-stress curve is typically an indicator of the initiation of DRX. Recrystallization is a softening mechanism and causes changes to the flow-stress curves of the material during deformation. The cobalt exhibited a strong double-yield phenomenon (Figure 58) at moderate and lower strain- rates. Within certain temperature and strain rate ranges, superalloys such as Waspaloy and Inconel 718 exhibited a single maximum peak to the flow-stress curve followed by softening due to recovery and recrystallization [62], but for other temperature and strain rates a double-yield was produced similar to the shape behavior demonstrated by the cobalt. Thus, the cobalt was similarly able to be fit to an empirical determination of the peak strain (Equation 9). In Equation 9, ep was the peak strain, a was a fitting constant,

−푄 휀 = 푎푑ℎ휀̇푚 푒푥푝 ( ) + 푐 (9) 푝 0 푅푇 d0 was the as-preheated diameter, edot was the strain rate, h was a grain growth exponent, m was a strain rate sensitivity, Q was the apparent activation energy, R was the gas constant, T was the test temperature, and c was another fitting constant. Using the Zener-

Hollomon parameter, Z, shown in Equation 10, the peak strain was written as Equation

11.

125

푄 푍 = 휀̇ 푒푥푝 ( ) (10) 푅푇 푛 푚 휀푝 = 푎푑 푍 (11) However, to utilize the data output from the finite element model, Equation 12 was fitted for the supersolvus peak strains in the general form of Equation 9. Of important note when switching between formulations was the exponent ‘m’ when applied to ‘Z’ translates to an exponent on the strain rate term and as a multiple to the value in the exponential term. While the apparent activation energy from §5.1.2 was Qapp = 462 kJ/mol, the effective activation energy appearing in the numerator of the exponential was m*Qapp.

69530 휀 = 3.227 ∗ 10−5푑0.5휀̇0.1505 exp ( ) 푓표푟 푇 > 1080 °퐶 (12) 푝 0 8.314푇 A temperature range was specified because deformation mechanisms governing peak strain may change with temperature and each temperature regime should be fitted— subsolvus, near solvus, and supersolvus. As this model was only intended demonstrate feasibility of modeling the single temperature regime of supersolvus behavior, only this single equation was derived.

6.1.2 Strain for 50% Dynamic Recrystallization

The next required component for the JMAK formulation for DRX was the strain for 50% dynamic recrystallization. This number was experimentally determined. EBSD maps from the as-deformed and quenched double-cone compression specimens were evaluated and correlated with their imposed strain using a specifically tailored finite element model. The observed sigmoidal dependence of the area fraction recrystallized versus input strain confirmed the traditional Avrami approach could be used (Equation 126

13). Here, the fraction recrystallized, X, was dependent on the input strain, e, the strain for 50% recrytallization, e0.5, and the Avrami exponent, n.

휀 푛 푋퐷푅푋 = 1 − 푒푥푝 [− 푙푛 2 ( ) ] (13) 휀0.5

For a given temperature and strain rate, e0.5 and n were constants. Once these are found, Equation 13 can be used. Despite being very rapid to recrystallize, lower fractions of recrystallization were still observed at low strains and this measurement was possible using the 3:1 upset double-cone tests. The strain to 50% DRX was determined by examining the double-cone compression test recrystallization data (Figure 75) and recording the corresponding strain where 50% DRX was observed for each test condition.

This data was then fit to an equation of the general form (Equation 14) where ‘a’ was a fitting constant, d0 was the as-heated grain size, h was a sensitivity factor, m was a strain rate sensitivity factor, Q was the activation energy, R was the gas constant, T was the temperature and c was an offset constant.

푄 휀 = 푎푑ℎ휀̇푚 푒푥푝 ( ) + 푐 (14) 0.5 0 푅푇 One possible solution for the strain for 50% recrystallization for the supersolvus wrought cobalt data was determined (Equation 15).

69530 휀 = 7.17 ∗ 10−4푑0.5휀̇01505 푒푥푝 ( ) for T > 1080 °C (15) 0.5 0 8.314푇 Again, this fit should be accomplished for each temperature regime as the mechanisms and recrystallization behavior may change with temperature. Similar to the peak strain, this equation was also only fit for the supersolvus deformation data.

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6.1.3 DRX as a Function of Imposed Strain

An expanded formulation of the Avrami formulation for DRX was used and included more tuning parameters, such as d and an offset for the critical strain, ec, required to initiate DRX (Equation 16).

