Deposition Kinetics of Titanium and Zirconium Diffusion Coatings on Nickel Microwires Via Pack Cementation

Deposition Kinetics of Titanium and Zirconium Diffusion Coatings on Nickel Microwires via Pack Cementation

A thesis submitted to the Division of Research and Advanced Studies of the University of Cincinnati in partial fulfillment of the requirements for the degree of

Master of Science

in the program of Materials Science and Engineering in the Department of Mechanical and Materials Engineering of College of Engineering and Applied Science University of Cincinnati, Ohio, USA

Ajith Achuthankutty

Committee Members: Dr. Ashley Paz y Puente (Chair) Dr. Matthew Steiner Dr. Sarah Watzman

Fall 2019

ACKNOWLEDGEMENT

My grad life, full of ups and downs, could not have been fulfilled without the guidance, support, and cooperation of so many people. Firstly, I thank god almighty for helping me with all the resources in this venture. Next, I am deeply indebted to my mentor and advisor, Dr. Ashley Paz y Puente, for giving me this fantastic opportunity to work in her lab. Her support and continued motivations were the major driving forces during my master’s program at the University of

Cincinnati. It was a little difficult for me initially to change track from mechanical to material science, but her consistent encouragement and training made it easy for me to hone my skills.

I am also grateful to her for supporting me with Teaching Assistantships, which helped to manage classes, improve my knowledge, presentation, and soft skills.

I am really thankful to Dr. Dinc Erdeniz for providing the training and installation of the lab equipment and giving me ideas and guidelines for this research. I would also like to acknowledge Dr. Matthew Steiner and Dr. Sarah Watzman for agreeing to be on my defense panel and supporting my work with valuable suggestions. I thank Dr. Melodie Fickenscher for helping and training me on SEM micrographs, elemental analysis data, and other characterization tools that were prominently used for this thesis work.

No words can express my respect and love towards my parents Mrs. Jayasree Nair and

Mr. Achuthankutty Nair and my siblings Arun Nair and Neethu Krishna. They were confident enough to send me abroad to achieve my dream and passion. My extreme gratitude to Arathi

Anil Sushma for helping through my hard times and motivating me in completing my degree

and finding a suitable job. She has been a pillar and a great role model to me both academically and otherwise.

I am really blessed to have such extremely supporting and helpful lab-mates. I thank

Arun Bhattacharjee, Pavan Ajjarapu, Haozhi Zhang and Safa Khodabakhsh for assisting me with my research and giving timely inputs and suggestions.

My life at Cincinnati would not have been a memorable one without my friends; Abhishek

Pai, Lavanya Varma, Anjaly Vijayan, Madhavan Sudhakar, Ketan Shah, Jose Joseph, Gibin

Raju, Smriti Jha, and Anantharaman Ashok. Finally, a special thanks to my classmates Pavan

Kandala, Suriya Sekar, Sakshi Satyanarayana and Biswajit Dalai for all the fun-packed times during the course.

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ABSTRACT

Pack cementation is a promising technique to synthesize hollow NiTi microtubes, which impart enhanced heating/cooling efficiency and, thus, faster shape memory response. Prior research investigating the effect of wire size on phase evolution and Kirkendall pore formation during pack titanization showed a significant difference between large (100 µm) and small (25

µm) diameter wires. The first objective of this research is to quantify the titanium deposition kinetics on intermediate-sized Ni wire (75 µm diameter) and determine the required coating time to achieve the desired near-equiatomic NiTi composition. It also explores an ex-situ study of the formation of a dual-pore structure that forms in some of the wires and provides insight for future in-situ investigations of the wire size effect on pack titanizing.

Even though binary NiTi shape memory alloys offer excellent biocompatibility, strength, and corrosion resistance, their application is limited to less than 100 ºC. Hence, there is a need to increase the transformation temperatures for high-temperature applications, especially for aerospace applications. Certain ternary alloying additions to NiTi form high temperature shape memory alloys (HTSMAs) as there is an associated increase in the austenite-martensite transformation temperature and, therefore, have attracted considerable attention from researchers. Ni-Ti-Zr is one such HTSMA that provides the additional advantages of weight and cost reduction as compared to its platinum (Pt), gold (Au) and hafnium (Hf) counterparts.

In the current research, attempts were made to produce Ni-Ti-Zr wires by co-depositing Ti and

Zr on Ni wires using a halide activated pack in a single step process. The samples gained Zr content, but not enough was deposited to increase the transformation temperature. Moreover,

a significant variation in the coating thickness was observed, resulting in a significant fluctuation in the composition.

Therefore, as a first step, deposition of Zr on Ni wires was conducted to assess the feasibility of a two-step coating process to achieve the desired ternary Ni-Ti-Zr composition.

The deposition kinetics, intermetallic phases formed, coating thicknesses, and composition were determined at 925 °C for various coating durations. The samples reached approximately

12-13 at. % Zr after 3 hours of coating. The as-coated wires, upon characterization via optical and laser microscopy, SEM, and EDX, displayed a core-shell structure with intermetallic phases which appeared to be Ni7Zr2 and Ni3Zr around the Ni core demonstrating the feasibility of the deposition process. The outcome of the present study will provide insight to further research regarding fabrication of these HTSMA microtubes, which can be utilized in the automotive, aerospace, and oil industries.

LIST OF FIGURES

Figure 1. (a) NiTi self-expandable neurosurgical stent (b) percutaneous aortic valve realized

with eNiTinol membrane [2]...... 5

Figure 2. Schematic of Shape Memory Effect: Upon cooling austenite transforms to twinned

martensite, which upon applying stress deforms and forms single variant

martensite. The original austenite phase is “remembered” upon heating [20]...... 7

Figure 3. Schematic illustration of mechanical response of pseudoelastic material [24]...... 9

Figure 4. Martensite start temperature as a function of Ni content (at.%) [3]...... 11

Figure 5. 3D printed Nitinol rectangular block using selective laser melting [7]...... 12

Figure 6. Particle size distributions and resulting the austenite finish temperatures [27] ...... 13

Figure 7. Schematic representation of the particle-based ink writing process [28] ...... 14

Figure 8. Binary NiTi phase diagram...... 15

Figure 9. Effect of ternary alloying additions on the transformation temperatures of Nitinol SMA

[14]...... 17

Figure 10. Aluminide coating on a Ni-base material [34]...... 19

Figure 11. Schematic showing the various constituents of the pack, the chemical reactions, and

the gas-phase and solid-state diffusion...... 20

Figure 12. Schematic of aluminum-rich coating formed on surface of a Ni wire cross-section

via pack cementation...... 21

Figure 13. Optical micrographs showing the growth of the coating during the aluminization

process on nickel wires for (a) 300 s and (b) 600 s. The inner core is γ-Ni and the

outer shell is δ-Ni2Al3 [37]...... 22 vi

Figure 14. Effect of titanization temperature (1000 °C) on porosity within and non-uniformity of

the coating [12]...... 23

Figure 15. Schematic of Kirkendall experiment showing Mo wires moving inward from their

original position, Kirkendall 1947 [39]...... 25

Figure 16. Schematic of (a) direct exchange mechanism, (b) ring mechanism, and (c) vacancy

mechanism of diffusion [39]...... 26

Figure 17. (a) Hypothetical binary A-B diffusion couple (b) SEM image of the Kirkendall voids

found at the SnPb solder/Cu pad interface after aging at 1508 ºC for three days

[28,42]...... 27

Figure 18. (a) Backscattered electron micrograph of a radial cross-section of a titanized Ni wire

following homogenization for 4 h at 925 °C (b) EDS profile confirming the uniform

near-equiatomic NiTi composition [9]...... 28

Figure 19. Tomographic images showing the continuity of the Kirkendall pore at three different

annealing times: (a) 2h (b) 4h (c) 8h (d) reconstructed cross-section, and (e) 3D

visualization of a magnified section of the pore [9]...... 29

Figure 20. Arrhenius plot of intrinsic diffusivities of nickel and titanium [43]...... 30

Figure 21. Backscattered electron micrographs for (a) 25 μm Ni wire titanized for 30 min (b)

100 μm Ni wire titanized for 8 h [11]...... 31

Figure 22. (a) SentroTech tube furnace with inert gas inlets and outlets for pack cementation

(b) Sentro Tech box furnace used for homogenization anneal of the samples..... 33

Figure 23. Allied MetPrep4 for semi-automatic polisher...... 34

Figure 24. Keyence VKX-250X 3D laser confocal microscope...... 34

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Figure 25. Thermofischer APREO SEM similar to the one used at AMCC...... 35

Figure 26. Plot of the phase area fraction as a function of the titanization time...... 37

Figure 27. Optical micrographs of cross-sections of NiTi wires titanized in 30 wt.% Ti pack for

(a) 15min, (b) 30 min, (c) 60 min, (d) 90 min, (e) 120 min, and (f) 150 min. The

intermetallic layers and pores formed are indicated and labeled...... 39

Figure 28. Kinetics of titanium gain (at. %) for NiTi wires titanized at 925 °C. The colored bands

designate the different phase fields on the binary phase diagram at 925 °C...... 41

Figure 29. Optical micrographs showing radial cross-sections of titanized 75 μm diameter Ni

wires homogenized at 925 °C for (a) 8 h and (b) 16 h. Wires were titanized at 925

°C for 150 min, with a pack change after 120 min...... 42

Figure 30. DSC curves for a) equiatomic NiTi and b) with 10 wt.% Hf [35]...... 44

Figure 31. Transformation temperature change in NiTi as a function of zirconium content [52].

...... 46

Figure 32. Optical micrograph of titanized 75 µm nickel wire coated for 90 min using TiC as the

filler. Some TiC particles are engulfed by the wires as the diffusion coating grows.

...... 47

Figure 33. (a) Micrograph showing black aggregates after 2 hours of titanization using Al2O3

at 1000 °C (b) intensity plot showing the presence of Al [12] ...... 48

Figure 34. Ellingham diagram for chlorides. The Gibbs free energy of formation of Zr and Ti at

925°C has been depicted [54]...... 50

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Figure 35. Micrographs of cross-sections after simultaneous Ti and Zr deposition using 66 wt.%

ZrO2 as the filler and 30 wt.% Ti powder as the source coated for (a) 60 min and

(b) 120 min at 925 °C...... 53

Figure 36. (a) and (b) show two different optical micrographs of the wires homogenized for 24

h after initially coating for 120 min using 66 wt.% ZrO2 and 30 wt.% Ti source

powder (images from Dinc Erdeniz)...... 54

Figure 37. The Co-deposition of Zr and Ti at 925 ºC for (a) 60 min and (b) 120 min...... 55

Figure 38. (a) The magnified BSE micrograph of a sample coated for 120 min. (b) The EDS

line profile for the co-deposited wires...... 55

Figure 39. Homogenized for 24 h after 120 min coating using 10 wt.% Ni-Zr pre-alloyed pack.

