Bond Improvement of Al/Cu Joints Created by Very High Power Ultrasonic Additive Manufacturing

THESIS

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

Adam G. Truog, B.S.W.E

Graduate Program in Welding Engineering

The Ohio State University

2012

Master's Examination Committee:

Professor Sudarsanam Suresh Babu, Advisor

Professor John C. Lippold

Adam G. Truog

2012

ABSTRACT

The extension of Ultrasonic Additive Manufacturing (UAM) to dissimilar materials allows for increased application in the aerospace, automotive, electrical and power generation industries. The benefit of UAM over standard ultrasonic welding is the ability to form complex geometries, such as, honeycomb structure and internal channels and also embed wires and sensors to create smart materials. UAM has had limited success bonding dissimilar materials and thus Very High Power Ultrasonic Additive

Manufacturing (VHP UAM), which increases the amplitude (from 26μm to 52μm) and normal force (from 2.5kN to 33kN), has been introduced to address this deficiency.

Al3003 and Cu110 dissimilar VHP UAM builds were heat treated at 350°C for ten minutes. A measure of maximum push-pin force revealed an improvement in the heat treated condition (from 23% to 49%) for all geometries. Intermetallic phase formation was noticed using the scanning electron microscope (SEM) backscatter detector. X-ray diffraction (XRD) was utilized to characterize the intermetallic layers through peak phase analysis. Al2Cu, AlCu and Al4Cu9 were found on the fracture surface of a heat treated build. It was determined that fracture occurred between the AlCu and Al4Cu9 intermetallic layers. High resolution SEM and fractal analysis were used to verify these findings.

Surface modification was evaluated as a method for improving bonding between dissimilar aluminum and copper welds. The copper foils were rolled with the sonotrode prior to welding, which increased the surface roughness from 0.175 Ra μm to 1.170 Ra

μm and then placed face down before welding. The maximum force during push-pin testing showed inconclusive results. Load versus displacement curves were analyzed and it was evident that modified structures exhibited a more energetic failure compared to as- welded builds.

A hypothesis was created to explain this phenomenon. It was expected that the peak load is a function of metallurgical bonding, while mechanical interlock requires more displacement for failure and is responsible for a more energetic failure. This indicates that surface modification led to an increase in mechanical interlock. This finding was supported by SEM fracture surface analysis where the amount of metallurgical bonding for the as-welded and surface modified builds was similar at 7.5% and 8.8% respectively. The surface modified builds, however, displayed 9% more flow morphology than the as-welded sample, indicating increased mechanical interlock.

In collaboration with researchers from the Mechanical Engineering department at

The Ohio State University, linear weld density (LWD) and area weld density (AWD) were correlated to both ultimate shear strength (USS) and ultimate transverse tensile strength (UTTS). It was found that no correlation between LWD and mechanical properties existed. AWD yielded a correlation between USS and percent bonded area, however no correlation was found for UTTS. Based on these findings, a new method was devised, SEM fracture surface analysis, to analyze the fracture surfaces in depth and

iii correlate the findings to mechanical strength. The initial findings of this method indicate a correlation between ductile fracture (expected to indicate metallurgical bonding) and both USS and UTTS.

To my parents, who taught me to walk

And to Liz, who inspired me to run

ACKNOWLEDGEMENTS

I would like to first acknowledge my advisor, Professor Suresh Babu. Thank you for your guidance, patience and continued dedication to furthering my education. Thanks also to my other committee member, Professor John Lippold. I have thoroughly enjoyed your lectures over the past years.

I express my gratitude to Mark Norfolk, Chris Conrardy, Gary Thompson and

Josh George of EWI for their help in defining and achieving my research goals. Thank you to Karl Graff for sharing your persistent and calculated approach to issues surrounding this thesis.

I would like to acknowledge the NSF I/UCRC Center for Integrative Materials

Joining Science for Energy Applications (CIMJSEA) for their sponsorship of this project.

Thank you to the fellow welding engineering graduate students for sharing knowledge, fruitful academic discussions and camaraderie. I would like to specifically thank Jeff Rodelas for your assistance with all things lab and electron microscopy related as well as your never ending support. Also, Samartha Channagiri for assistance with X- ray tomography. Thanks to Xiuli Feng for help with the TEM and Ryan Smith with

XRD. Finally, I extend my gratitude to my mentor, Sriraman Ramanujam, for creating the foundation for my success with this project.

Last of all, I would like to thank Margaret, Bill, Michael, Kathy, Liz and the rest of my friends and friends for your encouragement and positive thoughts. vi

Vita

September 5, 1986 ...... Born- Akron, OH U.S.A

2010...... B.S. Welding Engineering

The Ohio State University

Columbus, OH

2010 to present ...... Graduate Research Fellow

The Ohio State University

Columbus, OH

Publication

Hopkins, C.D., Wolcott, P.J., Dapino, M.J., Truog, A.G., Babu, S.S., Fernandez, S.A.,

Optimizing Ultrasonic Additive Manufactured Al 3003 Properties With Statistical

Modeling. Journal of Engineering Materials and Technology-Transactions of the

ASME, 2012. 134(1).

Field of Study

Major Field: Welding Engineering

vii

TABLE OF CONTENTS

ABSTRACT ...... ii

Vita ...... vii

List of Tables ...... xii

List of Figures ...... xiii

CHAPTER 1: MOTIVATION ...... 1

1.1 THESIS OUTLINE ...... 2

CHAPTER 2: BACKGROUND ...... 3

2.1 SOLID STATE BONDING ...... 3

2.1.1 Roll Bonding...... 4

2.1.2 Diffusion Bonding ...... 5

2.1.3 Ultrasonic Bonding ...... 8

2.1.4 Explosive Welding ...... 9

2.1.5 Friction Welding ...... 12

2.2 AL/CU SOLID STATE BONDING ...... 13

2.2.1 Explosive Welding ...... 13

2.2.2 Friction Stir Welding ...... 14 viii

2.2.3 Roll Bonding...... 15

2.2.4 Diffusion Bonding ...... 17

2.2.5 Ultrasonic Additive Manufacturing ...... 20

2.3 CHARACTERIZATION OF ULTRASONIC ADDITIVE MANUFACTURING 21

2.3.1 Linear Weld Density ...... 25

2.3.2 Scanning Electron Microscopy ...... 28

2.3.3 Tensile and Lap Shear Testing ...... 30

2.3.4 Peel Test ...... 32

2.3.5 Push-Pin ...... 35

CHAPTER 3: OBJECTIVES ...... 37

CHAPTER 4: FRACTURE SURFACE ANALYSIS OF ULTRASONIC ADDITIVE

MANUFACTURED AL 3003 ...... 38

4.1 INTRODUCTION AND MOTIVATION ...... 38

4.2 BACKGROUND ...... 39

4.2.1 Linear Weld Density ...... 39

4.2.2 Area Weld Density ...... 40

4.3 RESULTS AND DISCUSSION ...... 42

4.4 CONCLUSIONS ...... 46

CHAPTER 5: SURFACE MODIFICATION OF VERY HIGH POWER ULTRASONIC

ADDITIVE MANUFACTURED ALUMINUM AND COPPER STRUCTURES ...... 47

5.1 INTRODUCTION ...... 47

5.2 EXPERIMENTAL PROCEDURE ...... 50

5.3 RESULTS AND DISCUSSION ...... 59

5.4 CONCLUSIONS ...... 68

CHAPTER 6: HEAT TREATMENT OF VERY HIGH POWER ULTRASONIC

ADDITIVE MANUFACTURED (VHP UAM) ALUMINUM AND COPPER

STRUCTURES ...... 69

6.1 INTRODUCTION ...... 69

6.2 EXPERIMENTAL WORK ...... 72

6.2.1 VHP UAM of Al3003 and Cu110 ...... 72

6.2.2 Heat Treatment ...... 74

6.2.3 Mechanical Push-Pin Testing ...... 74

6.2.4 Metallography ...... 75

6.2.5 X-ray Diffraction ...... 76

6.2.6 Fractal Analysis ...... 76

6.3 RESULTS...... 77

6.3.1 Mechanical Testing of Builds ...... 77

6.4. DISCUSSION ...... 86

6.5. CONCLUSIONS ...... 90

CHAPTER 7: CONCLUSIONS ...... 91

CHAPTER 8: FUTURE WORK ...... 93

REFERENCES ...... 95

APPENDIX A: X-RAY TOMOGRAPHY OF VHP UAM BUILDS ...... 105

A.1 INTRODUCTION ...... 105

A.2 EXPERIMENTAL ...... 106

A.3 RESULTS ...... 106

A.4 DISCUSSION AND CONCLUSIONS ...... 113

A.5 REFERENCES ...... 113

APPENDIX B: LOAD VS. DISPLACEMENT CURVES ...... 115

List of Tables

Table 4.1. Area fraction of features measured on fracture surface using SEM ...... 44

Table 5.1. Measured surface roughness values for aluminum and copper foils...... 50

Table 5.2. Average fractal dimension for the as-welded and surface modified interfaces.

...... 65

Table 5.3. Area fraction of fracture morphologies measured from SEM images...... 66

Table 6.1. Specification of materials used in this investigation ...... 73

Table 6.2. Configuration of build geometries ...... 73

Table 6.3. Average fractal dimension for various interfaces. The lowest FD is found at the fractured layer between CuAl-Cu9Al4. The largest FD is found between the Al-

CuAl2 layers, significantly higher than the non-heat treated Al/Cu interface...... 85

xii

List of Figures

Figure 2.1. Schematic of the roll bonding process. Adapted from Li et al. [12]...... 5

Figure 2.2. Schematic of diffusion bonding set-up. Adapted from Zuruzi et al. [15] ...... 6

Figure 2.3. Stages of diffusion bond formation [16]...... 7

Figure 2.4. Wedge-Reed and later drive ultrasonic spot welding systems [23]...... 9

Figure 2.5. Schematic of explosive welding [25]...... 10

Figure 2.6. Typical wavy bond interface from explosive welding of titanium on steel [26].

...... 11

Figure 2.7. Schematic of friction welding. (a) One work piece is rotated while the other is held stationary. (b) The interfaces are brought together and axial force is applied which begins the upset process. (c) Rotation is stopped and upset process is complete[31]...... 12

Figure 2.8. Peel strength of roll bonded copper and aluminum laminates at various sintering temperatures [39]...... 16

Figure 2.9. Variation of peel force with intermetallic width [32]...... 17

Figure 2.10. SEM images taken with backscatter electron detector of diffusion bonded interface between aluminum and copper annealed for 10 minutes at (a) 400°C (b) 450°C and (c) 500°C [41]...... 19

Figure 2.11 Free energy (G) versus composition of the Cu-Al system at 500°C [41]. .... 19

xiii

Figure 2.12 Monte-Carlo simulations of electron trajectories for aluminum and copper at

20kV with 20nm probe diameter. Red trajectories denote backscatter electrons while blue lines are secondary electrons...... 21

Figure 2.13. Schematic of UAM process [43]...... 22

Figure 2.14. Schematic of a) UAM single transducer and b) VHP UAM double transducer “push-pull” design [6]...... 22

Figure 2.15. Fabrisonic LLC Soniclayer 7200 [44]...... 23

Figure 2.16. Ion beam induced secondary electron micrograph (IBISEM) illustrating the persistent oxide layer at the bond interface [43]...... 25

Figure 2.17. Cross section of UAM build where D1 indicates line defects, D2 shows parabola defects and D3 is point defects [46]...... 26

Figure 2.18. Near 100% linear weld density [46]...... 27

Figure 2.19. SEM images of interface between Metpreg® and Al3003 for different bonding conditions where a) shows no interfacial defects and b) has distinct voids at interface [3]...... 29

Figure 2.20. SEM image of SMA fibers with no visible distortion. Flow lines and weld interface are visible, adapted from Kong et al. [4]...... 30

Figure 2.21. Schematic for tensile test specimen, adapted from Hopkins et al. [5] ...... 31

Figure 2.22. Schematic for lap shear specimen, adapted from Hopkins et al. [5]...... 31

Figure 2.23. Schematic of peel test [45]...... 33

Figure 2.24. Load versus displacement curve with failure modes indicated, adapted from

Kong et al. [45]...... 34

xiv

Figure 2.25. a.) Push-pin testing setup and b.) schematic of push-pin testing sample, adapted from Zhang et al. [53]...... 36

Figure 4.1. LWD versus average USS and UTTS of Al3003 UAM specimens...... 39

Figure 4.2. Optical images of transverse tensile fracture surfaces (top surface) from

Experiment #2 a). before image processing and b). after threshold adjustment. Region I is damaged material cause by bonding or contact with previous surface. Region II is material unaffected by the UAM process...... 40

Figure 4.3. Percentage of bonded area on fracture surfaces versus mechanical strength of both UTTS and USS samples...... 41

Figure 4.4. Example SEM images representing shear ductile failure, ductile failure, flow, brittle shear, and machined surface...... 43

Figure 5.1. Push-pin testing setup in Gleeble™ ...... 51

Figure 5.2. Schematic of push-pin testing sample ...... 52

Figure 5.3. Schematic drawing indicating a decrease in apparent roughness, RL, correlating to an increase in segment measure length, η [59]...... 54

Figure 5.4 Fractal plot of profile roughness. DL, the fractal dimension is measured from the slope, α [59]...... 55

Figure 5.5 Joint Roughness Coefficient (JRC) vs. fractal dimension, adapted from Lee et al. [60] ...... 56

Figure 5.6. Representative SEM micrograph indicating machined, ductile and flow morphologies...... 58

Figure 5.7. Mechanical push-pin test results for the as-welded and surface modified condition ...... 60

Figure 5.8. Force vs. Displacement graph for as-welded and surface modified builds.

Percentages reflect relative area under curve normalized to the largest area (100%)...... 61

Figure 5.9. Schematic representation of delamination failure after peak load is achieved.

It is hypothesized that metallurgical bonding fails during peak loading (characterized by ductile fracture) and load is then transferred to mechanical interlock as extension continues until complete failure occurs...... 63

Figure 5.10. SEM micrographs of the a.) as-welded and b.) surface modified interface. 64

Figure 5.11. Low magnification fracture surface SEM image for a.) as-welded and b.) surface modified builds...... 67

Figure 6.1. a.) Push-pin testing setup in Gleeble™ and b.) schematic of push-pin testing sample. Adapted from Zhang et al. [53] ...... 75

Figure 6.2. Maximum force recorded for four different VHP UAM build geometries.

Percentages indicate amount of improvement in heat treated samples...... 77

Figure 6.3. Force vs. Displacement graph for as-welded and heat treated builds.

