Friction Stir Processing Nickel-Base Alloys

THESIS

Presented in Partial Fulfillment of the Requirements for the Degree Master of Science in the Graduate School of The Ohio State University

James R. Rule, B.S.W.E.

Graduate Program in Welding Engineering

The Ohio State University

2011

Master's Examination Committee:

Professor John C. Lippold, Advisor

Professor Sudarsanam Suresh Babu

James R. Rule

2011

ABSTRACT

The application of friction stir processing (FSP) on three Ni-base alloys, Alloy

625, Alloy 718, and Hastelloy X utilizing a tungsten-rhenium tool was investigated. A processing window for each alloy was defined in terms of travel speed and tool rotation rate. Process forces and thermal histories were successfully recorded for each alloy.

Peak process temperatures were recorded at 1150ºC for Alloy 625 and Hastelloy X and

1100ºC for Alloy 718 using embedded type K thermocouples. FSP microstructures were investigated using optical and electron microscopy. Significant grain refinement was experienced by each alloy with stir zone grain sizes of 6µm, 5µm, and 4µm for Hastelloy

X, Alloy 625, and Alloy 718, respectively, relative to starting grain sizes of 88µm, 26µm, and 44µm. Optical microscopy revealed streaking on the advancing side of the stir zone for each alloy. Scanning electron microscopy (SEM) and energy dispersive spectroscopy

(EDS) analysis showed these bands to be a result of tool wear. Hardness traverses show that all three alloys experience peak hardness outside the stir zone in the thermomechanically-affected zone. Additionally, it was observed that Alloy 718 and

Hastelloy X experienced hardening in the stir zone in comparison to the base metal whereas Alloy 625 underwent softening. Longitudinal sub-size tensile specimens extracted from stir zones showed that Hastelloy X experienced significant strengthening

ii due to FSP relative to the base material whereas Alloy 625 and Alloy 718 experienced virtually no improvement in strength in the FSP region relative to base material.

FSP was further investigated as a means to reduce the susceptibility of these alloys to heat-affected zone (HAZ) liquation cracking. Spot-varestraint testing was used to evaluate the effects of FSP on HAZ liquation in terms of total crack length (TCL) and maximum crack length (MCL). Testing showed a reduction in HAZ liquation susceptibility due to the reduction in HAZ grain size. Optical and electron microscopy evaluation revealed that HAZ liquation cracking resistance is enhanced in the FSP microstructure due to increased grain boundary area resulting in finer, more discontinuous networks of low melting eutectic and second phases.

The modified Gleeble® hot torsion test was utilized to simulate FSP microstructure and investigate shear stress and shear strain encountered by the material.

The thermomechanically-affected zone was successfully reproduced for all three alloys, however, stir zone microstructures could not be simulated. Shear stress and shear strain were calculated for each simulation. Examination of the stress-strain curves generated for each simulation showed that a strong correlation between dynamic recrystallization

(DRX) and the stress-strain signature. From observation of the curves it can be determined if DRX has occurred.

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To my bride to be

You truly are my better half.

Acknowledgments

First, I would like to thank my advisor Dr. John Lippold for providing the opportunity further my education and his guidance and support with this research.

Thank you to Jeff Rodelas for continued help with electron microscopy and friction stir knowledge. Also, thanks to the members of the Welding and Joining

Metallurgy Group and fellow welding engineering graduate students for both academic discussions and distractions.

Thank you to Dr. Sudarsanam Suresh Babu for his review of this document and as a member of my examination committee.

Thank you to Rollie Dutton and the Air Force Research Labs for funding of this research and to Pete Ditzel and Parker Hannifin for materials and insights which may have never been realized.

Finally, I would like to thank my soon to be wife, Sarah, for putting her dreams on hold while I attain mine and my parents for their love and support.

Project supported by the Air Force Research Lab through Universal Technology Corp.

Grant #09-S568-067-01-C1

Vita

May 1, 1986 ...... Born - Norwalk, OH U.S.A.

2009...... B.S. Welding Engineering

The Ohio State University

Columbus, OH

2009-Present ...... Graduate Research Associate

The Ohio State University

Columbus, OH

Field of Study

Major Field: Welding Engineering

TABLE OF CONTENTS

ABSTRACT ...... ii

Acknowledgments...... v

Vita ...... vi

List of Tables ...... xi

List of Figures ...... xiii

CHAPTER 1: INTRODUCTION ...... 1

CHAPTER 2: BACKGROUND ...... 4

2.1 NICKEL ALLOYS ...... 4

2.1.1 Ni-base alloy microstructure ...... 4

2.1.2 Ni-base alloy weldability ...... 6

2.2 FRICTION STIR PROCESSING ...... 13

2.2.1 Previous studies of friction stir welding/processing Ni-base alloys ...... 16

2.3 PHYSICAL SIMULATION ...... 28

CHAPTER 3: OBJECTIVES ...... 33

CHAPTER 4: RESULTS AND DISCUSSION ...... 34

4.1 FRICTION STIR PROCESSING NICKEL-BASE ALLOYS ...... 34 vii

4.1.1 Processing parameters and windows ...... 34

4.1.2 Microstructure ...... 38

4.1.3 Thermal History ...... 41

4.1.4 Process Forces ...... 43

4.1.5 Mechanical Properties ...... 44

4.2 FRICTION STIR PROCESSING NICKEL-BASE ALLOYS FOR

IMPROVING HEAT-AFFECTED ZONE LIQUATION CRACKING RESISTANCE

4.2.1 Spot-varestraint Testing ...... 46

4.2.2 Microstructure and Grain Size Effect on HAZ Liquation Cracking ...... 49

4.3 PHYSICAL SIMULATION OF FRICTION STIR PROCESS

MICROSTRUCTURE ...... 51

4.3.1 Thermal Control ...... 51

4.3.2 Microstructure ...... 53

4.3.3 Shear Stress-Strain Behavior ...... 54

CHAPTER 5: CONCLUSIONS ...... 56

REFERENCES ...... 58

APPENDIX A ...... 63

MICROSTRUCTURE AND MECHANICAL PROPERTIES OF THREE FRICTION

STIR PROCESSED NICKEL-BASE ALLOYS ...... 63 viii

A.1 ABSTRACT ...... 64

A.2 INTRODUCTION ...... 65

A.3 EXPERIMENTAL PROCEDURE ...... 70

A.4 RESULTS AND DISCUSSION ...... 74

A.4.1 Process Parameter Windows ...... 74

A.4.2 Thermal History ...... 78

A.4.3 Process Forces ...... 81

A.4.4 Microstructure ...... 90

A.4.5 Mechanical Properties ...... 99

A.5 CONCLUSIONS ...... 104

APPENDIX B ...... 109

APPLICATION OF FRICTION STIR PROCESSING FOR IMPROVING HEAT-

AFFECTED ZONE LIQUATION CRACKING RESISTANCE OF NICKEL-BASE

ALLOYS ...... 109

B.1 ABSTRACT ...... 110

B.2 INTRODUCTION ...... 110

B.3 EXPERIMENTAL PROCEDURES ...... 111

B.4 RESULTS AND DISCUSSION ...... 117

B.5 CONCLUSIONS ...... 140

B.6 REFERENCES ...... 141

APPENDIX C ...... 144

C.2 INTRODUCTION ...... 145

C.3 EXPERIMENTAL PROCEDURE ...... 149

C.4 RESULTS AND DISCUSSION ...... 155

C.4.1 Thermal History Acquisition ...... 155

C.4.2 Modified Gleeble Hot Torsion Test ...... 159

C.5 CONCLUSIONS ...... 174

C.6 RECOMMENDATIONS ...... 175

C.7 REFERENCES ...... 175

List of Tables

Table 2.1: Nominal compositions of three Ni-base alloys used in this study [7]...... 6

Table 4.1: List of selected FSP runs (non-inclusive)...... 37

Table 4.2: Summary of grain refinement due to FSP ...... 40

Table 4.3: FSP Parameters ...... 47

Table 4.4: Spot-varestraint test parameters...... 47

Table 4.5: EDS quantification of intergranular and matrix areas...... 50

Table A.1: Alloy compositions...... 71

Table A.2: List of selected FSP runs...... 75

Table A.3: Comparison of process forces, material strength, and heat input/index...... 89

Table A.4: Quantification of EDS scans showing matrix to be absent of tungsten and bright particles to consist primarily of tungsten...... 99

Table B.1: Material compositions...... 112

Table B.2: Friction stir processing parameters...... 112

Table B.3: Spot-varestraint test parameters...... 113

Table B.4: ThermoCalc® simulation results...... 123

Table B.5: EDS quantification of matrix and intergranular (IG) film regions resulting from spot-varestraint testing...... 132 xi

Table C.1: Alloy compositions...... 150

Table C.2: FSP parameters...... 153

Table C.3: Hot torsion test parameters...... 154

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List of Figures

Figure 2.1: The FCC unit cell (a) hard sphere representation, (b) reduced sphere representation. Adapted from Callister, Material Science and Engineering - An

Introduction, 7th Edition, 2007[9] ...... 5

Figure 2.2: Illustration of fusion weld and related regions: (A) fusion zone, (B) partially melted zone, (C) heat-affected zone, (D) base metal, (E) original joint edge of materials. 8

Figure 2.3: Mechanisms of HAZ liquation cracking (left) the segregation mechanism,

(right) the penetration and constitutional liquation mechanism. Courtesy J.C. Lippold .. 10

Figure 2.4: Effect of HAZ grain size on HAZ liquation cracking (Alloy 718) [14]...... 12

Figure 2.5: Effect of grain size on HAZ liquation cracking in terms of (a) total crack length and (b) maximum crack length (cast Alloy 718) [18]...... 12

Figure 2.6: Illustration of friction stir processing denoting the tool features (top) and the determination of advancing and retreating side (bottom)...... 14

Figure 2.7: Schematic of typical FSP showing four microstructurally distinct regions:

(A) stir zone, (B) thermomechanically affected zone, (C) heat-affected zone and (D) unaffected base material...... 15

Figure 2.8: Microhardness profile of Alloy 600 FSW reported by Ye et al (18)...... 17

Figure 2.9: Tensile properties of Alloy 600 FSW reported by Ye et al (18)...... 17

Figure 2.10: Transverse tensile test results for Alloy 600 FSW [23]...... 19

Figure 2.11: Microhardness profile of Alloy 600 FSW [23]...... 19 xiii

Figure 2.12: Summary of ferric sulfate-sulfuric acid corrosion test showing reactivation current ratios. Higher current ration denotes increased corrosion susceptibility [23]. .... 20

Figure 2.13: Effect of welding speed on stir zone grain size for Alloy 600[24]...... 21

Figure 2.14: Effect of welding speed on stir zone thermal history for Alloy 600[24]. ... 22

Figure 2.15: (left) Effect of welding speed on stir zone hardness for Alloy 600. (right)

Hall-Petch relationship between grain size and hardness for Alloy 600 [24]...... 22

Figure 2.16: Effect of welding speed on Alloy 600 tensile properties [24]...... 23

Figure 2.17: Thermal history for (a) Alloy 625 FSW and (b) Alloy 718 FSW (21; 22). 25

Figure 2.18: (a) Top view of Alloy 625 FSW, (b) banding structure observed in Alloy

625 SZ due to tool wear, (c) top view of Alloy 718 FSW, and (d) banding due to tool wear in SZ of Alloy 718 (21; 22)...... 25

Figure 2.19: Microhardness profiles of as welded and PWHT FSW in (a) Alloy 625 and

(b) Alloy 718 (21; 22)...... 26

Figure 2.20: Tensile testing summary for as welded and PWHT FSW (a) Alloy 625 and

(b) Alloy 718 (21; 22)...... 26

Figure 2.21: Optical micrographs showing response to HAZ liquation cracking in (a) friction stir processed material experiencing no cracking and (b) as cast Alloy 738 experiencing cracks varying from 20-110µm[27]...... 27

Figure 2.22: Type 310 SZ (a) actual FSW and (b) modified hot torsion test at 1000ºC and

1150 RPM[6]...... 30

Figure 2.23: Type 310 TMAZ (a) actual FSW and (b) modified hot torsion test at 1100ºC and 550 RPM[6]...... 31

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Figure 2.24: Type 310 HAZ (a) actual FSW and (b) modified hot torsion test at 1100ºC and 850 RPM [6]...... 31

Figure 2.25: Proposed shear stress-strain signatures indicating the occurrence of (a) discontinuous dynamic recrystallization and (b) continuous recrystallization in nickel and

Ni-base alloys. (X-axis for (a) and (b) are same scale) From Luton and Sellars[29]...... 32

Figure 4.1: Accustir FSW machine (left) and W-25Re tool profile (right)...... 35

Figure 4.2: Processing windows for Hastelloy X, Alloy 625 and Alloy 718 FSP. Open symbols represent parameters with visual discontinuities, filled symbols represent parameters with no visually detectable discontinuities...... 36

Figure 4.3: Optical macrographs of cross sectioned friction stir processed (i) Hastelloy X

(180RPM, 2IPM), (ii) Alloy 625 (100RPM,2.5IPM) and (iii) Alloy 718 (100RPM,

2RPM) (10X). Note feathered streaking on advancing side...... 39

Figure 4.4: Optical micrograph of (A) Hastelloy X base metal, (B) Hastelloy X TMAZ,

(C) Hastelloy X stir zone, (D) Alloy 625 base metal, (E) Alloy 625 TMAZ, (F) Alloy 625 stir zone, (G) Alloy 718 base metal, (H) Alloy 718, (I) Alloy 718 stir zone (500X)...... 40

Figure 4.5: Representative collected FSP thermal history (Alloy 625: 100RPM, 2.5IPM)

...... 41

Figure 4.6: Instantaneous cooling rates encountered during FSP Ni-base alloys...... 42

Figure 4.7: Cross section of thermocouple locations showing thermocouples were successfully incorporated into the stir zone during FSP. (Alloy 625: 100RPM, 2.5IPM)

...... 42

Figure 4.8: Summary of process parameter and force interaction (Alloy 625 FSP run 15).

...... 44

Figure 4.9: Microhardness traverses for FSP alloys...... 45

Figure 4.10: Tensile properties of alloys. Dark colors represent base metal and light colors represent FSP material. Wide bars represent tensile strengths and thin bars represent percent reduction of area ...... 46

Figure 4.11: Spot-varestraint maximum crack length results...... 48

Figure 4.12: Spot-varestraint total crack length results...... 48

Figure 4.13: Optical micrograph of Alloy 718 GTAW HAZ (left) base metal (right) FSP region (500X, oxalic etch)...... 49

Figure 4.14: Grain size effect on HAZ liquation (A) maximum crack length and (B) total crack length for Alloy 718...... 50

Figure 4.15: Reference temperature curve generated for Alloy 625 in the Gleeble...... 52

Figure 4.16: Expanded view of peak temperature profile encountered during hot torsion testing. Note sharp increase in temperature during the torsion event. Sharp drop in temperature occurs at the end of the torsion event due to activation of the He quench. .. 52

Figure 4.17: Optical micrograph of Alloy 625(A) FSP TMAZ and (B) hot torsion simulated TMAZ (500X, Lucas‟ reagent)...... 53

Figure 4.18: Optical micrograph of Alloy 718 (A) FSP stir zone and (B) hot torsion simulated stir zone (1000X, Oxalic etch)...... 54

Figure 4.19: Typical shear stress-strain curve generated for a specimen that does not undergo any significant microstructural changes...... 55

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Figure 4.20: Shear stress-strain curve generated for a specimen that experiences significant microstructural changes with evidence of dynamic recrystallization...... 55

Figure A.1: Illustration of friction stir processing denoting the tool features (top) and the determination of advancing and retreating side (bottom)...... 65

Figure A.2: Effect of friction stir welding travel speed on peak stir zone temperature and thermal history [9]...... 66

Figure A.3: Study by Song et al showing the effect of friction stir welding travel speed on Alloy 600 (left) grain size and (right) yield strength. Joint denotes specimen transverse to weld direction and stir zone denotes longitudinal. It was noted that joint specimens failed in the base metal [9]...... 67

Figure A.4: (a) Top view of Alloy 625 FSW, (b) banding structure observed in Alloy 625

SZ due to tool wear, (c) top view of Alloy 718 FSW, and (d) banding due to tool wear in

SZ of Alloy 718 [7,8]...... 68

Figure A.5: Microhardness profiles of as welded and PWHT FSW in (a) Alloy 625 and

(b) Alloy 718 [7,8]...... 69

Figure A.6: Tensile testing summary for as welded and PWHT FSW (a) Alloy 625 and

(b) Alloy 718 [7,8]...... 69

Figure A.7: W-25Re tool utilized for FSP...... 71

Figure A.8: Accustir friction stir welding machine...... 72

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Figure A.9: Schematic of thermocouple placement with relationship to tool profile

(dimensions in inches). A 304L backing plate with a through hole pattern and relief channels was utilized to avoid damaging thermocouple leads...... 73

Figure A.10: ASTM E8 sub-size tensile test specimen geometry [11]. Thickness of samples was 0.125 in...... 73

Figure A.11: Friction stir processing window established for Hastelloy X. Open circles denote parameters containing visual discontinuities...... 76

Figure A.12: Friction stir processing window for Alloy 625. Open circles denote parameters containing visual discontinuities. The diamond represents settings used by

Song and Nakata [7]...... 76

Figure A.13: Friction stir processing window for Alloy 718. Open circles denote parameters with visual discontinuities. The diamond represents settings used by Song and Nakata [8]...... 77

