1 Impact Welding and Impulse Shape Calibration of Nickel and Titanium

Impact Welding and Impulse Shape Calibration of Nickel and Titanium Alloys

Dissertation

Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy

in the Graduate School of The Ohio State University

Bhuvi Swarna Lalitha Nirudhoddi

Graduate Program in Materials Science and Engineering

The Ohio State University

2019

Dissertation Committee

Dr. Glenn S. Daehn, Advisor

Dr. Xun Liu

Dr. Stephen R. Niezgoda

Copyrighted by

Bhuvi Swarna Lalitha Nirudhoddi

2019

Abstract

High-temperature metallic materials such as nickel-based and titanium alloys are attractive as skin structures for aerospace vehicles. They can allow significant performance improvement and mass reduction in aircraft. However, there are substantial challenges in welding and forming them affordably for service. This project examines the use of impulse-based methods, as enabled by the vaporizing foil actuator method, for the impact welding and precise shaping of alloys Ni - 718, Ni - 625, Ni - 230, and Ti 6242.

The mechanical properties and weld microstructure of four similar and Six dissimilar VFA spot welding combinations are presented and analyzed. Microhardness measurements showed the absence of a heat affected zone (HAZ). The dissimilar Ni - Ni joints and Ni - Ti joints exhibited high loads to failure in lap-shear tests and show great potential for applications involving transition joints, repair welding, medical devices, and more. The VFA method is cheap, safe, fast, durable, and marks the advancement in the solid-state joining of dissimilar nickel and titanium systems.

Nickel alloys typically exhibit low springback during quasi-static forming processes. However, the large amounts of strain hardening that occurs during these operations often requires a second annealing stress relief operation. Titanium alloys are commonly known to exhibit high springback levels due to the high strength to stiffness ratios of titanium alloys. Sheet metals components are usually shaped by hot or

ii superplastic forming. This process is expensive and has long lead-times. This work examines an athermal process to relax or remove the residual stresses and elastic strains in sheet metals. All the materials explored, especially titanium showed significant improvements in shape conformance when processed through the VFA method.

Recent shock-based calibration studies have provided some insight into the previously unconfirmed mechanism of springback relief. The driving hypothesis for this physical phenomenon is that modest shock waves plastically relieve elastic residual stresses and result in target shape conformance. To better understand this mechanism,

VFA based shock processing experiments were used to change the curvature of pre- strained materials to a fully flat shape. It is speculated that the change in shape is a consequence of elastic stress relief caused by the propagation of planar shockwaves. A mechanics and shock physics based theory for shockwave interaction with residual stresses in a pre-strained sample is proposed. The possible pressures involved in this process were calculated from shock breakout velocity profiles captured by a photon doppler velocimetry system. The preliminary pressures are estimated to have satisfied the plastic yield criterion for this loading condition.

iii

Dedication

This dissertation is dedicated to Lord Sri Krishna, my beloved grandparents, parents, sister, brother-in-law, and nephew. I am eternally grateful for their boundless love,

encouragement, and sacrifice.

Sarvam Sri Krishnarpanam Astu

Acknowledgments

I would like to thank my advisor, Dr. Glenn Daehn, for his kindness, guidance, and generosity throughout my graduate career at Ohio State. I am very grateful for his patient explanations of complex phenomena and his undiminishing belief in my capabilities of understanding them. My family and I are forever in debt to his generosity and encouragement.

I am thankful for the mentorship of Dr. Anupam Vivek and Geoff Taber through my time at the Impulse Manufacturing Lab (IML). Their knowledge, expertise, and perspective knows no bounds and has been instrumental in pushing me to think beyond the confines of the known.

I would also like to thank my committee members Dr. Stephen Niezgoda and Dr.

Xun Liu for the support, patience, and expertise offered through this process.

I am very grateful to my past and present IML colleagues, particularly Brian

Ufferman, Angshuman Kapil, and Yu Mao. They have been crucial to my journey of navigating the world of materials science. I would also like to acknowledge Dr. Steve

Hansen for his detailed analysis and meticulous documentation of capturing shockwave breakouts. His efforts provided the groundwork for moving the technique forward. The microscopy and characterization analysis presented in this work would not have been

v possible without the help of Jianxiong Li, Taylor Dittrich, Claire Cary, Keely Shorter,

Chris Wagner, and Wayne Papageorge.

My dreams of pursuing a higher education would not have been possible without the compassion of Dr. Tom Shih from Purdue University, who told me to always look for the opportunity in defeat.

In these past ten years, I have been fortunate enough to find great friendship, support, and community, no matter where I lived in the US. I am especially grateful to

Dhruv Garg, Vindhya Tumati, Kevin Vuong, Marlina Triesjayanti, Jing-Wei Lee, my fellow WEGC sisters, and the kind devotees of ISKCON Columbus for their encouragement and kinship throughout my time here.

Finally, I would like to thank my family. It is impossible to put into words the amount of love and gratitude I feel towards them. No matter how many oceans and continents have parted us, I have always felt their love and support. It has been a blessing and a privilege to have been born to such humble parents and to have grown up with such an amazing sister. Thank you for believing in the little girl who dreamt of voyaging the cosmos.

Vita

2013 B.Sc., Aeronautical and Purdue University

Astronautical Engineering

2018 M.Sc., Materials Science and The Ohio State University

Engineering

2015 - present Graduate Research Associate Department of Materials Science

and Engineering, The Ohio State

University

Publications

B. P. Thurston, A. Vivek, B. S. L. Nirudhoddi, and G. S. Daehn, “Vaporizing foil actuator welding,” MRS Bulletin, vol. 44, no. 8, pp. 637–642, 2019.

P. Stechmann, David & Lim, Dasheng & Rotella, Saverio & Menon, Shankar & Nirudhoddi, Bhuvi. (2014). Design and Analysis of a High-Performance Hydrogen Peroxide Thrust Chamber Assembly. 50th AIAA/ASME/SAE/ASEE Joint Propulsion Conference 2014. 10.2514/6.2014-3500.

Fields of Study

Major Field: Materials Science and Engineering vii

Table of Contents

Abstract ...... ii Dedication ...... iv Acknowledgments...... v Vita ...... vii List of Tables ...... x List of Figures ...... xi Chapter 1. Introduction ...... 1 Chapter 2. Impact Welding of Nickel and Titanium Alloys ...... 5 2.1 Background and Motivation ...... 5 2.2 Literature Review...... 7 Similar Welding ...... 7 Dissimilar Welding ...... 17 2.3 Materials and Methods ...... 22 Materials ...... 22 Experimental Methods ...... 24 Mechanical Testing and Characterization ...... 31 2.4 Results and Discussion ...... 37 Similar Patch Welds ...... 39 Dissimilar Spot Welds: Nickel – Nickel ...... 53 Dissimilar Spot Welds: Nickel – Titanium ...... 62 2.5 Summary and Conclusions ...... 77 Chapter 3. Impulse Based Shape Calibration using the Vaporizing Foil Actuator Method ...... 83 3.1 Background and Motivation ...... 83 3.2 Materials and Methods ...... 84 Materials ...... 84

viii

Experimental Methods ...... 85 3.3 Results and Discussion ...... 89 3.4 Summary and Conclusions ...... 95 Chapter 4. Athermal Springback Relief: Investigation of Governing Mechanism ...... 97 4.1 Background and Motivation ...... 97 4.2 Literature Review of Impulse Based Athermal Springback Removal Methods ..... 98 4.3 Materials and Methods ...... 106 Materials ...... 106 Experimental Method...... 107 4.4 Results and Discussion ...... 117 4.5 Summary and Conclusions ...... 130 Chapter 5. Future Work ...... 133 Impact Welding ...... 133 Impulse Calibration ...... 135 Bibliography ...... 137 Appendix A. SEM Images of Dissimilar Ni - Ni and Ni - Ti Weld Interfaces ...... 146 Appendix B. PDV Flyer Velocity Plots ...... 155 Appendix C. Derivation of Uniaxial Strain Condition for Plastic Deformation...... 156 Appendix D. Velocimetry Data for Flattening Experiments ...... 158

List of Tables

Table 1. Similar and dissimilar VFA welding combinations of interest ...... 7

Table 2. (a) Mechanical properties and (b) Elemental weight compositions (%) of Nickel and Titanium alloys of interest ...... 23

Table 3. VFA Welding parameters for similar combinations ...... 30

Table 4. VFA Welding parameters for dissimilar Ni - Ni alloy combinations ...... 30

Table 5. VFA Welding parameters for dissimilar Ni - Ti combinations ...... 31

Table 6. Characterization methods used for assessing VFA welds ...... 36

Table 7. List of secondary objectives for joint mechanical and characterization tests, desired outcomes, and significance ...... 38

Table 8. Summary table of failure loads and modes for similar patch welds ...... 78

Table 9. Summary table of lap-shear failure loads and modes for similar spot welds ..... 79

Table 10. Summary table of failure loads and modes for dissimilar Ni - Ni spot welds .. 80

Table 11. Summary table of failure loads and modes for dissimilar Ni - Ti spot welds .. 81

Table 12. Thicknesses and strengths of VFA calibration metals ...... 85

Table 13. PDV + VFA flattening experimentation schedule ...... 110

Table 14. Summary of flatness angles, corresponding shock breakout velocities (SBV), and shock breakout pressures (SBP) for 2mm thick AA7075-T6 shock processed in the CU and CD configurations...... 127

List of Figures

Figure 1. Binary phase diagrams for major alloying elements of Ni - 718, Ni - 625, and Ni - 230 system: (a) Ni – Cr [30] ,and (b) Ni - Fe[31], (c) Ni – Nb [32], and (d) Cr – Fe [32] ...... 8

Figure 2. Additional binary phase diagrams for major alloying elements of Ni - 230 (a) Ni -W [34], and (b) Cr – W [34] ...... 10

Figure 3. Ti – Al binary phase diagrams for major alloying elements of Ti 6242 [57] .... 15

Figure 4. Binary phase diagrams for major alloying elements for dissimilar Ni – Ti joints. [32] ...... 18

Figure 5. Patch weld shape in a Ti 6242- Ti 6242 VFA lap weld sample ...... 24

Figure 6. (a) VFA Patch welding setup, (b) 0.08 mm (0.003") thick CP Aluminum Patch Foil with an active area of 76.2 x 25.4 mm (3" x 1"), and (c) top view of VFA mechanical testing specimen ...... 26

Figure 7. (a) VFA Spot Welding setup, (b) 0.08 mm (0.003") thick CP Aluminum Spot Foil with an active area of 10.16 x 12.7 mm (0.4" x 0.5"), and (c) typical spot lap welding sample ...... 27

Figure 8. (a) Deep spot die used for dissimilar Nickel – Nickel Spot Welds listed in Table 4, and (b) Shallow spot die used for Titanium – Nickel Spot Welds listed in Table 5 ..... 28

Figure 9. (a) Side view of patch weld lap-shear test [70], (b) Side view of patch weld peel test [70], (c) Front view of spot weld lap-shear sample and cross-section of weld, (d) Side view of modified coach peel configuration for Ni – Ti spot welds (f) coach or T- peel for Ni – Ni spot welds...... 33

Figure 10. Depiction of typical regions of measurement from Energy Dispersive X-Ray Spectroscopy (green) and Microhardness indents (dashed orange) across an impact weld interface...... 34

Figure 11. Typical cross-section of a VFA weld [12] ...... 39

Figure 12. Mechanical properties and characterization of Ni - 718 VFA patch weld (a) lap shear loads normalized by specimen width and representative failure mode, (b1-b2) peel loads normalized by sample width and categorized by failure modes, (c1-c3) light microscopy stitches of etched weld interface, and (c4) an inverted SEM image of the weld interface ...... 42

Figure 13. Mechanical properties and characterization of Ni - 625 VFA patch weld (a) lap shear loads normalized by specimen width and representative failure mode, (b1-b3) peel loads normalized by sample width and categorized by failure modes, and (c1) light microscopy stitch of etched weld interface, and (c2-c3) SEM and inverted SEM image of the weld interface ...... 46

Figure 14. Mechanical properties of Ni - 230 VFA patch weld (a) lap shear loads normalized by specimen width and representative failure mode, (b) peel loads normalized by sample width and categorized by failure modes, and (c) microhardness (HV) traverses across 230 - 230, 625 - 625, and 718 - 718 patch welds ...... 49

Figure 15. (a) Ti 6242 patch weld lap shear loads normalized by specimen width and representative failure mode, (b) microhardness (HV) traverses across and along Ti 6242- Ti 6242 weld interface, (c) Ti 6242- Ti 6242 spot weld lap shear loads, and (d) light microscopy stitch of etched weld interface ...... 51

Figure 16. Mechanical properties and characterization of Ni - 718 - 625 VFA spot weld (a) lap shear loads w.r.t parent strength and failure modes, (b) T - peel loads and failure modes, (c1 – c3) SEM + BSED images of weld interface, and (d) EDS: average composition (wt. %) of major alloying elements across weld interface taken at the different regions in Figure 10 ...... 55

Figure 17. Mechanical properties and characterization of Ni - 718 - 230 VFA spot weld (a) lap shear loads w.r.t parent strength and failure modes, (b) peel loads and failure modes, (c1 – c3) SEM + BSED images of weld interface, and (d) EDS: average composition (wt. %) of major alloying elements across weld interface taken at the different regions in Figure 10 ...... 57

Figure 18. Mechanical properties and characterization of Ni - 625 - 230 VFA spot weld (a) lap shear loads w.r.t parent strength and failure modes, (b) peel loads and failure modes, (c1 – c3) SEM + BSED images of weld interface, and (d) EDS: average composition (wt. %) of major alloying elements across weld interface taken at the different regions in Figure 10 ...... 60

Figure 19. Flyer velocity vs. distance plots for Ni - 718 and Ti 6242 in the spot-welding configuration at a foil vaporizing energy of 8 kJ ...... 63

Figure 20. Mechanical properties and characterization of 718 – Ti 6242 VFA spot weld (a) lap shear loads w.r.t parent strength and failure modes, (b) peel loads and failure xii modes, (c1 – c3) SEM + BSED images of weld interface for 718 - Ti (flyer) joints, (d) EDS: average composition (wt. %) of major alloying elements across weld interface taken at the different regions in Figure 10 for 718 – Ti (flyer) combinations, and (e1-e3) LM images of weld interface for Ti (target) – 718 joints ...... 65

Figure 21. Fractography and EDS scans of the 718 – Ti 6242 peel surfaces, (a) Image of the 718 (target) peel surface, (b) EDS scan of same surface, (c) Image of the Ti (target) peel surface, (d) EDS scan of same surface, (e) mode of failure for the 718 – Ti (flyer) sample, (f) mode of failure for the Ti (target) - 718 sample, (g) Image of the Ti (flyer) peel surface, (h) EDS scan of same surface, (i) Image of the 718 (flyer) peel surface .... 68

Figure 22. Mechanical properties and characterization of 625 – Ti 6242 VFA spot weld (a) lap shear loads w.r.t parent strength and failure modes, (b) peel loads and failure modes, (c1 – c3) SEM + BSED images of weld interface for 625 - Ti (flyer) joints, (d) EDS: average composition (wt. %) of major alloying elements across weld interface taken at the different regions in Figure 10 for 625 – Ti (flyer) combinations, and (e1-e3) SEM + BSED images of weld interface for Ti (target) – 625 joints...... 70

Figure 23. Fractography and EDS scans of the 625 – Ti 6242 peel surfaces, (a) Image of the 625 (target) peel surface, (b) EDS scan of same surface, (c) Image of the Ti (target) peel surface, (d) EDS scan of same surface, (e) mode of failure for the 625 – Ti (flyer) sample, (f) mode of failure for the Ti (target) - 625 sample, (g) Image of the Ti (flyer) peel surface, (h) EDS scan of same surface, (i) Image of the 625 (flyer) peel surface, and (j) EDS scan of same surface ...... 72

Figure 24. Mechanical properties and characterization of 230 – Ti 6242 VFA spot weld (a) lap shear loads w.r.t parent strength and failure modes, (b) peel loads and failure modes, (c1 – c2) SEM + BSED images of weld interface for 230 - Ti (flyer) joints, (d) EDS: average composition (wt. %) of major alloying elements across weld interface taken at the different regions in Figure 10 for 230 – Ti (flyer) combinations, and (e1-e3) SEM + BSED images of weld interface for Ti (target) – 230 joints...... 74

Figure 25. Fractography and EDS scans of the 230 – Ti 6242 peel surfaces, (a) Image of the 230 (target) peel surface, (b) EDS scan of same surface, (c) Image of the Ti (target) peel surface, (d) EDS scan of same surface, (e) mode of failure for the 230 – Ti (flyer) sample, (f) mode of failure for the Ti (target) - 230 sample, (g) Image of the Ti (flyer) peel surface, (h) EDS scan of same surface, (i) Image of the 230 (flyer) peel surface, and (j) EDS scan of same surface ...... 76

Figure 26. (a) Shape and dimensions of the flange die, and (b) springback in CP Titanium compared to the target punch shape ...... 86

Figure 27. (a) Insulated polyurethane punch, and (b) top view of patch foil and shaping punch alignment in the VFA calibration step ...... 87

xiii

Figure 28. Detailed VFA calibration assembly...... 87

Figure 29. Regions of measurement for a pre-formed and calibrated workpiece ...... 89

Figure 30. Comparison of average springback % in the curved and flat sections of the workpieces in the pre-forming stage (pre) vs. VFA calibration stage (post) ...... 91

Figure 31. Double-patch foil configuration used by Iriondo [82] to calibrate AA6061-T6 and CP Ti ...... 92

Figure 32. Microhardness traverses across flanged and curved section of pre-formed and VFA calibrated samples of (a) AA6061-T6, and (b) CP Ti ...... 94

Figure 33. (a) Specimens bent around a mandrel and then flattened with a steel plate (b) Specimens flatted using EM pulses of different voltages (Borrowed from Golovashchenko [83]) ...... 101

Figure 34 (a) Experimental setup of EH flattening test (b) Comparison of a strip of DP980, 1 mm thick, bent over a 1 in. radius and then flattened by a 1 MN clamping press, to a strip then calibrated by one EH pulse at 6.3 kV (borrowed from Golovashchenko et al., [84]) ...... 103

Figure 35. Curved to flat stress relief demonstrations conducted on CP Titanium, High Strength Boron Steel, and Inconel 625 using the VFA process ...... 105

Figure 36. Variation of shock pressure with particle velocity for AA7075 [99] ...... 107

Figure 37. (a) Sheet-metal air-bending operation and (b) method of measuring final springback angle in pre-bent specimens ...... 108

Figure 38. (a) Stacking and alignment depiction of VFA flattening experiments, (b) concave up (CU) stacking configuration, (c) concave down (CD) configuration, and (d) overall PDV + VFA fixturing setup for shock calibration experiments ...... 109

Figure 39. (a) VFA Calibration + PIPE setup for pressure estimation and (b) top view of AA6061 transducer sample after being indented with a file due to an 8kJ foil vaporization energy ...... 111

Figure 40. Illustrated principle of Photon Doppler Velocimetry (PDV) [101]...... 113

Figure 41. (a) Raw intensity data of the shock processed free surface, captured via PDV, (b) FFT processed intensite data that depicts the profile of a shock breakout and the likely peak pressure ...... 115

xiv

Figure 42. Adapated general shock-profile obtained using a Velocity Interferometer System for Any Reflector (VISAR) by TeiresiasSky/ CC BY-SA 4.0 (Desaturated from original) ...... 116

Figure 43. Influence of foil vaporization energy on the reduction of springback angle for AA7075 – T6 samples, shock loaded in either the concave up (CU) and concave down (CD) geometries: (a) 1 mm thick CU, (b) 1 mm thick CD, (c) 2 mm thick CU, (d) 2 mm thick CD, and (e) chart showing overall changes in flatness angle with energy ...... 119

Figure 44. Plane stress state in the cross-section of (a) blank after air-bending, (b) pre- bent workpiece constrained in the concave up geometry, and (c) pre-bent workpiece constrained in the concave down geometry. Also shown is the stress state of the shockwave itself ...... 122

Figure 45. Variation of Poission’s ratio with uniaxial strain yield factor ...... 126

Figure 46. Spatial pressure distributions for AA6061-T6 VFA shock processed through a urethane medium at 4.5 kJ, 6 kJ, and 8 kJ...... 129

Figure 47. SEM images of 718 (top) – 625 (bottom) weld interface taken at the regions shown by the red squares ...... 146

Figure 48. SEM images of 718 (top) – 230 (bottom) weld interface taken at the region shown by the red square ...... 147

Figure 49. SEM images of 625 (top) – 230 (bottom) weld interface taken at the region shown by the red square ...... 148

Figure 50. SEM images of 718 - Ti (flyer) weld interface taken at the region shown by the red square ...... 149

Figure 51. (a-c) SEM images of 625 - Ti (flyer) weld interface taken at the region shown by the red square, and (d) optical stitch of 625 – Ti weld showing region where debonding occurred (close to normal impact angle region) ...... 150

Figure 52. SEM images of 230 - Ti (flyer) weld interface taken at the regions shown by the red squares...... 151

Figure 53. Optical microscopy images of 718 - Ti (target) weld interface ...... 152

Figure 54. SEM images of 625 - Ti (target) weld interface ...... 153

Figure 55. SEM images of 230 - Ti (target) weld interface ...... 154

Figure 56. Flyer velocity profiles obtained using the PDV method for (a) Ni - 718, (b) Ni - 625, (c) Ni - 230, and (d) Ti - 6242 ...... 155 xv

Figure 57. Depiction of planar shock propagation through the thickness of the material ...... 156

Figure 58. Shock breakout profiles, corresponding velocities and pressures for 2 mm samples subjected to shock processing in the concave up (CU) geometry at different energy levels ...... 158

Figure 59. Shock breakout profiles, corresponding velocities and pressures for 2 mm samples subjected to shock processing in the concave down (CD) geometry at different energy levels ...... 159

Figure 60. Velocity profile for 1 mm sample subjected to shock processing in the concave up (CU) geometry at 8.5 kJ (no breakout) ...... 160

xvi

Chapter 1. Introduction

High-pressure shock-based manufacturing processes such as impact welding and impulse forming are of great importance to the manufacturing community due to the significant benefits that they offer over their conventional counterparts [1], [2].

Impact welding involves the high-speed collision of two materials resulting in a true metallurgical bond. This process can produce robust solid-state welds with strengths exceeding base metal values, involving little to no heat affected zone (HAZ) [3], [4].

Impulse forming, commonly known as high-speed or high velocity forming, involves the high strain rate deformation of materials into a target shape. This method has been used to combat the problem of springback in high strength materials and to create precision-shaped components [3], [5].

Traditional impulse methods rely on explosives, magnetic actuators, and pulsed lasers to generate the high-pressure pulses that are used for metalworking operations.

While these methods have been successfully implemented world over for manufacturing components for electronics and communication, oil and gas, chemical, automotive, and aerospace industries, they hold a small share in the global market compared to conventional fusion welding and quasi-static processes [3], [6], [7].

Explosives while cheap have safety and handling issues that limit their small- scale applicability. Magnetic pulse operations are restricted to conductive materials. The

1 coils used for magnetic pulse welding (MPW) and electromagnetic forming (EMF) are not suitable for repeated use, especially at high pressures. Both laser impact welding

(LIW) and laser shock forming (LSF) are capable of creating enormous pressures over small regions. LIW is restricted to small spot welding geometries and thin gauge workpieces. LSF is bounded by the penetration depth and area of the pulse. It is used as a means of forming thin gauge sheets by introducing compressive stresses to the workpiece surface and not by relieving the springback. These constraints thereby lower the level of precision of this process [1], [2], [4], [7].

