Sintering Behavior of Ni/Tic Cermet Scaffolds Fabricated Via Particle-Based Ink Extrusion 3D Printing

Sintering Behavior of Ni/TiC Cermet Scaffolds Fabricated via Particle-Based Ink Extrusion 3D Printing

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

By Kameswara Pavan Kumar Ajjarapu Bachelor of Technology in Metallurgical and Materials Engineering, Visvesvaraya National Institute of Technology Nagpur, India, 2017

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

Summer 2019

ABSTRACT

Scaffolds have a variety of applications from batteries to biomedical implants due to their low density and high surface area. Cermet scaffolds in particular are attracting significant attention in the recent past, due to combining the advantageous properties of both metals and ceramics.

Cermets have traditionally been used in wear resistant, high temperature resistant, high speed tooling applications. Cermet scaffolds such as Ni-YSZ, Ni-CGO, and Cu-infiltrated-YSZ find applications in solid oxide fuel cells while cermets of calcium silicate with Ti-55Ni and Ti-6Al-

4V are used for hard tissue replacements and other biomedical applications.

Fabricating such geometrically complex structures via traditional subtractive manufacturing techniques is difficult and even using powder bed additive manufacturing methods poses problems due to poor sintering and high residual stresses. In this study, we aim to assess the feasibility of producing Ni-TiC cermet scaffolds via a particle-based liquid ink extrusion 3D printing approach. The main objective of this thesis was to determine suitable annealing conditions

(i.e. time, temperature, atmosphere) to post-process the printed cermet scaffolds and understand the sintering behavior.

Ni-TiC (50-50 vol.%) particle inks were prepared by mixing DCM, DBP, 2-Bu and PLGA along with the Ni and TiC powders in the prescribed ratio. DBP served as a plasticizer and 2-Bu served as a surfactant while PLGA was used as a co-polymeric binder to hold the powder particles together. Scaffolds with dimensions of approximately 10-15 mm in diameter, 5-15 mm in height, and a square infill pattern, were printed and subjected to various heat treatment processes. The as- printed and sintered structures were then characterized using conventional metallography techniques to investigate microstructural and compositional changes as a function of time, temperature and environment to qualitatively understand the sintering behavior.

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PREFACE

This thesis has been submitted for the Master of Science degree at the University of

Cincinnati. The exploration portrayed herein was administered in the Department of Mechanical and Materials Engineering at the University of Cincinnati between August 2017 and June 2019.

I declare that, except where otherwise stated, this thesis is the consequence of my own work. No piece of this thesis has been submitted to the University of Cincinnati or any other

University for a degree or recognition or other capability.

Kameswara Pavan Kumar Ajjarapu June 2019

ACKNOWLEDGEMENTS

I would like to express utmost gratitude and sincere appreciation to my advisor, Dr. Ashley

Paz y Puente, for giving me an opportunity to work in her lab. Her co-operation and continued support were the major driving forces during my Master’s program at the University of Cincinnati.

Despite her busy schedule, she always made sure my questions were answered. I will always be thankful to her for believing in me and encouraging me to dream bigger. She has been a role model and an inspiration to me both academically and otherwise. Not for once did I regret working for

Dr. Ashley and every second I spent in the lab would be cherished.

I would also like to acknowledge Dr. Matthew Steiner for his immense help and support with my thesis and my academic goals. I cannot thank him enough for his invaluable suggestions, guidance and assistance over the years. I would like to thank Dr. Sarah Watzman for agreeing to be on my defense committee. Witnessing her work past office hours motivated me to work harder as a graduate student. I am also grateful to Dr. Melodie Fickenscher for training me on the SEM and for spending time answering my questions and helping me with the equipment.

I would consider myself blessed to have extremely supporting and helpful lab-mates. I thank Arun Bhattacharjee, Ajith Achuthankutty, Haozhi Zhang, Safa Khodabakhsh and Aditya

Patibandla for training me on the equipment and assisting me with my research. I would also like to acknowledge the CEAS graduate studies office, for provided me with an opportunity to work with amazing people such as Julie Muenchen, Barbara Carter, Julie Steimle, Eugene Rutz and

Miranda Barker. I would certainly be remiss if I did not give credit and thank my parents, Dr. A.

Prabhakara Rao and Dr. A.S. Lakshmi, and my sister, Pragjna Ajjarapu, for their constant support and patience. A special thanks goes out to my extended family, Phani Yenugu and Madhavi

Bhamidipati for their continuous emotional support and guidance.

LIST OF FIGURES

Figure 1: Ashby diagrams for (a) Young's modulus vs. strength (b) wear-rate constant vs. hardness.

The position of WC in these diagrams illustrates why cermets with WC ceramic

reinforcements are the most sought after cermets in the hard metal industry [9,10] ..... 5

Figure 2: Improvement of cutting speed with development of cutting tool materials over time [1]

...... 6

Figure 3: Commercial materials directly processed by AM (as of 2016) [38] ...... 11

Figure 4: Images of (a) ink (b) 3D-Bioplotter (c) tube furnace, representing the three stages in

fabrication of parts using extrusion-based 3D-printing with particle-based inks ...... 14

Figure 5: Chemical structure and function of the three solvents used in ink-preparation [57] .... 15

Figure 6: Chemical structure of the co-polymeric binder (PLGA) used [57] ...... 15

Figure 7: Schematic reproduced from [28] illustrating the evolution of particle packing during ink

preparation, 3D printing, and sintering ...... 17

Figure 8: Schematic illustration of the three stages of sintering: (a) green powder compact (b)

initial stage (b) intermediate stage (c) final stage [66] ...... 19

Figure 9: Schematic illustration of the two main phenomenon that occur during sintering -

densification and coarsening [67] ...... 20

Figure 10: Schematic illustrating the contact angle between a solid-liquid interface ...... 22

Figure 11: Schematic illustrating the oxidation of TiC leading to the formation of an oxide scale

and segregation of free carbon ...... 23

Figure 12: Kinetic oxidation curves of TiC (adapted from [75]) with the anomalous behavior

circled in red ...... 24

Figure 13: Photographs of the individual components used in ink preparation ...... 26

Figure 14: Image of a 3D-Bioplotter similar to the one used in RATE Lab to fabricate scaffolds

...... 27

Figure 15: Image of 3D-printed Ni-TiC cylindrical scaffolds (15x15x10 mm each) fabricated using

particle-based ink extrusion technique ...... 27

Figure 16: Carver laboratory cold press used to compress Ni-TiC (50-50 vol%) powders into small

compacts (or) pucks...... 28

Figure 17: Image of TGA Q5000 equipment similar to the one used to determine the pyrolysis

temperature ...... 29

Figure 18: Pictures of (a) printed origami TiH2 crane [58] and (b) 3D-printed W micro-lattices cut

using a razor blade [53] ...... 30

Figure 19: Images of (a) top-view and (b) cross-section of as-printed Ni-TiC samples cut using a

razor blade and (c) cut samples mounted on an alumina plate using PLGA ...... 31

Figure 20: Image of a 3D-printed scaffold heat-treated on an alumina plate coated with BN

showing the formation of a brown layer of titanium nitride after heat-treatment ...... 31

Figure 21: Pictures of (a) tube furnace and (b) box furnace that were used to pyrolyze and heat-

treat the 3D-printed scaffolds and cold-pressed pucks ...... 33

Figure 22: Allied MetPrep4 polisher used for sample preparation ...... 34

Figure 23: Keyence VKX-250X confocal microscope ...... 34

Figure 24: Image of a Thermofischer APREO SEM similar to the one used at AMCC ...... 35

Figure 25: Confocal images of (a) as-received Ni powder (b) as-received TiC powder ...... 36

Figure 26: Confocal image of (a) Ni particles and (b) TiC particles at the same high magnification

to compare the size of a single TiC particle to the size of Ni particles ...... 37

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Figure 27: Confocal image of a printed strand of Ni-TiC particle-based ink with Ni and TiC

particles embedded in a polymer binder matrix ...... 38

Figure 28: Confocal image of Ni-TiC particle-based ink illustrating how TiC particles hinder the

flowability of ink at lower extrusion pressure and higher print speed ...... 39

Figure 29: (a) Photograph of as-printed cylinder Ni-TiC scaffold (b) Optical image of top-view of

as-printed cylinder scaffold struts oriented at 90° to each other (c) longitudnal cross

section of struts (d) lateral cross section of struts ...... 40

Figure 30: Confocal images of Ni-TiC scaffold pyrolyzed at 300°C and sintered at 1100°C for (a)

14 hours and (b) 20 hours ...... 42

Figure 31: Confocal image of Ni-TiC scaffold pyrolyzed at 300°C and sintered at 1100°C for 20

hours in the tube furnace under Ar flow followed by an additional 96 hours in a box

furnace after encapsulating the sample under vacuum ...... 43

Figure 32: Photographs of sealed quartz tubes containing Ni-TiC samples before and after

operating at 1400°C for 4 hours to show the extent of bloating ...... 44

Figure 33: TGA analysis results representing decrease in weight% with (a) increase in temperature

and (b) increase in time ...... 45

Figure 34: BSE image of Ni-TiC scaffold pyrolyzed at 300°C for 1 hour and sintered at 1100°C

for 20 hours with spots where EDS analysis was performed ...... 47

Figure 35: EDS spot analysis results from spot 1 and spot 2 shown on the BSE image in Figure 34

...... 47

Figure 36: BSE image of Ni-TiC scaffold pyrolyzed at 300°C for 1 hour and sintered at 1100°C

for 20 hours under Ar + H2 flow in the tube furnace ...... 48

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Figure 37: Elemental maps of (a) nickel (b) titanium (c) carbon and (d) oxygen of a Ni-TiC sample

pyrolyzed at 300°C for 1 hour and sintered at 1100°C for 20 hours under Ar + H2 flow

in the tube furnace ...... 49

Figure 38: Image of a Ni-TiC puck, cold-pressed at a load of 5 metric tons ...... 50

Figure 39: Confocal images of (a) a 3D-printed Ni-TiC scaffold and (b) a cold-pressed Ni-TiC

puck heat treated at 300°C for 1 hour and 1100°C for 20 hours ...... 50

Figure 40: BSE image of cold-pressed puck heat treated at 300°C for 1 hour and 1100°C for 20

hours showing the locations on the sample where EDS spot analysis was performed 51

Figure 41: EDS spot analysis results of cold-pressed puck heat treated at 300°C for 1 hour and

1100°C for 20 hours ...... 51

Figure 42: Elemental maps of (a) nickel (b) titanium (c) carbon (d) oxygen of a Ni-TiC compressed

puck pyrolyzed at 300°C for 1 hour and 1100°C for 20 hours under Ar + H2 flow in the

tube furnace ...... 52

Figure 43: List of various titanium oxides formed at different temperatures according to XRD

anaylsis conducted by Voitovich and Pugach [75] ...... 54

Figure 44: BSE image of Ni-TiC scaffold pyrolyzed at 350°C for 1 hour and pre-sintered at 650°C

for 16 hours under Ar + H2 flow in tube furnace ...... 55

Figure 45: Elemental maps of (a) nickel (b) titanium (c) carbon and (d) oxygen of a Ni-TiC sample

pyrolyzed at 350°C for 1 hour and pre-sintered at 650°C for 16 hours under Ar + H2

flow ...... 55

Figure 46: Ellingham diagram to determine the stability of oxides at different temperatures ..... 57

Figure 47: Pictures of a scaffold sintered at 1300°C under vacuum after pyrolyzing and pre-

sintering heat treatment ...... 58

Figure 48: Different magnification confocal images of the cross-section of a droplet found on the

surface of a scaffold sintered at 1300°C under vacuum after pyrolyzing and pre-sintering