Figure A.14: Acquired thermal history for Alloy 625 FSP run 40 (100 RPM, 2.5 IPM). 79

Figure A.15: Thermocouple locations relative to stir zone for Alloy 625 FSP run 40. ... 79

Figure A.16: Instantaneous cooling rates encountered during FSP Ni-base alloys...... 80

Figure A.17: Thermal conductivities as a function of temperature [12-14]...... 81

Figure A.18: FSP run 22 showing how (top) process forces and torque are affected by

(bottom) process inputs of plunge depth and travel speed. (P denotes plunge series, dotted vertical lines denote when forward travel starts and stops, W denotes withdrawal of tool.) ...... 83

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Figure A.19: FSP run 9 showing (top) process forces and torque (bottom) process inputs of plunge depth and travel speed. (P denotes plunge series, dotted vertical lines denote when forward travel starts and stops, W denotes withdrawal of tool from material.) ...... 84

Figure A.20: FSP run 15 showing how (top) process forces and torque are affected by

(bottom) process inputs of plunge depth, tool rotation rate, and travel speed. (P denotes plunge series, dotted vertical lines denote when forward travel starts and stops, W denotes tool withdrawal.) ...... 86

Figure A.21: FSP run 29 showing how (top) process forces and torque are affected by

Figure A.22: FSP run 41 showing (top) process forces and torque (bottom) process inputs of plunge depth and travel speed. Torque data could not be collected. (P denotes plunge series, dotted vertical lines denote when forward travel starts and stops, W denotes tool withdrawal.) ...... 88

Figure A.23: Optical macrographs of cross sectioned friction stir processed (i) Hastelloy

X, (ii) Alloy 625 and (iii) Alloy 718 (10X). Note feathered streaking (dark etching) on the advancing side of the stir zone for all three alloys...... 91

Figure A.24: Optical micrograph of Hastelloy X (A) base metal with 88µm grain size and (B) stir zone with 6µm grain size (500X, oxalic etch)...... 92

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Figure A.25: Optical micrograph of Alloy 625 (A) base metal with 26µm grain size and

(B) stir zone with 5µm grain size (500X, Lucas' reagent). Note the NbC particles are unchanged from the base metal...... 92

Figure A.26: Optical micrograph of Alloy 718 (A) base metal with 44µm grain size and

(B) stir zone with 4µm grain size (500x, oxalic etch). Note the NbC particles are unchanged from the base metal...... 93

Figure A.27: Optical micrograph of FSP stir zone for (A) Hastelloy X, (B) Alloy 625, and (C) Alloy 718 (1000X). Particles in (B) and (C) are the NbC from the base metal. 94

Figure A.28: Optical micrographs of FSP TMAZ for (A) Hastelloy X, (B) Alloy 625, and (C) Alloy 718 (500X)...... 95

Figure A.29: High resolution SEM image of Alloy 625 stir zone (left) showing general appearance of streaks on advancing side (250X) and (right) higher magnification of bright phase (6000X)...... 96

Figure A.30: High resolution SEM image of particles that cause streaking in the advancing side of stir zones (Alloy 625, 20000X). Crosses show where EDS spot scan was taken for comparison of composition...... 97

Figure A.31: EDS count intensities for Alloy 625 (top) matrix and (bottom) particles. . 98

Figure A.32: FSP Hastelloy X hardness profile. (BM=base metal, T=TMAZ, SZ=stir zone.) ...... 101

Figure A.33: FSP Alloy 718 hardness profile. (BM=base metal, T=TMAZ, SZ=stir zone.) ...... 101

Figure A.34: FSP Alloy 625 hardness profile. (BM=base metal, T=TMAZ, SZ=stir zone.) ...... 102

Figure A.35: Microhardness profiles from Song and Nakata for as friction stir welded and post weld heat treated (a) Alloy 625 and (b) Alloy 718 [7,8]. Note the significant increase in hardness for the stir zone relative to the base metal...... 102

Figure A.36: Tensile properties of FSP and base metal for all three alloys. Darker shades are base metal, lighter shades are FSP metal. Wide bars represent ultimate stress properties and thin bars represent percent reduction of area...... 104

Figure B.1: W-25Re tool profile...... 112

Figure B.2: As FSP material (top) showing striations on surface. Surface machined FSP material eliminating striations (bottom)...... 114

Figure B.3: Schematic of spot-varestraint testing apparatus...... 115

Figure B.4: Illustration of gas tungsten arc spot weld placement relative to stir zone. . 115

Figure B.5: Example of crack measurement (20X)...... 117

Figure B.6: Optical micrograph showing grain size difference in (A) Hastelloy X base metal and (B) FSP stir zone (500X, 10% Oxalic), (C) Alloy 625 base metal and (D) FSP stir zone (500X, Lucas‟ Reagent), and (E) Alloy 718 base metal and (F) FSP stir zone

(500X, Oxalic)...... 119

Figure B.7: Comparison of maximum crack length...... 121

Figure B.8: Comparison of total crack length...... 121

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Figure B.9: Illustration of determination of temperature at which liquid begins to form using ThermoCalc® simulations...... 123

Figure B.10: Optical micrograph of Hastelloy X spot-varestraint HAZ in (left) base metal and (right) FSP region (200X). Arrows point to areas of intergranular film formation...... 125

Figure B.11: Optical micrograph of Alloy 625 spot-varestraint HAZ in (left) base metal and (right) FSP region (200X). Arrows point to areas of intergranular film formation. 125

Figure B.12: Optical micrograph of Alloy 718 spot-varestraint HAZ in (left) base metal and (right) FSP region (200X). Arrows point to areas of intergranular film formation. 126

Figure B.13: Hastelloy X SEM images of (A) liquid extension (light phase) from PMZ to crack opening (1500X) and (B) eutectic film formation (6000X). Crosses denote where

EDS spots were taken...... 126

Figure B.14: EDS intensities for Hastelloy X (top) matrix and (bottom) intergranular film regions...... 127

Figure B.15: Alloy 625 SEM images of (A) liquid extension (light phase) from PMZ to crack opening (1500X) and (B) eutectic film formation (6000X). Crosses denote where

EDS spots were taken...... 128

Figure B.16: EDS intensities for Alloy 625 (top) matrix and (bottom) intergranular film regions...... 129

Figure B.17: Alloy 718 SEM images of (A) liquid extension (light phase) from PMZ to crack opening (150X) and (B) eutectic film formation (2500X). Crosses denote where

EDS spots were taken...... 130

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Figure B.18: EDS intensities for Alloy 718 (top) matrix and (bottom) intergranular film regions...... 131

Figure B.19: HAZ liquation susceptibility response to grain size for Alloy 718 (from

Thompson et al) [8]...... 133

Figure B.20: Effect of grain size on HAZ liquation cracking in terms of (a) total crack length and (b) maximum crack length (cast Alloy 718) [10]...... 134

Figure B.21: Alloy 625 HAZ liquation crack response to grain size in terms of (A) maximum crack length and (B) total crack length...... 136

Figure B.22: Alloy 718 HAZ liquation crack response to grain size in terms of (A) maximum crack length and (B) total crack length...... 137

Figure B.23: Hastelloy X HAZ liquation crack response to grain size in terms of (A) maximum crack length and (B) total crack length...... 138

Figure B.24: HAZ liquation crack frequency comparison between alloys and base metal and FSP HAZ...... 139

Figure B.25: Arc weld HAZ grain size response to heat input. Solid lines represent FSP material and dashed lines represent parent material...... 140

Figure C.1: Modified Gleeble hot torsion test specimen geometry [1]...... 146

Figure C.2: Proposed shear stress-strain signatures indicating the occurrence of (a) discontinuous dynamic recrystallization and (b) continuous recrystallization in nickel and

Ni-base alloys. (X-axis for (a) and (b) are same scale) From Luton and Sellars...... 149

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Figure C.3: Thermocouple placement (left) view in direction of travel, (right) view of backside of processed plate embedded with type K thermocouples...... 151

Figure C.4: W-25Re tool profile (top), GTC AccuStir machine (bottom)...... 152

Figure C.5: Thermocouple placement for torsion tests...... 153

Figure C.6: Thermocouple locations for run 23 (Hastelloy X)...... 156

Figure C.7: Thermocouple locations for run 24 (Hastelloy X)...... 157

Figure C.8: Acquired thermal history for run 24...... 157

Figure C.9: Thermocouple locations in run 40 (Alloy 625)...... 158

Figure C.10: Acquired thermal history for run 40...... 159

Figure C.11: Example of common rupture experienced with Hastelloy X hot torsion tests

(scale is mm)...... 160

Figure C.12: Comparison of gauge center temperature and acquired thermal history from

FSP run 24 (Hastelloy X)...... 161

Figure C.13: Stress-strain curve for Hastelloy X hot torsion test (hastX3.d09). Note sample rupture occurrence at ~0.4 in/in strain evident by sharp drop in stress...... 161

Figure C.14: Optical micrograph of Hastelloy X (A) hot torsion test specimen unaffected base metal, (B) hot torsion test gauge center, and actual FSP stir zone (500X, oxalic etch)...... 162

Figure C.15: Comparison of thermal data for „hastx3.d13‟ hot torsion test with FSP run

24 (Hastelloy X) and shoulder temperature from a hot torsion specimen heated and cooled at the same rate without a torsion event...... 163

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Figure C.16: Stress-strain curve for Hastelloy X hot torsion test (hastX3.d13). Note sample shearing imitates at ~0.2 in/in strain evident by steady drop in stress and sample rupture occurrence at ~0.5 in/in strain evident by sharp drop in stress...... 164

Figure C.17: Optical micrograph of Hastelloy X (A) FSP stir zone and (B) hot torsion test stir zone (hastx3.d13) (1000X, oxalic etch)...... 165

Figure C.18: Optical micrograph of Hastelloy X (A) FSP TMAZ and (B) hot torsion test

TMAZ (hastx3.d13) (500X, oxalic etch)...... 165

Figure C.19: Thermal history comparison of Gleeble hot torsion test and actual FSP run for Alloy 625...... 167

Figure C.20: Expanded view of peak temperature encountered by Alloy 625 during hot torsion testing. Note sharp increase in temperature during the torsion event...... 167

Figure C.21: Stress-strain curve for Alloy 625 hot torsion test (625-6.d02)...... 168

Figure C.22: Optical micrograph of Alloy 625 hot torsion test (625-6.d02) (A) shoulder and (B) center of gauge section (500X, Lucas‟ reagent)...... 168

Figure C.23: Thermal history of Alloy 625 hot torsion test 625-6.d03 with comparison to

FSP run 40 and a control torsion specimen...... 170

Figure C.24: Stress-strain curve generated for Alloy 625 hot torsion test 625-6.d03. .. 170

Figure C.25: Optical micrograph of Alloy 625 (A) FSP stir zone microstructure and (B) hot torsion test 625-6.d03 simulated stir zone (1000X, Lucas' reagent)...... 171

Figure C.26: Optical micrograph of Alloy 625 (A) FSP TMAZ and (B) hot torsion test

625-6.d03 simulated TMAZ (500X, Lucas' reagent)...... 172

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Figure C.27: Thermal history of Alloy 718 hot torsion test 718-2.d08 with comparison to

FSP run 40 and a control torsion specimen...... 173

Figure C.28: Stress-strain curve generated for Alloy 718 hot torsion test 718-2.d08. .. 173

Figure C.29: Optical micrograph of Alloy 718 (A) FSP stir zone and (B) hot torsion test

718-2.d08 simulated stir zone (1000X, oxalic etch)...... 174

Figure C.30: Optical micrograph of Alloy 718 (A) FSP TMAZ and (B) hot torsion test

718-2.d08 simulated TMAZ (500X, oxalic etch)...... 174

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CHAPTER 1: INTRODUCTION

Friction stir processing (FSP) is a derivative of friction stir welding (FSW) where the rotating tool is limited to a partial penetration, non-joining action. This solid state process results in similar microstructural changes experienced in FSW due to the intense plastic deformation and simultaneous heating imparted on the material, typically referred to as dynamic recrystallization[1]. This process leads to highly refined, equiaxed grains that have been shown to provide various material property benefits, such as strength, ductility, corrosion resistance, etc[1]. FSP has an advantage over other material processing techniques that lead to similar grain refinement in that it can be very localized and selective.

Friction stir welding and processing of aluminum alloys has been widely studied since the invention of the process in 1991 by TWI [2]. The ability to successfully deploy this process on Ni-base alloys has only been recently realized. This is largely attributed to the development of specialized tool materials and machine size requirements to accommodate the higher process temperatures and forces required for friction stir welding and processing Ni-base alloys. Primarily, there are two groups of FSW/P tools for Ni-base alloys, tungsten base and polycrystalline cubic boron nitride (PCBN). Since the development of these tools, several ferrous alloys have been successfully friction stir welded and/or processed but application to Ni-base alloys has experienced limited research and success. 1 The focus of this study was to investigate the ability of current FSW equipment for FSP of three Ni-base alloys, Hastelloy X, Alloy 625, and Alloy 718. Process forces were recorded to gauge the minimum requirements needed to implement FSP for Ni-base alloys. Additionally, the interaction of process inputs (plunge depth, travel speed and tool rotation rate) and forces was investigated. Characterization of microstructure was performed to demonstrate what can be expected when FSP these materials. Evidence of tool wear within the microstructure was examined to further demonstrate the burden these materials have on current tooling and the machine. Furthermore, thermal histories were collected for each material to show that temperatures encountered when FSP Ni-base alloys are comparable to steels and to show that peak temperatures are still well below the melting point of the material.

An application for FSP Ni-base alloys was developed in the form of a pre-fusion welding treatment to increase resistance to heat-affected zone (HAZ) liquation cracking.

It is well known that Ni-base alloys suffer in terms of weldability due to a propensity for hot cracking during fusion welding. Many efforts have been made to reduce the solidification cracking tendencies of these alloys through arc manipulation in the form of pulsed current and arc stirring, but these techniques do not provide benefits for the heat- affected zone. Previous research has shown that reducing grain size and controlling phase distribution can significantly improve HAZ liquation resistance. Previous research used bulk forming techniques to achieve the small grain sizes. A patent by Ditzel et al[3] suggested that processing a material using friction prior to fusion welding can reduce hot cracking tendencies of Ni-base alloys and reduce residual stresses. FSP can be efficiently utilized in this regard as it is a localized process that can be selectively applied to areas of 2 the material that are to be fusion welded later in fabrication. The selective processing of small areas as opposed to bulk material has the potential to reduce fabrication costs and time and provide microstructure improvement only in locations where it is needed.

Quite recently, a method for physically simulating FSW microstructure was developed at Ohio State University by Dr. Seth Norton in the form of the modified

Gleeble hot torsion test[4]. Dr. Norton demonstrated that FSW microstructure of pure iron and HSLA-65 steel can be reproduced using this Gleeble based test. Furthermore, shear stress and shear strain can be collected to aid numerical model development of

FSW. Matthew Sinfeld has made improvements on this method by limiting the material to one revolution during the torsion event and using thermocouples to provide better control of heating and cooling of the material[5]. Sinfeld has successfully reproduced

FSW microstructures of HSLA-65 using these refinements. David Failla has made use of this test for simulating HSLA-65 and type 310 and 304L stainless steel FSW [6]. Much like Norton and Sinfeld, Failla was successful in simulating the HSLA-65 FSW microstructures but encountered issues with the stainless steels. His testing required the use of multiple test specimens at various settings to reproduce the three distinct regions of a FSW microstructure. This investigation used the modified Gleeble hot torsion test to physically simulate FSP microstructure of Ni-base alloys and suggest further refinements to more successfully apply the test method to a broader range of materials.

CHAPTER 2: BACKGROUND

2.1 NICKEL ALLOYS

A nickel base (Ni-base) alloy is one that has nickel as it primary constituent.

These alloys have found usage in a variety of industries, such as aerospace, petroleum, chemical, and power generation industries[7,8]. The alloys have been used in a variety of demanding environments from cryogenic to very high temperatures, fresh water and salt water to acidic solution containing vessels, steam and pressure vessels to electronic and biomedical components. It is evident that a wide variety of Ni-base alloys with different properties can be engineered for nearly any application. There is no clear classification system for Ni-base alloys based on number, evident from such names as Inconel 600,

Monel 400, Cabot 214, Rene 80, or Hastelloy C-22, however they are organized in the unified number system (UNS) under the N00001-N99999 series[7]. Often, the trade name alloys will be referred to simply as Alloy 600 (Inconel 600) or Alloy 214 (Cabot

214) when patents for the alloy expire [7].

2.1.1 Ni-base alloy microstructure

The primary crystal structure of all Ni-base alloys is face centered cubic (fcc), schematically shown in Figure 2.1, with atoms sharing sites at the corners of the unit cell as well as the face surfaces[7,9]. Three primary groups of Ni-base alloys have been established classified by their strengthening mechanisms; solid-solution strengthening,

4 precipitation strengthening/hardening, and oxide dispersion strengthening [9]. Generally, solid solution strengthening is done by atomic replacement (substitution) of one or more

Ni atoms in the lattice to induce some strain or deformation to the lattice thereby hindering motions of dislocations [9]. Precipitation strengthening is done by nucleation of an ordered second phase, typically in the form of γ‟ (Ni3(Al,Ti)) or γ” (Ni3Nb) which makes motion of dislocations in the material more difficult [7]. Oxide dispersion strengthening is quite similar to precipitation strengthening except that the second phase is an oxide, such as Alumina (Al2O3) or Yttria (Y2O3) and is introduced mechanically[7].