A recently introduced high-pressure metalworking technique called the vaporizing foil actuator (VFA) process offers significant advantages over these existing impulse base methods. The shockwave source for the VFA process is a rapidly expanding gas generated by the fast discharge of a large electrical current through a thin metallic conductor. As the resistive Joule heating from the high current quickly heats the actuator above the energy needed for its sublimation, the metal foil is converted into a high- energy gas which expands at high pressure. It is estimated that this method is capable of accessing a modest pressure window between 50 MPa - 5 GPa [8]–[12].

The VFA method is an agile and economical process that has been implemented in creating sound solid-state welds between a variety of structural alloys in both sheet and plate form. Its capabilities as a precision high-speed forming method have also been demonstrated with CP Titanium, Aluminum, and Steel alloys. A majority of the focus of previous research has been with alloys of interest to the automotive industry [9], [13]–

[17].

The work presented in this dissertation seeks to expand the existing knowledge base of the VFA method to industrial applications that rely on popular high-temperature

Nickel (718, 625, and 230) and Titanium 6242 alloys. These materials boast a variety of excellent mechanical properties. They were specifically designed to exhibit high creep, corrosion, and oxidation resistance at elevated temperatures while maintaining their high strengths.

These alloys were invented around the 1970s-1980s and therefore their weldability through fusion welding, and shape conformance through quasi-static or high temp forming methods have been researched and modified extensively. Despite these developments, the use of these materials is due to the high manufacturing costs associated with the similar and dissimilar joining and forming of components. The cost is a direct consequence of the complexities involved with customizing existing methodologies and equipment. Currently, the fabrication of parts made with these materials can only be accomplished by expert craftsmen or through expensive equipment [18]–[21].

As mentioned previously, the VFA process is a modular and inexpensive method capable of producing low-volume parts made of high strength nickel and titanium alloys.

The two themes of this research are impact welding (Chapter 2) and athermal springback calibration (Chapter 3 and Chapter 4) of Ti 6242, Ni 718, Ni 625, and Ni 230 alloys.

The VFA welding process was used to successfully weld 10 similar and dissimilar combinations of these nickel and titanium alloys. Chapter 2 presents the results of a high-level investigation into the room temperature mechanical properties of these welds.

The interface of the welds was analyzed using optical and scanning electron microscopy

3 methods. Out of the 10 combinations explored, the sturdy dissimilar Nickel-Nickel and

Nickel-Titanium joints are of great interest to the industry.

Chapter 3 presents a short study into the springback relief capabilities of the VFA impulse process. All four alloys were successfully calibrated to the target flange shape and showed minimal deviation in conformance. The compliance to the target shape was most drastic for Ti 6242, which is a near alpha titanium alloy, notorious for exhibiting high levels of springback during room temperature forming processes. It is hypothesized that the shape calibration is a result of an athermal residual stress relief mechanism possibly caused by the propagation of a high-pressure shockwave supplied through the foil vaporizing process.

Chapter 4 further explores this springback relief mechanism through a curved to flat shape calibration conducted with AA7075-T6. Similar to the method of stretching bent parts to make the through-thickness elastic strain similar, it is postulated that if a consistent in-plane plastic strain can be driven through the material, it will take the shape that it is currently held to. The work seeks to estimate the sustained pressures during the

VFA based shock processing step and compare it against the yielding criterion for the shock loading condition.

Finally, Chapter 5 provides an overall summary of the main findings of these experiments and presents an outlook on future directions for VFA based fabrication of

Nickel and Titanium alloys.

Chapter 2. Impact Welding of Nickel and Titanium Alloys

2.1 Background and Motivation

Aerospace vehicles typically consist of complex substructures that are expensive to fabricate and difficult to integrate into the primary assembly. Weight savings and response to rapid temperature change are crucial factors involved in material selection for the design of structural and engine components. Titanium alloys are preferred for aircraft components such as compressors, blades, discs, hot airframe skins, that require high specific strengths, corrosion and creep resistance at intermediate temperatures [22], [23].

Regions of aerospace vehicles that experience even higher service temperatures such as the turbine blades of a jet engine are primarily made of Nickel-based superalloys which offer high strengths, corrosion resistance, creep strength, and oxidation resistance [18],

[24].

Despite the high demand for titanium and nickel superalloys in high temperature and corrosive environments, fabrication and manufacturing of these materials are expensive due to the efforts of maintaining high properties after welding and shaping operations. Traditional fusion welding methods, particularly those used for dissimilar joining, can be especially complicated due to the special precautions that need to be taken for these metals. Newer technologies such as fiber laser welding and friction stir welding have shown some promise in joining these aerospace alloys but are expensive to purchase, and setup and often accompany problems in the form of weld cracking [25],

[26].

A new approach called Vaporizing Foil Actuator (VFA) welding has demonstrated much potential in the inexpensive welding of high strength materials. The

VFA process is a collision welding method that involves the oblique impact of a metal workpiece, referred to as the flyer, against a stationary. The flyer gets its acceleration from the rapid expansion of gasses produced during the electrically induced vaporization of a thin aluminum foil. This method is unique because it produces robust joints [1], [12] between sheet metal structures (both similar and dissimilar) and involves little to no heat affected zone. The solid-state nature of this high-speed welding process can maintain the original properties of the materials without the need for post weld heat treatments

(PWHT) to restore properties or remove residual stresses. Advances in this technology could lead to the potential reduction in material costs and allow for more complex or flexible designs in the transportation and medical industries.

Thus far, this method has been optimized and studied for the welding of automotive aluminum alloys and steels [9], [10], [15]–[17], [27], [28]. Recent interest from aerospace companies has motivated the use of VFA welding for studying the weldability of nickel alloys 718, 625, 230, and titanium 6242. Following a short literature review on properties and existing methodologies for similar and dissimilar joining of these materials, this chapter provides insight into the recent VFA ventures in creating high strength welds of the combinations listed in the table below.

Similar Dissimilar 718 – 718 Nickel- Nickel Titanium – Nickel 625 – 625 625 – 718 Ti – 625

230 – 230 230 – 718 Ti – 718

Ti – Ti 230 – 625 Ti – 230 Table 1. Similar and dissimilar VFA welding combinations of interest

2.2 Literature Review

Similar Welding

Nickel Superalloys

General Description

Alloy 718, commonly known as Inconel 718 or simply alloy 718, is a precipitation strengthened nickel-based superalloy developed in the 1960s. It is a Ni-Cr-

Fe alloy (relevant binary phase diagrams of elements above 5 wt. % in Figure 1) capable of maintaining high strength and corrosion resistance up at intermediate temperatures to

700°C (1300°F). The 718 (γ - FCC) Nickel matrix is precipitation strengthened by an ordered body-centered tetragonal phase of γ′′ (Ni3Nb) and γ′ (L12) intermetallic phase of

Ni3(Al, Ti). The sluggish precipitation, aging response, and transformation kinetics of the

γ′′ phase promote high weldability and stability of the 718 alloy. It stands out from the many varieties of superalloys due during the post-weld heat treatment. Due to these characteristics, Ni -718 has found many applications in nuclear reactors, chemical and

7 process industries, aerospace turbine engines, high-speed airframe components, wheels, high-temperature bolts and fasteners, and more [29].

Figure 1. Binary phase diagrams for major alloying elements of Ni - 718, Ni - 625, and Ni - 230 system: (a) Ni – Cr [30] ,and (b) Ni - Fe[31], (c) Ni – Nb [32], and (d) Cr – Fe [32]

Alloy 625 is a solid solution strengthened nickel-based superalloy that was developed in the 1950s to address the demands of high-strength steam-line piping materials. Like many other nickel-based superalloys, it maintains high strengths and resistance to corrosion, oxidation, and nitridation at elevated temperatures. Since its original introduction in 1962, it has seen a variety of applications in high temperature, cryogenic, and highly corrosive environments. It can be found in aircraft engine bleed ducts and bellows, airframe components, nuclear reactors, naval components, and more

[19].

These properties and applications are a result of alloying additions such as Cr,

Mo, and Nb into the face-centered cubic (FCC) nickel matrix. In addition to aiding in corrosion resistance through passivation, chromium also forms Cr23C6 carbides that promote creep resistance by nucleating at the grain boundaries around 815 oC to 980oC

(1500oF to 1800oF). Molybdenum acts as matrix stiffener through solid solution strengthening, promotes resistance to pitting corrosion, and provides creep resistance by forming M6C and MC carbides. Nb forms Ni3Nb which exists in both the γ″ and δ phase.

γ″- Ni3Nb a body-centered tetragonal phase which strengthens the matrix through coherency strain. δ- Ni3Nb, which typically forms due to long exposures to elevated temperatures, is incoherent with the matrix and might cause embrittlement. However, it can also contribute to hardening as a dispersant [19], [33].

Figure 2. Additional binary phase diagrams for major alloying elements of Ni - 230 (a) Ni -W [34], and (b) Cr – W [34]

The nickel 230 superalloy was developed in the 1970s by Dwaine L. Klarstrom to provide a cobalt-free material capable of maintaining high strengths and corrosion resistance in a high temperature environment. The alloy is majorly composed of Ni, Cr, and W by weight (relevant binary phase diagrams of elements above 5 wt.% in Figure 1a and Figure 2 ). The gamma phase matrix has an austenitic (FCC) structure that is solid solution strengthened by carbides that form due to the additions of tungsten and chromium. Primary tungsten carbides promote creep resistance and improve fatigue life and secondary chromium carbides (Cr23C6) improve both corrosion and creep resistance

[29]. Due to these mechanical properties, the 230 alloy is often found in nuclear reactors, rocket nozzles, gas turbine combustors and injectors, high temperature heat exchangers, and furnaces. The nature of these applications requires long lasting joints between and

10 within components, and hence, there is a need for strong welds capable of enduring these environments [35].

Welding Practices

Weldability was prioritized during the developmental stages of these superalloys

[33], [35], [36]. Given the demanding nature of their applications, an examination of any welds created with these alloys would need to exhibit minimal defects created due to the joining process

The most popular industry practice for joining nickel alloys 718,625, and 230 is

Gas Tungsten Arc Welding or GTAW [18]. GTAW, also known as Tungsten Inert Gas

(TIG) welding, is a fusion welding method that uses a consumable tungsten electrode.

During the process, the intense heat from the arc melts and fuses metals together. GTAW requires the use of inert shielding gases such as argon and helium to protect the electrode and the weld area from oxidation and contamination. Filler metals may or may not be used depending on the weldability of the desired metal combination.

Though this process has been applied extensively for joining all three alloys, it is associated with many welding defects that need particular precaution for rectification and prevention. The exposure to high fusion welding temperatures can cause the segregation of solutes in the HAZ leading to detrimental laves phase formation which deteriorate the mechanical properties of the base metal. The high heat input also increases the susceptibility to liquation cracking in the HAZ, solidification cracking in the weld metal during the final stages of solidification, and distortion due to induced residual stresses. A

11 significant amount of work has focused on removing weld cracking, distortion from fusion welding, and optimizing the fusion welding parameters and PWHTs to avoid or minimize these defects [37] [38]–[40]. As with most fusion welding methods, there also remains a concern for extensive cleanliness of the welding surface [41].

Recent progress in process control and technologies of other fusion welding methods such as Gas Metal Arc Welding (GMAW), Electron Beam Welding (EBW), and more recently Laser Beam Welding (LBW), has shown potential for minimizing fusion welding defects in nickel superalloys. Cornu et al., [42] optimized the parameters for joining 2 mm (0.08") thick 718 sheets in the solution treated, aged, and non PWHT states using an Nd: YAG laser. Lertora et al., [43] lap welded two different thicknesses of IN

718 sheets together using a CO2 laser for fatigue testing a prototype helicopter engine component.

Recent interest in laser welding of IN 625 has motivated the butt welding of 0.8 mm (0.03”) thick sheets using a High Power Fiber Laser (HFPL) [44]. Caiazzo et al., [45] were the first to use a CO2 laser to create an edge joint between 0.7 mm (0.027”) thick sheets. The welds while free of cracks still displayed porosity due to higher heat inputs.

Ernst [46] joined Ni - 230 blanks between 6mm - 19mm thick (0.24" - 0.75") using two different Ni-Cr-W-B based filler materials using GMAW and GTAW [39], [43]–[45].

Regardless, these welding issues remain persistent with all high temperature joining methods [19], [37], [41], [42], [47], [48]. Additionally, though laser welding offers many advantages over fusion welding, it remains an expensive process that requires a constant gas shielding method to prevent oxidation during welding.

Compared to fusion welding, a fractional amount of research has been conducted in solid-state welding methods of the nickel alloys. Mary and Jahazi [49] studied the

Linear Friction Welding (LFW) and processing conditions for joining IN 718 blocks.

Song and Nakata [50] and Damodaram et al., [51] both found that the mechanical properties of friction welded 718 rods improved after solution treating and ageing of the welds. Some of the recent work conducted through collaborative studies between Ding and Schneider [52], Williston [53], and Schneider et al.,[54], offers some insight into the parameters for the solid state joining of Ni - 230 through modified Friction Stir Welding

(FSW). Further information on advances of friction welding of nickel superalloys can be obtained through a review published by Chamanfar et al., [55].

Friction welding offers many advantages over fusion welding such as lower interfacial temperatures (below the melting point) which reduces the occurrence of solidification and liquation cracks, voids, pores, detrimental laves phases, HAZ, need for surface prep, flux, protective shielding gases, etc. However, joints made with Linear

Friction Welding and Inertia Friction Welding need to be heat treated to reduce residual stresses from the welding process and recover the hardness of the welded regions.

Additionally, post-weld machining of the flash associated with FSW processes is necessary [55].

One additional limitation factor in the joining of nickel alloys is cost. Despite the routine application of GTA to join sheet metal and forged plates made of nickel alloys, and friction welding of Inconel rods, Henderson et al., [18] and Chamanfar et al., [55] have stated that the technologies remain expensive and, in the case of GTAW, require

13 skilled workmanship. It is evident that a cost-effective solid-state welding method that requires minimum post welding treatments has a lot of potential for improvement joining nickel superalloys and, as will be described in the following sections, titanium alloys.

Titanium 6242

General Description

Ti-6242 (Ti-6Al-2Sn-4Zr-2Mo) is a near α titanium alloy, sometimes classified as an α+β alloy. It was developed in the late 1960s for hot section aeroengine components such as impellers, discs, turbines, etc. It has excellent mechanical strength, stability and creep resistance upto temperatures as high as 538°C (1000°F) [20], [56]. Figure 3 below shows the binary phase diagram for the Ti-Al system. Aluminum is a major alloying element in Ti 6242. It stabilizes the α (hexagonal closed packed - HCP) phase field while molybdenum is a β (body central cubic – BCC) stabilizer. Both tin and zirconium have neutral effects on phase stability, but they typically lower the α/β transformation temperature. Ti 6242 also contains a small percentage of silicon which promotes creep resistance [20], [56].

Figure 3. Ti – Al binary phase diagrams for major alloying elements of Ti 6242 [57]

Welding Practices

Ti 6242 has been classified as a fairly weldable alloy. Though conventional fusion welding methods such as GTAW, GMAW, EBW, and LBW have been employed previously, the ease of welding of Ti 6242 and other titanium alloys is challenging due to the reactive nature of titanium [58].

Titanium in general has a strong chemical affinity for oxygen causing the rapid formation of a stable oxide layer at room temperatures. Though this passivation behavior is responsible for its high corrosion resistance at elevated temperatures, it is a disadvantage for fusion welding processes. At temperatures over 500°C (930°F), the oxidation resistance of titanium decreases leaving it susceptible to embrittlement due to the interstitial dissolution of nitrogen, hydrogen, and oxygen. Therefore, inert

15 environments and shielding gases free of these interstitial elements is one of the key requirements for fusion welding titanium alloys [56], [58].

Brazing is a common solution to the oxygen embrittlement problem [59]. Brazing involves the joining of two or more metals by melting a filler material of a lower melting point into the joint. It has been used since the 1950s for both similar and dissimilar joining of titanium alloys [60]. It is widely used for aerospace, implants, and submarine applications where a high strength-to-weight ratio and corrosion resistance are important.

The selection of the brazing material is a crucial factor as the mechanical properties of the joint can be affected by the temperature and environment of the braze [61]. Though over a hundred brazing filler materials (BFM) have been discovered and widely used in industry, one of the biggest problems encountered is the negative effects of heating the titanium base metal over its β transus temperature. In the case of CP Ti and α-Ti alloys, especially, prolonged heating over 900°C (1650°F) has resulted in grain growth and reduced ductility [59]. Some other risks of brazing titanium alloys include stress concentrations in brazed joints, erosion of base metals, and formation of brittle intermetallics (such as TiCu2, TiAl3, TiNi2, etc., see Figure 3) in the joints. While the industrial recommendations for minimizing the adverse effects attributed to these problems exist, the post processing tends to drive up the time and cost of the brazing process [59], [61].

An essential factor for welding titanium is the cleanliness of the joint. A thorough list of precautions for welding titanium alloys can be found in the ASM technical by

Donachie [58]. The surface needs to be meticulously degreased from contaminants and

16 oils prior to welding. For fusion welding, depending on the thickness of the passivating oxide layer, the sheet metal workpieces need to be cleaned with solvents [58].

Diffusion bonding (DB), often accompanied by superplastic forming (SPF), is a common process used for joining titanium alloys meant for aerospace applications. It is a solid-state process in which the contaminant-free weld surfaces are held in together for long periods of time, under a moderate pressure and elevated temperature. As in the case of fusion welding, DB requires an inert environment to prevent the dissolution of oxygen and is often conducted in a vacuum chamber. Typically, diffusion bonded joints are stronger than fusion welded joints and possess properties equivalent to the base metal.

Unfortunately, DB is an expensive process and its applicability is limited by the stringent surface preparation, special equipment, and inert atmosphere requirements [58], [62],

[63].

Explosive welding is another promising solid-state process that has been previously used for welding Ti 6242, especially in a honeycomb form. While the high impact energy nature of this process diminishes both the inert environment and cleanliness requirements, the usage frequency is low due to safety and scalability issues

[3], [4].

Dissimilar Welding

There has been long standing interest in the dissimilar joining of nickel-based and titanium alloys. The aerospace industry, in particular, has shown much interest in the potential lightweighting and elevated temperature applications of these types of joints

[64], [65].

It is relatively easy to create dissimilar joints between non-age hardenable nickel- based superalloys using conventional fusion welding methods as long as the proper precautions (mentioned in previous sections) and adequate filler materials are used. With proper temperature monitoring and filler material selection, age hardenable alloys can be joined as well [41], [66]. The dissimilar welding of nickel alloys is typically used in repair welding scenarios and has been applied in the industry [41], [66]. While the need to improve joint strengths is not apparent, there is a definite need for cost reduction of these joining practices [18].

Figure 4. Binary phase diagrams for major alloying elements for dissimilar Ni – Ti joints. [32]

In contrast, the need for successful and strong joints between dissimilar nickel and titanium alloys is much greater.

Gorin [67] attempted to join nickel-based alloys to titanium alloys using GTAW but was unsuccessful in creating direct joints due to the brittle intermetallic compounds

(IMCs) that form at the welding temperatures. These IMCs were suppressed when a columbium or copper interlayer was used.

In 1973, Ells et al. [68] published a report comparing electroplating, brazing,

EBW, and EXW for joining a soft nickel alloy (Ni-201) to titanium alloys. EBW proved unsatisfactory for creating strong joints due to the lack of IMC suppression and the success of electroplating varied from one manufacturer to another. Brazing trials were conducted with five filler materials Cu, Al, Fe, Zr, and Zr-based alloys. While Cu displayed the highest impact strength, it was still considerably lower than that obtained by brazed Ti - Ti joints. The strengths of explosive welds between Ni-201 and three titanium alloys (Ti 64, Ti 6246, and IMI 550), despite showing the presence of TiNi and

TiNi3, were deemed satisfactory as long as a subsequent anneal operation is not required.

A major caveat of the explosive welds that were made was that the bonding was done by skilled workers at a commercial establishment, which affects the access to the technology.

Seretsky and Ryba [64] followed this by laser welding to join 1.6 mm (0.063”) thick sheets of pure nickel to pure titanium. Microstructural characterization revealed that about 75% of the weld interface was covered in cracks possibly caused by brittle intermetallics (TiNi3 and Ti2Ni). These welds were deemed unsuccessful and could not be improved by changing the laser power.

Since the publishing of these reports, actual progress in joining nickel to titanium alloys has been scant. Even with the latest technological advancements in the welding industry, only a handful of research has been published in the past three decades.

Min [69] failed to obtain direct joints between titanium rods (Ti 64 and Ti 6242) and Alloy 718 using friction welding. The joint was only possible when an Nb/Cu interlayer was used to act as both an IMC suppressor and a diffusion barrier.

Alemán et al. [65] successfully diffusion bonded Ti 6242 and IN 625 rods in a protective atmosphere at 900 C, 40 MPa pressure, and a 30 minute hold time. The purpose was to study the microstructures of the weld interface using advanced characterization techniques such as EDS, SEM, and STEM. The interface displayed a variety of intermetallic phases such as Ni3Nb, ηTiNi3, Cr (20% Mo), βCr2Ti, NiTi, TiO, TiNi, and

Ti2Ni close to the titanium side due to the faster diffusion rate of nickel into titanium. No report on joint strength or observation of weld defects was made.

Chen et al. [25] used continuous fiber laser welding to successfully join 2mm

(0.08”) thick sheets of Ti 64 to IN 718. While no reports on strength were made, it was suggested that it is possible to control the thickness of the intermetallics by modifying the weld parameters.

The most recent work in the welding of nickel to titanium was conducted by

Topolski et al. [26]. Pure Ti/Ni bimetal plates were explosively bonded, and bend tested in the original and annealed conditions. The non-annealed interface showed a 150 nm thick mixed Ti/Ni interlayer which evolved into three layers of intermetallics of a total 4

µm thickness. As a consequence, the annealed samples had lower bend strengths.

It should be noted that while brazing (and diffusion brazing) has been used as a means for creating dissimilar joints involving nickel and titanium alloys, it has been restricted to γ-TiAl alloys [59], [60] and not for the α/α+β Ti 6242 alloy which is the focus of this work.

In addition to the formation of brittle intermetallics during the welding process, conventional fusion welding methods have limited the applicability of Ni to Ti welding due to the cleanliness and inert chamber precautions. Solid-state joining using FSW and

DB has been limited by accessibility to manufacturers, equipment cost. EXW has demonstrated the highest potential in created successful, high strength joints but is also limited by access to equipment and safety precautions.

When optimized correctly, these welds have the capability of behaving as transition joints in jet engine compressors and gas turbine shafts. Ni - Ti welded materials can also potentially replace heavy nickel-based superalloy components and improve overall engine efficiency by reducing weight [69]. The aircraft industry clearly has a significant need for welded nickel-titanium structures.