...... 58

Figure 49: Phase diagram of Ni with stoichiometric TiC at equilibrium ...... 59

Figure 50: BSE image of a droplet found on the surface of a scaffold sintered at 1300°C under

vacuum after pyrolyzing and pre-sintering heat treatment ...... 60

Figure 51: Elemental maps of (a) nickel (b) titanium (c) carbon (d) oxygen of a droplet found on

the surface of a scaffold sintered at 1300°C under vacuum after pyrolyzing and pre-

sintering ...... 60

Figure 52: Confocal images of (a) as-pressed puck (b) as-printed scaffold (c) cold-pressed puck

heat treated at 350°C for 1 hour and at 650°C for 16 hours (d) 3D-printed scaffold heat

treated at 350°C for 1 hour and at 650°C for 16 hours ...... 61

Figure 53: Confocal images of (a,b,c): 3D-printed scaffold heat treated at 350°C for 1 hour, 650°C

for 16 hours, and 1250°C for 1 hour, 2 hours and 4 hours, respectively, and (d,e,f): cold

pressed puck heat treated at 350°C for 1 hour, 650°C for 16 hours, and 1250°C for 1

hour, 2 hours, and 4 hours ...... 62

Figure 54: Confocal image of (a) 3D printed scaffold and (b) Ni-TiC cold-pressed puck, heat

treated at 350°C for 1 hour, 650°C for 16 hours and 1250°C for 48 hours ...... 63

Figure 55: BSE image of 3D-printed scaffold heat treated at 350°C for 1 hour, 650°C for 16 hours

and 1250°C for 48 hours showing the locations on the sample where EDS spot analysis

was performed ...... 63

Figure 56: EDS spot analysis results of 3D-printed scaffold heat treated at 350°C for 1 hour, 650°C

for 16 hours and 1250°C for 48 hours ...... 64

Figure 57: Elemental maps of (a) nickel (b) titanium (c) carbon (d) oxygen of 3D printed scaffold

heat treated at 350°C for 1 hour, 650°C for 16 hours and 1250°C for 48 hours ...... 64

Figure 58: Plot of average relative density vs. sintering time for scaffolds and pucks sintered at

1250°C ...... 67

LIST OF TABLES

Table 1: Change in sublimation temperature of nickel with vapor pressure ...... 46

Table 2: Average relative density and standard deviation of pre-sintered and sintered Ni-TiC

scaffolds...... 66

Table 3: Average relative density and standard deviation of pre-sintered and sintered cold-pressed

Ni-TiC pucks ...... 66

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LIST OF ACRONYMS/ABBREVIATIONS

2-Bu 2-butoxy-ethoanol

AM Additive manufacturing

AMCC Advanced materials characterization center

ASTM American society for testing and materials

BSE Back-scattered electron

DBP Dibutyl phthalate

DCM Dichloromethane

DIW Direct ink writing

EDS Energy dispersive spectroscopy

HIP Hot isostatic pressing

MMC Metal matrix composite

PBF Powder bed fusion

PLGA Poly (lactic-co-glycolic acid)

SDS Shaping-debinding-sintering

SE Secondary electron

SEM Scanning electron microscopy

SLS Selective laser sintering

SPS Spark plasma sintering

TEAM Texture and elemental analytical microscopy

TGA Thermo-gravimetric analysis

YSZ Yttria stabilized zirconia

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Table of Contents

LIST OF FIGURES ...... vi

LIST OF TABLES ...... xii

LIST OF ACRONYMS/ABBREVIATIONS ...... xiii

1 INTRODUCTION ...... 1

2 BACKGROUND ...... 3

2.1 Metal Matrix Composites (MMCs) ...... 3

2.1.1 Introduction to cermets ...... 3

2.1.2 TiC-based cermets ...... 6

2.1.3 Traditional manufacturing techniques ...... 7

2.1.4 Cermet scaffolds...... 8

2.2 Additive Manufacturing ...... 9

2.2.1 Introduction to additive manufacturing ...... 9

2.2.2 Additive manufacturing of cermets ...... 11

2.2.3 Particle-based ink extrusion ...... 13

2.3 Sintering...... 18

2.3.1 Definition and driving force...... 18

xiv

2.3.2 Densification and wetting ...... 20

2.3.3 Sintering of TiC-based cermets ...... 22

3 EXPERIMENTAL DETAILS ...... 25

3.1 Ink Preparation ...... 25

3.2 Particle-Based Inks Extrusion ...... 26

3.3 Cold Pressing ...... 28

3.4 Thermo-Gravimetric Analysis ...... 29

3.5 Thermochemical Processing ...... 30

3.5.1 Sample preparation prior to heat treatment ...... 30

3.5.2 Heat treatment process ...... 32

3.6 Microstructure Characterization...... 33

3.6.1 Metallographic preparation of sample ...... 33

3.6.2 Confocal microscopy ...... 34

3.6.3 Secondary and back-scattered electron imaging ...... 35

3.6.4 EDS and elemental mapping ...... 35

4 RESULTS AND DISCUSSIONS ...... 36

4.1 Powder Characterization ...... 36

4.2 Ink Characterization and Printing ...... 38

4.3 Preliminary Heat Treatment Profile ...... 41

4.3.1 Effect of time on sintering behavior ...... 41

4.3.2 Effect of temperature and vapor pressure ...... 43

4.3.3 Effect of environment on oxidation behavior ...... 46

4.3.4 Oxidation in scaffolds vs. cold-pressed Ni-TiC puck ...... 49

4.4 Redefined Heat Treatment Profile ...... 53

4.4.1 Pre-sintering to prevent oxidation ...... 53

4.4.2 Effect of sintering temperature ...... 56

4.4.3 Effect of sintering time ...... 61

5 SUMMARY AND CONCLUSIONS ...... 68

6 FUTURE WORK ...... 70

REFERENCES ...... 74

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1 INTRODUCTION

The losses, both material and financial, resulting from wear and tear of tooling materials and other manufacturing equipment has been a huge cause of concern for over 100 years. This is more prevalent in industries such as mining, metallurgy, building material production and other manufacturing units, where the cost of replacing these wear parts can amount to about 40% of the life cycle cost of an equipment [1]. Therefore, the development of new reliable wear resistant materials has long been a top priority for these industries.

Cermets are a class of metal matrix composites (MMCs) containing ceramic reinforcements in a metal matrix. These materials combine the most favorable characteristics of ceramics and metals such as the high hardness and elastic modulus of ceramics and the ductility and toughness of the metals. TiC-based cermets, in particular, find widespread use in applications that demand wear resistance and corrosion resistance along with high strength, toughness, and low density. These are commonly used for tooling, grinding, and machining and in certain aerospace applications due to their high temperature resistance and low density.

A fundamental understanding of the correlation between processing, structure, properties and performance of a material system is the core of materials science and is commonly referred to as the materials tetrahedron. The performance of cermets mostly depends on their mechanical, tribological, and corrosion properties, which are significantly influenced by the composition and microstructure of these cermets. Historically, there has been an emphasis on the improvement of these performance characteristics to enhance cermet materials for various applications. Recently, there is also a growing demand for new cermet materials with complex geometries, especially at the micro and nano-scale. Intricately shaped structures, especially scaffolds find wide spread use in high temperature environments [2], energy absorption applications [3], and energy storage applications

such as fuel cells [4]. These intricately shaped scaffold structures are especially beneficial due to their low density and high surface area. However, fabricating intricate cermet structures at such scales with precision is a challenge, particularly via traditional subtractive manufacturing techniques, thus motivating researchers to explore various AM techniques as an alternative.

Extrusion-based techniques are thought to be the most beneficial [5] in fabricating such cermet structures and in this study, one such extrusion-based method has been employed to 3D print Ni-

TiC cermet scaffolds using particle-based liquid inks.

The objectives of this work were two-fold: (i) to 3D print stable scaffold structures of a two- powder mixture (Ni-TiC) using particle-based liquid ink extrusion technique and (ii) to post-process heat treat these as-printed scaffolds and study the sintering behavior as a function of time, temperature and environment. Therefore, an ink mixture containing Ni metal and TiC ceramic powders was prepared and 3D printed to fabricate scaffolds that were subsequently subjected to a range of heat treatments to understand the sintering behavior of these scaffolds under different conditions.

While the initial motivation and long term goal of this work was to be able to produce Ni-

TiC scaffolds for use in structural applications, it was observed that the conditions under which these scaffolds were sintered was not adequate to achieve a sufficiently high densification. This appears to be due to incomplete wetting of Ni and the inability of the Ni particles to fully infiltrate the space between TiC particles to form a continuous metal matrix. However, this study still demonstrates the feasibility of 3D printing multi-component mixtures of powders via particle-based ink extrusion and also provides important information regarding the selection of appropriate sintering parameters for Ni-TiC cermet scaffolds through discussing the compositional and microstructural effects of various annealing conditions.

2 BACKGROUND

2.1 Metal Matrix Composites (MMCs)

MMCs are comprised of a metal matrix combined with a secondary heterogeneous phase such that the properties of this composite exceed the properties of either of the individual phase constituents [6]. The properties and cost of manufacturing an MMC varies, mainly depending on the type of reinforcement (continuous or discontinuous) and the matrix material used [6].

Although various hard-metals and metallic alloys are commercially available and used for various applications, the demand for metal matrix composites with ceramic reinforcements is ever increasing [1,6,7]. This is mainly due to the fact that ceramics can impart their unique properties such as improved strength, temperature resistance, hardness and fatigue characteristics that a metallic matrix by itself cannot exhibit [8].

2.1.1 Introduction to cermets

A cermet is a specific type of MMC where the metal matrix is reinforced with a ceramic phase (typically carbides, nitrides or oxides) [8]. Such a composite of metal and ceramic exhibits a combination of metallic properties (ductility and toughness) and ceramic properties (hardness and elastic modulus) [7]. This combination of mechanical properties makes it a coveted material for high temperature use in industries such as aerospace, automotive, and high-speed tooling.

The ceramic reinforcement in cermets can be of various forms [8]:

1) long thin fibers with very high aspect ratio

2) platelets or whiskers

3) particulates with an aspect ratio ~ 1

Continuous reinforcements (like type 1 and type 2) produce excellent improvements in the properties of cermets, but the cost of manufacturing such cermets can be expensive [6]. However, cermets containing discontinuous reinforcements, like particulates, can have costs comparable to unreinforced metals with significantly better hardness and reasonably better toughness and strength

[6]. These are also easier to fabricate unlike cermets with continuous reinforcements where fiber breakup due to their anisotropic properties is a major problem [6,8]. For all of these reasons, cermets with ceramic particulates are much more sought after than those with fiber reinforcements.

The most ubiquitous examples of ceramic-metal composites and cemented carbides (fine particles of carbide cemented into a composite by a binder metal) are the WC-Co cermets. These have been in use for nearly 100 years and are still the most reliable ceramic-metal composites with a unique combination of elastic modulus, strength, hardness, toughness and wear resistance because of the properties of the WC ceramic reinforcement as shown in Figure 1 [9]. However in the recent past, research has been focused on manufacturing lighter weight cermet systems based on alternate carbides and nitrides to find an alternative to the much heavier WC-Co system [7].

Figure 1: Ashby diagrams for (a) Young's modulus vs. strength (b) wear-rate constant vs. hardness. The position of WC in these diagrams illustrates why cermets with WC ceramic reinforcements are the most sought after cermets in the hard metal industry [9,10]

2.1.2 TiC-based cermets

Throughout the history of materials used for high speed tooling applications, enormous efforts were put into finding a suitable replacement for tungsten, which was the main constituent in almost all hard metal alloys [1,6,7]. After years of research, titanium was deemed to be the best choice due to its ability to readily form hard ceramic phases of carbides, carbo-nitrides and borides

[11]. Titanium carbide based cermets first came to light in the late 1960s as shown in Figure 2 [12].