Percentages reflect relative area under curve normalized to the largest area (100%). The as-welded sample with a dotted line indicates more deformation during failure and differs from the other as-welded samples...... 79

Figure 6.4. Alternating aluminum and copper builds after push-pin testing in the a) as- welded b). as-welded with deformation (denoted by the dotted line in Figure 6.3) and c) heat treated condition...... 80

xvi

Figure 6.5. The effect of heat treatment on a dissimilar aluminum and copper VHP UAM interface. a.) As-welded interface. b.) Interface after 10 minute heat treatment at 350°C.

...... 81

Figure 6.6. X-Ray diffraction patterns from both copper and aluminum side of fractured heat treated sample...... 82

Figure 6.7. Aluminum and copper equilibrium phase diagram ...... 83

Figure 6.8. Fracture interface of heat treated Al/Cu build taken with high resolution

SEM. Fracture occurred in Cu9Al4 intermetallic layer...... 84

Figure 6.9. High magnification SEM backscatter image of heat treated interface. a.)

Original image. b.) Modified image with digitized interface...... 85

Figure 6.10 Free body diagram for two conditions, a) represents IMF below the critical thickness and results in a fully constrained interface. b) depicts a situation where the intermetallic layer has grown beyond the critical intermetallic thickness and is not constrained in the X direction...... 89

Figure A.1. Radiographic image of Cu sample 1...... 107

Figure A.2. Radiographic image of Cu sample 2...... 107

Figure A.3. Radiographic image of Cu sample 3...... 108

Figure A.4. Radiographic image of Cu sample 4...... 108

Figure A.5. Radiographic image of Cu sample 5...... 109

Figure A.6. Radiographic image of Cu sample 6...... 109

Figure A.7. Radiographic image of Cu sample 7...... 110

Figure A.8. Radiographic image of Cu sample 8...... 110 xvii

Figure A.9. Radiographic image of Cu sample 9...... 111

Figure A.10. Radiographic image of Cu sample 10...... 111

Figure A.11. Radiographic image of Cu sample 11...... 112

Figure A.12. Radiographic image of Cu sample 12...... 112

Figure B.1 Aluminum and copper baseplate load vs. displacement curve and fractured builds...... 116

Figure B.2. “All aluminum” and “all copper” load vs. displacement curves and fractured builds...... 117

Figure B.3 “Copper on top” and “aluminum on top” load vs. displacement curves and fractured builds. SEM image of sheared button shown in c)...... 118

Figure B.4 Heat treated “all aluminum” and “all copper” load vs. displacement curves and fractured builds...... 119

Figure B.5 14μm sonotrode and 700lbf Load vs. displacement curves and fractured builds in a “cu on top” configuration...... 120

xviii

CHAPTER 1: MOTIVATION

There has been interest in using Very High Power Ultrasonic Additive

Manufacturing (VHP UAM) to join dissimilar materials for energy applications, such as compact heat exchangers and solar collectors [1-3]. However, poor bond quality has been associated with multi material consolidation using UAM [2-5]. VHP UAM was created to improve bonding by increasing the maximum processing amplitude (from 26μm to

52μm) and normal force (from 2.5kN to 33kN). These increases have allowed for successful initial joining studies of previously difficult to join materials [6].

Ultrasonic additive manufacturing is a solid state process that offers the ability to create near net shape parts with unique geometries and dissimilar materials. Examples include: heat exchangers [7], solar collectors [2], honeycomb structure [2], embedding of optical, shape memory and sensing fibers [2], fiber reinforced metal matrix composites

[8], rapid tooling for injection molding [1] and structures with internal channels [1]. It is clear that UAM is a process that can be used in a wide array of applications and industries. Many research efforts have focused on expanding the process to new material combinations and optimizing process parameters for Al3003.

Despite attempts at optimizing bonding through process parameter manipulation, the best UAM builds have only a fraction of the bulk material properties [9]. It is clear that the current processing techniques are not suitable for the extension of this process to

1 new industries and applications. This research looks to address this issue by investigating the effect of pre- and post-processing techniques that are independent of standard processing parameters such as amplitude, normal force and travel speed. Two methods were investigated in this study, a pre-processing surface modification and post-processing heat treatment.

1.1 THESIS OUTLINE

The body of this thesis is composed of three individual papers submitted to various journals and conferences. Each section is composed of an introduction, experimental procedure, results/discussion and conclusion section, each specific to the research topic. The papers are as follows:

Chapter 4: Fracture Surface Analysis of Ultrasonic Additive Manufactured

Al3003, published in the Journal of Engineering Materials and Technology.

Chapter 5: Surface Modification of Very High power Ultrasonic Additive

Manufactured Aluminum and Copper Structures, submitted to the Trends in Welding

Research Conference.

Chapter 6: Heat Treatment of Very High Power Ultrasonic Additive

Manufactured (VHP UAM) Aluminum and Copper Structures, not yet submitted.

CHAPTER 2: BACKGROUND

2.1 SOLID STATE BONDING

An introduction to solid state bonding is necessary to understand the various advantages and disadvantages of Very High Power Ultrasonic Additive Manufacturing compared to other solid state processes. The following section will outline barriers to bonding, techniques to disrupt these barriers as well as an overview of specific solid state bonding processes.

Before the specifics of UAM can be investigated, the basics of solid state joining must be reviewed. There are three barriers to solid state bonding that inhibit nascent

(oxide and contaminate free) metal contact. Each of these must be eliminated in order to form a solid-state bond. The three barriers are: asperities, oxide layers and surface contaminates.

Asperities are high and low areas of a metal surface, often referred to as “peaks and valleys”. These undulations prevent complete contact between two parts. Oxides form on material surfaces and are generally hard and brittle. These films do not allow for intimate metal-metal contact and therefore inhibit solid state bonding [10]. Surface contaminates, such as grease, gas molecules and water vapor adhere to the surface by secondary bonding. These contaminates prevent metal-metal interaction and must be removed to achieve a solid state bond.

In order to reduce the barriers to solid state bonding a combination of surface preparation, plastic deformation and heat can be used. Each solid state welding process

3 uses a combination of these conditions to produce a joint. The following section will detail common solid state welding processes, indicate the specific mechanism that allows solid state bonding to occur and outline the benefits and disadvantages of each method.

2.1.1 Roll Bonding

A schematic representation of roll bonding is presented in Figure 2.1. In this process surface preparation of the material is performed to remove surface contaminates and then immediately passed through a rolling mill. During rolling the overall material thickness is reduced and significant deformation at the faying interface leads to the fracturing of the oxide layers and allows fresh metal surface interaction [11]. Heat can be applied during welding which reduces the yield stress of the material/s and ultimately requires less reduction in thickness to bond. This process is easily automated, has high productivity, is cost effective and able to join dissimilar materials at room temperature

[12]. The main limitation of roll bonding is weld geometry.

Surface Preparation Material Preparation

Rolling of Material

Stacking

Figure 2.1. Schematic of the roll bonding process. Adapted from Li et al. [12].

2.1.2 Diffusion Bonding

Diffusion bonding is a process in which faying surfaces are brought into intimate contact by using an applied pressure at elevated temperature. This allows for bond formation by atomic interdiffusion across the joint interface [13]. The predominant processing parameters are bonding temperature, bonding pressure and holding time [14].

Figure 2.2 depicts a basic diffusion bonding set-up where the materials to be joined are placed between two metal platens in an oven.

Oven Metal Platens

Specimens

Pull Rod

Figure 2.2. Schematic of diffusion bonding set-up. Adapted from Zuruzi et al. [15]

Mechanical intimacy of the faying surfaces and disruption/dispersion of the surface contaminates are two conditions that must be met to create a successful joint.

Figure 2.3 schematically pictures the stages of joint formation for diffusion bonding.

Stage 1 is asperity contact. Stage 2 involves the deformation of these asperities due to applied pressure and heat. The third stage includes grain boundary migration, recrystallization and pore size reduction. The final stage entails bulk diffusion, oxide and contaminant dissolution and pore elimination [16, 17].

Stage 1. Asperity contact Stage 2. Asperity deformation and interfacial boundary formation

Stage 3. Grain boundary Stage 4. Volume diffusion and migration and pore size reduction pore elimination

Figure 2.3. Stages of diffusion bond formation [16].

Diffusion bonding boasts many advantages, these include [18]:

 Low deformation which enables machined parts to be joined without distortion

 The ability to join large areas

 Smaller thermal gradients than fusion welding minimizing microstructural

gradients and residual stress [18].

 The ability to join metals in non conventional situations such as a low gravity

environment[19].

Despite many advantages, diffusion bonding is performed at temperatures above

0.75Ts, where Ts is the solidus temperature of the material. This leads to the following disadvantages [20]: 7

 Surface oxidation of the work piece is accelerated

 Grains grow rapidly during pre-heating

 Intermetallic formation between dissimilar materials

2.1.3 Ultrasonic Bonding

This process utilizes ultrasonic energy and pressure to induce an oscillating shear strain between the faying surfaces and produce a metallurgical bond [21]. The exact mechanism of bond formation for ultrasonic welding is not well understood. There are numerous proposed mechanisms for ultrasonic welding, these include recrystallization, plastic deformation, work hardening, breaking of oxide films, generation of heat by friction and even melting [21, 22]. It is generally accepted that localized slip and sub layer plastic deformation are desirable and interfacial slip is responsible for breaking up surface oxides, allowing for nascent metal interaction at asperity contact points [22].

The two most widely used systems for spot welding are the Wedge-Reed system and lateral drive system. Basic schematics for each system are pictured in Figure 2.4.

Both styles use a transducer to produce longitudinal vibrations from electrical power.

The Wedge-Reed system utilizes a wedge to amplify the vibrations from the transducer which are then transferred to the reed and sonotrode. In a lateral drive system, the vibrations are amplified by a booster and then applied to the samples through the sonotrode. During welding a normal force is applied to the part, this serves two functions. The first is to hold the work pieces in place and second is to impart plastic deformation during welding [23].

Wedge-Reed system

Lateral Drive system

Figure 2.4. Wedge-Reed and later drive ultrasonic spot welding systems [23].

The general advantages of solid state bonding apply to ultrasonic welding.

Shrinkage and distortion problems are avoided due to the absence of liquid formation

[22]. Welding can be performed without significant surface preparation because of the extreme deformation experienced at the interface. Lastly, this process allows for thin foils and wires to be joined to thicker material sections. Major limitations include joint geometry and material thickness.

2.1.4 Explosive Welding

Figure 2.5 shows a common explosive welding set-up. In this process explosives are used to accelerate a sheet of material (flyer plate) into another (parent plate) and form

9 a metallurgical bond between the two materials. As the plates crash together a thin layer of metal is removed from both surfaces (generally less than .05mm). This allows for perfectly cleaned metal surfaces to bond under high pressure [24]. A representative interface is pictured in Figure 2.6.

Sheet explosive Buffer Flyer plate Detonator

Parent plate Support AnvilOven

Figure 2.5. Schematic of explosive welding [25].

Figure 2.6. Typical wavy bond interface from explosive welding of titanium on steel [26].

Multiple theories exist to explain the wavy interface morphology created from explosion welding [27]. Discussion of these theories is beyond the scope of this review.

It is important to note, however, that the common principles of solid state bonding, such as the fracture of oxide layers, deformation and heat generation are all present, allowing for solid state bonding to occur.

Bond quality is dependent upon four parameters: collision angle, impact velocity, material properties and geometry of the flyer plate relative to the parent [28]. The main advantage of explosive welding is the vast array of materials that can be joined. Material selection is limited only by material ductility and strength [29]. Weld geometry is restricted to lap joints and thus the most common application for this process is cladding

[26].

2.1.5 Friction Welding

Friction welding produces coalescence of materials through the frictional heating of interfaces along with the application of pressure [30]. This process usually involves rotating one part against another. Friction welding is schematically represented in Figure

2.7. As the materials come in contact the interfacial temperature is raised. In most materials this will lower the yield strength and increase diffusion. As the parts are continually pressed together intimate contact at the faying surfaces occurs [31]. There are many variations of friction welding such as linear friction welding, friction stud welding and friction stir welding. While there are geometry restriction, friction welding is simple, easily automated and does not require surface preparation for oxide removal.

(a)

Friction phase (b)

Forging phase (c)

2.2 AL/CU SOLID STATE BONDING

Successfully bonding aluminum and copper presents significant challenges.

Aluminum and copper are incompatible metals because above 120°C intermetallic compounds can form which are typically low strength, brittle and posses high electrical resistance [32]. For this reason, fusion welding techniques are not applicable to join aluminum and copper [2, 32, 33]. Solid state processes are a logical candidate to join these materials because they avoid the liquid phase. However, some solid state processes operate at elevated temperatures and can form intermetallics at the interface without eclipsing the melting temperature. This section will provide an overview of aluminum and copper bonding with various solid state processes.

2.2.1 Explosive Welding

Gulenc studied the effect of explosive ratio, a major welding parameter described by the amount of explosive compared to the weight of the upper plate, on pure aluminum and copper plates [24]. As the ratio was increased the bonding interface became wavier.

It has been noted previously that excessive waviness has led to the formation of intermetallics in dissimilar steel bonding [34]. With optimized bonding parameters aluminum and copper plates can be joined successfully. The benefit of explosively welded aluminum and copper joints was also discovered by Veerkamp for use in high direct-current bus systems [35].

Magnetic Impulse welding is mechanistically similar to explosive welding but utilizes electromagnetic force to drive the flyer plate instead of explosives [36]. This

13 process was used to investigate the feasibility of joining aluminum and copper tubing

[36]. Welds were made with an outer copper tube and inner aluminum core, as well as the reverse of this geometry. It was determined that intermetallic compounds were formed in some cases, but not all. Intermetallic formation was the expected cause for lower weld strength.

2.2.2 Friction Stir Welding

Friction stir welding feasibility studies of joining aluminum and copper have been conducted [33, 37, 38]. Ouyang et al. attempted to create a butt joint between 6061 aluminum and pure copper plates [33]. The study concluded that direct welding of the two materials proved difficult due to the formation of intermetallic compounds, namely

CuAl2, CuAl and Cu9Al4 in the weld nugget. It was suggested to use an interlayer to avoid aluminum and copper intermetallic formation to produce sound welds. A second study by Elrefaey et al. looked to join commercially pure aluminum and copper plates in the lap joint geometry [37]. Successful joining was possible and it was noted that joints bonded at higher rotation speeds mainly fractured in a brittle manner, owing to the formation of intermetallic compounds.

Building on the findings of the previous studies, Abdollah-Zadeh et al. performed a weldability study to examine friction stir welded aluminum and copper lap joints and to evaluate the effect of the welding parameters on joint quality [38]. Results from tensile shear testing showed failure in the base material, indicating that high quality joints could be made over an array of processing parameters. It was concluded that higher rotational

14 speeds led to decreased mechanical strength. It was hypothesized that since higher rotational speed yields higher temperatures at the interface, nucleation and growth processes are accelerated and increased amounts of intermetallic compounds can form.