The recurrent theme in the literature review for nickel and titanium welding showed that there was a need for easier, faster, and cheaper solid-state welding. In one case such as dissimilar Ni - Ti joints, just the ability to obtain a successful joint can be regarded as a huge improvement. The following sections describe how the VFA welding process can be utilized for joining of then ten welding combinations mentioned in Table

1. After a description of the welding parameters and configurations that yielded a

21 successful joint, the work delves into the interface morphologies and room temperature strengths of these combinations.

2.3 Materials and Methods

Materials

The mechanical properties and elemental weight compositions of the materials used for the VFA welding study are listed in Table 2 (a) and (b). The Titanium and

Nickel-based alloys were obtained from TIMET and Haynes International respectively.

Based on the results from preliminary trial and error experiments, two different configurations of the VFA welding were employed for welding: 1) patch welding for similar combinations, and 2) spot weld for both similar and dissimilar Ni - Ni and Ni - Ti combinations. Further details regarding the process selection are presented in the following sections.

AMS YS (0.2% Material Heat Treatment Modulus Hardness Spec. Offset) Duplex 930 MPa 114 GPa TIMET 6242 4919 336 HV Annealed Solution Heat 416 MPa 200 GPa Ni - 718 5596 260 HV Treated Solution Heat 390 MPa 209 GPa Ni - 230 5878 250 HV Treated Solution Heat 426 MPa 208 GPa Ni - 625 5599 260 HV Treated (a)

(B) Near Alpha Ti (A) Nickel Superalloys Alloy Element Ni - 718 Ni - 625 Ni - 230 Element Ti 6242 Al 0.54 0.27 0.36 Al 6.03 B 0.004 - 0.004 C 0.006 C 0.052 0.022 0.11 Fe 0.03 Co 0.35 0.12 0.21 O 0.1 Cr 18 21.09 21.83 N 0.006 Cu 0.05 - 0.04 Mo 1.92 Fe 18.5 3.83 1.79 Si 0.084 Mn 0.24 0.25 0.49 Sn 2.03 Mo 3.05 8.78 1.24 Ti 87.63 Ni 53.2 70.62 59.53 Y <0.0004 P <0.005 0.008 0.006 Zr 4.19 S <0.002 <0.002 <0.002 Si 0.08 0.14 0.35 Ti 1.11 0.28 <0.01 Nb 5.01 3.381 - Ta <0.05 - - W - - 14.02 La - - 0.019 (b)

Table 2. (a) Mechanical properties and (b) Elemental weight compositions (%) of Nickel and Titanium alloys of interest

Experimental Methods

Patch Welding Setup

This VFA impact welding process gets its name from the race track shaped

“patch” weld shown in Figure 5. This is a result of the dog-bone shaped vaporizing foil

(Figure 6b) used for joining different metals. It resembles a seam weld and is the preferred method of joining for larger samples because it has a larger active area of 76.2 x 25.4 mm (3" x 1").

Figure 5. Patch weld shape in a Ti 6242- Ti 6242 VFA lap weld sample

The workpieces for patch welding were mechanically sheared into 0.51mm x

76mm x 90 mm (0.02" x 3" x 3.5"). The general bottom-up stacking order for the VFA 24 patch welding procedure for mechanical testing of the welds, shown in Figure 6a, is as follows: a patch foil insulated with a152 mm (6”) polyethylene sleeve is clamped on either side of the anvil that forms the base of the fixture. The flyer (moving workpiece) is cleaned with acetone and placed squarely on top of the foil. Following this, two G10 stand-offs are placed on top of the flyer, with a 31.75 mm (1.25") distance between them.

The target (stationary workpiece) is placed on top of the stand-offs so that it has a 50.8 mm (2”) overlap with the flyer (Figure 6c), and finally, a thick backing block made of impact resistant S7 tool steel is used to clamp the entire assembly together with the help of 12.7 mm (0.5") diameter alloy steel bolts. The base of the fixture that has two copper terminals that when connected to a capacitor bank, conduct the electrical discharge into the foil. The standoffs provide space for the flyer to accelerate up and an avenue for it to bend around and collide with the target at an oblique angle.

Figure 6. (a) VFA Patch welding setup, (b) 0.08 mm (0.003") thick CP Aluminum Patch Foil with an active area of 76.2 x 25.4 mm (3" x 1"), and (c) top view of VFA mechanical testing specimen

Spot Welding Setup

The spot-welding setup involves a few modifications to the welding fixture. The differences are as follows: 1) an insulated spot foil (Figure 7b) is used instead of a patch foil, and 2) a forming die (Figure 7a) is used to impart a standoff into the target material,

3) the same forming die aligned with the preformed target in the welding fixture where it acts as a backing block and also helps maintain the shape of the preformed standoff.

Figure 7. (a) VFA Spot Welding setup, (b) 0.08 mm (0.003") thick CP Aluminum Spot Foil with an active area of 10.16 x 12.7 mm (0.4" x 0.5"), and (c) typical spot lap welding sample

The pre-forming dies were lathed from shock resistant S7 tool steel and heat treated according to ASTM 4681 to achieve an average hardness of 46 on the Rockwell C scale. The 40 x 120 mm (1.57” x 4.72”) spot welding workpieces were mechanically sheared. Prior to welding, the targets were quasistatically indented with the standoff in a hydraulic press under a tonnage of 10 - 15 US Tons (9.07 - 13.61Tonnes) using the circular dies in Figure 8 and their corresponding punches. This allows for a built – in standoff to be introduced in the setup. Changing the depth and chamfer radius of the spot welding punch and die allows us to regulate the flyer impact angle and velocity. The shape and depth of the spot is different for the dissimilar Nickel – Nickel joints and the

Nickel – Titanium joints. The ductility of the nickel superalloys allows for the use of

27 sharper entry radii deeper spot dies (Figure 8a) for the similar and dissimilar Nickel –

Nickel spot welds. To account for the lower ductility displayed by Titanium in quasistatic deformation processes, the die used for the similar Titanium and dissimilar

Nickel-Titanium spot welds was redesigned to have a gradual deformation radius and reduced depth die (Figure 8b).

0.25” 0.04” (6.35 mm) (1.00 mm)

0.10” 0.06” (1.52 mm) 0.50” (2.54 mm) 0.50” (12.70 mm) (12.70 mm) (a) (b) Figure 8. (a) Deep spot die used for dissimilar Nickel – Nickel Spot Welds listed in Table 4, and (b) Shallow spot die used for Titanium – Nickel Spot Welds listed in Table 5

Acetone was used to removes oils, dirt, or debris from the preformed target and flyer surfaces. The two work pieces are held together using with insulative tape. The foil was insulated with a 152 mm (6”) polyethylene sleeve centered on its active area and aligned with the standoff using permanent marker lines. The foil, workpieces, and forming die are clamped together, as shown in Figure 7, and placed in a sealed steel chamber where the fixture is connected to a capacitor bank through the copper terminals.

As the capacitor bank discharges a quick current into the foil, the foil vaporizes, causing the resulting plasma to expand and accelerate the flyer upwards towards the target [8],

[11], [27]. The flyer deforms around the outer rim of the preformed spot and impacts the

28 target at an oblique angle (just as it did in the patch welding setup) to create a weld. The materials join near the rim of the recess during this impact welding procedure.

VFA Welding Parameters

The critical weldability parameters in impact welding are collision velocity and collision angle. These parameters can be varied by changing the foil vaporizing energy and standoff height between the flyer and target. As there is a scarcity in impact welding literature for the materials considered for this study, welding parameters for the similar welding combinations and dissimilar welding combinations were obtained by trial and error by varying the foil vaporizing energy and the standoff height. The parameters that yielded sound welds that endured both drop tests and mechanical cutting, are listed below. Each table lists the weld type dictated by the foil used, the flyer, target, foil vaporizing energy, the standoff type (G10 sheets for patch welds and preformed indents for spot welds), and standoff distance. Table 3 shows the parameters for creating welds between similar combinations of the materials presented earlier (Table 1). Table 4 and

Table 5 list the parameters for joining dissimilar combinations between the three nickel superalloys, and between each nickel superalloy and Ti 6242.

Foil + Standoff Weld Flyer Target Energy Foil Standoff Distance Type Type 718 718 Patch - 625 625 8 kJ Active Area 1.6 mm (0.06”) Patch 230 230 of 76.2 x G10 thick/31.75 (1.25”) 25.4 mm (3" mm apart Ti 6242 Ti 6242 9 kJ x 1") 718 718 2.54 mm (0.1") Spot - Active 625 625 deep Area of Preform (Figure 8a) Spot 230 230 8 kJ 10.16 x 12.7 15 UST mm (0.4" x 1.52 mm (0.06") Ti 6242 Ti 6242 0.5") deep (Figure 8b) Table 3. VFA Welding parameters for similar combinations

Foil + Standoff Standoff Weld Flyer Target Energy Foil Type Distance Type 718 230 Spot - Active 2.54 mm Area of 10.16 Preform Spot 625 230 9 kJ (0.1") x 12.7 mm 15 UST deep (0.4" x 0.5") 625 718 Table 4. VFA Welding parameters for dissimilar Ni - Ni alloy combinations

Foil + Standoff Flyer Target Energy Standoff Distance Weld Type Type Ti 6242 718 Preform Ti 6242 625 10 UST Ti 6242 230 1.52 mm (0.06") Spot 8 kJ deep 718 Ti 6242 Preform 625 Ti 6242 15 UST 230 Ti 6242 Table 5. VFA Welding parameters for dissimilar Ni - Ti combinations

Mechanical Testing and Characterization

Tensile Tests

To understand the joint integrity and efficiency of each weld combination, three different destructive mechanical tests were conducted on the welded samples: 1) lap shear test, 2) peel test, and 3) hardness test. An MTS 810 servo-hydraulic load frame was used for pulling the welded specimens at a crosshead speed of 0.1mm/s with zero preloading force. Figure 9 shows the testing configurations for both patch and spot welds. The 12.7 mm (0.5”) wide patch weld specimens from Figure 6c were gripped in the orientations pictured in Figure 9(a) and (b). For lap shear testing, the samples were gripped pulled as is without any modifications. The peel samples were prepared by bending the flyer to 900 and clamping the target end within a custom fixture gripped in the load frame. It was not

31 possible to peel the similar Ti 6242 patch welds as the base metal broke when bent to a right angle. A typical spot weld lap shear sample and its cross-sectional view are shown in Figure 9c. The overlap distance between the flyer and the target in this configuration is

40 mm (1.6”). This sample was not modified further for testing. To accommodate for the limited ductility of titanium, two different peel test configurations were used to measure the strength of the similar and dissimilar spot weld samples. Figure 9e shows the peel sample for a Ni – Ti joint. As the titanium fractured when bent at an angle of 900 at room temperature, the nickel superalloy end was bent instead to 1800 at a distance of 40 mm

(1.6”) from the edge closest to the weld (Figure 9d). This allows for the nickel side to be peeled from the titanium. The peel sample for the Ni – Ni joints were prepared by clamping the sample in a vice and bending both nickel ends to an angle of 900 with respect to the joint (Figure 9f) as is typical of the coach peel configuration.

Microhardness Measurements

Similar patch weld and dissimilar spot weld combinations were sectioned and mounted in conductive bakelite for hardness measurements and characterization of the weld interface. The purpose of these measurements is to verify the absence of a Heat Affected

Zone (HAZ) as is typical of the collision welding process.

Microhardness traverses were performed for each weld combination across the interface using a diamond indenter under a 200 gf load with a dwell time of 10 seconds. The

33 indented regions of measurement in and near the weld interface are depicted in Figure 10 with a dashed orange line. The distance between each measurement along the diagonal was 200 µm.

1 a

Alloy A b 2 3 c

4 d 5 Alloy B e

Figure 10. Depiction of typical regions of measurement from Energy Dispersive X-Ray Spectroscopy (green) and Microhardness indents (dashed orange) across an impact weld interface

Characterization: Light Microscopy, Scanning Electron Microscopy, Fractography

In order to better understand the sources of strength or weakness of the different welding combinations, the samples were observed under different microscopy techniques as listed in Table 6. The characterization techniques implemented provide insights into the following features: 1) Light Microscopy (LM) and Scanning Electron Microscopy

(SEM) reveal the shape and size of the bond line, and continuity or lack thereof of the weld line, 2) LM and SEM also show the presence, composition and concentration of 34 intermetallic compounds (IMC’s) and defects in the weld, 3) Energy Dispersive

Spectroscopy (EDS) across the weld interface provides the weight compositions of the different elements in and near the interface, 4) EDS on a mechanically fractured surface can also show material transfer if any.

The remnant edge pieces of the similar patch welds (Figure 6c), and the spot weld samples made for microscopy were sectioned using a mechanical shear or a resin bonded

SiC cut-off wheel. The sections were mounted in conductive bakelite and polished to a mirror finish using 600, 800, and 1200 grit pads using water, followed by a Chem-pad and colloidal silica. Metallographic samples of the 625 and 718 similar weld combinations were first observed under a Quanta SEM, and then electrolytically etched to be observed under an Olympus light microscope. While light microscopy was somewhat effective in observing the entire joint for dissimilar welds, Secondary Electron

& Back Scatter Electron Dispersive (BSE) imaging under an SEM was more effective in revealing detailed features of the interface.

Weld Characterization Flyer Target Etchant EDS Type Method Similar Combinations 10% Oxalic 718 718 LM and SEM Acid, 5 V N/A Electrolytic 10% Oxalic Patch 625 625 LM and SEM Acid, 5 V N/A Electrolytic 230 230 N/A N/A N/A Kroll's Ti 6242 Ti 6242 LM N/A Reagent Dissimilar Ni - Ni Combinations Line Across 718 230 SEM N/A weld Spot Line Across 625 230 SEM N/A weld 625 718 SEM N/A Across weld Dissimilar Ni - Ti Combinations Across weld and on Ti 6242 718 LM & SEM N/A fracture surface Across weld and on Ti 6242 625 LM & SEM N/A fracture surface Across weld Spot and on Ti 6242 230 LM & SEM N/A fracture surface On fracture 718 Ti 6242 LM N/A surface On fracture 625 Ti 6242 LM & SEM N/A surface On fracture 230 Ti 6242 LM & SEM N/A surface Table 6. Characterization methods used for assessing VFA welds

Line energy dispersive spectroscopy (EDS) scans were performed perpendicular to the weld interface for dissimilar combinations to obtain the primary compositions in and around the weld interface as shown by the green line in Figure 10. The different regions of focus are: 1) Alloy A rich, 2) Close to weld interface on alloy A rich side, 3)

Weld interface, 4) Close to weld interface on alloy B rich side, and 5) Alloy B rich.

Surfaces of the mechanically peeled dissimilar Ni - Ti welds were further observed in the

BSE mode under an Apreo SEM to understand the mode of fracture and qualitatively assess material transfer during the weld.

2.4 Results and Discussion

A majority of this work has been attempted for the first time and is very early on in the developmental stage. Therefore, the primary objective of this work was to verify the feasibility of the welding operation, i.e., whether or not a joint was created. The simplest way to assess this would be to use the results of the lap shear tests and microscopy of the weld interface. For the scope of this document, any weld that exhibits a non-zero lap shear failure load and displays discontinuous melting across the entire weld region is considered a successful joint.

The secondary objective is to report, and if possible, compare the joint strengths

(both lap and peel), hardness across the weld, typical and untypical features of the weld, and any weld defects such as cracks and voids. The desired outcomes are listed in the table below. Overall, it is a very high-level analysis by design.

Test Desired Outcome Significance Failure load exceeds Weld is stronger than the base metal Lap Shear parent strength (BM) Nugget Failure High adherence ability of parent Peel Nugget failure materials and material transfer for dissimilar welds (DW) Uniform across interface Microhardness and either at or above No heat affected zone (HAZ) base hardness No macro voids Lack of debonding Microscopy of the Lack of brittle intermetallics for No cracks weld interface DWs Discontinuous melting Robust joint Table 7. List of secondary objectives for joint mechanical and characterization tests, desired outcomes, and significance

It should be noted that every VFA weld presented in this document had a unwelded zone in the middle and should not be considered as a void or debonded region.

This is a typical feature in VFA based impact welds. Upon the vaporization of the foil, the middle region of the flyer is pushed upwards faster than the surrounding area [12],

[70]. It experiences a flat impact with the target leading to an unwelded region, as seen in

Figure 11. The lateral ends of the flyer, on either side of the unwelded region, will have an oblique impact angle. Under the correct impact velocity and collision angle, a successful weld is created. Should the appropriate conditions for wave formation be met

[8], [71], [72], interfacial waves will form at the joint. As the impact angles increases from the weld center towards the joint edge, the waves will gain in both amplitude and wavelength.

Figure 11. Typical cross-section of a VFA weld [12]

To help with clarity and flow of the discussion, this section is split into two major sub-sections to separate the results of the similar patch welds from those of the dissimilar spot welds. The findings of the similar patch welds are presented first. The though not unique combinations, the results will help support future developments in dissimilar patch welds containing one of these alloys as a joining component.

Similar Patch Welds

The results presented in the following sections represent the first trials conducted earlier on in the Nickel and Titanium VFA welding process. At the time of these trials, parallel efforts to optimize the VFA method for improved repeatability and mechanical properties were also taking place. The welding efforts were chronologically adapted to improvements in insulation strategies, fixturing, standoff design, and foil design.

Therefore, any initial observations of low strengths or undesirable failure modes in the similar patch welding results should not be taken at face-value as the enhancements in like-pair strengths are especially apparent in the spot-welding sections that follow. It is

39 expected that the impact angle for the Ni 230 welds was roughly 34 o and all other similar patch welds was 11o.

Overall, all the similar combinations accomplished the primary objective of a successful joint by consistently displaying non-zero failure loads and no signs of melting at the interface. Once the welding parameters were optimized, each combination welded easily. The results of the secondary objectives mentioned in Table 7 are discussed next.

718 – 718

The trends in lap shear load vs. displacement tests for the central pieces for the

718 - 718 patch welds from the are shown in Figure 12a. The load values are normalized by the width of the specimen (12.2 mm after water-jetting) instead of the area of the weld because it is difficult to consistently discern the full length of the weld (Figure 12(c1)) as waviness is not the only indicator of a bond [3]. To provide a base value for comparison, the normalized parent strength, obtained through a separate uniaxial tensile test, is provided. Also shown in this figure is a picture of a lap shear sample after the tensile test, the lap welds consistently failed in the base metal (BM). The onset of yielding for the lap welds is in close range to the yield point of the parent. After plastically yielding, the base metal exceeds strain hardens, causing the load to exceed that in the parent. The deformation occurs up to an average displacement of 12 mm (0.47”) where the material finally fails. The parent strain hardens for a more extended elongation in comparison and therefore has a higher ultimate tensile load. It is typical for the welded specimens to be limited in ductility [48] so this behavior is similar to that reported in literature. The

40 essential facet of these tests is that in lap shear, the weld is stronger than the base metal strength, which is a desirable feature of a robust joint.

Figure 12(b1-b2) shows the peel load vs. displacement curves for the 718 - 718 patch welds. The sizeable initial displacement shown by these samples is attributed to the extension displayed by the unwelded edge of the sample, clamped in the vice, as it bends into the correct angle for a peel test. The turning of the sample causes the low loads experienced until a 16 mm (0.63”) displacement, as there is no resistance to the force.

The curves are labeled according to the three types of failure modes that were seen during testing: 1) base metal on the target side (BM-T), 2) base metal on flyer side (BM-F), and

3) between parent and weld (BETWEEN).

Figure 12(b1) shows the sample curves, normalized by specimen width (12 mm –

0.47”) for which the failure load exceeded 1000 N. All but one of these samples failed in the base metal on the target side (side view of peel sample pictured, labeled ‘BM-T’) at roughly 30-44% of the maximum lap shear load. One sample failed partially in the first weld and then finally in the parent (pictured on the left labeled ‘Between’) at 26% of the

42 lap shear load. Some welds failed below 1000 N during peel tests, and the normalized curves are pictured in Figure 12 (b2). These welds majorly failed on the flyer side of the base metal (BM-F pictured on the left) at roughly 14-17% of the max lap shear load. One of these samples failed first in one of the two welded zones and then finally sheared on the flyer side of the base metal as well where the peak failure load fell into the same range as those that failed outside the weld. The purpose of these tests is to assess the integrity of the bond. As a majority of the samples failed outside the weld zone in the base metal, on either the flyer or the target side is another indication of a strong weld.

Figure 12(c1-c3) show electrolytically etched (10% oxalic acid) micrographs of the edge piece of the lap weld. Figure 12(c1) provides an overall view of the weld. The weld specimen consists of two welded zones on the left and right edges where the flyer impacts the target obliquely, and a central unwelded zone where the flyer impacts at a right angle to the target. Figure 12(c2) shows a magnified image of the left weld and

Figure 12(c3) shows the right weld. Both sides display a prominent bonded region with the wavy interface typical to impact welded metals [3], [8]s. Qualitative indicators of the high strengths of these welds are the continuous bond in the welded zone and the absence of voids, cracks, or microcracks. There is some evidence of a trapped jet that could have contributed to some of the weaker peel loads. However, it is less likely as a majority of the failure was in the base metal and not the weld. Other contributors to a sound weld are the lack of intermetallic compounds (IMCs), usually discernible by their contrasting color in dissimilar welds as seen in Figure 16. The black horizontal streaks parallel to the weld interface are likely niobium carbides, additional strengtheners in the 718 alloy. These

43 carbides have been reported to go into solution near the liquidus temperature during fusion welding, causing liquation cracking [73]. The location of these carbides in the surrounding matrix of the alloy and not in the weld itself are additional indicators of a good bond. Figure 12(c4) shows an SEM image taken near the weld interface that was manually recolored to black and white using Microsoft PowerPoint. The visible white streaks angled to the interface were initially suspected to be shear bands, a possibility in explosive metalworking of 718 [74]. However, when compared to etched images, this is likely caused by grains near the weld zone undergoing a mechanical shearing motion during the joining process. Compared to the metallographic image of Ti-6242 patch welds (red circle in Figure 15d), where the shear bands are very prominent, it is difficult to discern if the lines in the 718 patch weld are shear bands or just an artifact of the SEM.

Finally, the hardness traverse across the weld interface in Figure 14c displays no evidence of softening around the interface, common in fusion welds of alloy 718.

However, due to the severe plastic deformation at the welding interface, the hardness is almost double the original strength. Based on the application, this may or may not need a

PWHT to restore the properties of the weld.

625 – 625

The trend in lap shear loads for the 625 – 625 patch welds was similar to that of the 718 welds Figure 13a. Every lap shear specimen failed in the base metal and experienced yielding at the same load level as the parent. After yielding, the strain hardening in the base metal coupled with the strength of the weld, caused the load to

44 exceed that of the parent until an average displacement of 14 mm was reached, where the specimen finally broke. As mentioned previously, base metal failure indicates that the weld is stronger than the parent and is therefore the optimal mode.

The normalized peel test curves shown in Figure 13(b1-b3) are categorized by three failure types displayed: (b1) Partial failure in the parent and one weld zone, (b2) partial failure in only one weld zone, and (b3) full failure in both weld zones. The specimens in Figure 13 (b1) exhibited a max peak failure load between 23 – 26% of the maximum lap shear failure load. The second peak that follows the maximum load is similar to the partial weld failures seen in Figure 13(b2). It is possible that the maximum load in Figure 13(b1) was caused by the tearing of the parent material as shown. The peak loads that caused full failure in either one weld zone or both (Figure 13(b3)) are close to or higher than those samples that experienced partial failure in the parent and just one weld zone Figure 13(b1).