푛 휀 − 휀푐 푋퐷푅푋 = 1 − exp [−훽푑 ( ) ] (16) 휀0.5

For the original (Equation 13), d = ln (2) and the critical strain, ec = 0. Another way of writing ec is shown in Equation 17. The generally accepted value of a = 0.8, if it is used [60]. For this formulation it was not included due to lack of low-strain data to confirm whether a critical strain for the onset of recrystallization was observed. Other formulations for wrought materials did not include an offset term[41, 52], but was included for an ingot material [38].

휀푐 = 푎휀푃 (17) The Avrami exponent, n, was determined by taking the logarithm of Equation 13.

The result was plotted such that n was the slope of the line, or n = 1.18 (Figure 77).

128

2.5 2.0 y = 1.1755x - 0.4381 1.5

1.0 X)]}

- 0.5 0.0 -0.5

ln{ln[1/(1 -1.0 -1.5 -2.0 -2.5 -1.0 -0.5 0.0 0.5 1.0 1.5 2.0 2.5

ln (e/e0.5)

Figure 77. Recrystallization data plotted to determine, n, or the slope of the line.

Using the supersolvus constants determined from the mechanical testing and analysis, the final formulation correlating strain input to area fraction DRX becomes

Equation 18 and was the version used for this work.

휀 1.18 푋퐷푅푋 = 1 − exp [−0.697 ( ) ] 푓표푟 푇 > 1080 °퐶 (18) 휀0.5 6.1.4 DRX Grain Size

The final relationship to be determined for predicting dynamic recrystallization and the impact on microstructure evolution was the size of the dynamically recrystallized grains. In this empirical formulation, the DRX grain size is only a function of Zener-

Holloman parameter, but can be fit to Equation 19. Here, a is a constant, d0 is the as- heated grain size and h the sensitivity factor, e is the strain and n the related fitting factor,

129 edot is the strain rate, m is the sensitivity, Q is the apparent activation energy, R is the gas constant, and T is the temperature.

푄 푑 = 푎푑ℎ휀푛휀̇푚 푒푥푝 ( ) (19) 퐷푅푋 0 푅푇 The complexity was reduced such that 3 constants were needed, a, m, and Q. The value of m was determined by plotting ln (Z) versus ln(diameter) for grain sizes as a function of radial location and thus imposed strain (Figure 78). The slope of the fitted line on the data produced m = -0.0758 using an activation energy of 462 kJ/mol.

Decreasing the activation energy to 111,000 kJ/mole reduced m to approximately 0.0324.

2.25

1.75

ln (diameter) ln 1.5

1.25 32 33 34 35 36 37 ln (Z)

Figure 78. Plot of ln (Z) versus ln (diameter) in microns from the double cone upset testing. The slope provides the value of m for Equation 19.

The other constants were fit to the data such that the final dynamically recrystallized grain size in microns was reflected in Equation 20.

−35019 푑 = 84.75휀̇−0.0758 exp ( ) (20) 퐷푅푋 8.314푇 130

6.2 Grain Growth/Coarsening

After recrystallization is complete, grain growth can occur and is a significant driver in the microstructure evolution. The growth depends on temperature and time, and the presence and distribution of grain boundary pinning phases. Static grain growth, or growth without concurrent deformation, has been fitted using the empirical Equation 21, where dg is the diameter of the grown grain, d0 is the starting diameter, m is a growth

1 −푄 ⁄푚 푑 = [푑푚 + 푎푡 exp ( )] (21) 𝑔 0 푅푇 exponent, a is a scaling constant, t is the time, Q is the activation energy, R is the gas constant and T is the temperature. Short duration grain growth data was fit to Equation

21 using data from the heat treatments conducted as part of §4.3 and §5.3 and resulted in

Equation 22 where the units for diameter were microns, t was seconds, and T was Kelvin.

Growth exponents of up to 10 have been reported for steels [63] but value of 3 was used for wrought Waspaloy [52] and fit the data was reasonable for use here also.

1 −585000 ⁄3 푑 = [푑3 + 8.13 ∗ 1022푡 푒푥푝 ( )] (22) 𝑔 0 8.314푇 The fully-recrystallized grain size reached a stable, pinned size at times greater than 120 min and is represented by Equation 8. The interdendritic eutectic-type phase was limiting the growth instead of carbides or other secondary precipitated pinning phases like gamma-prime. Both 1149 °C and 1204°C produced the same pinned grain size after 2 hours, so further testing is needed to help differentiate the behavior for differences in temperature and using longer times.