...... 56

Figure 40. The EDS line profile for co-deposited samples after homogenization for 24 h. The

samples were brittle and broke during metallographic preparation...... 57

Figure 41. BSE micrographs for Ni wires co-deposited with Ti and Zr ( 20 wt. % Zr source

powder) for 120 min. The wires were highly porous...... 58

Figure 42. The diffusivity of zirconium in the Ni matrix (LDA Debye line) showing the diffusion

of Zr in Ni at 925 °C...... 59

Figure 43. Elemental maps of (a) nickel (b) zirconium (c) titanium (d) carbon on the zirconized

samples after 2 h of coating...... 63

Figure 44. SEM micrograph of Zr coated Ni wire for 120 min and corresponding EDS line scan

showing the varying composition of Zr and Ni...... 64

Figure 45. EDS spot analysis results of zirconized nickel wires for 120 min at 925 ºC...... 65

Figure 46. EDS spot analysis results of zirconized nickel wires for 120 min at 925 ºC...... 65

Figure 47. Binary Ni-Zr phase diagram showing the different phases formed at 925 °C...... 66

Figure 48. BSE micrographs representing the various layers formed and phase evolution in Ni

wires zirconized for (a) 30 min (b) 60 min (c) 90 min (d) 120 min (e) 150 min and (f)

180 min...... 67

Figure 49. Rate of deposition of Zr on Ni wires as a function of coating time...... 69

LIST OF TABLES

Table 1. The diameter of the NiTi wires as Ti is added as a function of coating time...... 38

Table 2. Composition from spot analysis of wires after co-deposition for 120 min ...... 56

Table 3. The diameter as Zr is deposited on Ni wires as a function of coating time ...... 68

Table 4. Presents the coating time and the average and standard deviation at.% of Zr

obtained, depicting the variation among the same batch of samples...... 69

LIST OF ACRONYMS/ABBREVIATIONS

BSE Back Scattered Electron

CVD Chemical Vapor Deposition

DM Detwinned Martensite

DSC Differential Scanning Calorimetry

EDS Energy Dispersive Spectroscopy

HTSMA High Temperature Shape Memory Alloy

NiTi Nickel Titanium

Ni-Ti-Hf Nickel Titanium Hafnium

Ni-Ti-Zr Nickel Titanium Zirconium

SEM Scanning Electron Microscopy

SMA Shape Memory Alloy

SME Shape Memory Effect

TM Twinned Martensite

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TABLE OF CONTENTS

ACKNOWLEDGEMENT ...... ii

ABSTRACT...... iv

LIST OF FIGURES ...... vi

LIST OF TABLES ...... xi

LIST OF ACRONYMS/ABBREVIATIONS ...... xii

1 Introduction ...... 1

2 Background ...... 5

2.1 Shape Memory Effect ...... 5

2.1.1 Introduction ...... 5

2.1.2 Stress-Induced Martensite (SIM) ...... 7

2.2 NiTi Shape Memory Alloys ...... 10

2.2.1 History and Applications ...... 10

2.2.2 Fabrication Techniques ...... 10

2.3 Effect of Ternary Alloying Additions to NiTi ...... 16

2.3.1 History and Applications ...... 16

2.4 Pack Cementation...... 18

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2.4.1 Introduction and History ...... 18

2.4.2 Pack Composition and Thermodynamics ...... 19

2.5 Solid-State Diffusion and Kirkendall Effect ...... 23

2.5.1 Introduction to Diffusion ...... 23

2.5.2 The Kirkendall Effect ...... 25

3 Fabricating NiTi Microtubes via the Kirkendall Effect ...... 28

3.1 Introduction and Literature ...... 28

3.2 Experimental Procedure ...... 32

3.2.1 Microstructure Characterization...... 33

3.3 Results and Discussion ...... 36

3.3.1 Coating Thickness and Phase Evolution ...... 36

3.3.2 Pore Formation ...... 38

3.3.3 Deposition Kinetics ...... 40

3.3.4 Homogenization ...... 42

4 Co-deposition of Zr and Ti on Ni Microwires ...... 43

4.1 Introduction and Literature ...... 43

4.2 Effect of Zirconium on NiTi ...... 44

4.2.1 Effect of Filler in Pack Titanizing ...... 46

4.2.2 Thermodynamics of the Reaction ...... 48

xiv

4.3 Experimental Procedure ...... 50

4.3.1 Homogenization ...... 52

4.4 Results and Discussion ...... 52

4.4.1 Zr Source: ZrO2 Powder ...... 52

4.4.2 Zr Source: 10 wt.% Ni-Zr Pre-alloyed Powder ...... 54

4.4.3 Zr Source: 20 wt.% Ni-Zr Pre-alloyed Powder ...... 57

5 Kinetics of Zr Deposition on Ni Microwires ...... 59

5.1 Introduction and Literature ...... 59

5.2 Experimental Procedure ...... 60

5.2.1 Image Analysis ...... 61

5.3 Results and Discussion ...... 62

5.3.1 Zirconizing and Phase Evolution ...... 62

6 SUMMARY AND CONCLUSIONS ...... 70

6.1 Fabrication of NiTi Microwires using the Kirkendall Effect ...... 70

6.2 Co-deposition of Zr and Ti on Ni Microwires ...... 71

6.3 Kinetics of Zr Deposition on Ni Wires ...... 71

7 FUTURE WORK ...... 73

8 REFERENCES ...... 74

xvi

1 Introduction

The shape memory effect is a unique behavior where a material “remembers” its pre- deformed geometry by undergoing a reversible phase transformation during heating and cooling cycles. Nitinol is one such shape memory alloy of nickel (Ni) and titanium (Ti), which possesses, biocompatibility, high-strength, and resistance to oxidation and corrosion. The addition of Ti to Ni increases the corrosion resistance and impedes any Ni leaching, which can otherwise result in allergic reactions and other detrimental health effects [1]. The performance and application of this material can be altered by changing the composition, porosity, microstructure (e.g. grain size), and applying thermal and mechanical treatment [2]. For example, in actuator applications, making NiTi porous provides improved actuation response because of an enhanced heating/cooling efficiency and, for biomedical applications, it aids in osseointegration of implants by allowing cell ingrowth.

The automotive and aerospace industries are aiming at replacing multi-component assemblies by providing single-piece, adaptive, lightweight, and multifunctional components that can safely be employed at high temperatures [3]. Shape memory alloys are a class of such smart materials that serve as a potential solution. However, the upper limit for the transformation temperature (TT) is less than 120 °C [3-4]. Hence there is a motivation to increase the TT above a specific limit (≥ 150 °C), which would have potential applications in the aerospace industry. Previous research demonstrates that ternary alloying additions to NiTi allow for the development of High Temperature Shape Memory Alloys (HTSMAs) [5]. Adding

Au, Pt, and Pd increases the phase transformation temperatures well above 150 °C, but the

cost of production with these elements is extremely high, which makes them less than ideal for commercial applications [6]. Alloying with less expensive elements like Hf and Zr seems promising while maintaining all of the merits of shape memory behavior. There is sufficient understanding of the effects of Hf additions to NiTi, but Ni-Ti-Zr alloys have not been extensively studied. However, Ni-Ti-Zr alloys provide a tremendous reduction in weight compared to their Ni-Ti-Hf counterpart, which makes it worthwhile to explore.

There is a significant challenge in the fabrication of NiTi alloys due to the composition sensitivity of the shape memory behavior. For instance, a small change of ± 1 at.% Ni leads to a drastic change in the martensite start temperature [3]. Methods like Vacuum Arc Remelting

(VAR) and Vacuum Induction Melting (VIM) are commercially used to produce NiTi alloys, but these processes require exceedingly high temperatures (>1000 ºC) and post processing, which makes it quite expensive. Additive manufacturing is a suitable method which has been recently investigated, but there is a considerable amount of oxidation during the melting process [7].

Moreover, these processes are limited when wires less than 100 μm in diameter are considered for producing woven structures and their subsequent weaving is intricate [8]. One such approach to overcome this is to deposit a Ti coating on a pure Ni wire woven structure using pack cementation (a simple Chemical Vapor Deposition (CVD) process), and subsequently homogenizing them via interdiffusion to obtain near equiatomic NiTi wires [9]. The feasibility of this process was studied for 50 μm diameter Ni wires and NiTi microtubes were synthesized by AE Paz y Puente et al. [10].

Literature regarding the pack titanization of large diameter (100 μm) and small diameter (25

μm) Ni wires to examine the effect of wire size shows a significant difference in Kirkendall pore 2

formation and phase evolution within the wires [11]. The initial objectives of this work were to assess the deposition kinetics during pack titanization of 75 µm Ni wires by evaluating coating thickness and phase evolution and to determine the coating time required to attain the desired near-equiatomic NiTi composition. This knowledge will provide insight for further in-situ studies on the mechanism of pore formation and the wire-size effect on the NiTi system.

While other researchers were studying the kinetics and fabrication of NiTi wires using zirconia filler material in the pack, they observed the final as-coated samples were contaminated with significant amounts of zirconium [12]. Elemental analysis revealed the presence of Zr with a corresponding drop in the amount of Ti. The initial indication based on the Ellingham diagram shows that Zr forms more stable chlorides than Ti and, thus, Zr deposition replaced Ti deposition in the system [13]. A similar approach was taken in the current research to assess the feasibility of using this in a controlled way to produce ternary

Ni-Ti-Zr SMAs. Hence, the objective of the second portion of this work was to co-deposit Ti and

Zr on the Ni microwires using pack cementation and, upon homogenization, obtain a single phase Ni(Ti,Zr) shape memory wire with an increased martensite start temperature. Initially, the source materials were chosen to be pure Ti and zirconia (ZrO2), respectively, but the oxygen liberated from zirconia reacted with titanium to form titania (TiO2), which led to a Ti deficiency in the system. Based on these preliminary results, the pack composition was modified to use pre-alloyed Ni-75 wt.% Zr powder as the source of Zr. It was observed that the co-deposition process led to a core-shell structure with enough Ti in the matrix such that, upon homogenizing, a central Kirkendall pore developed. However, the amount of Zr was not

sufficient to achieve the shape memory effect. Moreover, the composition of Zr was inconsistent among the samples processed in the same batch, with variations in microstructure, coating thickness, and final diameter obtained. The wires were very brittle and were porous on the surface following co-deposition, which made them difficult to handle for any further thermal and mechanical testing.

