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UNIVERSITY OF CALIFORNIA

LOS ANGELES

ZrB2-Based Composites for

Ultra-High-Temperature Applications

A dissertation submitted in partial satisfaction

of the requirements for the degree

Doctor of Philosophy in Materials Science and Engineering

By

Do Hwan Chung

2010

UMI Number: 3463935

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© copyright by

Do Hwan Chung

2010

The dissertation of Do Hwan Chung is approved.

Suneel Kodambaka

Daniel Yang

^ ^ ^ "V //£ /0#/<o

ta Jdnn/Ming Yafid Committee/Chair

University of California, Los Angeles

2010

Dedicated to my father, my mother, my wife, and my two sons.

in

TABLE OF C O N T E N T S

LIST OF FIGURES vii

LIST OF TABLES xi

ACKNOWLEGEMENTS xii

VITA xiv

PUBLICATIONS xvi

ABSTRACT xvii

Chapter 1 Ultra-high-temperature ceramics 1

1.1 Introduction 1

1.2 Applications of ultra-high-temperature ceramics 2

References 7

Chapter 2 Literature survey of ultra-high-temperature ceramics 8

2.1 Selection of ultra-high-temperature ceramics 8

2.2 Oxidation mechanism of ZrB2-SiC ceramics 9

2.3 Mechanical properties of ZrB2-SiC ceramics 12

2.4 Spark plasma sintering 14

2.5 Liquid phase sintering .........16

References 29

Chapter 3 ZrB2-nano SiC ceramics ............................................................................32

Abstract .....................................32

3.1 Introduction........................................................... 34

3.2 Experimental procedure....................................................................................35

3.2.1 Materials 35

3.2.2 Elastic property measurements 38

IV

3.2.3 Mechanical properties 38

3.2.4 Micropillar compression test 40

3.3 Results and discussion 40

3.3.1 Densities, phases, and microstructures 40

3.3.2 Elastic moduli, hardness and fracture toughness 43

3.3.3 Micropillar compression test of ZrB2-nano SiC composite materials 45

3.4 Conclusion 47

References 72

Chapter 4 Multiphase ZrC-ZrB2-SiC ceramics 74

Abstract 74

4.1. Introduction 76

4.2. Experimental procedures 78

4.2.1 Materials 78

4.2.2 Elastic moduli, hardness, and fracture toughness measurements 79

4.2.3 Thermal and electrical conductivity measurements 80

4.2.4 Micropillar compression test 81


4.3.1 Densities, phases, and microstructures ....82

4.3.2 Elastic moduli, hardness, and fracture toughness 85

4.3.3 Thermal conductivity 87

4.3.4 Electrical conductivity ............................................................................89

4.3.5 Micropillar compression tests of ZrC-ZrB2-SiC composites ....................91

4.4. Conclusion .......................................93

References ..........................119

V

Chapter 5 SCS 9a fiber-reinforced ZrB2-RB(Reaction-bonded) SiC composite 121

Abstract 121

5.1 Introduction 123

5.2 Experimental procedures 124

5.2.1 Fabrication 124

5.2.2 Phase analysis, density, and porosity measurements 125

5.2.3 Elastic properties, hardness, and fracture toughness measurements 126

5.2.4 Microstructure observations 126


5.3.1 Density, porosity, and phase 128

5.3.2 Elastic moduli, hardness, and fracture toughness 129

5.3.3 Microstructural analysis 130

5.4 Conclusion 132

References 146

Chapter 6 Conclusions 148

vi

LIST OF FIGURES

Figure 1-1. Re-creation of the Shockwave during re-entry 5

Figure 1-2. Belly-down attitude during re-entry 6

Figure 2-1. Several materials with melting temperatures approaching or above 3000°C

[1] 21

Figure 2-2. ZrB2-SiC phase diagram [12] 22

Figure 2-3. The schematic diagram of a proposed oxidation mechanism for ZrB2-SiC

ceramics, (a) unoxidized ZrB2-SiC (b) the initial response during heating (c)

evolution as the temperature approaches 1500°C (d) steady state at 1500°C

[13] 23

Figure 2-4. The cross-sectional microstructure of ZrB2-SiC ceramics [13] 24

Figure 2-5. Microstructure of ZrB2-SiC ceramics [17] 25

Figure 2-6. Schematic of the SPS process [19] 26

Figure 2-7. Basic mechanism of spark plasma sintering [20] 27

Figure 2-8. Three different steps of liquid phase sintering by Kingery 28

Figure 3-1. Classification of ceramic nanocomposites 54

Figure 3-2. X-ray diffraction patterns of each composite material (a) ZSN-0 (b)

ZSN-5(c)ZSN-10(d)ZSN-20(e)ZSN-30 55

Figure 3-3. Typical FE-SEM images of each sample 56

Figure 3-4. The change of average grain size of each component.............................57

Figure 3-5. Typical TEM microstructures of each sample ...............58

Figure 3-6. Typical SiC grain locations in ZrB2-SiC composite ...............................59

vii

Figure 3-7. Typical interface between ZrB2 grain and SiC grain (ZSN-5: HRTEM

image) 60

Figure 3-8. The TEM image and EDX result of second phase, Zr(0, B)x, in ZrB2

grain, ZSN-20 61

Figure 3-9. The change of hardness and fracture toughness of ZrB2-nano SiC

composite with SiC content 62

Figure 3-10. Typical cracking pattern of ZrB2-SiC composites (ZSN-0, ZSN-10, and

ZSN-20) 63

Figure 3-11. Fabricated micropillar of ZSN-0, ZSN-5, and ZSN-10 64

Figure 3-12. Stress vs. height change relation by micropillar compression test for

ZSN-0 65


ZSN-5 66


ZSN-10 67

Figure 3-15. ZSN-10 micropillar (a) before and (b) after compression 68

Figure 3-16. TEM microstructural images of ZSN-0 by the micropillar compression

test (a) Whole TEM laminar (b) Stacking fault (c) dislocation 69


test (a) Whole TEM laminar (b) Dislocations (c) Stacking fault 70

Figure 3-18. Fracture surface images of ZSN-5 by the micropillar compression test (a)

Fracture surface of whole micropillar (b) Intragranular fracture (c)

Intergranular fracture .............................................71

viii

Figure 4-1. Typical examples of recorded shrinkage curves during the SPS cycle for

ZZS-2, ZZS-3, and ZZS-5 101

Figure 4-2. X-ray diffraction patterns for each sample; (a) ZZS-1, (b) ZZS-2, (c)

ZZS-3, (d) ZZS-4, (e) ZZS-5, (f) ZZS-6, (g) ZZS-7, (h) ZZS-8 102

Figure 4-3. Typical FE-SEM images for the ZrEb-ZrC-SiC composites 103

Figure 4-4. Typical TEM images for the ZZS-1 composite 104

Figure 4-5. Typical TEM images for the ZZS-2 composite 105

Figure 4-6. Typical TEM images of interphase interface between (a) SiC and ZrE$2

and (b) ZrB2 and ZrC 106

Figure 4-7. Typical cracking behavior of ZrC-ZrB2-SiC composites 107

Figure 4-8. Current versus voltage measured at room temperature for the

ZrC-ZrB2-SiC composites 108

Figure 4-9. Fabricated micropillar of ZZS-1, ZZS-2, ZZS-6, and ZZS-7 109

Figure 4-10. Stress vs. strain (%) curve of ZZS-1 micropillar 110

Figure 4-11. Stress vs. strain (%) curve of ZZS-2 micropillar I l l

Figure 4-12. Intragranular fracture surface of ZZS-1 micropillar 112

Figure 4-13. Intragranular fracture surface of ZZS-2 micropillar 113

Figure 4-14. Stress vs. strain (%) curve of ZZS-6 micropillar 114

Figure 4-15. Stress vs. strain (%) curve of ZZS-7 micropillar................................ 115

Figure 4-16. Dimensional and shape change of ZZS-6 micropillar 116

Figure 4-17. Typical cracking behavior of ZZS-7 micropillar................................ 117

Figure 4-18. Typical TEM images of ZZS-6 micropillar after compression stress. .118

Figure 5-1. Detailed dimensions of SCS 9a fiber-reinforced composites for fracture

toughness 135

ix

Figure 5-2. FE-SEM images of SCS 9a fiber-reinforced composite; (a) Top surface

view, (b) Cross-sectional view 136

Figure 5-3. Typical X-ray diffraction pattern of SCS 9a fiber-reinforced ZrB2-RB SiC

composite 137

Figure 5-4. Typical indentation mark and cracking pattern of SCS 9a fiber composite.

138

Figure 5-5. FE-SEM images of SCS 9a fiber-reinforced composite fracture 139

Figure 5-6. Whole cross section view of SCS 9a fiber-reinforced composite 140

Figure 5-7. Typical FE-SEM images of cracks inside and outside the composite ... 141

Figure 5-8. FE-SEM images of residual carbon from reaction bonding and EDS result.

142

Figure 5-9. The FE-SEM image of the interface between fiber and matrix 143

Figure 5-10. Typical TEM microstructures and SAED (Selected Area Electron

Diffraction) of the composite 144

Figure 5-11. Typical TEM microstructural images and EDS results of the interface

between the fiber and the matrix; (a) Carbon coating layer of the fiber (b) ZrE$2

grain (c) ZrB2 and SiC agglomerate (d) ZrB2 and SiC grains 145

x

LIST OF TABLES

Table 2-1. Physical and mechanical properties with reducing SiC particle size of

ZrB2-SiC ceramics [17] (10, 1.4, 0.7, 0.45 in sample names are starting SiC

particle size of micron) 19

Table 2-2. Physical and mechanical properties of ZrB2-SiC ceramics with different

hot pressing time and temperature [18] (A, B: ZrB2 starting powder size; A:

6um, B: 2um, 1850, 1950, and 2050: temperature; °C, 45, 90, and 180: time;

min) 20

Table 3-1. Compositions of each sample 49

Table 3-2. Densities of each sample 50

Table 3-3. Average grain sizes of ZrB2-nano SiC composite materials 51

Table 3-4. Elastic properties of ZrB2-nano SiC composite materials 52

Table 3-5. Hardness and fracture toughness of ZrB2-nano SiC composites materials 5 3

Table 4-1. Compositions of ZrC-ZrB2-SiC composites 95

Table 4-2. Densities and elastic properties of ZrC-ZrB2-SiC composites 96

Table 4-3. The measured grain size of each component ZrB2, ZrC, and SiC for

ZrC-ZrB2-SiC composies 97

Table 4-4. Hardness and fracture toughness of ZrB2-ZrC-SiC composites 98

Table 4-5. Thermal properties measured at room temperature for the ZrC-ZrB2-SiC

composites 99

Table 4-6. Electrical properties measured at room temperature for the ZrC-ZrB2-SiC

composites................................................................................................. 100

Table 5-1. Gibbs free energy of formation of ZrB2, C, ZrC, and B4C at 298K, 1000K,

and 2000K [12] 134

xi

ACKNOWLEGEMENTS

This research is supported by the National Aeronautics and Space Administration

(NASA) Glenn Research Center in Cleveland Ohio. Narottam P. Bansal is the

program manager.

I would like to give my thanks to all of the committee members, Professor

Suneel Kodambaka, Professor Daniel Yang, and Professor Jenn-Ming Yang for the

satisfaction of this dissertation.

I wish to express my deepest gratitude to professor Jenn-Ming Yang, advisor,

who introduced me to the field of ultra-high-temperature ceramics and who has

guided my doctoral studies all the way. I would like to thank you for supporting and

encouraging me throughout these years. Your incredible enthusiasm and optimistic

dedication to science has impressed me deeply. I am especially grateful for your great

ability not only to share your academic knowledge, but also to help organize my

research in a comprehensive means. Your door has always been open, and I have

immeasurably treasured your advice and guidance, both in the scientific field and in

daily life. Without your help and support, it would have been impossible for me to

complete this dissertation.

I also want to thank my father and mother for struggling with all the

accompanying difficulties so that I could concentrate on my work and smoothly finish

xii

my studies abroad. Your endless support and understanding are my driving force. Last

but not least, I would like to express my deepest gratitude to my life partner, Min Jung,

who is the light of my life. Thank you so much for our ten-year journey together, for

all the tears and happiness we have had together, and also for your endless support,

encouragement and love. With you, life will never be boring.

All my friends are of great importance to me. All of you enrich my life and make

it more colorful and joyful. All those wonderful times are unforgettable and highly

appreciated. I wish to thank all of you as much as I can. I cherish your friendship and

consideration.

Xll l

V I T A

1992-1995 Completion of two school years in Chemistry

Daejeon University

Daejeon, Korea (R.O.K.)

1993-1994 Military Service, Korea (R.O.K.)

Corporal, Army

1996-1998 Bachelor of Science in Chemistry

Sejong University

Seoul, Korea (R.O.K.)

1999-2000 Bachelor of Science in Ceramic Engineering

Yonsei University


2001-2002 Master of Science in Ceramic Engineering


>003~2004 Master of Science in Chemical Engineering and

Materials Science

University of Southern California

Los Angeles, CA, USA

2005-2010 Doctor of Philosophy in Materials Science and

Engineering

University of California, Los Angeles

Los Angeles, CA, USA

XV

PUBLICATIONS

[1] Seung Hyun Lee, Do Hwan Chung, and Joon Keun Park, (2010), In-situ surface

stress and magnetic properties of the CoCrX(X=Pt, Ta)/CrTi bilayer thin films during

sputter-deposition, Current applied physics, forthcoming.

[2] Shu-Qi Guo, Yutaka Kagawa, Toshiyuki Nishimura, Dohwan Chung, Jenn-Ming

Yang, (2008), Mechanical and physical behavior spark plasma sintered ZrC-ZrB2-SiC

composites, Journal of the European ceramic society, Volume 28, ppl279-1285.

[3] Kihyun Yoon, Do Hwan Chung, Byungduk Yang, Jaehyuk Jang, Jongheui Kim,

(2003), Preparation and characteristics of PTFE (Polyteterafluoroethylene)

composites for microwave circuit board, Journal of the Korean ceramic society,

Volume 40, pp735-738.

[4] Jo Jung, Ho Sung Lee, Kihyun Yoon, Do Hwan Chung, and Byungduk Yang,

(2002), Manufacturing method of PTFE composite board, Korea Patent, Application

number 1020020078716.

XVI

ABSTRACT OF THE DISSERTATION

ZrB2-Based Composites for

Ultra-High-Temperature Applications

by Do Hwan Chung

Doctor of Philosophy in Materials Science and Engineering

University of California, Los Angeles, 2010

Professor Jenn-Ming Yang, Chair

This dissertation presents an investigation into the processing, microstructure,

and mechanical behavior of ZrB2-based composites for ultra-high-temperature

applications. Various forms of SiC including nano-sized particles, micron-sized

particles, and continuous fibers were used as reinforcement.

