Initial Weldability of High Entropy Alloys for High Temperature

Applications

Thesis

Presented in Partial Fulfillment of the Requirements for the Degree Master of Science in

the Graduate School of The Ohio State University

Alexander Charles Martin

Graduate Program in Welding Engineering

The Ohio State University

2019

Thesis Committee

Carolin Fink, Advisor

Antonio Ramirez

Copyrighted by

Alexander Charles Martin

2019

Abstract

High Entropy Alloys (HEA) are a new class of alloys that was introduced in the early 2000’s. These alloys are composed of five or more elements in near equiatomic ratios, with no single base element. HEAs have gained a lot of attention due to their unique or superior properties as compared to conventional alloys. However, there has been little attention paid to the welding metallurgy and weldability of HEAs. However, welding for manufacturing and repair is a key issue for structural engineering applications. This work aimed to establish an initial understanding of the welding metallurgy of HEAs, and identify any potential weldability issues with regard to weld cracking susceptibility in fusion welds. The outcomes of this initial evaluation were used to develop a methodology for rapidly screening the large compositional space of HEAs in order to find promising alloy compositions for weld applications, and ultimately to implement weldability in the early stages of HEA development.

The most commonly studied equiatomic AlCoCrCuFeNi HEA was determined to have very poor weldability, due to the positive mixing enthalpy of copper and a high hardness microstructure promoted by aluminum. An improved weldability was achieved by modifying the composition to Al0.5CoCrCu0.1FeNi. Pulsed laser welding was shown to eliminate HAZ liquation cracking for AlCoCrFeNiTi HEA and reduces softening in the

HAZ of Al0.5CoCrCu0.1FeNi HEA. Refractory HEA AlMo0.5NbTa0.5TiZr showed a very ii high susceptibility to porosity and brittle fracture, but a unique fusion zone microstructure with no cracking. A high-throughput screening based on thermodynamic modeling and experimental testing was developed in order to identify HEA compositions with promising weldability, and quickly reject alloy compositions with detrimental properties towards weldability.

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Dedication

I dedicate this work to my parents who have taught me to have a strong work ethic and achieve more than what is required of me, but most of all their loving support. Without that I would not be where I’m at today. Thank you Jill and Ron.

Acknowledgments

I would first like to thank my advisor Dr. Carolin Fink, for being the best advisor

I could have asked for. Without her expertise, guidance, strong work ethic, attention to detail, and willingness to go the extra mile for me; my research goals wouldn’t have been achievable. This has made the last 2 years such an enjoyable and greatly beneficial learning experience.

I would also like to thank the professors John Lippold, Antonio Ramirez, Wei

Zhang, and David Phillips for all of their guidance and advice throughout my time in the welding engineering program.

I would like to thank Ed Pfeifer and all the graduate and undergraduate students for their help and support in the labs.

I would like to thank my sponsor the American Welding Society for awarding me a fellowship and the Institute for material research who helped fund this research.

Lastly and most importantly I want to thank my girlfriend Daniella Morris for your love and always supporting me. I’m grateful to have you and none of this work would have been possible without you by my side.

Vita

May 1993…………………………………………………………...Born, Cincinnati, Ohio

May 2012…..……………………………….Olentangy Liberty High School Powell, Ohio

May 2017……………...…………….….B.S Welding Engineering, Ohio State University

May 2019……………………...……….M.S Welding Engineering, Ohio State University

Publications

Martin, A.C. and Fink, C. (2019): “Initial Weldability Study on Al0.5CrCoCu0.1FeNi High Entropy Alloy”, Welding in the World, Jan. 2019. https://doi.org/10.1007/s40194-019-00702-7

Fields of Study

Major Field: Welding Engineering

Table of Contents

Abstract ...... ii Dedication ...... iv Acknowledgments ...... v Vita ...... vi List of Tables ...... x List of Figures ...... xi Chapter 1. Introduction ...... 1 Chapter 2. Background ...... 4 2.1 High Entropy Alloys ...... 4 2.1.1 Proposed Core Effects of High-Entropy Alloys ...... 5 2.1.2 High-Entropy Alloy Families ...... 7 2.2 Properties, Microstructures, and Element Effects of the Most Common High- Entropy Alloys ...... 9 2.3 CALPHAD-based Thermodynamic Modeling of High-Entropy Alloys ...... 14 2.3.1 ThermoCalc database and reliability for HEA’s ...... 15 2.3.2 High-Throughput CALPHAD modeling ...... 16 2.4 Welding of High-Entropy Alloys...... 16 2.5 Factors Influencing Weldability ...... 20 Chapter 3. Objectives and Approach ...... 23 Chapter 4. Experimental Procedure ...... 27 4.1 Materials and Sample Preparation ...... 27 4.1.1 Heat Treatments ...... 28 4.1.2 Hot Rolling...... 28 4.2 Welding Procedures ...... 28 4.3 Cast Pin Tear Testing (CPTT) ...... 31 vii

4.4 Microstructure Characterization ...... 31 4.5 Physical and Mechanical Properties Characterization ...... 32 4.6 CALPHAD Modeling ...... 32 4.6.1 Scheil High-throughput Screening ...... 33 Chapter 5. Results and Discussion ...... 35

5.1 HEA System AlxCoCuyCrFeNi...... 35 5.1.1 As-melted Microstructures, Phases, and Micro-hardness ...... 35 5.1.2 Equilibrium and Non-Equilibrium Thermodynamic Calculations ...... 39 5.1.3 GTA Spot Weld ...... 42 5.1.4 Cast Pin Tear Test (CPTT) Results ...... 47

5.2 Comprehensive Weldability Evaluation of Al0.5CoCrCu0.1FeNi ...... 49 5.2.1 As-melted, Heat-treated, and Hot-rolled and Heat-treated Microstructures .... 50 5.2.2 Equilibrium and Non-Equilibrium Thermodynamic Calculations ...... 54 5.2.3 Gas Tungsten Arc Stationary and Linear Welds on As-Melted and Heat- Treated Alloy Condition ...... 56 5.2.4 Gas Tungsten Arc Butt Weld on Hot-rolled and Heat-Treated Alloy Condition ...... 58 5.2.5 Pulsed Laser Welds on Heat-Treated Alloy Condition ...... 59 5.2.6 Micro-Hardness across Gas Tungsten Arc and Laser Welds...... 62 5.2.7 Cast Pin Tear Test (CPTT) Results ...... 63 5.3 HEA System AlCoCrFeNiTi ...... 66 5.3.1 As-melted and Heat treated Microstructures ...... 66 5.3.2 Equilibrium and Non-Equilibrium Thermodynamic Calculations ...... 67 5.3.3 Gas Tungsten Arc Stationary and Linear Welds on As-Melted and Heat- Treated Alloy Condition ...... 69 5.3.4 Pulsed Laser Welds on Heat-Treated Alloy Condition ...... 74 5.3.5 Post Weld Heat-treatment of the Pulsed Laser Weld ...... 75 5.4 HEA System AlMoNbTaTiZr...... 77 5.4.1 As-melted Microstructures and Non-Equilibrium Thermodynamic Calculations ...... 77 5.4.2 Gas Tungsten Arc Stationary Welds on As-Melted Condition ...... 79 5.5 Composition Optimization for Solidification Cracking Resistance using High- Throughput Screening Calculations and Experiments ...... 81 5.5.1 Methodology for Improving the Weldability of HEAs...... 81 viii

5.5.2 High-Throughput Scheil Calculations, Solidification Temperature Range and Phase Maps ...... 84 5.5.3 Gas Tungsten Arc Stationary Welding, Solidification Microstructures and Micro-hardness on Bulk As-Melted Microstructures ...... 88 5.5.4 Cast Pin Tear Testing (CPTT) ...... 94 Chapter 6. Conclusions ...... 97

6.1 HEA System AlxCoCuyCrFeNi...... 97

6.2 Comprehensive Weldability Evaluation of Al0.5CoCrCu0.1FeNi ...... 98 6.3 HEA System AlCoCrFeNiTi ...... 100

6.4 HEA System AlMo0.5NbTa0.5TiZr ...... 101 6.5 Compositional Optimization for Solidification Cracking Resistance using High- Throughput Screening Calculations and Experiments ...... 101 Chapter 7. Recommendations and Future Work ...... 103 References ...... 105

List of Tables

Table 1: HEA compositions synthesized ...... 27 Table 2: Parameters for pulsed laser welding...... 30 Table 3: Shows a summary of all the performed welds for all the HEAs compositions in specific conditions...... 30 Table 4: Solidification structure and Vickers hardness of AlxCoCrCuyFeNi alloys...... 39 Table 5: Solidification start and finish data from Scheil calculations...... 41 Table 6: XEDS point analysis of dendrite matrix and interdendritic regions of as-melted Al0.5CoCrCu0.1FeNi alloy. Given as average of four individual measurements for each region...... 53 Table 7: Geometry of pulsed laser welds, measured on transverse cross-section...... 61 Table 8: The five chosen alloys composition and STR for experimental evaluation ...... 88 Table 9: Solidification morphology and micro-hardness in the fusion zone...... 94 Table 10: Results of cast pin tear testing (CPTT) at 1” mold length, shown as circumferential cracking %...... 95

List of Figures

Figure 1: HEA family classification with a few of common in their family ...... 8 Figure 2: Effect of Al amount on phases and hardness in AlxCoCrFeNi alloy [15] ...... 10 Figure 3: Hot hardness of AlxCoCrFeNi with Al from x=0-1.8 molar ratios [16] ...... 11 Figure 4: Yield strength vs. temperature for selected refractory HEAs compared to some nickel based superalloys [10]...... 13 Figure 5: Microstructure and phases of AlMo0.5NbTa0.5TiZr alloy [20]...... 14 Figure 6: EBSD of electron beam weld on Cantor HEA CoCrFeMnNi [26] ...... 17 Figure 7: Important factors that influence the weldability with respect to an alloy’s composition and material properties...... 21 Figure 8: Developed approach to explore the weldability of high entropy alloys...... 26 Figure 9: Clamping fixture used for GTA butt welds and the resulting weld...... 29 Figure 10: As-melted microstructure of AlCoCrCuFeNi alloy composition: a) optical micrograph in as-polished condition, b) SEM/BSE image of dendrite region, c) and d) SEM image and compositional profile (XEDS) across interdendritic region (arrow)...... 36 Figure 11: As-melted microstructures as a function of Cu and Al additions: a) AlCoCrCuFeNi, b) AlCoCrCu0.5FeNi, c) AlCoCrCu0.3FeNi, d) AlCoCrCu0.1FeNi, e) AlCoCrFeNi, and f) Al0.5CoCrCu0.1FeNi. Light optical micrographs etched with molybdenum solution...... 38 Figure 12: Equilibrium step-calculation (left) and equilibrium and Scheil solidification calculations (right) for equiatomic AlCoCrCuFeNi alloy composition...... 40 Figure 13: Scheil solidification calculations for AlCoCrCuyFeNi alloys with different Cu ratio...... 41 Figure 14: Calculated phase fractions at the end of solidification for AlxCoCrCu0.1FeNi alloys with different Al ratio (Scheil calculations)...... 41 Figure 15: Light optical micrographs from cross-sectioned autogenous GTA spot weld on equiatomic alloy composition AlCoCrCuFeNi: a) Overview of the fusion line and heat- affected zone (HAZ) region, and b) detail of HAZ cracking along Cu-rich interdendritic phase (yellow color). Molybdenum etchant...... 43 Figure 16: Image analysis of Cu-rich interdendritic phase (red colored) in autogenous gas tungsten arc (GTA) spot weld on equiatomic AlCoCrCuFeNi alloy composition...... 44 Figure 17: Autogenous gas tungsten arc (GTA) spot welds with cracks in the fusion zone and heat-affected zone (HAZ): a) AlCoCrCu0.5FeNi, b) AlCoCrCu0.3FeNi, c) AlCoCrCu0.1FeNi and d) AlCoCrFeNi...... 46

Figure 18: High magnification micrographs (LOM) of fusion zone (a, b) and heat- affected zone (c, d) cracking in the autogenous gas tungsten arc (GTA) spot welds: a) AlCoCrCu0.5FeNi, b) AlCoCrCu0.3FeNi, c) AlCoCrCu0.1FeNi and d) AlCoCrFeNi...... 46 Figure 19: Fracture surface of transgranular cracks in AlCoCrFeNi alloy (SEM micrograph) ...... 47 Figure 20: Cast pin tear test (CPTT) results of AlxCoCrCuyFeNi alloys with different molar ratio of Al (x = 0.5 and 1.0) and Cu (0 ≤ y ≤ 1.0)...... 49 Figure 21: Light optical micrographs of Al0.5CoCrCu0.1FeNi microstructures in different conditions: a) as-melted “button-shaped” ingot, b) cast pin sample (CPTT), c) heat- treated plate, and d) hot-rolled and heat-treated plate...... 51 Figure 22: Left: Bulk micro-hardness in Vickers of Al0.5CoCrCu0.1FeNi alloy in as- melted, heat-treated, and hot-rolled and heat-treated condition. Right: XRD patterns in as- melted and heat-treated condition...... 52 Figure 23: Backscattered electron (BSE) image and XEDS line scan across the interdendritic region of as-melted Al0.5CoCrCu0.1FeNi alloy...... 54 Figure 24: Left: Calculated equilibrium phases vs. temperature (property diagram) of nominal alloy composition Al0.5CoCrCu0.1FeNi. Right: Solidification calculation of phase formation sequence and solid phase fractions from equilibrium (dashed line) and Scheil simulation (solid line)...... 55 Figure 25: a) Cross-section of autogenous GTA stationary weld on as-melted “button- shaped” ingot with b) detail at the fusion line; and c) cross-section of autogenous GTA linear weld on 3 mm plate in heat-treated condition with d) detail at the fusion line. All optical micrographs...... 58 Figure 26: a) Cross-section of autogenous GTA square butt joint of 3mm plates in hot- rolled and heat-treated condition with b) and c) details at the fusion line. All optical micrographs...... 59 Figure 27: Cross-sections of laser welds on 3 mm plate in heat-treated condition: a) Parameter set #1, b) parameter set #2, and c) parameter set #3. All optical micrographs. Note the difference in magnification...... 61 Figure 28: a) Cross-section of laser weld using parameter set #1 on 3 mm plate in heat- treated condition, and b) detail of solidification cracking and porosity in the fusion zone. All optical micrographs...... 61 Figure 29: Micro-hardness traverses on GTA and laser welds performed on the different alloy conditions...... 63 Figure 30: Results of cast pin tear testing (CPTT) of Al0.5CrCoCu0.1FeNi alloy shown as circumferential cracking (in %) measured as a function of cast pin length (in inch). Error bars show maximum and minimum values...... 65 Figure 31: a) SEM-BSE image at fracture surface of cross-sectioned cast pin, and b) SEM image of actual fracture surface showing dendritic nature of cracking and evidence of interdendritic film formation...... 65 Figure 32: Light optical micrographs of Al10Ti6Cr8Fe15Co25Ni36 at% microstructures in different conditions: a) as-melted “button-shaped” ingot, b) over-aged plate [1200°C for 20 hr] ...... 67

xii

Figure 33: (Left) Calculated equilibrium phases vs. temperature (property diagram) of Al10Ti6Cr8Fe15Co25Ni36 at% alloy. (Right) Solidification calculation of phase formation sequence and solid phase fractions under equilibrium (dashed line) and Scheil condition (solid line)...... 68 Figure 34: a) Top surface of an autogenous GTA stationary weld on 3 mm thick plate of Al10Ti6Cr8Fe15Co25Ni36 at% alloy in as-melted condition, b) cross-section of the same weld, and c) detail at the fusion line...... 71 Figure 35: GTA linear weld on 3 mm plate of Al10Ti6Cr8Fe15Co25Ni36 at% alloy in over- aged condition [1200°C for 20 hr]: a) Top section overview, b) photograph of weld sample fractured for fracture analysis, c) cross-sections overview of the weld, and d) detailed view at the fusion line...... 73 Figure 36: GTA linear weld on 3 mm plate of Al10Ti6Cr8Fe15Co25Ni36 at% alloy in over- aged condition [1200°C for 20 hr]: a) cross-section of the fusion line region at the bottom of the weld with a HAZ liquation cracking, and b) detailed view of the HAZ with a liquation crack...... 74 Figure 37: Pulsed laser welds on 0.125” (3mm) plates of heat-treated in overaged condition [1200°C for 20 hr]: a) Top section overview, b) detailed view at the fusion line, c) cross-sections overview of the weld, and d) detailed view at the fusion line...... 75 Figure 38: Photomicrograph of the post weld solution heat treatment (1220°C for 30 min) on the pulsed laser welded Al10Ti6Cr8Fe15Co25Ni36 at% alloy...... 76 Figure 39: SEM-BSE micrographs of the AlMo0.5NbTa0.5TiZr in the as-melted “button- shaped” ingot condition...... 77 Figure 40: (Left) Solidification calculation of phase formation sequence and solid phase fractions from equilibrium (dashed line) and Scheil simulation (solid line) for AlMo0.5NbTa0.5TiZr alloy. (Right) Solidification calculation showing liquid phase composition vs. solid fraction...... 78 Figure 41: Cross-section of an autogenous GTA stationary weld on as-melted “button- shaped” ingot of AlMo0.5NbTa0.5TiZr alloy: a) LOM with the fusion zone at the top, and b) SEM-BSE micrograph of the fusion zone and HAZ...... 80 Figure 42: Schematic of proposed high-throughput weldability screening using computational and experimental tools...... 83 Figure 43: Heat maps of the Scheil STR’s at (85, 95, 99, and 100) % fractions solid (Fs) for Ni vs Al and overlaid are the phases present during solidification. The green circle is the STR of the Al0.5CoCrCu0.1FeNi HEA...... 85 Figure 44: Ni vs. Al phase map and acceptable STR range. The 5 yellow stars depict the alloys chosen for further weldability evaluation...... 87 Figure 45: Cross-sections of autogenous GTA stationary weld on as-melted “button- shaped” ingots for each alloys #1 to #5: a) Al7Co21.7Cr21.7Cu1Fe21.7Ni27 (#1), b) Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2), c) Al15Co21Cr21Cu1Fe21Ni21 (#3), d) Al15Co20Cr20Cu1Fe20Ni24 (#4), and e) Al11Co24Cr24Cu1Fe24Ni16 (#5). All light optical micrographs. Molybdenum etchant...... 90 Figure 46: Detailed cross-section of autogenous GTA stationary weld on as-melted “button-shaped” ingot at the fusion line. a) Al7Co21.7Cr21.7Cu1Fe21.7Ni27 (#1), b) Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2), c) Al15Co21Cr21Cu1Fe21Ni21 (#3), d) xiii