2.2.3 Roll Bonding

Roll bonding at room temperature has been found to avoid intermetallic formation

[12]. Despite this perceived advantage, recent roll bonding studies have applied a sintering heat treatment which was found to improve bond strength. Peng et al. studied the effect of rolling and sintering temperature on the interface and bond strength development of roll bonded aluminum and copper laminates [39]. The rolling temperature during bonding was increased from 350°C to 500°C and bond strength was evaluated via peel testing. All samples showed formation of intermetallic compounds.

The rolling temperature that produced the highest bond strength was 430°C. Structures rolled at 430°C were then sintered at temperatures between 350°C and 500°C, the results are shown in Figure 2.8. The peel strength of laminates generally increased from 400°C to 450°C and then decreased substantially. It was hypothesized that physical contact

(mechanical interlock) formed during the roll bonding stage and the formation of strong intermetallic phases at the interface after sintering were responsible for the improved mechanical properties.

Peel Strength (N/cm)StrengthPeel

Sintering Temperature ( C)

Figure 2.8. Peel strength of roll bonded copper and aluminum laminates at various sintering temperatures [39].

The development of intermetallic phases at the interface of roll bonded aluminum and copper laminates has been studied in detail [32, 40]. Figure 2.9 plots peel force versus intermetallic width and reveals a ductile to brittle transition at an intermetallic width of 2.5 μm [32]. This phenomenon accounts for the precipitous drop in peel strength found by Peng et al. as increased sintering temperature will affect the kinetics of intermetallic formation [38]. While the phenomenon is commonly noted, no mechanism

16 for intermetallic formation and improved bond strength in dissimilar materials has been proposed.

Ductile to brittle

intermetallic width Peeling force (N/cm) force Peeling

intermetallic width (μm)

Figure 2.9. Variation of peel force with intermetallic width [32].

2.2.4 Diffusion Bonding

Diffusion bonding of aluminum and copper is typically carried out between

300°C and 500°C with hold times anywhere from 10 to 290 minutes [17, 41]. As the

17 mechanism of diffusion bonding indicates, intermetallic formation is expected. Guo et al. studied the intermetallic phase formation of diffusion bonded aluminum and copper laminates at 400°C, 450°C and 500°C for varying times. It was found that three intermetallic phases form at the interface: Al4Cu9, AlCu and Al2Cu. These phases are shown in Figure 2.10 for various hold times. It was determined that Al2Cu was the first phase to form, followed by Al4Cu9 and finally AlCu. The reason for this order of formation was investigated. The effective heat of formation model, developed by

Pretorious et al. was successfully used to predict the formation of Al2Cu [42]. This model combines thermodynamics and kinetic theory to determine the effective heat of formation [42]. This factor, ∆H′, is composed of the standard heat of formation, ∆H , multiplied by the ratio of effective concentration of the limiting element, Ce, over C1, the concentration of the limiting element in the compound. A Gibbs free energy- composition diagram at 500°C was created, shown in Figure 2.11, which shows Al4Cu9 has the most negative driving force, resulting in the preferential formation of this phase over AlCu [41].

Figure 2.10. SEM images taken with backscatter electron detector of diffusion bonded interface between aluminum and copper annealed for 10 minutes at (a) 400°C (b) 450°C

and (c) 500°C [41]. G (kJ/mol) G

Composition Al (mole fraction) Figure 2.11 Free energy (G) versus composition of the Cu-Al system at 500°C [41].

2.2.5 Ultrasonic Additive Manufacturing

Few studies have attempted to bond aluminum and copper with Ultrasonic

Additive Manufacturing [1, 2]. Janaki Ram et al. performed a feasibility study using

Al3003 and Cu110 [2]. SEM micrographs revealed an interface that was described as looking “tight” at low magnification and revealed a thin porous layer on the aluminum side at high magnification [2]. It is worth noting that intermetallic formation was not discovered in images of the interface taken with the backscatter electron detector. This result suggests that Ultrasonic Additive Manufacturing may bond foils at temperatures low enough to suppress brittle intermetallic formation.

Energy Dispersive Spectroscopy (EDS) was performed across the interface. It was concluded that no diffusion occurred of copper into aluminum, however there was considerable aluminum diffusion into copper. This result is misleading for multiple reasons. First, the spot spacing used during the EDS line scan was 1.5μm. This distance is too large to pick up compositional differences across the interface that is less than 1μm.

Secondly, the spatial resolution of SEM EDS is governed by the beam/specimen interaction volume. A Monte-Carlo simulation was run at 20kv with 20nm probe diameter for both aluminum and copper. The results are pictured in Figure 2.12. It is evident that the interaction volume is on the order of 1μm-2μm and does not provide good resolution for an interface that is of the same scale. TEM EDS may be required to accurately determine the amount of diffusion across the Al/Cu interface.

Al 0 μm Cu 0 μm

548 μm 187 μm

1095 μm 374 μm

1643 μm 561μm

2191 μm 748 μm

-1640 μm -820 μm 0 μm 820 μm 1640 μm -560 μm -280 μm 0 μm 280 μm 580μm

Figure 2.12 Monte-Carlo simulations of electron trajectories for aluminum and copper at

20kV with 20nm probe diameter. Red trajectories denote backscatter electrons while blue lines are secondary electrons.

2.3 CHARACTERIZATION OF ULTRASONIC ADDITIVE MANUFACTURING

Ultrasonic Additive Manufacturing joins thin (100μm-300μm) foils in succession with ultrasonic energy. A computer numerical controlled (CNC) machining step is also incorporated so that material can be both added and subtracted during part formation.

Parts can be created based on a computer aided design (CAD) drawing. There are three main processing parameters: amplitude, normal force and travel speed. Other parameters include: platen temperature, foil and base plate surface roughness as well as sonotrode texture and roughness. Mechanistically UAM is no different from VHP UAM. VHP

UAM was developed so that harder and stronger materials could be processed by increasing the maximum amplitude (from 26μm to 52μm) and normal force (from 2.5kN to 33kN) [6]. A diagram in Figure 2.13 outlines the basic mechanics of UAM/VHP

UAM.

Oscillating and Force applied by sonotrode rotating sonotrode Oscillating unbonded foil Ultrasonic Vibration

Oxide Unbonded foil Layer Consolidated material Consolidated material Held stationary by anvil Base plate Figure 2.13. Schematic of UAM process [43].

Solidica Inc. manufactured the Form-ation™ UAM machine. This was the first commercial UAM machine. The Edison Welding Institute (EWI) produced a VHP UAM test bed where transducer design and material feasibility could be tested. This machine consisted of a VHP UAM dual transducer mounted to a modified 3-axis milling machine, a schematic of the dual transducer design next to a single transducer is pictured in Figure 2.14 [6]. Successful welding with the test bed led to the creation of Fabrisonic

LLC and the commercial availability of a large production VHP UAM machine, the

SonicLayer 7200, shown in Figure 2.15 [44].

a) b)

Figure 2.14. Schematic of a) UAM single transducer and b) VHP UAM double transducer “push-pull” design [6].

Figure 2.15. Fabrisonic LLC Soniclayer 7200 [44].

The UAM process is mechanistically very similar to ultrasonic welding. The generally accepted model for UAM bond formation has been outlined by Janaki Ram et al. and indicates two repeated and distinct stages of bonding [3]. First is the generation of contact points (contact stage), and second is the formation of bonds across the contact points (bond stage). The first condition is met when the sonotrode applies a normal force. As ultrasonic energy is applied through the sonotrode and it traverses the layers to be bonded, the application of normal and oscillating shear forces result in the generation

23 of dynamic interfacial stresses between the contact points. It is hypothesized that the oxide film at these contact points begins to crack and nascent metal is extruded. As welding continues the oxide layer is fractured and displaced, creating atomically clean metal surfaces, inducing metallurgical bonding [3].

Recent discoveries by Johnson et al. have challenged the theory of oxide fracture and dispersal [43]. Aluminum 3003 samples were examined using a scanning electron microscope/focused ion beam (SEM/FIB) machine to produce ion beam induced secondary electron micrographs (IBISEM‟s). Results from this technique indicated that after welding, a persistent surface oxide layer between the bonded foils remained. The oxide layer can be seen in Figure 2.16. It was hypothesized that mechanical interlock and bonding between aluminum and Al2O3 were responsible for bond formation in UAM builds.

Bottom of top foil layer

Top of bottom foil layer

Figure 2.16. Ion beam induced secondary electron micrograph (IBISEM) illustrating the persistent oxide layer at the bond interface [43].

2.3.1 Linear Weld Density

Linear weld density (LWD) is a measure of defect free interfacial length in UAM builds [45]. It is a determined by measuring the interface length without defects divided by the total interface length. It is commonly accepted that percent linear weld density correlates to the percent of bonding within a UAM structure [8, 45-47]. A thorough characterization of UAM bond defects has been performed by Janaki Ram et al. [46].

Three defect morphologies were observed: line, parabola and point. Figure 2.17 shows an example of each morphology. Attempts have been made to optimize process parameters based on linear weld density; a brief outline of these attempts is presented.

Figure 2.17. Cross section of UAM build where D1 indicates line defects, D2 shows parabola defects and D3 is point defects [46].

Kulakov and Rack investigated the effect of normal load, amplitude and travel speed on bond quality as characterized by the total liner weld density [48]. The repeatability of linear weld density was tested by taking five measurements of the same interface. There was less than 1% variability between the five measurements [48]. The same authors investigated the effect of surface damage left on the top foil surface from the sonotrode after welding on liner weld density [49]. An increasing linear relationship between surface damage and linear weld density was determined.

Kong et al. correlated bond strength measured by peel testing to linear weld density [45]. As amplitude and travel speed were increased it was found that in general

LWD and peel strength increased. However, there were multiple instances where peel

26 strength was high but linear weld density was low. It was concluded that low linear weld density did not necessarily indicate a poor weld [45].

Optimization of linear weld density was attempted by Janaki Ram et al. [46]. By adjusting amplitude, travel speed, normal force and substrate temperature the maximum

LWD achieved was 98%. It was concluded that while optimized parameters can improve linear weld density, defects cannot be eliminated completely with the current process and procedures. It was suggested that a surface machining step to level the top of the recently welded foil be added between welds. This technique, paired with optimal welding parameters produced samples with near 100% linear weld density; an example of a build with very high linear weld density is pictured in Figure 2.18.

Figure 2.18. Near 100% linear weld density [46].

The aforementioned studies indicate that a correlation exists between linear weld density and processing parameters. Some results, such as the contradictory findings by 27

Kong et al. indicate the need for linear weld density to be qualified with further microstructural and mechanical investigation.

2.3.2 Scanning Electron Microscopy

Many UAM studies have utilized a scanning electron microscope (SEM) to characterize UAM builds. SEM images have much higher resolution, magnification and depth of field than optical images which allows for higher definition pictures [50]. This technique has been used frequently during feasibility studies to make judgments on overall bond quality.

Several studies have investigated the viability of joining multiple materials using

UAM [1-4]. Scanning electron microscopy has been used to analyze the joint interface.

A study that joined Al 3003 to MetPreg®, a commercially available Al2O3 short fiber reinforced aluminum matrix composite, analyzed the joint interface between the materials for two different bonding conditions [3]. Figure 2.19 shows the SEM pictures for the two conditions. Excessive void formation indicative of poor bonding is shown in Figure

2.19b. The featureless interface was then described as “very well bonded” [3]. This assumption may not be true, however, as the only information that can be gained from

Figure 2.19a, is that no voids formed along the interfacial region. Mechanical testing or more in depth microstructural analysis is needed to accurately comment on bond quality.

a) b)

Figure 2.19. SEM images of interface between Metpreg® and Al3003 for different bonding conditions where a) shows no interfacial defects and b) has distinct voids at interface [3].

Shape memory alloy (SMA) fibers were embedded between aluminum foils in a study to assess the feasibility of fabricating adaptive composite structures. An SEM image, shown in Figure 2.20, clearly shows no distortion or deformation to the fibers.

The unbounded weld interface is visible along with plastic flow lines above and below the weld interface.

Weld interface

Flow lines

Figure 2.20. SEM image of SMA fibers with no visible distortion. Flow lines and weld interface are visible, adapted from Kong et al. [4].

2.3.3 Tensile and Lap Shear Testing

Both tensile and lap shear testing has been identified as techniques capable of collecting bulk strength properties of UAM builds [51]. Shear specimens have been built based on the ASTM standard test method for lap shear strength of sealants, ASM C 961-

06 [5, 9]. Due to the small size of tensile test specimens, shown in Figure 2.21, no standard was followed. Control tests of solid samples of the same material were performed to ensure the test was repeatable and not biased.

1.9 cm R 3.17 mm Tensile Load Tensile Load

.95 cm

Figure 2.21. Schematic for tensile test specimen, adapted from Hopkins et al. [5]

The lap shear specimens are designed so that the interface of a tape layer is oriented along the shear plane, as shown in Figure 2.22. Successful implementation of both methods has been noted by Hopkins et al. where optimal process parameters for aluminum and titanium UAM builds were determined based on lap shear and tensile strength data [5]. The main drawback of these testing methods is that they require many welded foils to create the test specimen. An overall thickness/height of 1.9 cm is needed for both tests which will require approximately 130 welded foils (each 150μm thick).

Shear plane Compressive force .95 cm Compressive force

3.8 cm Figure 2.22. Schematic for lap shear specimen, adapted from Hopkins et al. [5].

2.3.4 Peel Test

The peel test has been utilized in multiple UAM studies [8, 45, 47, 52]. This test is typically applied to adhesively bonded joints, however, it has been modified to test

UAM builds [51]. Testing is performed to British standard BS EN 2243-2:1991 which was designed to determine the strength of structural adhesives on clad metal sheets based on the maximum load a specimen can withstand under peeling action [47]. A schematic of the testing process is shown in Figure 2.23. To create a sample for testing the first layer is bonded to a support plate. The plate is necessary to prevent specimens from flexing in the rollers during loading [45]. A second foil is then joined on top of the first and tested.

Peeling apparatus

Support plate First foil

Roller

Second foil

Apply tensile load to pull the weld apart

Figure 2.23. Schematic of peel test [45].

Two failure modes have been observed during peel testing of UAM builds and are shown with respect to a load vs. displacement curve in Figure 2.24. The first mode fails with small amounts of extension and is characterized by a clear break at the beginning of the weld region. The second mode requires significantly more extension before failure and is characterized by a break further into the welded region with material tearing that gives the impression of „teeth‟ [45].

Clear break at beginning of weld region Break in welded region Load (N) Load with „teeth‟

Displacement (mm) Figure 2.24. Load versus displacement curve with failure modes indicated, adapted from

Kong et al. [45].