Three major investigations were conducted in this dissertation. First, the effect of

incorporating nano-sized SiC particles into ZrB2 was investigated. Spark plasma

xvii

sintering was used to consolidate nano-sized SiC/ZrB2 composite. The detailed

microstructure of the composite was analyzed using transmission electron microscope.

Micropillar compression test was also conducted. It was found that incorporation of

nano-sized SiC effectively hindered the grain growth of ZrB2. The second study

focused on the ternary ZrC-ZrB2-SiC ceramics. The fully densed ceramics were

prepared by spark plasma sintering. Elastic modulus, hardness, fracture toughness,

thermal conductivity, and electrical conductivity of ternary ZrC-ZrB2-SiC ceramics

were measured. It was found that the fracture toughness of ternary ZrC-ZrB2-SiC

ceramics is comparable to that of the ZrB2 ceramics and ZrB2-SiC ceramics. In

addition, micropillar compression tests revealed information about typical

longitudinal cracking behavior and generation of stacking faults. The third part of this

dissertation focused on studying the effect of incorporating continuous SiC fibers on

the microstructure and properties of ZrB2. The composite was consolidated by

conventional hot pressing method. The chemical reaction between fiber and matrix

materials was not observed based on thermodynamic calculation and TEM

microstructural analysis. The fracture toughness of composite was measured to be

four times higher than that of matrix materials. However, extensive matrix cracking

was observed due to mismatch in thermal expansion coefficient between the fiber and

xviii

matrix.

Finally, the challenge and future research needs in developing

ultra-high-temperature ceramics are discussed.

XIX

Chapter 1

Ultra-high-temperature ceramics

1.1 Introduction

Ultra-high-temperature ceramics (UHTCs) are a family of compounds that are

chemically and physically stable at high temperatures (e.g., above 2400 °C) and in

reactive atmospheres (e.g., monatomic oxygen) [1]. UHTCs are famous for

possessing some of the highest melting temperatures of known materials. In addition,

they are very hard, have good wear resistance, mechanical strength, and relatively

high thermal conductivities (compared to other ceramic materials). Because of these

characteristics, UHTCs are ideal for thermal protection systems, especially those that

require chemical and structural stability at extremely high operating temperatures [2].

Some of the earliest work on UHTCs was conducted by the company ManLab in

the early 1960s, under a research program funded by the Air Force Materials Lab

(AFML) [3,4]. Research on UHTCs was started to meet the need for

high-temperature materials that would allow the development of maneuverable

hypersonic flight vehicles. Since then, intermittent research has made some progress,

but several significant challenges remain in the use of UHTCs, and these materials

have not yet to be widely implemented [2]. As the interest in monolithic UHTCs has

l

risen again in the early 1990s, high costs of raw materials have led to many new

investigations into different methods of fabricating UHTCs, such as reactive hot

pressing and pressureless sintering by liquid infiltration and reaction [5,6].

1.2 Applications of ultra-high-temperature ceramics

One application of UHTCs is the hypersonic flight vehicle whose speed begins

at four to five times the speed of sound (Mach 4 or 5). The current desired speed of

this vehicle is Mach 6, which produces a surface temperature of about 1400 °C. In

addition, the materials of this vehicle should be sustainable for several hundreds of

hours in this temperature range and in an oxidizing atmosphere. UHTCs with

improved mechanical and thermal properties are required for hypersonic flight

vehicle parts such as airframe components, engine cowls, windows, propulsion

system components, and control surfaces [1].

A second application of UHTCs is the orbital/re-entry vehicle. The maximum

temperature of the nose cone and the leading edges of the wings is about 1650 °C

during re-entry. To reduce the large amount of heat generated, blunt edged designs

are employed, which produce the re-creation of the blunt-body shock wave that

develops during re-entry. To take a full advantage of this effect, the vehicle re-enters

2

the atmosphere with a belly-down attitude. The re-creation of shockwave and

belly-down attitude is shown in figure 1-1 and 1-2, respectively. However, the blunt

edged design of the vehicle give two major problems. The first problem is limited

maneuverability which is from a narrow 'abort-to-land' window at launch and limited

cross-range on re-entry. The second major problem is communication blackout period

which is due to barrier layer between shockwave and blunt body described in figure

1-1. As result of these drawbacks of blunt body design, the future re-entry vehicles

should be employed with sharp leading edges and trailing edges, which improve

aerodynamic performance by allowing for laminar flow over control surfaces. Thus,

the vehicles with sharp leading edges and trailing edges enable an increased

abort-to-land window on launch. The ability to fly faster at any given altitude also

increases the cross-range for landing. Improved materials are needed for leading and

trailing edges, a thermal protection system, other hot structure components, windows,

and control surfaces.

A final application is the propulsion system. Propulsion systems can include

air-breathing engines such as turbines, ramjets, and scramjets. Rocket motors are

currently the focus because they have a significant opportunity for UHTCs. However,

rocket motors have a very limited life expectancy because of the extreme

3

temperatures (over 3000 °C) and highly reactive environment (dissociating materials).

Therefore, the main interesting point in these applications is not whether a material

will fail, but how long it will last before failure.

4

•I .

-3r *

"SSfe

Sb

Figure 1-1. Re-creation of the Shockwave during re-entry

5

Figure 1-2. Belly-down attitude during re-entry

6

References

[1] William G. Fahrenholtz and Gregory E. Hilmas 2004 Draft of NSF-AFOSR Joint Workshop on Future Ultra-High Temperature Materials.

[2] Sylvia Johnson, Matt Gasch, and Mairead Stackpoole, 2009 Assessment of the

state of the art of ultra high temperature ceramics, NASA technical report

ARC-E-DAA-TN486.

[3] Kaufman, L. and Clougherty, E. V. 1966 Investigation of boride compounds for

very high temperature applications RTD-TRD-N63-4096, Part III, ManLabs Inc.

[4] Clougherty, E. V, Kalish, D. and Peters, E. T. 1968 Research and development of

refractory oxidation resistant diborides AFML-TR-68-190, ManLabs Inc.

[5] Woo, S. K., Kim, C. H. and Kang, E. S. 1994 Fabrication and microstructural

evaluation of ZrB2/ZrC/Zr composites by liquid infiltration J. Mat. Sci. [2]

5309-5315.

[6] Zhang, G, Deng, Z., Kondo, N., Yang J., and Ohji, T. 2002 Reactive hot pressing

of ZrB-2-SiC composites, J. Am. Ceram. Soc. [83] 2330-2332.

7

Chapter 2

Literature survey of ultra-high-temperature ceramics

2.1 Selection of ultra-high-temperature ceramics

The several materials with melting temperatures approaching or above 3000 °C

include several oxides, nitrides, carbides, borides, and refractory metals listed in

figure 2-1 [1]. In the case of real engineering applications, high melting temperature

is not the only criterion for materials selection. Oxidation resistance, strength at room

temperature or elevated temperature, thermal conductivity, thermal expansion, density,

fabricability, and cost are more important properties to determine optimal materials.

Applications I mentioned above will involve exposure to oxidizing fuels or aero

heating so that all non-oxide materials will undergo oxidation to form some

combination of solid, liquid, and gaseous reaction product. The oxides are reasonable

to consider for use in oxidizing environments, but these materials are not suitable for

those application because of poor thermal shock resistance due to high thermal

expansion and low thermal conductivity [2]. Several borides, carbides, and nitrides of

group IV and V elements are most suitable for these applications.

All borides, carbides, and nitrides listed have similar properties, such as high

melting point, moduli, and hardness due to strong covalent bonding. However, the

borides tend to have higher thermal conductivity compared to carbides and nitrides,

8

which give them better thermal shock resistance and make them the most ideal for

many high temperature applications [3,4,5]. Among the borides in figure 2-1, ZrE$2

and HfB2 have received the most attention because their oxidation resistance is

superior to the other borides, due to the stability of ZrC>2 and HfC^ scales that form on

these materials at elevated temperatures in oxidizing environments [1]. In addition,

the combination of other refractory phases such as SiC or MoSi2 improves the

strength and oxidation resistance [6]. In this study, ZrB2-based ceramics and

ZrB2-based ceramic composites will be discussed.

2.2 Oxidation mechanism of ZrB2-SiC ceramics

All the advantages mentioned above of borides (especially ZrB2 and HfB2) are

leading many researchers to investigate various properties. Even though the oxidation

mechanism is beyond the scope of this research, it will be discussed in detail in this

chapter since oxidation resistance of these materials is very important as well as

others and the most widely studied recently.

Pure ZrB2 crystal oxidized to Z1O2 crystal and B2O3 liquid in air between

700-1100 °C. Parabolic (diffusion controlled) kinetics are observed because it is

weight gain stage due to formation of B2O3 and Zr02 and kinetics are controlled by

9

the transport of oxygen through B2O3 [7]. The oxidation rate increases para-linearly

because overall rate of mass change is a combination of weight gain due to formation

of B2O3 and ZrC>2 and weight loss due to volatilization of B2O3 between 1100-1400

°C [8,9]. Above 1400 °C, the rate of evaporation of B2O3 is greater than its rate of

production, leaving a non-protective porous ZKD2 scale, which is unfavorable for the

applications. In these temperature ranges, rapid linear kinetics were observed because

mass gain due to formation of ZrC>2 is much greater than mass loss due to

consumption of ZrB2 [10].

However, the oxidation resistance of ZrB2 was improved by the addition of SiC

due to the formation of silica-rich scales on surfaces above 1100 °C. As the volatility

of SiC>2 scales is lower than that of B2O3 scales, diffusion controlled kinetics was

observed in these temperature ranges [6, 11].

William G. Fahrenholtz, et al. proposed oxidation mechanisms of ZrB2-SiC

ceramics to several steps up to 1500 °C. The first step is the so-called unoxidized

ZrB2-SiC stage. In this stage, SiC particles are uniformly dispersed in a ZrB2 matrix

and solid solution is not expected because the solid solution limit for each component

is low, shown in the ZrB2-SiC phase diagram in figure 2-2 [12, 13]. The second step

is the initial response during heating below -1200 °C and the main oxidation

10

mechanism is the oxidation of ZrB2 with reaction (1) in this stage. The parabolic

kinetics can be observed due to hindering of oxygen diffusion through the scale

composed of B2O3, ZrC>2, and SiC particles, which do not oxidize significantly. The

third step is evolution as the temperature approaches 1500 °C. In this stage, the

composition of the scale changes significantly from B2O3, ZrC>2, and SiC to ZrC>2 and

Si02 because oxidation of SiC is the dominant oxidation with reaction (2) and B2O3

start evaporating in this temperature range. Compared to pure ZrB2 kinetics in this

temperature range, parabolic kinetics can also be achieved by the SiC^-based scale

rather than B203-based scale. The last step is steady state at 1500 °C. The most

important characteristics of this stage are the generation of SiC depleted layer

between unoxidized ZrB2-SiC layers and a Zr02+Si02 layer and SiC<2 rich layer on

the top surface. The thickness of these two layers continuously increases because the

source of Si is the active oxidation of SiC, due to the low oxygen partial pressure in

the SiC depleted layer. Figure 2-3 shows the schematic diagram of a proposed

oxidation mechanism and figure 2-4 shows the oxidized cross sectional

microstructure of ZrB2-SiC ceramics [13].

As the temperature increases, the pressure in the SiC depleted region will

increase until it is high enough to rupture the scale or cause failure at one of the

11

interfaces in the layered structure. This may occur when the total pressure in the SiC

depleted region reaches -1.013 X 105 Pa, which is estimated to occur at -1775 °C

[14].

ZrB2(cr) + 5/202(g) -»• Zr02(cr) + B203(1) Reaction (1)

SiC(cr) + 3/202(g) -> Si02(l) + CO(g) Reaction (2)

2.3 Mechanical properties of ZrB2-SiC ceramics

As ZrB2 have high melting temperature and strong covalent bond characteristics,

the densification of ZrB2 powders requires very high temperature (2100-2300 °C) and

pressure-assisted sintering procedures [15]. However, these processing conditions

normally produce coarsening of the final microstructures, which leads to the

formation of microcracks due to thermal expansion mismatch during cooling. This

leads to degradation in the mechanical properties such as hardness and strength [16].

A lot of research related to high temperature properties such as oxidation

resistance and thermal shock resistance with various processing and combination of

additives have resulted in many achievements, but large amount of work has not been

carried out to improve mechanical properties and to understand fracture mechanism

scientifically in ambient temperature or high temperature.

12

Sumin Zhu, et al. studied the influence of silicon carbide particle size on the

microstructure and mechanical properties of ZrB2-SiC ceramics. The mechanical

properties (including physical properties) are listed in table 2-1 and typical

microstructures of ZrB2-SiC ceramics with four different SiC starting powder size

from lOum through 0.45um are shown in figure 2-5. The reduction of SiC grain size

leads to an increase in relative density, hardness, and flexural strength, but no effect

with reduction of SiC grain size are found on the modulus, Poisson's ratio, and

fracture toughness [17].

Alireza Rezaie, et al. researched the effect of hot pressing time and temperature

on the microstructure and mechanical properties of ZrB2-SiC. The physical and

mechanical properties of ZrB2-SiC ceramics with two different starting powder sizes

of ZrB2 (A: 6um, B: 2um), three different hot pressing temperatures, and time are

listed in table 2-2. As the hot pressing temperature and time increase, ZrB2 and SiC

grain sizes increase, but fracture toughness and flexural strength decrease. Modulus

and hardness are not related to the hot pressing temperature and time [18].

The results from the these two recent works indicate that SiC starting powder

size and sintering temperature and time play an important role in promoting physical

and mechanical properties of ZrB2-SiC ceramics.

13

2.4 Spark plasma sintering

Spark plasma sintering (SPS) is one of the powder consolidation methods with

applied pressure and pulsed DC current simultaneously. Powders are placed in a

graphite die and heating is affected by passing through the die and sample if the

sample is conductive while a pressure is applied on the powder. The schematic SPS

process is shown in figure 2-6 [19].

Sintering is realized by subjecting the green body to arc discharges generated by

pulsed DC current. An electrical discharge process takes place on a microscopic level

and accelerates material diffusion. The basic mechanism of neck formation by spark

plasma is shown in figure 2-7. When a spark discharge appears in a gap or at the

contact point between particles, local high temperature states of several to tens of

thousands of degrees centigrade are generated momentarily. This causes evaporation

and melting on the particle surfaces, and necks are formed around the area of contact

between particles [20]. These necks gradually develop and plastic transformation

progresses during sintering, resulting in a sintered compact of over 99% density.