Al15Co20Cr20Cu1Fe20Ni24 (#4), and e) Al11Co24Cr24Cu1Fe24Ni16 (#5). All light optical micrographs. Molybdenum etchant...... 93 Figure 47: SEM image of the Al14Co23.3Cr23.3Cu1Fe23.3Ni15 fracture surface from a pin with 100% cracking...... 96

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Chapter 1. Introduction

High-entropy alloys (HEA) are a new class of alloys that was introduced in the early 2000’s. These alloys have gained a lot of attention due to the unique properties that some alloy compositions exhibit, for example excellent specific strength, exceptional mechanical performance at high-temperature, and outstanding ductility and fracture toughness at cryogenic temperatures. High-entropy alloys are usually composed of five or more principal alloying elements, with concentrations in the range of 5 to 35 at% for each element and no single principle element. This results in a high mixing entropy, which is believed to stabilize random solid solutions in these alloys, and to prevent the formation of intermetallic phases [1]. In recent years, as more alloy compositions are being explored, high-entropy alloys have been shown to be more complex, and contain several phases and intermetallic compounds instead of a single solid solution phase.

Until very recently the majority of studies performed on high-entropy alloys have been of an exploratory nature aiming to provide a fundamental understanding of the relationship between composition, microstructure, and properties in these alloys [2].

Now, research is shifting towards a more application-based HEA design, where alloying elements are carefully selected based on the desired alloy properties, and multiple phases and tailored microstructures are deliberately introduced. This approach promises advanced high-entropy alloys with industrially relevant properties that offer new

solutions to challenges encountered in conventional alloys for high temperature applications, additive manufacturing, or cladding layers for corrosion and wear resistance. Welding is one of the most important manufacturing processes. In a large number of structural applications, the successful development and utilization of materials depend on their ability to be welded and joined. Therefore, welding and weldability will become a key issue in high-entropy alloys for structural engineering applications.

However, there has been little attention paid to the welding metallurgy and weldability of high-entropy alloys. In particular, fusion zone and heat-affected microstructures controlled by non-equilibrium microstructural evolution have not yet been studied in detail, but will have a significant impact on weld properties in these materials.

This work aimed to (a) establish an initial understanding of the welding metallurgy of high-entropy alloys, and (b) develop a methodology to implement considerations on the weldability of these alloys into the early HEA development process. These goals are an initial step towards the successful development and utilization of high-entropy alloys for structural engineering applications by enhancing weld properties and avoiding issues related to joining or weld repair of these alloys. This study investigates the welding metallurgy and weldability of three HEA systems that were proposed in the literature for use in high-temperature structural applications:

AlCoCrCuFeNi, AlCoCrFeNiTi and AlMoNbTaTiZr. Each alloy system was investigated to identify potential weldability issues with regards to weld cracking susceptibility, segregation behavior, and phase formation. Welding trials were performed on various material conditions: (a) as-melted ingots, (b) homogenized and heat-treated plates, and

pulsed laser beam welding (PLBW) processes. The resulting microstructures of the fusion and heat-affected zones were evaluated using a variety of characterization techniques

(LOM, SEM, EDX, XRD, DTA, and microhardness) and computational CALPHAD- based (CALculation of PHAse Diagrams) thermodynamic modeling. Weldability analysis addressed solidification cracking resistance using gas tungsten arc spot welding on small buttons, and cast pin tear testing (CPTT). Based on the results of this initial evaluation of the welding metallurgy and weldability of the selected three HEA systems, a methodology was developed that enables the screening of a large number of HEA compositions for their resistance to solidification cracking by high-throughput

CALPHAD-based solidification calculations and experimental testing techniques. The goal is to rapidly identify a compositional space for high-entropy alloys with development potential, and quickly reject alloy compositions with some critical deficiency with regard to weldability.

Chapter 2. Background

2.1 High Entropy Alloys

Most conventional alloys are based on one ‘base’ element. For example, steel is based on Fe, and aluminum alloys are based on Al. In order to improve the properties, different kind of alloying elements are added to the principal element, forming an alloy family. This micro-alloying approach has been used for new alloy development ever since antiquity, when bronze (90% Cu, 10% Sn) was the first alloy that impacted humanity. However, the number of elements in the periodic table is limited, thus the alloy families that can be developed are limited and so are the properties that can be achieved.

In sharp contrast, a new concept of alloy design was proposed in the early 2000’s, named high-entropy alloys (HEAs). These alloys are designed not from one or two ‘base’ elements, but from multiple elements at equiatomic or near equiatomic ratios. This multi- principal-element concept explores a new portion of the alloy design space, opening up a tremendous number of possible alloy compositions [3], [4].

High-entropy alloys are typically composed of five or more principal elements in near equiatomic ratios, whose concentrations fall in the range of 5 to 35 at%. These alloys are called “high-entropy alloys” because they have significantly higher mixing entropies than conventional alloys. Because of their unique multi-principal-element composition, HEAs can exhibit useful properties, some of them not seen in conventional

alloys. These include high strength and hardness, exceptional high-temperature strength, outstanding wear resistance, and good corrosion and oxidation resistance [1].

2.1.1 Proposed Core Effects of High-Entropy Alloys

In the first decade of HEA exploration, researchers have proposed four core effects to explain HEAs behavior: (1) the high entropy effect, (2) the lattice distortion effect, (3) the sluggish diffusion effect, and (4) the cocktail effect [5]. These hypotheses were made with the earliest publications [6], [7] at a time when very few compositions had been explored. By now research has shown that these effects are not applicable to most HEAs or not true at all. However, some may view these effects as theory’s rather than hypotheses.

The proposed “high entropy effect” [5] has been the signature concept of HEAs.

As the name suggests, high entropy alloys are based on the Boltzmann’s hypothesis of maximizing configurational entropy by increasing the number of elements added to the alloy system. The high configurational entropy was believed to stabilize random solid solutions in these alloys against the formation of intermetallic phases, which were considered often detrimental to alloy properties [1]. In recent years, many high entropy alloys have been shown to be more complex and contain several phases and intermetallic compounds instead of a single solid solution phase. Many studies have shown that high configurational entropy is not the dominating factor to stabilize single phase solid solutions, but other factors such as atomic size and packing density play a significant role

[8]. For this reason, HEAs are also known as multi-principal element alloys (MPEAs) or

complex concentrated alloys (CCAs). In this work they will be referred to as high- entropy alloys to avoid any confusion.

The lattice distortion effect is caused by various elements making up the crystal lattice with different atomic radius which imposes a local distortion to atoms occupying the lattice sites. This is known as solid solution hardening and is a common strengthening mechanism in conventional alloys. However, it has been claimed that the lattice distortion is more severe in HEAs [5]. The severe lattice distortion has also been reported to be the reason for low intensity X-ray diffraction peaks, high strength and hardness, and reduction in electrical and thermal conduction seen in these alloys [1]. However, this may not apply to all HEA compositions, and even though this may seem logical many claims are still lacking experimental quantification.

The sluggish diffusion effect claims that diffusion and kinetics are slower in

HEAs compared to conventional alloys. Due to the fact that different kind of atoms are occupying the lattice site, the energies change before and after an atom jumps to a vacancy. If an atom goes to a site with low energy, the atom may become trapped, and if an atom jumps to a high energy site, it is just as likely to jump back to the original site.

This has been used to explain the formation of nano-precipitates, which easily form in

HEAs, but slowly increase in size due to the low diffusion rates [9]. Another factor that may affect diffusion and kinetics is element activity level. Typically, elements that have a lower melting point are less active. This would make elements that have lower activity less likely to jump into vacancies and may retard phase transformations and nucleation.

However, most of these claims have been made on secondary observation, i.e. microstructure evolution and amorphous formation on solidification, and very few 6

diffusion experiments have been performed. There have also been studies that contradict these ideas, so that the sluggish diffusion effect is still a topic of debate [10].

Unlike the three other effects, the cocktail effect is not a hypothesis but rather the idea that good or unique material properties in HEAs can result from unexpected combinations of elements.

2.1.2 High-Entropy Alloy Families

One of the most attractive things about HEAs is the sheer endless combination of compositions, basically comprised of any elements on the periodic table. HEAs have been around for about 15 years, and researchers are only just scratching the surface of possible elemental combinations. As of now, there are hundreds or thousands of HEAs that have been studied, which is only a fraction of all possible alloy compositions. HEAs have been classified into different alloy families, with the best attempt done by Miracle and Senkov [1] in 2017. From the most commonly studied to the least studied alloy families. Presented in Figure 1 are their proposed classification which entails: 3d transition metals HEAs, refractory metals HEAs, light metal HEAs, and others, i.e. lanthanide transition metals, precious metals, and brasses with the most common alloys studied of that family.

Figure 1: HEA family classification with a few of common in their family

By far the most frequently studied HEAs are the 3d transition metals. The main elements used in this alloy family are Al, Co, Cr, Cu, Fe, Mn, and Ni with a minimum of four or five elements per alloy that can be at different atomic ratios. Other metal elements are frequently added, such as Mo, Nb, Si, Ta, Ti and V along with minor amounts of B,

C, and N. This HEA family can be divided into two different groups, the first is

CoCrFeMnNi, which is also known as the Cantor alloy. This group of alloys is known to consist of a single-phase FCC phase, and has remarkable ductility, fracture toughness, and strength at cryogenic temperatures. The second group in this HEA family is

AlCoCrCuFeNi, which typically has a BCC or BCC+FCC crystal structures. This alloy group has received wide attention due to the high temperature stability and strength retention up to 800°C that is seen in some alloy compositions [11]. As mentioned above, additional elements such as Mo, Nb, Si, Ta, and V are often added to improve properties.

Titanium has been added to the AlCoCrFeNi-Ti to form L12-Ni3 (Ti, Al) coherent nano- precipitates as a strengthening phase [12], which makes these alloys potential candidates for high temperature applications.

The refractory HEA family was proposed in 2010 by Senkov et al. [13] with the goal of developing alloys for high temperature structural applications. These alloys consist of Cr, Hf, Mo, Nb, Ta, Ti, V, W, and Zr. Alloys within this family are often based on MoNbTaW, MoNbTaTiZr, HfNbTaZr, CrMoNbTa or CrNbVZr element groupings, but have many derivatives with different atomic ratios and other added refractory elements. These alloys can also contain non-refractory elements such as Al or Si to decrease alloy density and improve properties [1].

Light metal or low density HEAs have densities around 2.67-5.59 g/cm3 and are made up of elements such as Al, Ca, Li, Mg, Sc, Si, Ti, and V. A challenging problem for this HEA family has been the strength to ductility tradeoff, but also the cost of these alloys is quite high. So far, research on light metal HEAs has been quite limited with few compositions that have been explored [14].

Research on other HEA families, i.e. lanthanide transition metals, precious metals, and brasses is extremely limited. These alloys are not relevant to this work and will therefore not be discussed further.

2.2 Properties, Microstructures, and Element Effects of the Most Common High-

Entropy Alloys

The equiatomic alloy AlCoCrCuFeNi is one of the most studied HEAs. This is due to its unique microstructure and mechanical properties at both room and high temperatures. Early on it was found that copper had a tendency to segregate during solidification due to its high mixing enthalpy with the other elements. For this reason, copper has been phased out of the majority of alloys currently being researched. 9

Aluminum (Al) in alloy AlCoCrFeNi has a big influence on microstructure and phase formation thus affecting the properties [11]. Figure 2 shows that as the amount of Al is increased, the hardness goes up due to Al stabilizing the BCC phase. In AlxCoCrFeNi the structure varies from single FCC phase at Al0–0.3 to mixed FCC and BCC phase at Al0.5–

0.9, and then finally to single BCC phase at Al0.9–3. Due to spinodal decomposition, the

BCC phase in these alloys is composed of disordered BCC (A2) enriched in Co, Cr, and

Fe, and ordered BCC (B2) enriched in Ni and Al [10].

Figure 2: Effect of Al amount on phases and hardness in AlxCoCrFeNi alloy [15]

The effect of annealing on the mechanical properties in AlxCoCrFeNi HEAs depends on the phase transformation involved. Different annealing temperatures can favor either the formation of the BCC or FCC phase. If BCC is favored, the alloys become harder and more brittle. When FCC is favored, the alloys soften and become

more ductile [2]. The hot hardness of AlxCoCrFeNi HEAs decreases slowly with increasing temperature below ~500oC, and decreases rapidly above 500oC. As shown in

Figure 3, Al0–0.3 alloys exhibit a much lower hardness, but these alloys do not show the

o hardness drop at temperatures above 500 C (shown in Figure 3). The Al0–0.3 alloys therefore exhibit a higher resistance to softening than alloys with different Al content.

Age hardening behavior was observed for Al0.9–1.0 alloys in the temperature range of 700–

800oC due to the formation of σ phase [16].

Figure 3: Hot hardness of AlxCoCrFeNi with Al from x=0-1.8 molar ratios [16]

Al0.5CoCrFeNi HEA has a dual phase FCC and BCC microstructure with a good balance of strength and ductility (1,220 MPa and 25% strain) respectively. The as-cast microstructure is primarily FCC with spinodal decomposition in the interdendritic regions

into BCC (A2 and B2) phases. Niu et al. [17] showed that when heat treated at 650°C for

0.5-8 hours, both the yield strength and the ultimate tensile strength increased with a slight loss in ductility. This was attributed to the precipitation of a nano-sized B2 phase in the interdendritic and dendritic regions. To achieve the same effect other elements such as titanium (Ti) or molybdenum (Mo) are added to AlxCoCrFeNi HEAs for precipitation strengthening. Manzoni et al. [18] used Ti and optimized the alloy composition

(Al10Co25Cr8Fe15Ni36Ti6, in at%) to produce a microstructure consisting of an ordered ’ phase embedded in a face-centered cubic solid-solution  matrix. At long time exposure to high temperatures above 900oC, needle-like B2 precipitates (NiAl) form, which was found to be very detrimental to mechanical properties. However, with optimal aging heat treatment the mechanical properties at room and high temperatures are superior to most

Ni-based superalloys commonly used for high temperature applications.

Refractory HEAs are also popular candidates for replacing Ni-based superalloys, some showing better properties in certain areas of application. Most of the refractory HEAs have a BCC crystal structure with a very high hardness and yield strength at both ambient and elevated temperatures (Figure 4).

Figure 4: Yield strength vs. temperature for selected refractory HEAs compared to some nickel based superalloys [10].

However, yield strength is not the only factor to consider when comparing refractory HEAs to other common high temperature alloys. Other considerations such as raw material cost, density, ductility at room temperature, and oxidation and corrosion behavior are some major drawbacks of the refractory HEAs. Senkov et al. [19] sought to reduce density and improve properties of an alloy containing MoNbTaZr with addition of lower density elements (Al and Ti). AlMo0.5NbTa0.5TiZr alloy was found to have the highest yield strength with a low density, but a major drawback was the low ductility at temperatures below 600oC. The microstructure of this alloy (Figure 5) consists of equiaxed grains of an ordered B2 matrix phase, coherent, nanometer-sized precipitates of a disordered BCC phase (~60% by volume) inside the grains, and coarse particles (~10%) of an ordered hexagonal phase at the grain boundaries [20]. The matrix B2 phase was

found to be enriched in Nb and Ta, whereas the HCP grain boundary particles were enriched in Al and Zr.