A distinct advantage of the peel test over other methods is that only two bonds are needed. This reduces the time and material required for testing and is highly conducive to rapid feedback of bond quality. The drawbacks of this testing process is that bond strength cannot be determined which does not make the results of this test comparable with other bond strength testing methods [53]. Also, the test is limited to low bond strength joints. Kong et al. discovered that peel strength increased with improved welding parameters to a point and then leveled off [8]. The plateau of peel strength did

34 not correspond to microstructural observations of linear weld density. It was determined that a critical peeling load exists and when exceeded will result in failure of the base metal adjacent to the weld region [8].

2.3.5 Push-Pin

Push-pin testing has been suggested as a method for determining UAM bond strength by Zhang et al. [53]. This testing set-up and specimen are schematically shown in Figure 2.25 and consists of pushing a rod through a machined hole that has been drilled to a specific interface in a UAM build. The main advantage of this process is the ability to test a specific interface which can be beneficial when joining dissimilar metals.

While no standard is set, tests have been completed with as few as 16 layers. It is hypothesized that this method can test single foil layers [53]. A finite element simulation is required to determine bond strength, which can then be compared to results from other test methods such as tensile and lap shear. Load versus displacement data is acquired during testing which can be a predictor of failure mode and energy.

b Fixture a All dimensions in mm Baseplate Push-pin

70.0 Ø 10.4 Force 25.4 UAM build

Direction of travel Baseplate Fixture

Figure 2.25. a.) Push-pin testing setup and b.) schematic of push-pin testing sample, adapted from Zhang et al. [53].

The shortcomings with the push-pin test method are similar to those of peel testing. A maximum push-pin load is characterized by a shear failure at the mouth of the machined hole. Otherwise, complete foil delamination occurs and signifies poor bond strength. Like peel testing, shear failure can occur without 100% bonding. This limits the ability of this test to determine mechanical properties of builds with high bond strength. This can be addressed, however, by modifying the diameter of the push-pin as well as the testing area. While it is suggested that single foils can be tested with this method, the shear strength of a single foil must be greater than the strength of the bonded area. Otherwise the push-pin will shear the foil resulting in a failed test.

CHAPTER 3: OBJECTIVES

1. Evaluate the feasibility of point count SEM fracture surface analysis as a method for

determining bond quality.

a) Define fracture surface morphologies.

b) Analyze fracture surfaces with defined technique.

c) Compare results with those of linear weld density and area weld density

measurements.

d) Correlate findings to mechanical testing.

2. Analyze surface modification of Cu foil in dissimilar Al/Cu weld

a) Perform mechanical push-pin testing and compare load vs. displacement curves.

b) Characterize interface and fracture surface morphology.

3. Evaluate the effect of a post-weld heat treatment on similar and dissimilar VHP UAM

builds.

a) Perform mechanical push-pin testing and compare load vs. displacement output.

b) Identify intermetallic phase formation between aluminum and copper welds.

c) Characterize interface morphology in dissimilar joints.

CHAPTER 4: FRACTURE SURFACE ANALYSIS OF ULTRASONIC ADDITIVE

MANUFACTURED AL 3003

4.1 INTRODUCTION AND MOTIVATION

The following chapter describes the microstructural analysis section of a collaborative paper by Hopkins et al. titled, “Optimizing Ultrasonic Additive

Manufactured Al 3003 Properties with Statistical Modeling” published in the Journal of

Engineering Materials and Technology1 [9]. The publication investigates the effects of processing parameters on the ultimate shear strength (USS) and ultimate transverse tensile strength (UTTS) of Al3003 UAM builds. Linear weld density (LWD) and area weld density (AWD) measurements of low magnification fracture surfaces were taken.

The two characterization techniques did not show a consistent correlation between bonded area and mechanical strength.

This discrepancy led to an in-depth investigation into the fracture surfaces of the mechanically tested builds in an attempt to quantify and characterize the features found at the fracture surface. This section will introduce the motivation for performing SEM fracture surface analysis and detail the technique and subsequent results.

1 This section outlines the authors contribution to the published paper, SEM fracture surface analysis, in detail. 38

4.2 BACKGROUND

4.2.1 Linear Weld Density

Linear weld density measurements were performed for each sample that was mechanically tested. Figure 4.1 plots the LWD results versus the mechanical data. It is evident that there is no trend between UTTS, USS and LWD.

Figure 4.1. LWD versus average USS and UTTS of Al3003 UAM specimens.

4.2.2 Area Weld Density

Optical analysis of the fracture surfaces for both USS and UTTS samples was performed. Similar examination has been performed previously with positive results [53,

54]. A representative low magnification fracture surface is shown in Figure 4.2a. The dark regions on the micrograph (Denoted with I) are textured areas considered to represent a bonded region. Conversely, the light areas (region II) have visible machining lines and are considered to be unbounded regions of unaffected material. Image J software was used to analyze the amount of dark region on the fracture surface by applying a fixed threshold to the images, shown in Figure 4.2b. The bonded area in black was then calculated. The results are plotted in Figure 4.3 [9].

Figure 4.2. Optical images of transverse tensile fracture surfaces (top surface) from

Figure 4.3 shows a correlation between percent bonded area and increasing mechanical strength for USS, however, no trend is apparent for UTTS. Due to these findings it cannot be concluded that the dark regions account for complete metallurgically bonded areas. Therefore, SEM was used to gain a better understanding of the fracture surface morphology and explain why there is a significant difference in strength, but not in percent bonded area.

Figure 4.3. Percentage of bonded area on fracture surfaces versus mechanical strength of both UTTS and USS samples.

4.3 RESULTS AND DISCUSSION

In an effort to fully quantify the fracture surface morphology at the interface between UAM foils and correlate fracture morphology to mechanical strength, SEM was utilized. Fifty randomly located SEM images were taken on the fracture surfaces of

Experiment #2 (low LWD, high mechanical properties) and Experiment #8 (high LWD, low mechanical properties). Secondary electron detection at 2000x magnification was used for all micrographs. A 17x20 grid of evenly spaced points was overlaid on each image and every point was characterized. Due to the subjective nature of the analysis, each point was characterized by two independent researchers (Truog and Wolcott). Five different features were identified on the fracture surfaces, original machined surface, ductile failure, shear ductile failure, flow and brittle shear. Figure 4.4 shows a representative fracture surface with individual morphologies indicated; a definition of each morphology is listed below.

(i) Machined Surface: Smooth, unaffected surface of upper foil. Machine

lines normally present with small, dark dots following machine direction.

Indicated no contact between the foil layers.

(ii) Ductile failure: Metallurgical bonding evidenced by the typical circular

cup/cone type fracture.

(iii) Shear ductile failure: Determined to be ductile failure viewed at an angle

not normal to the surface. Hypothesized to occur as the bonds are broken

during tensile testing, foils tend to peel from one another producing ductile

failure at a non-normal angle to the original foil surface.

(iv) Flow: Defined as texture produced due to foil to foil contact without

creating a true metallurgical bond. In this case, peaks of material in the

textured foil are hypothesized to press into the smooth upper foil creating

an impression. The peaks displace material which flows out around the

indentation as it is being pressed into.

(v) Brittle shear: Defined as sheared off region without any ductility

associated with its failure.

Figure 4.4. Example SEM images representing shear ductile failure, ductile failure, flow, brittle shear, and machined surface.

Metallurgical bonding is considered to be atomic bonding of two surfaces.

Ductile and shear ductile failure indicate metallurgical bonding as their fracture morphologies are consistent with this type of bonding. Brittle shear and flow areas represent areas of foil to foil contact that did not bond metallurgically. These areas may

43 have connected with similarly deformed areas on the bottom foil layer. This can lead to mechanical interlock, a type of bonding that is much weaker than metallurgical bonding.

The point counting results from both researchers were averaged and the area fraction of each feature calculated. The results are reported in Table 4.1.

Table 4.1. Area fraction of features measured on fracture surface using SEM

UTTS Machined Shear Ductile Flow Brittle Combined Exp # (Mpa) (%) ductile (%) (%) (%) shear (%) ductile (%) 2 31.5 17.9 13.9 4.2 63.2 0.9 18.1 8 12.7 24.9 11.8 1.5 58.5 3.3 13.3

The results of the area fraction analysis show two important trends. First, the area fraction of machined surface in the low UTTS sample (Experiment 8) is 7% higher than the high strength sample (Experiment 2). This suggests that there was less foil to foil contact compared to the high strength sample. Contact is a recognized stage of bonding necessary to create a solid state joint [3, 55]. Reduced contact will result in fewer joining opportunities along the interface. Secondly, the area fraction of combined ductile fracture in Experiment 2 is about 5% greater than in Experiment 8. This supports the hypothesis that higher mechanical strengths are a result of increased metallurgical bonding. It follows that lower mechanical strength is due to a lower area fraction of metallurgical bonding.

To review, Experiment 2 displayed a USS (16.5 MPa) and UTTS (31.5 MPa).

Experiment 8 had comparatively poor mechanical properties of USS (8.1 MPa) and

UTTS (12.7). It was expected that linear weld density would be higher for Experiment 2

44 than Experiment 8. Instead, it was found that Experiment 2 had a LWD of 67% compared to 85% for Experiment 8. This result showed that linear weld density did not correctly predict mechanical properties based on connected interfacial area.

Area weld density measurements were made based on low magnification fracture surfaces. Again, it was expected that a correlation between percent bonded area and mechanical strength would be found. Testing proved inconclusive and it was determined that the fracture surface exhibited more than two conditions of bonding, either bonded

(black) or unbounded (white). To fully understand the bonded areas and obtain a true bonded area measurement high magnification SEM pictures were taken. Fracture surface point count analysis determined that the high strength sample displayed more ductile and shear ductile regions and fewer machined regions than the low strength sample. This result was consistent with the expectation that improved mechanical properties relied on increased metallurgical bonding.

Further work is necessary to verify the statistical significance of the relationship between metallurgical bond area fraction and mechanical strength. Future applications of the SEM point count method include verifying optimized parameters and evaluating techniques for improving UAM bonding. A major drawback of this technique is the amount of time required for analysis and also the subjectivity inherent in identifying different features.

4.4 CONCLUSIONS

1) Interfacial microstructure analysis including linear weld density (LWD) and area

weld density (AWD) measurements did not correlate the determined bonded area to

mechanical strength.

2) SEM point count fracture surface analysis showed a relationship between bond area

and strength

3) A high area fraction of metallurgical bonding corresponded to a higher ultimate

transverse tensile strength.

CHAPTER 5: SURFACE MODIFICATION OF VERY HIGH POWER

ULTRASONIC ADDITIVE MANUFACTURED ALUMINUM AND COPPER

STRUCTURES2

5.1 INTRODUCTION

There has been interest in using Very High Power Ultrasonic Additive

Manufacturing (VHP UAM) to join dissimilar materials for energy applications [1-3].

However, a major issue associated with multi material consolidation using UAM is poor bond quality [2-5]. VHP UAM was created to address this issue by increasing the maximum amplitude (from 26μm to 52μm) and normal force (from 2.5kN to 33kN).

These increases have allowed for successful initial joining studies of previously difficult to join materials [6].

The main obstacle to bonding is the presence of surface oxides which prevents nascent surface interaction, and therefore, metallurgical bonding. It was previously believed that surface oxides were displaced into the local bulk material during UAM, however, it has been found that the surface oxide layer remains at the interface and is not significantly dispersed during processing [43]. An added degree of difficulty exists when joining dissimilar materials. Differences in mechanical properties will lead to unequal

2 Submitted to Trends in Welding Research Conference 47 amounts of deformation at the interface and can significantly affect the joint morphology

[56].

Similar obstacles are faced in other solid state bonding processes. To overcome these barriers, methods have been proposed to enhance bonding. The application of these methods to the UAM/VHP UAM process may significantly improve bonding and allow for increased application in a variety of industries. Some of the more common attempts at improving solid state bonding have included heat treatment [40], surface cleaning [12] and modifying surface roughness at the joining interfaces [15, 47, 57].

Heat treatment of roll-bonded aluminum and copper metal laminates was studied by Heness et al. [40]. It was hypothesized that a mechanical bond between metals is formed during rolling and that a sintering heat treatment would then produce a strong metallurgical bond. Laminates were heat treated at 450°C for increasing lengths of time.

It was found that peel strength increased during the formation of CuAl2 and Cu9Al4 intermetallics, while the formation of CuAl was associated with a decrease in peel strength.

Surface preparation is an important and necessary procedure when cold roll bonding [12]. Mechanical and/or chemical cleaning is performed immediately before welding to remove oxides, grease and dust. The effectiveness of this method has been proven for roll bonding and may be of great benefit if implemented in UAM.

Surface roughness has been studied for its effect on roll bonded structures.

Copper sheets were scratch bushed and then joined. It was found that peel force increased with increased surface roughness (Ra) [57]. A similar result was discovered

48 with diffusion bonding of Al6061 [15]. Two different grits of SiC paper were used to roughen the weld interfaces before bonding. The rougher surface was found to exhibit a markedly higher Ultimate Tensile Strength (UTS) than the smooth sample for all holding times [15].

The effect of sonotrode roughness on interfacial bond strength has been studied for UAM builds [47]. Welding was performed with both a worn and newly machined sonotrode with surface roughness of 3.44 Ra μm and 6.26 Ra μm respectively. Peel testing revealed that the newly machined sonotrode produced builds with higher peel strength compared to the worn sonotrode [47].

A previous study of dissimilar ultrasonic welding found that the softer of the two materials flows around the surface topography of the harder material [56]. A UAM study of aluminum and titanium bonding found that the asperities on the titanium surface did not undergo sufficient deformation to allow for nascent surface creation due to the hardness mismatch of the soft aluminum and harder titanium [5]. This study will focus on the effect of surface roughness modification of the foil interfaces in an effort to induce plastic deformation in the copper foil and promote metallurgical bonding across the interface. Push pin testing, fractal analysis and point count fracture surface analysis will be applied to determine the effect of surface roughness on aluminum and copper VHP

UAM builds.

5.2 EXPERIMENTAL PROCEDURE

Al3003 and Cu110 were joined in this study on an Al6061 base plate using a test bed VHP UAM machine provided by the Edison Welding Institute (EWI). The test bed featured a VHP UAM dual transducer design mounted to a modified 3-axis milling machine [6]. There is no tape feeding mechanism on the test bed and thus all foils were taped to the base plate and pulled taut with pliers during welding.

All builds made for this study were 14 total layers in the “Cu on top” geometry.