Research conducted by Keiffer et al [12] has led to the development of titanium carbo-nitride based cermets, which found potential in metal cutting applications.

Figure 2: Improvement of cutting speed with development of cutting tool materials over time [1]

Although Ti-(C,N) based cermets exhibited higher strength and toughness properties compared to TiC-based cermets, they were much more difficult to sinter and also possessed lower hardness and elastic moduli [13–15]. These properties made them unfavorable for wear applications

involving abrasion and erosion. The demand for materials for such applications has led to the development of titanium carbide based cermets [11]. The titanium carbide-nickel molybdenum system (commonly denoted as TiC-MoNi or TiC-Mo2C-Ni) in particular, proved to be extremely useful, especially due to its abrasive, erosion, and wear properties, which are comparable or, at times, even superior to WC-Co composites [16,17]. The beneficial mechanical properties of TiC such as high hardness (30 GPa), high elastic modulus (~400 GPa), high melting point (3160°C), and low density (4.95g/cc) in addition to the strong metallic bonding character that allows wetting of

TiC by liquid metals make it an attractive ceramic reinforcement phase in a cermet [8].

2.1.3 Traditional manufacturing techniques

Although TiC based cermets are known to be highly advantageous in various applications, there is still a dearth of literature on the manufacturing and processing of such materials. Cermets are commonly manufactured using traditional metallurgical techniques such as investment casting, pressure-less sintering (liquid phase sintering) and hot isostatic pressing (HIPing) [18–23]. The GE

710 program and ANL program for nuclear thermal propulsion (NTP) have adapted these techniques to manufacture cermets in the past [24]. These techniques require extremely high processing temperatures and long sintering times, which can lead to several problems including [25]:

i. Dissociation of gases at high temperatures creating splinters and unwanted pores.

ii. Degradation of mechanical properties due to processing at high temperatures for

long durations.

iii. Decrease in interfacial area of metal and ceramic, causing slower heat conduction.

In 2011, the Center for Space Nuclear Research (CSNR) at Idaho National Laboratory worked on processing cermets via Spark Plasma Sintering (SPS) [25]. This method seemed to be promising for the fabrication of fuel cermets at relatively low temperatures and shorter sintering

times compared to traditional sintering techniques [25]. However, this method still employed high pressures and high operating temperatures, although lower than traditional sintering techniques.

Moreover, the heating and cooling rates, as high as 200C per minute, employed in this technique induce residual stresses in the material which can be detrimental to the mechanical properties of the part produced. Although SPS produced highly dense cermet structures with reliable mechanical properties, there is still a need to develop a better, more convenient and cost-effective processing technique to fabricate cermets.

Later in 2016, Filippov et al conducted research on selective laser sintering of a cermet mixture of Ti and B4C [26]. However, the ceramic phases in the cermets have a high melting point and low thermal conductivity, which causes irregular thermal stresses [27]. Such an interaction of the laser with a metal-ceramic powder mixture leads to variations in local heat distribution and very high heating and cooling rates causing a build-up of residual stresses, which are detrimental to the mechanical properties of the structures and makes them highly susceptible to cracking. Therefore, if cermet structures are to be fabricated via AM techniques, it is necessary to isothermally heat and cool these structures to prevent residual stress development.

2.1.4 Cermet scaffolds

Traditional manufacturing methods to fabricate cermets are limited in various ways such as choice of material system, control over sintering parameters, and freedom to fabricate different shapes and structures. Scaffolds in particular, have gained enormous importance over the past decade due to their advantageous properties such as high specific stiffness, strength, damping, energy absorption and surface area [28–30]. 3D scaffold structures with high porosity and superior mechanical strength are widely used in biomedical applications due to their energy absorption properties [3]. Scaffolds also find applications in thermal interface materials which require the

structures to have high thermal conductivity and a low coefficient of thermal expansion for efficient heat dissipation [31]. Iron and nickel scaffolds are used in energy storage, emission control, catalyst support, and structural applications [29,30,32–34].

Cermet scaffolds such as nickel-yttria stabilized zirconia (Ni-YSZ), nickel-gadolinia doped ceria (Ni-CGO) and Cu-infiltrated-YSZ [35,36] were extensively studied for application in solid oxide fuel cells. There was also extensive research regarding new processing techniques to fabricate cermets of calcium silicate with Ti-55Ni and Ti-6Al-4V for hard tissue replacements and other biomedical applications [37]. Fabricating such cermet scaffolds containing a combination of materials at different size scales via traditional manufacturing techniques has always been a challenge.

Machining of cermets is also an extremely difficult process due to their high strength, hardness, and toughness, which are the same properties that make them a suitable material for cutting tools. Therefore, fabricating cermet parts with intricate shapes such as scaffolds using subtractive manufacturing is not feasible. In search of an alternative, the idea of processing cermets and cermet scaffolds using additive manufacturing technologies came about.

2.2 Additive Manufacturing

2.2.1 Introduction to additive manufacturing

Additive manufacturing (AM) is defined as the “process of joining materials to make parts from 3D model data, usually layer upon layer, in contrast to subtractive manufacturing and formative manufacturing methodologies” [38]. This approach, which initially was simply a technique to fabricate prototypes, has now evolved into an efficient and reliable process for manufacturing final products [39]. The “ASTM F42 – Additive Manufacturing” standards classify

AM into 7 broad categories or modalities [40,41]:

1. Binder jetting – Deposits binder in the form of droplets onto a powder bed to bind

the particles together. It requires post-processing to remove the binder and densify

the constituent powder [38].

2. Direct energy deposition – Uses thermal energy (in the form of laser or electron

beam) to melt and bind materials fed in the form of powder or a wire [38].

3. Material extrusion – Most popular AM technology due to its low cost of hardware.

Feedstock is extruded out of a nozzle and is deposited at room temperature or at

elevated temperature depending on the material that is fed [38].

4. Material jetting – Selective deposition of droplets of material onto a build platform.

Generally uses photosensitive thermoset polymers which cure upon deposition [38].

5. Powder bed fusion – Exposure of a powder bed to localized heating using a laser or

an electron beam to fuse/sinter particles together. Use of material powders as

feedstock makes it popular in manufacturing parts for service applications [38].

6. Sheet lamination – Multiple layers of sheet are bonded or adhered to each other

creating a stack and then cut to shape to fabricate the required part [38].

7. VAT photopolymerization – Liquid thermoset polymer resins are

photolithographically crosslinked to form a solid. It is one of the oldest commercial

AM processes in use [38].

Although there are various categories of AM processes available, there are not as many materials that can be processed via direct AM techniques (see Figure 3). Proper selection of material and feedstock type play a major role in fabricating structures using AM processes.

Figure 3: Commercial materials directly processed by AM (as of 2016) [38]

2.2.2 Additive manufacturing of cermets

Poor weld-ability, unreliable sintering characteristics, irregular melt pools, cracking on solidification, and a number of other such issues limit the process-ability of certain material systems by AM [38]. Ceramics and cermets are examples of such systems. Due to their combination of high melting point and low toughness, it is extremely difficult to process these systems via direct AM techniques [38,42]. Attempts were made to process alumina using direct energy deposition [43,44] and powder bed fusion techniques [45,46], but full densification could not be achieved. Most of the direct AM techniques used to process ceramics and cermets resulted in thermally induced cracking

[43–46].

Indirect AM processing techniques, however, make use of a binder to hold particles together while fabricating a part. This binder is usually transient in nature and can be converted or removed in a post-processing step [8]. A ceramic part produced using such indirect AM methods can be, after binder burnout, post-infiltrated to create near-fully dense ceramic composites. More recently, research was focused on developing indirect AM techniques that can process not just ceramic particles, but also a mixture of different types of materials to fabricate composite parts. Kitzmantel

Michael et al [5] studied the potential of extrusion based 3D printing of hardmetals and cermets in detail. Binder jet 3D printing was used extensively to fabricate various cermet materials including

TiC/steel composites [47], Inconel 625/transition metal carbide cermets [48] and Ni(x)Al(y)-TiC cermets [8]. For more than a decade, feedstocks containing metallic and ceramic powders were being developed and studied extensively which aided the fabrication of these materials via AM enormously. However, cermets and hard metals are a mixture of different powders with a range of particle sizes. This makes it difficult to fabricate cermets via the most common powder-bed AM techniques.

Fabrication of cermet structures using extrusion-based AM technology is highly beneficial compared to other AM routes mainly for the following reasons:

1. No loss of raw materials – This is especially beneficial in view of the expensive

cermet powders used compared to other metallic powders.

2. Controllable mixture of powders – Cermet powders often contain a mixture of

particles such as hard-phase carbides (TiC, WC etc.) and matrix powders (Ni, Co, Fe

etc.). These can separate if left unbound on a powder bed while sintering.

3. Isothermal heating and cooling – The SDS method (shaping-debinding-sintering)

that is associated with this technique prevents variations in local heating and high

cooling. This is a major issue with other AM techniques using laser or electron

beams, which in turn causes the buildup of residual stresses.

4. Multi-material deposition – Using extrusion-based techniques, multiple types of

materials can be printed with ease.

5. Low energy process – Extrusion based techniques are comparatively inexpensive and

do not require high cost and energy consumption unlike SLS, SLM and DED.

For these reasons, extrusion-based AM techniques are considered to be the most convenient and cost-effective solution for manufacturing cermet structures.

2.2.3 Particle-based ink extrusion

Like any other manufacturing process, AM technologies also largely rely on the type of feedstock, which is required to be in a state compatible with the process used. The major types of feedstock used by most of the AM technologies are filaments, sheets, resins and powders. Powders in particular, are the most widely available and commonly used feedstock for commercial applications. Particle-based ink extrusion makes use of these commonly available powders to fabricate material structures with high precision and improved mechanical properties.

This technique is a form of material extrusion process where inks, comprised of metallic, ceramic and/or polymeric powders dispersed in a solvent with a polymeric binder, are extruded out of a nozzle and solidified almost instantaneously through solvent evaporation and binder precipitation. This approach of printing a green body using AM techniques and sintering it via conventional metallurgical processes has been demonstrated to be beneficial in the fabrication of intricate cellular structures [28,49–56].

The fabrication of 3D structures using this technology involves a three-step process as shown in

Figure 4:

1. Preparation of particle-based liquid inks

2. Printing of inks by an extrusion-based technique

3. Thermochemical processing and sintering of as-printed structures

Figure 4: Images of (a) ink (b) 3D-Bioplotter (c) tube furnace, representing the three stages in fabrication of parts using extrusion-based 3D-printing with particle-based inks

The particle-based inks used for this process contain three main components: (i) a tri-solvent mixture of DCM, DBP and 2-Bu (ii) a co-polymeric binder (most commonly used polymer is PLGA)

(iii) powder (can be metal, polymer, or ceramic). The tri-solvent mixture assists in dispersing the powder particles and suspending them in the solution. Each of the three solvents serves a particular purpose. DBP acts as a plasticizer that improves the flowability of the ink. 2-Bu is a surfactant that forms a thin layer around the powders, thus preventing agglomeration of the particles [57]. DCM, on the other hand, is highly volatile (400 mm Hg at room temperature) and functions as an evaporant and also dissolves the PLGA that is used to bind the powder particles together [57]. The amounts of

DCM, DBP and 2-Bu that are required to make suitable inks were studied by Ahn, Dunand, Lewis et al [58]. For the purposes of this research, a 15:2:1 mass ratio mixture of DCM, 2-Bu and DBP

(Figure 5), respectively, was used to make the tri-solvent.