Figure 5: Microstructure and phases of AlMo0.5NbTa0.5TiZr alloy [20].

2.3 CALPHAD-based Thermodynamic Modeling of High-Entropy Alloys

Over the decades, as computational power increases and development thermodynamic databases become more reliable the use of CALPHAD (CALculation of

PHAse Diagrams) computational modeling is becoming more popular. Modeling can 14

provide information for an unknown or known alloy composition including the phases present, phase compositions, volume fractions, and transformation temperatures.

2.3.1 ThermoCalc database and reliability for HEA’s

Most thermodynamic databases are built on alloy systems that have one base element, which were developed to fit thermodynamic functions of experimental data from binary and ternary phase diagrams. High entropy alloys posse a different challenge due to the vast composition space would require quaternary and higher order systems, but obtaining experimental data would be extremally time consuming [21]. Instead, reliable

HEA databases can be obtained from the combination of extrapolating binary and ternary data [22]. Unfortunately, even the most developed thermodynamic databases lack in the complete thermodynamic description (full assessment) of all binary and, especially, ternary systems based on the elementary components present in the database. For example, the TCHEA3 database includes 26 elements that can form 325 binary and 2600 ternary systems, but only 294 binary and 136 ternary systems are assessed in the full range of composition and temperature. Therefore, complete thermodynamic description is currently unavailable for a huge number of quaternary (or higher order) alloy systems, which may have a potential practical interest [23]. Though current research suggests that the number and type of phases predicted are often reliable, but transformation temperatures, volume fractions and phase compositions are less accurate. As an example, a thermodynamic database for the AlCoCrFeNi system was developed by extrapolating binary and ternary systems to wider composition ranges, and phase diagrams predicted with this database agree well with experimental results [1]. However, calculations 15

become less accurate with more complex compositions or less studied HEAs. Current efforts are working towards developing new strategies to accelerate the discovery and validation of HEA’s enormous composition space [24].

2.3.2 High-Throughput CALPHAD modeling

High-throughput CALPHAD modeling is being used to accelerate the discovery of new HEAs and conventional alloys within a design window of desired microstructures and properties. Thousands of calculations are automated for a selected criterion to find alloy compositions that would require extensive time and money via the traditional trial and error method. The criteria often chosen for high throughput calculations for HEAs are: single-phase solid-solution alloys, number of phase alloys, type of phase, low density alloys, etc [24], [25]. High throughput calculations are also enriching knowledge of the types and numbers of phases formed. Both computational methods predict that the fraction of alloys that are single-phase solid solutions decrease as the number of elements increases, helping to resolve the high entropy hypothesis [1].

2.4 Welding of High-Entropy Alloys

Very little information has been published on the welding metallurgy and weldability of HEAs. However, there have been some studies regarding the cladding, hard facing, and additive manufacturing of different HEA compositions. The few studies that have focused on welding and joining have mainly been on the Cantor alloy, i.e.

CoCrFeMnNi. Wu et al. [26] at Oak Ridge National Labs were the first to do a weldability study in 2016 by performing electron beam welding on the Cantor alloy. An 16

autogenous butt weld was made on a 1.5 mm cold-rolled and annealed plate. No solidification or liquation cracking in the weld was reported. The Cantor alloy has a small solidification temperature range of ~60oC and low amounts of segregation, which reduces the susceptibly to solidification cracking. Figure 6 shows large elongated grains from the fusion line to the center of the fusion zone. Large grains in the fusion zone often reduce mechanical properties such as strength, ductility, and fracture toughness. However, despite the large grains, mechanical testing showed that the welded sample maintained nearly the same strength and ductility as the base metal for both cryogenic and room temperatures. These initial results show that the Cantor alloy has very promising weldability.

Figure 6: EBSD of electron beam weld on Cantor HEA CoCrFeMnNi [26]

A similar study by Jo et al. [27] from 2017 looked at laser welding and friction stir welding of the Cantor alloy. The laser welds showed similar results as the electron beam welds with no cracking and large columnar grains in the fusion zone. However, close to the fusion line the solidification microstructure was enriched in Mn, O, and S, which can reduce ductility. Interestingly, the fusion zone of the laser welds exhibited a higher microhardness as compared to the base metal, likely due to the small dendrite arm spacing and local compositional fluctuations. Friction stir welding of the Cantor alloy resulted in a large increase in hardness within the stir zone due to grain refinement. The tensile strength of both the friction stir welds and the laser welds was close to that of the base metal. The friction stir welds had a higher strength but lower ductility.

Also in 2017, Sokkalingam et al. [28] published on gas tungsten arc welding of

Al0.5CoCrFeNi HEA. This is the only other alloy for which published research on welding and joining is currently available. The same author published on laser beam welding of Al0.5CoCrFeNi HEA in 2018 [29]. For both studies, the cast alloy was forged into plates and then homogenized. Autogenous gas tungsten arc and laser welds were made. No solidification cracking or other defects were reported in the GTAW or laser welds. However, for both welding processes, a loss of hardness in the fusion zone was observed, and XRD results showed a decrease in the BCC (B2) phase. The GTA welds also showed a loss in strength and ductility as compared to the base metal.

Hard facing and cladding are common processes that utilizes local heat sources such as an arc or laser to apply a surface coating for hardness, wear resistance, and corrosion resistance. Multiple studies have shown HEAs have excellent mechanical and chemical properties such as high hardness, high mechanical strength, and high wear and 18

corrosion resistance [30], [31]. In 2008, Jie et. al. [32] used the AlCoCrNiMoFeSix HEA as a filler, which was deposited on low-carbon steel substrates using gas tungsten arc welding (GTAW) process. The hardness of cladding layers was around 885HV. The wear resistance was improved by increasing the Si content. However, throughout the paper the author did not mention any defects or cracking that may have occurred. Abed et. al. [33], hard faced a carbon steel substrate using GTAW with the Fe49Cr18Mo7B16C4Nb6 at%

HEA. The results were high-quality multi- layer deposits, free of cracking, with excellent metallurgical bonding to the substrate. The microstructure of the deposited layers consisted of a nanostructured matrix of α-Fe with an average hardness of 600-800 HV.

Multiple feasibility and research studies have been done using HEAs as cladding for better corrosion resistance. Most research has been done using the 3D transition metals

HEAs such as the AlxCoCrCuFeNi, Al2CoCrCuFeNiTix, and CoCrCuFeNi [34].

Unfortunately, most research only focuses on the corrosion aspect and doesn’t report on any cracking that may have occurred during cladding/welding. Additive manufacturing

(AM) of HEAs is another process that is very similar to welding and the use of HEAs is beginning to receive more attention to find new alloys suitable for AM. The 3D transition metals HEAs have been the most widely studied for various types of AM processes resulting in some with and without success [35]. Karlsson et. al. [36] studied the

AlCoCrFeNi HEA using powder bed selective laser melting, with the goal of establishing a process parameter window. Although, the author found it impossible to build crack- and pore-free samples with SLM. This can mainly be attributed to induced thermal stresses and to some extent segregation driven phase transformations during the building process.

However, with the addition of a pre-heating stage, thermal cracking could be minimized 19

and thereby enable SLM for AM of this alloy. Joseph et. al. [37] used direct laser fabrication (DLF) to produce bulk samples of three alloys based on the AlxCoCrFeNi (x=

0.3, 0.6 and 0.85) HEA system. The only cracking that was observed was in the first 5-8 mm of height adjacent the build plate. It was found with Al (0.6) alloy that the higher cool rates of direct laser sintering produced a Widmanstätten structure rather than a dendritic structure.

2.5 Factors Influencing Weldability

The term “Weldability” has been defined in many different ways and in different context, in short, the term describes “how easily a material can be jointed without defects”. Figure 7 shows a list of some key factors that influence weldability from a compositional and material property perspective. This list is by no means complete and demonstrates the complexity of incorporating these key factors in an early stage of HEA development. For that reason, this research will only focus on solidification cracking, liquation cracking, fusion zone and heat-affected zone softening, and any other observable defects.

Figure 7: Important factors that influence the weldability with respect to an alloy’s composition and material properties.

In the order of welding defects occurring, hot cracking such as solidification and liquation cracking will be the first defect to arise. So, it seems logical to identify the key factors in the order of their occurrence. Solidification cracking only occurs when there is tensile strain and liquid films along solidification boundaries at the end of solidification.

These liquid films are greatly influenced by the alloy’s composition. However, modifications to the composition can be used to affect the liquid present by changing the solidification temperature range (STR), fraction eutectic, number of low melting phases or solidification mode reducing the susceptibility to cracking. Over the years numerous amounts of experimental work has been done showing the relationship between the STR and cracking. Generally, alloys with lower STR’s are less prone to cracking due to the accumulation of less strain and formation of the large eutectic fractions. Although very commonly in nickel base alloys large STR’s and the amount of eutectic phases that form

at the end of solidification will cause cracking. To reduce this susceptibility engineers have found that increasing this eutectic phase even more promotes eutectic crack healing, which backfill eutectic liquid into the cracks. Another important factor is the solidification mode by solidifying with a BCC crystal structure alloys will have a high resistance to cracking.

Liquation cracking is often less common than solidification cracking and can sometimes be more easily avoided. These small cracks occur in the HAZ typically along the grain boundaries due to the remelting of a lower temperature phase. The melting of these phases can be from the metal being in the as-cast condition where the interdendritic regions melts due to segregation that occur on solidification during casting. Similarly, these same mechanisms can occur in multi-pass welding, which is referred to as weld metal liquation cracking. Another occurrence is when second phase particles undergo constitutional liquation and wet along the grain boundaries. Other properties of an alloy such as thermal conductivity, thermal expansion, grain size, and wetting behavior can further increase the likelihood of cracking.

Other detrimental effects due to welding can occur that will also limit HEAs to be effectively welded. Heat generated from welding can cause HAZ softening which reduces the strength in this region. Typically, there’re two ways that strength is lost in the HAZ; dilution of the precipitation strengthening phase or grain growth of an alloy that has been cold worked.

Chapter 3. Objectives and Approach

As outlined in the background section of this document, the HEA concept provides a much larger compositional range where alloys are no longer bound to a single base element. This shown in some alloy compositions to exhibit superior or comparable high temperature properties to common alloys used today. However, weldability has not been addressed for the HEA alloy compositions, which would limit their applications and cause manufacturing challenges in the future. Therefore, this research aims to do the following listed below.

1. Select the most promising structural engineering alloys from three different HEA

family’s for high temperature application.

2. Identify potential weldability issues with regard to weld cracking susceptibility,

segregation behavior and phase formation in fusion welds on high-entropy alloys

for high temperature applications.

3. Develop a fundamental understanding of welding metallurgy for HEAs and its

differences compared to conventional alloys.

4. Propose a methodology to evaluate the weldability of HEAs in order to implement

it as a material property in the early HEA development process.

The approach developed to explore the weldability of high entropy alloys is laid out in Figure 8. Four individual phases are tied together to gain a fundamental understanding 23

of the initial weldability of an HEA system, define a composition space with promising weldability, and quickly reject compositions with detrimental properties towards welding.

Phase (01) used a very simple initial evaluation of the basic welding behavior and sought to identify any weldability concerns. The HEAs compositions chosen all showed good properties at high temperatures but the alloys compositions were very different from one another. This was to cover a larger range of any weldability differences and similarities between the alloys. Since most of the basic physical metallurgy of high entropy alloys is rather unknown, the base metal microstructure was also characterized to better understand any welding metallurgy concerns. Alloys selected from literature, were subjected to GTA stationary “spot” welding on as-melted ingot microstructures, for a simple investigation. Equilibrium and Scheil CALPHAD modeling were done to confirm the results of experimental characterization techniques. Phase (02) characterized factors known to affect weldability in conventional alloys, such as the STR, hardness, phase formation and transformations, or element partitioning. These factors were determined and investigated as possible causes for cracking or defects that were observed in Phase

(01). Elements that attribute to these factors were identified, and compositional thresholds were established to mitigate poor weldability. In order to determine compositional thresholds GTA stationary “spot” welding on as-melted ingot microstructures of different alloy compositions was performed until no cracking or defects were observed in the welds. Phase (03) involved a more comprehensive study of the HEA compositions that were optimized in Phase (02) and showed promising weldability. Cast pin tear testing was used to find the cracking threshold. Based on this criterion the alloys were compared to conventional alloys (in this case mostly to Ni-base 24

alloys). Homogenizing, aging and phase altering heat treatments were utilized along with thermo-mechanical processing to achieve different material conditions for good mechanical properties, which might worsen the alloys weldability. Welding processes with vastly different heat input (pulsed laser and GTA welding) were used on the different material conditions to investigate the effect on weldability. Hardness mapping and microstructural characterization using SEM informed any changes due to welding in the heat-affected zone. Phase (04) entailed a refinement of the alloy composition to further improve weldability. This was done by a high throughput screening using computational and experimental tools. For this, the important factors that affect weldability in the alloy system were determined, and selection criteria for an improvement of weldability were selected. The computational screening using automated thermodynamic solidification modeling enabled the selection of few alloy compositions, whose weldability was then assessed by high-throughput experimental techniques. The results were also used to validate the thermodynamic modeling.

Figure 8: Developed approach to explore the weldability of high entropy alloys.

Chapter 4. Experimental Procedure

4.1 Materials and Sample Preparation

The selected HEA systems (AlCoCrCuFeNi, AlCoCrFeNiTi and AlMoNbTaTiZr) and their exact compositions which can be seen in Table 1. Each alloy was synthesized by arc melting of constituent elements (99.9 % pure) in a water-cooled copper hearth under argon atmosphere. The “button-shaped” ingots were remelted a total of three times to ensure complete mixing and improve chemical homogeneity. Melting time was minimized to avoid elements evaporating so the target composition was achieved. The final mass of the cast ingots was 10-20 grams depending on the final use. The 20 gram ingots were induction melted and gravity casted into a rectangular copper mold to produce thin plates with 63 x 13 x 3 mm and 63 x 6 x 6 mm dimensions.

Table 1: HEA compositions synthesized

Alloy System Alloys (molar) Alloys (atomic %) AlCoCrCuFeNi Al16.7Co16.7Cr16.7Cu16.7Fe16.7Ni16.7 AlCoCrCu0.5FeNi Al18.2Co18.2Cr18.2Cu9.1Fe18.2Ni18.2 AlCoCrCu0.3FeNi Al18.8Co18.8Cr18.8Cu5.6Fe18.8Ni18.8 AlCoCrCu0.1FeNi Al19.6Co19.6Cr19.6Cu1.9Fe19.6Ni19.6 AlCoCrFeNi Al20Co20Cr20Fe20Ni20 AlCoCrCuFeNi Al0.5CoCrCu0.1FeNi Al10.8Co21.7Cr21.7Cu2.1Fe21.7Ni21.7 Al0.4Co1.3Cr1.3Cu0.06Fe1.3Ni1.6 Al7Co21.7Cr21.7Cu1Fe21.7Ni27

Al0.8Co1.4Cr1.4Cu0.06Fe1.4Ni0.9 Al14Co23.3Cr23.3Cu1Fe23.3Ni15 Al0.9Co1.26Cr1.26Cu0.06Fe1.26Ni1.26 Al15Co21Cr21Cu1Fe21Ni21 Al0.9Co1.2Cr1.2Cu0.06Fe1.2Ni1.44 Al15Co20Cr20Cu1Fe20Ni24 Al0.6Co1.4Cr1.4Cu0.06Fe1.4Ni0.9 Al11Co24Cr24Cu1Fe24Ni16

AlCoCrFeNiTi Al0.6Co1.5Cr0.5Fe0.9Ni2.1Ti0.36 Al10Co25Cr8Fe15Ni36Ti6

AlMoNbTaTiZr AlMo0.5NbTa0.5TiZr Al20.8Mo8Nb20.8Ta8Ti20.8Zr20.8 27

4.1.1 Heat Treatments

The plates and buttons were subjected to different heat treatments for homogenization or to precipitate phases using a resistance or radiation furnace under an

o argon atmosphere. The Al0.5CoCrCu0.1FeNi alloy was heat treated at 650 C for 2.5 hours from the as-melted microstructure to precipitate the BCC (B2) as a strengthening phase

[17]. The Al0.6Co1.5Cr0.5FeNi2.6Ti0.5 alloy was homogenized at 1200°C for 20 hours then immediately air quenched, no other heat treatment was done before welding. A post weld heat treatment was conducted at 1200°C for 30 minuets on a laser weld that was homogenized prior to welding. To eliminate any thick surface oxide layers, all plates and buttons were ground down and finished using 400 grit paper prior to welding.

4.1.2 Hot Rolling

The Al0.5CoCrCu0.1FeNi alloy as-cast plates with 63 x 6 x 6 mm dimensions were subjected to hot rolling to refine the grain size of the microstructure. The plates was preheated to 800oC prior to each roll. The final dimensions after rolling were 127 x 13 x 3 mm, corresponding to a 50% reduction in thickness. The hot-rolled plates received a subsequent heat treatment at 650oC for 2.5 hours to precipitate the BCC(B2) as a strengthening phase. To eliminate any surface defects and/or thick oxide layers, all plates were ground down and finished using 400 grit paper prior to welding.