This consists of 7 layers of aluminum foil joined to the base plate with seven copper foils then joined to the aluminum. The single copper foil at the aluminum/copper interface was modified on the surface by rolling the sonotrode over the foil with 3.1 kN and no ultrasonic energy. The modified surface was then placed faced down on the aluminum build and welded as normal. The surface roughness of aluminum and copper foils was measured with a handheld stylus profilometer. The results are reported in Table 5.1. All welds were made with the same processing parameters: oscillation amplitude (34 μm), welding speed (35.5 mm/s) and normal force (5.5 kN). These parameters were modified from a previous VHP UAM study and have not been optimized for bond strength [7].

Table 5.1. Measured surface roughness values for aluminum and copper foils.

Material Ra (μm) Al3003 foil 0.061 Cu110 foil 0.175 Cu110 foil rolled with sonotrode 1.170

Push-pin testing, a method created for testing laminate structures by Zhang et al. was used to gather force vs. displacement data [53]. Both the maximum force to failure and area under the curve were determined. This testing method has been previously outlined in detail [53, 58]. Figure 5.1 illustrates the testing setup in the Gleeble™ 3800 thermal-mechanical simulator. Figure 5.2 shows a schematic of a push-pin sample with pertinent dimensions included. An extension rate of 0.2 mm/second was used.

Figure 5.1. Push-pin testing setup in Gleeble™

b Fixture All dimensions in mm

70.0 Ø 10.4 Force 25.4

Baseplate Fixture

Figure 5.2. Schematic of push-pin testing sample

To analyze the interface geometry between the as-welded and surface modified builds, multi-dimensional fractal analysis was used. In fractography, the fractal dimension (D) has been proposed as a parametric descriptor of interface/surface tortuosity [59, 60]. While the fractal dimension has been linked to surface complexity it is necessary to understand what the appropriate means of measuring the fractal characteristics of a fracture surface are and also how to interpret the data.

When measuring an interface, the measured length is directly related to the distance between measurements [61]. This concept is pictured in Figure 5.3. As the 52 segment length, η, decreases, the overall length and thus, apparent roughness (RL) will increase [59]. This concept is characterized by the Mandelbrot-Richardson equation:

R (η)=R η1-D L 1 L

Equation 1. Mandelbrot-Richardson equation where RL= total length, η= measuring segment length, R1= profile roughness parameter constant and DL= the fractal dimension.

Log-log plots of total length as a function of measuring segment length are known as “fractal plots”. A representative plot is shown in Figure 5.4 where the slope, α, yields

DL. Fractal analysis can be extended to multi-fractal analysis by the use of weighted distortion factors from which reliable measures of the fractal dimension, Ds, can be determined.

Segment length: η = η1

Segment length: η = η2

Segment length: η = η3

η1 < η2 < η3

RL(η1)> RL(η2) > RL(η1)

Figure 5.3. Schematic drawing indicating a decrease in apparent roughness, RL, correlating to an increase in segment measure length, η [59].

α =1-DL log R log

log η

Figure 5.4 Fractal plot of profile roughness. DL, the fractal dimension is measured from the slope, α [59].

A previous study by Lee et al. compared the Joint Roughness Coefficient (JRC), a measure of fracture surface roughness of rock fractures, to the fractal dimension. A guideline to the JRC is pictured in Figure 5.5. This method is based on visual comparison between representative fracture surfaces and the surface to be analyzed. The fractal dimension for each standardized JRC line was calculated and is shown next to each line. It is immediately evident that the fractal dimension for all lines is very close to

1, however, the fractal dimension increases predictably with rougher lines. This indicates that the fractal dimension measure is significant despite very small variations. 55

FD 1.00044 1.00168 1.00285 1.00397 1.00441 1.00564 1.00710 1.00805 1.00958 1.01343

Figure 5.5 Joint Roughness Coefficient (JRC) vs. fractal dimension, adapted from Lee et al. [60]

This investigation used multi-dimensional fractal analysis to compare the interface roughness between as-welded and surface modified builds. Details of this method have been previously investigated by multiple authors [62-64]. 10 SEM images at 250X magnification were taken of each sample at equidistant intervals across the interface.

Each picture was opened with Adobe Photoshop™ and the image was digitized by tracing the interface with a stylus on a tablet PC. The resultant black line on white 56 background was then analyzed using imageJ plug-in, FracLac 2.5. The resulting fractal dimension for each interface was used for comparison.

Fracture surface analysis was used to determine the amount of metallurgical bonding on the copper side of fractured push-pin samples. This technique was adopted from a previous study by Hopkins et al. that analyzed UAM builds with different mechanical properties, linear weld density (LWD) and area weld density (AWD) [9]. The analysis consists of a manual point count of fifty SEM micrographs from both the as-welded and surface modified fracture surfaces. A 19x17 grid was overlaid and each point was categorized as either the original machined surface, ductile fracture or flow morphology.

Definitions of these morphologies are listed below along with a representative micrograph depicting the different fracture morphologies in Figure 5.6.

(i) Machined Surface: Smooth, unaffected surface of upper foil. Machine lines

normally present with small, dark dots following machine direction. Indicated no

contact between the foil layers

(ii) Flow: Defined as texture produced due to foil to foil contact without creating a

true metallurgical bond. In this case, peaks of material in the textured foil are

hypothesized to press into the smooth upper foil creating an impression. The

peaks displace material which flows out around the indentation as it is being

pressed into.

(iii)Ductile: Metallurgical bonding evidenced by the typical circular cup/cone type

fracture

Figure 5.6. Representative SEM micrograph indicating machined, ductile and flow morphologies.

5.3 RESULTS AND DISCUSSION

Copper on top builds in the as-welded and surface modified conditions were mechanically tested using the push-pin method. The maximum load to failure for each build is plotted in Figure 5.7. It is evident that there is significant variability between tests. Due to the scatter of this testing parameter no conclusions of overall bond quality can be made. A more in depth study with optimized process parameters is required to evaluate the overall mechanical properties of these builds. Despite the lack of definitive mechanical data, load versus displacement curves from each test can be compared and related to interface and fracture morphology.

Figure 5.7. Mechanical push-pin test results for the as-welded and surface modified condition

The load versus displacement curves for all five builds are shown in Figure 5.8.

The area under the curve was determined and is reported as percentages below the peak of each curve. The percentages reflect the relative area under the curve normalized to the curve with the largest area, represented by 100%. The two surface modified builds show a more prolonged failure after reaching the peak load. This is evidenced by the relatively high area under the curve compared to the as-welded samples. The results are most 60 striking when comparing the as-welded and surface modified builds with similar peak loads. These curves are represented by dotted lines. With similar peak loads the build with surface modification displays more than twice the amount of area under the curve compared to the as-welded build. This indicates that the surface modified build required more energy to fracture.

63.8%

95.6% 100%

74.3%

32.5%

Figure 5.8. Force vs. Displacement graph for as-welded and surface modified builds.

Percentages reflect relative area under curve normalized to the largest area (100%).

A hypothesis has been created to explain the difference in energy required for total failure between the two build conditions. Figure 5.9 is a schematic representation of delamination failure after the peak load is achieved and will aid in this discussion. It is expected that the peak load is a function of the amount of metallurgical bonding present at the interface [9]. Due to the scatter in mechanical test data the load versus displacement curves cannot offer any insight into the relative amounts of metallurgical bonding. It follows that mechanical interlock is not responsible for the peak load, yet it must be overcome by deformation before complete failure can occur. It is hypothesized that during peak loading a significant portion of the metallurgical bonds fail. As extension continues, the load is transferred to mechanical interlock which continues to deform until complete failure. Figure 5.9 depicts the micro scale fracture after peak loading, where metallurgical bonds have already failed during peak loading and mechanically interlocked areas of the material are continuing to deform as the push-pin extends.

This realization then allows for the conclusion that surface modification increases the total amount of mechanical interlock allowing for a more energetic failure. Fracture surface analysis and fractal analysis were performed to further investigate the hypothesis of increased mechanical interlock among surface modified builds.

Interface condition Mechanical Interlock after peak load, before complete Machined Surface Ductile Fracture failure

Cu Al Fracture during testing

Push-Pin

Figure 5.9. Schematic representation of delamination failure after peak load is achieved.

To determine if the surface modification completed before welding affected the interface morphology, fractal analysis was performed. The value of the fractal dimension

(FD) has been proven to be directly proportional to interface roughness [60]. The fractal dimension will indicate how closely the interface can be modeled as a 1-D (FD=1.0) or 2- 63

D (FD=2.0) image. Ten low magnification pictures were taken across the entire interface length of both samples. Examples of the interfaces are pictured in Figure 5.10. The interfaces were digitized and then analyzed using imageJ plug-in FracLac 2.5.

Figure 5.10. SEM micrographs of the a.) as-welded and b.) surface modified interface.

Table 5.2 reports the averaged fractal dimension and standard deviation from the ten measurements. It is apparent that the as-welded sample has a very smooth interface

(FD= 1.0060) compared to the surface modified sample (FD=1.0230). The surface modified interface has a larger standard deviation because some parts of the interface are very flat, giving values similar to the as-welded condition, while others display a much wavier interface, as seen in Figure 5.10b. The appearance of increased roughness areas on the surface modified interface is associated with the modification of the copper foil. 64

Figure 5.10b shows indents into the copper foil from the sonotrode as well as flow regions that protrude into the aluminum foil. It is plausible that this flow structure, created by surface modification, provides both macro and micro mechanical interlock. It would follow that an interface with increased tortuosity may provide more opportunities for mechanical interlock, resulting in a more graceful failure.

Table 5.2. Average fractal dimension for the as-welded and surface modified interfaces.

Sample As-welded Surface Mod. Average FD 1.0060 1.0230 St. Dev. 0.0032 0.0179

Previous work has been done to characterize the fracture surface of a UAM bonded structure [9]. This study has laid the ground work to quantify specific fracture morphologies, such as: flow, ductile, shear ductile, brittle shear and machined. For this study only three morphologies were found: flow, ductile and machined.

A 19x17 grid was placed over 50 SEM images taken from random locations on both the as-welded and surface modified fracture surfaces. Each point was then characterized with the averaged results displayed in Table 5.3. It is expected that a ductile failure mode is representative of a metallurgical bond [9]. Thus, the amount of metallurgical bonding at the interface can be determined based on the relative amount of ductile fracture. The amount of ductile fracture was similar at 7.5% and 8.8% for the as- welded and surface modified builds respectively. The amount of machined and flow regions between the two builds were quite different. The surface modified build had

65 significantly more flow regions and fewer machined areas than the as-welded sample.

This finding supports the hypothesis that surface modification increased the amount of mechanical interlock as more flow regions would indicate increased aluminum and copper interaction.

Table 5.3. Area fraction of fracture morphologies measured from SEM images.

Morphology (%) Machined Flow Ductile As-welded 72.1 20.4 7.5 Surface Mod. 61.5 29.8 8.8

Low magnification fracture surface images of the copper foil are shown in Figure

5.11. There are noticeable divots on the surface modified surface that were created from the sonotrode during texturing. These areas appear to attract flow and ductile failure

(indicated by the lighter patches). The as-welded sample displays a more even distribution of flow and ductile areas. This indicates that surface modification may create areas that preferentially bond compared to the machined surface.

Figure 5.11. Low magnification fracture surface SEM image for a.) as-welded and b.) surface modified builds.

This finding yields the possibility that a more optimized surface modification technique may improve mechanical interlock and metallurgical bonding. Future work may also focus on combining surface modification and a post processing heat treatment.

Surface modification reduced the amount of machined surface which could provide improved intimate contact. This would lead to diffusion across the interface during heat treatment which has been shown to improve bond quality in a previous VHP UAM study

[65].

5.4 CONCLUSIONS

1. Peak force values obtained by push-pin testing of both as-welded and surface

modified builds was highly variable. While the processing parameters for the

different conditions were the same, they have not been optimized.

2. Surface modification was found to create a more energetic failure which was verified

by the area under the curves in the load vs. displacement graph. An increase in

mechanical interlock is hypothesized to be responsible for this difference.

3. Fractal analysis performed on the interfaces of both builds indicated a more tortuous

bonding path for the surface modified weld. A fractal dimension of 1.006 was

calculated for the as-welded sample and 1.023 for the surface modified.

4. Fracture surface analysis determined similar amounts of ductile failure for the as-

welded and surface modified builds of 7.5% and 8.8% respectively.

5. The as-welded fracture surface exhibited more machined and less flow regions than

the surface modified build. This difference may indicate more mechanical bonding in

the surface modified build which is responsible for a more energetic failure as

evidenced by the load vs. displacement curves.

CHAPTER 6: HEAT TREATMENT OF VERY HIGH POWER ULTRASONIC

ADDITIVE MANUFACTURED (VHP UAM) ALUMINUM AND COPPER

STRUCTURES

6.1 INTRODUCTION

Dissimilar metal joining using Very High Power Ultrasonic Additive

Manufacturing (VHP UAM) has many applications in the aerospace, automobile, ship building, nuclear, electronics and industrial machinery industries [1]. Despite the abundance of applications, the capabilities of Ultrasonic Additive Manufacturing (UAM) to join dissimilar materials are limited [2-5]. VHP UAM was developed to address this shortcoming by increasing the maximum amplitude (from 26μm to 52μm) and normal force (from 2.5kN to 33kN) so that harder and stronger materials could be processed [6].

Numerous UAM feasibility studies have been conducted for dissimilar metal welding; however no mechanical testing has been performed to correlate the microstructural findings to bond quality. In this paper aluminum and copper foils were joined using VHP

UAM and mechanically evaluated. The results were then correlated to microstructure.

Dissimilar material joining allows for multiple materials to be used in the creation of a finalized part. The benefits of multi-material structures include improved mechanical properties, reduced density, corrosion resistance as well as improved electrical and thermal conductivity [53, 66]. However, there are a variety of issues associated with the welding/joining of dissimilar metals. Intermetallic compound (IMC) formation, high residual stress and cracking are all problems commonly associated with

69 fusion processes [41, 67]. Due to the advantages of low heat input during welding, solid state processes are commonly chosen to weld multiple materials structures.

Al3003 has been the focus of many UAM studies due to its low hardness, availability and cost. For this study Al3003 was chosen because it has been so well documented in UAM research. Copper 110 was paired with aluminum because this material combination has been previously studied with many other solid state processes

[32, 33, 36, 40, 41, 68]. There are numerous practical applications of aluminum/copper dissimilar joints. Examples include solar collectors, semiconductor heat sinks, bimetallic heat exchanger tubes [2], high direct-current bus systems [32], as well as micro-electronic devices [69]. Joining aluminum and copper can be difficult even for solid state processes due to the formation of intermetallic compounds which can nucleate and grow when temperatures eclipse 120°C [32, 70]. These compounds are brittle and can be deleterious to mechanical and electrical properties [32]. For this reason, low heat input solid state processes such as cold rolling and ultrasonic welding are commonly used to avoid intermetallic formation when joining aluminum and copper.