The optical image of an etched sample in Figure 13(c1) shows the transition from an unwelded region to the welded region marked by a sequence of waves. As the duration of the etching was limited to prevent over-etching, the grain boundaries are not clear. It is important to note that there are no visible voids, cracks or visible intermetallic compounds in this region.

While the weld interface is not visible in the SEM image in Figure 13(c2), some light streaks angled to the weld line can be seen. The reverse SEM image in Figure

13(c3) reveals the weld line and some of these streaks. As suspected previously in the

718 - 718 lap welds, it is possible that they are shear bands, a possible occurrence during high strain deformation of metals. Further microscopy and etching would be necessary to confirm their presence.

Similar to the 718 - 718 joints, the hardness across the interface (Figure 14c) was well above the base value reported in Table 2 due to the severe strain hardening as a result of the plastic deformation that occurs at the weld interface during the collision welding process. Again, this weld might or might not require a PWHT depending on the application.

It is unclear why these welds are weaker in peel than the 718 welds despite being compositionally similar. It is likely that the VFA process parameters used for 625 are different and need to be studied further to obtain stronger joints. Nevertheless, most of the 625 – 625 samples experienced some form of a nugget failure which is an indicator of a robust weld.

230 – 230

While all the other combinations could be joined with minimal preparation, the surface of the 230 workpieces had to be sanded with 120 grit sandpaper to remove the protective chromium oxide layer that exists at room temperature. Though impact welding is characterized by the tendency to jet off oxides during the collision process, it is 47 possible that due to the extremely high oxidation resistance of Ni 230 [35] the surface layer was not removed completely leading to easy breakage of the joints.

Figure 14(a) shows the normalized lap shear failure loads for Ni 230 categorized by the failure mode. In most cases, the failure occurred in the base metal and exceeded the yield strength of the parent, indicating the high strength of the joint. The one specimen that failed in the weld was an edge sample of the patch weld (Figure 6c).

Typically, this region will experience a dip in impact velocity compared to the region close to the center [75]. The reduced velocity is likely to have caused a weakness in the joint. Regardless, though its overall failure load is below the parent, it was still close to the other base metal failure loads.

All three cases of the peel tests displayed the ideal failure mode in the base metal indicating strong adherence between the flyer and target. The hardness traverse across the weld showed steady trends without any drastic decrease in strength. Similar to the 718 -

718 and 625 - 625 case, the hardness was higher than the initial base value due to the work hardening caused by the impact-induced plastic deformation at the interface.

Due to difficulties in etching the alloy, the interface of the joint could not be characterized. However, overall, the joint displayed the most consistent properties in both lap and peel compared to the other systems. It is possible if the other nickel alloys are also welded using the same standoff height and distance, the joint strength would be higher and more consistent.

There were issues in repeatability for the 230 – 230 patch welds when a 1.6 mm

(0.06”) thick stand-off was used. This could probably be attributed to poor process

48 control. When a like-paired spot-welding configuration resulted in a high strength joint, it was suspected to be due to the result of a steeper impact angle facilitated by the depth and diameter of the pre-formed indent. Trials made by increasing the patch weld stand-off height from 1.6 mm (0.06") to 2.4 mm (0.1") and decreasing the distance between the standoffs from 31.75 mm (1.25") to 12.7 mm (0.5") made a great difference in repeatably creating a successful joint in the patch weld configuration and should be used at the welding parameters for future experiments with all three like-paired superalloys.

(a) (b)

WELD

500 (c) 400

300

200 230-230 625-625 718-718 100 Flyer Side Target Side Hardness (HV) Hardness 0 -0.6 -0.4 -0.2 0 0.2 0.4 0.6 Distance from weld interface (mm)

Ti 6242 – Ti 6242

Figure 15a shows the lap shear loads vs. displacement curves of Ti 6242-Ti 6242 patch welds compared to the tensile strength of the parent. Ideally, these loads would have been normalized by welded or cross-sectional area. However, it was difficult to visually estimate the full length of the welded zone in the Ni - Ni patch welds through optical microscopy. The width of the welded zone is determined by the specimen width after water jetting. So, for the sake of uniformity, both the weld and parent strengths have been normalized by sample width: 9.2 mm in the case of the welded specimen and 12.2 mm in the case of the tensile sample.

During mechanical testing of welded samples, failure through the base metal is preferred to that in the weld as this is indicates that the weld is stronger than the parent.

For the Ti-6242 patch welds, while most of the lap-shear specimens failed in the base metal (as shown in the figure), the average peak failure load was 11 % lower than the parent strength. Additionally, even though the parent by nature had very little tensile elongation (displacement), the overall extension shown by the welded samples was far less in comparison.

The patch weld sample was sectioned, mounted, and polished to a 1 µm surface finish. Microhardness traverses were first taken across the weld interface (Figure 15b) at diagonal intervals of 300 µm between measurements. Compared to the base metal hardness of 330 HV, the region surrounding the weld interface only displayed a 6% increase in value. Microhardness measurements taken parallel to the weld interface on both the target and flyer sides also showed very little variation from the base metal. 50

While this is a good indicator of the lack of a HAZ typical to fusion welds, it is also a contradiction of impact weld characteristics. Commonly, in impact or explosive welds, the region around the weld interface undergoes severe plastic deformation which leads to grain refinement and increased dislocation density which leads to a significant increase in micro-hardness. Clearly, there is a secondary softening mechanism involved that is nullifying the expected increase in hardness.

The same mounted samples were etched with Kroll’s Reagent to reveal the grain boundaries and morphology of the interface. A stitched image of one of the two welded zones can be seen in Figure 15d, along with magnified images of the pertinent regions.

The welded region is continuous and devoid of cracks, debonded zones. The edge of weld interface is decorated with angled streaks resembling adiabatic shear bands (ASBs), circled in red in Figure 15(d1). These ASBs are not visible in the in the wavy regions (d3) that formed closer to the unwelded zone at the middle of the weld. ASBs are regions of severe plastic deformation that are accompanied with thermal softening [76]. Shear bands have been reported as a common occurrence in severely strained Titanium alloys and are suspected to have a negative effect on creep properties making them undesirable [77],

[78].

Liu et. al, [70] stated that ASBs were more prominent at higher impact angles.

According to VFA welding velocimetry and collision angle analysis conducted by Lee et. al [75], along the direction of the weld (from center to edge), the impact velocity decreases and the impact angle increases. Therefore, the presence of the ASBs closer to the edge of the weld are justified. By modifying the VFA setup to achieve low angle welds, it is possible to achieve high lap shear strength without deteriorating the creep properties of the material.

A good example of strength optimization can be seen in Figure 15c. The tensile lap-shear loads obtained from Ti 6242-Ti 6242 pre-formed spot welds (12.7mm in width) have almost twice the normalized peak loads compared to the patch welds in Figure 15a.

As the target workpiece is pre-formed with a built-in standoff for the VFA spot welding

52 setup (Figure 7), the amount of deformation of the indent dictates the angle of impact of the flyer. It is very likely that the high springback Ti 6242 proved advantageous in decreasing the impact angle and therefore increasing the strength of the weld.

Unfortunately, due to time restrictions the Ti - Ti spot weld interfaces were not characterized. Future efforts involving VFA Ti welding need to take this aspect into consideration.

Regardless, the ability to obtain a high strength Titanium joint has been demonstrated using the VFA method. This marks a very significant advancement in the welding of Titanium alloys which is often expensive due to standard limitations dictated by the chemistry of the atmospheric (oxygen) and the cleanliness of the workpiece surface.

Dissimilar Spot Welds: Nickel – Nickel

The following sub-section describes the results of the dissimilar Ni - Ni spot welding experiments. Arcing issues notwithstanding, all three combinations satisfied the primary objective of a successful weld. Overall, the dissimilar Ni - Ni welds satisfied all the requirements for the secondary objectives in Table 7. Based on the geometry of the pre-forming die in Figure 8a, it is expected that the impact angle for each dissimilar Ni -

Ni weld was roughly 34o. Further details of this analysis are presented next for each combination. Appendix A has further SEM microscopy images for these combinations.

718 – 625

Most of the Ni - 718 - 625 spot welds showed high peak loads in the lap shear mode

(Figure 16a). Two out of three welds displayed nugget failure in the 718-base metal, pictured on the right in (Figure 16a). Unfortunately, one of the welds sheared in 718 around the imprint left by the circular die during the welding process, shown on the left of (Figure 16a), this curve is given as a lower baseline for reference. The peak load of the other two welds was 4% higher than the lap shear load of a similar 718 - 718 (parent material) lap weld, made using the same setup and conditions. The strain hardening displayed by these load-displacement curves was roughly the same, whereas the dissimilar welds showed a slightly higher elongation. Each of the similar 718 - 718 lap welds showed nugget failure in the parent indicating the high strength of the direct joint made between Ni - 718 and Ni - 625.

The results of the T - Peel loads, shown in (Figure 16b) also exhibited repetitive nugget failure in 718 (pictured on the left) suggesting a strong bond between the two base metals. One of the curves displayed a slightly higher peak caused by minor tearing of the parent metal in the nugget during the test. A similar failure mode can also be seen pictured in the 718 – 230 peel tests in Figure 17b.

(a) (b)

718

718 718

625

Nugget in 718 625 625 Sheared Nugget in around Die 718 Imprint

(c1) 718 (c2) 718 (c3) 718

625 625 625

(d)

718 Interface 625

The composition of these pockets is indicated by Region 3 of the line EDS scan taken across the weld interface (Figure 16d). The lack of a gradient in the Nb and Ti 55 compositions reveals that the intermediate phase is majorly comprised of Ni-Cr-Fe.

Analysis of the binary phase diagrams between these elements shows that around 1200oC

- 1300oC (2192oF - 2372oF), the melting ranges of 718 and 625, these elements exist in solid-solution with each other. The isolation of these pockets can be attributed to the compositional similarity between the 718 and 625 alloys. Intermetallic phases are usually inevitable within direct dissimilar due to the mixing of the major alloying elements.

Depending on the composition and intermetallics are brittle and serve as sites for crack initiation during mechanical loading. Therefore, isolated pockets are preferred as opposed to continuous layers to arrest crack propagation.

The high pressures and fast solidification rates experienced during impact welding can cause cracking in intermetallics. Furthermore, based on the welding literature of these metals, Ni - 625 is susceptible to solidification cracking [19], [41] at the interface. None of the high magnification scans revealed crack or voids either within or around the interface or in either base metal.

The absence of voids, cracks, debonding, and the continuity of the waves in addition to the nugget failure modes displayed during mechanical testing, are reliable indicators of a good weld with high joint efficiency.

718 – 230

(a) (b)

718 718

230 230 230

230 230 Nugget + Nugget

718 Tear in 718 in 718 718 WELD WELD Nugget (Arced) (Possibly Arced) in 230

(c1) (c2) 718 (c3) 718 718 230

230 230

(d)

718 Interface 230

One of the three 718 - 230 spot welds displayed nugget failure in the 230 base- metal during lap shear test, which exceeded the maximum load of the 230 - 230 similar spot weld by 1000 N (Figure 17a). The other two spot welds failed at very low peak loads within the weld. Examination of the base metal weld surfaces (pictured) of the latter two revealed that 1) the impact point was off-center, and 2) very little to no waviness was visible around the impact center. On the other hand, the three peel samples all showed repetitive nugget failure in the 718 (flyer) base metal, as shown in Figure 17 (b). One of the peaks is 410 N higher than the other two because of the partial tearing in the base metal pictured on the left as indicated by the red circle.

The SEM images of the 718 - 230 spot weld, pictured in Figure 17(c1-c3), display many qualities of a good impact weld, previously described in this document. The 230- side has tungsten and chromium carbides that appear as white spots against the light gray matrix and the 718- side has niobium and titanium carbides that also appear as white spots against its dark gray matrix. The waves are continuous, extend to the edge of the sample, and increase in amplitude in the horizontal direction from left to right. The edge of the weld closer to the center of impact has discrete, isolated intermixing (medium gray in Figure 17(c2)) on the crests and troughs of the waves. The quantity and size of mixed regions minimize to almost zero near the right edge of the sample, as shown in (c3). The line EDS scan across the 718 – 230 spot weld interface (Figure 17d) shows that the major elements at the interface (region 3 on the x-axis) are Ni, Cr, Fe, and W. The presence of the tungsten comes from the tungsten carbide in the form of a white spot engulfed in the mixed region in Figure 17(c2) meaning that the tungsten did not dissolve into the matrix.

As discussed in the previous section, Ni, Cr, and Fe form a solid solution around the melting temperature range of these alloys. No liquation or solidification cracks, and microfissures typical to welding of Ni - 230 were visible at the interface.

A few significant observations can be made regarding the 718 - 230 welds.

Despite the tear, this consistent nugget failure suggests that the weld is stronger than the base metal and that there is sufficient material transfer to create a bond between both dissimilar metals. The nugget failure occurs closer to the rim of the spot weld preform where the concentration of intermetallics is negligible. The three nugget failures in peel and the one in lap-shear suggest that there was possibly some level of experimental error in the two lap-shear tests that failed in the weld. Additionally, the lack of visible waves on the 3000 N specimen and the minimal extension of the waves in the 7000 N specimen suggests that the velocity of the flyer was not high enough to create a joint at the interface. The most likely cause of this is arcing in the foil caused by improper insulation.

During arcing, the current discharged from the capacitor bank finds an alternate conduction path preventing it from entirely going into the foil, thus reducing the efficiency of the vaporizing process and therefore the pressure from the expanding gasses.

Under optimal VFA welding parameters and conditions, the Ni - 718 and Ni - 230 can exhibit high dissimilar joint strengths while maintaining properties of the parent materials. If the experimental errors are disregarded, with proper insulation and alignment, the 718 and 230 alloys have much potential for being joined directly, without the need for a filler/interlayer or post-weld heat treatments.

625 – 230

(a) (b)

230 230

625 625

230 Nugget Nugget in 625 in 230 Nugget in 230

(c1) (c2) (c3) 625 625 625

230 230 230

(d)

625 Interface 230

The lap shear load-displacement curves for the Ni - 625 – 230 spot welds are shown in Figure 18 (a). Both the spot welds failed in the Ni - 230 base metal leaving behind a weld nugget (pictured). The average peak load of both these welds exceeded the

230 - 230 spot weld load by an average of 50 N. The general trends in stiffness and yielding resemble that of the 230-base metal. In the T-peel mode Figure 18 (b), all three welds that were tested experienced nugget failure in either the 625 or 230 base metal. The peak loads of and displacements of the welds that left a 625 nugget were approximately equal, while the weld with the nugget failure in 230 had a peak load that was 350 N higher. Regardless of the base metal type, the 625 – 230 spot welds demonstrated the ideal failure mode for these types of joints.

Part of the strength contribution comes from the features of the weld interface.

The joint interfaces show in Figure 18 (c1-c3) show white spots in the light gray matrix.

There is a thick continuous intermetallic layer, shown in medium gray between the 230 and 625 parent materials (Figure 18 (c1). This layer is within a short distance from the center of the flyer impact zone, where the metals are very likely to experience the highest interfacial temperatures and pressures. The intermetallics become more isolated closer to the edge of the sample, as shown in Figure 18 (c2) and (c3). These pockets help arrest any cracks that may form as they are weaker than the surrounding material. Similar to the

718 - 230 welds, the white spots within the medium gray intermetallics in Figure 18 (c1), are evidence that the maximum temperature at the interface was below the dissolution temperature of the primary tungsten carbides. At region 3 of the line EDS scan over the weld interface, (Figure 18 (d)), nickel and chromium have the highest weight

61 concentration. These elements exist in a solid solution near the individual melting temperatures of Ni - 625 and 230. The lack of weld defects (cracks and voids) within the intermediate phases are evidence that these phases are not as brittle as most other IMCs that form during dissimilar impact welds.

The Ni - 625 – 230 welds are capable of exhibiting high lap shear and peel loads with consistent nugget failure in the base metal demonstrating the high strengths of these welds.

Dissimilar Spot Welds: Nickel – Titanium

This section describes the results from mechanical testing and characterization of the Ni – Ti alloy combinations described in Table 5. Initially, the dissimilar Ni - Ti welds were made using titanium as the flyer due to its lower weight and hence capability of achieving a higher impact velocity. Additionally, when preformed at the same force, the nickel alloys had better shape conformation than titanium. However, these samples exhibited low failure loads compared to the base metal strengths. As the goal of this study was not only to optimize the VFA welding parameters to create a joint between these materials but also to obtain high strength welds, efforts to improve the joint efficiency were tried.

As part of the VFA welding parameter study, flyer velocity traces for Ni - 718 and

Ti 6242 for the spot-welding configuration were obtained using a Photon Doppler

Velocimetry (PDV) system (described in detail in Chapter 4). These curves, pictured in

Figure 19, showed that for the same foil vaporizing energy and standoff distance, both metals had roughly the same velocity, which meant that either could be the flyer

62 workpiece. As Ni - 625 (8.44 g/cc), and Ni - 230 (8.97 g/cc) are only slightly denser than

718 (8.22 g/cc), their impact velocity can be expected to be close to or slightly lower than

718 and Ti 6242. Therefore, welds with the titanium as the preformed target (15 UST) and nickel as the flyer were made. Preliminary lap shear tensile tests of a Ti-Ni bond showed that the strength of the weld exceeded that of the Ni - Ti bonds both in lap and peel modes. These load vs. displacement values are compared to the base metal spot weld of Ti - Ti in this section. Flyer velocity traces for both VFA patch and spot welds for all four materials can be found in Appendix B.

Figure 19. Flyer velocity vs. distance plots for Ni - 718 and Ti 6242 in the spot-welding configuration at a foil vaporizing energy of 8 kJ

Two types of figures were made to describe the strength related properties for each dissimilar weld, qualitatively. The first type is a combination of the lap-shear

63 curves, modified peel curves, SEM images and line EDS scans of the weld interface for the dissimilar joints. The designation ‘Ni – Ti (flyer)’ is used for welds made with the nickel alloy target and titanium flyer, and ‘Ti (target) – Ni’ is used for welds made with a titanium target and a nickel alloy flyer.

The second type of figure consists of SEM images and EDS scans of the peel fracture surfaces to visualize and quantify any material transfer between both workpieces, especially for specimens that failed in the weld. Appendix A has further SEM microscopy images for these combinations.

Based on the geometry of the pre-forming die in Figure 8b, it is expected that the initial impact angle for each dissimilar Ni - Ti weld with a Ti-flyer was roughly 11o.

However, analysis of samples that showed signs of misalignment, imply that this varied between 7.8o to 9.6o. Geometrical analysis of the Ti-Target joints put the initial impact angle estimate between 6o - 8o.

718 – Ti 6242

(a) (b)

WELD

Weld, Ti-Flyer PARTIAL NUGGET IN TI

(c1) (c2) Ti - Target

(e1)

Ti - Flyer Ti - Flyer

(c3) Ti - Target

(e2)

Ti - Flyer

(d) Ti - Target

(e3)

Figure 20. Mechanical properties and characterization of 718 – Ti 6242 VFA spot weld (a) lap shear loads w.r.t parent strength and failure modes, (b) peel loads and failure modes, (c1 – c3) SEM + BSED images of weld interface for 718 - Ti (flyer) joints, (d) EDS: average composition (wt. %) of major alloying elements across weld interface taken at the different regions in Figure 10 for 718 – Ti (flyer) combinations, and (e1-e3) LM images of weld interface for Ti (target) – 718 joints

Two out of three of the three 718-Ti (flyer) lap samples experienced interfacial failure in the weld (pictured top left of Figure 20a), and the third broke partially in the titanium flyer (bottom right of Figure 20b). The peak failure load ranged from 2000 to

8000 N with an average force of 6000 N. The two repeated failures at 8000 N could be indicative of the actual average peak load meaning that the 2000 N weld is an outlier, possibly an arced sample. As the partial failure occurred in the Ti flyer, these curves are compared to the lap shear strength of a Ti - Ti spot weld made using the same setup and conditions. It should be noted that the partial weld failure was a caused by an initial impact angle of 8o. Relative to the parent, all three welds failed well below the peak load of 11500 N. In contrast, the Ti (target) -718 weld failed at 13500 N. Even though this failure was in the weld interface, the bond held long enough to accommodate for the elongation to 0.6 mm beyond the parent.

All three 718-Ti (flyer) peel samples in Figure 20b failed in the Ti flyer, leaving behind a partial nugget (pictured on the bottom right). The failure loads ranged from 900

N – 1300 N, averaging at 1150N. From the initial rise in load at 13 mm, these samples further extended to a displacement of 10.7 mm. The Ti (target) - 718 peel sample extended to almost three times the displacement, peaked at 1200 N and ultimately failed in the weld at 1100 N.

SEM images of the weld interface show waves enveloped in intermetallics closer to the center of impact shows Figure 20(c3) that develop into continuous waves seen in

Figure 20(c1) and eventually lose amplitude as those seen in Figure 20(c2). The thickness of the intermetallics ranges anywhere from 4 – 37 µm.

Assuming the temperature at impact ranged from1000 C- 1300 C, the possible intermetallics and phases that can occur for the compositions in Figure 20d are TiNi,

TiNi3, TiFe, α-TiCr2, and β-Ti. TiCr2 and TiFe are laves phases that form either due to long term exposures to high temperatures or at the end of solidification. As the high temperatures at the impact weld interface last for a fractional amount of time due to the high cooling rate, the likelihood of the existence of these phases is very low. The

Chromium and Iron at the interface were likely detected due to their presence in solid solution within the 718 matrix.

Regular waves observed close to the center of impact (Figure 20(e1)) that flatten out closer to the edge of the spot weld. The intermetallics in the wavy regions are discontinuous and seem to penetrate onto the Ni side of the interface. The thickness of the intermetallics was 6 µm thick both near the wave and at the edge of the weld (Figure

20(e3). Thee continuous layer visible in Figure 20(e3), would be the brittle zone where fracture initiated, thus leading to full failure through the interface as was seen in both lap and peel setups.

The fracture surface of the 718 side of the 718-Ti (flyer), shows vivid and regular wavy features (Figure 21a). The EDS Scan to its right (Figure 21b) shows equivalent transfer of Titanium in the valleys of the waves of the more ductile 718. In contrast, the fractography taken of the titanium side of this same region (Figure 21g), seems to have wavy features with sharper peaks and flat facets, resembling brittle regions. However, it is littered with pockets of dimples which appear to be composed of both titanium and nickel (Figure 21h). In both cases, there is definite evidence of material transfer

67 indicating a strong bond strength between both base materials in those regions. However, the transfer is uneven and discontinuous which would correspond to complete or partial failure through the weld interface.

A higher magnification image of the titanium peel surface of the Ti (target) – 718 welds reveals the wider distribution of the 718 alloy onto the Ti target. The wavy features on the titanium side for this weld are more pronounced than the ones in (Figure 21g).

These features match the fracture surface of the 718 side. The EDS scan of the nickel surface is missing due to technical problems. Nevertheless, the larger displacements seen in the lap and peel modes attest to the mechanical and chemical interlocking between the two base metals and the high strengths imply that the bonding between the two base metals was sufficiently strong.