Figure 2.1: The FCC unit cell (a) hard sphere representation, (b) reduced sphere representation. Adapted from Callister, Material Science and Engineering - An

Introduction, 7th Edition, 2007[9]

Three alloys were used throughout this study including two solid-solution strengthened, Inconel (Alloy) 625 and Hastelloy (Alloy) X, and one precipitation strengthened, Inconel (Alloy) 718, nominal compositions listed in Table 2.1. Alloy 625

5 is an alloy strengthened by a combination of atomic replacement of lattice nickel atoms by others such as chromium, iron, manganese, and molybdenum, along with strong carbides in the form of niobium carbide (NbC) and M6C type where M is Cr, Mo. Alloy

X is strengthened in a similar manner to Alloy 625 except that there are no niobium carbides only M6C and M23C6 type carbides. Alloy 718 is primarily strengthened by precipitation of an ordered second phase, Ni3Nb (γ”), and can also contain NbC and M6C type carbides.

Table 2.1: Nominal compositions of three Ni-base alloys used in this study [7].

2.1.2 Ni-base alloy weldability

Ni-base alloys are generally considered weldable but many of the alloys require a high level of care and attention to accomplish a successful weld [7,8,10]. The weldability of these alloys has been studied for over half a century due to their widespread use in a variety of industries and environments [7,8,10]. In terms of fusion welding, most Ni-base alloys suffer from some form of hot cracking in the form of solidification cracking and/or heat-affected zone (HAZ) liquation cracking. Some alloys may suffer from warm cracking phenomena such as ductility-dip cracking and strain-age-cracking. Weldability studies continue today with the development of new Ni-base alloys and to further the fundamental understanding of contributing factors that lead to welding discontinuities.

6 Part of this study specifically deals with hot cracking in the form of HAZ liquation cracking, a form of hot cracking that affects many Ni-base alloys and, specifically, all three alloys used throughout this study[7,8,11,12,13,14]. To understand the mechanisms of HAZ liquation cracking, the basic metallurgical features of a fusion weld must be discussed. A typical fusion weld contains three microstructurally distinct regions, the fusion zone, the HAZ, and unaffected base metal, Figure 2.2[15]. The fusion zone is the region in a fusion weld where complete mixing of the two materials to be joined occurs due to melting. The HAZ is the region adjacent to the fusion zone where no mixing of materials occurs but partial melting of the base metal can occur.

Additionally, grain size and carbide morphology change in this region due to thermal excursion. The unaffected base metal is the remaining region about the HAZ that has fully preserved the original base metal microstructure. HAZ liquation cracking is limited to the HAZ region of a fusion weld and is associated with the partially melted zone

(PMZ) of this region directly adjacent to the fusion zone.

Figure 2.2: Illustration of fusion weld and related regions: (A) fusion zone, (B) partially melted zone, (C) heat-affected zone, (D) base metal, (E) original joint edge of materials.

8 Mechanistically, HAZ liquation can occur in at least two manners, segregation and penetration mechanisms, illustrated in Figure 2.3[7,12,13,16,17]. The segregation mechanism occurs due to previous segregation of phases to grain boundaries that have a low melting temperature relative to the matrix[7]. During fusion welding, these phases in the HAZ will liquate at a sub-solidus temperature experienced in this region. The liquated grain boundaries experience some level of stress due to the shrinkage of the weld pool during solidification in the adjacent fusion zone. The liquated boundary no longer has the strength to accommodate the stress and leads to crack formation. The penetration mechanism is a bit more complex than the segregation mechanism and requires the simultaneous action of the constitutional liquation mechanism. Constitutional liquation involves the presence of a second phase particle that, under thermal excursion in the

HAZ, reacts with the matrix to form some low melting phase surrounding the particle[12,16]. Simultaneously, motion of grain boundaries drag across these liquated particles and the liquid penetrates the boundary [16]. During solidification of the weld pool or cooling of the weldment, shrinkage stresses are imposed upon these boundaries.

The liquated boundaries cannot accommodate the stress and leads to the formation of cracks.

Figure 2.3: Mechanisms of HAZ liquation cracking (left) the segregation mechanism,

(right) the penetration and constitutional liquation mechanism. Courtesy J.C. Lippold

Fundamentally, it has been found that HAZ liquation cracking is influenced by three primary factors, material composition/condition, welding parameters, and restraint.

The majority of studies have been conducted on material composition and condition which include alloy element, phase, and grain size effects on HAZ liquation cracking.

Early research by Savage and Krantz dealt with Hastelloy X solute and impurity element

(S, Si, Mn) content in regards to hot cracking. It was found that elevated levels of sulfur increased the tendencies of Hastelloy X to HAZ liquation cracking. Additionally, it was demonstrated that to some extent, Mn and Si could help reduce cracking tendency by tying up S and deoxidizing the metal [12]. Furthermore, Savage and Krantz reconfirmed the concepts of constitutional liquation and its effects on HAZ liquation cracking through studying carbide morphology. It was determined that MC type carbides were susceptible to constitutional liquation and responsible for the cracks found in the HAZ whereas zirconium carbo-nitrides exhibited no tendency to constitutional liquation and subsequent 10 cracking [12]. Other studies by Duvall and Oczwarki have demonstrated that HAZ liquation cracking is attributed to secondary and eutectic phases that often are attributed to intentional alloy additions such as chromium, molybdenum, manganese, niobium, aluminum, titanium and carbon[16,17]. A study by R.G. Thompson et al shows that material condition in terms of initial grain size has significant influence on the extent of

HAZ liquation cracking. It was demonstrated that finer grain size improves the resistance of Alloy 718 to HAZ liquation cracking and that susceptibility, in terms of total crack length, increases linearly with grain size, Figure 2.4[14]. This relationship was further supported by I. Woo et al in a study of cast Alloy 718 grain size effect on HAZ liquation cracking, summarized by Figure 2.5, [18].

Figure 2.4: Effect of HAZ grain size on HAZ liquation cracking (Alloy 718) [14].

Figure 2.5: Effect of grain size on HAZ liquation cracking in terms of (a) total crack length and (b) maximum crack length (cast Alloy 718) [18].

12 2.2 FRICTION STIR PROCESSING

Friction stir processing (FSP) is a materials processing technique based on the basic principles of friction stir welding (FSW), where a non-consumable rotating tool containing a shoulder and pin is plunged into the surface of a work piece and translated along a planar direction perpendicular to the plunge, Figure 2.6[2,19,20]. The contact friction of the tool and the work piece along with the intense plastic deformation of material generates heating [1,20]. This localized heating softens the material around the tool allowing material movement as the tool rotates and translates along the travel direction. From the illustration in Figure 2.6 it is evident that some asymmetrical behavior can occur due to tangent velocity vectors of the rotating tool being in the same direction on one side of the processing region and opposing on the other. Thusly, terminology has been defined such that if the tangent velocity vector of the rotating tool and the travel direction are the same it is called the advancing side whereas if they are opposite it is termed the retreating side, Figure 2.6. FSP is a solid state process where virtually no melting of material occurs. A homogenous fine grained microstructure results due to the combined thermomechanical events that take place during FSP [1,20].

It has been demonstrated that refinement of microstructure has resulted in improved material performance such as strength, formability, corrosion, etc.[1,19,21]. The advantage of FSP over various other processes, such as accumulated roll bonding (ARB) and thermomechanical treatment (TMT), is that FSP is a localized, single step process whereas ARB and TMT affect the material in bulk and requires multiple steps to achieve the refined microstructure[19].

Figure 2.6: Illustration of friction stir processing denoting the tool features (top) and the determination of advancing and retreating side (bottom).

Like FSW, FSP microstructure contains four distinct regions, the stir zone (SZ), thermomechanically affect zone (TMAZ), heat-affected zone (HAZ), and unaffected base metal/material (BM), Figure 2.7[1,19,20]. The SZ is the region of processing that undergoes the most intense plastic deformation and experiences the highest temperatures.

The high density of dislocations resulting from plastic deformation with the presence of elevated temperature (~80% melting temperature (Tm)) leads to a simultaneous recovery and recrystallization phenomenon termed dynamic recrystallization (DRX) [20]. The

DRX mechanism typically yields a microstructure consisting of much finer grains than the BM. Due to asymmetrical behavior, microstructures away from the center of the SZ near the SZ/TMAZ boundary can differ between the advancing and retreating side. The

TMAZ is the region adjacent to the SZ on either side of the processed region containing evidence of original BM microstructure but in a very deformed state, typically elongated 14 grains in the direction of material flow. The HAZ lies outside of the TMAZ where no mechanical affects are detected but grain coarsening can occur due to thermal excursion.

The BM is fully preserved, unaffected parent material.

Figure 2.7: Schematic of typical FSP showing four microstructurally distinct regions:

(A) stir zone, (B) thermomechanically affected zone, (C) heat-affected zone and (D) unaffected base material.

Primary independent process variables for FSP consist of tool rotation speed, travel speed, and plunge depth. These primary process variables will affect heat generation rate, peak temperature, torque, and processes forces such as down force and path force. It has been established that faster tool rotation rates and slower travel speeds lead to higher peak temperatures and lower torques and forces whereas faster travel speed and slower tool rotation rate leads to lower peak temperatures and higher process torque and forces[20]. Tool design can affect the process with various pin and shoulder designs that can include threads, flats, and flutes along with various sizes of tools. Many studies have been conducted to show how tool design affects material flow, heat generation and power requirements but in summary, it is understood that larger tool volume plunged into

15 the material results in increased temperatures and power requirements to move that tool through the material[20].

2.2.1 Previous studies of friction stir welding/processing Ni-base alloys

Research of FSW/P Ni-base alloys is fairly limited and likely due to only recently having tools developed that can withstand process temperatures and forces required. The earliest study by Ye et al in 2006 investigated the microstructure and physical properties of 2 mm (0.079 in.) thick Alloy 600 friction stir welds[22]. Welding was performed at

400 rotations per minute (RPM) and a travel speed of 1.67 millimeters per second (mm/s)

(~3.95 inches per minute (IPM)) and 3º push angle. Temperature history of the backside of the center of the weld was recorded with a peak temperature of 800 ºC (~1472 ºF).

Microstructure of the stir zone was not fully resolved but shows some evidence of grain refinement in comparison to the original base metal grain size. Microhardness of the joint revealed that no significant softening or hardening occurred between the base metal,

HAZ, TMAZ, and SZ, Figure 2.8. Tensile samples were extracted transverse and parallel to the welding direction to investigate the strengths relative to base metal, results summarized in Figure 2.9. It was shown that the transverse tensile test had nearly the same strength as the base metal as failure occurred in the base metal away from the FSW joint. Additionally, the all FSW tensile specimen showed a slight strength increase over both the base metal and transverse specimens. It was noted that a decrease in elongation was experienced in the welded samples relative to base metal samples.

Figure 2.8: Microhardness profile of Alloy 600 FSW reported by Ye et al [22].

Figure 2.9: Tensile properties of Alloy 600 FSW reported by Ye et al [22].

17 An expanded study by Sato et al investigated the effect of microstructure on properties of FSW Alloy 600, including microstructure characteristics, mechanical properties and corrosion properties[23]. This study used 4.8 mm (~0.188 in.) thick Alloy

600 welded using a PCBN tool at 600 RPM and travel speed of 1 mm/s (~2.36 IPM) with a 3.5º push angle. Optical microscopy revealed a typical FSW containing a SZ, TMAZ,

HAZ, and BM however the SZ was subdivided into two regions being a „black zone

(BZ)‟ and a „grey zone (GZ)‟ due to etching behavior. It has found that the GZ and possibly the BZ had fine equiaxed grains (average diameter 22µm and 15µm respectively) as opposed to the relatively larger grains in the BM (average diameter

61µm). The difference between the GZ and BZ in the SZ is accounted for due to the high density of CBN particles in the BZ which pin grain boundaries, suppressing grain growth.

Transverse tensile tests of welded samples trend well with the previous study by Ye et al, showing that strength is slightly higher while a decrease in elongation occurs for the FSW sample, Figure 2.10. Microhardness profiles of welded samples also trend well with those in Ye‟s et al study, however it is noted that significant hardening occurs in the region of the BZ and is likely due to the very high hardness of CBN particles, Figure

2.11. Corrosion experiments were performed using a ferric sulfate-sulfuric acid test. It was demonstrated that the area of most severe corrosion was in a portion of the HAZ termed the area of greatest loss of grains (GLG). Within the SZ it was shown that the BZ was more susceptible to corrosion and is attributed to the presence of many CBN particles. The GZ experienced the least amount of corrosion. A summary of the corrosion testing is provided in Figure 2.12. A higher current ratio correlates to increased susceptibility. 18

Figure 2.10: Transverse tensile test results for Alloy 600 FSW [23].

Figure 2.11: Microhardness profile of Alloy 600 FSW [23].

Figure 2.12: Summary of ferric sulfate-sulfuric acid corrosion test showing reactivation current ratios. Higher current ration denotes increased corrosion susceptibility [23].

Further studies on Alloy 600 by Song, Fujii and Nakata investigated the effect of welding speed on microstructure and mechanical properties [24]. Welding was performed on 2 mm thick Alloy 600 plates at 400 RPM and travel speeds of 150, 200 and

250 mm/min (~5.9, 7.9, and 9.8 IPM) utilizing a tungsten carbide cobalt tool at a 3º push angle. It was demonstrated that as welding travel speed increased, SZ grain size decreased, shown in Figure 2.13. Temperature histories from the backside of the center of the welds showed that as welding speed is increased heat input and peak temperatures are decreased, Figure 2.14. Additionally, it was shown that microhardness increases in the stir zone as welding speed increases, Figure 2.15, attributed to the finer grain size and the Hall-Petch relationship between grain size and strength. To further the Hall-Petch relationship, tensile specimens were extracted transverse and parallel to the weld direction. It was shown that as welding speed increased (and grain size decreased) joint

20 and stir zone strength increased while elongation decreased, trending well with previous studies, Figure 2.16.

Figure 2.13: Effect of welding speed on stir zone grain size for Alloy 600[24].

Figure 2.14: Effect of welding speed on stir zone thermal history for Alloy 600[24].

Figure 2.15: (left) Effect of welding speed on stir zone hardness for Alloy 600. (right)

Hall-Petch relationship between grain size and hardness for Alloy 600 [24].

Figure 2.16: Effect of welding speed on Alloy 600 tensile properties [24].

Two further studies by Song and Nakata investigated the effect of friction stir welding and post weld heat treatment (PWHT) of Alloys 625 and 718[25,26]. FSW was performed using 2mm thick materials with a tungsten carbide cobalt tool at 200 RPM and a 3º push angle. Travel speeds used were 100 mm/min (~3.9 IPM) and 150 mm/min

(~5.9 IPM) for Alloy 625 and 718, respectively. Heat treatment of select welds was done in a vacuum atmosphere at 700ºC for 100 hours for Alloy 625 and 720ºC for 8 hours for

Alloy 718. Thermal histories were recorded for each of the alloys from the backside at the center line of the welds, Figure 2.17. Examination of microstructure showed that similar levels of grain refinement were realized by both materials in the stir zone. Base metal grain size for both alloys were between 5-20µm whereas the stir zones had grain sizes between 1-3µm. Further inspection of the stir zone microstructure revealed banding near the advancing side of the welds, Figure 2.18. It was determined through energy dispersive x-ray spectroscopy (EDS) that the banding region resulted due to tool wear.

23 Microhardness profiles for both as welded and PWHT conditions revealed that hardness in the stir zone is significantly higher than the base metal, Figure 2.19. It was observed the PWHT samples experienced a significant increase in hardness even though grain coarsening had taken place. This was due to precipitation of M23C6 carbides and γ”.

Additionally, it is observed that the stir zone remains the hardest region of the welds after

PWHT. Tensile specimens were extracted transverse and parallel to the welding direction. Testing showed that the both the transverse and parallel FSW specimens were stronger than the base material whereas elongation was reduced. Furthermore, it was observed that the heat treated specimens had even higher strengths and less elongation than the both the base metal and the as welded specimens, Figure 2.20.

Figure 2.17: Thermal history for (a) Alloy 625 FSW and (b) Alloy 718 FSW [25,26].

Figure 2.18: (a) Top view of Alloy 625 FSW, (b) banding structure observed in Alloy

625 SZ due to tool wear, (c) top view of Alloy 718 FSW, and (d) banding due to tool wear in SZ of Alloy 718 [25,26].

Figure 2.19: Microhardness profiles of as welded and PWHT FSW in (a) Alloy 625 and

(b) Alloy 718 [25,26].

Figure 2.20: Tensile testing summary for as welded and PWHT FSW (a) Alloy 625 and

(b) Alloy 718 [25,26].

An application study of FSP Ni-base superalloy Inconel 738 by Mousavizade et al investigated the effect of FSP on hot cracking as a pretreatment to laser welding[27].

FSP was performed on 5 mm (~0.197 in.) thick plates using a tungsten carbide tool rotated at 800 RPM with a travel speed of 50 mm/min (~1.97 IPM) at a 3º push angle with a depth of 1.4 mm (~0.055 in.). After FSP, laser welding was performed transverse

26 to the stir region using a pulsed Nd:YAG laser with a spot diameter of 1mm, pulse frequency of 40 Hz, pulse duration of 10 ms, and travel speed of 3 mm/s. Final specimens were sectioned along the laser welding direction transverse to the FSP region.