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6.3 Wrought JMAK Fit Model Results

The wrought microstructure model fits for dynamic recrystallization and grain growth were successfully developed. The fitted equations, as developed, were point solutions for specific input conditions. In addition to developing a spread-sheet tool, the

JMAK fits were successfully incorporated into a finite element framework [46] to enable the prediction of location specific recrystallization and growth behavior over a range of supersolvus temperatures. The double cone finite element simulations were re- accomplished with the microstructure data and successfully predicted differences in microstructure evolution behavior for the different testing conducted.

Comparison of the computed peak strain between the model and observed experimental flow stress data showed good agreement (Table 11). with the uncertainty introduced with the foil-stack for the upset compression specimens used in the fit. The peak strain was not explicitly used in the fit due to excluding the critical strain for recrystallization, so any error here did not propagate in the subsequent results.

Fit ep Experiment ep 1149 °C 0.01/s 0.037 0.035 1149 °C 0. 1/s 0.052 0.038 1204 °C 0.01/s 0.029 0.022 1204 °C 0. 1/s 0.041 0.032 Table 11. Comparison of peak strains of the JMAK fit with experiment.

Alternatively, the agreement of the JMAK fit with the experimental values for the strain to 50% recrystallization were good for the low temperature, high strain-rate case and for the high temperature, low strain rate case (Table 12). These were the slowest and fastest to recrystallize, respectively, bounding the behavior. The fit strongly under- 132 predicted e0.5 for the high temperature, low strain-rate case while over predicting the e0.5 for the low temperature, low strain rate case. The discrepancy was undesirable as the strain to 50% recrystallization anchors the sigmoidal curve in strain-space.

Fit e0.5 Experiment e0.5 1149 °C 0.01/s 0.39 0.47 1149 °C 0.1/s 0.55 0.53 1204 °C 0.01/s 0.31 0.29 1204 °C 0. 1/s 0.44 0.29 Table 12. Comparison of calculated JMAK fit values for strain to 50% recrystallization with experiment.

The JMAK model fit for fraction recrystallized overall was acceptable, mirroring many of the characteristics of the experimental data (Figure 79). Similar to the strain to

50% recrystallization data, the fit captured the bounding behavior curves including the low temperature, high strain rate and the high temperature low strain rate well. Also captured were the typical behavior of increasing fraction recrystallized with temperature and decrease in strain rate, as well as the observed changes in recrystallization rate with increasing strain.

133

1.0 0.9 0.8 0.7

0.6 × 0.01/s - Experiment 0.5 0.1/s - Experiment 0.4 ⸱ 0.01/s - JMAK Fit 0.3 0.1/s - JMAK Fit

0.2 1149 °C Fraction Recrystallized 0.1 1204 °C 0.0 0.0 0.5 1.0 1.5 2.0 Strain

Figure 79. JMAK fit versus experimental data for fraction recrystallized versus strain.

Unlike a typical sigmoidal curve with a slow initial rate of recrystallization followed by an increase in recrystallization rate, the fit for cobalt did not exhibit much of this behavior. The early onset of recrystallization precluded this behavior due to the decision to not include the critical strain for recrystallization in the formulation of the recrystallized fraction (Equation 16). Lack of recrystallization data for low imposed strains, e < 0.3, precluded an assessment of whether there was a true incubation or necessary imposed strain prior to the onset of dynamic recrystallization to warrant this term being included. Had it been included and fit to the standard Equation 17, the critical strain, ec, is typically on the order of 80% of the peak strain, ep [60]. For the low peak

134 strains, ep, observed for the cobalt, any critical strain offset was unlikely to make a substantive difference on recrystallization behavior. The fit determined for Waspaloy did not incorporate a separate critical strain term but rather observed the sigmoidal nature of the curve reflected by the minor offset to the strain required for recrystallization was captured in the sigmoidal-shape of the JMAK fit [41].

The fit over-predicted the recrystallized fraction for the low temperature, low strain rate. The fit also strongly under predicted the recrystallization behavior for the higher temperature, low strain-rate behavior. This was a result of a fitting compromise to make the formulation self-consistent. The strain rate fitting constant and the effective activation energy had to both be varied to fit the behavior, but were limited as they were dependent on each other. Changing the sensitivity term changed the spacing between the strain rate curves for a given temperature, whereas changes to apparent activation energy changed the overall spread between similar temperatures for a given strain rate. The trends in Figure 79 suggested that a higher apparent activation energy was needed while also requiring a smaller strain-rate sensitivity factor to decrease the change in behavior with strain rate. Unfortunately, due to the assumed derivation of these two factors, they were unable to be separately manipulated.