625 – Ti 6242

(a) (b)

PARTIAL NUGGET IN TI

WELD

(c1) (c2) Ti - Target

(e1)

Ti - Flyer Ti - Flyer

(c3) Ti - Target (e2)

Ti - Flyer

(d) Ti - Target (e3)

There was a lot of variability in the 625 – Ti (flyer) lap welds. Only two are pictured in Figure 22a, because one sample sheared around the circular backing die. The peak loads range from 7500 N – 9000 N and the displacement from 50 – 150% of the base metal value. Both samples failed through the weld interface and below the peak load in the parent (Ti - Ti spot weld) at 11500 N. The Ti (target) - 625 peaked slightly above the parent at 11550 N and failed across the weld interface with similar elongation.

The peel strengths for the 625 – Ti (flyer) welds show more repeatability compared to lap as seen in Figure 22b. The peak loads ranged from 450 N to 600 N. All three tests resulted in partial failures through the Titanium base metal. The displacement from the first plateau that begins at 5 mm, extended to 20-26 mm until final failure.

The Ti (target) - 625 peel sample showed much higher failure load of 950 N and extended from 10 mm to 39 mm before failing through the weld.

These trends can be explained by the microscopy of the weld interface and fractography of the peel surfaces. SEM images across the 625– Ti (flyer) revealed regions of regular wavy interfaces (Figure 22(c1-c2)) that occurred in the incline of the spot weld preform, but also clear regions of debonding as seen in Figure 22(c3) that occurred closer to the center of the weld. The waves in Figure 22(c3) indicate that while the two materials tried to shear into each other, the bond was too weak to accommodate the strain of impact. Additionally, the visibly thick intermetallics which varied from 2 – 20 µm, show noticeable cracks. The cracks which are a result of their brittle nature would have diminished the bonding as well. A macroscopic image of this weld showed that a wavy

71 region precedes this debonded region. The lack of symmetry in the left and right welded zones corresponds to misalignment between the die, target, and foil.

Based on the line EDS scan across the weld interface from Figure 22d, and the same reasoning used for the 718 - Ti (flyer) welds, the likely intermetallics that exist at region 3 are TiNi and TiNi3.

The Ti (target) – 625 weld interface shows discrete pockets of intermetallics which ranged from 2-7 µm in the wavy regions of Figure 22(e1-e2) and became a 17 µm thick continuous layer near the edge of the sample (Figure 22(e3)). Visible cracks going through the intermetallics at the edge of the weld in Figure 22(e3) resemble those seen previously in Ti (target) – 718. The edge would be the least ductile and weakest region of the weld and any fracture initiated in this region would have propagated through the weld.

EDS scans of the 625 – Ti (flyer) peel surfaces show the transfer of titanium onto the peaks of the wavy regions of 625 base metal (Figure 23b). In contrast, very little 625 can be seen in the valleys of the waves on the Ti side in Figure 23h. The distinct wavy features visible in both base metals would contribute to the ductility of the joint by the discontinuous transfer in the material would limit its strength.

A higher magnification image (Figure 23c) of the titanium peel surface of the Ti

(target) – 625 welds, shows irregularly shaped blocky islands with flat faces. This image was taken close to the edge of the weld sample and would fall somewhere between

Figure 22(e2) and Figure 22(e3). The nickel transferred onto this surface probably corresponds to a mixture of base metal and nickel tied up in intermetallics. These features are likely to have contributed to the dip in load around the 12 mm mark in the peel curve.

The 625 side of this same sample shows distinct wavy features away from the edge where significant regions of titanium transferred from the target. The continuous areas of titanium in the troughs of these waves corresponds has a strong correlation with the bond strength and overall ductility of the sample.

230 – Ti 6242

(a) (b)

WELD PARTIAL NUGGET IN TI

Ti - Target (c1) (c2) (e1)

Ti - Target Ti - Flyer Ti - Flyer (e2) (d)

Ti - Target (e3)

The 230 – Ti (flyer) demonstrated the lowest loads in lap shear compared to the other two superalloys. The peak loads in Figure 24a ranged from 3000 N to 4500 N and were well below the parent strength of the Ti - Ti spot. Both samples failed early, in the weld interface, which limited the displacement to 33 – 50% of the parent. The strength of the Ti (target) – 230 lap weld was 25 N below the max load of the parent spot weld. It exceeded the maximum displacement of the parent by 86% despite failing through the weld interface.

The 230 – Ti (flyer) peel samples showed similar trends in loads that varied from

500 N – 650 N with an average elongation of 7 mm (Figure 24b). A majority of these welds failed partially in the titanium base metal, which created some variance in the peak load. In contrast, the Ti (target) – 230 peel sample failed through the weld at 600 N after stretching to 27 mm.

The regular waves seen in the 230 – Ti (flyer) weld interface Figure 24(c1) show intermetallic compounds in the valleys of the waves that occasionally penetrate the wave itself. The irregular waves in Figure 24(c2) can constrain the intermetallics to the surroundings of the interface.

The line EDS scan in Figure 24d of the weld interface shows a 50-50 ratio by weight of nickel to titanium. Compared to all the other Ni - Ti welds, the 230- Ti had the highest amount of titanium at the interface which would also create the largest concentration of TiNi and Ti2Ni intermetallics.

While the interface of the Ti (target) – 230 displayed qualities that typically enhance the bond strength such as lack of voids and discontinuous intermetallics Figure

24(e1), it was the first to show cracks in the base metal (Figure 24(e2)), carbides (bottom left of Figure 24(e2), and the intermetallic layer (Figure 24(e3)). The other two combinations had shorter cracks, perpendicular to the base metal. In contrast, the cracks in the intermetallic layer of the Ti (target) – 230 interface were long and, angled to the base (Figure 24(e3)).

The cause of the cracks in the 230 base-metal in Figure 24(e2) is unclear.

Technically, the surrounding region above is more brittle and more likely to be decorated with defects. Based on their orientation w.r.t the weld line, is possible that these are

ASBs, however further microscopy is needed to confirm.

Fractography of the nickel side of the 230 – Ti (flyer) peel surface shows a highly cracked region in the top right corner (Figure 25a). The EDS scan in Figure 25b reveals a high concentration of titanium in this area. While the titanium side of this same sample

(Figure 25g) shows many flat facets typical of brittle fracture, there is a discernible transfer of the 230 base-metal onto the flyer as seen in Figure 25h.

In contrast, the titanium side of the Ti (target) – 230 surface show more ridge-like shapes which have both brittle and ductile regions (Figure 25c). The ridges might be the result of a strong bond which resisted the pulling motion of the peel test. While the wavy features on the nickel side (Figure 25i) are less discernible, the amount of titanium on the

230 side is more than twice the amount of nickel (Figure 25j).

2.5 Summary and Conclusions

The versatility and modularity of the VFA method have been implemented in creating ten similar and dissimilar joints between Nickel 718, Nickel 230, Nickel 625, and Titanium 6242. All ten types of weld structures were subjected to different mechanical tests and characterization techniques to investigate the strength and quality of the joints. The peel tests, in particular, are a new mechanical testing method for these

77 materials. They were especially useful for qualitatively examining the adherence of the similar joints and material transfer for dissimilar combinations.

It was found that all four similar Ni - Ni and Ti - Ti combinations were capable of producing ideal welds in the VFA patch weld geometry. These joints were easy to make and in lap-shear mode, the patch welds almost always displayed base-metal failure (Table

8), an excellent indicator of a sound weld. While the peel tests for the patch welds revealed mixed failure modes in all cases except Ni - 230, the failure loads are suitable for benchmarking the strengths for future optimization experiments. Hardness measurements across the weld interface showed the absence of a HAZ in these welds.

Microscopy of the weld interface showed wavy structures typical to impact welding, devoid weld defects such as cracks, debonding, and voids.

Similar Patch Weld Lap Shear Peel Common Common Flyer Target Avg. Load Failure Avg. Load Failure Mode Mode Base Metal - 718 718 4130 N Base Metal 1605 N Target 625 625 4170 N Base Metal 1045 N Mixed 230 230 4690 N Base Metal 1120 N Base Metal Ti 6242 Ti 6242 10670 N Base Metal N/A Table 8. Summary table of failure loads and modes for similar patch welds

The VFA spot weld geometry was suitable for joining all ten combinations and displayed many sound weld characteristics. It was especially optimal for obtaining robust dissimilar Ni - Ni and Ni - Ti joints.

Disregarding the cases involving poor insulation or workpiece misalignment, each dissimilar Ni - Ni combination experienced full nugget failure in the parent metal in both lap shear and peel modes (Table 10). In addition, the peak lap shear loads exceeded the strengths of like-paired spot welds (Table 9). SEM scans of the weld interface also revealed a continuous weld interface with wavy features decorated with melted regions of intermixing. By changing the impact velocity and angle it is possible to avoid melting, however for the case represented here, it is beneficial that they are discontinuous and free of cracks. These combinations displayed all the desired weld properties characteristic of strong sound joints.

Similar Spot Weld Lap Shear Flyer Target Avg. Load (N) Common Failure Mode 718 718 8035 Nugget 625 625 7235 Nugget 230 230 9120 Nugget Ti 6242 Ti 6242 11000 Nugget Table 9. Summary table of lap-shear failure loads and modes for similar spot welds

Dissimilar Ni - Ni Lap Shear Peel Spot Weld Common Common Flyer Target Avg. Load (N) Failure Avg. Load (N) Failure Mode Mode 718 230 9875 Nugget* 1250 Nugget 625 230 9460 Nugget 1715 Nugget 625 718 8340 Nugget 1230 Nugget Table 10. Summary table of failure loads and modes for dissimilar Ni - Ni spot welds

Obtaining high strength dissimilar joints between Nickel and Titanium alloys has been challenging to accomplish in the past due to the brittle nickel-titanium based IMCs that form at the weld interface. It has typically been achieved using expensive methods such as diffusion bonding which has limited the potential applications of these joints.

With some process optimization of the VFA process, it was possible to create dissimilar

Ni - Ti welds. These joints displayed mixed failure modes in both lap shear and peel

(Table 11). Lap shear loads made with a Ti-flyer were found to exceed the parent strengths even though these joints demonstrated interfacial failure. Fractography and

EDS scans of the peel surfaces showed both ductile and brittle facets that showed evidence of material transfer between both parent materials. While some of these failure modes were a consequence of workpiece misalignment, these welds still exhibited some critical findings that could potentially lead to improved joint efficiency.

Dissimilar Ni - Ti Lap Shear Peel Spot Weld Common Avg. Load Common Failure Flyer Target Failure Avg. Load (N) (N) Mode Mode 718 7835 Mixed 1150 Partial Nugget Ti 6242 625 8480 Weld 560 Partial Nugget 230 3820 Weld 565 Partial Nugget Weld/Material 718 13500 Weld 1200 Transfer Weld/Material 625 Ti 6242 11700 Weld 960 Transfer Weld/Material 230 11500 Weld 600 Transfer Table 11. Summary table of failure loads and modes for dissimilar Ni - Ti spot welds

The work presented here serves as a benchmark in lap-shear strength, peel strength, and microstructural features for the different nickel and titanium impact welding combinations. Future work in optimizing the failure modes and structure of the interface can build on this existing information. As these materials are typically used in high temperature and corrosive environments, the strength of these joints still needs to be assessed under the corresponding service conditions of the desired application. Further suggestions can be found in Chapter 5.

In addition to creating some unique joint combinations (718 - 230, 625 - 230,

6242 - 230, and 6242 - 718), through this work, it has been demonstrated that it is possible to create high strength solid-state joints between sheet metal structures of a near alpha titanium alloy and nickel superalloys. The joints retained the original mechanical properties of the parent materials and avoided contamination with carbon, oxygen,

81 hydrogen, and nitrogen, a big problem in current industrial welding practices with these alloys. While further exploring the behavior of these joints under the loading conditions of their specific application needs to be conducted, it has been shown that the VFA process has great potential for impact joining nickel and titanium alloys.

Chapter 3. Impulse Based Shape Calibration using the Vaporizing Foil Actuator Method

3.1 Background and Motivation

Sheet metal forming of simple profiles is an essential forming step in the industrial production of aerospace and automotive components. However, as the industry moves towards light weight and high strength materials, springback during the forming process of these formed parts has become a more significant problem. The precise bending of structural alloys proves to be difficult due to their insufficient formability and severe springback resulting from a high ratio of yield stress to elastic modulus. Springback results in the production of parts that are beyond the dimensional tolerances and therefore scrap, dramatically increasing manufacturing costs [79].

A standard method for springback control includes geometrical compensation where the workpiece is either over-bent using a mold designed for resolving springback or bent with additional stretching applied. However, these methods are not consistently valid for alloys with different mechanical properties or process parameters and can be destructive to the workpiece as high stressed are involved [2], [79], [80]. Pre-heating the tools, molds, and workpieces before bending, is another process that can significantly increase the formability of workpiece thus reducing the required bending pressure due to the elimination of the residual stresses inside workpieces after bending [81]. Hot forming requires additional equipment and temperature control that could affect the microstructure of the metal workpiece leading to weakened mechanical properties.

Impulse forming methods, such as explosive, electromagnetic, and electrohydraulic calibration have been proven to dramatically improve the dimensional accuracy of thin-walled

83 parts [3]. The high springback relief observed in the components manufactured through these processes is likely a direct result of the relief of elastic stresses due to the propagation of stress waves in the work pieces. However, each of these methods have their limitations. Explosive forming comes with the burden of safety issues, electromagnetic forming is not suitable for non- conducting materials, and electrohydraulic forming is still in the developmental stages [1], [7].

High-temperature metallic materials such as titanium and nickel-based alloys are attractive as skin structures for aerospace vehicles. They can allow significant performance improvement and mass reduction in aircraft. However, there are significant challenges in forming them affordably for service. This project examines the use of impulse forming methods, as enabled by the vaporizing foil actuator method, for the precise shaping of Ti 6242, Ni - 230,

Ni - 625, and Ni - 718. The effects of the shockwave on the microstructure, hardness, and springback relief is explored. The preliminary results show that this method has great promise for precision forming to net-shape dies.

3.2 Materials and Methods

Materials

The shape calibration capabilities of the VFA method was demonstrated using the six different high strength metals listed in Table 12. These metals exhibit a yield strength to elastic stiffness ratios (σy /E) between 0.002 – 0.008. Sheet metals with high σy /E ratios and lower thickness tend to exhibit higher levels of springback.

0.2% Yield Elastic Condition Material Thickness σy / E Strength (σy) Stiffness (E) Ti 6242 0.51 mm 930 MPa 114 GPa 0.008 Duplex Annealed AA 6061 1.00 mm 276 MPa 69 GPa 0.004 Aged to T6 CP 0.90 mm 340 MPa 102 GPa 0.003 Annealed Titanium Ni - 718 0.51 mm 416 MPa 200 GPa 0.002 Annealed Ni - 625 0.51 mm 426 MPa 208 GPa 0.002 Annealed Ni - 230 0.51 mm 390 MPa 209 GPa 0.002 Annealed Table 12. Thicknesses and strengths of VFA calibration metals

Experimental Methods

The springback removal was performed in two steps: 1) quasi-static pre-deformation of the workpiece to the target shape, 2) calibration to the target shape using the Vaporizing Foil

Actuator method. The pre-forming step provides a visual representation of the amount of springback for each material and prevents any gaps between the workpiece and the forming tools within the fixture. This step is crucial in isolating the impulse based springback removal process from high velocity forming/stamping or impact forming.

Pre-forming Step

Square 76.2 mm x 76.2 mm (3” x 3”) blanks of each material were mechanically sheared and stamped between a flange shaped steel punch & die, in a hydraulic press. The dimensions of the corresponding shaping die are shown in Figure 26.

Following the removal of the hydraulic pressure, each material displayed varied levels of springback (Figure 26b). This pre-formed workpiece was then fixtured between a tool steel die and a polyurethane punch (Figure 27a) bearing the same dimensions as the steel punch, for the

VFA shape calibration step. The punch was made by first sectioning a 25.4 mm (1”) radius 80 polyurethane rod and taping it to a 12.7 mm (0.5”) thick pad. Both the rod and pad have a shore hardness of 80A.

Figure 26. (a) Shape and dimensions of the flange die, and (b) springback in CP Titanium compared to the target punch shape

VFA Calibration Step

A 0.127 mm (0.005”) thick patch foil, previously introduced in Figure 6b, was used for removing the elastic springback from the workpieces. The bottom up stacking order of the assembly consists of the foil, polyurethane punch, pre-formed specimen, and tool steel die aligned together. A top view of the alignment between the foil and punch can be seen in Figure

27b. This assembly is held tightly between the top block and anvil of the VFA fixture, shown in

Figure 28.

12.7 mm (0.5”)

(a) (b)

Figure 27. (a) Insulated polyurethane punch, and (b) top view of patch foil and shaping punch alignment in the VFA calibration step

Figure 28. Detailed VFA calibration assembly

The copper terminals of the fixture were then connected to a capacitor bank. A current is discharged from the capacitor bank into the foil through the terminals. The resistive joule heating from the high current rapidly heats the actuator above its sublimation, converting the metal foil into a high-energy gas. At a critical point, the high pressure from the gas becomes a shockwave as it travels through the punch, into the workpiece, finally attenuating within the die/top block.

The passage of the shockwave results in the net shape calibration of the pre-formed samples.

Following the VFA step, after unclamping the fixture from capacitor bank, the samples could be removed by hand without the need of special protective gloves or equipment.

Data Collection

Top-view images of each flange edge (Figure 26b) were taken using an Olympus digital camera mounted on a vertical stand. The base of the samples was aligned and taped in place to a precision flat steel plate. The images were post processed in ImageJ. The angles of the flange w.r.t the horizontal (θ1 and θ2), height (h), and width (w) of the curved section (Figure 29) were measured and recorded for each sample preformed and calibrated sample.

Figure 29. Regions of measurement for a pre-formed and calibrated workpiece

Two of the six materials with mid-range of strength to stiffness ratio (CP Ti and

AA6061-T6) were sectioned, mounted, and polished to a 1 μm surface finish. Microhardness measurements were taken at 0.1 μm spacings along the horizontal direction in the curved and the flange sections of each sample. The results from the calibration experiments are presented and explained in detail in the following section.

3.3 Results and Discussion

The radius of the curved sections (r) was calculated for each sample edge using the radius of the arc formula below, where h is the height of the arc taken at the midpoint of the chord of width, w.

ℎ2 + �2 � = 2�

The radii were averaged for each material type over five samples. The average value was then compared to the target value of 1.60 mm (0.063") to obtain the percentage springback in radius. The left and right flange angles of each edge were averaged as well. This data is ordered based on materials which exhibited the highest to lowest σy/E ratios, as shown in Figure 30.

All materials showed higher elastic strains in curvature compared to the flange angle as only the curved region undergoes most of the plastic deformation during the "drawing" motion of the punch into the die during the quasi-static process.

The springback in curvature height influences the angle of the flange as demonstrated by the pre-formed Ti 6242. Despite having lower σy/E values, the springback in curvature seen in the nickel superalloys was equal to or on par with the CP Ti and aluminum alloy. The high strain hardening rate of these superalloys [29] would have increased the σy/E ratio over time due to the increase in strength, causing more springback. The smaller gauge thickness probably also played a role. Due to the complexities involved in springback prediction, further analysis will only be restricted to the reduction or removal of the elastic springback itself.

Almost all the metals experienced near-complete elimination in curvature and angular springback following the VFA calibration step. As these values are averaged over both edges of all five samples, some variation is expected. The post-VFA flange angles for CP Ti, AA6061-T6, and all three nickel-based superalloys were very close to 0o. The histograms in the chart in

Figure 30a have been slightly exaggerated to show the presence of actual datapoints. The visual depictions in Figure 30b support the values on the chart. The calibrated samples had a clean surface finish, displayed no evidence of cracking, tearing, or heating. The lack of variation in curvature radius and flange angle of the pre-formed samples before and after they were clamped within the fixture (prior to foil vaporization) suggests that no further deformation was introduced due to the clamping force of the VFA fixture.

60% Radius of Curvature - Pre Radius of Curvature - Post 50% Flange Angle - Pre Flange Angle - Post 40%

30%

20% Springback Springback %

10%

0% Ti 6242 AA6061 - T6 CP Ti 718 625 230 (a)

(b)

Figure 30. Comparison of average springback % in the curved and flat sections of the workpieces in the pre-forming stage (pre) vs. VFA calibration stage (post)

Both Ti 6242 and CP Ti also had some post-VFA springback remnant in the curvature, and flange in the case of Ti 6242. The alloying and processing of Ti 6242 modifies its nature compared to CP Ti, making it 2.73 times stronger. Therefore, there is a notable disparity in their shape conformance. The dramatic change in the shape of the Ti 6242 sample seen in Figure 30b is a strong indication of the effectiveness of the VFA method. The urethane pad thickness was

91 reduced from 12.7mm (0.5") to 3.2 mm (0.125") for the Ti 6242 samples. It was found that the higher thickness was causing a springforward motion in the flanges, which would have been a result of residual strains introduced by the pad. The role of the urethane material is to "shock-up” the incoming wave and keep it from early attenuation as has been seen in many other cases. Its nature seems to have had too much of an influence on the thin gauge Ti 6242. By reducing the thickness of the pad hanging over the flange section, the effect seems to have been nullified. In contrast, when an impact-resistant tool steel punch was used instead of the urethane punch, the reduction in springback was significantly lower for all the materials. It is possible that the steel was causing faster shock attenuation than the urethane resulting in a lower calibration pressure.

Figure 31. Double-patch foil configuration used by Iriondo [82] to calibrate AA6061-T6 and CP Ti

It should be noted that the data for CP Ti and AA6061-T6 was obtained from some previously unpublished work conducted by Iriondo [82] in which a double patch foil configuration, shown in Figure 31 was used. This current work was conducted after the original study and took advantage of recent improvements in foil design at the Impulse Research

Laboratory. The double patch foil behaves like a parallel circuit in which the vaporizing energy 92 is split between both foils. The pressure distribution would be wider, but overall, much lower in strength and uneven due to the overlap region. In contrast, the use of a thicker single foil accesses higher foil vaporization pressures at the same energy levels. It offers an even pressure distribution that majorly targets the curved section, where most of the calibration is needed. As opposed to the double-patch configuration, the thicker single patch yielded the best conformance results for the thinner materials.

The aluminum alloy samples were received in the aged stated while all the other metals were in the solution annealed state. As conventional automotive stamping methods find it difficult to plastically deform AA6061 in the T6 condition due to tearing or high springback levels, the material is typically formed in the T4 condition and heat treated to T6. According to the σy/E ranking in Table 12, AA6061-T6 should have the second highest springback %. Along with the Ni - 625 and 230 alloys, which displayed more springback in the pre-formed state than predicted due to their high strain hardening rates, the AA6061- T6 alloys showed full shape calibration after the VFA step. Owing to its low yield strength, it was also the material that required the lowest amount of calibration energy (6 kJ) compared to the other materials that ranged at 10 kJ.

Out of the three superalloys, the Ni - 718 alloy showed the most residual springback in curvature after the VFA step. This curvature is likely removable through process optimization probably remove. As these metals were being calibrated for the first time, the parameters for the

VFA experiments were obtained by trial and error. The nickel superalloys were treated the same due to similarities in their thicknesses and strength to stiffness ratios. The strengthening mechanism of the 718 alloys seems to have influenced the springback in the pre-formed sample, which would have required a slightly higher pressure for full conformance.