Given the difficulties posed with co-deposition of Ti and Zr, an alternative approach of depositing the two elements in separate steps was considered. Hence, as a first step, the pack zirconization of Ni wires was studied with the following objectives:

i. To determine the feasibility of fabricating binary Ni-Zr microwires using pack

zirconization of 50 µm diameter Ni wires

ii. To investigate phase formation and Kirkendall pore evolution during the Zr deposition

treatments

iii. To evaluate the Zr deposition kinetics (i.e. Zr gain as a function of coating time) and

determine an appropriate deposition time to achieve a final Zr content above 10 at.%.

Following the zirconization, once the desired amount of Zr is added, the wires can be homogenized and Ti can be deposited using the pack titanization procedure previously published [9] to reach the ternary Ni-Ti-Zr HTSMA composition.

2 Background

2.1 Shape Memory Effect

2.1.1 Introduction

For some alloys, “when a certain amount of force is applied, it deforms plastically, but can revert back to its original shape with the application of thermal energy. This property is termed as the Shape Memory Effect (SME), and the materials are known as Shape Memory

Alloys (SMAs)” [14]. There has been extensive research on these alloys as they have been used in various applications. For example, in bioengineering, they are being used to mend broken bones, for clogged arteries as stents as shown in Figure 1(a), and as dental braces

[15]. They are also being used in the trailing ends of helicopter blades for better precision and as actuators [1].

Figure 1. (a) NiTi self-expandable neurosurgical stent (b) percutaneous aortic valve realized with eNiTinol membrane [2]

SMAs respond to temperature changes, which eliminates the use of sensors, making the system lightweight and compact [16]. Also, these materials provide simple and frictionless systems, which reduces timely maintenance [17]. Because these materials respond to a temperature or a stress change, they are often called smart materials. This is the result of a solid-state austenite-martensite diffusionless phase transformation.

There are two main types of solid-state phase transformations:

a) Diffusional: In diffusional, long-range diffusion of atoms results in the formation of a new

phase [18]. This new phase has a different chemical composition compared to the parent

phase. Since the atomic movement is of long-range, it is time and temperature-

dependent.

b) Displacive: “This is also known as a diffusionless transformation, as the phase change

occurs without the long-range diffusion of atoms [18].” Instead, the atoms are rearranged

to a new stable structure. There is no change in the chemical composition as the atoms

are simply repositioned [19]. Hence, the type of transformation is called an athermal

transformation (i.e. independent of temperature). Reversible Phase Transformation

SMAs get this unique property due to this solid-solid reversible phase transformation.

This phase change is a martensitic transformation and, thus, requires no diffusion of atoms.

Instead, there is a directional shear of the high symmetry parent austenite phase at high temperature to a low symmetry daughter martensite phase at low-temperature when cooled

[3]. Figure 2 schematically depicts the various stages of transformation and the shape memory mechanism associated with the reversible phase transformation [20] .

Figure 2. Schematic of Shape Memory Effect: Upon cooling austenite transforms to twinned martensite, which upon applying stress deforms and forms single variant martensite. The original austenite phase is “remembered” upon heating [20].

The various steps involved are:

a. Upon cooling the austenite transforms into twinned martensite

b. Loading of twinned martensite causes a transformation to detwinned martensite, to

compensate for the stress

c. Heating the detwinned martensite transforms it back to austenite.

2.1.2 Stress-Induced Martensite (SIM)

This martensitic transformation can occur not only as a result of a temperature change, but also upon the application of stress [21]. The austenite-martensite transformation is 7

accompanied by a in crystal structure. When austenite (B2) transforms into martensite (B19’), there is no macroscopic change, but there is a change in volume (microscopic) to accommodate the change in the unit cell [22].

Now suppose, the material is in martensitic state at room temperature, it is in the twinned martensite (TM). When an external load is applied to this, it deforms elastically until a critical value of stress is approached [21]. At this stage, the TM starts detwinning to compensate for the strain produced due to the load and hence forms a Detwinned Martensite (DM). In this case, when it is unloaded and heated above the austenitic transformation temperature, it will transform into austenite and back to the TM when cooled due to the shape memory effect.

However, if the material is in the austenite phase at room temperature and a load is applied, it will deform elastically, and a thermodynamically favorable SIM will be formed at a critical value of stress. This mechanism propagates until the alloy exhibits a fully DM structure [23]. Upon unloading, the alloy will transform back to austenite, and this behavior is known as pseudoelasticity (due to its nonlinear stress-strain behavior).

Pseudoelasticity, also termed as superelasticity is an isothermal process where the austenite to martensite transformation occurs under an external load (stress-induced transformation). As the external load is removed, it reverts to the original austenite phase.

Figure 3 illustrates the superelastic behavior of an SMA [24].

Figure 3. Schematic illustration of mechanical response of pseudoelastic material [24].

• a-b represents the elastic deformation of austenite phase initially, which is stable at low

stresses.

• Point ‘b’ represents the minimum stress for the transformation to martensite.

• b-c represents the complete transformation of austenite to martensite. Typically, at 10%

strain, the transformation is complete.

• c-d represents the elastic deformation of the martensitic phase.

• d-e corresponds to the recovery of elastic strain in the martensitic phase when the load

is removed or deactivated.

• e-f exhibits transformation back to austenite.

• f-a is the recovery of elastic strain in the austenite phase.

2.2 NiTi Shape Memory Alloys

2.2.1 History and Applications

Nitinol alloys, which consist of near-equiatomic nickel and titanium and exhibit the shape memory effect, were first developed by William Buelher and Fredrick Wang during their research at the Naval Ordnance Laboratory in 1959. NiTi structures are also being widely used as stents, bone implants, mechanical actuators, cooling devices, and orthopedic applications as hard tissue replacement [1]. NiTi is chosen over other SMAs, such as AgCd, AuCd, NiAl,

CuAlNi, CuAlBe, CuSn, etc. due to its biocompatibility as it forms a passive TiO2 layer, excellent workability in the martensite phase, corrosion resistance, fatigue life, and stress hysteresis [1].

The application of porous NiTi for bone implants could provide excellent biocompatibility and an elastic modulus similar to that of bone, ranging from 2 GPa to 10 GPA, which are desired properties for osseointegration and to avoid stress shielding, respectively [15]. Making NiTi alloys porous can enhance heat exchange via a higher surface area leading to faster actuation response and low stiffness [11].

2.2.2 Fabrication Techniques

The exact composition of nickel in the Nitinol system has a drastic effect on the martensite transformation temperature. From Figure 4, NiTi typically consists of 49-51 at. % Ni.

A small change in the composition makes a significant difference in the transformation temperature and, hence, affects the shape memory and superelastic behavior, making it a challenge to produce NiTi alloys. The following section describes a few methods currently used to produce NiTi alloys.

Figure 4. Martensite start temperature as a function of Ni content (at.%) [3].

1. Vacuum Arc Remelting & Vacuum Induction Melting

Fabrication of NiTi alloys has historically been a tedious process because of the tight compositional control required and the reactivity of titanium. Two commonly used methods are the VAR and VIM techniques [12]. This is done by striking an electrical arc between the raw material and a water-cooled copper strike plate. Melting is done in a high vacuum, and the mold itself is water-cooled copper [25]. This is a commercially used process but is expensive and time-consuming because it requires a temperature > 1000 °C and a considerable amount of post processing.

Moreover, wires with a diameter of < 100 µm are challenging to manufacture with these techniques and, hence, a more economical way of fabrication needs to be developed.

2. Additive Manufacturing

i Selective Laser Melting

Additive manufacturing is defined as a “method of adding material layer by layer to form a desired 3- dimensional part” [26]. There have been some notable efforts in manufacturing

Nitinol parts using this process. One such attempt employed a laser-based additive manufacturing process, namely selective laser melting, where the metal powders are heated above the melting temperature using a laser source within an inert chamber, layer by layer to form the final part as shown in Figure 5 [7].

Figure 5. 3D printed Nitinol rectangular block using selective laser melting [7].

Some factors that make fabricating NiTi parts using this process difficult are:

i. The affinity of titanium to form oxides

ii. The volatility of nickel during the melting process

This would drastically affect the transformation temperature as both of these aspects will influence the final chemical composition.

The powder used in selective laser melting is typically prepared from a Nitinol ingot via gas atomization using an inert gas, which is later sieved to obtain particular particle size distributions [27]. Figure 6 shows four different powder samples with different particle size distributions and the resulting effects on the austentie finish temperatures [27]. It can be observed that there is some amount of contamination of oxygen, carbon, and nitrogen with the varying particle size distributions, which alters the austenite finish temperature.

Figure 6. Particle size distributions and resulting the austenite finish temperatures [27]

Further melting these powders to form the final 3D part increased the oxygen content with a fall in the nickel content. This process is auspicious and economical, but further intensive study has to be performed to strike the right composition and eliminate undesirable contaminants.

ii Particle-Based Ink Extrusion

Powder metallurgy has been a promising method to develop parts of interest with a very high amount of precision and superior properties. Particle-based ink writing is one related technique where metallic powders are mixed with a polymer binder and solvents and extruded through a nozzle to the desired shape. The extruded part solidifies nearly instantaneously as the solvents evaporate rapidly at room temperature. The as-printed green body is then pyrolyzed to remove the polymer binder and later sintered to densify the parts. The process of powder-based ink writing is summarized in the schematic in Figure 7 [28].

Figure 7. Schematic representation of the particle-based ink writing process [28]

The major factors that affect this process are:

i. Powder particle size and morphology

ii. The viscosity of the ink 14

iii. Shrinkage or volumetric change the sintering process

This methodology is currently being studied to develop Nitinol alloys.

3. Pack Cementation

Based on the phase diagram shown in Figure 8, the NiTi phase exists as a stable phase at room temperature. In 1947 Bastin and Rieck studied the formation of various intermetallic phases in Ni-Ti diffusion couples and reported the presence of the Ni(Ti), Ni3Ti, NiTi, NiTi2, and

Ti(Ni) phases, which are all present in the equilibrium phase diagram [29].

Figure 8. Binary NiTi phase diagram.

Pack cementation, a chemical vapor deposition-based diffusion coating technique, was used recently to create near-equiatomic NiTi wires below 100 µm in diameter with pores at the center. This process uses a halide salt that, at high temperature, is activated to form a metal halide gas which then diffuses to the substrate surface and deposits the metal. The history and the various steps involved in this process have been explained in section 2.3. This gas-phase alloying technique was used to deposit Ti onto Ni wires and then, after homogenizing the wire, to achieve a uniform NiTi composition. Kirkendall pores formed due to the difference in the diffusivities of Ni and Ti coalesce into a central channel, resulting in the formation of a microtube

[9-11]. For the case of hollow tubes, the surface area to volume ratio increases, thereby boosting the heating/cooling efficiency leading to faster response times of the thermal shape memory effect.