Since only the surface temperature of the particles rises rapidly by self-heating,

particle growth of the starting powder materials is controlled. Therefore, a precision

sintered compact is manufactured in a shorter time. At the same time, bulk fabrication

14

of particles with amorphous structure and nano-crystallization formation are now

possible without changing their characteristics [21].

Compared to conventional hot pressing methods, the most important difference

is the fast heating rate to as high as 1000 °C. Typically, the sample and the die are

heated by radiation from an enclosing furnace in the hot-pressing method. In contrast,

the die and sample are heated by Joule heating from a current passing through them if

the sample is conducting by the SPS method. However, in addition to providing heat,

the pulsed DC current has been assigned another role: that of creating plasma. The

plasma is proposed to cause a cleansing effect on the surface of the particles, leading

to sintering enhancement [19].

The advantages of the SPS method are: (a) low power consumption

(approximately one-fifth of Hot Pressing), (b) the absence of sintering aids, (c)

control of the thermal gradient (for functional graded materials (FGMs)), (d) selective

control of the density in specified regions, (e) accurate control of the porosity, (f)

single step sintering-bonding, (g) particle surface cleaning, (h) high heating rate and

(i) near-net-shape capability. The short sintering time is particularly suitable for: (a)

preserving initial powder grain size or nanostructure, (b) consolidating amorphous

materials, (c) improving bonding strength between particles and (d) controlling phase

15

reactions or decomposition (in the case of composites) [22].

Many improvements of mechanical and physical properties of materials by SPS

method were reported from the researchers who investigated ultra-high temperature

ceramics. Therefore, SPS is the most promising consolidation method for

ultra-high-temperature ceramics which require high temperature with fast heating rate

and pressure to maintain fine microstructure.

2.5 Liquid phase sintering

Liquid phase sintering (LPS) is a sintering process that involves liquid and solid

particles. The liquid can be present during the entire sintering process or during a

specific stage of the sintering cycle. They are two basic methods to obtain liquid

phase: (i) Using mixed powders of differing characteristics; and (ii) Taking advantage

of the interaction between two or more components or from the formation of a

eutectic liquid [23].

According to Kingery's model, the liquid phase sintering process has three

different steps that may in certain cases be partly overlapped. Figure 2=8 shows the

three steps of liquid phase sintering process by Kingery, et al. [24].

The first step, rearrangement, takes place right after liquid phase forms with

16

rapid partial densiflcation due to the capillary force exerted by the wetting liquid on

the surface of solid particles. The elimination of porosity takes place as the system

minimizes its surface energy. The most favorable features for rearrangement are: (i) a

homogeneous distribution of a congruently melting liquid that wets the particles

present; (ii) the solid particles are soluble in the liquid; (iii) a large solubility ratio.

The rate and extent of shrinkage depend upon the viscosity and quantity of liquid

phase formed and on its wetting properties [25].

The second step, the so-called solution precipitation step, occurs due to a

difference in solubility. This establishes a concentration gradient in liquid phase,

because small grains are more soluble that large grains. The atoms can be transported

from small grains to large grains by diffusion, so that the large grains grow at the

expense of the small grains. This process is coarsening or ripening. The two driving

forces of this step are the reduction of interfacial energy and the capillary force. This

process relies strongly on two critical steps: dissolution of solid into the liquid and

diffusion through the liquid. If the transport is limited by mass transfer from the

source to the sink, this process is diffusion-controlled; if the transport is confined by

interfacial dissolution or precipitation, the process is reaction-controlled [26-28].

The final step is microstructure coarsening and solid-state sintering. A solid

17

skeleton is formed, and grain coalescence occurs. The rigidity of the solid skeleton

hinders further rearrangement, although microstructure coarsening continues by

diffusion. In this case, the residual pores will enlarge if they contain entrapped gas

that results in swelling. The densification rate is greatly decreased during this part of

the sintering cycle [29].

18

Matei isil

ZS10

ZS1.-1

ZSO.-!

ZS0.45

Relativedensit} (%)

9".4

98.9

98.T

99.8

Avei age SiC gi nin size (Jim)

6.3±2 9

2.1±0."

1.6±0.7

1.0±0.4

V

0 16

0.15

0.15

0.16

E (GPvl)

4 "9±5

509±3

515±7

5 2 0 + -

HV0.2 (GPa)

1".5±0.4

19.1±1.0

19.3±0.6

20.7±1.0

KK (MPa m1'2)

4.5±0.1

4.3±0.3

4.2±0.2

4.6±0.1

<r(MPa)

389±45

805±"1

83"±116

909±136

Table 2-1. Physical and mechanical properties with reducing SiC particle size of

ZrB2-SiC ceramics [17] (10, 1.4, 0.7, 0.45 in sample names are starting SiC particle

size of micron).

19

Sample

A-1850-45

A-1950-45

A-2050-45

B-1850-45

B-1950-45

B-2050-45

B-2050-90

B-2050-180

AveZrB :

grain size (um)

2.1 ± 1.3

3.3 ± 1.5

3.7 ± 1.5

2.2 ± 1.2

2.5 ± 1.8

3.5 ± 2.0

3.8 ± 2.0

4.7 ± 3.0

Ave SIC grain size (um)

1.5 ± 0.7

2.5 ± 1.1

3.1 ± 1.3

1.2 ± 0.6

1.7 ± 0.8

2.0 ± 0.8

2.0 ± 0.9

2.7 ± 1.0

E (GPa)

503 ± 6

501 ± 1

503 ± 1

516+3

507 ± 3

505 + 2

508 ± 4

505 ± 1

H (GPa)

22 ± 2

22 ± 2

23 ± 2

20 ± 2

22 ± 2

23 ± 1

22 ± 1

22 ± 1

K IC(MPa.m1:)

3.9 ± 0.1

4.0 ± 0.2

4.3 ± 0.2

5.5 ± 0.3

5.2 ± 0.4

4.3 ± 0.2

4.2 ± 0.1

4.5 ± 0.2

<r(MPa)

888 ± 151

770 ± 133

720 ± 38

1063 ± 91

1060 ± 59

854 ± 88

850 ± 100

804 ± 73

Table 2-2. Physical and mechanical properties of ZrB2-SiC ceramics with different

hot pressing time and temperature [18] (A, B: ZrB2 starting powder size; A: 6um, B:

2(4.m, 1850, 1950, and 2050: temperature; °C, 45, 90, and 180: time; min)

20

Metals Oxides Borldes Carbides Nitrides

Material Family

Figure 2-1. Several materials with melting temperatures approaching or above 3000

°C[1]

21

3200

3000

"I 1 T 1 r

s3050° \

\ \ Liquid

2800 r

.2600

2400

2200

2000

L

2760°/-

I H / I

\ \

\ \ \ ' I

2270° \ / i-

J

J L__J L

(77%) M

I I H

j I i L _ J 1 0

2rB2

20 40 60

Mol % 100 SiC

Figure 2-2. ZrB2-SiC phase diagram [12]

22

> 2rB2 + SIC

ZK>2 + B2O3 + SiC

ZrB rSiC

9EL

'-1* ;v*™ • . ? v - * * v 7"; •7"r*,'~,|~

MrMMOT itei«MMfr Af t 11a

^-ZrO^SlOa

> ZrB rSiC

S i0 2

ZrO2 + S I0 2

ZrB2 (SiC-depIeted)

ZrB2 + SiC

Figure 2-3. The schematic diagram of a proposed oxidation mechanism for ZrB2-SiC

ceramics, (a) unoxidized ZrB2-SiC (b) the initial response during heating (c)

evolution as the temperature approaches 1500 °C (d) steady state at 1500 °C [13].

23

SiO;

Depleted of SiC

• ZrB2 -SiC

Figure 2-4. The cross-sectional microstructure of ZrB2-SiC ceramics [13]

24

Figure 2-5. Microstructure of ZrE$2-SiC ceramics [17]

(a) ZSIO (b) ZS1.4 (c) ZS0.7 (d) ZS0.45

25

Pressure control

Power Supply

Figure 2-6. Schematic of the SPS process [19]

26

Applied Pressure

Particle

Electron Flow Generation of Spark Impact Diffusion Bonding Plastic Deformation

Particle

Applied Pressure

Figure 2-7. Basic mechanism of spark plasma sintering [20].

27

Figure 2-8. Three different steps of liquid phase sintering by Kingery

(I) Rearrangement (II) Solution-precipitation (III) Microstructure coarsening and

solid state sintering.

28

References

[1] William G. Fahrenholtz, Gregory E. Hilmas, Inna G. Talmy, and James A.

Zaykoski (2007) Refractory diborides of zirconium and hafnium J. Am. Ceram. Soc,

Vol.90, Issue 5, pl347-1364

[2] E. Wuchina, E.Opila, M.Opeka, W. Fahrenholtz, and I. Talmy (2007) UHTCs:

Ultra-high temperature ceramic materials for extreme environment applications, The

electrochemical society Interface, p30-36

[3] Courtright, E. L., Graham, H. C., Katz, A. P. and Kerans, R. J. (1992) Ultra high

temperature assessment study- ceramic matrix composites AFWAL-TR-91-4061

Wright Patterson Air Force Base Ohio

[4] Culter, R. A. (1991) Engineering properties of borides ASTM Engineered

materials handbook, Vol 4-Ceramics and glasses, Schneider, S. J., Technical

chairman, p787-803

[5] Guillermet, A. F. and Grimvall, G (1991) Phase stability properties of transition

metal diborides Am. Inst. Phy. Conf. Proa, [231] p423-431

[6] W. C. Tripp, H. H. Davis, and H. C. Graham (1973) Effect of an SiC addition of

the oxidation of ZrB2, Ceramic Bulletin 52(8) p612-616

[7] R. J. Irving and I. G. Worsley (1968) Oxidation of titanium diboride and

zirconium diboride at high temperatures, J. Less-Common Metals, 16 [2] pl03-l 12.

[8] W. C. Tripp and H. C. graham (1971) Thermogravimetric study of the oxidation

of ZrB2 in the temperature range of 800-1500 °C, J. Electrochem. Soc, 118 [7]

pi 195-1199.

[9] F. Monteverde, A. Bellosi (2003) Oxidation of ZrB2-based ceramics in dry air J.

electrochem. Soc, 150 [11] p552-559.

[10] A. L. Chamberlain, W. G. Fahrenholtz, and G. E. Hilmas, Oxidation of ZrB2-SiC

ceramics under atmospheric and reentry conditions Refractory Appl. Trans., 1 [2]

29

pl-8.

[11] H. C. Graham, H. H. Davis, I. A. Kvernes, and W. C. Tripp (1971)

Microstructural features of oxide scales formed on zirconium diboride materials

pp35-48 in Ceramics in severe environments: Materials Science Research.

[12] S. S. Ordanyan, A. I. Dmitriev, and E. S. Moroshkina, (1989) Izv. Akad. Nauk

ASSR, Neorg. Mater.,25 [10] pi752-1755; Inorg. Mater. (Engl Transl.), 25 [10]

pl487-1489.

[13] W. G. Fahrenholtz, (2007) Thermodynamic analysis of ZrB2-SiC oxidation:

Formation of a SiC-depleted region, J. Am. Ceram. Soc, 90 [1] pl43-148

[14] M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, (2004) Oxidation-based materials

selection for 2000 °C+ hypersonic aerosurfaces: Theoretical considerations and

historical experience, J. Mater. Sci. 39 [19] p5887-5904.

[15] K. Upadhya, J. M. Yang, and W. P. Hoffman, (1997) Materials for ultra high

temperature structural applications, Am. Ceram. Soc. Bull., 76 [12] p51-56.

[16] C. Mroz, (1995), Titanium diboride, Am. Ceram. Soc. Bull., 74 [6] pl58-159.

[17] Sumin Zhu, W. G. Fahrenholtz, and G. E. Hilmas, (2007) Influence of silicon

carbide particle size on the microstructure and mechanical properties of zirconium

diboride-silicon carbide ceramics, J. Euro. Ceram. Soc, 27 p2077-2083.

[18] Alireza Rezaie, W. G. Fahrenholtz, and G. E. Hilmas, (2007) Effect of hot

pressing time and temperature on the microstructure and mechanical properties of

ZrB2-SiC, J. Mater. Sci., 42 p2735-2744.

[19] Z. A. Munir, U. Anselmi-Tamburini, and M. Ohyanagi, (2006) The effect of

electric field and pressure on the synthesis and consolidation of materials: A review

of the spark plasma sintering method, J. Mater. Sci., 41, p763-777.

[20] K. Lu, (2008), Sintering of nanoceramics, International Materials Review, 53 [1]

p21-38

30

[21 ] http ://www. scm-sps.com/e_htm/whatsps_e_htm/whatsps4_e.htm

[22] Salvatore Grasso, Yoshio Sakka, and Giovanni Maizza (2009) Electrical current

activated/assisted sintering (ECAS): a review of patents 1906-2008, Sci. Technol.

Adv. Mater. 10, 053001, pi-24

[23] R. M. German, (1985) Liquid phase sintering, Plenum press, New York and

London

[24] W. D. Kingery (1959) Densification during sintering in the presence of a liquid

phase. I. Theory., J. App. Phy 30 p301-306

[25] W. J. Huppmann and H. Riegger (1975) Modeling of rearrangement processes in

liquid phase sintering, Acta Metallurgica, Vol23, p965-971.

[26] I. M. Lifshitz, and V. V. Slyozov (1961) The kinetics of precipitation from

supersaturated solid solutions, J. Phy. Chem. Soli., 19 p35-50

[27] C. Wagner, (1961) Theory of precipitate change by redissolution, J. electrochem.,

65p581-591.

[28] A. J. Ardell (1971) Effect of volume fraction on particle coarsening. Theoretical

consideration. Acta metallurgica 20 p61-71

[29] W. D. Kingery, H. K. Bowen, and D. R. Uhlmann (1976), Introduction to

ceramics, Wiley Interscience, 2nd edition, Chapter 10.

31

http://scm-sps.com/e_htm/whatsps_e_htm/whatsps4_e.htm

Chapter 3

ZrB2-nano SiC ceramics

Abstract

The microstructure and mechanical properties of ZrB2 composites with

nano-sized SiC particles were examined. The composites were consolidated by spark

plasma sintering. The microstructure of the as-consolidated composites was examined

by field emission scanning electron microscopy and transmission electron microscopy.