Figure 5: Chemical structure and function of the three solvents used in ink-preparation [57]

Polylactic-co-glycolic acid (PLGA, 85:15 PLA-PLG by mass) has been used as the polymeric binder in previous work on 3D printing particle-based inks using a direct ink writing

(DIW) or extrusion technique [28,49,58]. This is a non-polar, biocompatible, and biodegradable polymer (Figure 6) that dissolves relatively easily in DCM and does not react with any of the other materials studied in this research.

Figure 6: Chemical structure of the co-polymeric binder (PLGA) used [57]

Powder components can be varied depending on the intended target composition to print

[57]. Although particle morphology affects the printability of inks, this process can still be used to print a variety of different materials irrespective of the particle chemistry [57]. However, it is very important to understand that this technique cannot be used to print materials whose powders are soluble in the solvents.

Once all of the components are added in the prescribed ratio to make the ink solution, the particles are only approximately 5% of the total suspension volume [28]. The solution is then vortexed and allowed to thicken before it is used for printing. As the solution thickens, the volume fraction of particles increases to about 15%-20%, when the ink reaches the desired viscosity of 30-

35 Pa.s and is ready to be printed [28,49].

During printing, once the ink is extruded out of the nozzle, the particle volume fraction goes up to ~ 60% [28]. This change is a result of the shear forces acting on the extruded struts causing mechanical compaction, leading to volume contraction and an increase in surface area [28]. This increased surface area increases the rate of evaporation of DCM and, hence, DCM evaporates almost instantaneously when the strands are printed, leading to precipitation of PLGA. Finally, as the polymer pyrolyzes and particles sinter, the volume fraction increases much further. Figure 7 shows a schematic that precisely portrays this sequence.

Figure 7: Schematic reproduced from [28] illustrating the evolution of particle packing during ink preparation, 3D printing, and sintering

These as-printed strands form an intermediate structure that needs to be thermo-chemically processed: pyrolyzed, reduced (if oxide powders are used and the metal is desired) and sintered. The first step of heat treating these structures involves pyrolysis or burning-off the polymeric binder.

The thermal and physical properties of PLGA (in various forms) has been studied extensively because of it biomedical applications [59–63]. TGA analysis of the PLGA used for this research provided information about the temperature at which pyrolysis should be carried out. Pyrolysis is followed by a reduction and sintering heat treatment. Thermochemical reduction and sintering temperatures are to be specifically tailored depending on the melting temperature, self-diffusion coefficient, thermodynamics and kinetics of oxide reduction, and the presence of additional alloying elements [49].

Thermochemical treatment of the as-printed structures causes noticeable volumetric and mass changes. However, the original architecture, structural integrity, and surface features are maintained in the sintered part. It is of importance to note that the samples do not warp or crack,

even after experiencing such drastic reduction in volume. The exact mechanism of this process is not yet clear. It is believed that the binder burnout, reduction, and sintering occur simultaneously rather than sequentially [49]. Moreover, the residual carbonaceous matrix from the pyrolyzed co- polymer might also be playing a role in acting as a temporary support for the overall structure [49].

Another theory suggests that “particle jamming” might be a possible mechanism by which this phenomenon occurs [49,64,65]. Irrespective of what the mechanism for shrinking without warping is, this reduction in volume and shrinkage in size can be used as an advantage to print structures with feature sizes smaller than achievable in the as-printed state.

2.3 Sintering

2.3.1 Definition and driving force

Sintering is defined as a thermal treatment of a powder or powder compact performed at an elevated temperature, usually below the melting temperature [66]. The basic phenomenon occurring during this process is densification and grain growth and the ultimate goal of sintering is to increase powder compact strength [66,67]. This process aims at fabricating reproducible and designed microstructures (viz. controlled grain size, sinter density, distribution of phases and pores) through control of sintering variables [67]. The final goal of microstructural control in most cases, is to achieve highest density with a fine grain structure [67].

Sintering processes generally occur in three different stages (Figure 8) [66]:

1. Initial stage: This stage is characterized by necking, grain boundary formation and

surface smoothening. Inter-particle contact area increases from 0 to 0.2 and the

density increases to 60-65%.

2. Intermediate stage: In this stage, isolated pore structures are created along with

considerable grain growth and densification. Pore channels are formed along three

grain edges and the density increases from 65% to 90%.

3. Final stage: This stage is characterized by pore shrinkage, pore closure and more

grain growth, thus increasing the density of the structure.

Figure 8: Schematic illustration of the three stages of sintering: (a) green powder compact (b) initial stage (b) intermediate stage (c) final stage [66]

The fundamental driving force for sintering is the reduction of the total internal interfacial energy. The total interfacial energy of a powder compact can be expressed as ɣA, where the specific interfacial energy is represented by ɣ and the total interfacial are of the compact is represented by

A. The reduction of total energy can be expressed by the following equation:

Δ(ɣA) = ΔɣA + ɣΔA

In the above-mentioned equation, the change in interfacial energy is due to densification and the change in interfacial area arises due to grain coarsening. This can be represented schematically as shown in Figure 9.

Figure 9: Schematic illustration of the two main phenomenon that occur during sintering - densification and coarsening [67]

2.3.2 Densification and wetting

Once the particle-based inks are 3D printed to form scaffolds, these as-printed structures must undergo thermo-chemical processing and be densified to attain the desired mechanical properties. Sintering is a crucial post-printing operation that assists in densification and reduces porosity in these printed scaffolds. However, there exists a trade-off between the sintering temperature, operating pressure, and oxidation of particles at such high temperatures, which will be

discussed in detail in the next section. It is highly unusual for ceramic particles, particularly a refractory carbide like TiC, to be densified using solid state sintering techniques [8]. These materials sinter at extremely high temperatures due to their high melting points. Research conducted by Zhang et al suggests that even under pressures of 30 MPa, sintering of TiC requires temperatures up to

1700℃ [8,68].

Densification of such high temperature ceramic materials can be achieved by binding the ceramic particles together using a metallic binder. This can be achieved by either solid state sintering or liquid state sintering of the metallic binder [8]. Solid state sintering occurs by necking and bulk mass transport. On the contrary, for liquid phase sintering to occur, there needs to be partial solubility of the solid phase in the liquid phase.

Ideally, it is expected that the metallic binder infiltrates the space between the ceramic particles and fills in the pores, binding the ceramic particles together to attain maximum densification. In processing techniques where no external pressure is applied, this infiltration of the metal matrix between ceramic particles is driven purely by capillary forces [8]. This is possible when the metallic phase is able to wet the solid ceramic phase present in the sample. This wetting behavior of one phase on another is typically described by a solid-liquid contact angle of less than

90 [8]. Contact angle is defined as the angle that the edge of the liquid phase makes with the solid surface that it is wetting (Figure 10). However, factors such as surface roughness of particles, pore shape, and volume fraction of pores can alter the contact angle such that a wetting to non-wetting transition occurs at angles greater than or less than 90° [8,69].

Figure 10: Schematic illustrating the contact angle between a solid-liquid interface

In metal-ceramic composites, great wetting can be achieved either by the formation of a strong chemical bond between the metal and the ceramic due to the formation of a reaction product at the interface or due to the dissolution of a part of the solid phase in liquid metal [8,69]. Transition metals with unfilled d-orbitals (viz. Fe, Ni, Co, Cr, Zr, Ti etc.) are known to wet transition metal carbides like TiC by dissolving the carbide and forming a solid solution [8,70]. The myriad of beneficial properties of TiC as a reinforcement, the extraordinary wetting behavior of Ni on TiC, and the numerous advantages and applications of Ni-TiC cermets motivated this investigation of the sintering behavior of Ni-TiC cermets 3D-printed via extrusion-based technique.

2.3.3 Sintering of TiC-based cermets

TiC-based cermets can be manufactured to attain transverse rupture strengths of up to

250,000 psi [71]. However, the properties of a cermet mainly depend on its microstructure, which in turn is largely dependent on the processing technique and the sintering parameters used to manufacture the cermet [71]. Sintering TiC-based cermets involves heat treating these samples at temperatures as high as 1400C. However, at such high temperatures, TiC tends to oxidize in the presence of atmospheric oxygen.

This phenomenon of oxidation of TiC at higher temperatures was extensively studied by investigators starting as early as the 1950’s [72,73]. Based on experimental findings [74], it was understood that diffusion in the scale occurs chiefly through its lattice defects. Defects in this structure include divalent oxygen vacancies and trivalent and tetravalent interstitial titanium ions, with oxygen vacancies dominating at low temperatures and high pressures, whereas interstitial ions dominate at high temperatures and low pressures as shown in Figure 11 [74,75]. Analysis of the oxide scale by Reichle and Nickl [76] formed during oxidation of a TiC single crystal suggests that scale is formed at the TiC/TiO2 interface due to oxygen diffusion at lower temperatures, but the significance of diffusion of titanium cations at the TiO2/gas interface increases substantially at higher temperatures [73].

Figure 11: Schematic illustrating the oxidation of TiC leading to the formation of an oxide scale and segregation of free carbon

Studies show that the kinetic oxidation curves of TiC specimens in the temperature range of

8000C-12000C were logarithmic at low temperatures and parabolic at high temperatures [75].

However, one characteristic feature of these oxidation curves is the anomalous position of its 9000C and 10000C oxidation isotherms (Figure 12). Data yielded from gasometric studies [75] attributes

this inhibition of the oxidation process with rise in temperature to sintering of the scale in the presence of carbon and not to the evolution of gaseous products such as CO and CO2.

Figure 12: Kinetic oxidation curves of TiC (adapted from [75]) with the anomalous behavior circled in red

It is important to understand that the presence of free carbon at the TiC/TiO2 interface and its diffusion along the grain boundaries and through the oxide lattice may give rise to certain phenomena that are not commonly observed during oxidation of pure metals, which include sintering of the scale and stabilization of oxide phases (such as anatase) that are not normally stable in this temperature range [75].

3 EXPERIMENTAL DETAILS

3.1 Ink Preparation

The main idea behind preparation of the ink is to suspend metal and/or ceramic particles homogenously in a viscous solution, held together by a polymeric binder. The inks that were used for these experiments were comprised of three components (Figure 13) [49]:

i. Ni and TiC powders mixed in a 50:50 volume ratio

ii. Polylactic-co-glycolic acid (PLGA) 82:18

iii. Graded tri-solvent mixture of dichloromethane (DCM), dibutyl phthalate (DBP)

and 2-butoxyethanol (2-Bu) [49] [77]

The graded tri-solvent mixture was prepared by mixing DCM, DBP and 2-Bu in the ratio of

15:2:1 by mass. This ratio of ink components was deemed to be useful based on previous research carried out by Taylor et al [28]. For every cm3 of Ni-TiC powder, 12.7g of tri-solvent mixture was used. This tri-solvent mixture was then separated into two equal portions. PLGA was added to the first half of the solvent mixture in a Nordson EFD 30 cc ink cartridge and vortexed until all the

PLGA was completely dissolved. A powder to PLGA ratio of 7:3 by volume was maintained as recommended by Taylor et al [28]. Ni powders (99.9%, nominal 3-7 µm, Alfa Aesar) and TiC powders (99.9%, nominal - 325 mesh, Alfa Aesar) mixed in the specified ratio were added to the second half of this tri-solvent mixture in a separate 50 cc Nordson EFS tube.

After all the PLGA dissolved, the powder suspension was poured into the ink cartridge. This solution is then thickened by prolonged vortexing due to evaporation of DCM. The aim is to achieve a viscosity of 30-35 Pa.s, which is comparable to the viscosity of Hershey’s chocolate syrup for reference. Once the ink achieves the required viscosity, it is ready to be printed and can be extruded out of a nozzle at optimum pressure and speed.

It was observed that the metal and/or ceramic powder particles tend to settle and agglomerate at the bottom of the cartridge over a period of time. This can be avoided by hand stirring and/or vigorously shaking the cartridge intermittently. Moreover, occasionally sonicating the ink at a temperature of about 40°C would help increase the viscosity of the ink since the rate of evaporation of DCM is higher at this temperature.