2.3 Effect of Ternary Alloying Additions to NiTi

2.3.1 History and Applications

Alloying Nitinol with a third element has been a great area of interest for researchers as there is a need for lightweight, high strength alloys that can be deployed at high temperatures.

The addition of lightweight elements like Aluminum (Al) and Magnesium (Mg) has improved the fatigue life of the alloys and reduced the weight of Nitinol for weight-sensitive applications [30].

Cobalt additions have increased the stiffness of the alloy, but have drastically decreased the shape memory response [30]. Copper (Cu) additions of around 20 at.% have notably decreased the thermal hysteresis, thus increasing the fatigue life of the SMA [31]. Despite the benefits, NiTiCu is not very commonly used as it is very difficult to machine.

Alloying with elements like Fe, Au, Pt, Pd, Hf, and Zr is known to alter the shape memory response and Transformation Temperatures (TTs) [32]. Fe is a martensite suppressant, meaning it brings down the TT. The addition of around 3 at.% of Fe will bring down the martensite start temperature to around - 40 °C [30]. Since the martensite start temperature of binary NiTi is less than 100 °C, it has limited use in high-temperature applications. Au, Pt, Pd,

Hf, and Zr are martensite enhancers as shown in Figure 9, as they increase the Ms temperature

[30]. This is mainly due to the precipitation of new phases whose presence in the microstructure can cause a shift in the TT [16].

Figure 9. Effect of ternary alloying additions on the transformation temperatures of Nitinol SMA [14].

Equiatomic TiAu exhibits a phase transformation from a B2 austenite to a B19 orthorhombic structure. Substituting Au for Ni increased the TT by 21 °C/at.%, which was reported by Eckelmeyer [16]. Similar behavior was observed with Pt, where initial additions 17

caused a drop in the TTs until it reaches 10-15 at.%. Once the B19 structure in TiPt becomes favorable, there is a linear increase in the TT with increasing Pt content and can reach up to

1000 °C. However, the high cost of these elemental additions have limited the applications of these alloys. Hence, there was significant research on less expensive alloying elements to enhance the high temperature properties of the NiTi alloys. Hf and Zr are two examples of such alloying additions considered, because of their low cost, decent ductility and workability, and excellent shape memory characteristics. The detailed effects of Zr additions are further discussed in chapter 4.

2.4 Pack Cementation

2.4.1 Introduction and History

Pack cementation, a simple chemical vapor deposition (CVD) process, is a method used to form diffusion coatings. This method is also known as the powder immersion technique and it provides the substrate material enhanced oxidation, corrosion, and wear resistance [33].

Initially, this technique was commonly used to deposit Al on Ni-based alloys, as shown in Figure

10 to form an NiAl intermetallic compound to impart hot corrosion resistance [34]. Later low alloy steels that were used in gas turbines were chromized (i.e. chromium was deposited on the surface) to provide resistance against aqueous corrosion [34].

A variety of materials, like Cr, Mo, and more recently Ti, Zr, and Hf [35], have been used to coat Ni and Co-based alloys. NiTi microwires showing shape memory properties can be fabricated using this process. With this diffusion coating, wires with a central pore can be developed because of the differences in the diffusivities of the components, as described

further in section 2.5 [9]. Compared to most other conventional techniques that require extensive and sophisticated equipment, temperatures greater than 1200 °C, and a high degree of post-processing to strengthen the parts, pack cementation is rather a simple and cost effective approach [33].

Figure 10. Aluminide coating on a Ni-base material [34].

The pack cementation process was first described by Van Aller in 1914 where Cu/Fe components, which had applications in steam power plant condensers, were embedded in a mixture of aluminum, ammonium chloride, and graphite and heated to 450 °C for 2 hours and later annealed at 700 °C [34].

2.4.2 Pack Composition and Thermodynamics

In this process, the pack is the mixture of various powders including the following three constituents:

a. a source metal or alloy powder containing the element to be deposited 19

b. a halide salt activator that reacts with the source to form a metal halide gas

c. a filler powder that occupies space such that the substrate can be buried

The pack cementation process can be divided into four different mechanistic steps which are schematically represented in Figure 11.

i. At high temperature, reaction between the activator and source, which determines the

vapor pressures of the active gaseous species in the pack

ii. The gaseous diffusion of the metal halide from the pack to the substrate surface

iii. Surface reactions of the metal halide at the substrate to deposit the coating elements

and form product vapor species

iv. Solid-state diffusion of the source elements into the substrate

Figure 11. Schematic showing the various constituents of the pack, the chemical reactions, and the gas-phase and solid-state diffusion.

The powders are weighed and mixed thoroughly to obtain a homogenous mixture with minimal agglomeration. In this process, the substrate material to be coated is embedded in the

pack within a crucible, which is then closed and placed inside a furnace under the flow of inert gas at the desired temperature. When the crucible is heated, a chemical reaction occurs where the source reacts with the activator to form metal halide vapors. A coating is formed via reduction reactions of the metal halide vapors at the substrate surface and subsequent solid- state diffusion between the source and the substrate species. This technique can be applied to nearly all geometries. For example, let us consider pack cementation of wire cross-sections as the current research is oriented towards developing microwires.

Figure 12 is a schematic representation of the aluminide coating formed on a Ni wire.

Here, aluminum (Al) is the source material, TiC the filler, and ammonium chloride (NH4Cl) the activator. At high temperature, the aluminum reacts with the ammonium chloride to form aluminum chloride and diffuses via the gas phase to the nickel surface where it is deposited.

Figure 12. Schematic of aluminum-rich coating formed on surface of a Ni wire cross-section via pack cementation.

When the source is available freely, the growth of the coating is controlled by solid-state diffusion, and the thickness is given by a parabolic law [36],

(1) 푦 = √푘 (푡 – 푡표) where y = thickness of the layer formed

k = parabolic growth rate constant

t = pack cementation time

to= delay time

The coating thickness is influenced by the composition of the pack, the coating conditions (i.e. temperature and the time) and is determined to quantify the deposition kinetics. Figure 13 shows the effect of coating time on the growth of the Ni2Al3 layer during aluminization of Ni wires.

Figure 13. Optical micrographs showing the growth of the coating during the aluminization process on nickel wires for (a) 300 s and (b) 600 s. The inner core is γ-Ni and the outer shell is δ-Ni2Al3 [37].

Temperature has a critical importance to the coating process as it is a diffusion-based approach. For instance, in the process of titanizing nickel to form equiatomic NiTi, a temperature of 1000 °C was used initially (this temperature was based on the aluminization of

nickel) [12]. Since NiTi2 is formed as the exterior layer during the process, and the melting point is 964 °C, there is the formation of liquid during coating at this temperature. This partial melting leads porosity and an uneven coating as shown in Figure 14.

Figure 14. Effect of titanization temperature (1000 °C) on porosity within and non-uniformity of the coating [12].

2.5 Solid-State Diffusion and Kirkendall Effect

2.5.1 Introduction to Diffusion

When a candle is lit, the fragrance of it drifts across the room. The same is the case with a drop of ink in water. “This movement of molecules from a region of higher concentration to a region of lower concentration is termed as diffusion” [38]. The possibility of atomic diffusion in solids was determined over 200 years ago. The process of diffusion in solids under the driving force of a concentration gradient is based on Fick’s first law with the assumption of steady state: 23

퐽 = −퐷 푑퐶/푑푥 (2) where J = diffusion flux, i.e., mass of material crossing a plane of unit area per unit time

C = concentration of diffusing element

x = the location within the host solid

D = diffusion coefficient

dC/dx = the concentration gradient driving force

Steady state conditions are not always met, however, so for non-steady state or transient scenarios, diffusion in solids and is described by Fick’s second law [18]:

훿퐶/훿푡 = −퐷 훿퐶2/훿푥2 (3)

where D is the diffusion coefficient and is given by the Arrhenius equation

퐷 = 퐷표 exp −푄/푅푇 (4)

Do = pre-exponential factor,

Q = activation energy for the diffusion process of interest,

R = universal gas constant,

T = temperature

The key takeaway is that diffusion is critically dependent on the temperature, time, and concentration. From the above equations, it is evident that an increase in any of these parameters would lead to an increase in the rate of diffusion [38].

2.5.2 The Kirkendall Effect

Kirkendall, with his student Smigelskas, designed a unique experiment in 1947. They took a brass (an alloy of Cu and Zn) block and wound it with molybdenum wires. This system was later electroplated with Cu as shown in Figure 15. It was then annealed at high temperature for several hours, and to their surprise, they noticed that the Mo wires appeared to move from the original position [34-35]. Because of faster diffusion of zinc into the copper than that of copper into the brass (i.e. the intrinsic diffusivity of Zn is ~ 2.5 times that of Cu), the Mo wires moved inward.

Figure 15. Schematic of Kirkendall experiment showing Mo wires moving inward from their original position, Kirkendall 1947 [39].

Before Kirkendall experiment, Zener had proposed a direct exchange mechanism of diffusion in which atoms exchange position with neighboring atoms as shown in Figure 16 (a), but this requires a very high activation energy, which was calculated to be around 25 eV/atom.

A ring mechanism was then proposed as depicted in Figure 16 (b), which again required high activation energies in the range of ~ 22 eV/atom. The Kirkendall experiment established that diffusion of substitutional lattice atoms involves defects that facilitate atomic jumps [40]. Thus, it was concluded that diffusion in a substitutional lattice occurs via the vacancy mechanism, as shown in Figure 16 (c).

Figure 16. Schematic of (a) direct exchange mechanism, (b) ring mechanism, and (c) vacancy mechanism of diffusion [39].

Now consider a diffusion couple by welding together hypothetical materials A and B, with inert markers placed at the weld junction as shown in Figure 17 (a). Since the diffusion fluxes of both materials are different, there will be a net flow of mass past the inert markers.

Because the flux of A is higher than that of B (i.e. |JA| > |JB|), A atoms move toward B much faster than B atoms can replace them, which results in a net flow of vacancies that [41]. The condensation of these vacancies can result in the formation of so-called Kirkendall pores near

the interface and within the side of the faster diffusing species, as shown in Figure 17 (b). The formation of pores can severely degrade the material properties, especially near the interface, but this concept can be harnessed to synthesize hollow nano and microtubes by condensing these voids into a single channel near the center of a wire [29].

(b)

Figure 17. (a) Hypothetical binary A-B diffusion couple (b) SEM image of the Kirkendall voids found at the SnPb solder/Cu pad interface after aging at 1508 ºC for three days [28,42].