If the scaling constraint between the strain rate sensitivity term and the apparent activation energy term was removed, the JMAK fit would be improved significantly.

This formulation of the fit required Qapp = 1757 kJ/mol and m = 0.1. While the value of m was reasonable, the apparent activation energy term was much higher than expected.

Quantitative comparisons between simulation and experiment at the equatorial mid-plane

135 of the double cone upset tests were promising (Figure 80). At the lower temperature, the model successfully captured the differences in experimental response for the different strain rates. At higher temperature, the experimental recrystallization data for both temperatures was similar and the model appeared to exaggerate the difference. The lack of difference in experimental data was attributed to a combination of variability in sample preparation, measurement, and analysis errors on a single upset sample and would benefit from additional data. Even so, the model was successful in demonstrating the capability but would benefit from additional data and analysis to determine why the original fit derivation was less successful.

1.0 0.9 0.8 0.7 0.6 × 0.01/s - Experiment 0.5 0.1/s - Experiment 0.4 ⸱ 0.01/s - JMAK Fit 0.3 0.1/s - JMAK Fit

0.2 1149 °C Fraction Recrystallized 0.1 1204 °C 0.0 0.0 0.5 1.0 1.5 2.0 Strain

Figure 80. Improved JMAK fit for DRX by removing constraints to apparent activation energy.

136

Utilizing the adjusted JMAK formulation within the finite element software, the imposed strain, fraction recrystallized, and average grain size were calculated for the double-cone upset tests (Figure 81). These results showed the imposed strain was slightly higher in the 0.01/s strain rate specimens than was the case for 0.1/s strain rate specimens. Also observed was higher overall dynamic recrystallized fractions with higher temperature, as well as increasing recrystallization for a given temperature with decreasing strain rate. The simulated starting grain size was specified at 40 m. Given

2.25

0.0 1.0

0.0 40.0

0.0

Figure 81. Half-symmetry double-cone upset simulation results for Von Mises strain, recrystallized fraction, and average grain size.

the higher fraction recrystallization in the higher temperature tests and lower strain rates, the smaller, as-recrystallized grains decreased the average grain size in areas where dynamic recrystallization was predicted more prevalent. The dead-metal zone at the 137 center-top and center-bottom formed during upset showed a larger fraction of the remnant, unrecrystallized grains and a corresponding larger predicted grain size.

Similarly, the outer diameter which saw less deformation and hence experienced less recrystallization also displayed a larger average grain size. These trends all agreed with the observations from experimental testing.

6.4 Comparison to a Wrought Nickel-Base Superalloy

The model fits developed for a gamma-prime strengthened cobalt superalloy were compared to those developed for wrought nickel-base superalloy, Waspaloy [41, 52].

Equations for both materials were summarized in Table 13. The nickel fit equations were developed in terms of Zener-Hollomon parameter, or temperature compensated strain rate, Z. The cobalt equations were modified and presented in the same manner so they could be easily compared.

138

Wrought Cobalt-Base Superalloy Wrought Nickel-Base Superalloy Supersolvus Dynamic Recrystallization Supersolvus Dynamic Recrystallization 푄 푄 푍 = 휀̇ 푒푥푝 ( ) 푍 = 휀̇ 푒푥푝 ( ) 푅푇 푅푇 For Q = 462 kJ/mol, T (K) For Q = 468 kJ/mol, T(K) −5 0.5 0.1505 −4 0.54 0.106 휀푝 = 3.227 ∗ 10 푑0 푍 휀푝 = 1.685 ∗ 10 푑0 푍 −4 0.3 0.1505 0.29 0.04 휀0.5 = 7.17 ∗ 10 푑0 푍 휀0.5 = 0.035푑0 푍

휀 1.18 휀 1.8 푋퐷푅푋 = 1 − exp [−푙푛2 ( ) ] 푋퐷푅푋 = 1 − exp [−푙푛2 ( ) ] 휀0.5 휀0.5

−0.0758 −0.0456 푑퐷푅푋 = 84.75푍 푑퐷푅푋 = 108.85푍 For Q = -462 kJ/mol, d (m) For Q = -468 kJ/mol, d (m)

Grain Growth – Short Time Grain Growth – Short Time 1 1 −푄 ⁄3 −푄 ⁄3 푑 = [푑3 + 8.13 ∗ 1022푡 푒푥푝 ( )] 푑 = [푑3 + 2 ∗ 1026푡 푒푥푝 ( )] 𝑔 0 푅푇 𝑔 0 푅푇 For Q = 585 kJ/mol, d (m), t (s), T(K) For Q = 595 kJ/mol, d (m), t(s), T (K) Table 13. Summary of microstructure fits for cobalt- and nickel-base [41] superalloys.