Examination of the HAZ of the laser weld in the as cast base metal and the FSP region revealed different response to liquation cracking. It was observed that the as cast material experienced HAZ liquation cracks varying from 20-110µm whereas the FSP material experienced no cracking, Figure 2.21. This difference in cracking response is attributed to the cast microstructure being cored dendrite containing several coarse and discontinuous secondary solidification phases, such as carbides and γ-γ‟ eutectics whereas the FSP microstructure contained very fine, equiaxed grains on the order of 5-

10µm in size with a finer carbide size and distribution.

27 2.3 PHYSICAL SIMULATION

Physical simulation of friction stir welding and processing is an aid to understanding the thermomechanical events experienced by the material. Recently, a

Gleeble based method to physically simulate FSW/P has been developed by Norton and further refined by Sinfield in the form of a modified hot torsion test[4,28]. This method was conceived by Norton by modifying a standard Gleeble hot torsion test sample to include a gauge section and a through hole in the center of the sample, Figure A.1 in appendix A, to achieve the cooling rates experienced during FSW. This modified hot torsion test was successfully implemented for simulating FSW microstructure in ingot iron and HSLA-65 steel[4].

Advancements to the Gleeble based test were made by Sinfield. Sinfield refined the test procedure by limiting the number of rotations to one as opposed to the multiple

(up to 12) rotations used in Norton‟s experiments[28]. Additionally, because Sinfield limited the material to a single rotation thermocouples could be used to monitor the simulation temperatures instead of an optical pyrometer. This allowed for better control of the temperature upon finding the optical pyrometer inaccurate below a particular temperature range. Sinfield also recognized that the strain and temperature was not uniform through the wall thickness of the sample and created a numerical model to visualize and more accurately calculate the shear strain and temperature gradient throughout the wall thickness[5,28]. From these refinements, Sinfield was able to successfully simulate all microstructurally distinct regions of a FSW in HSLA-65 steel.

From the investigation by Norton and Sinfield it has been established that the mechanical history/properties can be calculated, including the shear stress, τ, the shear 28 strain, γ, and the shear strain rate, . The shear stress was calculated using Equation 2.1, where T is the torque, r is the radius of the gauge section, and J is the polar moment of inertia of the gauge section. The polar moment, J, was calculated using Equation 2.2 where ro and ri are the outer and inner radius of the gauge section, respectively. The shear strain was calculated using Equation 2.3, where r is the radius of the gauge section,

Ѳ is the angular rotation in radians from a reference line parallel to the axis of the specimen, and l is the length of the gauge section. The shear strain rate was calculated using Equation 2.4 where RPM is the programmed rotations per minute. This equation is multiplied by 2π and divided by 60 to translate to radians per second.

Equation 2.1: Shear Stress[6].

Equation 2.2: Polar moment of inertia for gauge section[6].

Equation 2.3: Shear Strain[4].

Equation 2.4: Shear Strain Rate[28].

29 Recent use of this technique by Failla successfully simulated FSW microstructure of HSLA-65 steel and Types 310 and 304L stainless steel[6]. In order to successfully simulate each region of FSW, Fallia used various hot torsion samples to recreate each individual region, Figure 2.22 through Figure 2.24. It was noted that various simulation parameters needed to be used to recreate each region within a single FSW whereas

Norton and Sinfield were able to recreate all regions within a single simulation [4-6].

Figure 2.22: Type 310 SZ (a) actual FSW and (b) modified hot torsion test at 1000ºC and

1150 RPM[6].

Figure 2.23: Type 310 TMAZ (a) actual FSW and (b) modified hot torsion test at 1100ºC and 550 RPM[6].

Figure 2.24: Type 310 HAZ (a) actual FSW and (b) modified hot torsion test at 1100ºC and 850 RPM [6].

31 dynamic recrystallization occurs, Figure 2.25. Two signatures are shown, the first of which indicates discontinuous dynamic recrystallization by the periodic nature of the stress levels during steady state deformation. The second curve indicates continuous dynamic recrystallization due to the low variance of stress during steady state deformation. Furthermore, both signatures include a sharp rise to a peak stress at a critical strain before stress drops to the steady state regime during dynamic recrystallization. It is proposed that this peak stress at the critical strain level occurs due to the work hardening behavior of nickel and its alloys.

Figure 2.25: Proposed shear stress-strain signatures indicating the occurrence of (a) discontinuous dynamic recrystallization and (b) continuous recrystallization in nickel and

Ni-base alloys. (X-axis for (a) and (b) are same scale) From Luton and Sellars [29].

CHAPTER 3: OBJECTIVES

1. Evaluate application of FSP Ni-base alloys a. Determine processing windows for Hastelloy X, Alloy 625, and Alloy 718 b. Investigate the effects of travel speed, tool rotation rate, and plunge depth on process down force, path force, and torque c. Characterize FSP microstructure of each alloy. d. Evaluate effect of FSP on mechanical properties of each alloy. 2. Evaluate FSP as a material preparation technique prior to fusion welding for reducing Ni-base alloy susceptibility to HAZ liquation cracking. a. Use the spot-varestraint test as a method for simultaneous evaluation of FSP and parent material. b. Evaluate HAZ grain size affect on HAZ liquation cracking. c. Characterize phases contributing to HAZ liquation cracking. 3. Evaluate the modified Gleeble hot torsion test for use in simulating nickel alloy FSP microstructure. a. Compare modified Gleeble hot torsion test microstructure with actual FSP microstructure. b. Estimate shear stress experienced during FSP Ni-base alloys.

CHAPTER 4: RESULTS AND DISCUSSION

The experimental procedures, results and discussion can be found in the appendices of this thesis which includes papers written for submission to peer reviewed journals and proceedings. The first paper, Appendix A, is entitled “Microstructure and

Mechanical Properties of Three Friction Stir Processed Nickel-Base Alloys” describes process development of friction stir processing Ni-base alloys and characterization of resulting microstructure and mechanical properties. The second paper, Appendix B, is entitled “Application of Friction Stir Processing for Improving Heat-Affected Zone

Liquation Cracking Resistance of Nickel-Base Alloys” demonstrates and application for friction stir processing Ni-base alloys. The third paper, “Physical Simulation of Nickel-

Base Alloy Friction Stir Processing Microstructure” evaluates use of the modified

Gleeble hot torsion test for simulating Ni-base alloy FSP microstructure.

4.1 FRICTION STIR PROCESSING NICKEL-BASE ALLOYS

4.1.1 Processing parameters and windows

All FSP included in this research was performed on an Accustir (General Tool

Co., Cincinnati, OH) gantry style friction stir welding machine, utilizing a tungsten based tool alloyed with 25 weight percent (wt.%) rhenium (W-25Re), Figure 4.1. All processing was performed at a 3º push angle at various travel speeds, tool rotation rates, and plunge depths, Table 4.1. A dimensionless heat input was calculated for each run using Equation 4.1. The plunge depth is defined as the depth of the shoulder contact below the surface of the work piece. Process windows were established based on visual 34 inspection of the FSP runs, Figure 4.2. Acceptable runs contained no voids visible on the surface of the processed region or through the depth of the tool exit hole.

Figure 4.1: Accustir FSW machine (left) and W-25Re tool profile (right).

Equation 4.1: Heat input for friction stir welding and processing. (f is process efficiency,

σT is yield strength at temperature T, A is the area of the tool shoulder, ω is the tool

T rotation rate, Cp is the specific heat at temperature T, and U is the travel speed.)[20]

36 Table 4.1: List of selected FSP runs (non-inclusive).

37 4.1.2 Microstructure

Once processing windows were established for each material transverse cross sectioning of FSP runs were taken to investigate the resulting microstructures, Figure 4.3.

The macrographs reveal that FSP of each material results in three microstructurally distinct regions, the stir zone, TMAZ, and un-affected base metal. A HAZ was not observed in any of the materials. Additionally, it was observed that a feathered streaking occurred regularly on the advancing side of the stir zone for all three materials. A summary of resultant microstructures is shown in Figure 4.4. Grain size measurements were made for the stir zone and base metal using the Abrams three circle method[30] to quantify the level of grain refinement experienced by each alloy, Table 4.2.

Figure 4.3: Optical macrographs of cross sectioned friction stir processed (i) Hastelloy X

(180RPM, 2IPM), (ii) Alloy 625 (100RPM,2.5IPM) and (iii) Alloy 718 (100RPM,

2RPM) (10X). Note feathered streaking on advancing side.

Figure 4.4: Optical micrograph of (A) Hastelloy X base metal, (B) Hastelloy X TMAZ,

(C) Hastelloy X stir zone, (D) Alloy 625 base metal, (E) Alloy 625 TMAZ, (F) Alloy 625 stir zone, (G) Alloy 718 base metal, (H) Alloy 718, (I) Alloy 718 stir zone (500X).

Table 4.2: Summary of grain refinement due to FSP

40 4.1.3 Thermal History

Type K thermocouples were placed in a pattern such that the FSP tool would shear and then re-attach them into the stir zone. Resultant collected data was plotted to show complete thermal history for each thermocouple, revealing the heating and cooling rates along with peak temperature, represented by Figure 4.5. To compare cooling rates of the alloys a plot of the instantaneous cooling rate versus temperature was generated,

Figure 4.6. It was observed that all three alloys experienced nearly the same cooling rates. Cross sectioning thermocouple locations after FSP revealed that thermocouples were incorporated into the stir zone, Figure 4.7.

Figure 4.5: Representative collected FSP thermal history (Alloy 625: 100RPM, 2.5IPM)

Figure 4.6: Instantaneous cooling rates encountered during FSP Ni-base alloys.

Figure 4.7: Cross section of thermocouple locations showing thermocouples were successfully incorporated into the stir zone during FSP. (Alloy 625: 100RPM, 2.5IPM)

42 4.1.4 Process Forces

Process force feedback from the Accustir machine was collected for various FSP runs to investigate the effect of process settings on normal force, path force and torque.

A summary of these interactions is demonstrated by FSP run 15 on Alloy 625, Figure 4.8.

The plunge depth, travel speed, and tool rotation rate were changed individually during the run. It was observed that increasing the plunge depth resulted in proportional increases in both the normal and path forces, attributed to increasing the volume of material affected by the tool. Additionally, it was observed that decreasing travel speed resulted in proportional decreases in the normal and path forces, attributed to the increase in dimensionless heat input (Equation 4.1) allowing further softening of material in front of and in contact with the tool. Finally, it was observed that increase in tool rotation rate led to a decrease in torque, also attributed to the increase in dimensionless heat input, softening material in front of and in contact with the tool.

Figure 4.8: Summary of process parameter and force interaction (Alloy 625 FSP run 15).

4.1.5 Mechanical Properties

Microhardness traverses were performed on cross sections of each alloy, Figure

4.9. It was observed that peak hardness for all three alloys is experienced in the TMAZ due to residual plastic strain. Both Hastelloy X and Alloy 718 experienced moderate hardening in the stir zone relative to the base metal whereas Alloy 625 experienced

44 moderate softening. Softening in Alloy 625 may be attributed to residual cold work in the base metal when making this material.

Figure 4.9: Microhardness traverses for FSP alloys.

ASTM E8 sub-size tensile test specimens were produced from base materials and

FSP material. Testing revealed that both Alloys 625 and 718 experience marginal strength increases for FSP material relative to the base metal whereas FSP Hastelloy X experiences a 48% increase in strength relative to the base metal, Figure 4.10. It is likely that Alloys 625 and 718 did not experience significant strength increases from grain refinement due to the materials being supplied in either a partially aged condition or containing residual cold work due to incomplete annealing.

4.2 FRICTION STIR PROCESSING NICKEL-BASE ALLOYS FOR

IMPROVING HEAT-AFFECTED ZONE LIQUATION CRACKING RESISTANCE

4.2.1 Spot-varestraint Testing

Friction stir coupons of Hastelloy X, Alloy 625, and Alloy 718 were prepared using parameters found in Table 4.3. Upon completion, the surface of the coupons were machined such that striations resulting from FSP were removed, leaving a uniform surface for placement of a gas tungsten arc spot weld. The spot weld was placed such that it straddled the boundary between FSP and unaffected base metal as to provide a means of testing both material conditions simultaneously. Testing parameters are

46 displayed in Table 4.4. Results were evaluated in terms of maximum crack length

(MCL), Figure 4.11, and total crack length (TCL), Figure 4.12. It was found that Alloy

718 improved the most from the FSP pretreatment.

Table 4.3: FSP Parameters

Table 4.4: Spot-varestraint test parameters.

Figure 4.11: Spot-varestraint maximum crack length results.

Figure 4.12: Spot-varestraint total crack length results.

48 4.2.2 Microstructure and Grain Size Effect on HAZ Liquation Cracking

Gas tungsten arc spot weld HAZ microstructures were investigated to measure the grain size and characterize constituents responsible for sub-solidus cracking. An example of the difference in HAZ grain size between the FSP and base metal region is shown in

Figure 4.13. The MCL and TCL data were plotted as a function of grain size, example shown in Figure 4.14. It was observed that HAZ liquation susceptibility increases linearly with grain size. SEM and EDS analysis of liquid films present in the HAZ revealed that phases responsible for sub-solidus cracking include M23C6 and M6C carbides rich in Cr and Mo for Hastelloy X, NbC and M6C carbides rich in Mo and Cr for

Alloy 625 and Alloy 718, Table 4.5.

Figure 4.13: Optical micrograph of Alloy 718 GTAW HAZ (left) base metal (right) FSP region (500X, oxalic etch).

Figure 4.14: Grain size effect on HAZ liquation (A) maximum crack length and (B) total crack length for Alloy 718.

Table 4.5: EDS quantification of intergranular and matrix areas.

50 4.3 PHYSICAL SIMULATION OF FRICTION STIR PROCESS

MICROSTRUCTURE

The modified Gleeble hot torsion test was utilized to simulate FSP microstructure of Hastelloy X, Alloy 625, and Alloy 718. This test was developed at Ohio State

University for simulating friction stir welding microstructure and has experienced great success for Pure Iron and HSLA-65 steels. Limited success was realized for types 310 and 304L stainless steel.

4.3.1 Thermal Control

Previously recorded thermal histories for each alloy were used to program temperature profiles into the Gleeble 3800 thermomechanical simulator. Modified hot torsion bars were machined according to original specifications made by Dr. Seth Norton.

Torsion bars were prepared by mounting thermocouples at the center and shoulder of the gauged section. The center thermocouple was used as the control for temperature. Initial experiments were performed for each alloy without a torsion event to create a reference for the thermal gradient between the center and shoulder of the gauge, Figure 4.15.

These curves were used as references for experiments that had failed center thermocouples during torsion testing. It was found during testing that the torsion event induced an increase in temperature of about 20ºC due to adiabatic heating, Figure 4.16.

Programmed temperatures were adjusted such that the resultant increase during torsion did not exceed the peak temperatures experienced during FSP.

Figure 4.15: Reference temperature curve generated for Alloy 625 in the Gleeble.

52 4.3.2 Microstructure

Difficulty was experienced in recreating stir zone microstructures for all three Ni- base alloys. Use of refinements made to the modified Gleeble hot torsion test by Sinfeld to limit the specimen to one revolution proved insufficient for recreating any microstructure associated with FSP of Ni-base alloys. Testing was refined to include multiple revolutions in an attempt to increase the total amount of strain in the specimens.

In so doing, TMAZ microstructures for each alloy were successfully reproduced with similar grain size and structure, Figure 4.17. Stir zone microstructure size was still not realized but morphology was quite similar, Figure 4.18.

Figure 4.17: Optical micrograph of Alloy 625(A) FSP TMAZ and (B) hot torsion simulated TMAZ (500X, Lucas‟ reagent).

Figure 4.18: Optical micrograph of Alloy 718 (A) FSP stir zone and (B) hot torsion simulated stir zone (1000X, Oxalic etch).

4.3.3 Shear Stress-Strain Behavior

A major benefit in using the Gleeble based test is the ability to collect mechanical data for estimation of shear flow stress and strain encountered by the material during

FSP. Investigation of the shear stress-strain curves generated for modified Gleeble hot torsion testing of the three Ni-base alloys provided estimations of the flow stress and insight into recognizing the occurrence of dynamic recrystallization. A shear stress-strain curve for a sample that experienced no change in microstructure, Figure 4.19, was compared to a sample that had experienced significant microstructural change, Figure

4.20. It was established that a signature exists for the onset of dynamic recrystallization such that a peak stress is observed at a low critical strain followed by a significant drop in stress to a steady state value. The steady state value may oscillate about some mean value at a fairly regular interval signifying a constant competition between dynamic recrystallization and dynamic recovery.

Figure 4.19: Typical shear stress-strain curve generated for a specimen that does not undergo any significant microstructural changes.

Figure 4.20: Shear stress-strain curve generated for a specimen that experiences significant microstructural changes with evidence of dynamic recrystallization.

55 CHAPTER 5: CONCLUSIONS

1. Friction stir processing can be successfully employed on Ni-base alloys.

a. Processing windows were established for each alloy.

b. FSP results in significant reduction in grain size in the stir zone relative to

the base metal

c. Peak stir zone temperatures were 1150ºC for Hastelloy X and Alloy 625

and 1100ºC for Alloy 718

d. Increasing plunge depth leads to a proportional increase in normal force

and path force

e. Decreasing travel speed leads to a proportional decrease in normal force

and path force

f. Increasing tool rotation rate leads to a proportional decrease in torque.