4.2 Welding Procedures

Three different welding technique were performed throughout this research; (1)

Gas Tungsten Arc (GTA) spot welding (stationary welding torch), (2) GTA linear 28

welding, and (3) pulsed laser linear welding. The following setups and welding parameters were used: (1) Autogenous GTA spot welding was performed on “button- shaped” ingots (10 g) in as-melted condition using a stationary welding torch at 120 A and 11 V for 3 s at 5mm standoff distance. The cooling rate was measured by plunging thermocouples (Type C) into the solidifying weld pool. The measured cooling rate was about 200 °C/s (∆T = 1300-1100 °C). (2) Autogenous linear GTA welding was performed in the center of a 20 g plate (63 x 13 x 3 mm). Welding was done at 95 A current, 11 V voltage and 4.3 mm/s travel speed. Autogenous GTA butt welding was performed to join two 10 g plates (63 x 13 x 3 mm). The closed square-groove butt joint was done with one weld pass from each side welded at 100 A current, 10.5 V voltage and

3.0 mm/s travel speed. The plates were tightly clamped down as seen in Figure 9 to avoid distortion and increase restraint during welding.

Figure 9: Clamping fixture used for GTA butt welds and the resulting weld. 29

(3) Laser welding using a Trumpf Nd:YAG pulsed laser system was performed as an autogenous linear weld on a 20 g plate (63 x 13 x 3 mm). Three different parameter sets were used for various weld geometries. These are listed in Table 2.

Table 2: Parameters for pulsed laser welding.

Spot Spot Average Parameter Power Pulse Frequency Speed Diameter Overlap Power Set # (kW) (ms) (Hz) (mm/sec) (mm) (%) (W) 1 0.5 1 10 10 1.66 90 100 2 2.0 5 16 2 0.78 85 160 3 0.5 1 1 15 2.50 70 15

Table 3: Shows a summary of all the performed welds for all the HEAs compositions in specific conditions.

Alloy Material Condition Welding Process AlCoCrCuFeNi 10g button and As-cast GTA Spot Weld

AlCoCrCu0.5FeNi AlCoCrCu FeNi 0.3 10g button and As-cast GTA Spot Weld AlCoCrCu0.1FeNi AlCoCrFeNi 10g button and As-cast GTA Spot Weld GTA Linear Weld Al CoCrCu FeNi 20g plate and HT 0.5 0.1 Pulsed Laser Linear Weld 20g plate, Hot Rolled and HT GTA Linear Weld 10g button and As-cast GTA Spot Weld Al Co Cr FeNi Ti GTA Linear Weld 0.6 1.5 0.5 2.6 0.5 20g plate and HT Pulsed Laser Linear Weld AlMoNbTaTiZr 10g button and As-cast GTA Spot Weld

4.3 Cast Pin Tear Testing (CPTT)

The cast pin tear test (CPTT) was used to determine susceptibility to hot cracking.

The cast pin tear test is a self-restrained weldability test that utilizes an induction levitation coil to melt and feed the material of interest into pin-shaped copper molds of varying length. Hot cracking in the solidifying cast pin is induced by strain build up during solidification as a function of the mold length. Details on the test procedure and standard testing parameters that were used can be found elsewhere [38].

Alloy compositions were initially tested at a mold length of 0.5 inches with a total of three pins casted. Casting was done at a temperature of 1450°C and a chamber pressure of 3 psi (2.07 kPa). The amount of circumferential cracking on the surface of the cast pins was measured using a stereomicroscope. For some alloy compositions no cracking was observed at the 0.5 inch mold length, the mold lengths was then increased in increments of 0.125 inches till the cracking was detected.

4.4 Microstructure Characterization

The as-melted or heat-treated ingots and plates, welded samples and cast pins were sectioned and prepared for optical and electron microscopy by standard metallographic procedures. Polished surfaces were etched in molybdenum solution, containing 100 ml of each H2O, HCl and HNO3, and 3 g of molybdic acid (H2MoO4)

[39]. Samples were submerged in the solution for 3-5 s, then immediately rinsed with ethanol. Secondary electron (SE) and backscattered electron (BSE) imaging and compositional analysis were performed in a FEI Apreo FEG scanning electron microscope (SEM) equipped with X-ray energy dispersive spectroscopy (XEDS). Image 31

analysis was done for phase quantification using the Materials Image Processing and

Automated Reconstruction (MIPAR) software package [40].

4.5 Physical and Mechanical Properties Characterization

Microhardness mapping and single indentation was done to find the microhardness for all alloys and material conditions in accordance to ASTM standards.

Manual microhardness testing was done on a LECO M-400-H1used to find the bulk microhardness, samples with a hardness of around 500-600 Vickers used a 500 g load for

12 sec, and samples with a hardness of around 200-400 Vickers used a 200 g load for 12 sec. Microhardness mapper LECO LM-100AT was used for a hardness traverse on

Al0.5CoCrCu0.1FeNi alloy. Using a 200 g load for 12 sec and a spacing of 150 µm was done on all welded samples starting at the fusion line measuring towards weld and base metal. X-ray diffraction (XRD) was used to identify the crystal structure of phases present Al0.5CoCrCu0.1FeNi for the as-melted and heat-treated condition. The patterns were collected using Rigaku SmartLab with the diffracted beam graphite monochromator and a linear X’Celerator detector (Cu Kα radiation). The measurements were done at room temperature, with steps of 0.020 degree and a duration time of 0.5 s.

4.6 CALPHAD Modeling

Thermodynamic CALPHAD-based calculations were performed in Thermo-Calc version 2018a using TCHEA3 database in order to predict phase formation during heat- treatment (equilibrium step calculation), and during equilibrium and non-equilibrium

(Scheil) solidification, associated solidification path and solidification temperature range. 32

Scheil calculations hold the following assumptions: 1) local equilibrium at the planar solid-liquid interface, 2) no diffusion of substitutional elements in the solid phase, and 3) complete mixing in the liquid phase. Scheil solidification therefore represents the worst case in terms of micro segregation of elements to the grain and sub grain boundaries during solidification. In general, the more experimental solidification deviates from the equilibrium state, as typical in weld situations, the higher the possibility for additional phase formation at the final stages of solidification. All elements were included in the calculations. Scheil calculations were performed by decreasing the temperature stepwise from 2,500 °C in 1 °C intervals until the fraction of solid phase (푓푠) in the system reached

0.99. All available phases were allowed to form during equilibrium and non-equilibrium calculations, with the exception of HCP_ZN phase. Including this phase resulted in the formation of HCP-phase that was not observed experimentally.

4.6.1 Scheil High-throughput Screening

High-throughput Scheil calculations were performed using Python programming language, coupled with Thermo-Calc version 2018a and the most current database

TCHEA3. The Scheil calculations performed had the same assumptions as listed above in section 4.6. Python was used to setup the alloy composition ranges, set calculation conditions, run parallel computing, data handling of results, and data visualization. The compositions that were calculated varied, Al: 5-18 at% and Ni: 15-28 at% with increment size 0.5 at%, which yielded 651 unique alloys. Cu was set to 1 at% for each composition and the remaining compositional space was divided between the elements Co, Cr, and Fe.

All phases were accepted for the entire composition range except the HCP_ZN and 33

Al3Ni5 phases. The solidification temperature range was found at four different solid fractions (100, 98, 95, 85) percent. The predicted solidification phases for all compositions were considered at 100% fraction solid.

Chapter 5. Results and Discussion

5.1 HEA System AlxCoCuyCrFeNi

The equiatomic alloy composition AlCoCrCuFeNi was the starting point in this work for the welding metallurgy and weldability evaluation of this HEA system. This alloy was selected for the fact of it being the most commonly studied HEA in the literature for high temperature applications. A total of six alloy compositions in this system were investigated in order to study elemental effects by varying molar ratio of aluminum and copper. Metallurgical and weldability test results for the different alloy compositions are presented and discussed in the following subsections.

5.1.1 As-melted Microstructures, Phases, and Micro-hardness

Figure 10 (a) reveals the as-melted microstructure of equiatomic alloy composition AlCoCrCuFeNi. The light optical micrograph shows the dendritic structure in as-polished condition (no etching). Evidence of solute segregation is clearly distinguishable within the solidification structure. Chemical analysis (XEDS), shown in

Figure 10 (c)-(d), reveals enrichment of Cu in the interdendritic regions, indicating strong partitioning of this element to the liquid phase during solidification. All other constituent elements are depleted as compared to the dendrite composition. Strong partitioning of Cu has been reported for this alloy system due to its high positive mixing enthalpy with Fe,

Cr, Co and Ni [41].

Figure 10 (b) presents a representative backscatter electron micrograph of the dendrite microstructure. The dendritic regions solidify as body centered cubic (BCC) phase [41]. The complex modulated microstructure seen in Figure 10 (b) has been attributed to the occurrence of spinodal decomposition into ordered Al-Ni rich plates (B2) and disordered Cr-Fe rich BCC (A2) interplates [15], [41], [42]. The bright elongated constituents in Figure 10 (b) are several types of nano-sized Cu rich precipitates that have been previously reported for this alloy composition and were not further analyzed in this study [42]. The interdendritic regions, shown in Figure 10(c), were reported to be primarily composed of face-centered cubic (FCC) Cu rich phase, and some BCC plate- interplate structure formed by spinodal decomposition [42].

a) b)

25 µm

c) d) 70 Al K 60 Cr K Dendrite Fe K Interdendritic 50 Region

% Co K - 40 Ni K Cu K

30 Interdendritic Dendrite Element / at./ Element 20

0 0 1 2 3 4 5 6 Distance / µm Figure 10: As-melted microstructure of AlCoCrCuFeNi alloy composition: a) optical micrograph in as-polished condition, b) SEM/BSE image of dendrite region, c) and d) SEM image and compositional profile (XEDS) across interdendritic region (arrow). 36

Figure 11 shows the as-melted microstructures (optical micrographs) of

AlxCoCrCuyFeNi alloys with Al (x = 0.5 and 1.0) and Cu (0 ≤ y ≤ 1.0) in different molar ratio. Equiatomic alloy AlCoCrCuFeNi, shown in Figure 11(a), forms a cellular BCC dendritic structure with a high-volume fraction of Cu-rich FCC phase in the interdendritic regions. Decreasing Cu ratio in this alloy system reduces the amount of Cu-rich interdendritic phase, as shown in Figure 11(a)-(d). This is consistent with disappearing

FCC reflection peaks in X-ray diffraction on low copper compositions (y ≤ 0.5) [11]. It can also be seen that the solidification structure coarsens with decreasing molar ratio of

Cu in the alloy composition (Figure 11). The solidification mode shifts from a cellular dendritic to an equiaxed dendritic grain structure, as shown for AlCoCrCu0.1FeNi and

AlCoCrFeNi alloys in Figure 11(d) and (e), respectively.

Lowering the molar ratio of Al (x = 0.5) in the system results in a fine cellular dendritic microstructure with distinct interdendritic regions, compare AlxCoCrCu0.1FeNi alloys in Figure 11 (d) and (f). Al0.5CoCrCu0.1FeNi has a hardness of 208 HV, which is significantly lower than that of the other alloys in this study (> 500 HV, see Table 4).

This is due to the fact that the dendrite core regions in the low Al compositions (x ≤ 0.5) have an FCC structure, whereas BCC dendrites with an ordered structure form in all other alloy compositions [15]. It has been previously reported that aluminum facilitates the formation of BCC phase in this alloy system and enhances the ordering of the BCC structure [11] & [15]. Table 4 summarizes the observed solidification morphology and shows the measured hardness for all alloy compositions.

Figure 11: As-melted microstructures as a function of Cu and Al additions: a) AlCoCrCuFeNi, b) AlCoCrCu0.5FeNi, c) AlCoCrCu0.3FeNi, d) AlCoCrCu0.1FeNi, e) AlCoCrFeNi, and f) Al0.5CoCrCu0.1FeNi. Light optical micrographs etched with molybdenum solution.

Table 4: Solidification structure and Vickers hardness of AlxCoCrCuyFeNi alloys.

Alloys Solidification Morphology Hardness AlCoCrCuFeNi Cellular dendritic 526 ± 8

AlCoCrCu0.5FeNi Cellular dendritic 574 ± 8

AlCoCrCu0.3FeNi Cellular dendritic 577 ± 8

AlCoCrCu0.1FeNi Equiaxed dendritic 504 ± 5 AlCoCrFeNi Equiaxed dendritic 527 ± 7

Al0.5CoCrCu0.1FeNi Cellular dendritic 208 ± 5

5.1.2 Equilibrium and Non-Equilibrium Thermodynamic Calculations

From the equilibrium step-calculation in Figure 12 (left), phase formation under equilibrium conditions can be determined. Initially there are two liquid phases present, the first liquid (blue line) is predicted to solidify as BCC phase enriched in Co, Cr, and

Fe. The second liquid (red line) is enriched primarily in Cu but also slightly enriched with

Al and Ni. This liquid phase is stable down to slightly below 1000°C then freezes similar to a eutectic reaction. The completely solidified alloy is predicted to consist of an ordered

B2 phase enriched in Al and Ni, a Cu-rich order FCC phase, and Sigma phase. Figure 12

(right) presents the phase formation sequence during solidification under equilibrium conditions (black line), and during Scheil solidification of equiatomic alloy composition

AlCoCrCuFeNi. Solidification starts at about 1280°C. Two liquid phases are predicted.

The Cu-rich liquid phase persists down to low temperatures of about 950°C. The calculated Scheil solidification path indicates that the primary solidifying phase is a Cr-,

Co- and Fe-rich BCC phase. An ordered B2 phase enriched in Al and Ni forms at lower temperatures, and a Cu-rich FCC phase is predicted to form at the end of solidification.

Figure 12: Equilibrium step-calculation (left) and equilibrium and Scheil solidification calculations (right) for equiatomic AlCoCrCuFeNi alloy composition.

The Cu-rich liquid that is predicted under both equilibrium conditions and Scheil conditions (Figure 12) results in a large solidification temperature range (STR) for equiatomic composition AlCoCrCuFeNi of 300-350oC. Lowering the Cu ratio in the

AlCoCrCuyFeNi alloy system changes the shape of the predicted Scheil solidification curve, as shown in Figure 13. The curve is steeper in the final stages of solidification with decreasing Cu ratio, which indicates a smaller amount of liquid phase present at the end of solidification, and a narrower solidification temperature range. Table 5 compares the solidification start and finish temperatures, and the solidification temperature range for alloy compositions with varying Cu ratio (y = 1.0, 0.5, 0.3 and 0.1). AlCoCrCu0.1FeNi is characterized by the smallest solidification temperature range of 162oC.

1450

1400

1350

1300

C) 1250

Cu0.1 1200

Cu0.3

1150 Temperature ( Temperature Cu 1100 0.5 AlCrCoCuFeNi

1050 AlCrCoCu0.5FeNi Cu1 AlCrCoCu0.3FeNi 1000 AlCrCoCu0.1FeNi 950 0.00 0.10 0.20 0.30 0.40 0.50 0.60 0.70 0.80 0.90 1.00 Mole Fraction of Solid

Figure 13: Scheil solidification calculations for AlCoCrCuyFeNi alloys with different Cu ratio.

Table 5: Solidification start and finish data from Scheil calculations.

Solidification Finish Solidification Finish Solidification Temperature Solidification Start Alloys Temperature (fs=0.99) Temperature (fs=0.95) Range (STR) (fs=0.95) Temperature (°C) (°C) (°C) (°C) AlCoCrCuFeNi 1282 947 952 330

AlCoCrCu0.5FeNi 1304 953 960 345

AlCoCrCu0.3FeNi 1347 958 989 358

AlCoCrCu0.1FeNi 1387 974 1224 162

BCC_B2 AlCrCoCu0.1FeNi BCC_A2 FCC_L12 Al0.8CrCoCu0.1FeNi

Al0.5CrCoCu0.1FeNi

Al0.3CrCoCu0.1FeNi

0.0 0.2 0.4 0.6 0.8 1.0 Phase fraction at end of solidification Figure 14: Calculated phase fractions at the end of solidification for AlxCoCrCu0.1FeNi alloys with different Al ratio (Scheil calculations).

Figure 14 shows the effect of Al content (x = 1.0, 0.8, 0.5 and 0.3) on the type of phases present after Scheil solidification of AlxCoCrCu0.1FeNi alloys. It can be seen that a higher Al ratio promotes the formation of BCC phase in the solidified microstructure.

Al0.3CoCrCu0.1FeNi exhibits an almost completely FCC structure after Scheil solidification.