3 Fabricating NiTi Microtubes via the Kirkendall Effect

3.1 Introduction and Literature

Thus far, the gas-phase deposition and solid-state diffusion processes have been discussed. The imbalance in diffusivities of different atomic species result in a net flux of vacancies toward the faster-moving component, which can lead to the formation of Kirkendall pores. It was also mentioned that near-equiatomic porous NiTi has a significant advantage because of its low stiffness, higher surface area to volume ratio, excellent biocompatibility, and corrosion resistance.

Previously, pack cementation was used to deposit Ti on Ni 50 μm diameter wires to fabricate hollow NiTi microtubes [9]. Figure 18 (a) shows a cross-section of a Ni wire titanized for 2 h at 925 °C and subsequently homogenized for 4 h under vacuum, which results in a NiTi wire with uniform composition throughout, as shown in the EDS line scan in Figure 18 (b), along with a Kirkendall pore at the center. The as-titanized wires had a core-shell structure with an outer diameter of approximately 80 μm [9].

(a) (b)

Figure 18. (a) Backscattered electron micrograph of a radial cross-section of a titanized Ni wire following homogenization for 4 h at 925 °C (b) EDS profile confirming the uniform near- equiatomic NiTi composition [9]. 28

The images from an X-ray tomography experiment shown in Figure 19 revealed the presence of Kirkendall pores throughout the entire length of the wire within the field of view, which connected to form a continuous cylindrical channel [9]. For a wire with starting diameter of 50 µm, a single and continuous central pore of ~ 20 µm diameter was formed.

Figure 19. Tomographic images showing the continuity of the Kirkendall pore at three different annealing times: (a) 2h (b) 4h (c) 8h (d) reconstructed cross-section, and (e) 3D visualization of a magnified section of the pore [9].

Based on the plot in Figure 20 [43]Figure 20, the diffusion coefficient of Ni is much higher, approximately an order of magnitude higher, compared to Ti at 925 ºC and, hence, for the core-shell structure of a titanized Ni wire, Ni moves toward the surface faster than Ti does 29

toward the center, and due to the Kirkendall effect, internal pores are formed [27,29]. Three intermetallic phase layers, NiTi2, NiTi, and Ni3Ti, in order from the surface to inner Ni core, develop during the formation of the diffusion coating. This agrees with the binary phase diagram, which confirms the existence of these three phases at 925 °C [9,38].

Figure 20. Arrhenius plot of intrinsic diffusivities of nickel and titanium [43].

An investigation was performed previously on the NiTi system for different wire sizes, namely 25 μm, 50 μm, and 100 μm, to determine the wire size effect on the Kirkendall pore formation and evolution [11]. From the 50 μm samples the optimal titanization time to reach near-equiatomic NiTi composition was found to be 2 h. For 25 μm diameter Ni wires, the optimal

time should be less by a factor of 4, i.e. 30 min, and for 100 μm, it should increase by four times, i.e. 8 h [11].

Figure 21 shows the backscattered electron micrographs for the 25 µm and 100 µm diameter Ni wires titanized for 30 min and 8h, respectively. Both the samples exhibited a core- shell structure with two main pores. The 50 µm wires exhibited only one central pore, and no second pore was observed in this wire size. It is thought that at the early stages of coating, small pores are formed, which are left behind as the interfaces start moving inward. This inward movement leads to the formation of another pore around the core [11]. The early pores formed sinter out or are filled as Ti deposition increases.

Figure 21. Backscattered electron micrographs for (a) 25 μm Ni wire titanized for 30 min (b) 100 μm Ni wire titanized for 8 h [11].

The current study focuses on the kinetics of deposition of Ti on an intermediate wire size, i.e. 75 µm diameter Ni wires, to determine the following:

i. coating thickness as a function of time

ii. phase and Kirkendall pore formation and evolution

iii. titanization time required to achieve near-equiatomic NiTi composition

3.2 Experimental Procedure

Twelve Ni wires (99.99% purity) from Alfa Aesar of 75 µm diameter were cut and embedded in a pack containing 30 wt.% Ti (Alfa Aesar, - 325 mesh, 99.5% purity), 3 wt.% NH4Cl (Alfa

Aesar), and 67 wt.% TiC (Alfa Aesar, -325 mesh, 99.9% purity). The Ni wires were sandwiched between two thin layers (3 g each) of TiC powder as the Ti particles tend to sinter the pack together at high temperature, making the removal of the wires more difficult following coating.

The crucible was then pushed inside the tube furnace preheated to 925 °C under the flow of argon to activate the pack and allow deposition for 15 minutes. The crucible was then retracted back to the cold zone and kept under the argon flow for about 15 minutes. The crucible was removed from the furnace and the wires were removed from the pack and cleaned ultrasonically with acetone. Various coating times were investigated including 30 min, 60 min,

90 min, 120 min, and 150 min. The pack is exhausted after approximately 120 min and, hence, for the 150 min coating time the wires are removed after coating for 120 min and transferred to a new crucible with a fresh pack for an additional 30 min coating step.

The samples were then encapsulated in a capillary to homogenize so as to obtain single phase, near-equiatomic NiTi. To avoid oxidation, the capillary was vacuum encapsulated using

Mg ribbon. The capillary was then homogenized for 8hrs. and 16hrs. at 925 ºC by placing it in a box furnace as shown in Figure 22 (b). The wires are then taken out, cleaned and prepared for further characterization.

Figure 22. (a) SentroTech tube furnace with inert gas inlets and outlets for pack cementation (b) Sentro Tech box furnace used for homogenization anneal of the samples.

3.2.1 Microstructure Characterization

To characterize the as-coated microstructure, several of the wire samples are cross- sectioned, mounted, and polished. Two to three cross-sections from each sample were cut off and attached to a stainless steel sheet, which was later mounted in an epoxy mold. The mounted samples were mechanically ground using 320, 400, 600 and 800 grit SiC papers followed by cloth polishing with 6 µm, 3 µm, and 1 µm diamond solution and 0.05 µm alumina suspension [46]. The metallographic preparation was carried out using Allied MetPrep4, as shown in Figure 23.

Figure 23. Allied MetPrep4 for semi-automatic polisher.

The polished samples were examined using a Keyence VKX-250X confocal microscope with a 405 nm solid-state laser, as shown in Figure 24. The samples are imaged to study the diameter of and the layers formed within the as-coated samples. These results are used to estimate the total at.% Ti gain in the samples from the deposition process. Some of the intermetallic phases were not completely visible with the confocal microscope, e.g. the NiTi2 phase in the NiTi experiments. Hence, a more powerful tool, SEM, was used for more detailed characterization in terms of microstructure and phase identification.

Figure 24. Keyence VKX-250X 3D laser confocal microscope. 34

Secondary electron (SE) and backscattered electron (BSE) images were captured using a Thermofischer APREO SEM, as shown in Figure 25 in the Advanced Materials

Characterization Center (AMCC) at the University of Cincinnati. The samples were mounted in an epoxy mold, which is non-conductive. Therefore, the mount was coated with carbon before imaging. For the current experiment, a beam voltage of 10 kV and a beam current of 0.80 nA was used at 1000x magnification, and the brightness and contrast were adjusted accordingly to capture images of the as-coated samples.

Figure 25. Thermofischer APREO SEM similar to the one used at AMCC.

For the purposes of this research, BSE mode was used primarily as it provides Z- contrast. This aids in phase identification and provides information on local composition fluctuations. The overall composition of the cross-sections and the various intermetallic layers formed were analyzed using Energy Dispersive Spectroscopy (EDS) along with the Texture

and Elemental Analytical Microscopy (TEAM) software. When the samples are bombarded with electrons, x-rays emitted from the sample are detected. The spectrum of x-ray energy versus counts is evaluated to semi-quantitatively determine the elemental composition of the sampled volume [47]. A line scan over the entire diameter was performed, along with point analysis on various regions of interest, and the intensities of each element in the microstructure of the as- coated and homogenized samples were determined.

3.3 Results and Discussion

3.3.1 Coating Thickness and Phase Evolution

As the wires are titanized, they start exhibiting a core-shell structure. The high magnification optical image shown in Figure 27 has a representative core-shell structure. A thin

Ni3Ti layer surrounds an extensive Ni core after 15 min of coating. Adjacent to this are a small

NiTi layer and a NiTi2 layer. No pores are observed during the first 60 minutes of coating. Some black spots are seen in the images, which are TiC particles from the pack, which got engulfed during coating. Representative optical micrographs of the phase evolution and the growth of each intermetallic phase during the coating process for each time step are shown in Figure 27.

As titanization proceeds, Ni diffuses outward from the core and Ti inward from the surface. The thickness of NiTi2 increases radially outward whereas the NiTi layer grows radially inward. As the coating time increases there is enough Ti to diffuse into the wire, and there is a steady decrease in the Ni and Ni3Ti phases, which get consumed by the Ti-rich phases.

The thickness of the outer titanium-rich NiTi and NiTi2 layers after 150 min of coating was approximately 26 µm and 11 µm, respectively, which was anticipated to yield the desired

near-equiatomic NiTi composition upon homogenization. The area fraction curves shown in

Figure 26 (where 6-7 samples were mounted and measured) better explain the growth and shrinkage of each of the intermetallic layers. The outer layers of NiTi and NiTi2 grow as the Ni core and the Ni3Ti layer are consumed and shrink. At 150 min the pore around the Ni-Ni3Ti region grows bigger, as observed in Figure 27 (e).

0.7 Ni Ni3Ti NiTi NiTi2 0.6

0.5

0.4

0.3 AREA FRACTION AREA 0.2

0.1

0 0 20 40 60 80 100 120 140 160 TITANIZING TIME (MIN)

Figure 26. Plot of the phase area fraction as a function of the titanization time.

Table 1 reports the diameter of the as-coated wire with respect to time. It can be inferred that the rate of increase in diameter is high till the 90 min of coating and starts to slow down as the time increases, indicating that the pack is almost exhausted. Hence, for higher coating times, the wires are transferred to a new pack.

Table 1. The diameter of the NiTi wires as Ti is added as a function of coating time.

Coating Time (min) Diameter (µm)

15 78.36

30 86.57

60 94.33

90 107.22

120 118.80

150 125.87

3.3.2 Pore Formation

During the early stages of coating (around 60 min), discrete pores are formed between the Ni core and the Ni3Ti layer. As the time of coating increases, more Ni moves out of the core producing a net influx of vacancies by the Kirkendall mechanism, which coalesce to form a large pore. The pore marked as I in Figure 27 (d) represents the initial large coalesced pore that has formed. At this time the pore starts to separate from the core and, hence, the flux of

Ni from the core to the surrounding region is inhibited, resulting in variations in phase constituents and/or compositions ahead of and behind the pore.