Most UAM studies have been focused on the joining characteristics of Al3003 [2-

5, 9, 45, 51-53, 71]. Of the few papers published pertaining to dissimilar metal welding using UAM, the focus is on feasibility of joining multiple materials [2-5, 72]. Janaki

Ram et al. created dissimilar joints between Al3003 and four other foil materials: Al

2024, Inconel 600, Brass (Cu-30Zn) and SS 347 [3]. The weld interfaces were characterized using optical and scanning electron microscopes. These images were used to make determinations on bond quality based on physical discontinuities at the interface.

This method is sufficient for an introductory investigation of material and process feasibility; however, the performance of these joints has not been addressed.

Hopkins et al. performed a statistical analysis of grade 1 titanium and 1100-O aluminum composites joined with UAM [5]. The goal of the research was to identify the relationship between manufacturing parameters and bond quality. A Taguchi L16 orthogonal array was used for the experimental design which modified four parameters: normal force, amplitude, weld speed and foil layers. Interaction plots were created that detailed the effect of each parameter on ultimate transverse tensile strength (UTTS) and ultimate shear strength (USS). Microstructural analysis verified that the softer aluminum deformed to the surface topography of the harder titanium, a phenomenon originally observed by Joshi [56]. While conclusions between weld parameters and mechanical properties were determined, there was little emphasis on correlating the mechanical properties to microstructural effects.

A previous study by Janaki Ram, Johnson and Stucker investigated the possibility of fabricating multi-material parts out of aluminum and copper with the UAM process

[2]. SEM micrographs showed interfaces that were free from large physical discontinuities at low magnification. Upon closer inspection a porous layer on the Al side of the bond was visible, it was hypothesized that a difference in interdiffusivity was the root cause of these pores. Bonding was characterized by visual inspection of SEM secondary electron images where the aluminum on top configuration appeared “tighter” compared to the copper on top build. SEM EDS line scans were performed in the as- welded condition and it was concluded that there was considerable diffusion from

71 aluminum into copper (.3 to 1.7 Wt.%). One issue with SEM EDS is that the diffusion field at the interface is small compared to the electron interaction volume [73]. These results must then be considered qualitative. This initial study into Al/Cu bonding with the UAM process showed feasibility and also raised many questions. Mechanical testing of the aluminum and copper interface is needed to verify the bond quality.

Although feasibility studies are integral to the VHP UAM process or any new joining process, a more in depth investigation into the relationship between mechanical properties and microstructure at the bond interface is crucial to accelerate the deployment of VHP UAM. This study will detail the bond interface between aluminum and copper

VHP UAM builds using characterization techniques such as X-ray diffraction (XRD), fractal analysis and high resolution scanning electron microscopy (HR SEM) and correlates the findings to mechanical test results.

6.2 EXPERIMENTAL WORK

6.2.1 VHP UAM of Al3003 and Cu110

In this study Al3003 and Cu110 foils were bonded to an Al6061-T4 base plate.

The compositions and dimensions of these materials are listed in Table 6.1. The process parameters used for both Cu110 and Al3003 joints were: oscillation amplitude (34 μm), welding speed (35.5 mm/s) and normal force (5.5 kN). The parameters were not optimized for bond strength and were kept constant so that welding parameters did not provide an interaction with the mechanical testing. The parameters were modified from a

72 study conducted by Sriraman et al. found to produce satisfactory welds between Cu110 foils [7].

Table 6.1. Specification of materials used in this investigation

Material Nominal composition (Wt. %) Thickness Al 3003 (H18) Al-1.2Mn-.12Cu 150 μm thick foil Cu 110 99.95Cu 150 μm thick foil Al 6061 (T4) Al-1Mg-.6Si-.3Cu 1.27 cm thick plate

All UAM experiments were conducted using the Very High Power Ultrasonic

Additive Manufacturing (VHP UAM) test bed at the Edison Welding Institute (EWI).

The test bed features the VHP UAM dual transducer design mounted to a modified 3-axis milling machine. There is no tape feeding mechanism on the test bed and thus all foils were taped to the base plate and pulled taut with pliers during welding. Four different geometries were created and tested, each consisting of 14 total layers. The geometries,

“All Cu”, “All Al”, “Cu on top” and “Alternating Al/Cu” are described in Table 6.2.

Table 6.2. Configuration of build geometries

Geometry Configuration of 14 layer build All Al Baseplate / 14 layers Al3003 All Cu Baseplate/ 14 layers Cu110 Cu on top Baseplate/ 7 layers Al/ 7 layers Cu Alternating Baseplate/ 1 layer Al/ 1 layer Cu/ 1 layer Al…

6.2.2 Heat Treatment

Samples were heat treated in a Lucifer box furnace at 350°C for 10 minutes. All samples were placed on a metal grate and heat treated at the same time so that all samples were subjected to the same heating and cooling cycles. No shielding gas was used and samples were air cooled. A K type thermocouple was attached to a sample in the middle of the furnace to obtain a representative actual sample temperature.

6.2.3 Mechanical Push-Pin Testing

Push-pin testing, a method created for testing laminate structures by Zhang et al. was used to determine the maximum force to failure for VHP UAM builds [53]. Figure

6.1 shows the testing setup using the Gleeble™ 3800 thermal-mechanical simulator.

Sample preparation for the push-pin test consists of creating a constant sampling area that is 25.4 mm x 25.4 mm on a base plate that measures 25.4 mm x 72 mm. A 10 mm diameter steel push rod was used to apply force to the foil interface through an oversized

10.4 mm hole drilled through the base plate and to the desired interface. A displacement rate of .2 mm/s was used.

b Fixture a All dimensions in mm Baseplate Push-pin

70.0 Ø 10.4 Force 25.4 UAM build

Direction of travel Baseplate Fixture

Figure 6.1. a.) Push-pin testing setup in Gleeble™ and b.) schematic of push-pin testing sample. Adapted from Zhang et al. [53]

Samples were sectioned, mounted, polished and then optically measured to determine the drilling depth necessary to test the desired interface. The interface between the 7th and 8th layer was tested for all geometries. A force vs. displacement curve was created for each sample. The maximum force was then used to compare builds. Three builds of each geometry were made for repeatability

6.2.4 Metallography

A representative sample of each build geometry was investigated using metallographic techniques. Builds were sectioned perpendicular to the rolling direction and mounted in a 1 ¼” puck of Konductomet™. Samples were ground with silicon carbide paper through 800 grit and then polished using 9μm, 6μm, 3μm and 1μm

75 diamond paste on Kempads™. Vibratory polishing for 6 hours was the final step. The interface was examined using optical microscopy, scanning electron microscopy (SEM) and X-ray diffraction (XRD).

6.2.5 X-ray Diffraction

A heat treated and mechanically tested sample was analyzed using X-ray diffraction. The sample delaminated along the Al/Cu interface and both sides were analyzed using a XDS 2000 Scintag Inc. diffractometer, in a Bragg-Brentano diffraction geometry with Cu Kα radiation. The samples were cut to 25mm x 27mm surface dimensions. Scans were taken from 10 to 90 2θ to record low index peaks and then analyzed to determine phase content of the samples.

6.2.6 Fractal Analysis

Multi-dimensional fractal analysis was used to compare the surface roughness of heat treated and non-heat treated builds. 10 SEM images at 10,000X magnification were taken of each sample at equidistant intervals across the interface. Each picture was opened with Adobe Photoshop™ and the image was digitized by tracing the interface with a stylus on a tablet PC. The resultant black line on white background was then analyzed using imageJ plug-in, FracLac 2.5. The fractal dimension output from the multi-fractal analysis was then used for comparison.

6.3 RESULTS

6.3.1 Mechanical Testing of Builds

All five geometries were tested in the as welded and heat treated condition. The maximum force required during push-pin testing was determined and is reported in

Figure 6.2. The results indicate that the maximum force required for failure increased between 23% and 49% for the different geometries. This data shows that heat treatment of Al/Cu VHP UAM builds improves the maximum push-pin force needed for failure.

+43%

+49%

+23% +40%

Figure 6.2. Maximum force recorded for four different VHP UAM build geometries.

Percentages indicate amount of improvement in heat treated samples.

Load versus displacement curves were gathered for all tested samples. Curves for the alternating aluminum and copper build in the as-welded and heat treated condition are shown in Figure 6.3. It is immediately apparent that the area under the curve for the heat 77 treated samples is larger than that for the as-welded builds. The first peak for each build is within a very close range (1100N-1500N), however, the heat treated builds continue to increase load after this first peak whereas the as-welded samples tend to decline to failure. To gain a better understand for why this discrepancy in loading was occurring the builds were examined after testing and are shown in Figure 6.4.

Figure 6.4 a) is a typical failure in the as-welded condition. Delamination failure occurred with no noticeable deformation in the fractured sample. The as-welded sample that is drawn with a dotted line is shown in Figure 6.4 b). This individual sample did not follow the same failure load path as the other as-welded samples and began to increase loading after the initial peak, similar to the heat treated samples. Examination of the build after testing revealed a more deformed surface than the other two as-welded samples. A representative heat treated sample is shown in Figure 6.4 c). The top section of the build is deformed considerably.

Al 3003 H18 foil was used in this study. H18 indicates the fully hardened condition for this material. It is expected that heat treatment at 350°C for ten minutes induces recovery in the material, reducing strength and increasing ductility. This change in ductility accounts for the increased extension before failure in heat treated samples.

The improvement in maximum load for heat treated builds will be discussed in the following section.

100%

85.9%

69.2%

41.2%

18.6%

18.7%

Figure 6.3. Force vs. Displacement graph for as-welded and heat treated builds.

a) b)

Figure 6.4. Alternating aluminum and copper builds after push-pin testing in the a) as- welded b). as-welded with deformation (denoted by the dotted line in Figure 6.3) and c) heat treated condition.

6.3.2 Microstructural Analysis

Figure 6.5 a) pictures the interface between Al and Cu in the as-welded “Al on top” geometry with no discontinuities. The SEM image was taken with the backscatter electron (BSE) detector. Figure 6.5 b) pictures an “Al on top” build in the heat treated condition. The formation of intermetallic compounds at the interface can be visually confirmed. Contrast in the image is formed as a result of atomic number “Z” contrast.

The three shelves of intermetallics in Figure 6.5 b) change intensity based on their

80 composition, the brightest shelf is a copper rich intermetallic, while the darkest is aluminum rich.

a Al b

Figure 6.5. The effect of heat treatment on a dissimilar aluminum and copper VHP UAM interface. a.) As-welded interface. b.) Interface after 10 minute heat treatment at 350°C.

X-ray diffraction (XRD) was completed on both sides of a heat treated sample that delaminated at the Al/Cu interface. A peak height phase analysis was conducted to identify the phases on either side of the fracture and is pictured in Figure 6.6. On the copper side of the fracture surface copper and a significant amount of Cu9Al4 was found.

The aluminum fracture surface showed considerably more intermetallic formation.

Aluminum, CuAl2, CuAl and a small amount of Cu9Al4 were discovered. The only intermetallic common between the two sides was Cu9Al4. This suggests that the fracture occurred in the Cu9Al4 layer. The aluminum and copper equilibrium phase diagram in

Figure 6.7 shows CuAl2, CuAl and Cu9Al4 are all stable phases. The formation of CuAl2,

CuAl and Cu9Al4 intermetallics is consistent with literature [17, 33, 40, 41] 81

Al2Cu AlCu

Al4Cu9 Cu

) Al2O3 a.u

Intensity ( Intensity Cu side

Al side

Figure 6.6. X-Ray diffraction patterns from both copper and aluminum side of fractured heat treated sample.

Al Al

4 3

Cu Cu CuAl CuAl

Cu rt

Figure 6.7. Aluminum and copper equilibrium phase diagram

High Resolution Scanning Electron Microscope (HR SEM) images were taken with a backscatter detector to view the fracture path of a heat treated Al/Cu build. The atomic number “Z” contrast indicates three distinct intermetallic layers, shown in Figure

6.8. Fracture is shown to occur between CuAl and Cu9Al4 at the Cu/Al interface. This result is inconsistent with XRD analysis which found the Cu9Al4 intermetallic on both sides of the fracture surface. One possibility is that parts of the Cu9Al4 intermetallic layer fractured and stuck to the CuAl layer. A very small fracture region is shown in Figure

6.8 and is not representative of the entire interface. Serial sectioning could be performed in the future to verify this assumption.

CuAl2 CuAl

Cu9Al4 Cu

Figure 6.8. Fracture interface of heat treated Al/Cu build taken with high resolution

SEM. Fracture occurred in Cu9Al4 intermetallic layer.

6.3.3 Fractal Dimension Analysis

To determine if there was a relationship between interface roughness and mechanical properties, fractal analysis was conducted on the interface morphologies of both the heat treated and non-heat treated “Alternating Al/Cu” builds. The value of the fractal dimension (FD) has been proven to be directly proportional to interface roughness

[60]. In this study, the fractal dimension of both the heat treated and non-heat treated build interfaces were determined. Ten high magnification pictures were taken across the 84 entire interface length, an example is shown in Figure 6.9a. The interfaces were digitized and then analyzed using imageJ plug-in FracLac 2.5. An example of the digitized interface is pictured in Figure 6.9b. The ten resulting fractal dimensions for each interface were then averaged. The results are reported in Table 6.3.

a Al b

Figure 6.9. High magnification SEM backscatter image of heat treated interface. a.)

Original image. b.) Modified image with digitized interface.

Table 6.3. Average fractal dimension for various interfaces. The lowest FD is found at the fractured layer between CuAl-Cu9Al4. The largest FD is found between the Al-

CuAl2 layers, significantly higher than the non-heat treated Al/Cu interface.

Sample No HT 10 min. Heat Treatment

Interface Al-Cu Al-CuAl2 CuAl2-CuAl CuAl-Cu9Al4 Cu9Al4-Cu Average FD 1.029 1.077 1.047 1.034 1.037 St. Dev. 0.015 0.030 0.023 0.015 0.012

6.4. DISCUSSION

6.4.1 Surface Roughness Evaluation

It is clear that applying a heat treatment for ten minutes at 350°C to a VHP UAM aluminum and copper build improves the maximum push pin-force compared to as- welded structures. After performing microstructural analysis, the major difference between the heat treated and non-heat treated condition is the formation of intermetallic compounds. To explain if intermetallic formation is responsible for the improvement in push-pin force, fractal analysis was conducted to analyze the interface roughness.

XRD and HR SEM data suggest that the heat treated samples failed at the AlCu andCu9Al4 boundary. This interface exhibited the lowest fractal dimension of 1.034, indicating that this is the smoothest of the four visible interfaces formed after heat treatment. Crack deflection theory can be applied to the interface and may explain the reason why fracture occurred at this particular intermetallic layer. Interface roughness is the dominating parameter controlling crack propagation [74]. A rougher surface increases the crack path length and crack deflection; both of these factors require additional energy for a crack to propagate compared to a smooth interface. This theory can be used to explain why failure occurred in the smoothest interface. With the shortest crack path, the CuAl and Cu9Al4 interface required the least amount of energy to fracture.