The elastic constants of the composites were measured using an ultrasonic technique,

whereas the hardness and fracture toughness of the composites were determined using

an indentation measurement. The presence of intergranular and intragranular SiC

improves the sinterability and limits the grain growth of ZrB2 effectively until 10

vol% of nano-sized SiC particles. However, the grain size of SiC increased with an

increase of nano-sized SiC content due to agglomeration by the mixing process. In

addition, each phase of ZrB2 and SiC were confirmed by SAED (Selected Area

Electron Diffraction) and a second phase, Zr(0, B)x, was found due to a large amount

of heat during spark plasma sintering. The hardness of composites was increased with

an increase of nano-sized SiC content, but the fracture toughness was not related to

the content of SiC. The applied stress to micropillars of ZrB2 nano-sized SiC

32

composites produced many dislocations and stacking faults.

33

3.1 Introduction

Zirconium diboride and hafnium diboride have extremely high melting

temperatures (>3000 °C), high chemical stability, high electrical and thermal

conductivities, resistance to erosion/corrosion, and good mechanical properties which

make them suitable for ultra-high-temperature applications such as hypersonic flight

vehicle, atmospheric re-entry vehicle, and rocket propulsion systems [1]. However,

use of those single phase ceramics materials for these applications has limitations

because a passive layer composed of B203 start evaporating above 1000 °C, leaving

behind a non-protective porous Zr02 layer [2].

To overcome the poor oxidation resistance of these diborides, the second phase

such as SiC and MoSi2 were introduced due to the generation of a borosilicate

diffusion barrier layer above 1000 °C. In addition, the introduction of these second

phases improves the mechanical properties due to a liquid phase sintering effect [3,

4].

Among the second phases, SiC is the most widely studied by many researchers.

Sumin Zhu, et al. reported the influence of SiC initial particle size, in the range from

0.45urn to 10um, on the microstructure and mechanical properties of ZrB2-SiC

ceramics. The smaller SiC particle sizes led to improved densification, finer grain

34

sizes, higher hardness, and higher strength [5].

However, the introduction of nano-sized SiC starting powder into ZrE$2 ceramics

was not studied widely yet by researchers until present. Niihara, et al. proposed three

different types of ceramic nanocomposites shown in figure 3-1 [6]. Firstly,

intergranular nanocomposites can improve creep resistance due to grain boundary

pinning effect. Secondly, intragranular nanocomposites can improve transgranular

fracture toughness and strength. Thirdly, inter/intragranular nanocomposites can take

both effects mentioned above. Lastly, nano/nano composites can have superplasticity

due to grain boundary sliding [7].

In this study, nano-sized SiC particles were used to improve the mechanical and

physical properties of ZrB2-SiC ceramics and spark plasma sintering was used to take

an effect of nanocomposites as mentioned above.

3.2 Experimental procedure

3.2.1 Materials

The starting powders used in this study were: ZrB2 powder (Grade F, Japan New

Metals, Tokyo), average particle size -2.1 um and nano p-SiC powder (Sumitomo

Osaka Cement Co. Ltd., Osaka, Japan), average particle size ~30 nm and oxygen

35

content of 0.38 wt.%. In order to examine the effects of nano-particle content as well

as starting particle size, four batches of powder were prepared, containing nano (3-SiC

powder of 5, 10, 20 and 30 vol.%, respectively. In addition, the single-phase ZrB2

powder was prepared to compare the effect of SiC addition on mechanical properties.

Compositions of the ZrE$2-based composites with SiC are listed in Table 3-1. The

powder mixtures were ball-milled using SiC milling media and ethanol under 200

rpm for 24 h, and the resulting slurry was then dried under magnetic stirring to avoid

sedimentation. Before sintering, the dried mixtures were sieved through a metallic

sieve with 60-mesh screen size.

The powder mixture was put into graphite die lines with graphite foil and

densified using spark plasma sintering (SPS-1030, Sumitomo Coal Mining Co. Ltd.,

Tokyo, Japan). The sintering was performed at 1900 °C for 3 min under an external

pressure of 50MPa in an argon atmosphere. The temperature of the sample was

automatically raised to 600"C, and then was monitored by an optical pyrometer

through a hole opened in the die and automatically regulated to the final sintering

temperature with heating rate of ~300°C/min. The pressure was applied at room

temperature and held constant until the end of the sintering cycle. The load was

removed when the die temperature dropped below 1000°C with a cooling rate

36

of~ 600 °C/min. The final sintered specimen size was 10mm in diameter pellets with

a thickness of ~2.0 mm.

The density of the sintered composite compacts was measured by Archimedes

method with distilled water as a medium. The sintered composite pellets were then

polished with a diamond paste up to 0.5um. The morphology of the composites was

characterized by field emission scanning electron microscopy (FE-SEM, ZEISS

VP1550), and the crystalline phases were identified by X-ray diffractometry (XRD,

Panalytical X'Pert Pro). The grain size of each component, ZrB2 and SiC, was

measured directly from the FE-SEM images of sintered composites using image J

software. In order to estimate the average grain size of each component, five FE-SEM

images were used for direct measurement within the range of 200-300 population.

To investigate detailed microstructural information of sintered composites with

30nm SiC powders, randomly selected areas of ZSN-0, ZSN-5, ZSN-10, and ZSN-20

samples were observed by TEM (FEI-Philips CM300, FBI, Hillsboro, OR) with a

EDAX Energy Dispersive X-ray spectrometer (EDS) and a Si/Li detector super

ultrathln window. The TEM samples used in this study were prepared by Focused Ion

Beam (FIB Nova 600, FEI, Hillsboro, OR) with tungsten probe tips and four fingers

copper grid.

37

3.2.2 Elastic property measurements

The elastic moduli measurements of the composites were performed using an

ultrasonic technique (TDS 3052B, Tektronix Inc., Beaverton, OR USA) with a

fundamental frequency of 20 MHz. Young's modulus (E), shear modulus (G) and

Poisson's ratio (v) were calculated using the longitudinal and transverse sound wave

velocities measured in the composite specimens. The details of calculations are listed

below.

E = pV2 r3V2-4V2^

, V} - v2 , V v< v' J

(1)

G = pV,2 (2)

v = — - 1 (3) 2G

where p is the true density, V, and V, are the longitudinal and transverse sound

wave velocities, respectively. V, and V, are determined by

r,=£ (4) At,

At,

where h is the specimen thickness, At, and A/, are the elapsed times between

the pulse and the echo of the longitudinal and transverse waves, respectively [8]. The

accuracy of the soundwave velocity measurement was better than 1%.

3.2.3 Mechanical properties

38

The hardness and fracture toughness, K\c, of the composites were determined

using an indentation technique. The indentation tests were performed on the polished

surface of the specimens by loading with a Vickers microhardness indenter (AVK-A,

Akashi Co. Ltd., Yokohama, Japan) for 20 s in ambient air at room temperature. The

corresponding diagonals of the indentation, a, and crack sizes, c, were measured

using an optical microscope attached to an indenter. The indentation load of 49N was

used, and 10 indents were made for each measurement. For fracture toughness

calculations, the Palmqvist equation was used in this study becuase the ratio of l/a

were in the range of 0.25-2.5 from crack length measurement, (where l=c-a). The

Palmqvist equation is listed below:

KIC = r,(E/H)2/5P/(al m) (6)

where E is the Young's modulus (GPa), H is the hardness (GPa), P is the

indentation load (N), /is the crack length (|J.m), a is the half diagonal length (um), and

c is the total length including crack length and half diagonal length (pm). The

hardness, Hv, was calculated from

H = 1854.4-^=-, (7)

In the earlier work from Niihara, et al., the coefficient, r\, is obtained

experimentally. The value of r\ was 0.0089 and 0.0122 for l/a ratio varying in the

39

range of 0.25-2.5 and 1-2.5, respectively [9]. We have chosen r|=0.0122 because all

ratios of l/a were in the range of 1-2.5.

3.2.4 Micropillar compression test

To investigate deformation or fracture mechanism of ZrB2-SiC composites,

micropillars with ~5um diameter and ~20um height were prepared by Focused Ion

Beam (FIB Nova 600, FEI, Hillsboro, OR) on the surface of composites. The

obtained micropillars of each sample (ZSN-0, ZSN-5, and ZSN-10) were compressed

by the depth of 500nm from the sample surface with the flat punch tip of a

nanoindenter (MTS Nanoindenter/XP, MTS system, Eden Prairie MN). Applied stress

and displacement of samples were recorded to identify events such as fracture and

deformation. After compression tests, TEM laminar samples were prepared by the

Focused Ion Beam if the critical dimensional changes of micropillars were not found.

Microstructural information of micropillars from the prepared TEM samples was

investigated.

3.3 Results and discussion

3.3.1 Densities, phases, and microstructures

The measured densities and relative densities for various ZrB2-nano SiC

40

composites are summarized in table 3-2. It is evident that single-phase ZrB2 has the

lowest density (-96.1%), and densities of each sample slightly increase with increase

of SiC content from 96.1% to 98%. The low relative density of pure ZrB2 is due to its

strong covalent bond and low self diffusion [10]. As the SiC content of each

composite material increases, SiC improves sinterbility of composites more

effectively due to the liquid phase sintering effect and hindrance of ZrB2 grain

growth.

X-ray diffraction (XRD) patterns of each composite materials consolidated by

SPS are shown in figure 3-2. Only ZrB2 peaks are detected in ZSN-0. ZrB2 and SiC

peaks were detected in the other four samples (ZSN-5-ZSN-30). The intensity of SiC

peaks are increased with increasing SiC content in the composite material. Any other

second phases were not found and both ZrB2 and SiC phases were confirmed by XRD

in all samples.

Typical microstructural features of the ZrB2-SiC composites observed under

FE-SEM are shown in figure 3-3. The general microstructures of the ZrB2-SiC

composites are similar, consisting of the equiaxed ZrB2 (grey contrast) and SiC (dark

contrast) grains. The average grain sizes of the ZrB2 and SiC are also listed in table

3-3 and the typical tendency of average grain size of each component is shown in

41

figure 3-4. The average grain size of pure ZrB2 is -4.02 urn, which is significantly

higher than the starting powder (2.1 urn). Apparently, the grain growth of ZSN-0

occurred during spark plasma sintering. However, the average grain size of ZrE$2

decreases with increasing SiC content until 10 vol% SiC composite materials,

proving that SiC hinders ZrB2 grain growth. The average grain size of ZrB2 was not

decreased further after 10 vol% SiC content. It is due to the fact that the average grain

sizes of ZrB2 of ZSN-10, ZSN-20, and ZSN-30 are very similar and the average grain

sizes of ZrB2 of those samples which originated from - 2.1 um was measured to be

-1.7 um after ball milling and sintering. In addition, the average grain size of SiC

also increased with increasing SiC content from 1.07 to 1.60 um, which is

substantially higher than that of the starting powder (-30 nm). The mixing process

employed was not effective in breaking apart the agglomeration of nano-sized SiC

particles. As a result, clusters of nano-sized SiC particles were fused together during

spark plasma sintering to form SiC grains with a few microns in diameter. A more

effective mixing/dispersion technique to break apart the agglomeration of nano-sized

SiC particles needs to be developed to uniformly disperse the nano particles.

More detailed microstructural images by TEM and selected area electron

diffraction (SAED) of each sample (ZSN-0-ZSN-30) are shown in figure 3-5. Phases

42

and crystallographic directions of several grains in these images were confirmed by

EDAX and electron diffraction patterns. ZrB2 grains are dark contrast and SiC grains

are grey contrast in all images (opposite to FE-SEM images). At first, a small SiC

grain could be found in pure ZrB2 ceramic composite (ZSN-0); this SiC grain might

come from SiC milling media. The size of SiC grains increase from a few tens of

nanometers to several micrometers. Most SiC grains are of one crystallographic

orientation and a few SiC grains are of several crystallographic orientations in one

grain, which suggests that nano-sized SiC particles were fused together during spark

plasma sintering. Most SiC grains are located in interfaces between ZrB2 grains and

SiC grains or grainboundaries between ZrB2 grains, but a few SiC grains with the size

of several nanometers (30nm ~ 80nm) are found in ZrB2 grains in ZSN-5 and ZSN-30

(shown in figure 3-6). Typical interfaces between ZrB2 grains and SiC grains are

shown in figure 3-7. The thicknesses of interfaces are 1~2 nm, so those inter-phase

interfaces between two components are well bonded to each other. In addition, a

second phase, Zr(0,B)x, was found in ZSN-20 because large amount of heat during

spark plasma sintering might be localized in ZrB2 grain (shown in figure 3-8).

3.3.2 Elastic moduli, hardness and fracture toughness

The elastic constants measured in various ZrB2-SiC composites are listed in

43

table 3-4. The Young's modulus of single-phase ZrB2 was measured to be -499 GPa,

which is slightly higher than that reported in the literature (489 GPa) [11]. The results

indicated that Young's moduli of the ZrB2 composites with nano-sized SiC are not

sensitive to the compositions. Among the five compositions, ZSN-5 and ZSN-10 have

the best elastic properties and ZSN-20 and ZSN-30 have the worst properties because

the Young's modulus of silicon carbide is reported to be 415 GPa [12], which is lower

than that of zirconium diboride and silicon carbide agglomerate may cause

detrimental effect to elastic properties of composites materials.

Hardness and fracture toughness of ZrB2-SiC composites obtained from the

indentation technique are also summarized in table 3-4. The hardness of single-phase

ZrB2 was measured to be ~13 GPa, which is lower than the results in the literature (22

GPa)[13]. The porosity of the ZSN-0 might not be a main reason for the low hardness

value because the porosity value is -4%. It is the reason that low bonding strength

between ZrB2 grains to maintain coherent microstructure due to inherent ZrB2

characteristics such as strong covalent bond and low self-diffusion. However, the

hardness in the ZrB2-nano SiC ceramic increased with an increase of SiC content

(shown in figure 3-9). This effect is based on the basic rule of mixtures. The samples

with higher SiC content have the higher value of hardness because the hardness of

44

SiC (32 GPa) [12] is higher than that of ZrB2 (22 GPa) [13]. The fracture toughness of

all ZrB2-nano SiC composite were in the range of 6.8-7.5 MPa.m172 and these values

cannot be compared to others because many studies for fracture toughness of

ultra-high-temperature ceramics were conducted by methods. The results also

indicated that fracture toughness of the ZrB2 composites with nano-sized SiC is not

sensitive to the compositions shown in figure 3-9. A typical cracking pattern is shown

in figure 3-10. For the single-phase ZrB2 ceramic, the crack propagated primarily

along ZrB2 phase boundaries but the crack in the ZrB2 nano-SiC composite

propagated across the ZrB2 and SiC grains without being deflected along the grain

boundaries of the SiC grains. As a result, the incorporation of nano-sized SiC did not

impart additional resistance to crack propagation.

3.3.3 Micropillar compression test of ZrB2-nano SiC composite materials

Fabricated micropillars of ZSN-0, ZSN-5, and ZSN-10 are shown in figure 3-11.