As expected, when exercised the fits predict similar recrystallization behavior.

The nickel equation did not use a critical strain for recrystallization (Equation 17). After evaluating the equation fits, this term was not used for the cobalt either. As such, both materials were predicted to start recrystallization immediately with strain. Both predicted increased fraction recrystallized with increasing strain, but the rate of recrystallization was higher for higher temperatures in the cobalt. For grain growth, the data used to fit the terms showed very little grain growth for short times and is reflected in the scale of the constant but were otherwise determined similar. For that determination, an apparent activation energy was chosen for cobalt, 585 kJ/mol, to mirror the roughly 125% increase 139 documented in the Waspaloy between the activation energy for deformation and grain growth [41].

One of the most notable similarities observed was in the apparent activation energy for deformation between the two alloys, with Waspaloy at 468 kJ/mol while cobalt was experimentally determined at 462 kJ/mol. As expected, the 462 kJ/mol was higher than that found for deformation of pure cobalt, 254 kJ/mol [64, 65] and the apparent activation energy for self-diffusion in face-centered-cubic cobalt at 270 kJ/mol

[66]. With the apparent activation energy for deformation of cobalt determined experimentally, fitting of the cobalt model was made easier. Without this as a known constant, a significant potential error in both activation energy and fitting parameters would be possible as changes to one constant could offset or mask changes to the others in the formulation.

The fit for strain to 50 % recrystallization, e0.5, showed the largest difference to the formulation for Waspaloy both in fitting parameter for the Zener-Hollomon parameter, Z, and fitting constant. Given the other similarities observed, this may be attributed to the lack of true low-strain, low-area fraction recrystallized values. It was shown that the strain rates for the double-cone geometries at the low strain regions were under target (Figure 72), which would have the effect of confounding the fitting.

Otherwise, similarities were apparent for the fits. Functionally this observed similarity in model fits reflected the similar microstructure trends were observed in both cobalt and

Waspaloy.

140

A small difference observed was in the experimentally calculated Avrami- exponent, ‘n.’ The value determined for ’ cobalt superalloy was closer to n = 1.2 while the value for wrought Waspaloy at the same conditions was n = 1.8. Given the high fraction of dispersed intermetallic particles, it was expected that particle stimulated nucleation would factor strongly in the kinetics of the microstructure evolution. Goetz

[24] had shown that in coarse-grained nickel superalloys, PSN raised the effective

Avrami exponent to n = 3, or what would be expected for homogeneous site saturation

[67]. However, Cahn [68] had shown that for cases of high volume fractions recrystallized nucleation site saturation may be delayed and the Avrami exponent would drop from 4 closer to 1, as observed in this work. Also, prior work for austenitic stainless steels showed n=1.3 while exhibiting classical necklace recrystallization [69]. Thus, the n=1.2 Avrami exponent observed in the current work for fine-grain cobalt alloy was in- line with prior work and indicated nucleation at particles was not driving the behavior as thought but instead was dominated by traditional necklace recrystallization at grain boundaries with perhaps a small contribution from PSN and was supported by the significantly higher fraction of prior grain boundaries versus particles in the wrought condition. Further, it would be illustrative to measure low strain, low fraction recrystallization data and determine if inclusion brought the Avrami exponent more in line with the value determined for wrought Waspaloy.

The standard JMAK fit for cobalt and Waspaloy were compared and the fraction recrystallized plots graphically illustrated many of these differences. Using two supersolvus temperatures (T’ solvus +50 °C and T’solvus + 150 °C), two strain rates (0.01/s

141 and 0.1/s), and one starting grain size, fits for each material were plotted over a range of strain (Figure 82). Only minor differences in recrystallization rate/slope were observed for strains up to e=0.5. Beyond that, the recrystallization rate decreased in the cobalt whereas a similar decrease in rate did not occur until much later in the Waspaloy. The proximity of the curves to one another indicated that cobalt has a higher sensitivity to temperature and strain rate than Waspaloy exhibited, but that may be an artifact of the formulation. The initiation of DRX occurred much more rapidly for the cobalt than was observed for the Waspaloy. Despite this, the cobalt required higher strain input for conditions other than the high temperature, low strain rate to achieve the same fraction recrystallized.