The fractal dimension of the Al/Cu interface in the non-heat treated build and the

AlCu/Cu9Al4 boundary in the heat treated build are similar at 1.029 and 1.034 respectively. Therefore, the surface roughness theory for intermetallic interface failure

86 cannot be applied to explain the significant difference in mechanical properties between the heat treated and non-heat treated builds. This suggests that there is a more prominent mechanism that can rationalize the difference in mechanical properties between the two conditions, such as an improvement in metallurgical bonding.

6.4.2 Metallurgical bonding

Mechanical shear and tensile testing by Hopkins et al. discovered that the best mechanical properties achieved by UAM builds made of Al3003 were just 18% the USS and 16% the UTTS of the bulk material properties [9]. This establishes that while UAM builds may seem well bonded from optical interfacial pictures, in fact a very low percentage of metallurgical bonding is occurring between the surfaces, resulting in poor mechanical properties. The push-pin results in Figure 6.2 show a similar trend for VHP

UAM builds, where as-welded “all aluminum” and “all copper” builds exhibited only

40% and 54% of the bulk maximum push-pin force, respectively. Heat treatment of these VHP UAM builds shows a marked improvement. The heat treated “all aluminum” build had 60% of the bulk maximum push-pin force, a 20% improvement compared to the as-welded condition. The heat treated “all copper” builds yielded 78% of the bulk copper properties, a 24% improvement. It is hypothesized that heat treatment is responsible for the dramatic improvement in mechanical properties. While similar material combinations will not form intermetallics, the intermetallic formation in heat treated dissimilar metal welds confirmed that diffusion takes place across the interface.

The amount of metallurgical bonding is improved from heat treatment, and despite the

87 formation of intermetallics in dissimilar builds, an overall net gain is observed in bond strength.

Intermetallic formation is commonly assumed to be deleterious to mechanical properties and avoided in all circumstances [41]. The findings in this paper and others show specific instances where controlled intermetallic formation does not degrade mechanical properties, and has even been shown to improve the final microstructure [32,

40, 75, 76]. A ductile to brittle transition as intermetallic width increases past a critical width has been proposed to explain this phenomenon, however, the mechanism behind this theory has not been investigated. A theory for mechanical property improvement after intermetallic formation at the interface will be proposed in this work. It is understood, however, that an additional study is necessary to verify these assumptions.

The interfacial geometrical configuration of heat treated VHP UAM builds is very similar to that of brazing, where a thin layer of braze filler metal exists at the interface of a joint. It has been observed in brazing that a brazed sample can fail above the yield strength of the braze filler metal. This is attributed to a complex multi-axial stress state of the braze filler metal between the joined base materials. High levels of constraint do not allow the filler metal to deform plastically and account for the higher than normal yield strength, this situation is depicted in Figure 6.10a. A similar situation may be responsible for the improved mechanical strength in VHP UAM heat treated builds. The thin intermetallic layer formed at the interface is constrained during loading. Coupled with improved metallurgical bonding between the two foils, a net bond strength improvement is realized.

As the intermetallic width increases past the critical width, a condition depicted in

Figure 6.10b, the amount of restraint decreases, therefore altering the stress state such that deformation is no longer inhibited. The intermetallic layer can then deform independently from the bulk structure resulting in failure at lower loads.

P P σy a b σy σx σx

σy σy

P P

Load versus displacement curves that were captured during push-pin testing of the as-welded and heat treated “alternating aluminum and copper” show a significant increase in area under the curve and maximum failure force in the heat treated condition.

Analysis of the fractured samples found increased amounts of deformation in the heat treated builds. This was attributed to recovery in the cold worked aluminum 3003 H18 foil after heat treatment. It is suggested that a trade-off between the improvement of

89 metallurgical bonding due to heat treatment and the reduction of mechanical properties of the foil resulted in an overall improvement during mechanical testing.

6.5. CONCLUSIONS

1. Heat treated Al/Cu VHP UAM builds had improved mechanical push-pin

properties (23%-49% improvement) compared to as-welded builds.

2. Heat treated Al/Cu VHP UAM builds displayed more area under the load versus

displacement curve than as-welded builds. The curves were compared to the

fracture samples. The heat treated builds underwent significant deformation

compared to the as-welded builds, confirming a more energetic failure predicted

by the load versus displacement curves.

3. Fractal analysis indicated that the Cu9Al4/CuAl interface had the smoothest

interface in the heat treated sample (FD = 1.034). This is the interface that failed

during testing.

4. Fractal analysis concluded that changes in interface morphology did not account

for the improved mechanical properties in the heat treated condition. The fractal

dimension of the aluminum/copper interface without heat treatment was 1.029,

while the Cu9Al4/CuAl interface displayed a fractal dimension of 1.034.

5. X-ray diffraction analysis found that three intermetallic compounds were formed

during a 10 minute heat treatment at 350°C: CuAl2, CuAl and Cu9Al4.

CHAPTER 7: CONCLUSIONS

1) SEM point count fracture surface analysis was determined to be a feasible method for

establishing a relationship between metallurgically bonded area and mechanical

strength.

d) A standardized technique was devised for evaluation of SEM micrographs and

fracture surface morphologies were defined.

e) Linear weld density and area weld density measurements did not correlate bond

area to mechanical strength.

f) A higher area fraction of metallurgical bonding related to a higher ultimate

transverse tensile strength.

2) Surface modification of VHP UAM “cu on top” builds was performed. The findings

indicated:

g) Peak force values obtained by push-pin testing of both as-welded and surface

modified builds was highly variable. While the processing parameters for the

different conditions were the same, they have not been optimized.

h) Surface modification was found to create a more energetic failure which was

verified by the area under the curves in the load vs. displacement graph.

i) Fractal analysis performed on the interfaces of surface modified and as-welded

builds indicated a more tortuous bonding path for the surface modified weld.

j) Fracture surface analysis determined similar amounts of ductile failure for the as-

welded and surface modified builds of 7.5% and 8.8% respectively.

k) The as-welded fracture surface exhibited more machined and less flow regions

than the surface modified build. Increased mechanical bonding in the surface

modified build was cited as the reason for a more energetic failure.

3) Heat treatment of VHP UAM builds improved maximum push pin force for all

aluminum and copper build geometries.

l) Heat treated Al/Cu VHP UAM builds displayed more area under the load versus

displacement curve than as-welded builds.

m) After push-pin testing, the fractured samples in the heat treated condition

underwent significant deformation compared to the as-welded builds.

n) Fractal analysis indicated that the Cu9Al4/CuAl interface had the smoothest

interface in the heat treated sample. This is the interface that failed during testing.

o) X-ray diffraction analysis found that three intermetallic compounds were formed

during a 10 minute heat treatment at 350°C: CuAl2, CuAl and Cu9Al4.

CHAPTER 8: FUTURE WORK

1. Scanning electron microscopy of fracture surfaces analyzed with the point count

method proved to be a straight forward method for determining the area fraction

of various fracture morphologies. Future work is required to verify the statistical

significance of the relationship between metallurgical bond area fraction and

mechanical strength

2. The copper interface was modified prior to welding. A direct correlation between

surface modification and mechanical strength was not determined, however, it

was found that surface modification was responsible for a more energetic failure

and had a direct influence on the interface morphology. Two techniques are

proposed in the wake of this study that may affect bonding:

a. An in-depth investigation into various modification techniques as well as

amount of deformation/roughness to determine optimal surface roughness.

b. A surface preparation step directly before joining would remove the oxide

layer and any contaminates on the surface. While oxides form quickly,

this technique has been implemented with success in roll bonding

applications and may improve metallurgical bonding significantly.

3. A heat treatment at 350°C for 10 minutes was found to improve maximum push

pin force needed for failure in all geometries of aluminum and copper dissimilar

builds. The follow needs to be addressed:

a. Further studies investigating the relationship between mechanical strength

and post weld heat treatment must be analyzed for statistical significance.

b. Transmission electron micrography (TEM) analysis of the as-welded and

heat treated interfaces is needed. This would allow for the verification of

the characterization of intermetallic layers and EDS could determine the

amount of diffusion across the interface.

c. The application of heat treatment to different material combinations is of

interest to expand the application of UAM.

4. X-ray tomography was performed on a pure copper VHP UAM build.

Representative images from this analysis are shown in Appendix A. In the future,

this technique could be used to determine the volume fraction of voids as a

method for correlating void space to bond quality.

REFERENCES

1. Obielodan, J.O., et al., Multi-material bonding in ultrasonic consolidation. Rapid

Prototyping Journal, 2010. 16(3): p. 180-188.

2. Ram, G.D.J., D.H. Johnson, and B.E. Stucker, Ultrasonic Consolidation with

aluminum and copper. Virtual and Rapid Manufacturing, 2008: p. 603-610.

3. Ram, G.D.J., et al., Use of ultrasonic consolidation for fabrication of multi-material

structures. Rapid Prototyping Journal, 2007. 13(4): p. 226-235.

4. Kong, C.Y. and R.C. Soar, Fabrication of metal-matrix composites and adaptive

composites using ultrasonic consolidation process. Materials Science and

Engineering a-Structural Materials Properties Microstructure and Processing, 2005.

412(1-2): p. 12-18.

5. Hopkins, C.D., M.J. Dapino, and S.A. Fernandez, Statistical Characterization of

Ultrasonic Additive Manufacturing Ti/Al Composites. Journal of Engineering

Materials and Technology-Transactions of the Asme, 2010. 132(4).

6. Graff, K., Short, M., Norfolk, M. . A very high-power ultrasonic additive

manufacturing (VHP UAM) for advanced materials. in International Conference on

Additive Manufacturing. 2010.

7. Sriraman, M.R., S.S. Babu, and M. Short, Bonding characteristics during very high

power ultrasonic additive manufacturing of copper. Scripta Materialia, 2010. 62(8):

p. 560-563.

8. Kong, C.Y., R.C. Soar, and P.M. Dickens, A model for weld strength in

ultrasonically consolidated components. Proceedings of the Institution of

Mechanical Engineers Part C-Journal of Mechanical Engineering Science, 2005.

219(1): p. 83-91.

9. Hopkins, C.D., et al., Optimizing Ultrasonic Additive Manufactured Al 3003

Properties With Statistical Modeling. Journal of Engineering Materials and

Technology-Transactions of the Asme, 2012. 134(1).

10. Tylecote, R.F., SOLID PHASE WELDING OF METALS. Other Information: Orig.

Receipt Date: 31-DEC-69. 1968. Medium: X; Size: Pages: 343.

11. Quadir, M.Z., et al., Influence of processing parameters on the bond toughness of

roll-bonded aluminium strip. Scripta Materialia, 2008. 58(11): p. 959-962.

12. Li, L., K. Nagai, and F.X. Yin, Progress in cold roll bonding of metals. Science and

Technology of Advanced Materials, 2008. 9(2).

13. Shirzadi, A.A. and E.R. Wallach, New method to diffusion bond superalloys.

Science and Technology of Welding and Joining, 2004. 9(1): p. 37-40.

14. Mahendran, G., V. Balasubramanian, and T. Senthilvelan, Developing diffusion

bonding windows for joining AZ31B magnesium and copper alloys. International

Journal of Advanced Manufacturing Technology, 2009. 42(7-8): p. 689-695.

15. Zuruzi, A.S., H. Li, and G. Dong, Effects of surface roughness on the diffusion

bonding of Al alloy 6061 in air. Materials Science and Engineering a-Structural

Materials Properties Microstructure and Processing, 1999. 270(2): p. 244-248.

16. Society, T.A.W., Welding Handbook. 8th ed. Vol. 2. 1991, Miami.

17. Calvo, F.A., et al., Special features of the formation of the diffusion bonded joints

between copper and aluminium. Journal of Materials Science, 1988. 23(6): p. 2273-

2280.

18. Derby, B. and E.R. Wallach, Theoretical-Model for Diffusion Bonding. Metal

Science, 1982. 16(1): p. 49-56.

19. Derby, B. and E.R. Wallach, Joining methods in space: A theoretical model for

diffusion bonding. Acta Astronautica, 1980. 7(4–5): p. 685-698.

20. Somekawa, H., et al., Low temperature diffusion bonding in a superplastic AZ31

magnesium alloy. Scripta Materialia, 2003. 48(9): p. 1249-1254.

21. Li, L. and C. Zhang, Fundamentals of Ultrasonic Welding. Welding Fundamentals

and Processes. Vol. 6A. 2011, Materials Park, OH: ASM Handbook.

22. Chang, U.I. and J. Frisch, Optimization of Some Parameters in Ultrasonic Metal

Welding. Welding Journal, 1974. 53(1): p. S24-S35.

23. de Vries, E., Mechanics and mechanisms of ultrasonic metal welding. 2004, Ohio

State University.

24. Gulenc, B., Investigation of interface properties and weldability of aluminum and

copper plates by explosive welding method. Materials & Design, 2008. 29(1): p.

275-278.

25. Bahrani, A.S., T.J. Black, and Crosslan.B, Mechanics of Wave Formation in

Explosive Welding. Proceedings of the Royal Society of London Series a-

Mathematical and Physical Sciences, 1967. 296(1445): p. 123-&.

26. Cowan, G., O. Bergmann, and A. Holtzman, Mechanism of bond zone wave

formation in explosion-clad metals. Metallurgical and Materials Transactions B,

1971. 2(11): p. 3145-3155.

27. Reid, S.R., A discussion of the mechanism of interface wave generation in

explosive welding. International Journal of Mechanical Sciences, 1974. 16(6): p.

399-413.

28. Grignon, F., et al., Explosive welding of aluminum to aluminum: analysis,

computations and experiments. International Journal of Impact Engineering, 2004.

30(10): p. 1333-1351.

29. Babu, S.S., Solid State Welding Course. The Ohio State University, 2010.

30. Johnson, K., Interlaminar Subgrain Refinement in Ultrasonic Consolidation. 2008,

Loughborough University. p. 175.

31. Kautz, D.D., Fundamentals of Friction Welding. ASM Handbook. Vol. 6A. 2011,

Materials Park, OH: ASM Handbook.

32. Abbasi, M., A.K. Taheri, and M.T. Salehi, Growth rate of intermetallic compounds

in Al/Cu bimetal produced by cold roll welding process. Journal of Alloys and

Compounds, 2001. 319(1-2): p. 233-241.

33. Ouyang, J.H., E. Yarrapareddy, and R. Kovacevic, Microstructural evolution in the

friction stir welded 6061 aluminum alloy (T6-temper condition) to copper. Journal

of Materials Processing Technology, 2006. 172(1): p. 110-122.