The dimensions of micropillars in this figure vary due to different magnification in

order to show the configuration of micropillars. The stress and strain (%)

characteristics of ZSN-0, ZSN-5, and ZSN-10 are shown in figure 3-12, 3-13, and

3-14, respectively. The study of micropillar or nanopillar compression tests about

single crystal metal such as Au, Cu, Ni, and even Si had similar results to the

45

universal scaling law, which can be explained briefly that smaller micropillars or

nanopillars have higher stresses [14, 15]. However, the maximum stress values of

micropillar compression tests of pure ZrE$2 and ZrB2-SiC ceramics did not follow the

universal scaling law. The maximum stress of pure ZrE$2 micropillar was 15.5 GPa,

which is slightly higher than the normal microhardness test (~13GPa). And, the

micropillar of ZSN-5 was broken after the stress of ZSN-5 reached only -2.74 GPa. It

is interesting to note that the maximum stress of ZSN-10 was fairly low (-0.33 GPa)

because the flat punch tip might be improper contact with the top surface of the

micropillar so that it is slightly tilted in one direction (indicated with arrow in figure

3-15). The TEM images of ZSN-0 and ZSN-10 and the fracture surface of FE-SEM

images of ZSN-5 are shown in figure 3-16, 3-17, and 3-18, respectively. Several

small dislocations and stacking faults which go through whole lamella were found in

the ZSN-0 sample. The dislocations might be produced during the milling process

because they are very small and located inside of one grain, but stacking faults went

through whole lamella so that they were produced in compression stress. Also,

several discrete regions are found in stress-strain curve in figure 3-12. This

stress-strain curve might be explained for production of stacking fault in ZSN-0. In

the ZSN-10 TEM lamella, lots of dislocations and small stacking faults were found

46

compared to ZSN-0 lamella. These two defects might also be produced by

compression stresses because the stress-strain curve shape is an inward curve;

perhaps some part of energy from the applied stress was used for the generation of

dislocations and stacking faults. The fracture mode observed from ZSN-5 micropillar

was mixed inter/intragranular fracture, which is the same result as the fracture surface

reported [16].

3.4 Conclusion

The processing and mechanical behaviors of single-phase ZrE$2 and ZrB2

with nano-sized SiC particles were investigated. The mixing process employed was

not effective in breaking apart the agglomeration of nano-sized SiC particles. As a

result, clusters of nano-sized SiC particles were fused together during sintering to

form SiC particles sub-micron in diameter. These sub-micron-sized particles were

distributed primarily along the grain boundaries. However, some nano SiC particles

were embedded within the grains. The presence of intergranular and intragranular SiC

improves the sinterability and limits the grain growth of ZrB2. Further improvement

in properties may be achievable through a more uniform dispersion of SiC nano

particles. The hardness of ZrB2 with nano-sized SiC was increased with increase of

47

SiC content, but the fracture toughness of composites was not sensitive to

composition. The phase of ZrB2 and SiC in each composite was confirmed SAED and

the second phase, Zr(0, B)x, was found in ZrB2 grain due to localized large amounts

of heat during spark plasma sintering by the TEM investigation. The applied stress to

micropillars of ZrB2 nano-sized SiC composites produced lots of dislocations and

stacking faults.

48

Samples

ZSN-0

ZSN-5

ZSN-10

ZSN-20

ZSN-30

Compositions (vol.%)

ZrB2

100

95

90

80

70

SiC

0

5

10

20

30

Table 3-1. Compositions of each sample

49

Samples

ZSN-0

ZSN-5

ZSN-10

ZSN-20

ZSN-30

Theoretical Density 3

(g/cm )

6.09

5.95

5.80

5.52

5.23

Measured Density 3

(g/cm )

5.85

5.78

5.66

5.37

5.13

Relative Density

(%TD)

96.1

97.1

97.5

97.4

98.0

Table 3-2. Densities of each sample

50

Materials

ZSN-0

ZSN-5

ZSN-10

ZSN-20

ZSN-30

Grain size of ZrB2 (urn)

4.02±1.31

2.86±0.77

1.71±0.56

1.76±0.63

1.64±0.7

Grain size of SiC (u.m)

N/A

1.07±0.84

1.17±0.82

1.39±0.82

1.60±1.07

Table 3-3. Average grain sizes of ZrB2-nano SiC composite materials.

51

Samples

ZSN-0

ZSN-5

ZSN-10

ZSN-20

ZSN-30

Elastic Properties

G (GPa)

219

225

223

200

200

E (GPa)

499

517

513

455

455

V

0.14

0.15

0.15

0.14

0.14

Table 3-4. Elastic properties of ZrB2-nano SiC composite materials

52

Sample

ZSN-0

ZSN-5

ZSN-10

ZSN-20

ZSN-30

Hardness

(GPa)

13±0.59

16.64 ±0.79

17.72 ±0.71

17.63 ±0.65

20.30 ±1

Toughness

(MPa.m1/2)

7.19±0.52

7.47±0.57

7.25±0.6

6.78±0.24

7.36±0.52

Table 3-4. Hardness and fracture toughness of ZrB2-nano SiC composites materials

53

\ / •

« V — / # / • y .

^ \ • / '

^ — < • •

,« / • \ a

(a) Inter type (b) Intra type

(c) Intra/inter type (d) Nano/nano type

Figure 3-1.Classification of ceramic nanocomposites.

54

7000

6000

5000

T H 4000

CO a & 3000

" H 2000

1000

0

'aufcMi MMMMM tabnn (e),

Uw*»

WW

*mm*

• 2B 2

I JL

•I

An * >

MKM

M »

•Mn

, f t t I M M W ,

•At M M

U M

IIJLII

I I A

mmmJjkm

L J L.

10 2 0 3 0 4 0 5 0 6 0 7 0 8 0 9 0 100

29(cfegTee)

igure 3-2. X-ray diffraction patterns of each composite material (a) ZSN-0 (b)

ZSN-5 (c) ZSN-10 (d) ZSN-20 (e) ZSN-30

55

Figure 3-3. Typical FE-SEM images of each sample.

56

5

^—^

E 4

E ^^^ 0 3 N

'55 C 2

2 O 1

n

~

-1

-

-#-ac

•̂

(

i

\ - T

J • •

• •

i i i

0 10 20 30

vu %of ac

Figure 3-4. The change of average grain size of each component.

57

Figure 3-5. Typical TEM microstructures of each sample.

58

Figure 3-6. Typical SiC grain locations in ZrB2-SiC composite

59

TJF" -*" •<

Figure 3-7. Typical interface between ZrB2 grain and SiC grain (ZSN-5: HRTEM

image)

60

Figure 3-8. The TEM image and EDX result of second phase, Zr(0, B)x, in ZrB2

grain, ZSN-2(

61

CO a.

30

25

20

Q)

(0

10

5 -

0

Hardness Toughness

j _ _L J _ -1_

CM

Q_

(I) (/)

3 O

ZSN-0 ZSN-5 ZSN-10 ZSN-20 ZSN-30

Sample

Figure 3-9. The change of hardness and fracture toughness of ZrB2-nano SiC

composite with SiC content.

62

Bt * • .•¥ -, -•:."•-** v- 4-umV

Figure 3-10. Typical cracking pattern of ZrB2-SiC composites (ZSN-0, ZSN-10,

ZSN-20).

63

• 'f ?«• ' * -"

SI •iif^iBJnai.ii7iiiiMiit:fffii

Figure 3-11 Fabricated micropillar of ZSN-0, ZSN-5, and ZSN-10

64

16

14

12

(GP

a)

00

o

V) 6

2

0

.

_

• ZSN-0

/ f

Jr /

^r j

i i i i i

0 1 2 3 4

AL/LO (%)


ZSN-0

65

3.0

2.5 •*™*>s

CO Q- 2.0 CD N ^ ^

<*> 1 5

CD s_

CO 1 0

0.5

0.0

• ZSN-5

- r

-

i i i i

•

M

/

f • I

10 20 30 40

AL/L(%)

50


ZSN-5

66

0.35

J 1 I I L

0.0 0.5 1.0 1.5 2.0 2.5

AL/L (%)


ZSN-10

67

0.30

0.25 s-^.

^ 0.20

CD ^ 0.15 co CO |D 0.10 !_ - t—>

CO 0.05

0.00

-

-

-

Figure 3-15. ZSN-10 micropillar (a) before and (b) after compression

68


test (a) Whole TEM laminar (b) Stacking fault (c) dislocation

69


test (a) Whole TEM laminar (b) Dislocations (c) Stacking fault

70

Figure 3-18. Fracture surface images of ZSN-5 by the micropillar compression test

(a) Fracture surface of whole micropillar (b) Intragranular fracture (c) Intergranular

fracture

71

References



Vol.90, Issue 5, pi347-1364

[2] A. L. Chamberlain, W. G. Fahrenholtz, and G. E. Hilmas, Oxidation of ZrB2-SiC

ceramics under atmospheric and reentry conditions Refractory Appl. Trans., 1 [2]

pl-8.

[3] Jiecai Han, Ping Hu, Xinghong Zhang, and Songhe Meng, (2007) Oxidation

behavior of zirconium diboride-silicon carbide at 1800 °C, Scripta Materialia 57,

p825-828.

[4] Ronald Loehman, Erica Corral, Hans Peter Dumm, Paul Kotula, and Raj an

Tandon, (2006) Ultra high temperature ceramics for hypersonic vehicle applications,

SAND 2006-2925

[5] Sumin Zhu, W. G. Fahrenholtz, and G. E. Hilmas, (2007) Influence of silicon

carbide particle size on the microstructure and mechanical properties of zirconium

diboride-silicon carbide ceramics, J. Euro. Ceram. Soc, 27 p2077-2083.

[6] Koichi Niihara (1991) New design concept of structural ceramics; ceramic

nanocomposites, The ceramic society of Japan, 99, [10], p974-982

[7] A. Mukhopadhyay and B. Basu, (2007), Consolidation-microstructure-property

relations in bulk nanoceramics and ceramic nanocomposites: a review, International

materials reviews, Vol52, No5, p257-288.

[8] O. Yeheskel, and O. Tevet, (1999), Elastic Moduli of Transparent Yttria, J. Am.

Ceram. Soc, 82, pl36-144

[9] K. Niihara, R. Morena, and D. P. H. Hasselman, (1982), Evaluation of KiC of

brittle solids by the indentation method with low crack-to-indent ratios, J. Mater. Sci.

Lett., I , p l 3 4 6 .

72

[10] A. Bellosi, F. Monteverde, D. D. Fabbriche, and C. Melandri, (2000),

Microstructure and Mechanical Properties of ZrB2-based Ceramics, J. Mater. Proc.

Man. Sci., 9 [2], pl56-170.

[11] R. A. Cutler, (1991), Engineering properties of borides in ceramics and glasses:

Engineered materials handbook, vol 4, Edited by S. J. Schneider Jr. ASM

International, Materials Park, OH, pp787-803.

[12] R. G. Munro, (1997), Materials properties of a sintered alpha-SiC, Journal of

physical and chemical reference data, vol 26, ppl 195-1203.

[13] C. Morz, (1994), Zirconium diboride, Am. Ceram. Soc. Bull., 73 [6] p 141-142.

[14] Michael D. Uchic, Dennis M. Dimiduk, Jeffrey N. Florando, William D. Nix,

(2004) Sample dimensions influence strength and crystal plasticity, Science, 305,

p986-989.

[15] R. Dou and B. Derby, (2009), A universal scaling law for the strength of metal

micropillars and nanowires, Scripta materialia, 61, p524-527.

[16] Frederic Monteverde and Alida Bellosi, (2005), Development and

characterization of metal diboride based composites toughened with ultra fine SiC

particulates, Solid state sciences, 7, p622-630.

73

Chapter 4

Multiphase ZrC-ZrB2-SiC ceramics

Abstract

The effects of composition on mechanical, thermal and electrical properties of

ZrC-ZrB2-SiC multiphase composites were examined. The composites were

consolidated by spark plasma sintering. The microstructure of the as-consolidated

composites was examined by field emission scanning electron microscopy and

transmission electron microscopy. The elastic moduli of the composites were

measured using the longitudinal and transverse soundwave velocities measured,

whereas the hardness and fracture toughness of the composites were determined using

an indentation measurement. The results indicated that the shear modulus was in the

range 180 to 225 GPa and the Young's modulus was in the range 434 to 517 GPa. The

ranges of hardness and fracture toughness values were measured to be 18.79 to 21.50

GPa, and 4.69 to 6.08 MPa m1/2, respectively. On the other hand, the thermal and

electrical conductivities of the ZrC-ZrB2-SiC composites were measured at room

temperature by a nanoflash technique and a current-voltage method, respectively. The

thermal conductivities for the composites were in the range 38.25 to 98.25 W (m K)'1.

The electrical conductivities of the multiphase composites were in the range

0.916xl04 to 4.521xl04 (Q cm)"1. The applied stress to micropillar of ZrC-ZrB2-SiC

74

composites produced long cracks from surface to bottom without deflecting along

grain boundaries.

4.1. Introduction

Diborides and carbides of zirconium (ZrB2 and ZrC) have extremely high

melting points (>3000 °C), high thermal and electrical conductivities, chemical

inertness against molten metals, and great thermal shock resistance [1-4]. The unique

combinations of mechanical and physical properties make them attractive candidates

for structural applications at ultra-high temperatures. However, the use of these

single-phase ceramics materials for high-temperature structural applications is limited

by their poor oxidation, thermal shock and ablation resistance as well as poor damage

tolerance. A composite approach has been successfully adopted in order to improve

the oxidation and ablation resistance of single-phase ceramics. For example, the

addition of second phase such as SiC to ZrE$2 results in a composite with

improvement of strength and better oxidation, thermal shock and ablation resistance

[5-8]. The improvement of oxidation and ablation resistance is believed to arise from

the formation of coherent passivating oxide scales on the surface. More recently, the

ZrC-ZrB2-SiC multiphase composite system has been demonstrated to have superior

resistance to ablation or conversion than do the corresponding ZrB2/SiC composites

under an arc-jet environment [8]. The ZrB2-30ZrC-10SiC (vol %) composites have

been successfully consolidated using both hot-pressing and spark plasma sintering [9,

76

10]. Furthermore, this composite exhibited high strength with low scattering of

strength up to 1500°C, these components are also both thermodynamically and

chemically stable at high temperature because the intergranular reaction in the

composite is absent [10]. However, the effects of separate components on

mechanical, thermal and electrical properties of ZrC-ZrB2-SiC composites are not

yet systematically studied. The compositional dependence of these properties is very

important for designing ZrC-ZrB2-SiC composites in sustained thermomechanical

applications.