5.1.3 GTA Spot Weld

Figure 15 (a) shows the fusion line and heat-affected zone (HAZ) region of an autogenous gas tungsten arc (GTA) spot weld on an as-melted button shaped ingot of equiatomic alloy composition AlCoCrCuFeNi. Weld solidification occurs by epitaxial nucleation and growth at the fusion line. The fusion zone exhibits a finer dendritic structure as compared to the as-melted ingot microstructure, caused by the much higher cooling rate of the stationary spot weld as compared to the button-melted ingot. Copper is heavily partitioning to the liquid during weld solidification. This results in Cu-rich FCC phase in the interdendritic regions of the fusion zone, as reported for the as-melted ingot microstructure. In the heat-affected zone, small cracks (≤ 500 µm) were observed adjacent to the fusion line, see Figure 15 (a). Higher magnification micrographs revealed that the crack path proceeds along the Cu-rich interdendritic regions of the as-melted ingot microstructure, see Figure 15 (b). No cracking was observed in the fusion zone of the gas tungsten arc spot weld.

Figure 15: Light optical micrographs from cross-sectioned autogenous GTA spot weld on equiatomic alloy composition AlCoCrCuFeNi: a) Overview of the fusion line and heat-affected zone (HAZ) region, and b) detail of HAZ cracking along Cu-rich interdendritic phase (yellow color). Molybdenum etchant.

Image analysis on etched light optical micrographs was performed to measure the volume fraction (assumed to be equal to measured area fraction) of Cu-rich interdendritic phase in the fusion zone, heat-affected zone and the as-melted ingot microstructure, as seen in Figure 16. It was noticed that the fraction of the segregated Cu-rich phase in the

HAZ seemed to be higher than in the base metal ingot microstructure, or the fusion zone of the GTA spot weld. There was a ~5% increase in Cu-rich interdendritic phase in the

HAZ compared to the ingot microstructure. It is hypothesized that this in due to the high temperatures reached in the HAZ during GTAW, which resulted in partial melting and enhanced Cu segregation to the interdendritic regions during re-solidification. The fusion zone showed a ~5% decrease of Cu-rich interdendritic phase relative to the ingot microstructure. The fusion zone experienced much higher cooling rates compared to the button-melted ingot which resulted in a much finer dendritic structure. This makes an accurate image analysis of the fraction of Cu-rich interdendritic phase much more

challenging and potentially less accurate when using the same magnification as was done in this study.

Figure 16: Image analysis of Cu-rich interdendritic phase (red colored) in autogenous gas tungsten arc (GTA) spot weld on equiatomic AlCoCrCuFeNi alloy composition.

Figure 17 shows light optical macrographs of gas tungsten arc spot welds on as- melted ingots as a function of Cu ratio in AlCoCrCuyFeNi alloys (y = 0.5, 0.3, 0.1 and 0).

The equiatomic alloy composition (y = 1.0) exhibited small liquation cracking in the

HAZ of the GTA spot weld (see Figure 15). As the Cu ratio was reduced to 0.5 and 0.3 molar ratio respectively, severe macro-cracking (up to 2.5 mm crack length) can be seen to occur in both the fusion zone and the HAZ, some of them extending into the button- melted ingot microstructure (Figure 17, a & b). Upon higher magnification in Figure 18

(a & b), it can be seen that the cracking is interdendritic, and Cu-rich phase is apparent on the edges on both sides of the cracks.

In Figure 17 (c & d) the Cu ratio was reduced to 0.1 and 0 molar ratio, which resulted in a significant reduction of the overall cracking in the GTA spot welds. The observed cracks occurred in the fusion zone, and the HAZ, some of them extending into

the as-melted ingot microstructure. At higher magnification in Figure 18 (c & d), it can be seen that the cracking is runs transgranular through the dendritic structure. No Cu-rich phase was observed along the crack edges. Both suggests that Cu is not contributing to the observed cracking in the low/no-Cu alloy compositions. The fracture surface (SEM image) of transgranular HAZ cracking for AlCoCrFeNi can be seen in Figure 19. The surface releveled a quasi-cleavage fracture mode rather than smooth features commonly seen for cracking associated with liquid phases. Quasi-cleavage fracture is often associated with a hard brittle microstructure. In fact, both alloy compositions have a very high micro hardness exceeding 500 HV (see Table 4). The high hardness of these compositions is a result of the BCC and B2 phases, which are stabilized by the aluminum content (Alx, x = 1.0). The thermal strain experienced during the GTA spot welding was high enough to cause brittle cracking, since the high hardness microstructure was much less able to accommodate the strain.

Figure 17: Autogenous gas tungsten arc (GTA) spot welds with cracks in the fusion zone and heat-affected zone (HAZ): a) AlCoCrCu0.5FeNi, b) AlCoCrCu0.3FeNi, c) AlCoCrCu0.1FeNi and d) AlCoCrFeNi.

Figure 19: Fracture surface of transgranular cracks in AlCoCrFeNi alloy (SEM micrograph)

5.1.4 Cast Pin Tear Test (CPTT) Results

The CPTT results are presented in Figure 20 for all alloy compositions. Tests were done at a single pin length of 0.5 inch. Cracking is shown as the average circumferential cracking observed over all tested cast pin samples for an alloy composition. The equiatomic composition AlCoCrFeNi showed 100% circumferential cracking for all three tested samples. This indicates a very high susceptibility to weld hot cracking, since 0.5 inch pin length does not constitute severe test conditions for cracking in the cast pin tear test. For comparison, conventional Ni-based alloys, which are generally considered somewhat susceptible to hot cracking, typically exhibit cracking not up to pin lengths of around 1 inch and higher. Complete (100%) circumferential cracking was also obtained for alloy compositions with decreased Cu molar ratio of y = 0.5 and

0.3. Based on the presented characterization of the solidification microstructure and the computational solidification modeling, this cracking is associated with the segregation of

Cu during solidification forming a low melting liquid phase at the end of solidification along grain boundaries and interdendritic regions. Liquid films were observed on the fracture surface of the cracking in the cast pin samples and the Cu-rich phase along the interface of the crack. Cracking in the CPTT decreased for alloy compositions with a Cu molar ratio of y = 0.1. Metallographic examination of the cracks in cross-sectioned cast pins did not show any signs of segregated Cu-rich phase along the crack path, or any smooth features present on the fracture surface. Alloy composition AlCoCrFeNi without any Cu (y = 0) showed a much lower percentage of average circumferential cracking, but exhibited a large standard deviation (see Figure 20). For both, AlCoCrFeNi and

AlCoCrCu0.1FeNi alloys (y = 0 and 0.1), metallographic examination of the cast pin samples revealed that the cracking is transgranular. This observation matched what was seen in the GTA welded samples for these alloy compositions, where cracking was attributed to the high hardness, brittle microstructure.

Based on the obtained characterization, modeling and cast pin results, two additional alloy compositions in the AlCoCrCuFeNi HEA system were prepared:

CoCrFeNi (no Al and Cu), and Al0.5CoCrCu0.1FeNi (lowering the Al ratio). Table 4 shows that the hardness in the AlCoCrCuFeNi is about 500-600 HV0.5 independent of the Cu ratio. However, the hardness decreases to around 138 HV for alloy composition

CoCrFeNi (no Al and Cu). This can be attributed to a decrease in the brittle BCC and B2 phases, and the formation of primarily FCC phase. However, it was acknowledged hardness this low would not have high enough strength for engineering applications and a 48

balance of FCC and BCC was needed. For this reason no characterization or cast pin tear testing was done for alloy composition CoCrFeNi (no Al and Cu). To increase the mechanical properties Al was increased to (Al = 0.5) molar ratio, which increased the hardness to around 208 HV. Figure 20 shows that with decreased Al ratio, alloy composition Al0.5CoCrCu0.1FeNi showed no cracking in the cast pin samples, hence exhibits a much higher weld cracking resistance than all other tested alloy compositions in this HEA system.

100% 100% 100% 97% 100%

80%

60% 47%

40%

20% Average Circumferential Cracking Circumferential Average 0% 0% Cu (1) Cu (0.5) Cu (0.3) Cu (0.1) Cu (0) Cu(0.1), Al (0.5)

Figure 20: Cast pin tear test (CPTT) results of AlxCoCrCuyFeNi alloys with different molar ratio of Al (x = 0.5 and 1.0) and Cu (0 ≤ y ≤ 1.0).

5.2 Comprehensive Weldability Evaluation of Al0.5CoCrCu0.1FeNi

The previous section presented results on the elemental effects of Al and Cu on the weldability of equiatomic alloy composition AlCoCrCuFeNi. It was shown that a reduction of the Al and Cu ratio led to a decrease in weld cracking susceptibility. Alloy composition Al0.5CoCrCu0.1FeNi showed promising CPTT results with zero cracking at 49

0.5 inch pin length, and a lower hardness microstructure. The following sections show results of a more comprehensive weldability evaluation of this alloy composition.

5.2.1 As-melted, Heat-treated, and Hot-rolled and Heat-treated Microstructures

Figure 21 shows the solidification microstructures (optical micrographs) of

Al0.5CoCrCu0.1FeNi alloy in different conditions: a) as-melted “button-shaped” ingot, b) cast pin sample (CPTT), c) heat-treated plate, and d) hot-rolled and heat-treated plate. All micrographs show a two-phase microstructure with dendritic solidification and distinct interdendritic regions due to segregation during solidification. The dendrite matrix has been reported in a similar alloy composition (Al0.5CoCrFeNi) to be composed of face centered cubic (FCC) phase, while the interdendritic regions are of body centered cubic

(BCC) crystalline structure [15]. The difference in cooling rate upon solidification between the “button-shaped” ingot microstructure and the cast pin and plate microstructures is obvious from the difference in dendrite arm spacing. Also, as shown in

Figure 21(b), less amount of interdendritic phase is observed in the faster cooling rate solidification microstructure.

Figure 21 (c) shows the solidification microstructure after heat treatment at

650 °C for 2.5 hours. From the optical micrograph no, obvious microstructural changes were observed as compared to the as-melted condition. However, the effect of the heat treatment becomes obvious from the micro-hardness results shown in Figure 22 (left).

The bulk hardness of the alloy increases from 208 HV in as-melted condition to 287 HV after heat treatment. As reported by Niu et al. [17], this increase is due to the formation of nano-sized BCC (B2) precipitates upon heat treatment in the dendrite matrix and the 50

interdendritic regions. Another significant increase in hardness to 429 HV is observed in the hot-rolled and heat-treated condition. From Figure 21 (d) it can be seen that the conducted hot rolling procedure did not cause recrystallization in the microstructure, simply a deformed dendritic-interdendritic structure can be observed. This might have caused an increase in dislocation density and in the precipitation response during the subsequent heat treatment (650 °C for 2.5 hours) and led to the increase in hardness.

a b

c d

Figure 21: Light optical micrographs of Al0.5CoCrCu0.1FeNi microstructures in different conditions: a) as-melted “button-shaped” ingot, b) cast pin sample (CPTT), c) heat-treated plate, and d) hot-rolled and heat-treated plate.

Figure 22: Left: Bulk micro-hardness in Vickers of Al0.5CoCrCu0.1FeNi alloy in as-melted, heat- treated, and hot-rolled and heat-treated condition. Right: XRD patterns in as-melted and heat- treated condition.

Figure 22 (right) shows the XRD pattern of the as-melted and heat-treated alloy.

In as-melted condition, the alloy exhibits mainly FCC crystalline structures. However, the asymmetry of the FCC peak at 43.4 degree indicates a small amount of BCC phase present. In the heat-treated alloy, a mixture of FCC and BCC crystalline structures was observed. The BCC peak can be clearly seen at 44.4 degree. No ordered BCC (B2) crystalline structures were observed in the XRD results after heat treatment. This is probably due to the overall small amount of these strengthening nano-sized precipitates.

Scanning electron microscopy (SEM) was used to characterize the elemental distribution in the FCC dendrite matrix and BCC interdendritic regions of the as-melted solidification microstructure. As shown in Table 6, the dendrite composition is close to the nominal alloy composition, whereas the interdendritic regions are enriched in Al and depleted of Co and Fe. In contrast to alloy compositions with a higher Cu content (see

Section 5.1), no enrichment in Cu in the interdendritic regions was observed. The SEM-

BSE image presented in Figure 23 (left) shows a close-up view of the interdendritic 52

region, which displays a periodic, fine-scale structure of bright and dark interconnected phases. This is consistent with the modulated plate-interplate structure of disordered BCC phase and ordered BCC (B2) phase that has been reported to form by spinodal decomposition in this alloy system [43] & [44]. Line scan analysis further revealed that the region along the edge of the dendrite phase, i.e. the darker interdendritic area in

Figure 23, is even more enriched in Al and Ni as compared to the overall interdendritic region. This is consistent with previously reported results of HEAs with a similar composition. During solidification of Al0.5CoCrCu0.1FeNi alloy, the FCC dendrites form first while Al and Ni partition to the interdendritic region, which finally solidifies as disordered BCC phase, and further decomposes by a spinodal decomposition mechanisms

[45] into the modulated plate-interplate structure shown in Figure 23. Al and Ni atoms are continuously repelled from the dendrite matrix and accumulate as an Al- and Ni-rich layer between the dendrite and interdendritic regions.

Table 6: XEDS point analysis of dendrite matrix and interdendritic regions of as-melted Al0.5CoCrCu0.1FeNi alloy. Given as average of four individual measurements for each region.

Element (at %) Al Ni Co Cr Fe Cu Nominal composition 10.87 21.74 21.74 21.74 21.74 2.17 Dendrite matrix 11.28 20.39 20.86 23.90 21.50 2.73 Interdendritic region 20.28 21.64 16.03 22.05 15.63 2.70

Al Ni Cu Cr Fe Co 30

12 Atomic percent (at%) percent Atomic 6

Interdendritic Dendritic

7 0 Distance Figure 23: Backscattered electron (BSE) image and XEDS line scan across the interdendritic region of as-melted Al0.5CoCrCu0.1FeNi alloy.

5.2.2 Equilibrium and Non-Equilibrium Thermodynamic Calculations

Figure 24 presents the calculated equilibrium property diagram and solidification calculations from equilibrium and Scheil simulation, which were developed to predict phase formation during heat treatment and (weld) solidification of Al0.5CoCrCu0.1FeNi alloy. The property diagram in Figure 24 (left) shows the mole fraction of stable phases as a function of temperature. It can be seen that at the heat treatment temperature of

650 °C the phase balance is predicted to be a mixture of Cr-, Fe- and Co-rich FCC phase and some amount of Al- and Ni-rich BCC (B2) phase, the latter being reported to be the strengthening phase in this alloy composition [17]. A Cr-rich sigma phase is also predicted, but its precipitation was not observed experimentally in the heat-treated alloy and is not reported in the available literature. The formation of sigma phase might be overestimated in the used thermodynamic database (TCHEA3).

Figure 24: Left: Calculated equilibrium phases vs. temperature (property diagram) of nominal alloy composition Al0.5CoCrCu0.1FeNi. Right: Solidification calculation of phase formation sequence and solid phase fractions from equilibrium (dashed line) and Scheil simulation (solid line).

Figure 24 (right) presents the phase formation sequence during equilibrium solidification (dashed line) and Scheil solidification simulation (solid line). It can be seen that FCC and Ni- and Al-rich BCC (B2) phase form under equilibrium conditions from the liquid phase. The Scheil simulation suggests that Ni- and Al-rich BCC (B2) phase and disordered Cr- and Fe-rich BCC phase form under non-equilibrium conditions in the interdendritic regions during the final stages of solidification. The solidification temperature range (STR) from the Scheil simulation is about 141°C, measured from the start of the first solid phase formation (FCC) at 1346°C down to the temperature at which a fraction solid of 0.95 is reached during solidification (1205 °C). The STR has been shown to be an indicator for the solidification cracking resistance of an alloy during welding. The predicted STR of Al0.5CoCrCu0.1FeNi alloy is narrow as compared to compositions with a higher Cu content (≥ Cu0.3), which have a calculated STR > 300°C

(see Section 5.1). This is due to strong partitioning of Cu to the interdendritic regions during solidification, which results in the formation of a low-melting Cu-rich phase that can lead to solidification cracking in the fusion zone and heat-affected zone liquation cracking. The liquid phase composition from the Scheil simulation on

Al0.5CoCrCu0.1FeNi alloy indicates partitioning of Cu into the final liquid during solidification. However, no separate Cu-rich phase is predicted and no experimental evidence of a Cu-rich interdendritic phase was found.