1.0 0.9 0.8 0.7

0.6 0.01/s 0.5 0.1/s 0.4 Waspaloy ' + 50 °C 0.3 Waspaloy ' + 150 °C

0.2 Cobalt ' + 50 °C Fraction Recrystallized 0.1 Cobalt ' + 150 °C 0.0 0.0 0.5 1.0 1.5 2.0 Strain

Figure 82. JMAK fits for wrought ' cobalt and Waspaloy for similar test conditions.

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Industrially, the JMAK fit of the experimental data indicated the gamma-prime strengthened cobalt alloy could require less thermomechanical processing to achieve a more uniformly recrystallized microstructure over a wider range of deformation inputs than would be expected with a similar nickel-base alloy. The result would be less cost and time for manufacturing while maintaining or improving mechanical properties due to reductions in microstructure variability.

Overall, the JMAK-style empirical microstructure models attempted were successfully fit and implemented for a Co-Al-W -prime strengthened cobalt superalloy.

The equations fit using the experimentally-determined data were still representative of the correct trends but were less accurate than the fits accomplished where activation energy and the strain rate sensitivity factor were allowed to vary independently and to questionable values. This illustrated a key limitation of the JMAK fitting technique where the result may be acceptable, but may not be physically representative of the process or the material behavior. Despite that limitation, the fit models showed reasonable agreement with experimental results and were a valid first attempt at modeling supersolvus microstructure evolution in an alloy system where limited models existed previously.

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Chapter 7: Conclusions and Future Work

Industrially relevant thermomechanical processing of a Co-Al-W gamma-prime strengthened superalloy was completed. After characterization of the starting ingot material, three major thermomechanical processing stages were accomplished: ingot homogenization, ingot conversion, and wrought processing.

(1) Ingot homogenization – The required homogenization time was determined using

subscale coupons cut from the starting ingot. Homogenization of the ingot

material was accomplished in stages using a hot isothermal press unit.

(2) Ingot conversion – High temperature flow-stress behavior of the coarse-grain

ingot material was determined using cylindrical, hot isothermal compression

testing. The ingot material was successfully converted using a combination of

extrusion and post-extrusion recrystallization heat treatment.

(3) Wrought processing – High temperature flow-stress behavior of fine-grain

converted material was determined using cylindrical, hot isothermal compression

testing. Specially designed, double-cone compression test specimens were

successfully tested with subsequent grain growth heat treatments to provide data

on microstructure response as a function of imposed strain, strain-rate, and

temperature.

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Subsequent to the thermomechanical processing, JMAK-style microstructure models were developed using the experimental process and microstructure information extracted from the material manufactured during the effort. Fits for dynamic recrystallization and grain growth behavior were successfully incorporated into a commercial finite element software package and reflected acceptable behavior.

7.1 Conclusions

Several conclusions were drawn from executing an ingot metallurgy thermomechanical processing route for a novel gamma-prime strengthened Co-Al-W superalloy and the development and implementation of an empirical fit microstructural model. Comparisons to a traditional gamma-prime strengthened nickel-base superalloy,

Waspaloy, were also informative.

(1) Ingot metallurgy was demonstrated a viable thermomechanical processing route

for this novel class of gamma-prime strengthened cobalt-base superalloys, similar

to gamma-prime strengthened nickel superalloys. A full processing sequence,

complete with industrially relevant homogenization, ingot conversion, and billet

processing were all successfully completed and were shown as a viable means of

producing an alloy of this kind into a useable form for consideration in high

temperature turbine components.

(2) The homogenization time scaled with the square of the primary dendrite arm

spacing. The high tungsten content of this particular alloy and the low diffusivity

of tungsten required multiple days of homogenization time that, while feasible,

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would not be industrially palatable – particularly if done in an oxidizing

environment. Increasing the ingot size to production diameters required for

turbine components would increase the diffusion distance and increase the

homogenization time further.