Figure 13: Photographs of the individual components used in ink preparation

3.2 Particle-Based Inks Extrusion

All 3D-printed structures were fabricated using a 3D-BioPlotter from EnvisionTEC as shown in Figure 14. The ink that was prepared, as described in the previous section, was extruded at pressures ranging from 0.5 bar to 4.5 bar using a nozzle with a cylindrical metallic tip of diameter ranging from 200 µm to 400 µm. Cylindrical samples 10-15 mm in diameter and 5-15 mm in height with a strut spacing of 1-2 mm were 3D printed at a speed of 10 mm s-1 and a pressure of 2 bar. The values for extrusion pressure and print speed were determined by extensive experimentation with this particular particle-based ink composition via a trial and error method to obtain a quality print.

Each subsequent layer was rotated by 90⁰ from the previous layer as shown in Figure 15 for reference. These cellular structures were designed using EnvisionTEC custom software.

Figure 14: Image of a 3D-Bioplotter similar to the one used in RATE Lab to fabricate scaffolds

Provided that the inks had an appropriate viscosity, where tip clogging occurs if the viscosity is too high and spreading or sagging occurs if the viscosity is too low, the samples solidify upon printing and maintain their structure. However, a hold time of 5 seconds between layers was used ensure enough time to allow DCM to evaporate, making the strands stable. These printing parameters must be adjusted based on the material system and the structures that are being printed to achieve.

Figure 15: Image of 3D-printed Ni-TiC cylindrical scaffolds (15x15x10 mm each) fabricated using particle-based ink extrusion technique

3.3 Cold Pressing

To determine how Ni-TiC powder compacts sinter compared to the printed scaffolds of Ni-

TiC containing the PLGA binder, the same Ni and TiC powders used in the inks were mixed in the specified 50:50 volume ratio and were compacted at 150 MPa by applying a load of 5 metric tons using a Carver laboratory press (Figure 16). Small cylindrical compacts containing Ni and TiC were produced, with a diameter of 10 mm and height ranging from 1-3 mm. These cold-pressed compacts of Ni-TiC were subjected to the same heat-treatment processes as the 3D-printed scaffolds for comparison purposes.

Figure 16: Carver laboratory cold press used to compress Ni-TiC (50-50 vol%) powders into small compacts (or) pucks

3.4 Thermo-Gravimetric Analysis

As-printed samples that are fabricated by extrusion of Ni-TiC ink contain not just Ni and

TiC powders. They also have the polymeric binder (PLGA) and solvents (DCM, DBP and 2-Bu).

These as-printed samples are then heat treated to pyrolyze the polymeric binder and solvents before transitioning to a higher temperature heat treatment to sintering the metallic and/or ceramic powders.

Thermo-gravimetric analysis (TGA) provides information about the temperature at which these solvents and polymer burn off. TGA was conducted on the pure PLGA (82:18), the tri-solvent mixture, and the PLGA dissolved in the tri-solvent mixture to confirm whether the pyrolysis temperature of 300C reported in literature is sufficient to burn off all of these components of the ink. Each of these samples was heated at a constant rate of 5C per minute on a TGA Q5000 (Figure

17) to measure how the weight changes as a function of temperature.

Figure 17: Image of TGA Q5000 equipment similar to the one used to determine the pyrolysis temperature

3.5 Thermochemical Processing

3.5.1 Sample preparation prior to heat treatment

Ni-TiC as-printed samples that were fabricated using the DIW technique have a unique advantage. They can be folded, stretched, cut or otherwise manipulated to modify the geometry from the as-printed condition. Ahn et al made intricate origami structures (Figure 18 (a)) by folding printed sheets of TiH2 ink [58]. Figure 18 (b) illustrates how Calvo et al cut W micro-lattices that were 3D printed using DIW[53].

Figure 18: Pictures of (a) printed origami TiH2 crane [58] and (b) 3D-printed W micro-lattices cut using a razor blade [53]

The Ni-TiC scaffolds that were printed for the purpose of experimentation in this work were cut in a similar fashion. Each scaffold was cut into approximately equal halves (Figure 19) and mounted on an alumina plate with a small amount of PLGA dissolved in DCM underneath to act as a temporary glue to hold the samples in place while inserting them into the tube furnace for pyrolysis and sintering heat treatments.

Figure 19: Images of (a) top-view and (b) cross-section of as-printed Ni-TiC samples cut using a razor blade and (c) cut samples mounted on an alumina plate using PLGA

Literature states that coating the alumina plate with boron nitride would help in reducing the friction expected due to shrinkage of the sample [53]. However, in this case, boron nitride reacted with the TiC particles to form titanium nitride as shown in Figure 20. As mentioned, these following samples used in the remainder of the work were mounted on the alumina plate coated with a layer of PLGA dissolved in DCM. Since PLGA and DCM evaporate during the pyrolysis step, these should not affect the sintering behavior of the samples.

Figure 20: Image of a 3D-printed scaffold heat-treated on an alumina plate coated with BN showing the formation of a brown layer of titanium nitride after heat-treatment

3.5.2 Heat treatment process

Initially, the as-printed samples were subjected to a two-step heat treatment profile. Such a heat treatment process is necessary to ensure the samples contained only metallic and ceramic powders before being subjected to sintering as opposed to the as-printed samples containing the polymeric binder and solvents along with the Ni and TiC. The solvents used, DCM, 2-Bu, and DBP, have boiling points of 40°C, 171°C and 340°C, respectively.

The heat treatment cycle begins with pyrolyzing the as-printed samples to burn off the polymer and the solvents. Prior research carried out on DIW scaffolds of various metals and alloys employed pyrolysis temperatures ranging from 300°C to 450°C [28,49,51–56,58]. Initially, the as- printed samples were pyrolyzed at 300°C. However, TGA was conducted on PLGA and PLGA dissolved in the tri-solvent to confirm if this pyrolysis temperature was sufficient. Based on the TGA results, the pyrolysis temperature was modified such that all of the as-printed samples were heated at a constant rate of 5°C per minute to 350°C and held at that temperature for 60 minutes. Once all the solvents and polymeric binder were burned off, the samples were subjected to a second stage higher temperature sintering heat treatment. Various sintering temperatures were implemented, taking into consideration the oxidation and sintering behavior of TiC and Ni. The entire two-stage heat treatment was carried out in a tube furnace (Sentro Tech Corp. 1200C-3.5-12 High Temperature

Tube Furnace, Figure 21 (a)) in an Ar + 2% H2 atmosphere.

After analyzing preliminary results based on the aforementioned heat treatments, it was determined that the samples were oxidizing, despite using the Ar + 2% H2 atmosphere. Therefore, some samples were pyrolyzed and pre-sintered in the tube furnace under Ar + 2% H2 and then encapsulated in quartz capsules under vacuum with a small piece of Zr foil, which acts as an oxygen getter. These encapsulated samples were then heated to higher temperatures in a box furnace

(SentroTech Corp. 1600C-666 High Temperature Box Furnace, Figure 21(b)) and quenched in water after various sintering times. The exact temperatures to and time for which these samples were heated are detailed in the results section of this thesis.

(a) (b)

Figure 21: Pictures of (a) tube furnace and (b) box furnace that were used to pyrolyze and heat- treat the 3D-printed scaffolds and cold-pressed pucks

3.6 Microstructure Characterization

3.6.1 Metallographic preparation of sample

Once these cut samples are heat treated, they are mounted and polished using conventional metallographic sample preparation techniques for further characterization. More specifically, these samples were mounted using epoxy resin and hardener and were mechanically ground using 320,

400, 600 and 800 grit emery paper followed by cloth polishing with 6, 3, and 1 µm diamond solution and 0.05 µm alumina suspension. Each of these steps was carried out on an Allied MetPrep4 as shown in Figure 22.

Figure 22: Allied MetPrep4 polisher used for sample preparation

3.6.2 Confocal microscopy

Once the samples are mounted and polished, they are examined using a Keyence VKX-250X confocal microscope with a 405 nm solid state laser (Figure 23). The samples are imaged after each heat treatment to understand how each cycle affects the microstructure. Software available with the confocal microscope was used to analyze the dimensions of the particles and the 3D height map of the sample, which qualitatively correlates with the relative hardness of each particle type.

Figure 23: Keyence VKX-250X confocal microscope

3.6.3 Secondary and back-scattered electron imaging

SE and BSE images were captured using a Thermofischer APREO SEM (Figure 24) in the

Advanced Materials Characterization Center (AMCC) at the University of Cincinnati. A beam voltage of 10 kV and a beam current of 0.80 nA was used while adjusting the brightness and contrast to capture images of the heat-treated samples.

Figure 24: Image of a Thermofischer APREO SEM similar to the one used at AMCC

3.6.4 EDS and elemental mapping

The composition of particles in the heat-treated samples was analyzed via EDS using an

EDAX Octane Elite detector combined with Texture and Elemental Analytical Microscopy (TEAM) software. The polished samples were carbon coated prior to analysis. The beam voltage required to accurately detect any particular element should be around twice the value of the peak obtained from the K-line of that element. However, given that EDS is a semi-quantitative tool at best, particularly considering the fact that this study is examining materials containing carbon and oxygen (i.e. light weight elements) quantification was not performed and only elemental spectra and mapping are reported to discuss relative intensities of each element in the microstructure of sintered samples.

4 RESULTS AND DISCUSSIONS

4.1 Powder Characterization

Feedstock plays a major role in any manufacturing process and is especially important in

AM techniques. For the purpose of these experiments, the feedstocks used were Ni and TiC powders.

The size distribution and morphology of these powders were significantly different, as shown in

Figure 25, which presents optical micrographs of polished cross-sections of the Ni and TiC powders.

Scaffolds printed via extrusion of particle-based inks are expected to sinter better if the particles in the ink are closely packed and the metal binder forms a continuous network in the space between the ceramic particles. With this in mind, a large difference in particle size between the Ni and TiC powders was intentionally chosen with Ni being the smaller of the two as they could easily fill in the gaps formed by the uneven packing of larger irregularly shaped TiC particles. The Ni particles were used in the as-received condition with a nominal size distribution of 3-7 µm and the TiC particles were first sieved for 90 minutes to achieve a powder size in the range of 25-45 µm.

(a) (b)

Figure 25: Confocal images of (a) as-received Ni powder (b) as-received TiC powder

It was observed that the Ni particles are roughly spherical in shape, as can be seen in Figure

26, although agglomeration of nickel powders was evident based on the micrograph. This was assumed to be mainly due to the extremely small size of these particles. However, the agglomeration can be remedied by suspending these particles in 2-Bu which is one of the components of the tri- solvent used in making the ink. 2-Bu acts as a surfactant and forms a thin layer over the powder particles preventing agglomeration and acts as a medium for dispersing them.

(a) (b)

Figure 26: Confocal image of (a) Ni particles and (b) TiC particles at the same high magnification to compare the size of a single TiC particle to the size of Ni particles

The TiC particles, however, are significantly larger compared to the size of Ni powder as expected and intended when choosing this feedstock. They are irregularly shaped, faceted, and appear to contain pores or inclusions in the particles, which could either be carbon-rich regions or processing defects in the powder particles.

4.2 Ink Characterization and Printing

The ink is prepared by vortexing all of the components comprising the ink in the right proportions. Extrusion-based printing of intricate structures such as scaffolds, that contain multiple overhanging structures and strands, relies heavily on the properties of the particle-based inks that are used to fabricate them.

Various factors such as the rheology, powder particle size distribution, particle packing, particle to PLGA ratio, and type of powder particles used affect the sintering behavior and densification of the final component. However, for the purpose of this study, these parameters were kept constant as this work was conducted to assess the feasibility of using this technique to print cermets. However, the printed strands were characterized using a confocal microscope to observe the powder distribution in the polymer binder in the as-printed condition as shown in Figure 27.