2. FSP can be successfully utilized as a base metal treatment prior to fusion welding

to reduce susceptibility to HAZ liquation cracking

a. HAZ liquation cracking susceptibility decreases linearly with grain size

for both maximum and total crack length.

b. Relative susceptibility of the alloys can be ranked as

Hastelloy X

3. Physical simulation of FSP microstructure for Ni-base alloys using the modified

Gleeble hot torsion test has not been fully realized

a. Ni-base alloys require more total strain for reproduction of microstructure

b. The TMAZ of all three alloys can be successfully simulated

56 c. Shear stress-strain curves can be used to determine if dynamic

recrystallization has taken place during hot torsion testing.

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62 APPENDIX A

MICROSTRUCTURE AND MECHANICAL PROPERTIES OF THREE

FRICTION STIR PROCESSED NICKEL-BASE ALLOYS

63 A.1 ABSTRACT

Friction stir processing (FSP) of three Ni-base alloys, Alloy 625, Alloy 718, and

Hastelloy X utilizing a tungsten-rhenium tool was investigated. A processing window for each alloy was defined in terms of travel speed and tool rotation rate. Process forces and thermal histories were successfully recorded for each alloy. Peak stir zone temperatures were recorded at 1150ºC for Alloy 625 and Hastelloy X and 1100ºC for Alloy 718 using embedded type K thermocouples. FSP microstructures were investigated using optical and electron microscopy. Significant grain refinement was experienced by each alloy with stir zone grain sizes of 6µm, 5µm, and 4µm for Hastelloy X, Alloy 625, and Alloy

718, respectively, relative to starting grain sizes of 88µm, 26µm, and 44µm. Optical microscopy revealed streaking on the advancing side of the stir zone for each alloy.

Scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS) analysis showed these bands to be a result of tool wear. Hardness traverses show that all three alloys experience peak hardness adjacent to the stir zone in the thermomechanically- affected zone. Additionally, it was observed that Alloy 718 and Hastelloy X experienced slight hardening in the stir zone in comparison to the base metal whereas Alloy 625 underwent slight softening. Longitudinal sub-size tensile specimens extracted from stir zones showed that Hastelloy X experienced significant strengthening due to FSP relative to the base material whereas Alloy 625 and Alloy 718 experienced virtually no improvement in strength in the FSP region relative to base material. The FSP specimens from Hastelloy X material experienced ultimate tensile strengths near that of base metal

Alloy 625 and Alloy 718. For both base metal and FSP specimens, the relative ranking of the ultimate tensile strengths was found to be Hastelloy X

Friction stir processing (FSP) is a materials processing technique based on the basic principles of friction stir welding (FSW), where a non-consumable rotating tool consisting of a shoulder and pin is plunged into and translated across the surface of a work piece, Figure A.1 [1-3]. A homogenous, fine grain microstructure results due to the combined thermomechanical events that take place during FSP [3,4]. It has been demonstrated that refinement of microstructure has resulted in improved material performance such as strength, formability, corrosion, etc.[1,4-10]. The advantage of FSP over various other processes, such as accumulated roll bonding (ARB) and thermomechanical treatment (TMT), is that FSP is a localized, single step process whereas ARB and TMT affect the material in bulk and requires multiple steps to achieve the refined microstructure [5].

Figure A.1: Illustration of friction stir processing denoting the tool features (top) and the determination of advancing and retreating side (bottom).

65 Primary independent process variables for FSP consist of tool rotation speed, travel speed, and plunge depth. These primary process variables will affect heat generation rate, peak temperature, torque, and processes forces such as down force and path force. In general, it has been established that faster tool rotation rates and slower travel speeds lead to higher peak temperatures and lower torques and forces whereas faster travel speed and slower tool rotation rate leads to lower peak temperatures and higher process torque and forces [3]. A previous study by Song et al confirmed that increasing travel speed significantly decreases stir zone temperature, Figure A.2 [9].

Furthermore it was established that increased travel speed yields smaller grain size providing increased strength within the stir zone for Alloy 600, Figure A.3.

Figure A.2: Effect of friction stir welding travel speed on peak stir zone temperature and thermal history [9].

Further studies by Song and Nakata investigated the effect of FSW and post weld heat treatment (PWHT) of Alloy 625 and Alloy 718 [7,8]. Examination of microstructure showed that similar levels of grain refinement were realized by both materials in the stir zone. Base metal grain size for both alloys were between 5-20µm whereas the stir zones had grain sizes between 1-3µm. Further inspection of the stir zone microstructure revealed banding near the advancing side of the welds, Figure A.4. It was determined through energy dispersive x-ray spectroscopy (EDS) that the banding region resulted due to tool wear. Microhardness profiles for both as welded and PWHT conditions revealed that hardness in the stir zone is significantly higher than the base metal, Figure A.5. It was observed the PWHT samples experienced a significant increase in hardness even though grain coarsening had taken place. This was due to precipitation of M23C6 and

67 M6C carbides and γ”. Additionally, it is observed that the stir zone remains the hardest region of the welds after PWHT. Tensile specimens were extracted transverse and longitudinal to the welding direction. Testing showed that both the transverse and longitudinal FSW specimens were stronger than the base material whereas elongation was reduced. Furthermore, it was observed that the heat treated specimens had even higher strengths and less elongation than the both the base metal and the as welded specimens, Figure A.6.

Figure A.4: (a) Top view of Alloy 625 FSW, (b) banding structure observed in Alloy 625

SZ due to tool wear, (c) top view of Alloy 718 FSW, and (d) banding due to tool wear in

SZ of Alloy 718 [7,8].

Figure A.5: Microhardness profiles of as welded and PWHT FSW in (a) Alloy 625 and

(b) Alloy 718 [7,8].

Figure A.6: Tensile testing summary for as welded and PWHT FSW (a) Alloy 625 and

(b) Alloy 718 [7,8].

69 A.3 EXPERIMENTAL PROCEDURE

¼” X 3” X 6” Hastelloy X, Alloy 625, and Alloy 718 plates, compositions listed in Table A.1, were subjected to FSP utilizing a tungsten-25 weight percent rhenium (W-

25Re) tool (Figure A.7) at a 3º push angle. FSP was carried out on an Accustir (General

Tool Co. Cincinnati, OH) gantry style friction stir welding machine, Figure A.8. Select coupons of each alloy were embedded with type K thermocouples to record thermal histories during processing, schematic shown in Figure A.9. The thermocouple pattern was chosen based on the profile of the tool and to provide insight into possible asymmetrical temperature behavior in the stir zone due to velocity vector mismatch between the advancing and retreating side with respect to the welding direction and possible increased heat near the shoulder of the tool due to friction. Process forces were recorded by the friction stir welding machine. Transverse sections of FSP‟d plates were mounted polished with a final step on a vibratory polisher in 0.05µm colloidal silica fluid. Hastelloy X and Alloy 718 samples were electrolytically etched using 10% oxalic acid at 0.53Amp/cm2 for 10-20 seconds. Alloy 625 samples were electrolytically etched using Lucas‟ reagent (150mL HCl, 50mL lactic acid, 3g oxalic acid) at 2V. Micrographs were taken utilizing an Olympus GX51 microscope and macrographs were taken using a

Nikon SMZ1000 stereo microscope outfitted with a Nikon Digital Sight DS-L1 digital photomicrograph system. Samples subjected to electron microscopy were left un-etched after final polishing and utilized a Quanta 200 scanning electron microscope (SEM) outfitted with and EDAX energy dispersion spectrometer (EDS). Higher resolution images utilized a Sirion field-emission gun (FEG) SEM. Microhardness was performed using an AMH43 (Leco) automated microhardness system with a Vickers diamond 70 indenter. ASTM E8-09 sub-size tensile specimens [11], Figure A.10, were extracted from base material and longitudinal to the FSP region. Testing was carried out on an

MTS 810 load frame at a rate of 0.05 in/sec.

Table A.1: Alloy compositions.

Figure A.7: W-25Re tool utilized for FSP.

Figure A.8: Accustir friction stir welding machine.

Figure A.9: Schematic of thermocouple placement with relationship to tool profile

(dimensions in inches). A 304L backing plate with a through hole pattern and relief channels was utilized to avoid damaging thermocouple leads.

Figure A.10: ASTM E8 sub-size tensile test specimen geometry [11]. Thickness of samples was 0.125 in.

73 A.4 RESULTS AND DISCUSSION

A.4.1 Process Parameter Windows

Process windows were established for each alloy by varying travel speed and tool rotation rate and recording which settings resulted in a friction stir process visually free of discontinuities such as top side voids or wormholes present in the keyhole left from withdrawing the tool from the material, Table A.2. The processing window for Hastelloy

X was found to lie between 125-180 RPM and 1.5-2 IPM, Figure A.11. Alloy 625 demonstrated a wider processing window between 100-145 RPM and 1-3 IPM, Figure

A.12. This window is at considerably lower RPM and travel speed than the 200 RPM and 3.94 IPM reported by Song and Nakata [7]. These differences in parameters are likely due to the use of a smaller tool of different composition and thinner sheets by Song and Nakata. The Alloy 718 processing window is quite narrow lying between 100-110

RPM and 1.5-2 IPM, Figure A.13. Again the window is much lower than parameters used by Song’s and Nakata’s [8] study of FSW of Alloy 718 and is likely due to their use of thinner material and smaller tool of different composition.

74 Table A.2: List of selected FSP runs.

Figure A.11: Friction stir processing window established for Hastelloy X. Open circles denote parameters containing visual discontinuities.

Figure A.12: Friction stir processing window for Alloy 625. Open circles denote parameters containing visual discontinuities. The diamond represents settings used by

Song and Nakata [7].

Figure A.13: Friction stir processing window for Alloy 718. Open circles denote parameters with visual discontinuities. The diamond represents settings used by Song and Nakata [8].

To compare process parameter windows of the materials the dimensionless heat index was calculated for each processing run, included in Table A.2. The dimensionless heat index was calculated according to work by Nandan et al (Equation A.1) [3] where f is process efficiency, σT is the yield stress at temperature T, A is the area of the tool

T shoulder, ω is the tool rotation rate, Cp is the specific heat at temperature T, κT is the thermal conductivity at temperature T, and U is the travel speed. For simplicity, an efficiency value of 1 was used throughout the calculations. From the dimensionless heat index it was shown that Alloy 718 required higher heat than Alloy 625 and Hastelloy X with a heat index window between 20100 and 39300. Hastelloy X required relatively intermediate heat with the narrowest heat index window between 15800 and 22800.

Alloy 625 accommodated a range of heat with a heat index window between 10900 and

28300. 77 Equation A.1: Dimensionless heat index for friction stir welding and processing [3].

A.4.2 Thermal History

Once processing windows were established, a nominal setting was selected for each material based on appearances and repeatability. These settings were used for acquiring thermal histories of FSP; 180 RPM and 2 IPM for Hastelloy X, 100 RPM and

2.5 IPM for Alloy 625, and 100 RPM and 2 IPM for Alloy 718. The thermal history of the stir zones were successfully captured and consistent for each material, represented in

Figure A.14. Sectioning of these FSP runs shows clearly that thermocouples were incorporated into the stir zone, Figure A.15. An 1150ºC peak stir zone temperatures was recorded for both Hastelloy X and Alloy 625 while Alloy 718 experienced an 1100ºC peak stir zone temperature.

Figure A.14: Acquired thermal history for Alloy 625 FSP run 40 (100 RPM, 2.5 IPM).

Figure A.15: Thermocouple locations relative to stir zone for Alloy 625 FSP run 40.

79 To compare cooling rates of the alloys during FSP, the instantaneous cooling rate was calculated and plotted as a function of temperature, Figure A.15. It was observed that all three alloys experienced nearly the same cooling rates although the dimensionless heat index for each material is quite different. Plotting the thermal conductivity as a function of temperature demonstrates how the thermal conductivity may correct for these different dimensionless heat indexes, Figure A.16. It is clearly demonstrated that

Hastelloy X and Alloy 718 have very similar thermal conductivity however Alloy 625 thermal conductivity is significantly lower. The low thermal conductivity of Alloy 625 should provide a slower cooling rate than Hastelloy X and Alloy 718 however the heat index for Alloy 625 (12600) is much lower than the heat index for Hastelloy X (22800) and Alloy 718 (20100).

Figure A.16: Instantaneous cooling rates encountered during FSP Ni-base alloys.

Figure A.17: Thermal conductivities as a function of temperature [12-14].

A.4.3 Process Forces

Process force data was collected from the Accustir FSW machine to observe how process inputs of travel speed, tool rotation rate, and plunge depth affect processing in terms of normal force, path force, and torque. Hastelloy X run 22 had changes made to the plunge depth and travel speed, Figure A.18. It was observed that increasing the plunge depth created a proportional increase in normal force, path force, and torque.

These proportional increases can be attributed to increased volume of material being affected by the tool. It was also observed that decreasing travel speed led to a proportional decrease in normal force and path force whereas torque was only slightly reduced. The decreases in normal force and path force due to decreased travel speed can be attributed to allowing more time to generate heat, further softening material in front of and in contact with the tool. The resultant softened material would require less torque to

81 rotate the tool. Hastelloy X run 9 had no changes made to parameters during processing.

It was observed the processing is fairly steady-state for set parameters, Figure A.19.

Figure A.18: FSP run 22 showing how (top) process forces and torque are affected by

(bottom) process inputs of plunge depth and travel speed. (P denotes plunge series, dotted vertical lines denote when forward travel starts and stops, W denotes withdrawal of tool.)

84 Similar FSP runs were made for Alloy 625 and Alloy 718 in which parameters of travel speed, tool rotation rate, and plunge depth were varied to observe how process forces reacted, Figure A.20 and Figure A.21. It was observed that plunge force and travel speed had similar effects to those experienced by Hastelloy X; increase in plunge depth led to proportional increases in normal force and path force and travel speed decreases led to proportional decreases in normal and path force. Again the proportional increase in normal and path force due to increased plunge depth can be attributed to increasing the volume of material being affected by the tool. Also, decreasing forces proportional to decreasing travel speed can be attributed to increased heating in front of the tool softening material in front of and in contact with the tool. The effect of changing tool rotation rate was noticed most in the torque output. As tool rotation rate increased torque decreased proportionally attributed to increasing heat generation which softens the material requiring less torque to spin the tool. Process forces for a FSP run in Alloy 625 with no changes to parameters were recorded to show the steady-state nature of the process, Figure A.22.

Figure A.20: FSP run 15 showing how (top) process forces and torque are affected by

(bottom) process inputs of plunge depth, tool rotation rate, and travel speed. (P denotes plunge series, dotted vertical lines denote when forward travel starts and stops, W denotes tool withdrawal.)

Figure A.21: FSP run 29 showing how (top) process forces and torque are affected by

(bottom) process inputs of plunge depth, tool rotation rate, and travel speed. (P denotes plunge series, dotted vertical lines denote when forward travel starts and stops, W denotes tool withdrawal.)

To understand how material properties and process parameters interact, the steady state normal and path force and torque were collected, shown in Table A.3. Additionally, the process parameters were listed in terms of the previously discussed dimensionless heat index and the heat input calculated according to Lienert et al [15] (Equation A.2) where f is the process efficiency, ω is the tool rotation rate, τ is the torque output, and U

88 is the travel speed. Comparison of the dimensionless heat index and the heat input are in good agreement showing that Hastelloy X was processed at the highest heat input/index and Alloy 625 at the lowest, with values for Alloy 718 being close to those of Hastelloy

X. Comparing the material room temperature yield tensile stress it is demonstrated that the strengths are ordered as Hastelloy X

Alloy 625 process forces and torque was very close to values obtained for Alloy 718.

This observation is accounted for by the heat input/index. It was shown that the heat input/index for Alloy 625 is much lower than Alloy 718. This lower heat input may require more force and torque to move the tool through the material. For similar heat input/index values, a material with a higher yield stress will require higher process forces and torque as demonstrated by the differences between Hastelloy X and Alloy 718 (Table

A.3).

Equation A.2: Friction stir welding heat input equation [15].

Table A.3: Comparison of process forces, material strength, and heat input/index.

89 A.4.4 Microstructure

All three alloys exhibited microstructures similar to friction stir welds and processes found in other materials. Each contained distinct microstructures for the stir zone, thermomechanically-affected zone (TMAZ) and the base metal, Figure A.23.

There were no noticeable heat-affected zones in any of the alloys. It was noticed for all three alloys that a feathered streaking occurred on the advancing side of the stir zones

(dark etching) and is thought to be a result of tool wear. To quantify the level of grain refinement resulting from FSP, base metal and stir zone microstructure was evaluated in terms of grain size using the Abrams three circle method [16]. Hastelloy X base metal average grain diameter was measured to be 88µm while the stir zone average grain diameter was on the order of 6µm, Figure A.24, a reduction of 93%. Alloy 625 base metal average grain diameter was measured to be 26µm with a stir zone grain diameter of

5µm, Figure A.25, an 81% reduction. Alloy 718 base metal average grain diameter was measured as 44µm with a stir zone grain size of 4µm, Figure A.26, a 91% reduction. It was noticed for both Alloy 625 and Alloy 718 that the necklace structure of the base metal due to rolling with irregular distribution of carbides and second phases (dark etching phases) became equiaxed and more evenly distributed in the stir zones, however

NbC (dark etching phase) particle size is preserved. Higher magnification images of the stir zone were obtained to more clearly show the grain size and distribution, Figure A.27.