Microhardness measurements were taken for an aged and an annealed material that fell into the mid-range of strength to stiffness ratios. The traverses taken across the thicknesses of the pre-formed and calibrated AA6061-T6 and CP-Ti specimens, along with the sampling locations are shown Figure 32 below. There was minimal variation in the hardness levels between the pre- formed and VFA processed states. The average difference was 2% for the AA6061-T6 flange sections and 1 % in the curved section. In the case of CP Ti, the variation was slightly higher

(7% flange and 6% curve) but still insignificant compared to typical shock hardening or thermal softening values [5].

Figure 32. Microhardness traverses across flanged and curved section of pre-formed and VFA calibrated samples of (a) AA6061-T6, and (b) CP Ti

Based on the accuracy of the shape conformance in Figure 30 and the trends in Figure 32 the following observations can be made: 1) The calibration effect is independent of the material, its alloyed nature, the underlying strengthening mechanism, the crystal structure, etc. All of them

94 were calibrated to the required shape and the only differences were fixturing related while the general working principle remained the same. 2) The absence of drastic spikes or depreciation in hardness suggests that thermal effects were negligible or non-existent during the process.

Therefore, it can be assumed that an athermal mechanism is responsible for the removal of springback from these materials.

Studies conducted on the calibration of high strength steels using electromagnetic and electrohydraulic methods [6], [83], [84] in the past suggest that the stress waves generated from these processes resulted in the relaxation of springback within the workpieces. Based on the high flyer velocities for VFA welding process [1] and the intermediate pressures reported for low foil vaporizing energies in the VFA metalworking process [11], the probability of achieving the critical conditions for the formation of shockwaves that exceeds the elastic limit of the material is high. It is hypothesized that the shape conformance seen in the VFA calibration process is caused by the shockwaves generated from the rapid vaporization of the foils. These shockwaves were transferred into the workpiece via the urethane medium and relieved the residual stresses within the pre-formed material. This concept will be further explored in Chapter 4.

3.4 Summary and Conclusions

The VFA method has been proven to be an effective way to calibrate high strength materials to a net shaped die. It was used to eliminate springback in shape for six materials:

AA6061-T6, CP Ti, Nickel 718, Nickel 625, and Ni 230 of varying thicknesses. The aluminum workpiece required 6kJ of input energy for full calibration while it took 10kJ input energy for calibrating the nickel and titanium sheets.

Variation in hardness trends between the pre-formed and VFA processed samples were negligible. The lack of spikes in hardness suggests the absence of shock hardening likely 95 indicated a medium - low pressure incoming shockwave. Additionally, no adverse softening suggests an athermal mechanism is responsible for the calibration process.

As opposed to a steel punch, a urethane based punch acted as a medium to increase the strength of the incoming pressure pulse thereby reducing the foil vaporizing energy requirements.

The high-pressure pulse, generated from the rapid vaporization of the foils, was transferred into the workpiece via the urethane medium. It is hypothesized that this shockwave caused the removal of residual stresses in the workpiece and nearly eliminated springback.

While the mechanism of springback relief is still being verified (Chapter 4), it is clear that the VFA method offers many advantages for shape calibration. It does not require heat cycles for removing springback. The process eliminates the need for designing dies for countering springback. The agile VFA equipment requires less real estate than a mechanical press. Unlike electromagnetic forming, this process does not require conductive workpieces for being effective. The VFA springback relief method has potential for net-shaping low volume components made of high strength materials.

Chapter 4. Athermal Springback Relief: Investigation of Governing Mechanism

4.1 Background and Motivation

During metalworking operations, different regions within the metal experience different strain levels due to 1) discrepancy in strength between the phases in the alloys, 2) uneven distribution of strain caused by the die shape, and 3) temperature gradient. If these pockets of strains remain after the removal of the external force, they lead to internal or residual stresses in the metal [85]. Residual stresses can occur due to different processing steps, but they are especially common when localized heating is involved. For example, in fusion welding, the area surrounding the weld (HAZ) experiences a very sharp rise in temperature, which has a dissipating effect on the surrounding area. Due to uneven cooling rates, regions of dissimilar temperature gradients experience different changes in dimensions which lead to residual stresses in the HAZ and the region surrounding it.

Residual stresses are commonly undesirable because they negatively affect corrosion resistance, strength, and fatigue life. They often occur during metal forming operations and are the prime cause of the springback phenomenon. Springback is the tendency of a material elastically recover to its original shape after an external load is removed. The degree to which the metal springs back is proportional to the level of yielding and hence the amount of residual stresses remaining in the material. The causes and prevention of springback have been widely researched over the years, but models are often material dependent and not universal [86]. Since they are unavoidable in metal forming processes, it is necessary to eliminate or reduce them.

Over the years, metal forming and joining processes have relied on thermally activated

97 relaxation. The underlying mechanism of these processes is relief of elastic strain through dislocation motion activated by high temperatures (solution annealing temperature) which typically leads to an undesirable decrease in material strength [87]. Additionally, thermal processing is expensive, time consuming, and not always applicable. For example, precipitation hardened materials such as AA6061 are often shaped at a lower T4 temper and then heat treated up to a T6 temper. Hence, there is a need for athermal stress relaxation methods that can achieve similar if not better results.

4.2 Literature Review of Impulse Based Athermal Springback Removal Methods

Three impulse based methods: explosive, electromagnetic, and electrohydraulic, have been implemented in the past to shape metal alloys and remove residual stresses. The works specifically dedicated towards stress relief are reviewed here.

Explosive Relief

Explosives were the earliest and most common sources of high-energy pulses. The idea of using explosive methods to generate shockwaves for residual stresses relief can be traced back to the 1970’s, and it is possible to use some of this groundwork to understand the athermal relief mechanism in shock-induced calibration.

Petushkov [88] studied the effect of explosive treatment on welded joints and proposed a general residual stress relief mechanism based on the superposition of shockwaves and residual stresses leading to an overall decrease in stresses. Schmidt and Shockey [89] used explosive treatments on butt-welded joints made of A36 steel and used a hole-drilling method to measure residual stresses before and after the treatment. An overall decrease of longitudinal tensile residual stresses and an uneven change in transverse residual stresses was observed. Pruemmer

[90] studied the theory of superposition of tensile and compressive shockwaves and residual stress states by comparing a simple model to experimental results. The model predicted that the chosen shock loading condition would lead to decrease in local compressive stresses from 200

MPa to 50 MPa. This theory was tested using a 20 mm thick layer of carbonite explosive was detonated parallel to the butt-weld line of an ST-37 (ASTM A36) steel joint (200x100x20 mm).

XRD stress measurements were taken for the {211} lattice planes across the weld line starting at the center before and after the explosive treatment. Prior to the shock treatment, the highest residual stress value of ~250 MPa was found at 20 mm from the center of the weld line, and it decreased by 50 MPa after the shock treatment. A trend of overall residual stress reduction was observed along the measurement line, and the values were well within the prediction limits. The model made overly simplified assumptions of only one type of shockwave and residual stress interaction. Also, the XRD stress measurements were only conducted along one line across the weld, there is no information about the stresses along the weld line, on the other side, or inside the weld interface. Though this experiment is a good attempt at backing this theory, more information is needed. Pruemmer and Petushkov’s work was used by several Russian researchers, however, all the published works are in a different language and hence difficult to review.

Since its proposal, some work has gone into applying explosive treatments on low carbon steel welded structures in an effort to improve fatigue life and resistance to SCC [91], [92]. But the process has not taken off compared to thermal or vibratory stress relief methods. This is possibly due to the extra level of safety measure required. Additionally, very little numerical simulation work [92], [93] can be found in this area and even though the work states that the

99 predictions match the measurements, the high strain rate models used in the simulations are still not accurate with respect to the actual behavior of the materials.

Electromagnetic Relief

Electromagnetic forming (EMF) is a non-contact impulse or high-speed forming process.

It uses a pulsed magnetic field to apply Lorentz forces that calibrate conductive materials. A very detailed explanation of the process of electromagnetic forming and joining and a description of different technologies can be found in a review by Psyk and Tekkaya [2]. Though this method boasts excellent shape calibration, the actual mechanism of springback elimination is still unclear. In the past 15 years, Golovashchenko [83] was the first to propose that the elastic wave generated during electromagnetic forming propagates multiple times within a strained workpiece and relieves residual stresses in the process. Golovashchenko tested this theory for springback relief in automotive materials such as aluminum (AA6111-T4) and high strength steels (BH210,

DP500, DP600). Flat metal specimens are first plastically deformed into a U-shape around a mandrel as shown in

Figure 33(a). After the bending force is removed, the elastic portion of the plastic deformation (residual stress) drives the strip to springback to its flat state. This bent strip is forcefully flattened by a steel plate; however, upon removal of the external force, the strip will try to revert to its bent state because of residual elastic stresses introduced into the strip. When a sample in this state of bending is subject to EMF through varied voltages and a number of discharges, the sample’s shape moves toward the original flat state, shown in

Figure 33(b). Since the only difference between the original flat sample and the final curved sample is the presence of residual stresses due to processing, it is safe to assume that the EMF

100 waves discharged into the sample acted to remove the residual stresses. Similar observations were made when higher strength steels were tested.

(a) (b)

Figure 33. (a) Specimens bent around a mandrel and then flattened with a steel plate (b) Specimens flatted using EM pulses of different voltages (Borrowed from Golovashchenko [83])

The EMF process was further tested [94] with aluminum alloys AA6111-T4 and AA5754 and for the steels BH210, DP500, and DP600 for a more complicated U-channel shape.

Unfortunately, no actual data nor any evidence of numerical simulations were presented for this shape, so it is difficult to comment on the underlying mechanism and the amount of springback elimination.

No further development of probable mechanism theories were proposed until Iriondo et al., [6]. In this paper, three theories for non-thermal springback relief mechanisms in EMF were proposed. The first theory is that the pressure directly caused by the electromagnetic pulse is briefly localized in some regions of the workpiece where it exceeds the yield strength of the material such that the stored elastic strains are eliminated causing plastic deformation. The second theory relies on the impact generated by the high collision speed that exceeds the yield strength and leads to plastic deformation. The third theory considers the joule heating or high current density induced decreases in flow stress aka the electro-plastic effect. Interestingly,

101

Golovashchenko’s [83], [94] prior work was not explicitly considered while it is most likely the mechanism at play in the first theory. Also, springback relief has been observed in EMF processes that do not use impact processing so the likelihood of the second theory being a dominating effect is low.

A lot of work has been done in advancing the EMF technology, but unfortunately, the role of the EMF pulse in removing the strain fields and hence aiding in the calibration is not addressed. It is briefly mentioned as an underlying mechanism but never explored as in-depth as the magnetoplastic effect mentioned earlier through modeling or experiments.

Electrohydraulic Relief

The electrohydraulic (EH) effect is a complex phenomenon that occurs when a high voltage is discharged through electrodes into a liquid medium [95]. This discharge creates a fast- expanding high-pressure plasma channel that leads to the production of shockwaves. The technology that leverages this effect has been researched and applied in the metal forming industry since the early 1950’s [96], [97]. The increase in material formability [98] and excellent shape calibration [84], [95] make it a very appealing sheet-metal forming process.

Golovashchenko [84] was the first to propose the idea of springback calibration through shockwave stress relief demonstrated using the electrohydraulic method. The process can be viewed as the canceling of stresses resulting from the interaction between normal, through thickness, compressive stresses from the shockwaves and the existing tensile residual stresses.

This method of stress relief is advantageous because it eliminates the introduction of new strains and stresses from any bounce back effects, so it is important to make sure that there is very little to no gap between the blank and the die.

102

The mechanism of the pulsed EH effect was studied by flattening pre-strained metal samples. Residual stresses were introduced into blanks made of 6111-T4 Al (0.9 mm thick) by deforming them into a U-Shape that caused springback upon removal. The deformed blanks were clamped into their target shapes using a clamping tool. When shockwaves created through electrohydraulic pulses are directed at the clamped blanks, the additional elastic strains created through the clamping are converted into permanent plastic strains leading to shape calibration.

(b (a) )

Golovashchenko et. al, [84] further demonstrated the relief effect through electrohydraulic flattening tests, similar to the electromagnetic flattening tests that had been conducted previously [2005]. The purpose of the EH flattening experiments was to quantitatively measure the change in springback in a simple and inexpensive manner. This experiment tested the pulse calibration effect on initially flat strips of 305 x 51 mm DP980 sheets of two different thicknesses (1 mm and 1.4 mm) that were bent to an angle of 90o against tools of two different

103 bend radii (12.7mm and 25.4mm). Each bent sample was then compressed under a 1 MN hydraulic press. This compression step 1) shows the amount of springback that can be removed by quasi-static methods, and 2) reduces the required energy for pulsed calibration. After the quasi-static flattening step, the sample was clamped flat between a steel plate and an EH chamber, with the convex surface facing down towards the water-filled cavity (Figure 34a).

Because of the way the samples are clamped in the EH step, it is easier to decouple the impact effects from the impulse effects. A pulse of 4 kJ resulted in maximum flattening in samples of both thicknesses. Pulses beyond this energy level lead to a spring-forward motion because it is possible that new residual stresses were introduced into the surface layer. A general trend of decrease in springback with an increase in energy was seen for samples of both thicknesses and radii. It is interesting to note that even though a higher initial springback angle (by ~4o) was observed in the thinner sample when bent against the tool, the energy that leads to maximum flattening was very close to that for the thicker sample (Figure 34b).

Golovashchenko [84] proposed that the stress relief in both EM and EH cases was a consequence of the high-pressure waves traveling through the thickness of the workpiece, eliminating the residual stresses in their path.

Some preliminary trials (Figure 35) for the same EM and EH flattening experimental setup described above, were conducted using the VFA method on CP Ti, Ti 6242, Ni – 230, Ni –

625, Ni – 718, and boron steel. The results not only showed that the calibration effect is independent of the material composition but also that it was independent of the high-pressure delivery method. If the mechanism commonalities between the VFA method and the EH method are considered, it can be justified that a shockwave based residual stress relief phenomenon is taking place.

104

Figure 35. Curved to flat stress relief demonstrations conducted on CP Titanium, High Strength Boron Steel, and Inconel 625 using the VFA process

There are a few gaps in the research on the shock-based calibration mechanism that the work in this chapter seeks to expand on. Firstly, it is unclear if the orientation of workpiece w.r.t the incoming shockwave affects the level of springback relief. Secondly, a proper verification of the mechanism would require the measurement of the shock pressure causing the stress relief event. Finally, it is important to understand the actual criterion for plastic deformation based on the boundary conditions of the system. If it is in fact a shock-based relief mechanism, based on previous research [5], the workpiece could be assumed to be in a state of uniaxial strain during shock loading. The yielding factor would be determined from analyzing this condition.

The work described in this chapter seeks to use the VFA method to expand on the flattening experiments conducted by Golvashchenko [83], [84] and provide satisfactory answers to the research gaps mentioned above. The following sections describe the experimental methodologies and data analysis for the PDV + VFA flattening experiments. Some insight into the sensitivity of the calibration with the orientation of the sample is also given. A brief analysis of the thickness of the sample vs. the reduction in springback and the likely shock pressures is also provided. And finally, the results are compared against the uniaxial yield criterion and analyzed accordingly.

105

4.3 Materials and Methods

Materials

Initial trials of the VFA based flattening experiments showed that the calibration ability is widely applicable to different high strength materials (Figure 35). As the investigation of the calibration mechanism relies on capturing velocimetry information and relating it to thermodynamic shockwave parameters, it is important to isolate the study to metals that generate information capable of being captured by the existing

Photon Doppler Velocimetry (PDV) system at the research facility.

Shock breakout capturing experiments conducted with AA3003, Cu 1100, mild steel, 304 stainless steel, and JSC 1500 steel revealed that AA3003 had the best breakout resolution. However, as AA3003 has a low strength to stiffness ratio, it is not the best material to for testing the calibration method.

Tests conducted with a higher strength AA7075-T6 alloy showed not only its capability to demonstrate pronounced springback relief but also provide reasonable signal reflectivity for the data collection. Additionally, the shock parameters of this material are readily available through the Los Alamos National Laboratory (LANL) shock database.

Figure 36 shows the linear dependence of particle velocity with shock pressures under 5

GPa.

Therefore, two different thickness of AA7075-T6 were used to study shock-based calibration mechanism. 42 x 101.6 x 1 mm (1.65” x 4” x 0.04”) and 42 x 101.6 x 2.23 mm (1.65” x 4” x 0.09”) blanks were mechanically sheared.

106

Figure 36. Variation of shock pressure with particle velocity for AA7075 [99]

Experimental Method

Sample Pre-straining

Due to the lack of access to a press-brake, the blanks were prepared for VFA calibration by using the air-bending setup illustrated in Figure 37 instead. Each blank was carefully bent between a semi-circular punch and die with a 25.4 mm (1”) radius. The specimens were quasi-statically deformed in the hydraulic press until a 90o bend was reached between the overhanging ends of the workpiece. Once removed from the semi- circular punch and die, the samples were placed between two precision flat steel blocks and flattened at a force of 9.1 Tons (10 US Tons).

107

Figure 37. (a) Sheet-metal air-bending operation and (b) method of measuring final springback angle in pre-bent specimens

For accurate tracking of the evolution in flatness, the bend angle (θ) was measured for each sample after the quasi-static bending and quasi-static flattening steps.

The side of each specimen was traced in graphing paper. Tangents drawn along the edges, against the curved region, were used to locate the center of curvature of the bend.

With this point and the two tangents as reference, a protractor was used to measure the springback in the bending angle (Figure 37b).

VFA - Calibration

The general bottom up stacking order of the shock-based calibration step (similar to Figure 38d) is as follows: insulated foil, 12.7 mm (0.5”) thick polyurethane pad, pre- strained sample, and flat backing block (without a PDV port). To ensure a symmetrical pressure distribution, the stack is prepared separately by carefully aligning each component along the vertical and horizontal centerlines (Figure 38a), and then clamping

108 the assembled stack within the fixture. The calibration mechanism was tested in two different configurations, one where the pre-formed sample is placed with the concave side up (CU) and one with the concave side down (CD) as shown in Figure 38b and

Figure 38c.

As a majority of the elastic residual stresses are concentrated in the bent region of the workpiece, a foil with an active length and area matching these dimensions was used for the flattening experiments. Starting from 3 kJ, the foil vaporizing energy was either incrementally increased or decreased by 0.5 - 2 kJ until a full or maximum reduction in curvature was obtained. The thickness of the foil was varied from 0.025 – 0.127 mm 109

(0.001” – 0.005”) depending on the energy used. Thinner foils are more efficient around

1- 4 kJ. The parameters of the VFA flattening experiments are listed in Table 13 below.

The bend angle (θ) for each sample was measured after the VFA flattening step as described previously.

Sample Sample Foil Thickness Energy PDV? Thickness Orientation (mm) (kJ) 0.076 (0.003”) 5 Yes 0.076 (0.003”) 7 No Concave Up 0.076 (0.003”) 8.5 No 1 mm 0.127 (0.005”) 9 No 0.076 (0.003”) 6 No Concave Down 0.127 (0.005”) 10 No 0.127 (0.005”) 14 No 0.025 (0.001”) 1 Yes 0.051 (0.002”) 3 Yes Concave Up 0.051 (0.002”) 4.5 Yes 0.076 (0.003”) 5 No 2 mm 0.051 (0.002”) 4.5 Yes 0.076 (0.003”) 7 Yes Concave Down 0.076 (0.003”) 8 Yes 0.127 (0.005”) 12 Yes 0.127 (0.005”) 16 Yes Table 13. PDV + VFA flattening experimentation schedule

110

Upper Bound Pressure Estimate with PIPE

Figure 39. (a) VFA Calibration + PIPE setup for pressure estimation and (b) top view of AA6061 transducer sample after being indented with a file due to an 8kJ foil vaporization energy

Brune and Hansen [100] calibrated three transducer materials, AA5052, AA6061-

T6, and SS304 and obtained a linear relationship between the file indentation width and the applied quasi-static pressure. The same method, known as the Profile Indentation

Pressure Estimation or PIPE, was applied in conjunction with the VFA calibration setup

111 to obtain a lower estimate on the pressures experienced by the workpiece due to the inbound shockwave.

The bottom-up stacking arrangement of this PIPE + VFA calibration setup

(Figure 39a) is described as follows. The same foil shown in Figure 38, was placed below a 35 x 31.75x 12.7 mm (1.7"x 1.25"x 0.5") thick polyurethane pad with an 80A shore hardness. The file is aligned and placed above the polyurethane pad. Finally, an AA6061-

T6 transducer workpiece of matching dimensions is placed on top of the foil with the reflective side facing down.

Three separate experiments were conducted at 4.5 kJ, 6 kJ, and 8 kJ. The distance between each indentation was measured using a microscope across each sample, as shown by the zoomed in image in Figure 39b. These indentation widths were related to pressure using the approximation:

Pressure|AA6061 = 1.27 x indentation width

Lower Bound Pressure Estimate with PDV

Measuring Shock Breakout Velocity

The measurement of the shock pressures experienced by the workpiece during the

VFA step are crucial for understanding the mechanism of springback relief. As of date, the direct measurement of these transient pressures is difficult. However, it is possible to calculate the shock pressure values using a combination of free surface velocity measurements and known Rankine-Hugoniot conditions for the materials of interest.

112

Free surface velocities of the work pieces undergoing shock processing were obtained using a Photon Doppler Velocimetry system (PDV) built at the Impulse

Manufacturing Lab.

Figure 40. Illustrated principle of Photon Doppler Velocimetry (PDV) [101]

A PDV system is a device capable of measuring the velocity of a moving object.

It consists of a laser source, detector, two interferometers and a probe, illustrated in

Figure 40 . A laser beam emitted from the source passes through the interferometer and splits into two beams. One beam is directed to a high-speed detector and the other beam is focused through a probe onto the target surface which has a reflective surface finish.

As the surface moves, it scatters the incoming light waves which are directed back through the probe into the detector. The detector records the raw intensity data of both the initial unscattered beam and the scattered beam. The data is post processed using the

113 spectrogram function in MATLAB which uses Fast Fourier Transformations to separate the shift in frequency and phase between the two superposed beams.

Separate VFA flattening experiments coupled with a single channel Photon

Doppler Velocimetry (PDV) system, were conducted to obtain surface velocity information. These experiments are also listed in Table 13. As the setup is time consuming, PDV + VFA flattening experiments were mainly conducted at the energy levels where the change in angle was drastic in VFA only experiments.

The original flattening setup was slightly modified to include a top block with a

6.35 mm (0.25”) diameter that serves as a workpiece viewing aperture (Figure 38d). For best signal acquisition, the bent region of the workpiece was coarsely sanded with 600 grit paper in a crisscross pattern until a reflective surface was obtained.

The axial alignment of the foil, sample, and viewing port are crucial for data collection. Additionally, the careful positioning of the 250 mm laser focuser with top

(free) surface of the sample ensures that only the normal velocity of the surface is collected. A 10 GHz data collection rate yielded good shock breakout results in the post- processing. Once the raw PDV data was obtained for the different flattening energies, it was filtered and processed according to steps in the next section.

114

Data Processing

The detector of the PDV system records the time averaged intensity data of the top surface of the metal undergoing the VFA shock process. This raw data, pictured in

Figure 41a, consists of frequency information of both the original and the doppler shifted light beams. The frequency distribution of the top surface can be extracted from this raw data by using a Fast Fourier Transformation (FFT) operation. The built-in MATLAB spectrogram function was used for the post-processing operation. The syntax of the

MATLAB spectrogram function used is as follows: spectrogram (x, window, noverlap, f, fs, ’yaxis’), where x is the input data. The input arguments that yielded the best results are: window = 3000, noverlap = 1500, f = 3000, fs = 10E9.