On the other hand, spark plasma sintering (SPS) is one of the most recent

processing techniques developed for densifying ceramic materials, including poorly

sinterable compounds [11, 12]. One advantage of SPS concerning ceramics is that the

grain growth of starting materials is restricted, as a considerable shorter sintering time

(within several minutes) is required compared to hot-pressing or hot isostatic

pressing, thereby retaining the fine and homogenous grains. Another is to enhance

densification of poorly sinterable ceramics, as a result of the spark discharge

generated between powders as well as the presence of the electrical field under pulsed

direct current. Previous studies in ZrB2-based ceramic materials showed that SPS

enhanced densification and refined microstructure in very short processing cycles [10,

77

13]. This is attributable to the presence of the electrical field during SPS which

caused faster diffusion due to the enhanced migration speed of ions [14].

In the present study, ZrC-ZrE^-SiC composites with different compositions were

consolidated by spark plasma sintering. The effects on physical properties,

mechanical properties with microstructural analysis, thermal properties, and electrical

properties of ZrC-ZrB2-SiC composites are discussed in this chapter. In addition,

fracture mechanisms based on micropillar compression test are presented.

4.2. Experimental procedures

4.2.1 Materials

The starting powders used in this study were: ZrB2 powder (Grade F, Japan New

Metals, Tokyo), average particle size —2.12 |im, ZrC powder (Grade F, Japan New

Metals), average particle size «2.32 um; and a-SiC powder (Grade UF-15, H.C.

Starek, Berlin, Germany), average particle size «0.5 jina. In order to examine the

effect of composition on the mechanical, thermal and electrical properties, eight

series of ZrC-ZrB2-SiC compositions were prepared in this study. These compositions

are shown in table 4-1. The powder mixtures were wet-mixed using SiC milling

media and ethanol for 24 h, and then dried in oven. Before sintering, the dried

78

mixtures were sieved through a metallic sieve with 60-mesh screen size. The powder

mixture was put into a graphite die lines with graphite foil and densified using spark

plasma sintering (SPS-1030, Simitomo Coal Mining Co. Ltd., Tokyo, Japan). The

sintering was performed at 1950°C for 2 min under a pressure of 50 MPa in argon

atmosphere. Changes in temperature and sintering displacement were recorded by a

computer during the entire sintering process. Final sintered specimen size was 10 mm

in diameter pellets with a thickness of -2.0 mm. After removing the surface of the

sintered compact to avoid contamination from the graphite die, the densities of the

sintered composite compacts were measured by the Archimedes method with distilled

water as medium. X-ray diffraction (XRD) was used for crystalline phase

identification of the composites. The grain size of each component, ZrB2, ZrC, and

SiC, was measured directly from the field emission scanning electron microscopy

(FE-SEM) images of sintered composites using image J software. In order to estimate

the average grain size of each component, five FE-SEM images were used for direct

measurement within the range of 200-300 population. Microstructure of the

composites was observed by FE-SEM and transmission electron microscopy (TEM).

4.2.2 Elastic moduli, hardness, and fracture toughness measurements

The elastic moduli measurements of the composites were performed using an

79

ultrasonic technique (TDS 3052B, Tektronix Inc. Beaverton, OR USA) with a

fundamental frequency of 20 MHz. Young's modulus, E, shear modulus, G, and

Poisson's ratio, v, were calculated using the longitudinal and transverse soundwave

velocities measured in the composite specimens. The details of calculations were

listed in chapter 3. The hardness and fracture toughness, KIC, of the composites was

determined using an indentation technique. The indentation tests were performed on

the polished surface of the specimens by loading with a Vickers microhardness

indenter (AVK-A, Akashi, Co., Ltd., Yokohama, Japan) for 20 s in ambient air at

room temperature. The corresponding diagonals of the indentation and crack sizes

were measured using an optical microscope attached to the indenter. The indentation

load of 9.8 N was used, and ten indents were made for each measurement. The

fracture toughness, Klc, of composites were calculated from the Palmqvist equation

[15].

4.2.3 Thermal and electrical conductivity measurements

The thermal diffusivity, a, of the composites was measured on a disk-shaped

specimen with a diameter of 10 mm and thickness of ~2 mm using the nanoflash

technique (LFA447/2-4N, Nanoflash, NETZSCH-Geratebau GmbH, Postfach,

Germany). The flash source is a Xenon flash lamp operating in the wavelength range

80

of 0.15 urn to 2 um. Prior the measurements, the samples were coated with a colloidal

graphite spray in order to enhance the absorption of the Xenon light pulse energy and

the emission of IR radiation to the temperature detector. Also, the heat capacity, C ,

was measured with alumina as the reference material. All of the measurements were

performed in ambient air at room temperature. Subsequently, the thermal

conductivity of the composites, kc, is determined from thermal diffusivity, heat

capacity, and density of the composites according to the following equation, [16]

kt=pCpa (1)

where p is the density of the composites, Cp is the heat capacity of the

composites, and a is the thermal diffusivity of the composites.

Moreover, the electrical conductivity measurements of the composites were

performed using the four-wire probe at room temperature. A power supply (Model:

6220, Keithley, Cleveland, Ohio, USA) and digital multimeter (Model: 2182

Nanovoltmeter, Keithley) were used to measure the IV characteristics of the samples.

4.2.4 Micropillar compression test

ZrC-ZrB2=SiC composites micropillars with ~5 um diameter and -20 um height

were prepared by Focused Ion Beam (FIB Nova 600, FEI, Hillsboro, OR) on the

surface of composites in order to investigate fracture behavior or effect of

81

compression stress on ZrC-ZrB2-SiC composites. The prepared micropillars of each

sample (ZZS-1, ZZS-2, ZZS-6, and ZZS-7) were compressed to the depth of 500 nm

from the sample surface with the 10 \m\ diameter flat punch tip of a nanoindenter

(MTS Nanoindenter/XP, MTS system, Eden Prairie MN). An optical microscope with

high magnification (X 500) in the nanoindenter was used to place the flat punch tip

precisely to avoid tilting the micropillars during compression. Applied stress and

displacement of samples were observed to identify events such as fracture and

deformation. Cracks, fracture surface, or dimensional changes of micropillars were

observed by FE-SEM after compression tests were finished. In addition, detailed

microstructural observation of compressed micropillars was conducted by TEM.


4.3.1 Densities, phases, and microstructures

The shrinkage curves obtained during the SPS cycle for the various composite

materials consolidated by SPS are shown in figure 4-1. The shrinking behavior of

ZrC-ZrB2-SiC (ZZS) is almost the same for the studied compositions regardless of

component content. The measurable shrinkage was observed at temperatures ranging

from 1500 to 1600°C, depending upon the compositions. For examples, the onset

82

temperature of densification was determined to be ~1530°C for ZZS-1, ~1580°C for

ZZS-2, ~1550°C for ZZS-3, and ~1540°C for ZZS-5. During subsequent

densification, all of the compositions showed almost the same shrinkage rate with

time. The main part of the densification occurred within a period of ~2 min, whereas

conventional hot-press typically required hours of densification at substantially higher

temperatures to yield fully densified compacts. The measured densities and relative

densities for the various ZrC-ZrB2-SiC composite materials consolidated by SPS are

summarized in table 4-2. The theoretical densities of the composites were calculated

according to the rule of mixtures. It is evident that all the composites were almost

fully densified (relative density > 98%) and their relative density was almost the same

regardless of composition. This is attributed to the fact that the same shrinkage

behavior was observed during SPS for all of the studied compositions.

X-ray diffraction patterns for the various ZrC-ZrB2-SiC composite materials

consolidated by SPS are presented in figure 4-2. Although the peaks of ZrB2, ZrC and

SiC phases showed the different intensity with compositions, only ZrB2, ZrC and SiC

phases were detected in every case. This suggests that an intergranular reaction did

not occur at the grain boundaries during sintering. Typical microstructural features of

the ZrC-ZrB2-SiC composites observed under FE-SEM is shown in figure 4-3. The

83

general microstructures of the ZrC-ZrB2-SiC composites are similar, consisting of the

equiaxed ZrB2 (grey contrast), ZrC (bright contrast), and SiC (dark contrast) grains.

In the case of the ZZS-1 composition, ZrB2, ZrC, and SiC particles are homogenously

and individually present in the isolated locations because the added ZrB2, ZrC, and

SiC amounts are the same for this composition. For other compositions, however, the

primary component in the composite, such as ZrB2 particles for ZZS-2, ZZS-6, and

ZZS-7, ZrC particles for ZZS-3, ZZS-4, and ZZS-8, and SiC particles for ZZS-5,

formed an interconnected skeleton structure with the other two components

homogenously dispersed in it. The average grain size of each component (ZrC, ZrB2,

and SiC) about ZrC-ZrB2-SiC composites was listed in the table 4-3. The average

grain size of ZrB2 was in the range of 1.25 to 1.75 um, that of ZrC was in the range of

0.99 to 1.39 um, and that of SiC was in the range of 0.59 to 1.04 um. It is noted that

the milling process was effective to reduce initial powder size such as ZrB2 (-2.12

pm) and ZrC (-2.32 pm) but not helpful to reduce average grain size from small

initial powders like SiC (-0.5 um).

Typical TEM images for ZZS-1 and ZZS-2 are listed in figure 4-4 and figure

4-5, respectively. The contrast of each component grain is opposite that of each

component grain in FE-SEM images: ZrB2 (grey contrast), ZrC (dark contrast), and

84

SiC (bright contrast). Phases and crystallographic information of all three

components, ZrC, ZrB2, and SiC, were confirmed by SAED, but no second phase was

found inside the grains (compared to ZrE$2-nano sized SiC composites). Typical

interphase interfaces between ZrB2-ZrC and ZrC-SiC are shown in figure 4-6. Any

TEM interface images with high resolution mode could not be obtained due to large

TEM lamellae thickness, but interphase interfaces were very sharp and narrow so that

any second phase cannot be expected in the interfaces.

4.3.2 Elastic moduli, hardness, and fracture toughness

The elastic moduli measured in the various ZrC-ZrB2-SiC compositions

consolidated by SPS are listed in table 4-2. From this table, it is found that the shear

and Young's moduli are related to component content. In the case of the same volume

percent of ZrB2, ZrC, and SiC-containing composition (ZZS-1), the shear and

Young's moduli are 205 GPa and 477 GPa, respectively. Then, both the elastic moduli

increased with increasing ZrB2 as well as SiC addition, but decreased with increasing

ZrC addition. The highest shear and Young's moduli were measured in ZZS-2

composition, and their values are 225 GPa and 517 GPa, respectively. The lowest

shear and Young's moduli were measured in the ZZS-3 composition, and their values

are 180 GPa and 435 GPa, respectively. In contrast, Poisson's ratio remains almost

85

constant for the studied compositions regardless of component content.

The hardness and fracture toughness of the ZrC-ZrB2-SiC composites

obtained from an indentation technique are summarized in table 4-4. Note that for the

ZZS-5 composition the fracture toughness was not listed in this table because it is

difficult to measure the crack length accurately. The ranges of hardness and fracture

toughness values were measured to be 18.8 to 21.51 GPa, and 4.69 to 6.1 MPam1/2,

respectively. The results indicated that both the hardness and fracture toughness of the

composites are dependent on the composition. The compositional dependence of

hardness and fracture toughness may be associated with the complex residual stress

state that develops during cooling from the processing temperature due to the thermal

expansion mismatch among ZrE$2 (CTE: 6.5 ppm/°C), ZrC (CTE: 7.1 ppm/°C) and

SiC (CTE: 4.7 ppm/°C). A typical cracking pattern is shown in figure 4-7. The crack

propagated across the ZrC, ZrB2 and SiC grains without being deflected along the

grain boundaries.

The fracture toughness of hot-pressed ZrE$2 , SiC was reported to be 2.3-3.1

MPa- m1/2,[17] and 3.0-4.3 MPa- m1/2,[18] respectively. The fracture toughness

values of ZrB2-SiC (10-30 % SiC) composites were reported to be between 4.1-5.3

MPa- m1/2 [8]. Therefore, the fracture toughness values of the ZrC-ZrB2-SiC

86

multiphase composites are comparable to that of the single-phase ceramics and

ZrB2-SiC composites.

4.3.3 Thermal conductivity

The measured heat capacities, thermal diffusivities, and the calculated

thermal conductivities of the various ZrC-ZrB2-SiC composites consolidated by SPS

are summarized in table 4-5. From this table, it is obvious that the heat capacity and

the thermal diffusivity decreased with increasing ZrC amount, but increased with

increasing SiC and ZrB2 additions. The heat capacity was in the range of 0.5 to 0.62

Jg'K" , showing a compositional dependence. The thermal diffusivity was in the

range of 12.36 to 30.46 mm2 s"1. It is evident that the compositional dependence is

stronger for the thermal diffusivity than for the heat capacity. This strong dependence

of thermal diffusivity on composition suggests that the thermal conductivity of the

ZrC-ZrB2-SiC composition is dominated by the heat flow in the composites. In the

case of ZZS-1, the same volume percent of ZrB2, ZrC, and SiC containing

composition, the thermal conductivity measured was 72.64 Wirf'K~'. The thermal

conductivity then decreased with increasing amount of ZrC. In particular, the ZZS-3

composition showed the low heat capacity as well as the lowest thermal diffusivity, in

turn resulting in the lowest thermal conductivity in the studied compositions. The

87

thermal conductivity dropped from 72.64 Wirf'KT1 for the ZZS-1 composition to

38.25 Wm_1K~l for the ZZS-3 composition, for approximate loss of 50%. On the

other hand, the thermal conductivity of the ZrC-ZrB2-SiC composites increases with

increasing ZrB2 as well as of SiC amount. The improvement of thermal conductivity

is more substantial for increasing SiC than for increasing ZrB2. The ZZS-5

composition, 50 vol.% SiC containing composite, showed the highest heat capacity as

well as the highest thermal diffusivity, which in turn resulting in the highest thermal

conductivity among all the materials. The thermal conductivity increased from 72.64

Wm^K"1 for the ZZS-1 composite to 92.85 WnT'K"1 for the ZZS-5 composition

material, for approximate increase of 30%.