Figure 27: Confocal image of a printed strand of Ni-TiC particle-based ink with Ni and TiC particles embedded in a polymer binder matrix

The nickel particles, being extremely small, are not clearly visible in the as-printed strand.

It is also observed that the particles are not closely packed as anticipated with the polymer matrix separating them, which is one of the primary reasons for incomplete sintering of the scaffolds upon heat treating them. Because of the presence of the PLGA, the TiC particles are also somewhat obscured, but they are clearly present and intersperse within the printed strand.

Moreover, it can also be observed that the large size and irregular shape of the TiC particles tends to reduce the flowability of the ink. This is can be observed in Figure 28 where long streaks of polymer are missing in regions behind TiC particles present in the ink when increasing the speed of printing and reducing the pressure of extrusion.

Figure 28: Confocal image of Ni-TiC particle-based ink illustrating how TiC particles hinder the flowability of ink at lower extrusion pressure and higher print speed

After the ink has reached the desired viscosity following sufficient vortexing and sonication and once the print parameters (i.e. speed and pressure) are optimized, it is then extruded from the

nozzle to print cylindrical and cubical scaffolds 10-15 mm in diameter/length with a height of approximately 5-15 mm with strands at approximately 1-1.2 mm apart (Figure 29). The exact diameter of the strands depends on the diameter of the nozzle from which the particle inks are extruded, the extrusion speed, and the extrusion pressure. Metallic tips with a diameter of 200-400

µm were used.

(a) (b)

Figure 29: (a) Photograph of as-printed cylinder Ni-TiC scaffold (b) Optical image of top-view of as-printed cylinder scaffold struts oriented at 90° to each other (c) longitudnal cross section of struts (d) lateral cross section of struts

4.3 Preliminary Heat Treatment Profile

Particle-based extrusion 3D printing has been and is still being studied by researchers from various institutions. Literature shows that the most common heat treatments used involved pyrolyzing the as-printed samples at 300°C for 1 hr under constant gas-purging to get rid of the polymer and all the solvents, followed by a sintering step at temperatures specific to the material system in use. Both these steps were carried out sequentially in a tube furnace with a constant flow of argon or hydrogen.

In an aim to replicate the heat treatment conditions used by previous authors for the pyrolysis step and to be able to maintain a constant gas flow, the tube furnace previously mentioned was used to process the scaffolds. However, due to limitations of this tube furnace the maximum sintering temperature achievable was 1100C. Given that sintering is a diffusion-based process, which follows an Arrhenius behavior temperature dependence, longer duration anneals were used at this relatively low temperature and additional heat treatments were conducted in a box furnace to be able to reach higher sintering temperatures as discussed in the following sections.

4.3.1 Effect of time on sintering behavior

Considering the limitations on the tube furnace and the heat treatment conditions used by previous researchers, the initial batch of 3D printed samples were pyrolyzed at 300°C for 1 hour under a constant flow of Ar + H2 followed by an annealing treatment at 1100C for 14 hours and 20 hours, respectively.

A confocal image of a sample annealed at 1100°C for 14 hours (Figure 30 (a)) shows TiC particles sparsely distributed with Ni particles forming necks and starting to sinter. TiC has a high hardness, comparable to that of the SiC particles used for grinding, which leads to the Ni and TiC particles polishing at different rates as can be seen based on the topological features of the TiC

particles in the confocal images. A confocal image of a sample annealed for 20 hours, as shown in

Figure 30 (b), appeared to have higher relative density than those sintered for 14 hours as expected, but the morphology and polishing rate of what was thought to be the TiC particles also changed significantly suggesting some kind of reaction took place during the additional anneal time.

Figure 30: Confocal images of Ni-TiC scaffold pyrolyzed at 300°C and sintered at 1100°C for (a) 14 hours and (b) 20 hours

In an attempt to increase the densification of the scaffold, these samples were further sintered at 1100°C for 96 hours in a box furnace after being encapsulated in a quartz tube under vacuum.

This process of encapsulation and heat treating in a box furnace was used so that an Ar flow was not required for such a long duration anneal. Samples annealed for 96 hours in the box furnace were observed under a confocal microscope, as shown in the micrograph in Figure 31, to determine the extent of densification.

Figure 31: Confocal image of Ni-TiC scaffold pyrolyzed at 300°C and sintered at 1100°C for 20 hours in the tube furnace under Ar flow followed by an additional 96 hours in a box furnace after encapsulating the sample under vacuum

Confocal images of these samples clearly show that the powder particles are not completely sintered and there is still significant porosity within the sample. The bright phase was assumed to be sintered nickel particles and the dark grey phase was originally thought to be titanium carbide.

Everything else between the powder particles is epoxy from the mount in which the samples are held, as labelled in Figure 31. Based on the lack of sintering, even after the extended duration anneal, a higher temperature sintering heat treatment was chosen to assess how much more the relative density could be increased.

4.3.2 Effect of temperature and vapor pressure

To study the effect of temperature on sintering behavior and relative density, samples that were pre-sintered at 1100°C for 20 hours were encapsulated again under vacuum and heat treated at

1400°C for 4 hours in a box furnace. This temperature was chosen considering the fact that nickel

melts at 1455°C. It was expected that annealing at this significantly higher temperature (increase of

300°C from 1100°C) would drastically improve the densification. However, upon conducting the experiment, it was observed that the quartz tubes in which these samples were encapsulated bloated, as shown in Figure 32, and the samples themselves became extremely fragile and essentially crumbled upon removal.

Figure 32: Photographs of sealed quartz tubes containing Ni-TiC samples before and after operating at 1400°C for 4 hours to show the extent of bloating

Such bloating of the capsule should only occur in the event that there is a buildup of gases inside the quartz tube, thereby increasing the pressure inside the capsule. Considering the different components of the as-printed samples (i.e. PLGA, solvents, nickel particles and TiC particles), the possible reasons for this bloating were expected to be one of the following:

1. Evaporation of PLGA

2. Evaporation of Ni

3. Oxidation of TiC and evolution of CO/CO2

Although previous researchers performed pyrolysis at 300°C for 1 hour to burn off all the polymer and solvents, the pyrolysis temperature for the polymer used in this study had not been

determined experimentally. To verify the pyrolysis temperature, pure PLGA, PLGA dissolved in

DCM and PLGA dissolved in the tri-solvent were analyzed using TGA as shown in Figure 33.

(a) (b)

Figure 33: TGA analysis results representing decrease in weight% with (a) increase in temperature and (b) increase in time

Results from the TGA analysis (Figure 33) indicate that all the PLGA (irrespective of whether or not it is dissolved in solvents) pyrolyzes at 350°C after heating at a constant rate of 5°C per minute for at least 60 minutes. This is the same heating used when heating the as-printed samples in the tube furnace. Therefore, the pyrolysis heat treatment was modified to 350°C for 1 hour before proceeding to the sintering step to ensure complete burn-off of the polymer and solvents.

Additionally, the results of the TGA seem to rule out the possibility that swelling of the quartz capsules was caused by the evaporation of PLGA or solvents as all of the PLGA and solvents should certainly burn-off in the tube furnace before encapsulating the sample.

Therefore, the reason for formation of gas within the capsule is narrowed down to two other potential options, the Ni and TiC powders. Titanium carbide has a very high melting point (3160°C) while nickel has a relatively low melting point of 1455°C and sublimes at 2837°C at atmospheric pressure. However, since the as-printed samples are encapsulated under vacuum prior to heat treating them, these samples are no longer at atmospheric pressure (760 Torr), but are at less than

10-2 Torr. Literature shows that the sublimation temperature of Ni can drop to as low as 1371°C if operated at pressures lower than 10-3 Torr as shown in Table 1 [78].

Table 1: Change in sublimation temperature of nickel with vapor pressure

Vapor pressure (Torr) Sublimation temperature (°C)

10-1 1679

10-2 1510

10-3 1371

10-4 1257

Although the exact pressure in the quartz capsule was not known, a vacuum of at least 10-2

Torr was achieved prior to sealing it. To confirm that the source of the vapor was the nickel powder, the same pure nickel powder that was used in the ink preparation was encapsulated in a quartz tube under vacuum and subjected to the same heat treatment profile (1400°C for 4 hours). The capsule swelled in a similar fashion indicating that the vapor came from these powders likely from Ni evaporation combined with water vapor from the surface of the very fine powder particles.

Therefore, annealing of such Ni-TiC scaffolds should not be performed at 1400°C or higher when operating under vacuum.

4.3.3 Effect of environment on oxidation behavior

Although it was observed that sublimation of nickel at 1400°C under vacuum causes the encapsulated quartz tube to bloat, one other major cause of concern is the oxidation of TiC. These

Ni-TiC scaffolds were initially annealed in a tube furnace under a continuous flow of Ar + H2 gas instead of pure Ar gas, with the 2% H2 in the gas serving as a getter to prevent oxidation of the

powder particles. However, it was observed that only 2% hydrogen in the argon was not sufficient to prevent oxidation of TiC.

Samples sintered in this environment showed a distinct oxide scale on the surface with segregation of free carbon near the center. Backscattered electron images, as shown in

Figure 34, were captured and EDS spot analysis was conducted on the two main phase constituents of the microstructure to determine the approximate compositions. Two EDS spectra are shown in Figure 35 that are representative of spots 1 and 2 in the micrograph in Figure 34.

Figure 34: BSE image of Ni-TiC scaffold pyrolyzed at 300°C for 1 hour and sintered at 1100°C for 20 hours with spots where EDS analysis was performed

Figure 35: EDS spot analysis results from spot 1 and spot 2 shown on the BSE image in Figure 34

Results from the spot analysis indicate that the bright phase is indeed Ni as expected. The dark grey phase however, contained high levels of oxygen along with titanium and carbon, while all the space between these particles is epoxy from the mount. To understand the oxidation behavior of

TiC in greater detail, elemental mapping of these microstructures was carried out as shown in Figure

36 and Figure 37. These elemental maps qualitatively display regions of the microstructure where a particular element is present. A higher concentration of an element in a particular area is denoted by a more intense color and lower concentration is denoted by the lack of color in any given region.

These maps give insight into the oxidation of TiC, showing that an oxide of titanium is formed on the surface of the TiC particles, liberating free carbon (graphite) near the center (as shown in Figure

36). These results indicate that the TiC particles completely oxidize when heat treated in tube furnace at 1100°C for 20 hours.

Figure 36: BSE image of Ni-TiC scaffold pyrolyzed at 300°C for 1 hour and sintered at 1100°C for 20 hours under Ar + H2 flow in the tube furnace

(a) (b)

(c) (d) Figure 37: Elemental maps of (a) nickel (b) titanium (c) carbon and (d) oxygen of a Ni-TiC sample pyrolyzed at 300°C for 1 hour and sintered at 1100°C for 20 hours under Ar + H2 flow in the tube furnace

4.3.4 Oxidation in scaffolds vs. cold-pressed Ni-TiC puck

To verify that the oxidation was not related to the presence of the solvents and polymer binder in the printed scaffold, Ni-TiC compacts, as shown in Figure 38, were fabricated by compressing pucks containing Ni and TiC powders in a 50:50 volume ratio to a load of 5 metric tons using a cold-press and were subjected the same heat treatment to compare the oxidation behavior. These heat-treated pucks were then mounted, polished, and observed under a microscope to analyze the microstructure. Figure 39(a) and (b) show a comparison of the two microstructures obtained within the scaffold and the puck, respectively.