34. Acarer, M., B. Gülenç, and F. Findik, The influence of some factors on steel/steel

bonding quality on there characteristics of explosive welding joints. Journal of

Materials Science, 2004. 39(21): p. 6457-6466.

35. Veerkamp, W.E., Copper-to-aluminum transitions in high direct-current bus

systems. Industry Applications Society 42nd Annual Petroleum and Chemical

Industry Conference - Record of Conference Papers, 1995: p. 187-195.

36. H. Hokari, T.S., K. Kawauchi and A. Muto, Magnetic impulse welding of

aluminium tube and copper tube with various core materials. Welding International,

1998. 12(8): p. 619-626.

37. Elrefaey, A., M. Takahashi, and K. Ikeuchi, Microstructure of Aluminum/Copper

Lap Joint by Friction Stir Welding and Its Performance. Journal of High

Temperature Society, 2004. 30(5): p. 286-292.

38. Abdollah-Zadeh, A., T. Saeid, and B. Sazgari, Microstructural and mechanical

properties of friction stir welded aluminum/copper lap joints. Journal of Alloys and

Compounds, 2008. 460(1-2): p. 535-538.

39. Peng, X.K., G. Heness, and W.Y. Yeung, Effect of rolling temperature on interface

and bond strength development of roll bonded copper/aluminium metal laminates.

Journal of Materials Science, 1999. 34(2): p. 277-281.

40. Heness, G., R. Wuhrer, and W.Y. Yeung, Interfacial strength development of roll-

bonded aluminium/copper metal laminates. Materials Science and Engineering a-

Structural Materials Properties Microstructure and Processing, 2008. 483-84: p.

740-742.

41. Guo, Y.J., et al., Intermetallic phase formation in diffusion-bonded Cu/Al

laminates. Journal of Materials Science, 2011. 46(8): p. 2467-2473.

42. Pretorius, R., T.K. Marais, and C.C. Theron, Thin film compound phase formation

sequence: An effective heat of formation model. Materials Science Reports, 1993.

10(1–2): p. 1-83.

43. Johnson, K., et al., New discoveries in ultrasonic consolidation nano-structures

using emerging analysis techniques. Proceedings of the Institution of Mechanical

Engineers Part L-Journal of Materials-Design and Applications, 2011. 225(L4): p.

277-287.

44. White, D.R., Ultrasonic consolidation of aluminum tooling. Advanced Materials &

Processes, 2003. 161(1): p. 64-65.

45. Kong, C.Y., R.C. Soar, and P.M. Dickens, Optimum process parameters for

ultrasonic consolidation of 3003 aluminium. Journal of Materials Processing

Technology, 2004. 146(2): p. 181-187.

46. G.D. Janaki Ram, Y.Y., J. George, C. Robinson and B.E. Stucker, Improving

Linear Weld Density in Ultrasonically Consolidated Parts. 2006: p. 692-708.

100

47. Li, D.Z. and R. Soar, Influence of sonotrode texture on the performance of an

ultrasonic consolidation machine and the interfacial bond strength. Journal of

Materials Processing Technology, 2009. 209(4): p. 1627-1634.

48. Kulakov, M. and H.J. Rack, Control of 3003-H18 Aluminum Ultrasonic

Consolidation. Journal of Engineering Materials and Technology-Transactions of

the Asme, 2009. 131(2).

49. Kulakov, M. and H.J. Rack, Surface damage during ultrasonic consolidation of

3003-H18 aluminum. Rapid Prototyping Journal, 2010. 16(1): p. 12-19.

50. Goldstein, J.I., et al., Scanning Electron Microscopy and X-Ray Microanalysis. 3rd

ed. 2003, New York: Plenum Publishers.

51. Schick, D.E., et al., Microstructural Characterization of Bonding Interfaces in

Aluminum 3003 Blocks Fabricated by Ultrasonic Additive Manufacturing. Welding

Journal, 2010. 89(5): p. 105s-115s.

52. Friel, R.J., et al., The effect of interface topography for Ultrasonic Consolidation of

aluminium. Materials Science and Engineering a-Structural Materials Properties

Microstructure and Processing, 2010. 527(16-17): p. 4474-4483.

53. Zhang, C.B., L.J. Li, and A. Deceuster, A Method for Bond Strength Evaluation for

Laminated Structures with Application to Ultrasonic Consolidation. Journal of

Materials Engineering and Performance, 2009. 18(8): p. 1124-1132.

54. Hahnlen, R.M., Development and Characterization of NiTi Joining Methods and

Metal Matrix Composite Transducers with Embedded NiTi by Ultrasonic

Consolidation. 2009, Ohio State University.

101

55. Stucker, B.E., G.D.J. Ram, and Y. Yang, Effect of process parameters on bond

formation during ultrasonic consolidation of aluminum alloy 3003. Journal of

Manufacturing Systems, 2006. 25(3): p. 221-238.

56. Joshi, K.C., Formation of Ultrasonic Bonds between Metals - Is Based on

Dislocation Processes Which Are Manifested as a Unique Softening Phenomenon.

Welding Journal, 1971. 50(12): p. 840-&.

57. Abbasi, M. and M.R. Toroghinejad, Effects of processing parameters on the bond

strength of Cu/Cu roll-bonded strips. Journal of Materials Processing Technology,

2010. 210(3): p. 560-563.

58. Schick, D.E., Characterization of Aluminum 3003 Ultrasonic Additive

Manufacturing, in Material Science Engineering. 2009, The Ohio State University:

Columbus.

59. Drury, W. and A.M. Gokhale, Measurement and Interpretation of Fracture Surface

Fractal Dimension. Metallography: Past, Present and Future, 1993. 75: p. 295-310.

60. Lee, Y.H., et al., The Fractal Dimension as a Measure of the Roughness of Rock

Discontinuity Profiles. International Journal of Rock Mechanics and Mining

Sciences & Geomechanics Abstracts, 1990. 27(6): p. 453-464.

61. Stiny, G., Fractals - Form, Chance, and Dimension - Mandelbrot,Bb. Environment

and Planning B-Planning & Design, 1977. 4(2): p. 248-250.

62. Lopes, R. and N. Betrouni, Fractal and multifractal analysis: A review. Medical

Image Analysis, 2009. 13(4): p. 634-649.

102

63. Roberts, A.J. and A. Cronin, Unbiased estimation of multi-fractal dimensions of

finite data sets. Physica A: Statistical Mechanics and its Applications, 1996. 233(3–

4): p. 867-878.

64. Roach, D.E. and A.D. Fowler, Dimensionality Analysis of Patterns - Fractal

Measurements. Computers & Geosciences, 1993. 19(6): p. 849-869.

65. Truog, A.G. and S.S. Babu, Heat Treatment of Very High Power Ultrasonic

Additive Manufactured (VHP UAM) Aluminum and Copper Structures.

Unpublished Research, 2012.

66. Matthews, S.J., et al., Dissimilar Metals, in Welding Handbook: Metals and their

Weldability, J.R. Condra, Editor. 1982, American Welding Society: Miami, FL. p.

514-587.

67. Liu, H.J., et al., Microstructural characterisation and mechanical properties of

friction stir welded joints of aluminium alloy to copper. Science and Technology of

Welding and Joining, 2011. 16(1): p. 92-99.

68. Ueoka, T. and J. Tsujino, Welding characteristics of aluminum and copper plate

specimens welded by a 19 kHz complex vibration ultrasonic seam welding system.

Japanese Journal of Applied Physics Part 1-Regular Papers Short Notes & Review

Papers, 2002. 41(5B): p. 3237-3242.

69. Hentzell, H.T.G., R.D. Thompson, and K.N. Tu, Interdiffusion in Copper-

Aluminum Thin-Film Bilayers .1. Structure and Kinetics of Sequential Compound

Formation. Journal of Applied Physics, 1983. 54(12): p. 6923-6928.

103

70. Braunovic, M. and N. Aleksandrov, Intermetallic Compounds at Aluminum-to-

Copper and Copper-to-Tin Electrical Interfaces. Electrical Contacts - 1992 :

Proceedings of the Thirty-Eighth Ieee Holm Conference on Electrical Contacts,

1992: p. 25-34.

71. Babu, S.S. and R.R. Dehoff, Characterization of interfacial microstructures in 3003

aluminum alloy blocks fabricated by ultrasonic additive manufacturing. Acta

Materialia, 2010. 58(13): p. 4305-4315.

72. Yang, Y., G.D.J. Ram, and B.E. Stucker, Bond formation and fiber embedment

during ultrasonic consolidation. Journal of Materials Processing Technology, 2009.

209(10): p. 4915-4924.

73. Taban, E., J.E. Gould, and J.C. Lippold, Dissimilar friction welding of 6061-T6

aluminum and AISI 1018 steel: Properties and microstructural characterization.

Materials & Design, 2010. 31(5): p. 2305-2311.

74. Ogawa, T., K. Tokaji, and K. Ohya, The Effect of Microstructure and Fracture

Surface-Roughness on Fatigue-Crack Propagation in a Ti-6al-4v Alloy. Fatigue &

Fracture of Engineering Materials & Structures, 1993. 16(9): p. 973-982.

75. Rodelas, J., G. Hilmas, and R.S. Mishra, Sinterbonding cobalt-cemented tungsten

carbide to tungsten heavy alloys. International Journal of Refractory Metals & Hard

Materials, 2009. 27(5): p. 835-841.

76. Date, M., et al., Ductile-to-brittle transition in Sn-Zn solder joints measured by

impact test. Scripta Materialia, 2004. 51(7): p. 641-645.

104

APPENDIX A: X-RAY TOMOGRAPHY OF VHP UAM BUILDS

A.1 INTRODUCTION

Interfacial characterization methods such as linear weld density and visual observation with a scanning electron microscope are limited because they only gather information from a cross sectional plane within an overall UAM build. Therefore, void density measurements are limited to area calculations which must be averaged and estimated for the entire build. X-ray tomography was investigated in this study, as this technique can produce 2-D and 3-D recreations of metallic parts [77]. This technique is non-destructive and can show microstructural features such as inclusions, cracks and voids [78].

X-ray tomography has been used for medical applications for over forty years with a typical resolution of 300μm. Continual development of the process has allowed for resolution to improve [78]. Medium resolution micro-tomography has a resolution limit of 8μm, where some of the most advanced high resolution micro-tomography machines have <1μm resolution. The voids in a typical UAM build are on the order of microns, thus high resolution microtomography is needed.

X-ray tomography is the process of combining information from many radiographs, each taken with a different orientation of a sample in front of the detector

[78]. It is then possible to compute the attenuation coefficient at each point of the sample 105 from the set of radiographs. This information is used to create a reconstruction of the scanned part in 3-D.

A.2 EXPERIMENTAL

Both the copper and aluminum side of a VHP UAM dissimilar metal weld were analyzed with X-ray tomography. The processing parameter for all foils was the same: oscillation amplitude (34 μm), welding speed (35.5 mm/s) and normal force (5.5 kN).

These parameters were modified from a previous VHP UAM study and have not been optimized for bond strength [7]. A diamond saw blade was used to cut a cube from each sample, 1.5 mm on each side.

A Skyscan 1172D X-Ray μCT instrument was used to obtain two dimensional cross section tomography scans of both aluminum and copper samples. The instrument uses a lab source with a Tungsten target operating in the voltage region from 20-100kV, power of 10W, X-Ray focal spot size less than 5um and CCD Detector 11 megapixel resolution. The highest resolution capability of the machine is reported at .8um [79]. An operating voltage of 30kV and an Aluminum filter were used.

A.3 RESULTS

Representative pictures taken with X-ray tomography from a VHP UAM copper build are shown below.

106

500 μm

Figure A.1. Radiographic image of Cu sample 1

500 μm

Figure A.2. Radiographic image of Cu sample 2

107

500 μm

Figure A.3. Radiographic image of Cu sample 3

500 μm

Figure A.4. Radiographic image of Cu sample 4

108

500 μm

Figure A.5. Radiographic image of Cu sample 5

500 μm

Figure A.6. Radiographic image of Cu sample 6

109

500 μm

Figure A.7. Radiographic image of Cu sample 7

500 μm

Figure A.8. Radiographic image of Cu sample 8

110

500 μm

Figure A.9. Radiographic image of Cu sample 9

500 μm

Figure A.10. Radiographic image of Cu sample 10

111

500 μm

Figure A.11. Radiographic image of Cu sample 11

500 μm

Figure A.12. Radiographic image of Cu sample 12

112

A.4 DISCUSSION AND CONCLUSIONS

A movie was created of the images as they travel through the sample. The copper build had significant void formation throughout the entire sample. Figure A.3 shows the cross section with few voids and is representative of the best interface bonding condition that was observed throughout the sample. Figure A.5 is representative of a particularly dense void region in the sample. No conclusions of bond quality or process capability can be made from this analysis as parameters were not optimized and void region has not been correlated to mechanical or interfacial analysis.

This research is presented here to demonstrate the use of X-ray micro-tomography on UAM builds. It is possible that this technique could be used to calculate void volume within a sample. Other works within this thesis have shown that cross-sectional interfaces that appear consolidated can be misleading. It is quite possible that this technique would simply be an extension of linear weld density, which does not accurately describe bond quality. Further investigation of this fascinating technique is required.

A.5 REFERENCES

1. Olurin, O.B., et al., The investigation of morphometric parameters of aluminium foams using micro-computed tomography. Materials Science and Engineering a-

Structural Materials Properties Microstructure and Processing, 2002. 328(1-2): p. 334-

343.

2. Maire, E., et al., On the application of X-ray microtomography in the field of materials science. Advanced Engineering Materials, 2001. 3(8): p. 539-546.

113

3. Sriraman, M.R., S.S. Babu, and M. Short, Bonding characteristics during very high power ultrasonic additive manufacturing of copper. Scripta Materialia, 2010. 62(8): p. 560-563.

4. Bruker. SkyScan 1172 high-resolution micro-CT. 2012; Available from: http://www.skyscan.be/products/1172.htm.

114

APPENDIX B: LOAD VS. DISPLACEMENT CURVES

115

a) b)

Figure B.1 Aluminum and copper baseplate load vs. displacement curve and fractured builds.

116

a) b)

Figure B.2. “All aluminum” and “all copper” load vs. displacement curves and fractured builds.

117

a) b)

Figure B.3 “Copper on top” and “aluminum on top” load vs. displacement curves and fractured builds. SEM image of sheared button shown in c).

118

a) b)

Figure B.4 Heat treated “all aluminum” and “all copper” load vs. displacement curves and fractured builds.

119

a) b)

Figure B.5 14μm sonotrode and 700lbf Load vs. displacement curves and fractured builds in a “cu on top” configuration.

120