Figure 27. Optical micrographs of cross-sections of NiTi wires titanized in 30 wt.% Ti pack for (a) 15min, (b) 30 min, (c) 60 min, (d) 90 min, (e) 120 min, and (f) 150 min. The intermetallic layers and pores formed are indicated and labeled.

Beyond this point, Ni starts to diffuse through the Ni3Ti layer to NiTi region, giving rise to a second core-shell structure, as shown in Figure 27 (e). A second pore starts to form, marked as II, around the Ni core. As the coating time reaches 150 min, these pores start to grow as the Ni core is further consumed. Based on these results, it is evident that a dual-pore structure develops in the intermediate wire size as well.

3.3.3 Deposition Kinetics

The overall composition of the as-coated wires was estimated by measuring the average remaining diameter of the Ni core and the thickness of the shells/layers. For each coating time,

6-7 cross-sections were measured from different wire specimens and the micrographs for each section were analyzed using ImageJ. The physical properties of density and molecular weight were taken from the ASM database for the binary NiTi alloy. The number of atoms in each phase is given by:

푛 = 푁푎 ∗ (휌) ∗ 푉 / (푀푤) (5)

Where n = number of atoms in phase

ρ = density of the phase in mg/µm3

V = volume in µm3

Mw = molecular weight in mg

23 Na = Avogadro’s number (6.022*10 atoms/mole)

The results were then used to compute the approximate at. % of Ti gained for each coating time using the following expression: 40

푎푡. % 푇푖 = 푛푢푚푏푒푟 표푓 푇푖 푎푡표푚푠/ 푡표푡푎푙 푛푢푚푏푒푟 표푓 푎푡표푚푠 (6)

The plot in Figure 28 represents the Ti gain as a function of coating time, which follows a linear trend. The wires reach the desired near-equiatomic NiTi composition when coated for 150 min.

The average final composition was measured as Ni-48.5Ti with a standard deviation of +/- 0.56.

Figure 28. Kinetics of titanium gain (at. %) for NiTi wires titanized at 925 °C. The colored

bands designate the different phase fields on the binary phase diagram at 925 °C.

3.3.4 Homogenization

The 75 µm wires, which were titanized for 150 min, exhibited an average as-coated diameter of approximately 125 µm. These wires were subjected to homogenization for 8 h and

16 h at 925 °C. Representative optical micrographs are shown in Figure 29. After 8 and 16 h of homogenization both pores remain intact and appear to be stable. After sixteen hours of homogenization, there is complete consumption of the Ni core and the intermetallic shells and single phase NiTi is achieved. The homogenized samples have a near-equiatomic NiTi composition

Figure 29. Optical micrographs showing radial cross-sections of titanized 75 μm diameter Ni wires homogenized at 925 °C for (a) 8 h and (b) 16 h. Wires were titanized at 925 °C for 150 min, with a pack change after 120 min.

4 Co-deposition of Zr and Ti on Ni Microwires

4.1 Introduction and Literature

To study the effect of ternary alloying additions on the TT, a characterization tool called a Differential Scanning Calorimeter (DSC) is used. DSC is a thermal analysis instrument which measures the physical properties of a given material as a function of temperature or, in other words, it measures the energy absorbed or released as the sample is heated or cooled [48]. In a DSC analysis, the heat flow from the furnace to the sample of interest is measured relative to the heat flow to the reference pan. Both the sample pan and reference pan are similar, with the reference being empty.

Initially, both the pans are at the same temperature, but as the sample goes through a series of physical changes, heat is either added or dissipated. The process where energy is absorbed is termed as an endothermic reaction, and when energy is released is termed as an exothermic reaction. Figure 30 (a) shows the DSC curves for an equiatomic NiTi wire, which was as received, and Figure 30 (b) shows the change in transformation curve with 10 wt.% Hf added to it. In the NiTi sample, there is a sharp peak for both the transformations.

It was determined that for this binary NiTi alloy, the As temperature was around 14 °C and the

Ms occurred around 0 °C. With the introduction of Hf, the heating and cooling curves are drastically shifted. The peaks were observed to be of lower intensity with significant peak broadening. The As was recorded around 147 °C and the Ms temperature was around 162 °C

[35].

Figure 30. DSC curves for a) equiatomic NiTi and b) with 10 wt.% Hf [35].

4.2 Effect of Zirconium on NiTi

From the previous research on Ni-Ti-Zr, this HTSMA showed less favorable shape memory characteristics compared to its Ni-Ti-Hf counterpart [49]. There was also a downfall in the strain recovery, ductility, and thermal stability, which makes this system difficult. In the past, HTSMAs with (Ti+Zr) or (Ti+Hf) rich compositions were considered to be more favorable because, as explained earlier, as the Ni content increases to Ni-rich compositions there is a dramatic decrease in the TT [16]. But, a lower Ni content makes the alloys prone to Ti-rich phases in the ternary alloys like Ti2Ni and Ti4Ni2Ox precipitates, which would make the alloys brittle and difficult for drawing [50]. However, recent research has been focused more on Ni-rich compositions of the alloys where many nanoscale precipitates are formed, known as the H phase [16]. Aging these Ni-rich alloys have notably increased their stability and SME of

Ni-Ti-Hf alloys due to the precipitation of the Ni4(Ti,Hf)3 phase, which strengthens the material

[51]. 44

Significant advantages of Zr in NiTi compared to Hf are:

i. Hf costs almost ten times as much as Zr

ii. There is a large weight reduction in the Ni-Ti-Zr system as compared with Ni-Ti-Hf,

which would be beneficial for weight-sensitive applications.

Figure 31 shows the effect of zirconium additions to Nitinol on the martensite start temperature. There is a sudden fall in the TT as a small amount of Zr is added to it. Zhongjie et al. explained that the addition of Zr to the NiTi matrix converts the B2 FCC structure to B’19 monoclinic, which results in the lowering of the TT [52]. As the amount of Zr increases above

10 at.%, there is a drastic increase in the martensite start temperature as the author confirms that it promotes the formation of the B19 orthorhombic crystal structure. This change in crystal structure increases the TT, and this has been validated using XRD analysis, where the B19 phase was identified for 15 at.% Zr in the sample [53].

In Ni-Ti-Zr alloys, as the concentration of Zr increases, numerous secondary phases start to form, especially in the (Ti+Zr) rich compositions [16]. When Zr content is higher than 7 at.%, laves phases are formed and the Ni(Ti,Zr)2 intermetallic phase forms as the Zr concentration increases above 10 at.%. The presence of these Ni(Ti,Zr)2 precipitates reduces the mechanical properties, transformation temperatures, and the stability of the alloy. For a Ni- rich alloy, Ni7(Ti,Zr), Ni10(Ti, Zr)7, or NiZr form depending upon the Zr content, and as the Zr content increases the fully recoverable strain decreases [53]. Since the volume fraction of the secondary phases is considerably high, these Ni-Ti-Zr alloys are less stable compared to their

Ni-Ti-Hf counterparts [16].

350

300 C)

° 250

200

150

100 Ms Temperature ( Temperature Ms

0 0 5 10 15 20 25 30 Zr Content (at. %)

Figure 31. Transformation temperature change in NiTi as a function of zirconium content [52].

4.2.1 Effect of Filler in Pack Titanizing

The filler is the addition to the pack mixture to retain the volume of the pack and act as a binder during the coating process. Typically, they should remain inert and should not interfere with the source and the halide activator during the process, which otherwise would affect the deposition kinetics and the final composition of the samples.

Titanium carbide (TiC) has been widely used recently as the filler for the pack titanization process since it is inert and stable and does not contaminate the NiTi samples. Although TiC particles can be seen engulfed near the surface of the cross-section presented in Figure 32, they do not affect the chemistry of the NiTi wires [7]. Before the use of TiC, Al2O3 was used as the filler, as is typical of pack cementation processes in general. Figure 33 (a) shows a

micrograph of a cross-section pack titanized using Al2O3 filler for 2 h at 1000 °C under the flow of argon. Some black aggregates were observed on the surface of the as-coated samples, and the EDS scan revealed the presence of Al in it. The scan included Al, along with Ni and Ti, and shows a substantial drop in Ti at the expense of Al pick up, thus affecting the final composition

[12].

Figure 32. Optical micrograph of titanized 75 µm nickel wire coated for 90 min using TiC as the filler. Some TiC particles are engulfed by the wires as the diffusion coating grows.

The following chemical reaction occurs during titanization [12]:

2퐴푙2푂3 + 3푇푖퐶푙4 → 4퐴푙퐶푙3 + 3푇푖푂2

2퐴푙2푂3 + 6퐶푙2 + 3푇푖 → 퐴푙퐶푙3 + 3 푇푖푂2 + 2퐴푙2푂3 (6)

Later, zirconia was used as the filler instead of alumina. A similar result was observed during the experiment. The sample picked up Zr at the expense of Ti.

푍푟푂2 + 2퐶푙2 + 푇푖 → 푍푟퐶푙4 + 푇푖푂2

푍푟푂2 + 푇푖퐶푙4 → 푍푟퐶푙4 + 푇푖푂2 (7)

Figure 33. (a) Micrograph showing black aggregates after 2 hours of titanization using Al2O3 at 1000 °C (b) intensity plot showing the presence of Al [12]

4.2.2 Thermodynamics of the Reaction

The Gibbs free energy is the driving force for the aforementioned chemical reactions to occur. A negative value suggests that the thermodynamic reaction is favorable and occurs spontaneously without an additional energy input [13]. The Gibbs free energy is expressed as:

훥퐺° = 훥퐻 − 푇 훥푆 (8) where ΔH = the enthalpy, i.e. the actual energy absorbed or liberated during the reaction

ΔS = the entropy, i.e. the measure of disorder

T = temperature

An Ellingham diagram is a series of straight-line plots with a positive slope between the

Gibbs free energy, ΔG°, and temperature, with the ΔH and ΔS values being constant unless a change in the phase occurs. These plots help to determine the stability of a compound across a range of temperatures. Figure 34 shows the Ellingham plot with the Gibbs free energy change for chlorides.

At 925 °C, the ΔG° for Ti is around - 84 kcal/mole, whereas that of Zr is - 90 kcal/mole.

This indicates that the zirconium chloride formation is more favorable than titanium chloride.

Hence, the deposition of Zr in a Ni wire is possible using a halide activator. Therefore, this study proposes to fabricate ternary Ni-Ti-Zr wires using the pack cementation approach. The main objective is to examine the feasibility of simultaneous Zr and Ti deposition, also known as co-deposition, and to determine the amount of Zr and Ti in the wires. Optimizing the process is difficult as the actual kinetics of deposition are unknown. Thus, packs of different Zr concentration were studied to determine the deposition rates and the Zr content in the as- coated wires.