5.2.3 Gas Tungsten Arc Stationary and Linear Welds on As-Melted and Heat-

Treated Alloy Condition

Figure 25 shows a cross-section of a stationary GTA weld made on a “button- shaped” ingot in as-melted condition. No cracking was observed in the fusion zone or heat-affected zone (HAZ). The weld was repeated on another ingot under the same welding conditions with no cracking. Some small crater cracks can be observed in Figure

25 (a) due to the abrupt shut off of the arc and shrinkage in the center of the weld. Crater cracks are not uncommon in autogenous welding and most commonly observed in the end crater of linear welds. The close-up view in Figure 25 (b) shows epitaxial growth from the fusion line that resulted in columnar grains, some of which reach all the way up to the surface of the weld. The fusion zone exhibits a finer dendritic structure as compared to the as-melted ingot microstructure due to a much higher cooling rate of the liquid weld pool. It is recognized that metallic alloys are not typically welded in as- melted (or as-cast) condition. However, this setup was chosen since it simulates a worst- case scenario seen in repair welding or multi-pass processing. 56

A cross-section of the linear GTA weld performed on a 3 mm plate in heat-treated condition is shown in Figure 25 (c). As shown in Figure 22 (left), heat treatment at 650°C for 2.5 hours results in a strengthening of the alloy. Welding in heat-treated condition is significant towards the weldability of this alloy. The alloy is at a higher strength, which might increase the intrinsic restraint during welding, and might make it more susceptible to weld cracking. Furthermore, the strengthening is due to nano-sized precipitates that might get dissolved at elevated temperature exposure close to the fusion line, which might cause a locally softened region in the weld heat-affected zone (HAZ). It can be seen from the etching contrast in Figure 25 (c), that the linear GTA weld resulted in a large HAZ that is about 1.5 mm wide on the surface and reaches through the whole thickness of the plate. The close-up view in Figure 25 (d) shows a narrow, but distinct region adjacent to the fusion line that appears to exhibit much less interdendritic phase as compared to the rest of the plate HAZ microstructure further away from the weld. The high temperatures reached in this region during welding likely caused a partial dissolution of the BCC interdendritic phase into the dendrite matrix. This is also apparent from the calculated property diagram in Figure 24 (left), which indicates that only FCC and BCC (B2) phase are present at temperature right below the melting temperature under equilibrium condition. The reason that this dissolution region is not present in the stationary GTA welds is probably the much lower heat input as compared to the linear

GTA weld. No cracking was observed in the fusion zone or heat-affected zone of the linear GTA weld.

a c Fusion Zone Fusion Zone

HAZ

b Fusion Zone

d Fusion Zone

HAZ

Figure 25: a) Cross-section of autogenous GTA stationary weld on as-melted “button-shaped” ingot with b) detail at the fusion line; and c) cross-section of autogenous GTA linear weld on 3 mm plate in heat-treated condition with d) detail at the fusion line. All optical micrographs.

5.2.4 Gas Tungsten Arc Butt Weld on Hot-rolled and Heat-Treated Alloy Condition

Figure 26 shows a cross-section of the autogenous square butt GTAW joint made on 3 mm plate in the hot-rolled and heat-treated condition. The goal of achieving a full penetration weld was not met, as shown in Figure 26 (a). A gap of about 0.24 mm was observed between the two weld passes. The cross-section shows the size of the HAZ, which was more narrow (about 0.5 mm) than what was observed for the linear GTA weld. This is probably due to the two plates joined together for this weld forming much more of a heat sink than the single plate the linear weld was performed on. The close-up views in Figure 26 (b) and (c) show the partial dissolution of interdendritic phase in the high temperature heat-affected zone that was also observed in the linear GTA weld. No cracking was found in the fusion zone or HAZ of the weld. At the start of the second weld pass, a small centerline crack of about 1 mm in length formed but did not proceed 58

into the rest of the weld. This kind of solidification cracking occurs in most FCC alloys and is often prevented using a start tab, pulsing, or by reducing the travel speed. Butt welds on 304 stainless steel, which were done as a reference using the same welding parameters and plate dimensions, resulted in severe solidification cracking in the center of both weld passes.

a 1st Weld Pass c Fusion Zone HAZ

2nd Weld Pass b Fusion Zone

HAZ HAZ

Figure 26: a) Cross-section of autogenous GTA square butt joint of 3mm plates in hot-rolled and heat-treated condition with b) and c) details at the fusion line. All optical micrographs.

5.2.5 Pulsed Laser Welds on Heat-Treated Alloy Condition

Cross-sections of the linear pulsed laser welds on 3 mm plate in the heat-treated condition are shown in Figure 27. The three different parameter sets used resulted in very different weld sizes and geometries (Table 2 & Table 7). The weld shown in Figure

27 (a) was made in keyhole mode in order to achieve the deepest weld penetration.

Depending on plate thickness, this mode would be used for joint welding. Laser beam welding is a high energy density process. This results in refined solidification

microstructure in the fusion zone, and very narrow HAZ region. Using optical microscopy, no microstructural difference that would indicate the width of the HAZ was observed as compared to the plate microstructure further away from the fusion line. In particular, no partial dissolution of interdendritic phase adjacent to the fusion line as observed in the GTA welds (see Figure 25 and Figure 26) was readily apparent in the laser weld. Porosity in the fusion zone and several small solidification cracks were found in some of the cross-sections obtained from this weld. Pulsed laser welding in keyhole mode is more susceptible to solidification cracking due to the high cooling rates, which generate higher strains during welding [46]. Porosity is another common problem in pulsed laser welding in keyhole mode and might further facilitate solidification cracking in the vicinity of the pores, as shown in the cross-section in Figure 28 (a) and (b). Both cracking and porosity can be reduced or prevented by a ramp down pulse shape along with optimized laser welding parameters [47]. However, the used laser system did not allow this to the extent necessary to obtain defect free welds. The laser weld in Figure 27

(b) was done in conductive mode to eliminate the porosity by increasing the width to depth ratio but keeping a constant fusion zone area as for the keyhole mode weld. This resulted in a very stable weld process, and a much wider and shallower weld geometry.

No solidification cracking or porosity was observed in the fusion zone. Finally, in an attempt to mimic laser processing as would be done in an additive laser powder bed fusion process, the laser parameters were varied to obtain a very small weld pool size.

The single pass weld shown in Figure 27 (c) has due to its small size very high cooling rates and an extremely refined microstructure in the fusion zone. No weld defects were observed in any of the obtained cross-sections. 60

Figure 27: Cross-sections of laser welds on 3 mm plate in heat-treated condition: a) Parameter set #1, b) parameter set #2, and c) parameter set #3. All optical micrographs. Note the difference in magnification.

Table 7: Geometry of pulsed laser welds, measured on transverse cross-section. Fusion Average Width Depth Ratio Parameter Set # zone area Power(W) (mm) (mm) (width/depth) (mm2) 1 100 1.36 1.06 1.28 0.69 2 160 2.58 0.41 6.29 0.72 3 15 0.53 0.09 5.89 0.03

ba b

Solidification Cracking

Porosity

Figure 28: a) Cross-section of laser weld using parameter set #1 on 3 mm plate in heat-treated condition, and b) detail of solidification cracking and porosity in the fusion zone. All optical micrographs. 61

5.2.6 Micro-Hardness across Gas Tungsten Arc and Laser Welds

The diagram shown in Figure 29 compares the evolution of hardness from the fusion zone across the HAZ and into the base material for all performed GTA welds, and the laser weld in keyhole mode laser welds (parameter set #1, see Table 2). All measured traverses were aligned at the fusion line to facilitate a better comparison. It can be seen that the hardness in the fusion zone is at about the same level for all GTAW welds, i.e. independent from the processing condition of the alloy (as-melted, heat-treated, or hot- rolled and heat-treated). This is due to the heat source producing about the same solidification microstructure no matter the initial alloy condition. The fusion zone hardness of the keyhole laser weld can be seen to be much higher as compared to the

GTA welds. This is believed to be due to the very refined solidification microstructure produced by this high energy process. It can be seen that the hardness in the HAZ varies considerably as a function of initial alloy condition and used welding process. The stationary GTA weld on the as-melted condition (orange line) does not show any decrease in HAZ hardness as would have been expected, since fusion zone and initial alloy condition are both in a very similar as-solidified condition. The slightly finer dendritic structure in the fusion zone does not have a significant effect on hardness.

Linear GTA welding on the heat-treated alloy (grey line) resulted in a very wide HAZ, which is consistent with the cross-section in Figure 25 (c). The HAZ hardness reaches the hardness level of the initial heat-treated alloy (287 HV) at a distance of about 2.8 mm from the fusion line. The high heat input of this weld probably caused the dissolution of the strengthening phase in the dendrite matrix and interdendritic regions. On the contrary, the laser weld on the same alloy condition (yellow line) does not exhibit a HAZ hardness 62

drop due to its high energy density characteristics. Actually, the fusion zone has a lower hardness in this case, since the weld metal does not exhibit any precipitation strengthening after solidification. The GTA butt weld (blue line) exhibits the largest drop of hardness in the fusion zone and HAZ as compared to the initial hot-rolled and heat- treated alloy condition. The HAZ is much narrower as compared to the linear GTA weld, probably due to the larger heat sink of this joint setup.

400 Rolled-GTAWhot-rolled/heat-treated, GTAW butt as-melted, GTAW stationary 350 Cast-GTAW Fusionline Aged-GTAWheat-treated, GTAW linear

300 Aged-Laserheat-treated, Laser linear hardness hardness (HV)

- 250 Micro 200 Fusion Zone

150

150 450 600 900 300 750

-300 -150 -450

1200 1350 1650 2100 2400 2850 1050 1500 1800 1950 2250 2550 2700 Distance (µm) Figure 29: Micro-hardness traverses on GTA and laser welds performed on the different alloy conditions.

5.2.7 Cast Pin Tear Test (CPTT) Results

Results of the cast pin tear test are presented in Figure 30. Opposite to what was done for several different alloy compositions in section 5.1, alloy Al0.5CoCrCu0.1FeNi was tested over a range of different pin length in order to assess cracking susceptibility.

Cracking in the cast pins is shown as the average circumferential cracking observed over all tested pins as a function of the pin length. As can be seen, cracking in the cast pin

samples increases with increasing pin length due to the increase in intrinsic restraint during solidification in the copper mold. A common criteria to rank different materials in terms of their susceptibility to solidification cracking is the last pin length that can be cast without cracking, i.e. the lower cracking threshold (LCT). For Al0.5CoCrCu0.1FeNi alloy, the cracking threshold is at a pin length of 0.625 inch. Figure 31 shows that the cracking occurs along interdendritic regions that solidify last during solidification. The work on the AlCoCrCuFeNi compositions with a high Cu content (≥ Cu0.3) exhibit severe cracking

(i.e. 100% circumferential cracking) at a pin length of 0.5 inch (see section 5.1). Based on the results, Al0.5CoCrCu0.1FeNi alloy can be considered to be much more resistant to solidification cracking, which corresponds to the crack-free GTA welds. Typically weld metals with LCT values below 1 inch are considered somewhat susceptible to solidification cracking, which might explain the small solidification cracks observed in the keyhole pulsed laser weld on Al0.5CoCrCu0.1FeNi.

100

40 Circumferential Cracking Cracking % Circumferential 20

0 0.375 0.625 0.875 1.125 1.375 1.625 Pin length (in) Figure 30: Results of cast pin tear testing (CPTT) of Al0.5CrCoCu0.1FeNi alloy shown as circumferential cracking (in %) measured as a function of cast pin length (in inch). Error bars show maximum and minimum values. a b Interdendritic

Interdendritic

Figure 31: a) SEM-BSE image at fracture surface of cross-sectioned cast pin, and b) SEM image of actual fracture surface showing dendritic nature of cracking and evidence of interdendritic film formation.

5.3 HEA System AlCoCrFeNiTi

The Al10Ti6Cr8Fe15Co25Ni36 at% alloy was chosen for its popularity in literature as an HEA composition that was optimized for high temperature applications, exhibiting mechanical properties and other properties superior to Ni-based superalloys used for aircraft turbines and power generation. Similar to Ni-based superalloys this HEA is strengthening by a gamma prime phase.

5.3.1 As-melted and Heat treated Microstructures

Figure 32 shows the microstructures (optical micrographs) of

Al10Ti6Cr8Fe15Co25Ni36 at% alloy in two different conditions: a) as-melted “button- shaped” ingot, and b) over-aged [1200°C for 20 hr] plate. The as-melted micrograph shows a microstructure with dendritic solidification and distinct interdendritic regions due to segregation during solidification. From EDS analysis the interdendritic region was found to be enriched in Al and Ti, while the dendrite exhibits higher amounts of Co and

Cr. It has been reported that both the dendrites and interdendritic regions contain gamma prime precipitates, which form on cooling [18]. Figure 32 (b) shows large precipitates enriched in Ni, Al, and Ti along grain boundaries and intragranularly. The EDS results

[Ni(~40%), Al(~15%), and Ti(~8%)] suggest that these large precipitates are gamma prime, which matches results found in literature on the same alloy composition [18]. The matrix consists of an FCC phase and small gamma prime and BCC (B2) precipitates [18].

The B2 precipitates are not seen in the micrographs in Figure 32 due to their small size.

Figure 32: Light optical micrographs of Al10Ti6Cr8Fe15Co25Ni36 at% microstructures in different conditions: a) as-melted “button-shaped” ingot, b) over-aged plate [1200°C for 20 hr]

5.3.2 Equilibrium and Non-Equilibrium Thermodynamic Calculations

Figure 33 presents the calculated equilibrium property diagram and solidification calculations from equilibrium and Scheil simulation, which were developed to predict phase formation during heat treatment and (weld) solidification. The property diagram in

Figure 33 (left) shows the mole fraction of stable phases as a function of temperature from 650oC -1350oC. It can be seen that the three primary phases present are gamma, gamma prime, and BCC (B2). No in-depth work was performed in this research to validate the predicted phases at different heat-treatment temperatures. The calculation is in agreement with the gamma and gamma prime phases observed in the over-aged material condition using SEM/EDS in this study (Figure 32 b), and matches quite reasonably what can be found in current literature [48]. Not shown in Figure 33 (a), but a

Cr-rich Sigma phase was predicted at temperatures less than 650oC. This phase has not been reported in literature to be experimentally observed. The initial formation of the

BCC B2 phase at around 1220oC might be over predict in the calculations because when

homogenized at 1200-1220oC only gamma was present with no B2 in the microstructure or XRD [48].

Figure 33 (right) presents the phase formation sequence during equilibrium solidification (dashed line) and Scheil solidification simulation (solid line). For the Scheil solidification it can be seen that FCC rich in Ni, Co, and Fe was the first phase to solidify at around 1320oC. BCC (B2) phase forms upon further solidification at around 1240oC followed by a Ti-rich phase at the very end of solidification. The STR was found to be

225oC an 248oC at 95% and 100% fraction solid, respectively. As reported in section

5.3.1, the interdendritic region of the as-melted microstructure is enriched in Al and Ti, and depleted of the remaining elements. The Scheil simulation indicates partitioning of Ti out of the dendrite core and into the final liquid during solidification. Though for Al the

Scheil simulation shows the opposite with little at the end of solidification.

5.3.3 Gas Tungsten Arc Stationary and Linear Welds on As-Melted and Heat-

Treated Alloy Condition

Figure 34 shows LOM macrographs of the top surface and cross-section of a stationary GTA weld made on 3 mm thick plate of Al10Ti6Cr8Fe15Co25Ni36 at% alloy in as-melted condition. From the top surface view no cracking was observed in the fusion zone or heat-affected zone (HAZ). Epitaxial growth can be seen from the fusion line that resulted in long columnar grains, some of which reach all the way to the center or the surface of the weld. A small crater crack can be observed in Figure 34 (b) due to the abrupt shut off of the arc. Porosity in the center of the plate stems from the casting process and is not associated with the weld. Crater cracks are not uncommon in autogenous welding and most commonly observed in the end crater of linear welds. The close-up view in Figure 34 (c) shows the fusion zone region of the cross-section of the weld. The fusion zone exhibits a coarser dendritic structure as compared to the as-melted casted plate microstructure due to a much slower cooling rate of the liquid weld pool.

The as-cast microstructures in sections 5.1 and 5.2 were much coarser because they were casted into a button shape on top of a copper crucible. Whereas the

Al10Ti6Cr8Fe15Co25Ni36 at% alloy was casted into a copper mold to produce a thin plate resulted in much higher cooling rates. In the high temperature HAZ of this weld three distinct zones standout in the microstructure. Zone 1, which is adjacent to the fusion zone reaches the highest temperatures and appears to show partial melting of the interdendritic region. Zone 2 shows a narrow, but distinct region that appears to exhibit much less interdendritic phase as compared to the rest of the HAZ microstructure further away from the weld. The high temperatures reached in this region during welding likely caused 69

partial dissolution of the interdendritic phase into the dendrite matrix. Zone 3 appears to have no changes to the interdendritic phase from the micrograph in Figure 34 (c), but likely had small precipitation and growth of the gamma prime phase. It is recognized that metallic alloys are not typically welded in as-melted (or as-cast) condition. However, like with the previous alloys this setup was chosen since it simulates a worst-case scenario seen in repair welding or multi-pass processing.

Figure 34: a) Top surface of an autogenous GTA stationary weld on 3 mm thick plate of Al10Ti6Cr8Fe15Co25Ni36 at% alloy in as-melted condition, b) cross-section of the same weld, and c) detail at the fusion line.