The TMAZ microstructure is similar for all three alloys, Figure A.28. It was clearly observed that prior base metal grains were affected thermally and mechanically during processing due to the elongation of the grains and an apparent substructure within parts of these grains. 90

Figure A.23: Optical macrographs of cross sectioned friction stir processed (i) Hastelloy

X, (ii) Alloy 625 and (iii) Alloy 718 (10X). Note feathered streaking (dark etching) on the advancing side of the stir zone for all three alloys.

Figure A.24: Optical micrograph of Hastelloy X (A) base metal with 88µm grain size and (B) stir zone with 6µm grain size (500X, oxalic etch).

Figure A.25: Optical micrograph of Alloy 625 (A) base metal with 26µm grain size and

(B) stir zone with 5µm grain size (500X, Lucas' reagent). Note the NbC particles are unchanged from the base metal.

Figure A.26: Optical micrograph of Alloy 718 (A) base metal with 44µm grain size and

(B) stir zone with 4µm grain size (500x, oxalic etch). Note the NbC particles are unchanged from the base metal.

Figure A.27: Optical micrograph of FSP stir zone for (A) Hastelloy X, (B) Alloy 625, and (C) Alloy 718 (1000X). Particles in (B) and (C) are the NbC from the base metal.

Figure A.28: Optical micrographs of FSP TMAZ for (A) Hastelloy X, (B) Alloy 625, and (C) Alloy 718 (500X).

95 An Alloy 625 FSP specimen was observed in a high resolution scanning electron microscope to further investigate the feathered streaking that occurs on the advancing side of the stir zone microstructures. It is apparent that the phase that creates the streaking is of high atomic number due its high z-contrast (brightness) using the back- scatter detector, Figure A.29. At higher magnification, it was observed that the phase is discontinuous and appears as discrete particles, Figure A.30. EDS analyses of the particles were compared to the surrounding matrix and were found to have a high intensity count for tungsten, Figure A.31. Analysis of the intensity maps revealed that the particles are primarily composed of tungsten, Table A.4, and obviously a result from tool degradation.

Figure A.29: High resolution SEM image of Alloy 625 stir zone (left) showing general appearance of streaks on advancing side (250X) and (right) higher magnification of bright phase (6000X).

Figure A.31: EDS count intensities for Alloy 625 (top) matrix and (bottom) particles.

98 Table A.4: Quantification of EDS scans showing matrix to be absent of tungsten and bright particles to consist primarily of tungsten.

A.4.5 Mechanical Properties

Hardness traverses revealed somewhat conflicting results with previous research.

Traverses for Hastelloy X and Alloy 718 showed a slight increase in hardness values for the stir zone relative to the base metal, Figure A.32 and Figure A.33. Alloy 625 shows a slight decrease in hardness for the stir zone relative to the base metal, Figure A.34. All three alloys exhibited peak hardness in the TMAZ, due to work hardening from residual strain due to plastic deformation in this region. Previous studies on Alloy 625 and 718 by, Song and Nakata show that the FSW joint in the as welded condition for both alloys experiences peak hardness in the stir zone with a significantly higher hardness value than the respective base metal, Figure A.35 [7,8]. The much finer stir zone grain size (1-3µm) experienced by Song and Nakata may contribute to the higher hardness values.

Additionally, softening in the stir zone experienced by Alloy 625 may be attributable to prior base metal conditions such as an incomplete or low temperature solution annealing in order to meet grain size requirements resulting in residual strain from cold work.

Furthermore, the microhardness profiles may provide insight into the DRX mechanism resulting from friction stir processing these materials. Luton and Sellars have shown in early studies with hot torsion testing that two DRX process may take place,

99 discontinuous DRX and continuous DRX [17]. Continuous DRX is characterized as a process where complete dynamic recrystallization occurs such that new grains form as dislocations are introduced into the material. Discontinuous DRX is characterized as a process where dynamic recrystallization is incomplete such that new grains form when critical density of dislocations occurs followed by inducing new dislocations in the newly formed grains. The stir zone softening encountered by Alloy 625 may suggest that, for these particular process parameters, the small grains resulted due to a continuous DRX behavior and posses very little dislocation networks within these grains. On the other hand, for the process parameters used for Hastelloy X and Alloy 718 may have induced a discontinuous DRX mechanism to form the smaller grain size in the stir zone, resulting in finer grains with a fairly high density of dislocations networks resulting in hardening in the stir zone.

100

Figure A.32: FSP Hastelloy X hardness profile. (BM=base metal, T=TMAZ, SZ=stir zone.)

Figure A.33: FSP Alloy 718 hardness profile. (BM=base metal, T=TMAZ, SZ=stir zone.)

101

Figure A.34: FSP Alloy 625 hardness profile. (BM=base metal, T=TMAZ, SZ=stir zone.)

102 Sub-size tensile specimens were extracted from base metal and FSP regions for all three alloys. Testing revealed that friction stir processed Alloy 625 and Alloy 718 experience marginal strength and percent reduction of area increases compared to the base metal, Figure A.36. However, it was observed that friction stir processed Hastelloy

X experienced a substantial strength (48%) and percent reduction of area (50%) increase relative to the base metal. This resultant increase in strength for Hastelloy X gives it an ultimate tensile strength close to that of Alloys 625 and 718 base metals. The marginal strength increases for Alloy 625 may be due to the base metal not being fully annealed resulting in residual strain in the base metal. Additionally, both Alloys 625 and 718 had a preserved carbide particle size between the base and FSP metal whereas Hastelloy X did not preserve carbide size. These particles aid in the pinning of grain boundaries and make dislocation motion more difficult. If grain boundary is increased but carbide particle size remains the same, redistribution of stresses along more grain boundary area is the only strengthening mechanism. However, if both grain and carbide particle size is reduced, along with finer distribution of the carbides, strengthening is achieved by both increased pinning agents and redistribution of stress along more grain boundary area.

103

A.5 CONCLUSIONS

Friction stir processing was successfully applied to three Ni-base alloys, Hastelloy X,

Alloy 625 and Alloy 718.

1. Processing windows utilizing the W-25Re tool were established for all three

alloys.

2. Processing results in significant grain refinement with stir zone grain sizes of

6, 5, and 4µm for Hastelloy X, Alloy 625 and Alloy 718, respectively.

3. Tool wear is encountered by all three materials and is consistently observed

on the advancing side of the stir zone.

104 4. Hastelloy X and Alloy 718 experience slight hardness increases in the stir

zone, attributable to the finer grain size and distribution. All three alloys

experience peak hardening in the TMAZ due to work hardening from residual

strain due to plastic deformation.

5. Thermal histories were successfully recorded for settings central to the

processing window with peak temperatures for Hastelloy X and Alloy 625 of

1150ºC and 1100ºC for Alloy 718.

6. Investigation of process force and torque feedback shows a clear relationships

between material properties, process parameters and machine force and torque

output.

 Increasing plunge depth, increases the volume of material affected

by the tool proportionally increasing forces.

 Decreasing travel speed decreases process forces attributed to

increasing the heat generated which allows softening of material in

front of and in contact with the tool.

 Increasing tool rotation rate decreases torque proportionally due to

increased heating by the tool, softening material around and in

contact with the tool.

 Materials with higher yield strength require more force and torque

for a given heat input.

7. Hardness traverses do not indicate a clear relationship between grain size and

hardness value

105  Peak hardness for all three alloys was experienced in the TMAZ

due to residual strain from plastic deformation in this region.

 Hastelloy X and Alloy 718 experienced hardening in the stir zone

relative to the base metal.

 Alloy 625 experienced softening in the stir zone relative to the

base metal and may be due to the base metal having residual strain

resulting from incomplete annealing to control base metal grain

size.

8. Tensile testing did not show a clear relationship between grain size and

ultimate tensile strength

 Hastelloy X showed the most improvement in tensile properties

due to FSP with a strength increase of 48% and a reduction of area

increase of 50%

 Both Alloys 625 and 718 showed marginal improvements in

strength and reduction of area due to FSP, attributed to base metal

condition and conservation of carbide particle size and distribution.

106 A.6 REFERENCES

[1] R.S. Mishra, M.W. Mahoney, S.X. McFadden, N.A. Mara, and A.K. Mukherjee,

"HIGH STRAIN RATE SUPERPLASTICITY IN A FRICTION STIR PROCESSED

7075 Al ALLOY," Scripta Materialia, vol. 42, pp. 163-168, 2000.

[2] W.M. Thomas et al., "Friction-stir butt welding," 9125978.8, December 1991.

[3] R. Nandan, R. DebRoy, and H.K.D.H. Bhadeshia, "Recent advances in friction-stir

welding - Process, weldment structure and properties," Progress in Materials Science,

vol. 53, pp. 980-1023, 2008.

[4] R.S. Mishra and M. Mahoney, "Friction Stir Processing," in Friction Stir Welding and

Processing, R.S. Mishra and M. Mahoney, Eds. Materials Park, OH, U.S.A.: ASM

International, 2007, pp. 309-350.

[5] Z.Y. Ma, "Friction Stir Processing Technology: A Review," Metallurgical and

Materials Transactions A, vol. 39A, pp. 642-658, 2008.

[6] Y.S. Sato, P. Arkom, H. Kokawa, T.W. Nelson, and R.J. Steel, "Effect of

microstructure on properties of friction stir welded Inconel Alloy 600," Materials

Science and Engineering A, vol. 477, pp. 250-258, 2008.

[7] K.H. Song and K Nakata, "Effect of precipitation on post-heat-treated Inconel 625

alloy after friction stir welding," Materials and Design, vol. 31, pp. 2942-2947, 2010.

107 [8] K.H. Song and K. Nakata, "Microstructural and mechanical properties of friction-

stir-welded and post-heat-treated Inconel 718 alloy," Journal of Alloys and

Compounds, vol. 505, pp. 144-150, 2010.

[9] K.H. Song, H. Fujii, and K. Nakata, "Effect of welding speed on microstructural and

mechanical properties of friction stir welded Inconel 600," Materials and Design,

vol. 30, pp. 3972-3978, 2009.

[10] F Ye, H Fujii, T Tsumura, and K Nakata, "Friction stir welding of Inconel alloy

600," Journal of Materials Science, vol. 41, pp. 5376-5379, 2006.

[11] ASTM International, Physical Testing Standards and Mechanical Testing Standards,

2009.

[12] Haynes International, Inc., Hastelloy X Alloy, 1997.

[13] Special Metals Corporation, Inconel Alloy 625 Technical Bulletin, January 2006.

[14] Special Metals Corporation, Inconel Alloy 718 Technical Publication, September

2007.

[15] T.J.,Sellwag, W.L. Jr., Gould, J.E. Lienert, "Comparison of Heat Inputs: Friction Stir

Welding vs. Arc Welding," in Abstracts of the 2002 AWS Convention, 2002.

[16] ASTM International, Standard Methods for Determining Grain Size in Materials,

2010.

[17] M.J. Luton and C.M. Sellars, "Dynamic recrystallization in nickel and nickel-iron

alloys during high temperature deformation," Acta Metallurgica, vol. 17, pp. 1033-

1043, 1969.

108 APPENDIX B

APPLICATION OF FRICTION STIR PROCESSING FOR IMPROVING HEAT-

AFFECTED ZONE LIQUATION CRACKING RESISTANCE OF NICKEL-BASE

ALLOYS

109 B.1 ABSTRACT

Friction stir processing (FSP) was evaluated as a means to reduce the susceptibility of two solid-solution strengthened Ni-base alloys, Alloy 625 and Hastelloy

X, and Ni-base superalloy, Alloy 718, to heat-affected zone (HAZ) liquation cracking.

FSP of the base metal resulted in a large reduction in grain size (average grain diameter less than 10µm). These grains contained highly tortuous grain boundaries with a very fine distribution of second phases. Cracking resistance was evaluated using the spot- varestraint test. Testing showed a reduction in cracking susceptibility relative to the parent material due to the reduction in HAZ grain size. Optical microscopy evaluation revealed that HAZ liquation cracking resistance is enhanced due to increased grain boundary areas resulting in finer, more discontinuous networks of low melting eutectic and second phases.

B.2 INTRODUCTION

Friction stir processing is a derivative of friction stir welding (FSW) where the rotating tool is limited to a partial penetration, non-joining action. This solid-state process results in microstructures similar to FSW due to the same dynamic recrystallization phenomenon resulting from intense plastic deformation at high temperature [1]. Typical advantages of the resultant microstructure are finer grain size, stress redistribution, and second phase redistribution.

HAZ liquation of heat resistant Ni-base alloys has been studied for more than 40 years [2]. The mechanism of this phenomenon has yet to be fully understood but qualitative mechanisms have been proposed, primarily, the segregation, penetration, and constitutional liquation mechanisms [3-5]. All mechanisms describe a process where 110 solute and/or eutectic constituents promote local melting and wet the grain boundaries in the partially melted zone (PMZ) of the HAZ. These liquated boundaries possess insufficient strength necessary to accommodate shrinkage stresses encountered across the fusion boundary and/or extrinsic restraint placed on the weld [3]. Alloys widely reported in the literature that experience HAZ liquation cracking are Alloy 625, Hastelloy X, and

Alloy 718 [2-16].

Fundamentally, HAZ liquation cracking is affected by three factors: material composition and condition, welding parameters, and restraint [3]. With regards to material condition, base metal microstructure with fine grain size provides better resistance to HAZ liquation than coarse-grained material [7, 8]. Controlling composition to limit the amount of low melting constituents, such as carbides, also improves resistance [4-6]. However, this is not a typical approach as these constituents are the result of intentional additions to the alloy to improve mechanical properties, corrosion resistance, oxidation resistance, etc. Restraint is typically difficult to control and quantify and cannot fully be eliminated due to shrinkage of weld metal in the fusion zone during solidification. This shrinkage imparts some level of stress on the adjacent HAZ near the fusion boundary leading to cracking. The objective of this research is to evaluate FSP as a local grain refinement process as a method to reduce HAZ liquation cracking susceptibility of Alloy 625, Hastelloy X, and Alloy 718.

B.3 EXPERIMENTAL PROCEDURES

Three different Ni-base alloys were used in this study, Alloy 625, Hastelloy X, and Alloy

718. Test coupons were produced from wrought plate of each alloy measuring ¼” X 3”

X 6” nominally. The compositions of these alloys are listed in Table B.1. Test coupons 111 were then subjected to friction stir processing along the 6” length of the material. FSP was carried out on a gantry style GTC Accustir friction stir welding machine utilizing a

W-25Re tool with truncated cone geometry, Figure B.1. FSP parameters for each alloy can be found in Table B.2.

Table B.1: Material compositions.

Figure B.1: W-25Re tool profile.

Table B.2: Friction stir processing parameters.

112 Upon completion of FSP, the coupons were machined to remove the surface roughness, as shown in Figure B.2, to reduce any geometric cracking effect during spot- varestraint testing. Machined coupons were then subjected to spot-varestraint testing as a means to evaluate HAZ liquation cracking susceptibility [9], Figure B.3. Coupons were aligned in the testing fixture such that the spot weld straddles the FSP region as to get a simultaneous test of parent and FSP‟d material, Figure B.4. Spot-varestraint test parameters are found in Table B.3. Furthermore, ThermoCalc® was employed utilizing the Ni-base alloy database TTNI7 to predict low melting constituents during on heating.

Table B.3: Spot-varestraint test parameters.

113

Figure B.2: As FSP material (top) showing striations on surface. Surface machined FSP material eliminating striations (bottom).

114

Figure B.3: Schematic of spot-varestraint testing apparatus.

Figure B.4: Illustration of gas tungsten arc spot weld placement relative to stir zone.

115 Upon completion of the spot-varestraint test, coupons were lightly hand polished with 800 grit SiC paper, in order to remove any oxide layer and to enhance definition of cracks. Inspection of the spot welds was done in a Nikon stereo microscope at 20X magnification. Cracking was measured using a Nikon SMZ1000 stereo microscope outfitted with a Nikon Digital Sight DS-L1 digital photomicrograph system. Crack lengths were measured perpendicular to the fusion boundary of the spot weld, Figure B.5.

Further metallographic preparation included mounting and polishing of plan view sections of GTA spot welds to a final step using a vibratory polisher in a solution of

0.05µm colloidal silica. Optical microscopy specimens were etched electrolytically in either 10% oxalic acid solution at 0.53Amp/cm2 for Hastelloy X and Alloy 718 or Lucas‟ reagent (150mL HCl, 50mL lactic acid, 3g oxalic acid) at 2V for Alloy 625. Electron microscopy specimens remained un-etched.

116

Figure B.5: Example of crack measurement (20X).

B.4 RESULTS AND DISCUSSION

Investigation of FSP regions revealed considerable reduction in grain size relative to the base metal for all three alloys, Figure B.6. Using the Abrams three circle procedure [17] for determining grain size it was found that the base metal of Hastelloy X exhibits an average grain diameter of 88µm whereas the stir zone has an average grain diameter of 6µm. Similar microstructure refinements were observed in both Alloy 625 and 718, Figure B.6. The magnitude of grain size reduction is not as large as what is observed in Hastelloy X, due to these alloys having smaller initial grain size, however grain sizes in the stir zones are similar. Again using the Abrams three circle method for determining grain size it was determined that the base metal average grain diameter was

117 26µm and 44µm whereas the average FSP stir zone grain diameter was 5µm and 4µm for

Alloy 625 and Alloy 718, respectively. It was observed that both Alloy 625 and Alloy

718 exhibited a rolled microstructure in the base metal with grains elongated in the direction of rolling. Furthermore, it was noticed that Alloy 718 had an irregular distribution of grain sizes and secondary phases (dark etching constituents). Friction stir processing of both materials not only resulted in the finer grain size but a finer, more regular distribution of grain size and secondary phases, with a conserved NbC particle size between the base metal and stir zone.