925 °C

Figure 34. Ellingham diagram for chlorides. The Gibbs free energy of formation of Zr and Ti at 925°C has been depicted [54].

4.3 Experimental Procedure

The main idea behind this process is to simultaneously deposit two both Ti and Zr, making the coating a one-step process, which reduces the time and cost of fabrication. The

first attempt was to further investigate the use of ZrO2 as the filler and varying the corresponding pack composition. The pack initially consisted of the following:

i. 67 wt.% ZrO2 filler/source powders (< 44 μm)

ii. 30 wt.% Ti as source powder

iii. 3 wt.% NH4Cl activator powder

A more reactive Zr-contianing powder, a pre-alloyed Ni-75 wt.% Zr was chosen as the Zr source powder along with pure Ti, to compare the deposition kinetics. The pack constituents were as follows:

i 57 wt.% inert TiC filler powders (< 44 μm)

ii 30 wt.% Ti and 10 wt.% Ni-75 wt.% Zr as source powders

iii 3 wt.% NH4Cl activator powders

These powders, all procured from Alfa Aesar, were mechanically mixed for at least 45 min. From this, 44 g of the pack was weighed and used for the process. Approximately 22 g of this mixture was poured into an alumina crucible and then all wires were immersed evenly spaced in the pack mixture. A thin layer (~ 3g) of TiC was added both on the bottom and top of the wires to prevent sintering of the pack directly surrounding the wires to aid in removal following coating. The crucible was then closed with an alumina lid and inserted into a preheated tube furnace, which is water-chilled to prevent premature activation of the pack [9].

From the binary NiTi phase diagram, the melting point of NiTi2 is at 964 C, and some liquid may have been forming above this temperature during coating in previous experiments [12]. In order to avoid this, the experiments were carried out at 925 C.

After sealing the tube furnace and flushing with Ar for ∼15 min, the crucible was inserted into the hot zone of the furnace [9]. The samples were coated initially for 60 min. The crucible was then retracted back to the cold zone and kept under the argon flow for about 15 minutes.

The crucible was then removed and the wires were taken out of the pack and cleaned ultrasonically in acetone. The same procedure was repeated for a pack composition with 20 wt.% pre-alloyed Ni-Zr powders.

4.3.1 Homogenization

To homogenize the as-coated wires, they were embedded in TiC powder within a stainless- steel foil and wrapped, which was inserted into a quartz tube and vacuum-sealed. The wires were subjected to heat treatment at 925 ºC for 24 hours in a box furnace. The tube was then water quenched, and the wires were taken out, cleaned in acetone, cross-sectioned, and mounted. The samples were polished down to 0.5 µm for microstructure characterization

4.4 Results and Discussion

4.4.1 Zr Source: ZrO2 Powder

The composition of the pack plays a significant role in the deposition kinetics and the final composition of the wires. The deposition of Ti using the ZrO2 filler leads to a core-shell structure. In the current study, the exact phase constituents and phase evolution will not be studied, as the intention is simply to check the feasibility of the process and the average concentration of Ti and Zr. After the initial 60 min of coating, as shown in Figure 35 (a), the wires picked up very little Ti. A small amount of Zr was also picked up with an average

concentration of around 1.5-2 at.%. As the time of coating was increased to 120 min, shown in

Figure 35 (b), the coating thickness grew at the expense of the Ni core.

Figure 35. Micrographs of cross-sections after simultaneous Ti and Zr deposition using 66 wt.% ZrO2 as the filler and 30 wt.% Ti powder as the source coated for (a) 60 min and (b) 120 min at 925 °C.

Therefore, the samples did not pick up enough Ti to reach 50Ni-50(Ti,Zr). After homogenizing the samples for 24 h, shown in Figure 36, it was evident that the cross-sections did not reach the single-phase Ni(Ti,Zr) composition required to impart the shape memory properties. The possible reasons for this behavior could be (a) the affinity of Ti to form oxides

(Eq. 7) and (b) thermodynamic favorability of Zr to form chlorides easier than Ti.

Figure 36. (a) and (b) show two different optical micrographs of the wires homogenized for 24 h after initially coating for 120 min using 66 wt.% ZrO2 and 30 wt.% Ti source powder (images from Dinc Erdeniz).

4.4.2 Zr Source: 10 wt.% Ni-Zr Pre-alloyed Powder

The BSE micrographs of the Ti and Zr coated Ni wires are presented in Figure 37. After the initial 60 min of coating, shown in Figure 37 (a), the wire had an average diameter of 64.35

µm. Elemental analysis was performed on these wires, and the average composition of Ti was around 17 at.%.

A longer duration coating of the wires was conducted for 120 min, which substantially increased the diameter of the wire as shown in Figure 37 (b), with the Ni core shrinking due to the outward diffusion of Ni from the core. A high magnification EDS scan was performed on the sample, as shown in Figure 38, to qualitatively study the concentration profile. The samples picked up a significant amount of Ti, with the average concentration being 45-48 at.%. The concentration of each of the layers has been tabulated in the Table 2. The samples from the

same batch showed a considerable variation in the deposition rate, and the final composition of Zr varied from 4 to 7 at.%.

a b

Figure 37. The Co-deposition of Zr and Ti at 925 ºC for (a) 60 min and (b) 120 min.

100 NiL ZrL TiK

Atomic% 40

0 0 10 20 30 40 50 Distance, µm

Figure 38. (a) The magnified BSE micrograph of a sample coated for 120 min. (b) The EDS line profile for the co-deposited wires.

Table 2. Composition from spot analysis of wires after co-deposition for 120 min

Nickel Concentration Titanium Concentration Zirconium Concentration Spot # (at. %) (at. %) (at. %) Spot 1 68.43 29.79 1.78

Spot 2 74.74 20.25 5.01

Spot 3 58.73 36.66 4.61

The as-coated wires were homogenized for 24 h. The Ni from the core moves radially outward with Ti and Zr moving towards the center. Due to the Kirkendall effect, a central pore was formed, and the cross-section had a final average Zr concentration of ~ 4-5 at.%, as shown in Figure 39 and Figure 40 respectively. The samples were highly brittle, which hampered the metallographic preparation. Moreover, the samples showed a notable variation in the coating thickness and, hence, the final composition within the same batch of wires.

Figure 39. Homogenized for 24 h after 120 min coating using 10 wt.% Ni-Zr pre-alloyed pack.

70 Ti Zr Ni 60

30 Atomic%

0 0 10 20 30 Distance µm Figure 40. The EDS line profile for co-deposited samples after homogenization for 24 h. The samples were brittle and broke during metallographic preparation.

4.4.3 Zr Source: 20 wt.% Ni-Zr Pre-alloyed Powder

To increase the amount of Zr in the samples, a higher concentration of source was chosen. A similar core-shell structure was observed, as shown in Figure 41. The deposition was much faster, with Ni core being consumed and diffusing out to the surface. Small Kirkendall pores were also observed around the Ni core. The samples had a significant variation in the diameters, ranging from 83 µm to 102 µm leading to a more considerable variation in the final concentration Moreover, all the sample wires from the same batch were highly porous, which made further processing difficult and unambiguous.

a b

Figure 41. BSE micrographs for Ni wires co-deposited with Ti and Zr ( 20 wt. % Zr source powder) for 120 min. The wires were highly porous.

Since there is an extreme variation in the coating thickness and the final concentration, it is essential to study the binary NiTi and Ni-Zr systems individually and further processing it. As

NiTi system has been thoroughly studied, it would be vital to study the behavior of Zr in Ni wires using halide activated pack.

5 Kinetics of Zr Deposition on Ni Microwires

5.1 Introduction and Literature

In order to fabricate ternary Ni-Ti-Zr shape memory microwires, the co-deposition approach previously described was proposed to reduce the time and cost of the process.

However, the samples were inconsistent with respect to the diameter, coating thickness, and the final composition. Another approach would be to first deposit Zr and then add Ti via a second pack cementation step. Therefore, it is necessary to study the deposition of Zr on Ni wires to determine the various stages of core-shell formation and the feasibility of using the pack cementation process for this purpose.

The Ni-Zr binary system proves beneficial in developing thermal barrier coatings on superalloys [55]. From the thermodynamics of the coating processes as explained in the previous chapter, formation of the chlorides of Zr are more favorable compared to those of Ti.

However, since Zr is a large atom and has a high melting point, it has a lower diffusivity in Ni.

Based on Figure 42, at 925 °C, the diffusivity of Zr in Ni is in the range of 10-12 m2/s.

Figure 42. The diffusivity of zirconium in the Ni matrix (LDA Debye line) showing the diffusion of Zr in Ni at 925 °C. 59

The current chapter focuses on the binary Ni-Zr system and fabrication of Ni-Zr wires using the pack cementation process. This would provide insight into how much zirconium can be added to the nickel wires and the various intermetallic phases formed. Since the NiTi system has been studied thoroughly, it is crucial to check if pack cementation is a feasible process for the Zr deposition. The study objectives are as follows:

i. Determine the kinetics and uniformity of zirconization of Ni wires

ii. Identify the layers formed in the as-coated condition

iii. Final Zr content in the as-coated samples

5.2 Experimental Procedure

The experimental procedure remains the same from the previous two chapters, with the exception of a change in the pack constituents and their composition. First, twelve Ni wires

(99.99% purity, from Alfa Aesar), each 4 cm long with 50 μm initial diameter, were cut. Next, the pack was prepared and consisted of the following:

i. 87 wt.% inert TiC filler powders (< 44 μm)

ii. 10 wt.% Ni-75 wt.% Zr pre-alloyed source powders (< 44 μm, 99.5% purity)

iii. 3 wt.% NH4Cl activator powders

These powders, all procured from Alfa Aesar, were mechanically mixed for at least 45 min.

From this, a total of 50 g of the pack was weighed and used for the process. Approximately 25 g of this mixture was poured into an alumina crucible and then all wires were immersed evenly spaced in the pack mixture. The remaining 25 g of the pack was added on top of the wires and

the crucible was closed with an alumina lid and inserted into a preheated tube furnace. The assembly was then moved to the hot zone at 925 °C for 30 min under the Ar environment. The crucible was then moved to the cold zone and removed from the furnace after 15 min. The wires were then removed from the crucible and cleaned ultrasonically in acetone. The coating process was conducted for different durations including 60 min, 90 min, 120 min, 150 min, and

180 min to study the rate of deposition of Zr on the Ni wires.