Figure 35 and shows the surface-section of the autogenous linear GTA weld made on 3 mm plate of Al10Ti6Cr8Fe15Co25Ni36 at% alloy in the over-aged condition [1200°C for 20 hr]. In Figure 35 (a & b) cracking can be seen in both the fusion zone and HAZ on the surface of the weld. Numerous cracks occurred in the partially melted HAZ, and further away from the fusion line at the edge of the HAZ along grain boundaries, see

Figure 35 (c). The cracks seen in the fusion zone are perpendicular to the direction of welding. These cracks might have initiated in front of the weld pool in the HAZ during the initial heating of the weld. The cracks were welded over but were still present in the

HAZ, as the fusion zone solidifies the cracks propagate into the fusion zone. Analysis of the fracture surface using SEM near the fusion line shows evidence of liquid film presence in the high temperature HAZ. However, the fracture surface in Figure 35 (d), which was taken in the HAZ further away from the fusion line with no liquid film present. This region shows cracking along the grain boundaries with very large precipitates on the fracture surface. This indicates that solid-state cracking might also play a role by embrittling the grain boundaries. Further characterization of the cracking, and welding of this alloy composition in different heat treatment conditions are necessary to explore the cracking mechanisms.

Figure 35: GTA linear weld on 3 mm plate of Al10Ti6Cr8Fe15Co25Ni36 at% alloy in over-aged condition [1200°C for 20 hr]: a) Top section overview, b) photograph of weld sample fractured for fracture analysis, c) cross-sections overview of the weld, and d) detailed view at the fusion line.

Figure 36 shows the fusion zone region near the bottom (cross-section) of the same weld as the one shown in Figure 35. In the high temperature HAZ, a partially melted region can be seen adjacent to the fusion boundary, similar to what was seen in the GTA spot weld in Figure 34. This partially melted region is likely associated with constitutional liquation of the large precipitates throughout the grains and along the grain boundaries. It is unclear whether the cracks started from grain boundaries liquation from melting of the precipitates outside the partially melted region or in partially melted

region. Interestedly, the cracks were found only in the HAZ and not to continue into the fusion zone.

Figure 36: GTA linear weld on 3 mm plate of Al10Ti6Cr8Fe15Co25Ni36 at% alloy in over-aged condition [1200°C for 20 hr]: a) cross-section of the fusion line region at the bottom of the weld with a HAZ liquation cracking, and b) detailed view of the HAZ with a liquation crack.

5.3.4 Pulsed Laser Welds on Heat-Treated Alloy Condition

LOM micrographs of a pulsed laser weld on 3 mm plate of

Al10Ti6Cr8Fe15Co25Ni36 at% alloy in the over-aged condition are shown in Figure 37. To lower the high heat input of the GTAW process, pulsed laser welding was used to prevent

HAZ liquation cracking. Welding parameters (set 1) from Table 2 was used which was in the keyhole mode and reached depth of 1.2 mm with none or very little trapped porosity.

Figure 37 (a & b) show the microstructure of the laser weld top surface with no cracks seen in the HAZ or the fusion zone. Figure 37 (d & c) show the very refined solidification microstructure in the fusion zone, along with a very narrow HAZ. Optical microscopy revealed no distinct microstructural changes in the HAZ. In particular, no partial dissolution or melting of the precipitates adjacent to the fusion line as observed in the

GTAW. No solidification cracking was observed in the weld, which is quite common for both pulsed laser welding and nickel-based alloys with high gamma prime [49]. The initial results show that pulsed laser welding maybe a suitable process choice for sound welding of Al10Ti6Cr8Fe15Co25Ni36 at% alloy.

Figure 37: Pulsed laser welds on 0.125” (3mm) plates of heat-treated in overaged condition [1200°C for 20 hr]: a) Top section overview, b) detailed view at the fusion line, c) cross-sections overview of the weld, and d) detailed view at the fusion line.

5.3.5 Post Weld Heat-treatment of the Pulsed Laser Weld

A post weld heat treatment was done 1220oC for 30 minutes in order to solutionize the weld and base metal microstructures. Figure 38 show optical micrographs of the weld cross-section after the heat treatment. The top micrograph shows significant 75

grain growth in the base metal, with some grains reaching up to 2-4 mm in diameter. A heat treatment at 1220oC dissolves all the precipitates which pinned the grain boundaries and allowed for the rapid grain growth. Interestingly, the sample exhibited severe intergranular cracking in the fusion zone and the base metal around the weld. No conclusion can be drawn from this single heat-treatment experiment, but the result indicate that a strain-age cracking mechanism might be cause of this cracking, similar to what is seen in gamma prime strengthened Ni-base superalloys.

Figure 38: Photomicrograph of the post weld solution heat treatment (1220°C for 30 min) on the pulsed laser welded Al10Ti6Cr8Fe15Co25Ni36 at% alloy.

5.4 HEA System AlMoNbTaTiZr

The AlMo0.5NbTa0.5TiZr alloy was chosen from literature as an HEA composition that was optimized for extremely high temperature applications, lower density, and high compressive strength.

5.4.1 As-melted Microstructures and Non-Equilibrium Thermodynamic

Calculations

Figure 39 shows the microstructure (SEM-BSE micrographs) of

AlMo0.5NbTa0.5TiZr alloy in the as-melted “button-shaped” ingot condition. Dendritic solidification and distinct interdendritic regions can be observed due to elemental segregation during solidification. Figure 39 (right), shows a higher magnification image of the dendrite and interdendritic region. The darker, interdendritic region was found to be enriched in Al and Zr using EDS analysis in the SEM.

Figure 39: SEM-BSE micrographs of the AlMo0.5NbTa0.5TiZr in the as-melted “button-shaped” ingot condition.

Figure 40 presents solidification calculations from Scheil simulation, which were used to understand the weld solidification microstructure. Figure 40 (left) presents the phase formation sequence during equilibrium solidification (dashed line) and Scheil solidification simulation (solid line). It can be seen that on Scheil solidification BCC (B2) phase is the first phase to form at around 1890oC. Next the calculation predicts Sigma phase formation, followed by an Al-Zr-rich phase during the final stages of solidification.

Solidification is complete at around 1350°C. The STR at both 100 % and 95 % solid fraction is around 535oC due to isothermal solidification starting at 86 % solid fraction.

The calculated STR for this alloy was extremely large and may be caused by the difference in each elements melting temperatures [Al=660oC, Ta=2980oC]. Figure 40

(right) shows the mole percent of elements in the liquid phase during solidification. It can be seen that both Al and Zr strongly partition out of the dendrite core and into the final liquid during solidification. The results resemble what was found in EDS analysis with the enrichment of Al and Zr in the interdendritic regions.

Figure 40: (Left) Solidification calculation of phase formation sequence and solid phase fractions from equilibrium (dashed line) and Scheil simulation (solid line) for AlMo0.5NbTa0.5TiZr alloy. (Right) Solidification calculation showing liquid phase composition vs. solid fraction. 78

5.4.2 Gas Tungsten Arc Stationary Welds on As-Melted Condition

Figure 41 shows the cross-section of a stationary GTA weld made on an as- melted “button-shaped” ingot of AlMo0.5NbTa0.5TiZr. In Figure 41 (a) large amounts of porosity and cracking can be seen in the as-melted ingot, i.e. not related to the GTA weld.

Using the same cleaning and arc melting procedure as used for the other HEAs evaluated in this work, AlMo0.5NbTa0.5TiZr alloy is very susceptible to porosity during arc melting.

The large cracks that can be seen in Figure 41 (a) in the as-melted ingot were not further investigated in this research but may have occurred due to a brittle microstructure. The bulk hardness is very high with about 700 HV. The GTA spot weld can be seen near the top of the ingot. It is very shallow with only about 0.4 mm of penetration at the center.

No cracks or porosity were observed in the fusion zone or HAZ. Figure 41 (b) is an SEM-

BSE at higher magnification of the weld region with the fusion zone near the top and the

HAZ at the bottom. Dendrite arm spacing can be seen to be more refined throughout the fusion zone as compared to the as-melted ingot microstructure. Typically, the fusion line can be easily identified due to epitaxial nucleation off the base metal, however little or no dendrites or grains can be seen nucleating at the fusion line. This microstructure suggests that the solidification mode in the fusion zone is entirely equiaxed dendritic due to large constitutional supercooling or heterogeneous nucleation off metallic oxides. From both the top surface and cross section view no cracking was observed in the fusion zone or heat-affected zone (HAZ). No further compositional optimization of this alloy system, and no further investigation into other welding processes or weldability was done.

Figure 41: Cross-section of an autogenous GTA stationary weld on as-melted “button-shaped” ingot of AlMo0.5NbTa0.5TiZr alloy: a) LOM with the fusion zone at the top, and b) SEM-BSE micrograph of the fusion zone and HAZ.

5.5 Composition Optimization for Solidification Cracking Resistance using High-

Throughput Screening Calculations and Experiments

5.5.1 Methodology for Improving the Weldability of HEAs

This work aimed to propose a methodology for incorporating weldability into an early HEA development process. It focuses on avoiding hot cracking during welding based on the fundamentals of welding metallurgy and decades of experience on what makes alloys like aluminum, nickel, and stainless steel weldable. For example, one of the key factors influencing weld solidification cracking is the liquid present at the end of solidification. Compositional modifications can be used to affect the liquid present by changing the solidification temperature range (STR), amount of liquid phase at the end of solidification, fraction eutectic, number of low melting phases, or primary solidification mode. A compositional space could be determined, to meet certain limits of these criteria to increase the resistance to weld cracking. Within this composition space of “good” weldability (i.e. weld cracking resistance), changes in alloy composition could be made to increase other material properties. This approach would yield in hundreds to thousands of possible alloy compositions that need to be screened and evaluated - a huge time and cost commitment. However, achievements in CALPHAD-based computation and accuracy of thermodynamic and kinetic databases enable the implementation of high- throughput calculations to screen a vast number of alloy compositions at reasonable time and cost. The goal using this method is to rapidly identify a compositional space for alloys with development potential, and quickly reject alloys with some critical deficiency with regard to weldability. 81

Solidification cracking was chosen as the primary weldability concern due to the fact that it is a major issue in fusion welding, and has been previously reported in the

HEA system (AlCoCrCuFeNi) that is the main focus of this research. The factors or criteria influencing the solidification cracking and the limits for each one to mitigate cracking were based on the results obtained for the AlCoCrCuFeNi system during its weldability evaluation (section 5.1 of this work), in addition to knowledge from welding of conventional alloys, where it seems appropriate. Figure 42 shows the process for high- throughput weldability screening and evaluation of high entropy alloys with resistance to solidification cracking. Different criteria would be devised as a function of intended application, initial or base alloy and welding processes.

The first part that was setup were the conditions and variables for the

“ThermoCalc program”. The nominal composition of the alloy, two composing elements of interest and their composition range are selected. The Al0.5CoCrCu0.1FeNi alloy showed promising weldability in terms of resistance to solidification cracking in the preformed experimental work (section 5.2). The following process aimed for a further optimization with regard to cracking resistance, but with regard to the use of this alloy in structural application mechanical properties should also be optimized.

Al0.5CoCrCu0.1FeNi was chosen as the baseline, and Al and Ni were varied the due to their effect on the STR, on phase formation, and on the composition of the liquid phase at the end of solidification. It should be point out that the current work only varies two elements in the alloy composition. The process of selecting the main influencing elements may be much more complicated depending on the alloy, or weldability issue. A

Python script was developed to run hundreds to thousands of Scheil calculations in 82

ThermoCalc software depending on composition range and increment size. The calculations were then compiled into 3D-plots of the predicted STR, and 2D-maps of the phases formed during solidification. Errors would arise the first few times these calculations were done, such as unreasonable solidus and liquidus temperatures or phases that needed to be rejected. These errors were addressed by modifying the Python script, and the calculations were done again in an iterative process. Once adequate results had been achieved, five alloy compositions were chosen for experimental testing and validation. Predicted and experimental data were compared. In case of significant deviations between the data, either the Python script was modified, or the selection criteria were changed. If the data matched well, and promising weldability was achieved

(0% cracking in CPTT, and no cracking in GTAW), the alloy composition would go on linear GTA and laser welding and may be recommend for good weldability.

Figure 42: Schematic of proposed high-throughput weldability screening using computational and experimental tools.

5.5.2 High-Throughput Scheil Calculations, Solidification Temperature Range and

Phase Maps

Figure 43 are four heat maps of the Scheil STR’s at (85, 95, 99, and 100) % solid fractions. The reason for having different the STRs at different fraction solid is to better represent solidification during welding. In all four fraction solid heat maps the

Al0.5CoCrCu0.1FeNi (green circle) STR is around average and there are many other compositions with lower STR. The STR at 95 % and 85 % fraction solid both have the highest STRs between 10-14.5 at% Al. Compositions with Al greater or less than this range, the STR decreases gradually to the left but Al % greater than 14.5 % there is a large drop in the STR. This large drop can be associated with the initial phase that solidifies going from FCC to BCC. Ni doesn’t have as large of an effect on the STR as

Al, but at Ni amounts less than 20% do show to decrease in the STR with increasing Al

% and stop around 14%. The STR at 100 % and 98 % fraction solid both have very similar results with very little differences. However, one distinct difference for the 98 % fraction solid was at Al less than 7-7.5 % the STR decreased by nearly 200oC. The difference between the two would imply a steep slope in the Scheil (temperature vs. fraction solid) curve. As discuss in the background chapter Al stabilizes the BCC order and disorder phases which can be seen to still hold true for the Scheil solidification heat maps in Figure 43. Ni however stabilizes the discorded FCC phase and in the top left corner of the heat map this region FCC is the only phase present during solidification.

Figure 43: Heat maps of the Scheil STR’s at (85, 95, 99, and 100) % fractions solid (Fs) for Ni vs Al and overlaid are the phases present during solidification. The green circle is the STR of the Al0.5CoCrCu0.1FeNi HEA.

Presented in Figure 44 is the Ni vs Al Scheil results of the solidification phases and acceptable STR (in blue). The different phase regions that are outlined in the map as

“FCC”, “FCC+B2”, and “B2+FCC” are the phases that form on solidification for the corresponding alloy compositions. The difference between “FCC+B2” and “B2+FCC” is that the phase listed first is the initial phase that forms at the liquidus temperature and represents the majority of the volume fraction. The region labeled as “Acceptable STR

o area” depicts STR values less than Al0.5CoCrCu0.1FeNi HEA [269 C @ 95%] at different solid fractions. Though the Al0.5CoCrCu0.1FeNi alloy showed promising resistance to solidification, CPTT results showed that there was still room for improvement in terms of cracking resistance. That is why an STR of less than 20oC was chosen as the limit for this 85

criterion. Figure 44 shows five yellow stars, which depict the alloy compositions chosen for further experimental weldability evaluation. The following selection criteria were used:

1) Compositions with an acceptable STR,

2) Compositions with a balance of (brittle vs. ductile) microstructure (primarily

determined by formation of B2 phase), and

3) Light considerations of mechanical properties.

It was deemed important to explore all of the acceptable STR ranges shown in Figure 44, even if there may be other weldability or mechanical properties concerns. For example, alloy #1 has a predicted fully FCC microstructure on solidification. In the welding metallurgy world, it is commonly known that a primary FCC solidification mode will increase the likelihood of solidification cracking due to the slow diffusion in solid increasing partitioning. This is especially true when impurity elements are introduced into the alloy composition which often happens outside of the laboratory environment.

Nonetheless because this is new welding metallurgy territory all possibility should be reviewed so nothing is overlooked. Alloys #2, #3, and #4 were all selected to meet criteria 2) but also had a low STR as seen in Table 8. The formation of the strengthening

BCC (B2) in this alloy system is primarily a function of aluminum content with higher amounts of aluminum increasing the amount of B2 phase. This is why these alloys are located on the left side of this STR range region (Figure 44) to keep it aluminum percentage as low as possible. However, an increasing nickel content increases the amount of disordered FCC, which helps to reduce the hardness of the solidification microstructure. Alloy #5 was selected for both criteria 1) and 2), but the exact 86

composition was selected mainly based on criteria 3). Reducing the aluminum content is often beneficial to prevent an extremely hard, brittle microstructure, but reducing aluminum too much and the alloy will become too ductile with a low yield strength. This is why alloy #5 is located on the right side of that STR range region (Figure 44) in order to maximize aluminum amount for strength, but avoiding the formation of too much BCC

(B2) phase in the B2+FCC region.

Figure 44: Ni vs. Al phase map and acceptable STR range. The 5 yellow stars depict the alloys chosen for further weldability evaluation.

Table 8: The five chosen alloys composition and STR for experimental evaluation

Co, Cr, Fe Alloy # Al at% Ni at% at% Cu at% STR 95% 1 7 27 21.67 1 155oC 2 14 15 23.33 1 97oC 3 15 21 21 1 94oC 4 15 24 20 1 90oC 5 11 16 24 1 212oC

5.5.3 Gas Tungsten Arc Stationary Welding, Solidification Microstructures and

Micro-hardness on Bulk As-Melted Microstructures

Figure 45 (a, b, c, d, & e) shows cross-sections of stationary GTA welds made on

“button-shaped” ingots for alloys Al7Co21.7Cr21.7Cu1Fe21.7Ni27 (#1),

Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2), Al15Co21Cr21Cu1Fe21Ni21 (#3),

Al15Co20Cr20Cu1Fe20Ni24 (#4), and Al11Co24Cr24Cu1Fe24Ni16 (#5) in the as-melted condition. The as-melted solidification microstructure for each alloy has very large columnar grains, which are on the scale of millimeters in the vertical direction. For each alloy, the fusion zone is outlined in red dashed lines. All fusion zones are approximately the same size with a weld penetration of about 1 mm. The main objective was to use the

GTA stationary welds on the as-melted ingots as a quick and easy screening tool to assess the susceptibility to cracking in fusion welding of the selected alloy compositions.