(3) The macroscopic flow behavior of the coarse-grain ingot material during

deformation was similar to coarse-grain work on a production-scale cast

Waspaloy ingot. Strong deformation anisotropy was observed that correlated

with the starting ingot texture prior to upset. With lack of boundaries to nucleate

recrystallization in the coarse grain ingot material, particle stimulated nucleation

at the intermetallic phase in the cobalt, similar to particle stimulated nucleation at

carbides in Waspaloy, increased the rate of recrystallization observed during ingot

conversion. The workability permitted successful extrusion at multiple extrusion

ratios, and nominally full recrystallization was achieved with a short post-

extrusion heat treatment.

(4) The shape of the isothermal flow curve behavior was similar between the ingot

and wrought materials, but the wrought materials exhibited a higher strength in all

conditions. The similarity in flow curves suggested similar mechanisms

governing hardening and softening in both ingot and wrought materials and

compare well with the behavior exhibited by Waspaloy.

(5) The double-cone geometry designed for the present work provided a range of

imposed strain within one sample. Using this strategy reduced the number of

required high temperature mechanical tests which saved material, time, and

146

significant costs. While the imposed strain gradient was roughly linear across the

radius of the test specimen, the strain-rate gradient was not. As such, this type of

testing may not be well suited to materials that exhibit a high degree of strain-rate

sensitivity to microstructure response. Confounding the strain rate with imposed

strain effects that change with radial location made interpretation of the evolution

behavior more difficult as variability increased with increasing strain rate.

(6) The finite element simulations for each of the thermomechanical processes were

highly sensitive to friction coefficient used. The friction coefficient impacted

calculated strain and strain-rate predictions for the empirical models, which in-

turn altered the JMAK formulation fits.

(7) The fraction dynamic recrystallization followed the expected trend of increased

recrystallization with increasing temperature and decreasing strain rate. The

average recrystallized grain size did not change with imposed strain but did

increase with temperature. These findings were consistent with observations of

recrystallization in Waspaloy. Experimentally, the recrystallization behavior for

cobalt was more sensitive to temperature than strain rate, but the JMAK fit model

predicted high sensitivity to both strain rate and temperature.

(8) The developed JMAK-style empirical fit model was successful in reflecting the

supersolvus microstructure evolution due to dynamic recrystallization and grain

growth. The fitting constants were comparable to those used for a nickel-base

superalloy, Waspaloy. Specifically, the Avrami exponents for the wrought cobalt

were comparable with those observed for nickel-base alloys undergoing similar

147

thermomechanical processes. The demonstrated success indicated determining

similar empirical fit models for metadynamic and static recrystallization would

also be worthwhile.

(9) The presence of the intermetallic phase did not diminish the utility of the current

work to demonstrate the feasibility of an ingot metallurgy approach for

thermomechanical processing of gamma-prime strengthened cobalt superalloys.

The particles increased the rate of recrystallization and improved grain refinement

during ingot conversion and were considered advantageous similar to carbides in

nickel superalloys for those reasons. However, the intermetallic phase may be

deleterious to mechanical properties and compositions should be revised in future

work to understand its impact and possibly avoid formation during solidification.

(10) The oxidation behavior of the Co-Al-W-Ta quaternary alloy used in this work at

hot working temperatures was poor. Steps were taken during the research to

mitigate the impact of the oxidation during various thermomechanical processes

and preparation, but using an alloy with such strong reactivity would not be

industrially prudent. Strong reactivity with other substances like borosilicate-

containing lubricants were also discouraging, but changes to the alloy should

reduce that effect.

(11) While not reported in literature of similar compositions, magnetic effects in this

alloy were strong. This was successfully countered by strategies including thinly

sectioning samples for electron microscopy or by using commercially available

demagnetizing hardware.

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(12) The demonstrated similarity of the behaviors exhibited by the gamma-prime

strengthened cobalt alloy to a structurally similar gamma-prime strengthened

nickel alloy, Waspaloy, was promising. Waspaloy is a historical alloy and well

understood industrially, so similarity with the cobalt suggests that currently

available and conventional machines, processes, and tooling could be used for

manufacturing. Utilizing updated, well-accepted computational modeling tools

for the nickel-base alloys will enable process- and location-specific simulation

and optimization which will enable more rapid transition and implementation of

this novel alloy class.

7.2 Future Work

While this work was able to demonstrate the potential for traditional hot-working of cobalt-base superalloys and the applicability of existing empirical microstructure evolution models, several opportunities for future work in the area remain to further the development of gamma-prime strengthened cobalt superalloys and transition to an eventual system.