(a) (b)

SBV Shock Breakout

Figure 41. (a) Raw intensity data of the shock processed free surface, captured via PDV, (b) FFT processed intensite data that depicts the profile of a shock breakout and the likely peak pressure

115

An example of the FFT processed frequency data is shown in Figure 41b. The velocity of the surface is a function of this signal frequency and the laser wavelength

(1550 nm) per equation.

Using the frequency (f) information at the peak break-out value, and the wavelength (λo) of the incoming beam (1550 nm) as inputs in the Doppler equation, the free surface velocity (v) can be calculated as:

�� = � � 2

Eq 1. Doppler equation for velocity based on wavelength and frequency

Figure 42. Adapated general shock-profile obtained using a Velocity Interferometer System for Any Reflector (VISAR) by TeiresiasSky/ CC BY-SA 4.0 (Desaturated from original)

116

By comparing the free surface frequency distribution in Figure 41b, obtained using the PDV method, against a generic shockwave profile obtained using the VISAR method in Figure 42, a shockwave can be distinguished from lower stress waves by the steep rise in velocity, known as a “shock breakout”, as indicated in Typically, in VISAR based shock loading experiments, the first “knee” in this profile indicates the transition from an elastic to an elastic plastic state, known as the Hugoniot Elastic Limit (HEL). As the HEL is a material property rigorously determined through specific Split Hopkinson experiments, it is very likely that the pressure at the “knee” for the springback calibration experiments is below that value [102]. Assuming that the free velocity profiles obtained for the VFA calibration experiments are correct, this first peak velocity will be referred to as the shock breakout velocity (SBV) for the remainder of the document.

Using the pressure vs. particle velocity relationship for AA7075-T6 provided in

Figure 36, the SBV (twice the particle velocity) can used to calculate the free surface shock breakout pressure (SBP), in the direction of shock propagation. The VFA flat shape calibration and free surface velocimetry results are presented and analyzed in the next section.

4.4 Results and Discussion

Flattening results

The variation springback in the workpiece with increasing foil vaporization energy is pictorially represented in Figure 43. The images of the side views of the workpieces displayed in Figure 43(a-d) have been cropped to show only the angled end.

117

These samples represent two different thicknesses of AA7075 - T6 that have been subject to an impulse treatment in either the concave up (CU) or the concave down (CD) configuration, with respect to the incoming pressure wave. An overall representation of the trends in springback angle with the input energy is provided in Figure 43e.

The initial angle in the workpiece introduced by the air-bending and quasi-static flattening steps are denoted by the acronyms “PB” - Pre-Bent and “PF” - Pre-flattened.

Each of the VFA treated samples are labeled with the processing energy which has the general form of “xx kJ”. The maximum processing energy is limited to the maximum discharge energy of the capacitor bank which is 16 kJ.

The “flatness angle”, defined here as the springback angle between one overhanging end of the workpiece and the horizontal axis (dashed yellow line), of each shock-processed specimen is targeted to 0o. This angle implies that a fully springback relieved flat shape, non-discernible from an as-received blank, has been achieved.

118

1 mm CU PB 1 mm CD PF PB PF 5 kJ 6 kJ 7 kJ 10 kJ 8.5 kJ 14 kJ 9 kJ (a) (b) 2 mm CU 2 mm CD PB PB

PF PF 7 kJ 1 kJ 8 kJ 3 kJ 12 kJ 4.5 kJ 16 kJ (c) (d)

(e)

The angles at 0 kJ are averaged over 5 specimens whereas those at the maximum springback removal condition are averaged over 2 specimens. However, it should be noted that each individual data point in between 0 kJ to the maximum energy for that case, represents only one experimental value. This was a consequence of restrictions in time and material availability. Therefore, only generalized qualitative observations have been made.

Based on Figure 43e it appears that the broad effect of increasing the input energy, and thus increasing the input pressure, is to reduce the angle of springback in all the four cases shown in Figure 43(a-d). There is a definite difference in the level of springback relief against thickness (1 mm vs 2 mm) and orientation of the workpiece (CU vs. CD) with respect to the incoming pressure wave. These observations are summarized and discussed below.

Thickness effects

For the same sample orientation, a thicker sample takes less energy to calibrate to a flat shape compared to the thinner sample (Figure 43 a & c).

Chronologically, the 2 mm thick workpieces were calibrated in the concave up configuration before the 1 mm thick samples and achieved a flat shape at 4.5 kJ. As the thinner workpiece displayed higher springback in the unprocessed condition (after quasi- static deformation but before shock treatment), it was intuitively estimated that an energy slightly higher than 4.5 kJ would be needed for calibration.

120

At a discharge energy level of 5 kJ in the concave up configuration, the thinner sample showed a 26o reduction in springback from its unprocessed state and which was not fully relieved to 0o until an energy of 9 kJ was tested.

This general trend of higher springback in thinner samples is absent in the concave down configurations. All the data points follow a roughly linear trend.

The influence of thickness contradicts the trends observed by Golovashchenko

[2014] who observed that at the same energy level the relief in springback is greater in the thinner sample of DP 980 electrohydraulically processed in the concave up configuration. However, it seems to agree for the concave down geometry.

At this point, the differences in observed thickness trends is unclear. It is possible that this is a result of experimental error and would need further experimentation to resolve. As impulse based springback relief is more dependent on the interactions between the tensile and compressive strains of the material and that of the incoming shockwave, it is possible that the sample thickness factor is not of much importance for determining the mechanism and only influences general processing parameters.

Therefore, the effect of orientation is of higher interest for understanding the mechanism of impulse based shape calibration.

Orientation Effects

For the same sample thickness, the ability to achieve a fully flat shape was only observed in the concave up orientation (Figure 43 a & b). Not only does the concave

121 down orientation take more pressure to show a reduction in curvature but also the effect seems to become marginally smaller with increasing energy.

As mentioned previously, the effect of orientation on residual stress elimination relies on the interaction between the strain states of the material and shockwave. This interaction is qualitatively explained with the figure below.

Tensile

Compressive

(a)

(Easy) (Hard)

(Hard) (Easy)

Shockwave Shockwave

(b) (c)

Consider the cross-section of the sheet metal blank undergoing an air-bending operation. During the operation, for the case of uniform loading, the stresses at the 122 centerline will be neutral. The outer fibers of the membrane are in a state tension and the inner fibers will be in a state of compression. If the elastic limit of the material is exceeded, plastic flow will ensue on both the outer and inner fibers. However, the centerline region will still be elastically strained. If the bending load is removed, the induced stresses will try to relieve themselves by straightening the workpiece. This changes the sign of the stress, leading to a compressive strain state on the outer fiber

(shown in green in Figure 44a) and a tensile strain state on the inner fiber (shown in red in Figure 44a). As the quasi-static flattening operation still showed residual springback

(Figure 43a), the strain state of the material is not much different from that illustrated in

Figure 44a.

Figure 44b shows the strain state of the pre-formed workpiece constrained in the concave up configuration and Figure 44c shows the state of the concave down configuration. An incoming shockwave will subject the material to a uniaxial strain [5],

[103] in the thickness direction (+z) that has the form:

0 0 0 (Eq.2) �� = [0 0 0 ] 0 0 −��

Only the deviatoric (plastic component) of the strain will cause deformation. The deviatoric component (εdev) can be calculated by removing the hydrostatic strain component (εhyd) from the total strain (εtot) as follows:

(Eq.3) �� = �� − �ℎ��

� + � + � −� � = � � � = � (Eq.4) ℎ�� 3 3

123

� � 0 0 3 � 0 � 0 (Eq.5) �� = 3 −2� 0 0 � [ 3 ]

Therefore, the plastic strain component of the shockwave will have a negative component in the +z direction and a positive component in the +x and +y as shown in

Figure 44 (b) and (c). As the shock compresses the workpiece, it stretched the material in the lateral directions. This can be considered analogous to the mechanism of springback relief during the bending under tension metal forming operation. For these operations it has been empirically determined that a 0.1% plastic strain in the plane of the sheet is sufficient in removing a majority of the residual stresses. Similarly, the deviatoric component of a shockwave can easily eliminate existing stresses as well, creating a uniform distribution of stresses across the thickness of the material.

The shockwave dissipates as it moves through the thickness of the material from the urethane facing surface to the steel anvil. In the case of the concave up orientation

(Figure 44b), the strong shock first encounters and relieves the opposite signed compressive stress in the plane of the sheet. The last area to interact with the weakened shock will be the tensile stress state on the anvil side, which is easy to relieve due to the same sign of the shock.

In the case of the concave down configuration (Figure 44c), the full strength of the shock first encounters tensile stresses on the urethane side. The compressive stresses on the anvil side will oppose the weakened shock’s ability to extend the material in the

124 plane. Therefore, the concave down geometry is significantly more difficult to calibrate than the concave up.

Uniaxial strain condition

The drastic shape change reported earlier in this section indicates that the workpiece experienced springback relief as a consequence of the impulse operation. As the shockwave in an impulse operation typically places the material in a state of uniaxial strain, and permanent deformation resulting from it would need to satisfy a minimum yielding criterion. This criterion was proposed by Jones and Graham [104] and the key idea extends as further back as Bancroft et al., [105].

The uniaxial strain condition for a shockwave propagating through a material is given by equation 11 (a detailed derivation can be found in Appendix C).

−� ��−�+� � = � ( ) (Eq 6) � � �−�

The corresponding pressure at the onset of yielding is provided by:

�−� � = � ( ) (Eq 7) �|�� −��

Figure 45 shows a pictorial representation of the coefficient of yield stress in the above equation against Poisson’s ratio (ν). For metals, ν typically ranges from 0.3 – 0.33.

Therefore, the uniaxial pressure must exceed the uniaxial yield stress by a factor of about

1.75 – 2 to meet the yielding condition during shockwave propagation.

125

Figure 45. Variation of Poission’s ratio with uniaxial strain yield factor

AA7075 - T6 has an approximate yield stress of 450 MPa, an ultimate strength of

520 MPa, and Poisson’s ratio of 0.33. Using pressure criterion listed above, the minimum pressure that promotes plastic deformation in the uniaxial condition is 886 MPa.

Bounding the Pressure

Velocimetry Data

Table 14 summarizes the variation in springback angle with input energy for 2 mm thick AA7075-T6 samples processed in both the CU and CD geometries. It provides the matching free surface SBVs, measured from the PDV + Flattening experiments, and their corresponding SBPs, obtained by interpolating the linear pressure vs. particle velocity relationship shown in Figure 36. The velocimetry data for the 4.5 kJ shot in the

CD geometry showed no visible shock break out, therefore this pressure value has been 126 left unpopulated. The actual FFT processed velocimetry profile for these cases can be found in Appendix D.

Sample Energy Flatness Springback Shockwave SBV SBP Orientation (kJ) Angle Reduction breakout? (m/s) (MPa) 1 10 60% Yes 44.94 324 Concave Up 3 3 88% Yes 66.45 474 4.5 0 100% Yes 77.8 561 4.5 -16 35% No 71.66 N/A 7 -13 47% Yes 92.47 666 Concave 8 -12 51% Yes 101.227 729 Down 12 -11 55% Yes 83.01 598 16 -7 72% Yes 65.04 469 Table 14. Summary of flatness angles, corresponding shock breakout velocities (SBV), and shock breakout pressures (SBP) for 2mm thick AA7075-T6 shock processed in the CU and CD configurations

As each SBV and SBP value in this table is a single experimental data point, only qualitative observations will be discussed here.

For the CU geometry, the breakout pressure follows the expected model and increases with input energy. The pressure at the maximum calibration energy of 4.5 kJ was 561

MPa.

In the CD configuration, the SBP first increases and then decreases. The pressure observed at the maximum calibration energy of 16 kJ is 469 MPa. This trend follows a new model that is yet unclear.

127

As previously discussed, the strength of the shockwave dissipates as it propagates through the thickness of the workpiece. Therefore, the actual shock pressure that causes the observed shape change is larger than the breakout pressures captured.

Considering only the energies that accomplished a drastic change in % flatness, 3 kJ and

4.5 kJ for the CU configuration and 16 kJ in the CD configuration, the average pressure exiting the workpieces was ~ 500 MPa. Therefore, this value could represent a lower bound in the actual strength of the initial shockwave.

PIPE Data

The PIPE method introduced by Brune [100] and furthered by Hansen [106] puts an upper bound on the pressure of the incoming shockwave. The pressure approximation is based on the hardness theory, proposed by Tabor [107], which suggests that there is a linear relationship between quasi-static plastic deformation and work hardening created by the indentation process and the uniaxial yield strength of the material.

Using the PIPE method, the estimated pressures range from 400 – 700 MPa at 4.5 kJ to 800 – 950 MPa at 8 kJ (Figure 46) for AA6061-T6 transducer material that was shock processed through a urethane medium using the VFA method.

Combining the findings from the VFA flattening experiments, the PDV breakout experiments and the PIPE experiments, for the experiments that saw a drastic change in curvature, it is possible that the pressure of the incoming shockwave could have ranged from 700 MPa to above 1 GPa. These sustained shock pressures would have been large

128 enough to meet the uniaxial strain yielding criterion, relieve the existing elastic residual stresses, and calibrate the workpiece to the shape it is currently being held to.

Figure 46. Spatial pressure distributions for AA6061-T6 VFA shock processed through a urethane medium at 4.5 kJ, 6 kJ, and 8 kJ.

The pressures estimated using the uniaxial strain criterion correspond to an as- received workpiece that has not been plastically strained through quasi-static deformation processes. However, the workpieces that subsequently underwent VFA shock processing had previously been deformed quasi-statically through bending and flattening operations.

Parts of the material would have already experienced some elastic stress relief. Therefore,

129 the actual pressure to cause full yielding of the workpiece would have been slightly different than that estimated by the uniaxial strain condition.

The pressure values estimated using the PDV method are not as reliable as they only represent one data point. Furthermore, the full profile of the processed velocimetry data is not entirely understood and would need to be studied further. It is believed that the

PIPE method provides a more accurate value of a lower bound on the pressure caused by the incoming shockwave as it has been tested rigorously for repeatability.

4.5 Summary and Conclusions

The mechanism of impulse-based shape calibration was investigated using the

VFA method and a theory for impulse based stress relief with respect to the sample orientation was proposed.

Workpieces were quasi-statically deformed to introduced elastic residual stresses into the plane of the material. The workpieces were then subject to VFA shock processing in a concave up (CU) and concave down (CD) geometry. The input foil vaporizing energy levels incremented until a fully flat specimen with no residual curvature was obtained. Two different thicknesses of AA7075-T6 were the subject of these studies.

A general decrease in curvature with increasing energy was observed in all four cases. It was harder to VFA flatten workpieces held in the CD configuration as opposed to the CU configuration.

130

It is speculated that this is due to the dissipative effect of the shockwave traveling through the material. It is preferable for it to first encounter compressive stresses (the case in CU) as these are harder to relieve and therefore need the full strength of the shockwave.

The thicker samples required lesser pressure for curvature relief in comparison to the thinner samples. The reason is unclear.

Free surface velocity measurements obtained using a photon doppler velocimetry system (PDV) for the 2 mm thick samples under VFA shock is speculated to have captured shock breakouts. It is hypothesized that the workpieces sustained pressures ranging from 800 MPa to above 1 GPa. Samples held in the CU configuration seem to comply with the uniaxial strain yielding criteria for this material. The breakout results obtained for the CD geometry are scattered.

Further experiments to capture higher resolution breakouts are needed. PVDF or

Magnanin pressure gauges are better suited for high pressure experiments. The velocity profiles currently captured with the PDV system have some peculiar features that are not yet understood. Both the experimental reproducibility and experimental measurement of

P(t) are currently tricky as the exact form of pressure pulse that is developed initially is still difficult to estimate. Furthermore, the attenuation of the shockwave through the material, and the effect of discontinuities at shock impedance interfaces have yet to be explored.

Regardless, it has been shown that the VFA method is a powerful and simple way to induce such pressures over much larger areas than laser shock peening or shot peening

131 can. By impulse loading the materials to pressure is 1.75 - 2 times the flow stress can induce significant plastic deformation at the shock front. The propagation of this shockwave can reduce elastic strain gradient significantly. This impulse process has several potential applications including residual stress removal in welds & additively manufactured components, forming components from net-shape dies, correcting shape in incrementally formed parts, and more.

132

Chapter 5. Future Work

The feasibility of impact welding and impulse shape calibration processes for nickel and titanium alloys has been demonstrated throughout the research presented in this dissertation. The studies conducted in both cases are unique and still need further investigation. This chapter provides a summary of unanswered questions and hopes to prompt further research.

Impact Welding

Compared to conventional welding methods, the VFA welding process is still quite novel. Continuous process improvements in foil design, fixturing, welding configurations, and bank efficiency are ongoing. Parallel work is also being conducted in understanding the relationship between the macroscopic weld properties against microscopic weld features such as the formation and implication of interfacial waves, and the composition and size of the mixed regions.

In the case of the similar patch welds, all the combinations displayed the properties of a sound joint, and any variation can be attributed to poor process control.

According to literature, for the case of like-paired joints, there seems to be much interest in welding precipitation hardening nickel alloys in an already aged state. Recent efforts

[42] have found it difficult to obtain joint strengths higher than the base metal strength.

133

Another aspect that needs to be explored further is the ideal process parameters for avoiding the formation of adiabatic shear bands (ASB) in the Ti 6242 – Ti 6242 case.

ASBs typically lower the strength of the joint, so it is ideal to avoid their formations.

As mentioned in the discussion section of Chapter 2, misalignment and arcing in the foil caused issues with repeatability in mechanical testing, especially for the dissimilar welds. In general, more data points are necessary to narrow down the variation in joint strength. As all these alloys are typically used at elevated temperatures, the endurance of the welds in the prospective service conditions would help determine the failure loads and necessary failure modes. Estimating the weld area is critical for comparing the strengths obtained to values reported in the literature.

The furthest work needs to be done with dissimilar Ni - Ti welding. Though the three combinations were successfully welded, the findings for one system could not be translated or related to the other. If only one system is studied, for example 718-Ti, it might help narrow down the research questions that need to be asked. These questions can then be tested for the other combinations to understand the trends or lack thereof in properties.

Another aspect of the dissimilar Ni - Ti welds was the stark difference in failure loads obtained when the titanium was used as a target instead of a flyer. Evidently, the titanium target provides an even shallower impact angle due to the springback in the pre- forming stage. However, it is unclear if the increase in strength was a result of changing the impact angle or because of better process control. One of the restrictions of using titanium as a flyer during the Ni - Ti patch weld trials was the height and steepness off

134 the standoff. It is possible that those welding trials might yield different results if titanium was used as a target instead. It would be interesting to know if switching the flyer metal for the dissimilar Ni - Ni welds also affects the strengths drastically. The impact factor of any improvements for these joints can be very significant for the aerospace industry.

Based on the literature review, it appears that joining nickel and titanium alloys to steel is of great interest to the aerospace and nuclear community. Studying the feasibility of directly joining these dissimilar materials would also be of interest.

Impulse Calibration

The significance of the work presented on the VFA impulse calibration method is vast. Development of the method could lead to application such as non-thermal residual stress relief in fusion welds, tailored pressure delivery for shape calibration, and more.

However, there remain some gaps in the analysis that need to be further explored.

The variation of the level of impulse calibration against the sample thickness is inconclusive.

Actual measurements of the pressures developed during the propagation and upon exit from the material surface are not understood. It might be possible to provide another

135 estimate of the pressure by using materials such as zirconium or SS304 that exhibit twin formation upon shock loading. As the minimum stress required to induce twin formation is known for each case, it is possible to estimate the sustained pressure. If the workpieces are thick enough, it is possible to use the same experiments to understand the attenuation of the shockwave through the material as the energy dissipates. PVDF and Magnanin gauges placed on both sides of the shock processed workpiece have the potential to measure incoming and outgoing pressures as well.

Another aspect that needs to be explored is the shock impedance caused by the discontinuities in the propagation path, namely the gaps of air that exist between the interfaces and the effect of the fixturing material properties on the strength of the wave.

Compared to the impact welding section, the path forward for the impulse calibration method has more clarity. The uncertainties described here can be more readily tackled and easily related to the previously obtained results.

136

Bibliography

[1] A. Vivek, G. A. Taber, J. R. Johnson, S. T. Woodward, and G. S. Daehn, “Electrically driven plasma via vaporization of metallic conductors: A tool for impulse metal working,” J. Mater. Process. Technol., vol. 213, no. 8, pp. 1311– 1326, 2013.

[2] V. Psyk, D. Risch, B. L. Kinsey, A. E. Tekkaya, and M. Kleiner, “Electromagnetic forming - A review,” J. Mater. Process. Technol., vol. 211, no. 5, pp. 787–829, 2011.

[3] T. Z. Blazynski, Explosive Welding , Forming and Compaction. 1983.

[4] S. H. Carpenter and R. H. Wittman, “Explosion Welding,” Annu. Rev. Mater. Sci., vol. 5, no. 1, pp. 177–199, 2003.

[5] E. Meyers, A.M, Lawrence, Shock Waves and High-Strain-Rate Phenomena in Metals. 1980.

[6] E. Iriondo, M. A. Gutiérrez, B. González, J. L. Alcaraz, and G. S. Daehn, “Electromagnetic impulse calibration of high strength sheet metal structures,” J. Mater. Process. Technol., vol. 211, no. 5, pp. 909–915, 2011.

[7] H. Wang and Y. Wang, “High-Velocity Impact Welding Process: A Review,” Metals (Basel)., vol. 9, no. 2, p. 144, 2019.

[8] A. Vivek, B. C. Liu, S. R. Hansen, and G. S. Daehn, “Accessing collision welding process window for titanium/copper welds with vaporizing foil actuators and grooved targets,” J. Mater. Process. Technol., vol. 214, no. 8, pp. 1583–1589, 2014.

[9] B. C. Liu, A. Palazotto, A. Vivek, and G. Daehn, “Impact welding of wrought and additively manufactured 15-5 PH stainless steel,” 2018.

[10] B. Liu, A. Vivek, and G. S. Daehn, “Joining sheet aluminum AA6061-T4 to cast magnesium AM60B by vaporizing foil actuator welding: Input energy, interface, and strength,” J. Manuf. Process., vol. 30, pp. 75–82, Dec. 2017.

[11] A. Vivek, R. C. Brune, S. R. Hansen, and G. S. Daehn, “Vaporizing foil actuator used for impulse forming and embossing of titanium and aluminum alloys,” J. Mater. Process. Technol., vol. 214, no. 4, pp. 865–875, 2014.

[12] A. Vivek and G. S. Daehn, “Vaporizing foil actuator: A versatile tool for high 137

energy-rate metal working,” Procedia Eng., vol. 81, no. October, pp. 2129–2134, 2014.

[13] B. C. Liu, A. N. Palazotto, A. Nassiri, A. Vivek, and G. S. Daehn, “Experimental and numerical investigation of interfacial microstructure in fully age-hardened 15- 5 PH stainless steel during impact welding,” J. Mater. Sci., 2019.