Figure 38: Image of a Ni-TiC puck, cold-pressed at a load of 5 metric tons

(a) (b)

Figure 39: Confocal images of (a) a 3D-printed Ni-TiC scaffold and (b) a cold-pressed Ni-TiC puck heat treated at 300°C for 1 hour and 1100°C for 20 hours

The exact same morphology of individual oxidized TiC particles like within the scaffolds was not evident in the compressed puck because of the compaction. However, there was a continuous matrix of the dark grey phase, similar to that in the scaffold and likely TiO2 (rutile) based on [75], surrounding and engulfing the nickel. Sintering of TiO2 at this temperature was previously reported by Voitovich et al [73,75,79] during their study on the oxidation mechanisms of TiC. Such low temperature sintering of TiO2 was explained by the presence of free carbon driving the sintering process. To confirm the presence of oxygen in this phase, EDS analysis and elemental mapping

were performed on these heat-treated compressed pucks with the backscattered electron image shown in Figure 40 and the EDS spectra for spots 1 and 2 in Figure 41.

Figure 40: BSE image of cold-pressed puck heat treated at 300°C for 1 hour and 1100°C for 20 hours showing the locations on the sample where EDS spot analysis was performed

Figure 41: EDS spot analysis results of cold-pressed puck heat treated at 300°C for 1 hour and 1100°C for 20 hours

It is evident from EDS spot analysis of the heat-treated puck that the bright phase is indeed nickel rich and that the dark grey phase has a high concentration of oxygen and titanium. Elemental mapping was also performed on these microstructures to understand the distribution of each element as shown in Figure 42.

(a) (b)

Figure 42: Elemental maps of (a) nickel (b) titanium (c) carbon (d) oxygen of a Ni-TiC compressed puck pyrolyzed at 300°C for 1 hour and 1100°C for 20 hours under Ar + H2 flow in the tube furnace

Elemental mapping clearly indicates that the bright phase is indeed nickel and shows the presence of abundant amounts of titanium and oxygen in the regions corresponding to the dark grey phase with limited carbon present. Again, literature suggests that such drastic densification and

sintering of titanium oxide is accelerated by the presence of free carbon segregated during oxidation of TiC [75].

Based on this, it is clear that the oxidation was not limited to scaffolds fabricated via extrusion-based 3D printing, but also occurs in the cermet compacts fabricated via conventional powder metallurgy cold compaction techniques. Therefore, the heat treatment conditions must be modified to be able to sinter 3D-printed Ni-TiC cermet scaffolds which are more dense and structurally robust while preventing the oxidation of the TiC present.

4.4 Redefined Heat Treatment Profile

Based on the results presented thus far, there is a need to redesign the thermochemical treatment for this material system. From the data obtained from the TGA of PLGA dissolved in the solvents, it was found that all the polymer and solvents burn-off at 350°C. Hence, the pyrolsis step of the heat treatment was changed to 350°C at a constant rate of 5°C per minute for 60 minutes to make sure all the polymer and solvents burn off completely.

Although a high temperature sintering step is necessary to provide stability and densification, the preliminary results showed oxididation of TiC at temperatures as low as 1100°C. Therefore, it is necessary to change the second step of the heat treatment to prevent oxidation while maintaining the structural integrity of the scaffold.

4.4.1 Pre-sintering to prevent oxidation

While the high temperature anneal at 1100C under Ar + H2 flow in the tube furnace resulted in oxidation, if the scaffold was removed from the furnace only after the pyrolysis step it would be fragile and difficult to handle for any further processing. Therefore, the scaffolds need to be pre- sintered to a brown body state such that Ni particles have sufficiently sintered to maintain the structural integrity so they can be handled while preventing oxidation of the TiC. X-ray diffraction

phase analysis conducted by Voitovich and Pugach [79] detected oxide phases in TiC at temperatures as low as 700°C (Figure 43). Based on this, the pre-sinter heat treatment step was chosen to be conducted at 650°C to avoid oxidation while obtaining as much densification as possible.

Figure 43: List of various titanium oxides formed at different temperatures according to XRD anaylsis conducted by Voitovich and Pugach [75]

After testing this heat treatment for different intervals of time, it was determined that heating the sample to 650°C for a minimum of 16 hours produced a stable pre-sintered scaffold structure.

Therefore, the samples were first pyrolyzed at 350°C for 1 hour and then pre-sintered at 650°C for

16 hours to create brown body scaffolds that do not break upon handling. These brown body samples were then analyzed using EDS elemental mapping to confirm that no oxidation of TiC took place using this revised pyrolysis and pre-sinter heat treatment. The BSE image and elemental maps are shown in Figure 44 and Figure 45, respectively, for reference.

Figure 44: BSE image of Ni-TiC scaffold pyrolyzed at 350°C for 1 hour and pre-sintered at 650°C for 16 hours under Ar + H2 flow in tube furnace

(a) (b)

Figure 45: Elemental maps of (a) nickel (b) titanium (c) carbon and (d) oxygen of a Ni-TiC sample pyrolyzed at 350°C for 1 hour and pre-sintered at 650°C for 16 hours under Ar + H2 flow

These maps are a stark contrast to those reported for the samples annealed at 1100°C and show that the TiC particles remain intact and have not oxidized during pre-sintering. Therefore, the redefined heat treatment worked as expected and provided a brown body scaffold that can be further processed to reach a higher density.

4.4.2 Effect of sintering temperature

Once the Ni-TiC as-printed scaffolds are pyrolyzed and pre-sintered to form delicate brown bodies, they need to be further sintered to achieve higher densification. However, this sintering step cannot be performed in the tube furnace under Ar + H2 atmosphere due to the oxidation of TiC under these conditions. Because heat treating these scaffolds at any temperature higher than 700°C in a non-reducing atmosphere has proven to be deleterious due to oxidation, these pre-sintered brown bodies were carefully transferred into a quartz tube and encapsulated under vacuum. However, encapsulating the brown-body alone in vacuum would still leave the scaffold exposed to trace amounts of oxygen present inside the capsule. Based on the Ellingham diagram, as shown in Figure

46, zirconium oxide has a lower Gibbs free energy than titanium oxide and is also stable at high temperatures. Therefore, zirconium should be effective as an oxygen getter for this system and, hence, a small piece of zirconium foil was added to each capsule.

Figure 46: Ellingham diagram to determine the stability of oxides at different temperatures

With the pyrolysis treatment, pre-sintering treatment and the environment in which sintering would be carried out being determined, the only parameter remaining was to determine the temperature at which final sintering was to be performed. Previous research carried out by Taylor et al [28] shows that pure Ni scaffolds fabricated via the particle-based ink extrusion technique used here were sintered at 1300°C to attain a densification of more than 98%. However, the experiments carried out by Taylor et al were under a reducing atmosphere (H2 furnace) and not under vacuum.

As mentioned earlier, annealing these samples encapsulated in a quartz tube at 1400°C resulted in bloating of the capsules, but the slightly lower temperature of 1300°C should be low enough to prevent Ni evaporation under the vacuum level achieved during encapsulation [78] . Therefore, the scaffolds were encapsulated under vacuum along with Zr foil and annealed at 1300°C for 1, 2, and

4 hours.

Although full densification was not expected, the scaffolds were expected to sinter much more significantly at such a high temperature compared to 1100°C. However, following the heat treatment and removal from the capsules, the sintered samples were extremely fragile and would disintegrate upon handling. There were, however, what appeared to be metallic droplets that appeared on the surface of the scaffold as can be seen in the two photographs in Figure 47.

Figure 47: Pictures of a scaffold sintered at 1300°C under vacuum after pyrolyzing and pre- sintering heat treatment

After careful metallographic preparation of one of these droplets, they were imaged using a confocal microscope and it was observed that they contained a lamellar microstructure suggesting the presence of a eutectic as shown in Figure 48.

Figure 48: Different magnification confocal images of the cross-section of a droplet found on the surface of a scaffold sintered at 1300°C under vacuum after pyrolyzing and pre-sintering

A 50:50 volume ratio of Ni-TiC translates to approximately 64.25 wt.% nickel. At this composition of nickel, there is a eutectic between Ni and TiC at a temperature of 1270°C according to the pseudo-binary phase diagram shown in Figure 49 [71].

Figure 49: Phase diagram of Ni with stoichiometric TiC at equilibrium

One of these droplets was characterized using SEM to analyze the finer details of the microstructure and elemental mapping was performed to determine the elemental distribution as shown in Figure 50 and Figure 51. The results from the elemental mapping of these droplets indicate that the dominant phase is primary Ni, which does not contain any titanium, and the lamellae are made up of alternating Ni and TiC. This microstructure is what would be expected based the presence of a Ni-TiC eutectic given the ink composition and is, therefore, in agreement with the phase diagram reported in literature.

Figure 50: BSE image of a droplet found on the surface of a scaffold sintered at 1300°C under vacuum after pyrolyzing and pre-sintering heat treatment

(a) (b)

Figure 51: Elemental maps of (a) nickel (b) titanium (c) carbon (d) oxygen of a droplet found on the surface of a scaffold sintered at 1300°C under vacuum after pyrolyzing and pre-sintering

4.4.3 Effect of sintering time

Because of the presence of the eutectic and liquid formation during sintering at 1300°C when operating under vacuum, subsequent sintering anneals were reduced to a temperature of 1250°C.

Both pre-sintered brown bodies and cold-pressed pucks were encapsulated in vacuum along with a strip of Zr foil and were annealed at 1250°C for 1 hour, 2 hours, and 4 hours for comparison. It should be noted that, because these pucks are compressed unlike the as-printed scaffolds which are extruded at low pressures with PLGA separating the powders, a significant difference in densification is expected. This is evident even after the pre-sintering step as shown in Figure 52.

(a) (b)

Figure 52: Confocal images of (a) as-pressed puck (b) as-printed scaffold (c) cold-pressed puck heat treated at 350°C for 1 hour and at 650°C for 16 hours (d) 3D-printed scaffold heat treated at 350°C for 1 hour and at 650°C for 16 hours

Following anneals for 1, 2 and 4 hours at 1250°C, it is clear that the relative density of the puck is drastically higher than that of the scaffolds based on the micrographs shown in Figure 53 for each time step. It can also be observed that the densification in both the scaffolds and the pucks did not vary drastically with increase an in time from 1 to 4 hours. To determine if a much higher densification could be achieved in this system, one scaffold and one cold-pressed puck of Ni-TiC were each encapsulated in different quartz tubes under vacuum along with Zr foils and were annealed at 1250°C for 48 hours, the results of which are shown in Figure 54.

(a) (b) (c)

(d) (e) (f)

Figure 53: Confocal images of (a,b,c): 3D-printed scaffold heat treated at 350°C for 1 hour, 650°C for 16 hours, and 1250°C for 1 hour, 2 hours and 4 hours, respectively, and (d,e,f): cold pressed puck heat treated at 350°C for 1 hour, 650°C for 16 hours, and 1250°C for 1 hour, 2 hours, and 4 hours

Figure 54: Confocal image of (a) 3D printed scaffold and (b) Ni-TiC cold-pressed puck, heat treated at 350°C for 1 hour, 650°C for 16 hours and 1250°C for 48 hours

Densification of the cold-pressed puck appears to be only slightly better than in the samples sintered for shorter times. However, the additional sintering time did not appear to influence the densification of the scaffold. To verify that such a long duration anneal did not lead to oxidation of the TiC, these scaffolds were characterized via SEM as shown in Figure 55 through Figure 57.