118

(500X, Oxalic).

119 The susceptibility to HAZ liquation cracking was determined in terms of maximum crack length (MCL) and total crack length (TCL). The summary of the results is shown in Table B.4, comparing parent material behavior to that of FSP material.

Testing showed the greatest overall improvement from FSP was experienced by Alloy

718 with the highest percentage of improvement in terms of both MCL and TCL for all tested strains, Figure B.7 and Figure B.8, respectively. It was also shown that Alloy 625 experienced moderate improvement and Hastelloy X had the least improvement. It is well known that Nb contributes to the low melting constituents for both Alloy 625 and

718 in the form of carbides (NbC) and delta phase (Ni3Nb). Low melting constituents for

Hastelloy X are other carbides (M23C6, M6C) and complex phases (Laves and P phase)

[2-4]. It is also observed that relative susceptibility of each alloy to HAZ liquation cracking can be ranked as Hastelloy X

120

Figure B.7: Comparison of maximum crack length.

Figure B.8: Comparison of total crack length.

121 It has been established that susceptibility of Ni-base alloys to solidification cracking can be ranked according to solidification temperature range (STR) and that this range is strongly influenced by composition [3]. A similar relationship between composition and temperature could be developed for HAZ liquation cracking tendencies.

Investigation of the thermodynamic simulations performed using ThermoCalc® revealed differences between the alloys in terms of a temperature difference during heating between when the alloy just begins to form liquid and when it is fully liquid, Table B.4 and Figure B.9. It is observed that the fraction of the melting temperature at which liquid begins to form for the alloys can be ranked as Hastelloy X

Correlation of the spot-varestraint results with these thermodynamic simulations is in good agreement in that it shows that Hastelloy X is the least susceptible while Alloy 718 is the most susceptible.

122

Figure B.9: Illustration of determination of temperature at which liquid begins to form using ThermoCalc® simulations.

Table B.4: ThermoCalc® simulation results.

Plan view sections of the GTA spot welds were examined using scanning electron microscopy (SEM) and energy dispersive spectroscopy (EDS) to evaluate observed intergranular phases observed optically, Figure B.10, Figure B.11, and Figure B.12.

Intergranular films observed in Hastelloy X samples were very fine and extended from 123 the PMZ towards the crack tips, Figure B.13. EDS spot evaluation of the matrix and the intergranular films were performed to compare elemental intensity peaks, Figure B.14. It was observed that the intergranular films had higher peak intensities in Cr and Mo than the matrix, indicating that Cr23C6 and Mo6C were responsible for liquation. Intergranular films observed in Alloy 625 were found to be somewhat thicker than those found in

Hastelloy X, Figure B.15. EDS spot evaluation of the matrix and intergranular films showed increased intensities for Mo and Nb, Figure B.16. Intergranular films observed in Alloy 718 were similar in morphology and content to Alloy 625, Figure B.17. EDS spot evaluation revealed increased intensities for Mo and Nb, Figure B.18, signifying formation of Mo6C and either the niobium rich NbC or delta phase. Quantification of the peak intensities provided the relative weight percent of each element, Table B.5.

124

Figure B.10: Optical micrograph of Hastelloy X spot-varestraint HAZ in (left) base metal and (right) FSP region (200X). Arrows point to areas of intergranular film formation.

Figure B.11: Optical micrograph of Alloy 625 spot-varestraint HAZ in (left) base metal and (right) FSP region (200X). Arrows point to areas of intergranular film formation.

125

Figure B.12: Optical micrograph of Alloy 718 spot-varestraint HAZ in (left) base metal and (right) FSP region (200X). Arrows point to areas of intergranular film formation.

Figure B.13: Hastelloy X SEM images of (A) liquid extension (light phase) from PMZ to crack opening (1500X) and (B) eutectic film formation (6000X). Crosses denote where

EDS spots were taken.

126

Figure B.14: EDS intensities for Hastelloy X (top) matrix and (bottom) intergranular film regions.

127

Figure B.15: Alloy 625 SEM images of (A) liquid extension (light phase) from PMZ to crack opening (1500X) and (B) eutectic film formation (6000X). Crosses denote where

EDS spots were taken.

128

Figure B.16: EDS intensities for Alloy 625 (top) matrix and (bottom) intergranular film regions.

129

Figure B.17: Alloy 718 SEM images of (A) liquid extension (light phase) from PMZ to crack opening (150X) and (B) eutectic film formation (2500X). Crosses denote where

EDS spots were taken.

130

Figure B.18: EDS intensities for Alloy 718 (top) matrix and (bottom) intergranular film regions.

131 Table B.5: EDS quantification of matrix and intergranular (IG) film regions resulting from spot-varestraint testing.

A previous study by Thompson et al correlates the susceptibility of HAZ liquation cracking in terms of TCL for Alloy 718 to HAZ grain size [8]. It was demonstrated that finer HAZ grain size has a significant effect in reducing susceptibility to HAZ liquation

[8]. Additionally, it was shown that the response of susceptibility to grain size is linear,

Figure B.19. This relationship was furthered by Woo et al in a study of grain size effect on HAZ liquation susceptibility of cast Alloy 718 [10]. In this study Woo et al found a linear response to grain size for both total crack length and maximum crack length,

Figure B.20.

132

Figure B.19: HAZ liquation susceptibility response to grain size for Alloy 718 (from

Thompson et al) [8].

133

Figure B.20: Effect of grain size on HAZ liquation cracking in terms of (a) total crack length and (b) maximum crack length (cast Alloy 718) [10].

Plan view sections of the GTA spot welds were prepared to measure the HAZ grain sizes in the FSP and base metal. Due to the narrow size of the HAZ the linear intercept method was used to determine the average grain size. Trends of grain size effect on HAZ liquation for Alloys 625 and 718, Figure B.21 and Figure B.22, respectively, follow previous studies quite well in terms of both MCL and TCL. It is evident that both alloys realize a decrease in susceptibility as grain size decreases.

Hastelloy X material experiences a similar trend to Alloy 625 and Alloy 718 for reduction of the MCL due to finer grain size, Figure B.23 (A). It is observed that

Hastelloy X shows a more significant response to grain size at higher strains, demonstrated by the steepening of the slope of the trend line for 7% strain. However, the

TCL response behavior of Hastelloy X is much different than Alloys 625 and 718, Figure

B.23 (B). It is seen that TCL increases as grain size decreases for 3% and 5% strains 134 whereas TCL decreases with decreasing grain size for 7% strain. This behavior is attributed to the difference in the frequency of cracks observed in the HAZ of the FSP and base metal regions, Figure B.24. It is observed that Hastelloy X experiences a significant increase (>10) in crack frequency for both 3% and 5% strain in the FSP portion of the GTAW HAZ whereas both Alloy 625 and Alloy 718 experience an increase in frequency of less than 7 cracks. This significant increase in crack frequency in the FSP HAZ region of contributes considerably to the total crack length in Hastelloy

X spot-varestraint tests.

135

Figure B.21: Alloy 625 HAZ liquation crack response to grain size in terms of (A) maximum crack length and (B) total crack length.

136

Figure B.22: Alloy 718 HAZ liquation crack response to grain size in terms of (A) maximum crack length and (B) total crack length.

137

Figure B.23: Hastelloy X HAZ liquation crack response to grain size in terms of (A) maximum crack length and (B) total crack length.

138

Figure B.24: HAZ liquation crack frequency comparison between alloys and base metal and FSP HAZ.

A drawback to the spot-varestraint test is the total amount of heat delivered to the test piece due to the high arc current and long hold times required to generate a stable weld pool. For this study, an arc current of 180 amps with a 20 second hold time was utilized to form the spot weld resulting in large HAZ grain sizes. To better simulate real field welding, linear GTAW were made along the FSP/base metal interface to track HAZ grain size response to arc weld heat input. It was observed that HAZ grain size response to arc weld heat input in the base metal is quite different from the FSP material, shown in

Figure B.25. It was observed that base metal HAZ grain sizes were larger than FSP HAZ grain sizes for all heat inputs. Additionally, it was noted that the Alloy 718 and Hastelloy

X base metal HAZ grain size experienced rampant grain growth between heat inputs of

24.4kJ/in and 28.8kJ/in. whereas the FSP HAZ grain size experienced limited growth due 139 to the increase in heat input. From these trends it is possible that field welding will benefit more from grain refinement than what the spot-varestraint test indicates.

Figure B.25: Arc weld HAZ grain size response to heat input. Solid lines represent FSP material and dashed lines represent parent material.

B.5 CONCLUSIONS

1. Resistance to HAZ liquation cracking experienced in Ni-base alloys can be improved by refining grain size.

2. Friction stir processing is a viable technique to achieve grain refinement and can be employed locally to regions of the material to be joined later by fusion welding processes.

3. For this study, the susceptibility can be relatively ranked using the spot-varestraint data as Hastelloy X

140 4. ThermoCalc® can be utilized to determine susceptibility of an alloy to HAZ liquation cracking based on temperature difference between initial melting and fully liquid state.

5. Maximum crack length in the HAZ decreases with decreasing grain size.

6. Alloys 625 and 718 decreasing total crack length with decreasing HAZ grain size.

7. Hastelloy X experiences an increase in total crack length as HAZ grain size decreases and is accounted for by the significant increase in crack frequency in the FSP region of the GTAW HAZ.

B.6 REFERENCES

[1] Mishra, R.S. and Mahoney, M., Friction Stir Processing, in Friction Stir Welding

and Processing, R.S. Mishra, M. Mahoney, Editors. 2007 ASM International:

Materials Park, OH USA. Pg. 309-350

[2] Savage, W.F. and Krantz, B.M. 1966. An investigation of hot-cracking in

Hastelloy X, Welding Journal, 45(1): 13s-25s

[3] DuPont, J.N., Lippold, J.C. and Kiser, S.D., Welding Metallurgy and Weldability

of Nickel-Base Alloys. John Wiley & Sons, Inc., 2009

[4] Duvall, D.S. and Owczarki, W.A. 1967. Further heat affected zone studies in heat

resistant nickel alloys, Welding Journal, 46(9): 423s-432s

[5] Owczarki, W.A., Duvall, D.S. and Sullivan, C.P. 1966. A model for heat affected

zone cracking in nickel base superalloys, Welding Journal, 45(4): 145s-155s

[6] Pepe, J.J. and Savage, W.F. 1967. Effects of constitutional liquation in 18Ni

maraging steel weldments, Welding Journal, 49(12): 411s-422s 141 [7] Thompson, E.G. 1969. Hot cracking studies of Alloy 718 weld heat-affected

zones, Welding Journal, 48(2): 70s-79s

[8] Thompson, R.G., Cassimus, J.J., Mayo, D.E., and Dobbs, J.R. 1985. The

relationship between grain size and microfissuring in Alloy 718, Welding Journal,

64(4): 91s-96s

[9] Savage, W.F., Nippes, E.F. and Goodwin, G.W. 1977. Effect of minor elements

on hot-cracking tendencies of Inconel 600, Welding Journal, 56(8): 245s-253s.

[10] Woo,I., Nishimoto, K., Tanaka, K., Shirai, M. 1999. Effect of grain size on heat

affected zone cracking susceptibility. Study of weldability of Inconel 718 cast

alloy (2nd Report). Quarterly Journal of Japan Welding Society, 17(4): 534-542

[11] Thompson, R.G. and Genculu, S. 1983. Microstructural evolution in the HAZ of

Inconel 718 and correlation with the hot ductility test. Welding Journal, 62: 337s-

346s

[12] Baeslack, W.A. III and Nelson, D.E. 1986. Morphology of weld heat-affected

zone liquation in cast Alloy 718. Metallography, 19: 371-379

[13] Baeslack, W.A. III, West, S.L., and Kelly, T.J. 1998. Weld cracking in Ta-

modified cast Inconel 718. Scripta Metallurgica, 22: 729-734

[14] Lin, W., Nelson, T., Lippold, J.C., and Baeslack, W.A. III. 1992. A study of the

HAZ crack susceptible region in Alloy 625. International Trends in Welding

Science and Technology, Gatlinburg, TN U.S.A., June 1-5, 1992: 695-702 142 [15] Homma, H., Nishi, T., and Kohira, K. 1980. Application of the Tigamajig Test

to the Study of Hot-Cracking in Weldments. Journal of the Japan Welding

Society, 49(5): 290-296

[16] Nakuo, Y. and Shinozaki, K. 1990. Hot Cracking Susceptibilities of Nickel-

Base Superalloy Weld Zones. High Temperature Materials for Power

Engineering, Liege, Belgium, September 24-27, 1990: 1461-1470

[17] ASTM Standard E112, 2010, “Standard Test Methods for Determining Average

Grain Size.” ASTM International. West Conshohocken, PA. 2010. DOI:

10.1520/E0122-10

143 APPENDIX C

PHYSICAL SIMULATION OF NICKEL-BASE ALLOY FRICTION STIR

PROCESSING MICROSTURCTURE

144 C.1 ABSTRACT

The modified Gleeble hot torsion test was utilized in an attempt to simulate the friction stir processed microstructure of three Ni-base alloys; Hastelloy X, Alloy 625 and

Alloy 718. The simulation temperatures were based on actual thermal cycles measured by imbedded thermocouples during friction stir processing of these alloys. Peak process temperatures for Hastelloy X and Alloy 625 were determined to be approximately

1150ºC and 1100ºC for Alloy 718. The peak temperature and cooling rates were programmed into the Gleeble 3800 thermomechanical simulator to simulate the stir zone, thermomechanically-affected zone, and heat-affected zone microstructures. The thermomechanically-affected zone was successfully simulated using this technique, but the stir zone microstructure could not be accurately reproduced. Shear stress and strain rates as a function of temperature were determined for each material using the hot torsion simulation.

C.2 INTRODUCTION

Friction stir welding and processing are extremely complex thermo-mechanical processes and there is strong motivation to computationally simulate the microstructure that evolves from these processes. A major setback has been the experimental determination of both thermal and mechanical histories for a given friction stir weld due to the complexities of the process. Recently, a methodology has been developed to physically simulate friction stir microstructures with the ability to extract thermal and mechanical history of the material during testing. This method is called the modified

Gleeble hot torsion test [1,2] and was developed at the Ohio State University.

145 The test involves use of an annular sample with a 1inch (25 mm) gauge length

(Figure C.1), which is resistively heated to a desired temperature and rotated in torsion a particular number of revolutions at a given rotational speed (RPM), and then cooled at a programmed rate [1]. Temperature is controlled and monitored by a thermocouple placed at the center of the gauge section [2,3]. Additionally, thermocouples are placed at the shoulder of the gauge section such that the temperature gradient can be recorded and thermal control can be maintained should the center thermocouple detach during testing

[2,3].

Figure C.1: Modified Gleeble hot torsion test specimen geometry [1].

In addition to simulating the microstructure developed by friction stir processing, elevated temperature mechanical properties can be calculated including, shear stress (τ), shear strain (γ), and shear strain rate . The shear stress is calculated using Equation

146 C.1, where T is the torque, r is the radius of the gauge section, and J is the polar moment of inertia of the gauge section. The polar moment, J, was calculated using Equation C.2 where ro and ri are the outer and inner radius of the gauge section, respectively. The shear strain was calculated using Equation C.3, where r is the radius of the gauge section,

Ѳ is the angular rotation in radians from a reference line parallel to the axis of the specimen, and l is the length of the gauge section. The shear strain rate is calculated using Equation C.4 where RPM is the programmed rotations per minute. This equation is multiplied by 2π and divided by 60 to translate to radians per second.

Equation C.1: Shear Stress [3]

Equation C.2: Polar moment of inertia for gauge section [3]

Equation C.3: Shear Strain [1]

Equation C.4: Shear Strain Rate [1]

Previous studies have been conducted on pure iron, HSLA-65, and Type 310 and

304L stainless steel [1-4] . Recent results by Failla have demonstrated difficulties in reproducing FSW microstructure in Type 310 and 304L stainless steels [3]. Unlike the 147 experience of Norton and Sinfield with HSLA-65, Failla‟s experiments required use of multiple samples at different combinations of temperature and RPM to reproduce each microstructurally distinct region of a FSW [1-3]. It has been demonstrated by Norton,

Sinfield, and Failla that a single hot torsion specimen can reproduce each microstructurally distinct region in HSLA-65 and Armco Iron [1,2]. The present investigation was conducted to evaluate if the modified Gleeble hot torsion test could be used to successfully simulate friction stir processed (FSP) microstructures of three Ni- base alloys Hastelloy X, Alloy 625, and Alloy 718.