[14] X. Liu, S. Lan, and J. Ni, “Experimental study of Electro-Plastic Effect on Advanced High Strength Steels,” Mater. Sci. Eng. A, vol. 582, pp. 211–218, 2013.

[15] A. Kapil, T. Lee, A. Vivek, R. Cooper, E. Hetrick, and G. Daehn, “Spot impact welding of an age-hardening aluminum alloy: Process, structure and properties,” J. Manuf. Process., vol. 37, pp. 42–52, Jan. 2019.

[16] S. Su, S. Chen, Y. Mao, J. Xiao, A. Vivek, and G. Daehn, “Joining Aluminium Alloy 5A06 to Stainless Steel 321 by Vaporizing Foil Actuators Welding with an Interlayer,” Metals (Basel)., vol. 9, no. 1, p. 43, Jan. 2019.

[17] T. Lee, Y. Mao, R. Gerth, A. Vivek, and G. Daehn, “Civilized explosive welding: Impact welding of thick aluminum to steel plates without explosives,” J. Manuf. Process., vol. 36, pp. 550–556, Dec. 2018.

[18] M. Henderson and D. Arrell, “Nickel based superalloy welding practices for industrial gas turbine applications,” Sci. Technol. Weld. Join., vol. 9, no. 1, pp. 13– 21, 2004.

[19] S. D. K. John N. Dupont, John C. Lippold, Welding Metallurgy and Weldability of Nickel-Base Alloys. 2009.

[20] J. C. Williams and E. A. Starke, “Progress in structural materials for aerospace systems,” Acta Mater., vol. 51, no. 19, pp. 5775–5799, 2003.

[21] J. D. Beal, R. Boyer, D. Sanders, and T. B. Company, “Forming of Titanium and Titanium Alloys,” ASM Handb. Metalwork. Sheet Form., vol. 14B, pp. 656–669, 2006.

[22] R. R. Boyer, “Titanium for aerospace: Rationale and applications,” Adv. Perform. Mater., vol. 2, no. 4, pp. 349–368, 1995.

[23] J. P. Immarigeon, R. T. Holt, A. K. Koul, L. Zhao, W. Wallace, and J. C. Beddoes, “Lightweight materials for aircraft applications,” Mater. Charact., vol. 35, no. 1, pp. 41–67, 1995.

[24] A. Fuji, Y. Horiuchi, and K. Yamamoto, “Friction welding of pure titanium and pure nickel,” Sci. Technol. Weld. Join., vol. 10, no. 3, pp. 287–294, 2005.

138

[25] H. C. Chen, A. J. Pinkerton, and L. Li, “Fibre laser welding of dissimilar alloys of Ti-6Al-4V and Inconel 718 for aerospace applications,” Int. J. Adv. Manuf. Technol., vol. 52, no. 9–12, pp. 977–987, 2011.

[26] K. Topolski, P. Wieci??ski, Z. Szulc, A. Ga??ka, and H. Garbacz, “Progress in the characterization of explosively joined Ti/Ni bimetals,” Mater. Des., vol. 63, pp. 479–487, 2014.

[27] A. Vivek, S. Hansen, J. Benzing, M. He, and G. Daehn, “Impact Welding of Aluminum to Copper and Stainless Steel by Vaporizing Foil Actuator: Effect of Heat Treatment Cycles on Mechanical Properties and Microstructure,” Metall. Mater. Trans. A Phys. Metall. Mater. Sci., vol. 46, no. 10, pp. 4548–4558, Oct. 2015.

[28] S. R. Hansen, A. Vivek, and G. S. Daehn, “Impact Welding of Aluminum Alloys 6061 and 5052 by Vaporizing Foil Actuators: Heat-Affected Zone Size and Peel Strength,” J. Manuf. Sci. Eng., vol. 137, no. 5, p. 51013, Sep. 2015.

[29] J. R. Davis, ASM specialty handbook: nickel, cobalt, and their alloys. 2000.

[30] “Nickel-Chromium Phase Diagram,” 2000. [Online]. Available: http://www.calphad.com/pdf/Ni_Cr_Phase_Diagram.pdf.

[31] “Iron-Nickel Phase Diagram,” 2008. [Online]. Available: http://www.calphad.com/iron-nickel.html.

[32] T.B. Massalski -Binary alloy phase diagrams. V2.” .

[33] H. L. Eiselstein and D. J. Tillack, “The Invention and Definition of Alloy 625,” pp. 1–14, 1991.

[34] A. F. Guillermet and L. Östlund, “Experimental and theoretical study of the phase equilibria in the fe-ni-w system,” Metall. Mater. Trans. A, vol. 17, no. 10, pp. 1809–1823, 1986.

[35] D. L. Klarstrom, “Haynes 230,” 2009.

[36] L. M. Pike, “100+ Years of Wrought Alloy Development at Haynes International,” in 8th International Symposium on Superalloy 718 and Derivatives, 2014.

[37] M. Agilan, S. C. Krishna, S. K. Manwatkar, E. G. Vinayan, D. Sivakumar, and B. Pant, “Effect of Welding Processes (GTAW & EBW) and Solutionizing Temperature on Microfissuring Tendency in Inconel 718 Welds,” Mater. Sci. Forum, vol. 710, pp. 603–607, 2012.

[38] C. M.J., “The Welding and Solidification Metallurgy of Alloy 625,” Weld. J., pp. 139

49–56, 1987.

[39] C. Fink, D. Keil, and M. Zinke, “Evaluation of hot cracking susceptibility of nickel-based alloys by the PVR test,” Weld. World, vol. 56, no. 7–8, pp. 37–43, 2012.

[40] E. K. D. Passos, J. T. de Assis, V. I. Monine, R. S. Gonzaga, and J. da C. Payão Filho, “X-Ray Diffraction Technique for Residual Stress Measurement in NiCrMo Alloy Weld Metal,” Adv. Mater. Sci. Eng., vol. 2018, pp. 1–10, 2018.

[41] R. E. Mortland, J.E. , Evans, R.M., and Monroe, “WELDING AND BRAZING OF NICKEL-BASE ALLOYS,” 1972.

[42] D. Cornu et al., “Weldability of superalloys by Nd: YAG laser,” Weld. Int., vol. 9, no. 10, pp. 802–811, 1995.

[43] E. Lertora, C. Mandolfino, and C. Gambaro, “Mechanical Behaviour of Inconel 718 Thin-Walled Laser Welded Components for Aircraft Engines,” Int. J. Aerosp. Eng., vol. 2014, pp. 1–9, 2014.

[44] D. M. Janicki, “Fiber laser welding of nickel based superalloy Inconel 625,” Laser Technol. 2012 Appl. Lasers, vol. 8703, no. January 2013, p. 87030R, 2014.

[45] F. Caiazzo, V. Alfieri, F. Cardaropoli, and V. Sergi, “Investigation on edge joints of Inconel 625 sheets processed with laser welding,” Opt. Laser Technol., vol. 93, pp. 180–186, 2017.

[46] S. C. Ernst, “Weldability Studies of Haynes® 230 Alloy Boron has a potent detrimental effect on weld metal solidification cracking and microfissuring during multipass GTA and GMA welding.”

[47] A. . Lingenfelter, “The Welding metallurgy of Nickel Alloys in Gas Turbine Components,” Join. Repair Gas Turbine Components, 1997.

[48] J. Gordine, “Some Problems in Welding Inconel 718,” Weld. Res. Suppl., vol. November, p. 480–s, 1971.

[49] C. Mary and M. Jahazi, “Multi-scale analysis of IN-718 microstructure evolution during linear friction welding,” Adv. Eng. Mater., vol. 10, no. 6, pp. 573–578, 2008.

[50] K. H. Song and K. Nakata, “Microstructural and mechanical properties of friction- stir-welded and post-heat-treated Inconel 718 alloy,” J. Alloys Compd., vol. 505, no. 1, pp. 144–150, 2010.

[51] R. Damodaram, S. Ganesh Sundara Raman, and K. Prasad Rao, “Effect of post- 140

weld heat treatments on microstructure and mechanical properties of friction welded alloy 718 joints,” Mater. Des., vol. 53, pp. 954–961, 2014.

[52] B. H. Ding, R.J., Schneider, J., & Walker, “Advances in Solid State Joining of Haynes 230 High Temperature Alloy,” NASA Tech. Reports Serv., 2010.

[53] D. H. Williston, “Comparison of welding processes for Haynes 230 superalloy,” 2013.

[54] J. A. A. Schneider, D. Williston, T. L. L. Murphy, C. Varner, J. Hawkins, and B. Walker, “Solid state joining of nickel based alloy, Haynes 230,” J. Mater. Process. Technol., vol. 225, pp. 492–499, 2015.

[55] A. Chamanfar, M. Jahazi, and J. Cormier, “A Review on Inertia and Linear Friction Welding of Ni-Based Superalloys,” Metall. Mater. Trans. A Phys. Metall. Mater. Sci., vol. 46, no. 4, pp. 1639–1669, 2015.

[56] L. Williams, James C. and Gerd, Titanium. .

[57] “Titanium-Aluminum Phase Diagram,” 2008.

[58] M. J. Donachie, Titanium, 2nd ed. ASM International, 2000.

[59] A. E. Shapiro and T. Brazing, “Brazing of Conventional Titanium Alloys,” Welding, Brazing Solder., vol. 6, pp. 1–25, 1993.

[60] C. Jos, “Joining of γ -TiAl Alloy to Ni-Based Superalloy Using Ag-Cu Sputtered Coated Ti Brazing Filler Foil,” pp. 1–14, 2018.

[61] A. E. Shapiro and Y. A. Flom, “Brazing of Titanium at Temperatures below 800°C: Review and Prospective Applications,” http://titanium- brazing.com/publications/DVS-Manuscript_1020-Copy2-19-07.pdf, pp. 1–24, 2007.

[62] J. Sieniawski and M. Motyka, “Superplasticity in titanium alloys,” J. Achiev. Mater. Manuf. Eng., vol. 24, no. 1, pp. 123–130, 2007.

[63] F. Viana, A. S. Ramos, M. T. Vieira, M. F. Vieira, and S. Sim, “Reaction-assisted diffusion bonding of TiAl alloy to steel,” vol. 171, pp. 73–82, 2016.

[64] J. Seretsky and E. R. Ryba, “Laser Welding of Dissimilar Metals: Titanium To Nickel,” Weld. J., vol. 55, no. 7, p. 208s–211s, 1976.

[65] J. Aleman, B , Gutierrez, I, and Urcola, J, “Interface microstructures in the diffusion bonding of a titanium alloy Ti 6242 to an INCONEL 625,” Metall. Mater. Trans. A, vol. 26, no. 2, pp. 437–446, 1995. 141

[66] R. . Evans, “Joining of Nickel-Base Alloys,” 1962.

[67] L. Gorin, “Welding titanium alloys to nickel-base alloys,” Weld. Prod., 1964.

[68] C. E. Ells et al., “JOINING NICKEL TO TITANIUM,” pp. 2527–2534, 1973.

[69] M. Kuo, “Dissimilar Friction Welding of titanium alloys to Alloy 718,” The Ohio State University, 1993.

[70] B. Liu, “Joining Dissimilar Structural Alloys by Vaporizing Foil Actuator Welding: Process Conditions, Microstructure, Corrosion, and Strength,” 2016.

[71] J. B. Ribeiro, R. Mendes, and A. Loureiro, “Review of the weldability window concept and equations for explosive welding,” J. Phys. Conf. Ser., vol. 500, no. 5, p. 52038, May 2014.

[72] D. Jaramillo, A. Szecket, and O. T. Inal, “On the transition from a waveless to a wavy interface in explosive welding,” Mater. Sci. Eng., vol. 91, pp. 217–222, Jul. 1987.

[73] A. Mitchell, “Primary Carbides in Alloy 718,” Superalloy 718 Deriv., pp. 161– 167, 2012.

[74] R. M. Feng, J., Vieira, R., Thorne, R.J., Pelloux, “Localized Shear Band in Explosively Bonded Alloy 718/Copper/Alloy 718 Laminates,” in Superalloy 718 - Metallurgy and Applications, 1989, p. 43.

[75] T. Lee, S. Zhang, A. Vivek, G. Daehn, and B. Kinsey, “Wave formation in impact welding: Study of the Cu–Ti system,” CIRP Ann., vol. 68, no. 1, pp. 261–264, 2019.

[76] M. A. Meyers, G. Subhash, B. K. Kad, and L. Prasad, “Evolution of microstructure and shear-band formation in α-hcp titanium,” Mech. Mater., vol. 17, no. 2–3, pp. 175–193, Mar. 1994.

[77] J. Liu, Q. Fan, H. Cai, and F. Wang, “Underlying mechanism of periodical adiabatic shear bands generated in Ti–6Al–4V target by projectile impact,” Mater. Des., vol. 87, pp. 231–237, Dec. 2015.

[78] Y. Yang, Z. Xinming, L. Zhenghua, and L. Qingyun, “Adiabatic shear band on the titanium side in the Ti/mild steel explosive cladding interface,” Acta Mater., vol. 44, no. 2, pp. 561–565, Feb. 1996.

[79] R. H. Wagoner, J. F. Wang, M. Li, and T. Ohio, “Springback,” no. c.

[80] S. Woodward, C. Weddeling, G. Daehn, V. Psyk, B. Carson, and A. E. Tekkaya, 142

“Production of low-volume aviation components using disposable electromagnetic actuators,” J. Mater. Process. Technol., vol. 211, no. 5, pp. 886–895, 2011.

[81] A. K. Ghosh and C. H. Hamilton, “Superplastic Sheet Forming,” ASM Handb. vol.14B - Metalwork. Sheet Form., p. 345, 2004.

[82] B. Iriondo, Edurne., Nirudhoddi, “Unpublished work on Impulse Shape Calibration using VFA.”

[83] S. F. Golovashchenko, “Springback Calibration Using Pulsed Electromagnetic Field,” AIP Conf. Proc., vol. 284, no. May 2015, pp. 284–285, 2005.

[84] S. F. Golovashchenko, A. J. Gillard, A. V. Mamutov, and R. Ibrahim, “Pulsed electrohydraulic springback calibration of parts stamped from advanced high strength steel,” J. Mater. Process. Technol., vol. 214, no. 11, pp. 2796–2810, 2014.

[85] I. Toten, Howes, “Handbook of Residual Stress and Deformation of Steel,” Handb. Residual Stress Deform. Steel, vol. http://www, pp. 54–69, 2002.

[86] R. H. Wagoner, H. Lim, and M. G. Lee, “Advanced issues in springback,” Int. J. Plast., vol. 45, pp. 3–20, 2013.

[87] M. R. James, “Relaxation of Residual Stresses an Overview,” Adv. Surf. Treat., no. Vol 4, pp. 349–365, 1987.

[88] V. A. T. V.G. Petushkov, A.G. Bryzgalin, “Mechanism of Reslaxation of Residual Stresses in Explosion Loading of Welded Joints,” no. 614.

[89] C. G. Schmidt and D. A. Shockey, “Reduction of Residual Stresses in Weldments with Explosive Treatments,” Weld. Res. Suppl., no. December, pp. 4–7, 1992.

[90] R. Pruemmer, “Residual Stress Relief Treatment by Shock Waves,” Met., vol. 52, no. 10–11, pp. 633–634, 1998.

[91] J. L. Garcia-Jacomino, J. Burgos Sola, A. Cruz-Crespo, M. Alvarez Luna, and J. Garcia Arteaga, “Use of explosives in the reduction of residual stresses in the heated zone of welded joints,” Weld. Int., vol. 24, no. 12, pp. 920–925, 2010.

[92] K. X. Liu, J. X. Zhang, K. Zhao, X. J. Li, and K. Zhang, “Mechanism of explosive technique on relieving welding residual stresses,” Chinese Phys. Lett., vol. 22, no. 3, pp. 744–746, 2005.

[93] J. Zhang, K. Liu, K. Zhao, X. Li, Y. Liu, and K. Zhang, “A study on the relief of residual stresses in weldments with explosive treatment,” Int. J. Solids Struct., vol. 42, no. 13, pp. 3794–3806, 2005.

143

[94] S. Golovashchenko, “Electromagnetic Forming and Joining for Automotive Applications,” Ichsf, pp. 201–206, 2006.

[95] S. F. Golovashchenko, N. M. Bessonov, and A. M. Ilinich, “Two-step method of forming complex shapes from sheet metal,” J. Mater. Process. Technol., vol. 211, no. 5, pp. 875–885, 2011.

[96] E. Bruno, “High-velocity forming of metals,” 1968.

[97] R. S. and E. R. A. Davies, “Developments in high speed metal forming,” 1970.

[98] V. S. Balanethiram and G. S. Daehn, “Hyperplasticity: Increased forming limits at high workpiece velocity,” Scr. Metall. Mater., vol. 30, no. 4, pp. 515–520, Feb. 1994.

[99] P. R. L. and I. N. L. V. E. Kogan, “Shockwave Database.” [Online]. Available: http://www.ihed.ras.ru/rusbank/.

[100] R. C. Brune, S. R. Hansen, A. Vivek, J. M. Sosa, and G. S. Daehn, “Profile indentation pressure evaluation method for impulse manufacturing technologies,” J. Mater. Process. Technol., vol. 248, no. October 2016, pp. 185–197, 2017.

[101] J. R. Johnson et al., “Coupling experiment and simulation in electromagnetic forming using Photon Doppler Velocimetry,” Steel Res. Int., vol. 80, no. 5, pp. 359–365, 2009.

[102] C. L. Mader Program Manager Terry R Gibbs, J. F. Barnes Bobby G Craig William E Deal, and J. D. Richard Dick James N Johnson Elizabeth Marshall Charles E Morris Timothy R Neal Suzanne W Peterson Raymond N Rogers Melvin T Thieme Jerry D Wackerle John M Walsh, Los Alamos Series on Dynamic Material Properties Lasl Data Center for Dynamic Material Properties Technical Committee. 1980.

[103] M. A. Meyers, Dynamic Behavior of Materials. 1994.

[104] R. Jones, E and Graham, “Shear Strength Effects on Phase Transition ‘Pressures’ Determined From Shock-Compression Experiments,” in Accurate Characterization of the High-pressure Environment, 1971, p. 229.

[105] D. Bancroft, E. L. Peterson, and S. Minshall, “Polymorphism of iron at high pressure,” J. Appl. Phys., vol. 27, no. 3, pp. 291–298, 1956.

[106] S. R. Hansen, “Vaporizing Foil Actuator Process Parameters: Input CharactVaporizing Foil Actuator Process Parameters: Input Characteristics, Energy Deposition, and Pressure Outputeristics, Energy Deposition, and Pressure Output,” 2018. 144

[107] D. Tabor, “A simple theory of static and dynamic hardness,” Proc. R. Soc. London. Ser. A. Math. Phys. Sci., vol. 192, no. 1029, pp. 247–274, 1948.

145

Appendix A. SEM Images of Dissimilar Ni - Ni and Ni - Ti Weld Interfaces

(a) (b) (c) Figure 47. SEM images of 718 (top) – 625 (bottom) weld interface taken at the regions shown by the red squares

146

(a) (b) (c)

Figure 48. SEM images of 718 (top) – 230 (bottom) weld interface taken at the region shown by the red square

147

(a) (b) (c) (d)

Figure 49. SEM images of 625 (top) – 230 (bottom) weld interface taken at the region shown by the red square

148

(a) (b) (c) (d)

Figure 50. SEM images of 718 - Ti (flyer) weld interface taken at the region shown by the red square

149

150

(a) (b) (c)

Figure 52. SEM images of 230 - Ti (flyer) weld interface taken at the regions shown by the red squares

151

Figure 53. Optical microscopy images of 718 - Ti (target) weld interface

152

Figure 54. SEM images of 625 - Ti (target) weld interface

153

Cracks in Tungsten Carbides

Polishing Residue

Figure 55. SEM images of 230 - Ti (target) weld interface

154

Appendix B. PDV Flyer Velocity Plots

Figure 56. Flyer velocity profiles obtained using the PDV method for (a) Ni - 718, (b) Ni - 625, (c) Ni - 230, and (d) Ti - 6242

155

Appendix C. Derivation of Uniaxial Strain Condition for Plastic Deformation

The figure below represents an elastic perfectly-plastic material with no strain or directional hardening. The material has a finite thickness subjected to uniaxial compression due to a planar shockwave propagating through the thickness along the +z direction. This criterion was proposed by Jones and Graham [104] and the key idea extends as further back as Bancroft et al., [105].

Figure 57. Depiction of planar shock propagation through the thickness of the material

Assuming that the shockwave does not alter the dimensions of the material in the radial direction, perpendicular to its direction of propagation, the workpiece would be isolated to a state of uniaxial strain.

Therefore, the stresses (��,�) and strains (��,�) in the radial direction would be:

(Eq 8) �� = �� = �

(Eq 9) �� = �� 156

Assuming that the material is a linear isotropic solid, it would obey the following generalized Hooke’s laws:

� � = (� − � (� + � )) (Eq 10) � � � � �

� � = (� − � (� + � )) (Eq 11) � � � � �

� � = (� − � (� + � )) (Eq 12) � � � � �

Where E is the Elastic Modulus and � is the Poisson’s ratio

Substituting the boundary condition above, these are reduced to:

−� ��−�+� � = � ( ) (Eq 13) � � �−�

The goal of this work is to understand the stress condition for elastic to plastic transition during shockwave propagation. The Tresca yield criterion also known as the maximum shear stress condition serves as a conservative estimate for the minimum yielding condition.

According to the Tresca yield criterion:

(Eq 14) �� (��) = �� − ��

Which in this case is:

(Eq 15) �� = �� − ��

As �� = 0,

�−� � = � ( ) (Eq 16) �|�� −��

Equation 10 represents a conservative estimate of the minimum stress that promotes plastic deformation.

157

Appendix D. Velocimetry Data for Flattening Experiments

The frequency/velocity profiles for the flattening experiments in section 4.3 are provided below with the corresponding shock breakout velocities (SBV) and shock breakout pressure (SBP) wherever a breakout was observed.

Both Figure 58 and Figure 59 show the breakout profiles for the 2 mm samples held in the concave up (CU) and concave down (CD) geometries and shocked at different energies. These profiles correspond to the pressures shown in Figure 43e.

1 kJ 3 kJ 4.5 kJ

66.45 m/s 77.8 m/s 44.94 m/s 478.67 MPa 560.42 MPa 323.72 MPa

(a) (b) (c)

Figure 58. Shock breakout profiles, corresponding velocities and pressures for 2 mm samples subjected to shock processing in the concave up (CU) geometry at different energy levels

158

4.5 kJ 7 kJ 8 kJ

92.47m/s 666.1 MPa 101.23 m/s 71.66 m/s 729.18 MPa

(a) (b) (c) 12 kJ 16 kJ

83.01 m/s 597.95 MPa

65.04 m/s 468.51 MPa

(d) (e)

Figure 59. Shock breakout profiles, corresponding velocities and pressures for 2 mm samples subjected to shock processing in the concave down (CD) geometry at different energy levels

Figure 60 shows the velocimetry data for the sole PDV+Flattening experiment that was conducted on the 1 mm thick sample held down in the CU configuration and shocked at 8.5 kJ. The slight slope in the profile at the 26.5 µs mark indicates that the shock wave breakout was not captured.

159

8.5 kJ

39.26 m/s

Figure 60. Velocity profile for 1 mm sample subjected to shock processing in the concave up (CU) geometry at 8.5 kJ (no breakout)

160