(1) Additional alloy development work is warranted to improve the alloy for

properties desirable for use as turbine engine disks. The alloy composition

selected for this work was a published, generic composition that, at the time, was

considered novel and suitable for further research work. However, compared to a

modern nickel-base superalloy, the relatively simple quaternary composition left

many properties, such as oxidation behavior, with significant room for

149

improvement. Due to the pace of research and significant interest in this area,

much progress has already been made in improving the viability of cobalt

superalloys through compositional changes especially in light of the cost volatility

and sourcing of cobalt itself.

(2) While the work presented here was significant in that it demonstrated a

representative processing flow path for a model cobalt superalloy, process and

tool development work is still needed. A major constraint on testing for this

program was the safe capability limitations of the tooling and will be a similar

constraint for industry. As was experienced during the course of this work, not all

of the supply chain was ready to go hotter. Lubricants, coatings, tooling,

furnaces, etc. all have challenges that need addressed before there can be

widespread adoption. The large infrastructure supporting today’s state-of-the-art

nickel superalloys will need to be replicated for even higher temperatures and will

be a significant undertaking.

(3) The higher recrystallized fractions were promising for showing propensity for

dynamic recrystallization, but unfortunately meant a lack of lower fraction

recrystallization data to quantify the lower strain region of the empirical fit.

Completing lower reduction testing would provide this missing information and

extend the applicable strain range of the fits and provide insight into the initial

nucleation behavior of the material, whether driven by particles or grain

boundaries. This, coupled with a more thorough investigation of the apparent

activation energy would be useful in reducing uncertainty of the empirical fits.

150

(4) Some gamma-prime strengthened alloys contain less desirable phases, similar to

the Co-Al-W-Ta alloy in this work. Characterization and quantification of the

metastable interdendritic eutectic phase and the sensitivity to changes

composition would be useful to reducing its occurrence. If unable to reduce the

occurrence of the interdendritic phase, an understanding of the size and

morphology stability through thermal and thermal-mechanical sequences would

be useful for assessing impact to mechanical properties.

(5) Higher strain-rate mechanical and microstructure response data would be useful to

facilitate future research on solid state joining, or higher strain-rate forming

methods used later in disk or component manufacture.

(6) The magnetic properties of cobalt make it popular for use in magnets, and is

among its primary uses. Precautions were taken to reduce the impact of the

magnetic tendencies of the material during characterization like utilizing thin

sections to reduce magnetic field and utilizing demagnetizing equipment prior to

examination in the electron microscopy equipment. The high Curie temperature

suggests that while not magnetic during high temperature forming operations, the

material could still exhibit magnetic tendencies at expected service temperatures.

This brings the possibility of utilizing a cobalt superalloy disk as a true multi-

functional material, perhaps for on-board power generation natively within the

turbine instead of relying on auxiliary power units. Weight, complexity, and

performance benefits make this worth additional investigation.

151

(7) Development of a better friction model or additional testing would allow more

representative formulation of the friction conditions for the foil and lubricants

used during this work and implementation in the finite element framework. This

would reduce the uncertainty in the analysis and better reflect the high

temperature test behavior of cobalt.

(8) Early during the project, the wavelength dispersive spectroscopy (WDS)

equipment in the characterization facility experienced a significant system failure

and was not available for detailed spatial-composition mapping. Conducting a

more thorough investigation of the segregation of elements and their distribution

and mobility as a function of time with heat treatments would have gone a long

way to reducing the uncertainty in the original analytical determinations for

homogenization time. As ingots get larger having a detailed analysis would be

beneficial to understanding the homogenization behavior.

(9) The upper temperature limit of the mechanical testing was largely chosen out of a

need to utilize existing tooling and instrumentation proven up to 1232 °C. The

cobalt mechanically is capable of much higher. Procurement of higher-

temperature tooling, such as silicon-nitride or silicon-carbide, in large-enough

sizes for die inserts would allow similar mechanical evaluation at much higher

temperatures and investigation of the impact of higher temperatures on

microstructure evolution.

(10) Lastly, as the current work only demonstrated the capability, it would be

worthwhile to investigate any future selected alloy by expanding to sub-solvus

152 evaluations, near solvus evaluations, and full-supersolvus temperature ranges.

This data would fully populate the temperature space of expected thermomechanical processes, the resultant properties, and could be used to develop robust microstructure models on a future generation gamma-prime strengthened cobalt superalloy.

153

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