5.2.1 Image Analysis

The area fraction of each layer was measured using the image analysis software

ImageJ. Using SEM micrographs, which clearly show the distinct layers after each coating time, the number of atoms in a particular phase can be calculated by the following equation:

푛 = 푁푎 ∗ (휌) ∗ 푉 / (푀푤) (9) where n = number of atoms in phase

ρ = density of the phase in mg/µm3

V = volume in µm3

Mw = molecular weight in mg

23 Na = Avogadro’s number (6.022*10 atoms/mole)

Once the area was calculated using ImageJ, the physical properties of each phase were taken from the ASM database for calculation. The results were then used to compute the at.% of Zr using the following:

%푍푟 = 푛푢푚푏푒푟 표푓 푍푟 푎푡표푚푠/ 푡표푡푎푙 푛푢푚푏푒푟 표푓 푎푡표푚푠 (10)

5.3 Results and Discussion

5.3.1 Zirconizing and Phase Evolution

To understand the growth of intermetallic phases in each stage, the wires were characterized using high magnification backscattered images. The micrographs show a core- shell structure upon zirconization. Elemental mapping was performed on the backscattered image of the as-coated samples to check for the presence of Zr. Figure 43 shows the elemental map after 120 min of coating. As stated before, Zr deposits slower compared to Ti because its diffusivity is lower in Ni. Hence, only a small amount of Zr gets deposited after the initial 30 min of coating. Nickel has a higher diffusivity compared to zirconium and, hence, diffuses from the core to the shell faster than Zr does radially inward. The elemental map shows Ni throughout the cross-section with Zr in the shells. The samples showed a small amount of Ti, which was contributed by TiC as it was engulfed during the coating process. The carbon in the map is from the epoxy resin that was used to embed the samples for characterization.

Figure 43. Elemental maps of (a) nickel (b) zirconium (c) titanium (d) carbon on the zirconized samples after 2 h of coating.

A high magnification image of the cross-section with a corresponding EDS line scan, as shown in Figure 44, reports the composition of the shell layers. The shells appear to be the Zr2Ni7 and

Ni3Zr phases from the Ni core to the surface, respectively, based on the compositions according to the binary Ni-Zr phase diagram. For reference the EDS spectra corresponding to

Spot 1 and Spot 2 labeled in Figure 44 are shown in Figure 45 and Figure 46.

Figure 44. SEM micrograph of Zr coated Ni wire for 120 min and corresponding EDS line scan showing the varying composition of Zr and Ni.

From the EDS profile shown in Figure 44, the inner layer has an average composition of approximately 20.6 at.% Zr and the outer layer has an average composition of 24.9 at.% Zr, which correspond most closely to the Zr2Ni7 and Ni3Zr phases on the phase diagram. However, given that EDS is only semi-quantitative and there are other intermetallic phases shown on the phase diagram close in composition, further work needs to be conducted using XRD or TEM to confirm this phase identification based on the crystal structure and is part of the future work.

Spot 1

Figure 45. EDS spot analysis results of zirconized nickel wires for 120 min at 925 ºC.

For the remainder of this chapter regarding the calculation of Zr content as a function of deposition time it was assumed that the phases were correctly identified by EDS and should not drastically affect the Zr content calculated because it is the fact that all of the phases are so similar to each other in terms of composition that is making them difficult to distinguish.

Spot 2

Figure 46. EDS spot analysis results of zirconized nickel wires for 120 min at 925 ºC.

Figure 47. Binary Ni-Zr phase diagram showing the different phases formed at 925 °C.

Figure 48 presents the growth of each layer as a function of time. Initially, after 30 min of coating, there is only a small amount of zirconium that is deposited on Ni wires. The SEM micrograph in Figure 48 (a) shows a large Ni core surrounded by a thin layer of Zr2Ni7 (~ 3 µm wide) and a very thin outer layer of Ni3Zr (~ 1 µm wide). The final average diameter of the wires was measured at 52.30 µm. As the coating time increases, there is a significant amount of Zr and enough time for it to diffuse into the core. The outer Ni3Zr shell grows radially outwards and the Zr2Ni7 layer grows radially inwards at the expense of the Ni core. Approximately 12-13 at. % Zr was calculated after 180 min of coating, which is enough for the current study, as a

concentration of greater than 10% at. % Zr would start to shift the transformation curves to a higher temperature range in the NiTi system.

Figure 48. BSE micrographs representing the various layers formed and phase evolution in Ni wires zirconized for (a) 30 min (b) 60 min (c) 90 min (d) 120 min (e) 150 min and (f) 180 min.

The diameters of wires after each coating time and the thicknesses of each shell has been tabulated in Table 3. After 180 min of coating the samples showed an average diameter of 70.23 µm. The wires manifested a significant deviation in the coating thickness, which impacts the final composition of the wires.

Table 3. The diameter as Zr is deposited on Ni wires as a function of coating time

Coating Time (min) Diameter (µm)

30 52.30

60 55.41

90 59.01

120 62.44

150 68.72

180 70.23

The average composition with the standard deviation in each step has been plotted as a function of time in Figure 49. The plot between Zr and the coating time shows almost a linear behavior (r^2 = 0.95). Table 4 reports the average composition along with the minimum and maximum fluctuations. The samples showed an average composition of 12.3 at. % after 180 min of coating with a standard deviation of +/- 0.47. The kinetic study sheds light on the feasibility of the zirconization process and the diffusion behavior of Zr in the Ni matrix. This would be beneficial to future work on the fabrication of ternary Ni-Ti-Zr HTSMAs.

4 Zirconium Content (at .%)(at Content Zirconium

0 0 50 100 150 200 Time of Coating (min)

Figure 49. Rate of deposition of Zr on Ni wires as a function of coating time.

Table 4. Presents the coating time and the average and standard deviation at. % of Zr obtained, depicting the variation among the same batch of samples.

Coating Time Average S. D (min) at. % Zr 30 1.4 0.66

60 2.7 0.64

90 4.3 0.71

120 6.3 0.41

150 8.8 0.56

180 12.3 0.47

6 SUMMARY AND CONCLUSIONS

6.1 Fabrication of NiTi Microwires using the Kirkendall Effect

The pack cementation process was successfully implemented to synthesize NiTi microwires. The main objective of this study was to determine the kinetics of deposition of titanium on 75 µm Ni wires and the time required to reach the near-equiatomic NiTi composition. The wires were coated for approximately 150 min to reach a composition of Ni-

48.5Ti. Upon titanization for 150 min at 925 °C, a core/tri-shell structure of Ni/Ni3Ti/ NiTi/NiTi2 was observed.

The thickness of each shell was calculated, and after coating at 150 min, the wires grew to a final diameter of 125 µm. Kirkendall pores were also observed in the center of the wire due to the differences in the diffusivities of Ni and Ti. As the coating time increased from 90-120 minutes, two kinds of pores were observed:

I. Crescent-like shaped pore between the Ni3Ti and NiTi

II. Pore are formed between Ni3Ti and the Ni core, which on homogenization, forms a

continuous central pore [based on unpublished work performed by a collegue].

The wires, when homogenized for 16 h, resulted in a near-equiatomic NiTi composition, with the circular and crescent-shaped pores. Although a single phase NiTi with near equiatomic composition was observed across most of the cross-sections, there was some fluctuation in the final composition of Ti ( 48-51 at.%), due to the initial coating variations [9].

6.2 Co-deposition of Zr and Ti on Ni Microwires

The addition of Zr to the NiTi system increases the transformation temperature of the shape memory wires. A single step co-deposition process of both Zr and Ti on Ni microwires to produce ternary Ni-Ti-Zr wires reduces the time and cost of synthesizing the HTSMA.

Several conclusions were drawn from this study.

I. The use of ZrO2 as the source leads to very sluggish deposition of Zr and Ti. Also, the

samples were deficient in Ti as Ti oxide (TiO2) forms and consumes Ti from the matrix.

II. The as-coated samples using 10wt.% Ni-75 wt.% Zr pre-alloyed powders as the Zr

source picked up around 4-5 at.% Zr with the total Ti concentration varying between 44

and 46 at.%.

III. Upon homogenization for 16 h, a central Kirkendall pore was observed. The samples

had an inconsistent coating thickness, which led to considerable variation in the overall

composition within the same batch.

IV. To increase the Zr content in the wires, the composition of the Zr source was increased

to 20 wt.%. The samples showed high surface porosity and were brittle, making further

processing and testing cumbersome.

6.3 Kinetics of Zr Deposition on Ni Wires

The inconsistency in the co-deposition process led to the decision to study the deposition of Zr on Ni wires using the halide activated gas-phase deposition approach. The main objective of this work was to determine if Zr could be added to Ni microwires (50 µm) and

to assess the deposition kinetics and phase formation and evolution as a function of coating time. Several conclusions were drawn from this work:

I. The samples gained approximately 12-13 at.% Zr upon coating for 180 min. From the

literature, a Zr content above 10 at.% leads to a significant increase in the transformation

temperatures.

II. The samples grew to a final diameter of 70.23 µm after 3 h coating time. The deposition

kinetics followed a linear trend, with some variation in the coating thicknesses and,

hence, slight fluctuations in the final composition.

III. Upon the zirconization of Ni wires for 180 min at 925 °C, a core/two-shell structure of

Ni/Ni7Zr/Ni3Zr which was in agreement with the binary Ni-Zr phase diagram.

7 FUTURE WORK

Titanization of intermediate-sized (75 m diameter) wires exhibited a dual pore structure. The current study revealed the time taken to reach the near-equiatomic NiTi composition, which on further homogenization for 16 h, resulted in a single phase hollow NiTi wire. The ex-situ kinetic study will help in the further in-situ study of the mechanism of pore formation and the continuity of pores.

The addition of Zr to Ni wires using the pack cementation process was demonstrated as a viable option to add Zr toward synthesizing Ni-Ti-Zr HTSMA. However, much work still needs to be done to produce wires with the desired shape memory properties. Firstly, an XRD study needs to be performed to quantitatively evaluate the phases, determined from the EDS plot and the phase diagram. Below are a few suggestions and the possible future work to fabricate and test ternary Ni-Ti-Zr HTSMAs.

I Zirconize NiTi Wires (Two-step Deposition Process)

NiTi microtubes of near-equiatomic composition were successfully produced with a central pore (and secondary pore). Adding Zr to these homogenized microtubes using the pack composition of 10 wt.% Ni-75 wt.% Zr source powder, 3 wt.% NH4Cl, and 87 wt.% TiC for various time periods could be a possible option to obtain the desired ternary composition.

II Titanize NiZr Wires

Similar research can be conducted by titanizing the homogenized Ni-Zr wires fabricated in this work using the pack composition of 30 wt.% Ti, 3 wt.% NH4Cl, and 67 wt.% TiC to assess

the phase formation and evolution. Studying the behavior of Ti coatings on binary Ni-Zr wires through extensive characterization would serve as a foundation for the fabrication of HTSMAs.

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