The modified Gleeble hot torsion test is based on traditional hot torsion testing which has been used for many years to determine hot workability of materials. Early research by Luton and Sellars utilized the hot torsion test to investigate dynamic recrystallization for pure nickel and nickel-iron alloys [5]. It was demonstrated that the shear stress-strain curves generated during these tests can be used to determine if dynamic recrystallization occurs, Figure C.2. Two signatures are shown, the first of which indicates discontinuous dynamic recrystallization by the periodic nature of the stress levels during steady state deformation. The second curve indicates continuous dynamic recrystallization due to the low variance of stress during steady state deformation. Furthermore, both signatures include a sharp rise to a peak stress at a critical strain before stress drops to the steady state regime during dynamic recrystallization. It is proposed that this peak stress at the critical strain level occurs due to the work hardening behavior of nickel and its alloys.

148

Figure C.2: Proposed shear stress-strain signatures indicating the occurrence of (a) discontinuous dynamic recrystallization and (b) continuous recrystallization in nickel and

Ni-base alloys. (X-axis for (a) and (b) are same scale) From Luton and Sellars.

C.3 EXPERIMENTAL PROCEDURE

Hastelloy X, Alloy 625, and Alloy 718, compositions given in Table C.1, were subjected to FSP with an array of embedded type K thermocouples, Figure C.3.

Thermocouples were embedded in this configuration to survey the thermal history of the entire processed zone, including possible differences associate with the advancing and retreating side. FSP was performed using a GTC AccuStir machine utilizing a W-25Re tool with a truncated cone geometry (Figure C.4) using the process parameters shown in

149 Table C.2. Thermal history was recorded from the thermocouples using an instruNet

Model 100 (Hastelloy X and Alloy 625) and an Omega OMB-DAQ55 (Alloy 718) data acquisition system.

Once representative thermal histories were recorded, simulation was performed using the modified Gleeble hot torsion test specimen geometry, Figure C.1, and the

Gleeble 3800 outfitted with the torsion mobile conversion unit (MCU). Type K thermocouples were placed at the center and shoulder of the gauge section on the outer radius, Figure C.5. Temperature was controlled using the center thermocouple while the thermocouple at the shoulder provided insight into the thermal gradient and history along the sample axis. Cooling rates were controlled using Helium gas as an external and internal quench medium at a flow rate of 80 CFH. Parameters used for the modified

Gleeble hot torsion tests are given in Table C.3.

Table C.1: Alloy compositions.

150

Figure C.3: Thermocouple placement (left) view in direction of travel, (right) view of backside of processed plate embedded with type K thermocouples.

151

Figure C.4: W-25Re tool profile (top), GTC AccuStir machine (bottom).

152 Table C.2: FSP parameters.

Figure C.5: Thermocouple placement for torsion tests.

153 Table C.3: Hot torsion test parameters.

Completed friction stir processed plates were subject to sectioning and metallographic examination. Plates embedded with thermocouples were sectioned across the line of thermocouples to expose the exact thermocouple locations. Additionally, completed hot torsion specimens were sectioned along the axis so that the entire gauge length was in a single mount. All mounts were polished to 1µm diamond followed by a 2 hour final polish using 0.05µm colloidal silica in a vibratory polisher. Upon final polish,

Hastelloy X and Alloy 718 specimens were electrolytic etched using 10% oxalic acid solution at 0.582Amps/cm2. Alloy 625 specimens were electrolytic etched using Lucas‟

154 Reagent at 1.5V and 1Amp for 10-25 seconds. Inspection of etched samples utilized an

Olympus GX51 optical microscope.

C.4 RESULTS AND DISCUSSION

C.4.1 Thermal History Acquisition

Sectioning of plates with embedded thermocouples revealed the actual location of the thermocouples relative to the stir zone. For the Hastelloy X plates, it was found that there was misalignment of the friction stir tool and both the advancing side and stir zone thermocouples ended up in the stir zone, while and the retreating side thermocouple was at a distance of 0.20” and 0.13” from the edge of the stir zone for run numbers 23 and 24, respectively, as illustrated in Figure C.6 and Figure C.7. Temperature data captured for these runs provided consistent thermal profiles and peak stir zone temperatures of

~1150ºC, Figure C.8. Alloy 625 plates had much improved alignment such that all thermocouples ended up in the stir region, Figure C.9. The thermal data again gave consistent thermal histories and peak stir zone temperatures, ~1150ºC, Figure C.10.

Alloy 718 runs experienced similar misalignment to Hastelloy X runs where embedded thermocouples were not incorporated into the stir zone. Additionally, some thermocouples had failed to record any data during friction stir processing. However, stir zone temperatures were able to be captured as at least two thermocouples on separate runs became incorporated into the stir zone. Peak temperatures experienced in the stir zone were consistent between the two runs, ~1100ºC.

155

Figure C.6: Thermocouple locations for run 23 (Hastelloy X).

156

Figure C.7: Thermocouple locations for run 24 (Hastelloy X).

Figure C.8: Acquired thermal history for run 24.

157

Figure C.9: Thermocouple locations in run 40 (Alloy 625).

158

Figure C.10: Acquired thermal history for run 40.

C.4.2 Modified Gleeble Hot Torsion Test

Thermo-mechanical data were acquired for each modified hot torsion test as a function of time including, temperature at the center and shoulder of the gauge section, torque, and rotation. Upon completion of a test, the temperature profile for the gauge center was checked with respect to the acquired history from actual FSP runs. Stress strain curves were generated along with calculation of the strain rate using Equations 1-4.

It was found that Hastelloy X was the most difficult material to test given that all samples ruptured during testing, as shown in the example in Figure C.11. Few samples were able to collect temperature at the gauge center throughout the duration of the test. Most samples lost center thermocouples during rupture. An example of how the programmed temperature correlated with the acquired thermal data from FSP run 24 for Hastelloy X is

159 shown in Figure C.12. It was found that a temperature increase initiates at the onset of torsion and reaches its peak at the end of the torsion event. This temperature change has been consistently recorded at approximately 20ºC. Upon inspection of the sample after testing it was found that the sample ruptured. Investigation of the stress strain curve,

Figure C.13, shows clear evidence of the rupture event due to a precipitous drop in stress, noted at about 0.4 in/in strain. An average stress of 20 ksi was obtained between 0.04 and 0.4 in/in strain. Observed microstructure through the length of the gauge section reveals a reduction in grain size near the center of the gauge length relative to unaffected base metal near the shoulder, Figure C.14. A transition region was not observed in this test sample.

Figure C.11: Example of common rupture experienced with Hastelloy X hot torsion tests

(scale is mm).

160

Figure C.12: Comparison of gauge center temperature and acquired thermal history from

FSP run 24 (Hastelloy X).

Figure C.13: Stress-strain curve for Hastelloy X hot torsion test (hastX3.d09). Note sample rupture occurrence at ~0.4 in/in strain evident by sharp drop in stress.

161

Figure C.14: Optical micrograph of Hastelloy X (A) hot torsion test specimen unaffected base metal, (B) hot torsion test gauge center, and actual FSP stir zone (500X, oxalic etch).

Refinements to the Hastelloy X torsion tests included lowering the programmed peak temperature to account for the temperature increase affect during torsion, changing the rotation rate to provide a lower strain rate and increasing the number of revolutions to accumulate more strain into the system. Again, it was found that the sample ruptured during the torsion event. Furthermore, the center thermocouple disconnected at the onset of torsion providing no temperature information for the cooling phase of the test.

However the thermocouple at the shoulder remained intact and provided some insight in

162 regards to the cooling rate experienced, Figure C.15. Comparison of the temperature recorded by the shoulder thermocouple for specimen „hastx3.d13‟ with recorded shoulder temperature recorded for a sample ran at the same heating and cooling rate shows that

„hastx3.d13‟ cooled at a faster rate than programmed.

Figure C.15: Comparison of thermal data for „hastx3.d13‟ hot torsion test with FSP run

24 (Hastelloy X) and shoulder temperature from a hot torsion specimen heated and cooled at the same rate without a torsion event.

The stress-strain curve generated for this hot torsion specimen shows when rupture occurred, noted by the sharp drop in stress at 0.5 in/in strain, Figure C.16.

Additionally, it is shown that stress spikes to a value of 40 ksi at the onset of torsion and quickly drops to a steady state value of 31 ksi averaged from 0.035 to 0.2 in/in strain. At

0.2 in/in strain a steady drop in stress begins, signifying the initiation of specimen shearing leading to rupture at 0.5 in/in strain. The stress-strain curve obtained for this

163 sample more closely resembles previous studies of DRX in nickel alloys under torsion loading [6-8]. Historically, the spike in stress and subsequent suppression of stress indicates the onset of the DRX mechanism [6,8].

The microstructure observed in the „hastx3.d13‟ specimen (Figure 23) is closer to the actual FSP microstructure. Comparison of FSP stir zone and microstructure encountered near the center of the gauge section of the specimen shows the level of grain refinement is not fully realized but is much closer than previous specimens, Figure C.17.

Grain size for the hot torsion specimen in this zone was found to have an average grain diameter of 17µm, about three times the size of actual FSP stir zone microstructure.

Additionally, a transition region was observed with structure similar to the 164 thermomechanically-affected zone (TMAZ), Figure C.18. This region contained pockets of elongated grains with substructure of smaller grains at the grain boundaries.

Figure C.17: Optical micrograph of Hastelloy X (A) FSP stir zone and (B) hot torsion test stir zone (hastx3.d13) (1000X, oxalic etch).

Figure C.18: Optical micrograph of Hastelloy X (A) FSP TMAZ and (B) hot torsion test

TMAZ (hastx3.d13) (500X, oxalic etch).

Although Alloy 625 specimens more readily survived the hot torsion test, actual

FSP microstructures were still difficult to simulate. Temperatures recorded at the center

165 of the specimen gauge length follow the thermal history of FSP run 40, however there is considerable deviation in the on cooling portion of the test between 35 and 70 seconds at one point giving a difference of nearly 110ºC, Figure C.19. Again, a thermal spike of about 20ºC was encountered at the onset of torsion, Figure C.20. The stress-strain curve shows a peak stress of 33 ksi at the initiation of torsion with an average steady-state stress of 25 ksi between 0.05 and 0.55 in/in strain, Figure C.21. The subtle decrease in stress following the peak stress indicates DRX has not occurred for this test. Microstructures observed for Alloy 625 specimen „625-6.d02‟ did not match the FSP microstructure.

Throughout the gauge length, no appreciable difference in microstructure was observed.

It was found that the shoulder microstructure and the center microstructure grain size and morphology were the same, Figure C.22, confirming that DRX did not occur.

166

Figure C.19: Thermal history comparison of Gleeble hot torsion test and actual FSP run for Alloy 625.

Figure C.20: Expanded view of peak temperature encountered by Alloy 625 during hot torsion testing. Note sharp increase in temperature during the torsion event.

167

Figure C.21: Stress-strain curve for Alloy 625 hot torsion test (625-6.d02).

Figure C.22: Optical micrograph of Alloy 625 hot torsion test (625-6.d02) (A) shoulder and (B) center of gauge section (500X, Lucas‟ reagent).

The test procedure was then modified to include 2 full revolutions at increased strain rate to introduce more strain into the system. Due to the increased strain rate the center thermocouple detached providing only on heating control during the test.

168 Shoulder thermocouple temperature was compared with a control specimen with the same heating and cooling rate without a torsion event. Cooling rates according to the shoulder thermal history are similar to the control specimen up to 55 seconds, Figure

C.23. From 55 seconds to the end of the test, cooling rates were faster than the control specimen. The stress-strain curve shows peak stress of 40 ksi and an average steady-state stress, between 0.055 and 0.77 in/in strain, of 25 ksi, Figure C.24. It is also shown that sample rupture begins at 0.77 in/in strain, evident by the sharp drop in stress. The curve generated for this revised test indicates DRX has occurred, evident from the significant and rapid increase in stress at the onset of torsion followed by a precipitous drop in stress.

169

Figure C.23: Thermal history of Alloy 625 hot torsion test 625-6.d03 with comparison to

FSP run 40 and a control torsion specimen.

Figure C.24: Stress-strain curve generated for Alloy 625 hot torsion test 625-6.d03. 170 Observed microstructure for the revised Alloy 625 test more closely resembled the FSP microstructure, but the grain size of the torsion specimen (16µm) was still about three times that of FSP stir zone grain size (5µm), Figure C.25. Additionally, a transition region was observed with structure similar to the TMAZ, Figure C.26. It is clear that the increased strain and strain rate aided in enhancing the DRX kinetics, more closely simulating microstructures resulting from FSP. This result is in good agreement with previous studies by Zhou and Baker in regards to resultant microstructure due to total applied strain [6].

Figure C.25: Optical micrograph of Alloy 625 (A) FSP stir zone microstructure and (B) hot torsion test 625-6.d03 simulated stir zone (1000X, Lucas' reagent).

171

Figure C.26: Optical micrograph of Alloy 625 (A) FSP TMAZ and (B) hot torsion test

625-6.d03 simulated TMAZ (500X, Lucas' reagent).

Alloy 718 reacted similarly to Alloy 625 during torsion testing. It was found that using a single revolution caused no appreciable changes to microstructure. Accordingly testing was changed to include four complete revolutions keeping the heating and cooling rates the same as well as peak temperature. The center thermocouple detached during testing therefore temperature was referenced by comparison of the shoulder thermocouple and a control specimen shoulder thermocouple, Figure C.27. It was observed from the comparison that the specimen cooled considerably faster than the control sample. The stress-strain curve shows that the sample ruptured during testing at 1.05 in/in strain, evident by the abrupt drop in stress, Figure C.28. A peak stress of 35 ksi was experienced at the onset of torsion with an average steady-state stress of 20 ksi between

0.05 and 1.05 in/in strain. Observed microstructure for the revised Alloy 718 test more closely resembled the actual FSP microstructure with an average grain diameter of12µm relative to the FSP stir zone grain size of4µm, Figure C.29. Additionally, a transition region was observed with structure similar to the TMAZ, Figure C.30.

172

Figure C.27: Thermal history of Alloy 718 hot torsion test 718-2.d08 with comparison to

FSP run 40 and a control torsion specimen.

Figure C.28: Stress-strain curve generated for Alloy 718 hot torsion test 718-2.d08.

173

Figure C.29: Optical micrograph of Alloy 718 (A) FSP stir zone and (B) hot torsion test

718-2.d08 simulated stir zone (1000X, oxalic etch).

Figure C.30: Optical micrograph of Alloy 718 (A) FSP TMAZ and (B) hot torsion test

718-2.d08 simulated TMAZ (500X, oxalic etch).

C.5 CONCLUSIONS

1. Thermal histories were successfully recorded for friction stir processing of

Hastelloy X, Alloy 625, and Alloy 718. Peak temperature for Hastelloy X and Alloy

625 were recorded at 1150ºC and for Alloy 718 at 1100ºC.

2. Modified Gleeble hot torsion testing of the three alloys failed to fully simulate the stir

zone microstructures encountered in FSP material; however the thermomechanically- 174 affected zones were readily reproduced. The number of revolutions and strain rate

are key factors influencing the ability of the modified Gleeble hot torsion test to

simulate friction stir microstructure in Ni-base alloys.

3. Increasing the number of revolutions from 1 to 2 and 4 induced more strain into

specimens providing more nucleation sites for new grains enhancing the kinetics of

the DRX mechanism.

4. Evidence that the DRX mechanism has taken place is indicated by the stress-strain

curve evident by a rapid spike in stress at the onset of torsion followed by a

significant decrease in stress to a steady-state value.

C.6 RECOMMENDATIONS

Further revisions to the modified Gleeble hot torsion test should be made specifically for Ni-base alloys to include more revolutions in order to induce enough strain for DRX to occur. Other low stacking fault energy materials should be subjected to modified Gleeble hot torsion testing to investigate the effect stacking fault energy has on required test parameters for successful simulation of friction stir microstructure. It is recommended that compressive force be added to the test method in order to reduce instances of sample rupture during torsion and to better simulate the compound compression-torsion action a friction stir welding/processing tool imposes on materials

C.7 REFERENCES

[1] S. Norton, Ferrous Friction Stir Weld Physical Simulation, 2006.

[2] M.F. Sinfield, Advancements in Physical Simulation and Thermal Acquisition

Techniques for Ferrous Alloy Friction Stir Welding, 2007.

175 [3] D.M. Failla, Friction Stir Welding and Microstructure Simulation of HSLA-65 and

Austenitic Stainless Steels, 2009.

[4] M.F. Sinfield, "Physical Simulation of Friction Stir Weld Microstructure of High

Strength, Low Alloy Steel (HSLA-65)," in 7th International Symposium on Friction

Stir Welding, Awaji Island, Japan, 2008.

[5] M.J. Luton and C.M. Sellars, "Dynamic recrystallization in nickel and nickel-iron

alloys during high temperature deformation," Acta Metallurgica, vol. 17, pp. 1033-

1043, 1969.

[6] L.X. Zhou and T.N. Baker, "Effects of strain rate and temperature on deformation

behaviour of IN718 during high temperature deformation," Materials Science and

Engineering A, vol. 177, pp. 1-9, 1994.

[7] J.P. Sah, G.J. Richardson, and C.M. Sellars, "Grain-Size Effects during Dynamic

Recrystallization of Nickel," Metal Science, vol. 8, pp. 325-331, 1974.

[8] M. Ueki, S. Horie, and T. Nakamura, "Factors affecting dynamic recrystallization of

metals and alloys," Materials Science and Technology, vol. 3, no. 5, pp. 329-337,

1987.

[9] Dynamic Systems Inc., Gleeble 3500/3800 Operations Manual, 2006.

[10] Dynamic Systems Inc., Gleeble 3500/3800 Options Reference Manual, 2001.

176