It is known that the thermal conductivity of the composites is dependent on

the thermal conductivity of the components and the interfacial thermal resistance of

the grain boundaries. The thermal conductivity of SiC is higher than that of the ZrB2

and ZrC materials, [1, 19, 20] and ZrC has the lowest thermal conductivity among the

ZrB2, ZrC, and SiC components [1, 19]. This implies that the increasing SiC and ZrB2

content should improve thermal conductivity of ZrC-ZrB2-SiC ceramics. Conversely,

increasing ZrC content should decrease thermal conductivity. This effect is closely

linked to the amount of SiC, ZrB2 and ZrC additions as well as to the distribution

88

because they influenced the heat flow resistance through the components and the

interfaces. In the case of high ZrC-containing ZrC-ZrB2-SiC (ZZS-3, ZZS-4, and

ZZS-8), the ZrC was the pristine phase and the other two phases (ZrB2 and SiC) were

dispersed in it. This structure characteristic led to increased resistance for the heat

flow through the components and their interfaces, compared to ZZS-1. This thermal

resistance was enhanced with increasing amount of ZrC. In the case of high SiC or

ZrB2-containing ZrC-ZrB2-SiC (ZZS-2, ZZS-5, ZZS-6, and ZZS-7), on the other

hand, the ZrC particles were embedded in a SiC or ZrB2 matrix. This characteristic

structure of SiC or ZrB2 particles formed in ZrC-ZrB2-SiC composites should

enhance its heat transport capability, because it could provide a route of higher heat

flow. In particular, SiC has the highest thermal conductivity among ZrB2, SiC and

ZrC. Thus, in the present study, the addition of SiC or ZrB2 improved heat capacity as

well as heat transport, resulting in high thermal conductivity.

4.3.4 Electrical conductivity

In figure 4-8, the two examples of the current-voltage relations measured at

room temperature for the various ZrC-ZrB2-SiC composites consolidated by SPS are

presented. It is clear that the current increased linearly with voltage for every case, i.e.

a linear relationship between current and voltage. This indicated that good ohmic

89

contacts have been realized between the measured composite samples and the

electrode. Additionally, the slope of the current-voltage curve is related to the

compositions: high ZrB2 content resulted in a lower slope. This indicated that the

resistance is reduced with increasing ZrB2 in the ZrC-ZrB2-SiC compositions. The

measured electrical resistivity and conductivity of the various ZrC-Z'rB2-SiC

composites consolidated by SPS are summarized in table 4-6. The electrical

conductivity of the ZrC-ZrB2-SiC composites was measure to be in the range of

0.916xl04 to 4.521xl04 Q ' W 1 . In the case of ZZS-1, 33.3 vol.% ZrB2-33.3 vol.%

ZrC-33.3 vol.% SiC composite, the measured electrical conductivity was 1.606xl04

fr'cirf'. The electrical conductivity improved with ZrB2 addition. The highest

electrical conductivity was measured in the ZZS-2 composition composite and the

value was 4.521 xlO4 fT'cm"1. On the other hand, the electrical conductivity decreased

with increasing ZrC and/or SiC contents. The lowest electrical conductivity was

measured in the ZZS-4 composition composite. Although bulk ZrB2 and ZrC are

located in the electrical conductivity range of conductors, the electrical conductivity

of ZrC was significantly lower than that of ZrB2 [19]. Thus, the decrease in electrical

conductivity due to ZrC addition is a result of the lower electrical conductivity of

ZrC. In addition, it was found that the addition of SiC further reduced the electrical

90

conductivity of the ZrC-ZrB2-SiC composites because SiC is a semiconductor.

However, the lowest electrical conductivity was not observed in the 50 vol.%

SiC-containing composites, but in the 30 vol.% SiC-containing composite. This is

because of high ZrB2 content for the former (30 vol.% ZrB2) compared with the latter

(15 vol.% ZrB2). This indicated that ZrB2 addition is important for improving the

electrical conductivity of ZrC-ZrB2-SiC composites. Although the high ZrC and/or

SiC-containing ZrC-ZrB2-SiC composites exhibited lower electrical conductivity as

compared with high ZrB2-containing composites, the electrical conductivities of all

the ZrC-ZrB2-SiC composites investigated in this study are still within the range of

conductors. As a result, electrical discharge machining can be used for all the

ZrC-ZrB2-SiC composites.

4.3.5 Micropillar compression tests of ZrC-ZrB2-SiC composites

Typical FE-SEM images of fabricated micropillars about ZZS-1, ZZS-2, ZZS-6,

and ZZS-7 are listed in figure 4-9. Both ZZS-1 and ZZS-2 micropillars did not follow

the universal scaling law mentioned in chapter 3 and these two micropillars were

broken at fairly low maximum stress with 4.38 GPa for ZZS-1 and 1.39 GPa for

ZZS-2 shown in figure 4-10 and figure 4-11, respectively. Typical fracture surfaces of

ZZS-1 and ZZS-2 micropillars were shown in figure 4-12 and figure 4-13,

91

respectively. Most grains of both micropillars exhibited intragranular fracture

behavior rather than intergranular fracture behavior. The stress-strain (%) curve of

ZZS-6 and ZZS-7 are listed in figure 4-14 and figure 4-15, respectively. The

maximum stress of ZZS-6 and ZZS-7 reached 3.3 GPa and 4.42 GPa, respectively.

However, the micropillar of ZZS-6 was not broken, but small dimensional and shape

changes were noticed by FE-SEM images after applied compression stress in figure

4-16. The micropillar was slightly tilted to one direction (indicated by arrow mark in

the image). The images taken after applied compression stress for the ZZS-7

micropillar are shown in figure 4-17. The image was captured right before complete

fracture. It is noted that the crack did not0 start from the bottom edge of the micropillar

but from the top surface of the micropillar in the longitudinal direction. Several TEM

microstructures are listed in figure 4-18. In these images, stacking faults and small

dislocations were found in the number 1 area and the number 2 area in the TEM

lamella, respectively. The dislocations were very small and located in one grain so

that they were produced from the milling process, but stacking faults might be

produced by applied compression stress because stacking faults were connected

through several grains. These results were very similar to the ZSN system discussed

in chapter 3.

92

4.4. Conclusion

(1) The ZrC-ZrB2-SiC composites were consolidated by SPS at 1950°C for 2

min under a pressure of 50 MPa, and the composites were almost fully dense

regardless of composition. The microstructure of the composites consisted of the

equiaxed ZrB2, ZrC and SiC grains. Primary component phase formed a short net-like

structure, and the other secondary phases were homogenously dispersed in it.

(2) The shear modulus of the ZrC-ZrB2-SiC composites was in the range 180

to 225 GPa, and the Young's modulus was in the range 434 to 517 GPa, depending on

composition. Poisson's ratio was almost the same for all the studied compositions.

The ranges of hardness and fracture toughness values were measured to be 18.8 to

21.51 GPa, and 4.69 to 6.1 MPa- m1/2, respectively.

(3) The thermal conductivity of the ZrC-ZrB2-SiC composites decreases with

increasing ZrC content, but it increases with increasing SiC and/or ZrB2 content. The

measured thermal conductivities are in the range 38.25 to 87.59 Wm"'K"', depending

on component content.

(4) The electrical conductivities of the ZrC-ZrB2-SiC composites increased

with increasing ZrB2 content, conversely; the electrical conductivity decreased with

increasing SiC and/or ZrC content. The measured electrical conductivities were in the

93

range 0.916xl04 to 4.521xl04 O ' W 1 .

(5) The applied compression stress to ZZS-1 and ZZS-2 micropiUars led to

complete fracture with maximum stress of 4.38 GPa and 1.39 GPa, respectively. The

longitudinal cracking behavior was observed in the ZZS-7 micropillar. In addition,

stacking faults were generated by applied stress (by TEM investigation).

94

Materials

ZZS-1

ZZS-2

ZZS-3

ZZS-4

ZZS-5

ZZS-6

ZZS-7

ZZS-8

Compositions (Vol.%)

ZrB2

33.3

70

15

15

30

55

55

30

ZrC

33.3

15

70

55

20

15

30

55

SiC

33.3

15

15

30

50

30

15

15

Table 4-1. Compositions of ZrC-ZrB2-SiC composites.

95

Materials

ZZS-1

ZZS-2

ZZS-3

ZZS-4

ZZS-5

ZZS-6

ZZS-1

ZZS-8

Theoretical

Density

(g/cm3)

5.51

5.85

6.14

5.70

5.01

5.51

5.93

6.06

Measured

Density

(g/cm3)

5.44

5.76

6.05

5.65

4.94

5.44

5.86

5.97

Relative

Density

(%TD)

98.7

98.5

98.5

99.1

98.6

98.7

98.8

98.5

Elastic Properties

G (GPa)

205

225

180

192

206

215

211

196

E (GPa)

477

517

435

449

486

500

496

457

V

0.17

0.16

0.18

0.17

0.18

0.16

0.18

0.17

Table 4-2. Densities and elastic properties of ZrC-ZrB2-SiC composites.

96

Sample

ZZS-1

ZZS-2

ZZS-3

ZZS-4

ZZS-5

ZZS-6

ZZS-7

ZZS-8

ZrB2 (um)

1.33±0.61

1.30±0.54

1.75±0.61

1.27±0.49

1.25±0.52

1.53±0.55

1.71±0.77

1.31±0.60

ZrC (um)

1.09±0.48

1.21±0.46

1.39±0.65

1.33±0.60

0.99±0.43

1.20±0.41

1.35±0.62

1.36±0.51

SiC (um)

0.78±0.36

0.82±0.36

0.60±0.25

0.59±0.25

1.04±0.55

0.74±0.31

0.64±0.26

0.58±0.26

Table 4-3.The measured grain size of each component ZrB2, ZrC, and SiC for

ZrC-ZrB2-SiC composies

97

Sample

ZZS-1

ZZS-2

ZZS-3

ZZS-4

ZZS-5

ZZS-6

ZZS-7

ZZS-8

Hardness (GPa)

19.20 ± 0.96

21.51 ± 1.32

19.56 ± 1.00

18.80 ± 0.78

20.40 ± 1.94

19.64 ± 0.74

19.40 ± 1.21

19.61 ± 0.75

Toughness (MPa.m1/2)

6.10 ± 0.75

6.09 ± 0.53

4.69 ± 0.26

5.52 ± 0.33

N/A

5.76 ± 0.42

5.04 ± 0.37

5.63 ± 0.24

Table 4-4. Hardness and fracture toughness of ZrB2-ZrC-SiC composites

98

Samples

ZZS-1

ZZS-2

ZZS-3

ZZS-4

ZZS-5

ZZS-6

ZZS-1

ZZS-8

Heat capacity

Cpdg'K-1)

0.58

0.54

0.51

0.55

0.62

0.58

0.53

0.50

Thermal

diffusivity

a(mm2s_1)

22.98

27.69

12.36

16.82

30.46

28.20

23.64

16.33

Thermal

conductivity

Kc (Wm-'K"1)

72.64

85.63

38.25

51.77

92.85

89.02

73.73

49.07

Table 4-5. Thermal properties measured at room temperature for the ZrC-ZrE$2-SiC

composites

99

Samples

ZZS-1

ZZS-2

ZZS-3

ZZS-4

ZZS-5

ZZS-6

ZZS-7

ZZS-8

Electrical

resistivity

R(10"5Ocm)

6.226

2.212

9.384

10.917

8.282

3.333

3.094

5.682

Electrical conductivity

( lOtoW1)

1.606

4.521

1.066

0.916

1.207

3.002

3.231

1.760

Table 4-6. Electrical properties measured at room temperature for the ZrC-ZrB2-SiC

composites

2.0

0 -

~! 1—I 1 1 1 1 1 1 1 1 — | 1 T - , ,"l 1 1 1 1 1—I 1 1 1 1 1 ] 1 1—1 1 "

(a)

E 1-5 E

• 4 - j

c f? 1.0 <D O 03 CL .«a 0.5 Q

- 1800

-I—I—I 1—I—I—I—1—I—I—1—1—I—I—I—I—I—I—1—I—I—I—I—1—I 1—I—I—I—I—u-

2000

1000 0 50 100 150 200 250 300

Sintering Time (second)

Figure 4-1. Typical examples of recorded shrinkage curves during the SPS cycle for

ZZS-2, ZZS-3, and ZZS-5.

101

o to

N N GO •

y

"2 N N GO

i*. 'o1

N N GO i

<̂» ^ •£> N N GO •

0 \

N N GO

| - J J ^ N N GO

0 0

31 era' c Of

4^ i

X 1

3 a.

5? P3

o o" 3

•a

C/5

O

8 =r « P 3 HT

1? N N GO

^ ^̂ ^ N N GO

ro o Q. CD (Q CD CD

ro o

CO o

o

a i o

CD o

--4 o

cx> o

*? ^ ' /

: |

f L f

-r i~ ' . f ?

.

t f :f i

.

*"

I f • f cz *"- c-c c c cr I L r f—

to

Intensity, (a.u)

S- 5 J S f ( ^

5|ua «mmmM&¥mm«*^&vm

Figure 4-3. Typical FE-SEM images for the ZrB2-ZrC-SiC composites

103

Figure 4-4. Typical TEM images for the ZZS-! composite

104

Figure 4-5. Typical TEM images for the ZZS-2 composite

105

>*~ " W i d C D )

Figure 4-6 Typical TEM images of interphase interface between (a) SiC and ZrB2

and (b) ZrB2 and ZrC

106

Figure 4-7. Typical cracking behavior of ZrC-ZrB2-SiC composites

107

20

. — s .

> O —̂ X <D

O >

15

10

5

U

-5

-10

-15

-20 _ j i—i i _

—I—i—I—i—i—i—r-

•ZZS-4 •ZZS-6

-150 -100 -50 0 50

Current (mA)

100 150

Figure 4-8. Current versus voltage measured at room temperature for the

ZrC-ZrB2-SiC composites

108

Figure 4-9. Fabricated micropillar of ZZS -1, ZZS-2, ZZS-6, and ZZS-7

109

0 10 20 30 40 50 60 70

AL/L(%)

Figure 4-10. Stress vs. strain (%) curve of ZZS-1 micropillar

110

(TJ Q_ O

</5

1.4

1.2

1.0

0.8

0.6

0.4

0.2

0.0 -

• •

• ZZS-2

. , . .

•

/

/

- H r— 0 10 20 30 40

Al_/L(%)

50 60


111

Figure 4-12. Intragranular fracture surface of ZZS-1 micropillar

112

Figure 4-13. Intragranular fracture surface of ZZS-2 micropillar

113

3.5

3.0

2.5 ^ ^

S. 2.0 O "**-•

w 1.5 w 0 i _

CO 1 0

0.5

0.0 i

• ZZS-6

#

/

/I it /I

H If # /

/ /

yj 0.00 0.05 0.10

AL/L (%)

0.15 0.20


114

5

1 I I I L

0.00 0.05 0.10 0.15 0.20 0.25 0.30

AL/L (%)


115

CO Q . CD

</) w <D !_ -*—•

CO

4

3

2

1

0

Figure 4-16. Dimensional and shape change of ZZS-6 micropillar

116

* • _ » 2*

Figure 4-17. Typical cracking behavior of ZZS-7 micropillar

117

Figure 4-18. Typical TEM images of ZZS-6 micropillar after compression stress.