Figure 55: BSE image of 3D-printed scaffold heat treated at 350°C for 1 hour, 650°C for 16 hours and 1250°C for 48 hours showing the locations on the sample where EDS spot analysis was performed

Figure 56: EDS spot analysis results of 3D-printed scaffold heat treated at 350°C for 1 hour, 650°C for 16 hours and 1250°C for 48 hours

(a) (b)

Figure 57: Elemental maps of (a) nickel (b) titanium (c) carbon (d) oxygen of 3D printed scaffold heat treated at 350°C for 1 hour, 650°C for 16 hours and 1250°C for 48 hours

Although EDS can only be used as a semi-quantitative measurement of composition, from these results, it again appears that TiC has not oxidized even after sintering at 1250°C for 48 hours in vacuum along with a sacrificial material. However, it can also be observed that the densification is not appreciable and seems to have reached a threshold value and does not increase any further sintering time. To quantitatively measure the change in densification as a function of sintering time, the relative density was calculated for each sintering time at four different regions of the sample using ImageJ. The confocal micrographs were segmented to distinguish between material and porosity within both the scaffolds and pucks that were pyrolyzed, pre-sintered, and subsequently sintered at 1250°C for 1 hour, 2 hours, 4 hours and 48 hours.

Table 2 and Table 3 list the average relative density and standard deviation of pre-sintered and sintered scaffolds and pucks, respectively. It is observed that the average density of cold-pressed

Ni-TiC pucks is much higher (approximately 20-30%) than the average density of the 3D printed scaffolds. As previously mentioned, this is to be excepted because the cold-pressed pucks are comprised of pure Ni and TiC powders that were subjected to loading as high as 5 metric tons, which compact the powder particles so that there is much higher contact between them than in the extruded ink with PLGA separating them. Due to this cold-pressing process, the average density of the as- pressed pucks is much higher than the as-printed samples in the initial stage. Figure 58 shows a plot of relative density as a function of sintering time for both the as-printed and cold-pressed pucks to compare the behavior.

Table 2: Average relative density and standard deviation of pre-sintered and sintered Ni-TiC scaffolds

Sample Avg. relative density % Standard deviation

Pre-sintered Scaffold 33.4 5.2

Annealed at 1250°C for 1 hour 49.1 2.1

Annealed at 1250°C for 2 hours 50.5 1.5

Annealed at 1250°C for 4 hours 50.4 3.1

Annealed at 1250°C for 48 hours 56.0 1.8

Table 3: Average relative density and standard deviation of pre-sintered and sintered cold-pressed Ni-TiC pucks

Sample Avg. relative density % Standard deviation

Pre-sintered Puck 68.5 1.6

Annealed at 1250°C for 1 hour 72.0 1.1

Annealed at 1250°C for 2 hours 77.0 0.9

Annealed at 1250°C for 4 hours 79.4 2.1

Annealed at 1250°C for 48 hours 80.9 1.1

Figure 58: Plot of average relative density vs. sintering time for scaffolds and pucks sintered at 1250°C

This plot between the average relative density and sintering time (Figure 58) indicates that after approximately 10 hours, the densification does not increase appreciably with increased anneal time at the same temperature. Moreover, we can also observe the rate of sintering in the scaffolds is extremely high initially compared to the rate of sintering of the cold-pressed Ni-TiC pucks.

However, it can be observed that the density of 3D printed scaffold structure that was sintered for

48 hours does not even reach the initial density of pre-sintered cold-pressed puck. This shows that the maximum densification achievable for the Ni-TiC cermet scaffolds is nowhere near sufficient under these conditions, but that these scaffolds need to be sintered at higher temperatures and/or pressures to get anywhere near a fully dense structure.

5 SUMMARY AND CONCLUSIONS

Ni-TiC cermet scaffolds with a 50:50 volume ratio of Ni to TiC were 3D printed using a particle-based ink extrusion technique. The ink used has three main components of the Ni and TiC particle base, a PLGA polymer binder, and a tri-solvent mixture. It was observed that attaining an appropriate viscosity of ink is extremely important in printing and maintaining the structure of these intricately shaped scaffolds. Extrusion pressure, print speed and other print parameters also play a major role in fabricating such scaffolds and these parameters are to be calibrated depending on the material system that is being printed.

The as-printed structures were subjected to various heat treatment conditions to understand the sintering behavior of these Ni-TiC scaffolds. Efforts to sinter Ni-TiC scaffolds in Ar + H2 atmosphere did not yield promising results. Only 2% hydrogen present in the cover gas proved to be insufficient in preventing TiC particles from oxidizing. The oxidation of TiC at higher temperatures gave rise to the formation and segregation of free graphite. This free graphite was reported in literature to aid the sintering of the titanium oxide formed, thereby essentially producing

Ni-TixOy cermets instead of Ni-TiC cermets as desired.

To prevent TiC from oxidizing, there was a need to sinter these scaffolds under vacuum, in the presence of a sacrificial material that can oxidize preferentially. However, since the as-printed structures contained not just the powder particles, but also considerable amounts of PLGA and solvents, it was important to perform a pyrolysis heat treatment in a tube furnace before sintering the scaffolds in vacuum. Therefore, as-printed scaffolds were pyrolyzed at 350°C in a tube furnace under a constant flow of Ar + H2 gas followed by pre-sintering at 650°C for 16 hours in the same tube furnace before encapsulating them in vacuum along with a strip of Zr foil, which acts as the sacrificial material.

These pre-sintered samples encapsulated in a quartz tube under vacuum were then annealed at different temperatures (1400°C, 1300°C and 1250°C) for different time intervals (1, 2, 4, and 48 hours). However, it should be noted that encapsulating the scaffolds in vacuum limits the maximum temperature at which these scaffolds can be sintered successfully. The quartz tube in which these pre-sintered samples were encapsulated swells when operated at a temperature of 1400°C, likely occuring due to the sublimation of Ni and/or vaporization of water remaining on the powder surfaces, producing gas which increased the pressure inside the capsule causing it to bloat.

Annealing the samples at 1300°C under vacuum was also unsuccessful since at this temperature, Ni was found to form a eutectic with TiC. This eutectic results in the formation of a liquid phase, which under vacuum is wicked out of the scaffold and form droplets on the surface upon cooling. Annealing these scaffolds at a slightly lower temperature (1250°C) did not produce any such defects in the samples. The Ni particles sintered enough to maintain the structural integrity of the scaffolds when they were sintered at 1250°C for 1 hour, 2 hours, 4 hours and 48 hours respectively. However, there was only an incremental gain in densification observed upon sintering for the longest duration of 48 hours at this temperature and the final relative density was still less than the initial density of a pre-sintered cold-pressed puck of only Ni and TiC.

This study provides preliminary results required to 3D print Ni-TiC cermet scaffolds using particles-based liquid inks via extrusion. This technique could be used to fabricate stable scaffold structures of multi-component systems, especially cermets,that are difficult to sinter via traditional metallurgical techniques and other more conventional powder bed AM techniques. However, a more thorough investigation varying pre-processing, printing, and post-processing parameters needs to be conducted to make this truly a feasible approach to fabricating such materials. Some promising approaches and alternatives to move this research forward are discussed in the next section.

6 FUTURE WORK

Ni-TiC cermet scaffolds fabricated using an extrusion-based technique are relatively stable and maintain their shape, demonstrating that this technique of 3D printing using particle-based inks is a viable option. However, much work needs to be done to improve the structure and mechanical properties of these scaffolds if the long-term goal is to be able to use them in industrial applications.

Below are a few suggestions for future work that could be carried out with the goal of achieving higher densification and to fabricating scaffolds possessing significantly better properties. These can be classified into three areas of research and are as follows:

1. Modifications to the ink:

In this study, a large difference in the size of the Ni and TiC particles was maintained

with the assumption that the smaller Ni particles would infiltrate the gaps between the larger

TiC particles, eventually sintering and encompassing the TiC particles. However, this size

difference did not appear to have been beneficial in terms of the final achievable

densification. It is necessary to study the effect of the size and morphology of TiC particles

on the sintering behavior of these Ni-TiC cermet scaffolds.

Moreover, the PLGA and tri-solvent mixture in which the powder particles are

dispersed burns off during the pyrolysis step. This leads to the formation of pores in the as-

printed struts, thereby creating larger gaps between the particles, which makes it more

difficult for the Ni powder to sinter and surround the TiC particles. Therefore, there is a need

to experiment with the powder to PLGA ratio or particle volumetric loading used in the ink

to minimize the amount of polymer in the as-printed struts while still being able to hold the

structure together prior to sintering.

A more promising ink can be prepared using TiC particles coated with appropriate

amounts of Ni on their surface. Such coated TiC particles would help increase the green

body packing density, thereby promoting better sintering. The Ni coating on the surface of

TiC particles would also aid in forming a continuous matrix of Ni, as no TiC particles would

be directly in contact with each other, thereby improving the densification of the sintered

strut through solid-state sintering of Ni only.

2. Modifications to the post printing heat treatment

The purpose of fabricating Ni-TiC scaffolds via extrusion-based 3D printing

technique is to develop a more efficient and simple process of creating intricate structures

without the use of expensive and sophisticated equipment. However, pre-sintering these

scaffolds for 16 hours and then sintering them for a few hours more while encapsulating in

vacuum is not the most efficient way to fabricate such scaffolds. For this purpose, it is

important to understand the sintering behavior under reducing atmosphere such a hydrogen

furnace or a high temperature vacuum tube furnace to eliminate the pre-sintering step of heat

treatment and conduct the complete heat treatment cycle in one furnace.

Heating these 3D printed scaffolds in a high temperature tube furnace that can reach

1400°C would allow the entire heat treatment cycle to be conducted in one furnace.

However, since TiC oxidizes at higher temperatures, these scaffolds are to be sintered in an

environment that prevents oxidation. For this purpose, Ar + H2 gas with higher amount of

H2 should be experimented with. Sintering these scaffolds under vacuum would also reduce

the oxidation of TiC. However, under vacuum, the highest temperature at which these

scaffolds can be sintered is limited. To overcome this difficulty, the scaffolds can be

pyrolized at 350°C for 1 hour while constantly purging Ar + H2 gas, to make sure all the

gases produced due to pyrolysis escape. After pyrolysis, vacuum is created in the tube

furnace to make sure all the gases in the tube furnace (including oxygen) are flushed out of

the tube furnace. Once the tube furnace is under vacuum, Ar + H2 is then purged in the tube

to make sure the tube is at atmospheric pressure. Although the tube is now being purged by

Ar + H2 gas, there would still be trace amounts of oxygen in the gas. To prevent further

oxidation, the scaffolds can be held in a closed crucible with Zr foil.

Moreover, it would also be beneficial to study the densification of the scaffolds when

sintered under pressure at higher temperatures. At high enough temperatures (> 1300°C

based on the work conducted here), a eutectic forms between the Ni and TiC. Having such

a liquid present could be beneficial such that liquid phase sintering occurs if the wetting of

the liquid is sufficient. Literature states that in the absence of a driving force, due to poor

wetting behavior, the use of an external applied force may be used to overcome the negative

capillary pressure, thereby forcing the metallic binder to infiltrate the pores between the

ceramic particles. Hot isostatic pressing (HIPing) is one such technique, where a high-

pressure gas is used at high temperature to infiltrate the binder material between the ceramic

particles [11]. This process uses isostatic pressure and, hence, the strength of the green body

does not need to be very high and the intricate geometry of the scaffolds can be maintained.

3. Modifications to the material system

The goal of fabricating Ni-TiC cermets for the purpose of this study was to

understand the feasibility of manufacturing such intricate structures via extrusion-based 3D

printing and to observe the sintering behavior of these cermet scaffolds. The cermets used in

various industrial and structural applications, however, contain more additives and a plethora of different binders and ceramic reinforcements depending on the applications in which these cermets are used.

Addition of Mo and Nb have been proven to increase the wetting of Ni on TiC particles [21,59]. Moreover, a combination of metals such as Ni-Co can be used as a binder to bind various ceramic reinforcements such as TiC, Ti(C,N) or WC [80,81]. Once the right parameters to manufacture dense structures are determined, there are a variety of other material systems that could be fabricated via this technique.

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