However, no cracking was observed for any of the alloys in the fusion zone, HAZ, or the as-melted ingot microstructure. This result indicated that these alloy compositions may be quite resistance to cracking. However, only one GTA stationary weld was performed for

each alloy composition, and multiple samples should be done in the future to confirm these results.

Figure 45: Cross-sections of autogenous GTA stationary weld on as-melted “button-shaped” ingots for each alloys #1 to #5: a) Al7Co21.7Cr21.7Cu1Fe21.7Ni27 (#1), b) Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2), c) Al15Co21Cr21Cu1Fe21Ni21 (#3), d) Al15Co20Cr20Cu1Fe20Ni24 (#4), and e) Al11Co24Cr24Cu1Fe24Ni16 (#5). All light optical micrographs. Molybdenum etchant.

Figure 46 shows the same GTA spot welds, but at higher magnifications at the fusion line. The solidification microstructure in the fusion zone is cellular or cellular dendritic for all alloys except for the Al11Co24Cr24Cu1Fe24Ni16 alloy (#5). The fusion zone in Figure 46 (e) does not show a dendritic solidification microstructure but appears to be a two phase lamellar structure possibly due to solid-state transformations on cooling.

Compared to the Al0.5CoCrCu0.1FeNi alloy the interdendritic regions of these five alloys are much less pronounced, which may be associated with a decrease in elemental segregation to the interdendritic regions during solidification. The HAZ microstructure appears quite different for each alloy, but the solidification microstructure disappears in the high temperature HAZ near the fusion line for alloys Al7Co21.7Cr21.7Cu1Fe21.7Ni27

(#1), Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2), Al15Co21Cr21Cu1Fe21Ni21 (#3), and

Al15Co20Cr20Cu1Fe20Ni24 (#4), see Figure 46 (a-d). Dissolution of the interdendritic phase and other phases close to the fusion line may be beneficial for preventing or reducing the partial melting of these phases. This is likely the reason no liquation cracks were seen in the HAZ for any of the alloys. Phase changes and/or precipitation can also be seen in different regions throughout the HAZ for all alloys. However, no in-depth characterization of the HAZ microstructures was performed in this research.

Previous results in this work of GTA stationary welds on the AlCoCrCuFeNi alloy system in the as-melted condition showed that increasing Al content leads to a significant increase in hardness, which led to brittle cracking in the fusion zone and around the weld region. Some of the selected alloys have 14-15 at% Al, which led to a high hardness of about 550 HV (Table 9), but no brittle cracking was seen in any of the

GTA welded ingots. Solidification cracking in absent for all welds possibly due to the 91

narrow STR of the alloys. However, these welds should be repeated, and full-scale welding needs to be done in future work to confirm the results. The

Al11Co24Cr24Cu1Fe24Ni16 alloy (#5) exhibits a hardness of about 471 HV, which is higher than expected considering the Al content of 11 at%. Similar HEA compositions with

11 at% Al, but 21% Ni have been reported with a hardness between 200-300 HV [15].

This could indicate that the reduction in Ni stabilizes less of the FCC phase, which is what leads to the higher hardness. The alloy with the lowest hardness is

Al7Co21.7Cr21.7Cu1Fe21.7Ni27 (#1) with about 160 HV. This alloy has the highest Ni content and lowest Al content. Another observation of the hardness but was not extensively study, was the softening in the high temperature HAZ for both alloys with

15% Al at%. In Figure 46 (c & d), Widmanstatten plates (white color) can be seen in the

HAZ, which is likely the growth of FCC phase [15] and causes the reduction in hardness in the HAZ.

Figure 46: Detailed cross-section of autogenous GTA stationary weld on as-melted “button- shaped” ingot at the fusion line. a) Al7Co21.7Cr21.7Cu1Fe21.7Ni27 (#1), b) Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2), c) Al15Co21Cr21Cu1Fe21Ni21 (#3), d) Al15Co20Cr20Cu1Fe20Ni24 (#4), and e) Al11Co24Cr24Cu1Fe24Ni16 (#5). All light optical micrographs. Molybdenum etchant.

Table 9: Solidification morphology and micro-hardness in the fusion zone.

# Alloys (at%) Solidification Morphology Hardness 1 Al7Co21.7Cr21.7Cu1Fe21.7Ni27 Cellular dendritic 160 ± 5 2 Al14Co23.3Cr23.3Cu1Fe23.3Ni15 Cellular dendritic 555 ± 13 3 Al15Co21Cr21Cu1Fe21Ni21 Cellular dendritic 555 ± 7 4 Al15Co20Cr20Cu1Fe20Ni24 Cellular dendritic 541 ± 12 5 Al11Co24Cr24Cu1Fe24Ni16 Lamellar 471 ± 11

5.5.4 Cast Pin Tear Testing (CPTT)

Cast pin tear testing was done at a single mold length of 1 inch. Conventional Ni- based alloys that show no cracking in the CPTT at this length have shown good resistance to solidification cracking [38]. Table 10 shows the results as circumferential cracking % measured in the two cast pin samples that were tested for each alloy at 1 inch pin length. It can be seen that that Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2) and

Al15Co20Cr20Cu1Fe20Ni24 (#4) exhibited cracking in at least one of the two tested pins.

This would indicate that these alloys are somewhat susceptible to solidification cracking.

However, alloys with this high hardness have never been tested in the CPTT machine outside this work. The Al14Co23.3Cr23.3Cu1Fe23.3Ni15 alloy (#2) shows 0% and 100% cracking in the two tested pins which that big of a difference should have never occur. Its hypothesized that small solidification cracks or casting defects initially occur on the casted pins surface which acts as a stress concentrator. As the metal continues to cool, tensile strain builds up. Once high enough stresses are reached at the crack tip, the crack easily propagate through the brittle microstructure of the pin. Figure 47 shows the fracture surface of a cast pin of Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2) with a smooth fracture surface near the top-outer surface of the pin. The rest of the fracture surface shows a cleavage fracture mode. This might be an indication that the CPTT is not a suitable 94

option for alloys with such high hardness, since it might drastically affect the results. The

Al7Co21.7Cr21.7Cu1Fe21.7Ni27 (#1) alloy had a much lower hardness and exhibited no cracking in any of the pins. Both Al15Co21Cr21Cu1Fe21Ni21 (#3) and

Al11Co24Cr24Cu1Fe24Ni16 (#5) alloys also exhibited no cracking in the CPTT despite having a high hardness, which shows that these alloys too have promising resistance to solidification cracking. More CPTT should be done for these alloy compositions, but these alloys show very good initial resistance to solidification cracking.

Table 10: Results of cast pin tear testing (CPTT) at 1” mold length, shown as circumferential cracking %. Average # Alloys (at%) Pin #1 Pin #2 Cracking % 1 Al7Co21.7Cr21.7Cu1Fe21.7Ni27 0 0 0 2 Al14Co23.3Cr23.3Cu1Fe23.3Ni15 0 100 50 3 Al15Co21Cr21Cu1Fe21Ni21 0 0 0 4 Al15Co20Cr20Cu1Fe20Ni24 30 4 17 5 Al11Co24Cr24Cu1Fe24Ni16 0 0 0

Figure 47: SEM image of the Al14Co23.3Cr23.3Cu1Fe23.3Ni15 fracture surface from a pin with 100% cracking.

Chapter 6. Conclusions

This work explored the welding metallurgy and potential weldability issues with regards to weld cracking susceptibility, segregation behavior, and phase formation for

AlCoCrCuFeNi, AlCoCrFeNiTi and AlMoNbTaTiZr high entropy alloys. An initial characterization of the welding metallurgy of HEAs was performed. From this knowledge, a methodology was developed that enables the screening of a large number of

HEA compositions for promising alloy compositions with resistance to weld solidification cracking using high-throughput CALPHAD-based solidification calculations and experimental testing techniques.

6.1 HEA System AlxCoCuyCrFeNi

1. The equiatomic HEA (AlCoCrCuFeNi), showed poor weldability with large

cracks in the HAZ that were characterized as liquation cracking due to low

melting Cu-rich interdendritic regions in the arc-melted ingot microstructure.

2. Weldability testing revealed, that Cu segregation to the grain boundaries and

interdendritic regions during solidification promotes liquation and solidification

cracking in the fusion zone and HAZ when the Cu content is greater than 0.1

molar ratio.

3. Aluminum promotes a high hardness microstructure, which leads to brittle

fracture (transgranular cracking) in the fusion zone and HAZ.

4. Cracking in autogenous GTA spot welds on arc-melted “button-shaped” ingots

was for all tested alloy compositions in good agreement to CPTT results. Lots of

cracks in the fusion zone and HAZ of the GTA welds and in the cast pin samples

were observed at high Cu content, but compositions below 0.3 Cu ratio showed

much less cracking.

5. Decreasing Cu and Al content in the AlCoCrCuFeNi HEA system mitigates both

hot cracking (liquation and solidification cracking) and brittle cracking.

6.2 Comprehensive Weldability Evaluation of Al0.5CoCrCu0.1FeNi

1. The solidification microstructure was comprised of FCC dendrites and

interdendritic disordered BCC and ordered BCC (B2) phase. Al and Ni partition

to the liquid phase during solidification. A fine-scale modulated structure forms in

the interdendritic BCC regions. An Al- and Ni-rich layer is observed along the

edge to the dendrite matrix. No Cu-rich interdendritic phase was observed.

2. Predictions of Scheil calculations in Thermo-Calc using the current HEA database

(TCHEA3) are in good agreement with the weld solidification microstructure

observed experimentally. The solidification temperature range (STR) from the

Scheil simulation is about 270oC at 95% solid fraction. Calculated equilibrium

phase formation resembles the precipitation of a strengthening BCC (B2) phase

during heat treatment. An increase in hardness was experimentally observed due

to the precipitation of the BCC(B2) phase.

3. The GTA welds were free of solidification cracking or heat-affected zone (HAZ)

liquation cracking for all alloy conditions and welding setups. Partial dissolution

of the interdendritic phase of the initial arc-melted ingot microstructure was

observed in the high temperature HAZ adjacent to the fusion boundary. This

might explain the absence of HAZ liquation cracking in the welds, since the

formation of low-melting liquid might be inhibited in the absence of the

interdendritic phase close to the fusion line.

4. Welding on the heat-treated alloy resulted in HAZ softening due to the dissolution

of the strengthening BCC (B2) phase in the dendrite matrix and interdendritic

regions.

5. The laser weld on the heat-treated alloy showed no drop in hardness or apparent

microstructural changes in the HAZ. No HAZ liquation cracking was observed in

keyhole welding mode. Several cross-sections exhibited trapped porosity,

probably due to inadequate welding parameters. In conductive welding mode, no

solidification or liquation cracking was observed.

6. Results of the cast pin tear test (CPTT) showed a cracking threshold of

0.625 inch, which indicates that Al0.5CoCrCu0.1FeNi alloy is much more resistant

to solidification cracking as compared to AlCoCrCuFeNi alloys with a higher Cu

content.

7. The alloy is characterized by an overall good weldability using GTAW and laser

beam welding processes. 99

6.3 HEA System AlCoCrFeNiTi

1. The solidification microstructure was comprised of primarily FCC dendrites and

interdendritic FCC and Gamma prime phases. The interdendritic region was

found to be enriched in Al and Ti, which partition to the liquid phase during

solidification.

2. In the overaged condition (1220°C for 30 min), large precipitates formed along

grain boundaries and intragranularly that were enriched in Al, Ni, and Ti. These

are thought to be large Gamma prime precipitates.

3. Predictions of Scheil calculations in Thermo-Calc using the current HEA database

(TCHEA3) are in good agreement with the weld solidification microstructure

observed experimentally. Calculated equilibrium phase formation resembles both

Gamma prime and BCC (B2) phase during heat treatment. Equilibrium

calculations over predicted the BCC (B2) phase near the solidus which should of

only had a single FCC phase.

4. Autogenous GTA spot welds on arc-melted ingots showed no solidification or

liquation cracking. Autogenous GTA spot welds performed on the over-aged

alloy condition showed numerous liquation cracks in the partially melted HAZ.

Intergranular cracking was also observed further away from the fusion line.

5. Autogenous pulsed laser welding performed on the over-aged alloy condition

showed no solidification or liquation cracking, probably due to the decrease in

heat input.

100

6. Post weld-solutionizing heat treatment done at 1200oC for 30 minutes resulted in

severe grain growth and severe intergranular cracking in the fusion zone and the

base metal around the weld.

6.4 HEA System AlMo0.5NbTa0.5TiZr

1. The solidification microstructure showed BCC (B2) dendrites and an

interdendritic region enriched in Al and Zr.

2. Scheil calculations in Thermo-Calc using the current HEA database (TCHEA3)

predicted Al and Zr to partition into the liquid phase during solidification. The

predicted STR was 535oC at 95% solid fraction.

3. During arc-melting of “button-shaped” ingots this alloy composition was found to

be very susceptible to porosity and brittle cracking at the center of the ingot.

4. Autogenous GTA spot welds on the as-melted ingot microstructure showed no

cracking in the fusion zone or HAZ. The fusion zone exhibited shallow

penetration and an equiaxed dendritic solidification morphology with no epitaxial

nucleation from the fusion line.

6.5 Compositional Optimization for Solidification Cracking Resistance using High-

Throughput Screening Calculations and Experiments

1. A methodology was developed to optimize HEA compositions for weldability

using high-throughput CALPHAD-based solidification calculations and

experimental testing techniques. Focus in this research was to optimize for

solidification cracking resistance during fusion welding for structural engineering 101

applications. The methodology can be adapted to other weldability issues,

welding processes and applications.

2. A Python program was developed, which enables automated Scheil solidification

calculations in Thermo-Calc on a 671 alloy compositions. Another Python script

was developed to compile the obtained data for visualization in to order to enable

the selection of promising alloys with regard to good weldability.

3. High-throughput weldability screening enabled the selection of five different

alloy compositions based on generated heat maps that depict calculated STR and

phase balances for a large number of alloy compositions with varying Al and Ni

content.

4. Autogenous GTA stationary welds performed on the as-melted microstructure for

all selected alloy compositions served as a high-throughput experimental tool to

screen the weldability of the alloys. No cracking was observed in the fusion zone

or HAZ in any of the alloy compositions.

5. Cast pin tear test results of the Al7Co21.7Cr21.7Cu1Fe21.7Ni27 (#1),

Al15Co21Cr21Cu1Fe21Ni21 (#3) and Al11Co24Cr24Cu1Fe24Ni16 (#5) showed no

cracking at 1-inch mold length. Alloys Al14Co23.3Cr23.3Cu1Fe23.3Ni15 (#2) and

Al15Co20Cr20Cu1Fe20Ni24 (#4) did experience cracking at the same pin length.

However, CPTT may not be a suitable option for alloys with high hardness above

400 HV which will drastically affect the results.

6. Reducing the STR below 200oC was found to be an effective way in reducing

cracking in the GTA stationary welds and (CPTT) pins.

102

Chapter 7. Recommendations and Future Work

1. No mechanical testing was done for any of the HEAs in this work. In the future,

tensile testing should be implemented as a part of the weldability study for both

ambient and high temperatures.

2. DSC and other thermal techniques should be used to validate the solidus and

liquidus temperatures from the Scheil calculations. XRD analysis should also be

utilized to validate the predicted solidification phases.

3. This work primarily focused on hot cracking, however, solid-state cracking might

have occurred in the Al10Ti6Cr8Fe15Co25Ni36 at% alloy. Weldability testing for

strain-age cracking should be looked at in the future so composition modification

can be implement into HEA development.

4. The five alloys from section 5.5, both linear GTA and pulsed laser welding should

be done on these alloys and possibly implemented into the high-throughput

screening method.

5. Exploring other HEAs to optimize their composition for solidification cracking

using the high-throughput CALPHAD screening method. The

Al10Ti6Cr8Fe15Co25Ni36 at% alloy would be a great option to research next by

varying Al, Ni, and Ti.

6. Cast pin tear testing (CPTT) is a great test for high-throughput screening, but this

work found some issues that still need to be addressed, as listed below.

103

• How do alloys with high hardness affect cracking length and what’s the

maximum hardness that can be used?

• Should 1-inch pin length be deemed resistant to solidification cracking for

HEAs or other non-nickel-based alloys? Do differences in coefficient of

thermal expansion affect the strain, which ultimately effects circumferential

cracking length?

• Better standardization for counting circumferential crack length (what area of

the pin should cracks be counted, should cracks be counted if they occur from

casting defects, if cracks occur at the bottom of the pin should it be rejected)?

104

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