118

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[20] Y. Takeda, (1988), Development of high-thermal-conductive SiC ceramics. Am.

Ceram. Soc. Bull, 67(12), pl961-1963.

120

http://www.jnm.co.jp

Chapter 5

SCS 9a fiber-reinforced ZrB2-RB(Reaction-bonded) SiC composite

Abstract

SCS 9a fiber-reinforced ZrB2-RB (Reaction-bonded) SiC composite was

examined. SCS 9a fiber-reinforced ZrB2-RB SiC composite was fabricated by hot

pressing with reaction bonding process of SiC after tape casting with slurry and

woven SCS 9a fibers. Density was measured by pycnometer and porosity was

calculated from measured density. Hardness and fracture toughness of matrix

materials were measured by indentation technique, but fracture toughness of the

composite was measured by four point bending method. At last, microstructural

analysis was conducted by FE-SEM with EDS and TEM with EDS and SAED

(Selected Area Electron Diffraction). The porosity was 10.40% due to SiC reaction

bonding process and the large cracks (including micro cracks were observed; these

are due to the thermal expansion mismatch between the fiber and matrix materials.

This high porosity and large cracks with micro cracks affected the hardness and

fracture toughness of the composites. However, the fracture toughness of the

composite was improved compared to the fracture toughness of the matrix itself

because of the effect of fibers. No chemical reaction was found by thermodynamical

calculation, SEM, and TEM analysis in the interface between the fiber and the matrix

121

materials. In addition, complex SiC structures such as twin boundaries and stacking

faults were observed due to the reaction bonding process of SiC.

5.1 Introduction

Zirconium diboride has a high melting temperature (>3000 °C), chemical

stability, high electrical and thermal conductivities, resistance to erosion/corrosion,

and good mechanical properties, making it promising candidate for

ultra-high-temperature applications [1]. Addition of a silicon carbide to zirconium

diboride enhances the oxidation resistance and limits the zirconium diboride grain

growth due to borosilicate diffusion barrier and liquid phase sintering effect,

respectively [2, 3].

However, the use of ultra-high-temperature ceramics without reinforcements has

limitations due to low fracture toughness and poor thermal shock resistance [4]. Thus,

the incorporation of a fiber reinforcement phase into ZrB2-SiC ceramics is

indispensible in order to improve fracture toughness, thermal shock resistance, as

well as to lower the composite density [5]. Among various ceramic fibers for high

temperature applications, SiC fiber with low oxygen content and small diameter has

good thermomechanical reinforcement capability [6].

Reaction bonding of ceramics has various advantages, which include low raw

materials costs, near net-shape tailorability, low-to-zero shrinkage capability, and

glass-phase-free grain boundaries for many technical and high performance

123

applications. Among these advantages, the low-to-zero shrinkage capability makes

most reaction forming techniques suitable for the fabrication of composites [7].

Levine, et al. studied ZrB2-SiC ceramics, ZrB2-SiC-C, and SiC fiber-reinforced

ZrB2-SiC composites. The flexural strengths of SiC fiber-reinforced composites were

130 MPa, 101 MPa, and 84.5 MPa at room temperature, 1127 °C, and 1327 °C,

respectively. The flexural strength of composites were fairly low due to high porosity

(up to 30%) from incomplete densification and large amount of micro cracks from the

difference between coefficient of thermal expansion of matrix and that of fiber.

However, the composites were not separated into two separate pieces due to crack

bridging by the fibers [8].

In this study, SiC fiber-reinforced ZrB2-RB (Reaction Bonding) SiC composites

were consolidated by a hot pressing method. The physical, mechanical, and

microstructural analyses are discussed in this chapter.

5.2 Experimental procedures

5.2.1 Fabrication

The starting powders used in this study were ZrB2, Si, and carbon powders and

the fibers used in this study were SCS 9a fibers (Specialty Materials Inc, Lowell,

124

MA), which is composed of a 33 um carbon core and 46 um SiC outer layer. The first

step to prepare composite billets with size of 10 cm by 7 cm was mixing the three

powders to slurry state with composition of 20 vol% SiC and 80 vol% ZrE$2. The SCS

9a fibers were wound on the four-sided mandrel with eight-layer thickness to achieve

20 vol% SiC fiber content. The wound fibers were coated with the matrix slurry by

tape casting. The prepared green billet was dried in oven at 110 °C, and then the dried

billet was trimmed and loaded into the hot pressing die. The die was placed on a

support fixture and loaded into an inert gas furnace. The furnace was heated from

room temperature to 525 °C at 0.55 °C/min and held for one hour. The die was then

loaded into the hot-press and processed at reaction bonding temperature (1450-1650

°C) and under a vacuum of 4.3X10"3 Torr for 75min. No extrusion of fiber of matrix

was observed upon removal from the die.

5.2.2 Phase analysis, density, and porosity measurements

X-ray diffraction was used in order to confirm the crystalline phases, especially

reaction-bonded SiC. The density of fabricated composite was measured directly by

pycnometer with helium gas as a medium for accurate. The porosity was calculated

from apparent density, and measured density by direct measurement from the

pycnometer. The simple equation is that porosity is (1-measured density/apparent

125

density) X100 [9]. For density and porosity, five samples were cut from the

composite billet.

5.2.3 Elastic properties, hardness, and fracture toughness measurements

Elastic properties could not be measured by ultrasonic technique due to

difficulties in measuring the velocity of sound waves. Young's modulus was

measured by nano-indentation technique (MTS Nanoindenter/XP, MTS system, Eden

Prairie MN). Hardness and fracture toughness of matrix materials were measured by

microhardness tester as described in previous chapters. The load for measurement

was 49N and ten tests were made in the matrix part. The single edge pre-cracked

beam tested in accordance with ASTM C 1421 [10] was used to determine the

fracture toughness (KiPb) of SiC-fiber-reinforced ZrE$2-SiC composites with four test

samples. All samples were polished with wet silicon carbide papers (grits 400, 600,

and 1200, successively) on all surfaces except end planes, and the sharp edges were

chamfered. The equations for calculations of fracture toughness are listed below and

the detailed dimensional information is shown in figure 5-1.

• Iph f P \S - S 110"

max L o / J BWV2

3 [a I W f 2[l-a/W]3 (1)

l3A9-0.6S\a/W] + l35\a/W]2)\a/W](l-\a/W]} / = 1.9887-1.326[a/^]--L L _ J l. Ul i i _ J £ (2)

L J {\ + [a/W] 2

126

5.2.4 Microstructure observations

In order to analyze the structural information of SiC fibers and matrix materials,

a whole cross sectional view of the composites was obtained by optical microscope

with CCD camera (Carl Zeiss Microimaging Inc. Thornwood NY). A total of 18

(6 X 3) pictures were taken of different areas with X 50 magnification so that all fiber

layers structures should be included. The elemental analysis was conducted by

FE-SEM with EDS for finding unknown elements in the composites. The crack

density and grain size of each component were investigated by FE-SEM images with

Image-J software. In order to estimate average crack density and average grain size of

each component, ten FE-SEM images were used. In addition, fracture mechanism of

matrix and composites were analyzed by FE-SEM images.

To investigate detailed microstructural information of hot-pressed composites

with SiC fibers, randomly selected areas of composites matrix and interface between

SiC fibers and matrix were observed by TEM (FBI-Philips CM300, FEI, Hillsboro,

OR) with EDAX Energy Dispersive X-ray spectrometer (EDS) and Si/Li detector

super ultrathin window. The TEM samples used in this study were prepared by

Focused Ion Beam (FIB Nova 600, FEI, Hillsboro, OR) with tungsten probe tips and

four fingers copper grid.

127


5.3.1 Density, porosity, and phase.

. The measured density, the apparent density, and the theoretical density of

hot-pressed SCS 9a fiber-reinforced ZrB2-20 vol% SiC composites billet were 3.98

g/cm3, 4.44 g/cm3, and 4.96 g/cm3, respectively. The measured densities of ZrB2-nano

SiC ceramics were in the range of 5.13-5.85 g/cm3 and those of ZrC-ZrB2-SiC

ceramics were in the range of 4.94-6.05 g/cm3. The density of the composite is lower

than that of two matrix ceramics because the density of the fiber (2.8 g/cm3) is

relatively lower than that of the other two components (ZrB2: 6.09 g/cm3, SiC:

3.21 g/cm3). The porosity of the composite from direct calculation of densities was

10.4%, which is higher than that of ZrB2-nano SiC ceramics and that of

ZrC-ZrB2-SiC ceramics. It is possible that the large amount of pores is from

reaction-bonded SiC due to volume contraction after the reaction. The FE-SEM

images of top surface view and cross sectional view are shown in the figure 5-2.

X-ray diffraction patterns for the composite consolidated by HP (Hot Press) are

presented in figure 5-3. Although the peaks of ZrB2 and SiC phases showed the

different intensity, only ZrB2 and SiC phases were detected in the top surface of the

composite and the cross section of the composite. In addition, the difference between

128

intensity of SiC peaks of the top surface of the composite and that of the cross section

of the composite was found because SiC outer layers of the fibers make stronger SiC

peaks. However, we speculate that the small carbon core of the fiber could not be

detected by X-ray diffraction.

5.3.2 Elastic moduli, hardness, and fracture toughness

The measured elastic modulus by nano indentation technique was 266±17 GPa,

which is lower than that of ZrB2-nano SiC ceramics and that of ZrC-ZrB2-SiC

ceramics. It is difficult to compare with the other two ceramics due to the different

fabrication methods and raw materials, but porosity of sintered materials makes their

mechanical and physical properties decrease exponentially [11].

The hardness and the fracture toughness for the matrix part were 7.58±0.89 GPa

and 1.44±0.26 MPa-m1/2, respectively. These values are also lower than that of

ZrB2-nano SiC ceramics and that of ZrC-ZrB2-SiC ceramics due to the large amount

of pores in the matrix part. Typical indentation marks and cracking patterns are shown

in figure 5-4. The cracks initiated from the corner of the diamond indentation pattern

and propagated across the ZrB2 and SiC grains without being deflected along the

grain boundaries of the ZrB2 grains and the SiC grains.

One of the purposes of the fiber-reinforced composite is to improve the fracture

129

toughness of materials. Thus, four point bending test with pre-cracked beam was

conducted for measurement of fracture toughness in order to confirm the effect of the

fiber-reinforced composite. The measured fracture toughness of SiC-fiber-reinforced

composite was 6.01 ±1.86 MPa-m1/2. The fracture toughness was similar to the

ZrC-ZrB2-SiC fracture toughness and lower than the ZrB2-nano SiC fracture

toughness. However, the fracture toughness of the composite was more than four

times improved if compared to that of the matrix only. In addition, it is interesting to

note that the composite beam did not separate into two separate pieces due to a crack

bridging effect by the fibers. The FE-SEM images of composite fracture are shown in

figure 5-5. The crack initiated from the edge of a pre-crack and several fibers were

fractured. Delamination of composite was observed instead of a cut through the

composite.

5.3.3 Microstructural analysis

The whole cross section view of SiC fiber-reinforced composite is shown in

figure 5-6. The composite structure is 8 layers with 0° 90° direction fiber arrays.

However, the array of fibers inside the composite was not straight compared to the

first green billet. It is possible that fiber swimming occurred during fabrication. Many

large cracks and micro cracks were found inside and outside the composite. Typical

130

FE-SEM images of the composite cracks are shown in figure 5-7. The crack density

of the composite was l~3/mm2, but the size of most cracks was very large. These

cracks are due to the matrix thermal expansion coefficient being greater than that of

the fiber and thus leading to cracking upon cooling from the processing temperature

[8]. In addition, small black particles were observed in the matrix part of the

composite shown in figure 5-8. The black particles turned out to be carbon by EDS

result with FE-SEM in figure 5-8. These carbon particles could be from residual

carbon after reaction bonding of silicon and carbon.

The interface between SiC fiber and matrix material is critical for the

fiber-reinforced composite in order to improve fracture toughness. A FE-SEM image

of the interface between fiber and matrix with high magnification (X 20000) is in

figure 5-9. The outer carbon coating layer of fiber and the ZrB2-SiC matrix part were

well bonded to each other. First of all, the reaction between fiber and matrix was

considered. The most possible is between ZrB2 and carbon, which is listed below.

2ZrB2 + 2C -> 2ZrC + B4C Reaction (1)

The standard Gibbs free energy of formation of each component in reaction (1) at

different temperatures (298K, 1000K, and 2000K) are listed in table 5-1. Based on

basic calculations of Gibbs free energy for each component in reaction (1), the

131

reaction (1) has low possibility because the Gibbs free energy of reaction (1) is

42.189 kcal/mol, which is a positive value.

In order to analyze the detailed interface between the fiber and the matrix, TEM

microstructures were investigated. Typical TEM microstructures and SAED (Selected

Area Electron Diffraction) of ZrB2 and SiC grain are shown figure 5-10. In addition,

it is noted that several SiC grains from the reaction bonding process have complex

structures. The complex stacking faults were found in the region of part 2 and two

twin boundaries were observed in the region of part 3 in figure 5-10. As no complex

structures were found in the ZrB2 grains, the complex SiC grain structures could be

from the reaction bonding process. The detailed TEM microstructure images of

interface between the fiber and the matrix are shown in figure 5-11. The desirable

interface of the composite is not well-bonded interface with chemical reaction

between the fiber and the matrix. The several regions in the interface part were

confirmed by the EDS and diffraction, which is not indicated in figure 5-11. It seems

that only carbon coating layer of the fiber, ZrB2 grain, ZrB2 and SiC agglomerate, and

ZrB2 with SiC grain were found in the interface without chemical reaction.

5.4 Conclusion

132

The SCS 9a fiber-reinforced ZrB2-RB SiC composite was investigated. The hot

pressing method with reaction bonding of SiC resulted in 10.40% porosity due to the

SiC reaction bonding process and large cracks due to the thermal expansion mismatch

of the fiber and the matrix materials, and hence the low hardness and fracture

toughness properties of the matrix. However, the fracture toughness of the composite

itself was improved compared to the fracture toughness of matrix materials. In

addition, it is noted that the composite did not separate into two separate pieces by the

four point bending test.

The microstructural analysis of the interface between the fiber and the matrix

materials was conducted. Based on thermodynamical calculations with Gibbs free

energies and TEM investigations with electron diffraction and EDS analysis, no

chemical reaction was observed in the interface between the fiber and the matrix

materials. In addition, complex SiC grain structures were observed, such as twin

boundary and stacking faults.

The SCS 9a fiber-reinforced ZrB2-RB SiC composite is a good approach for a

composite with ultra-high temperature applications, but several processes are needed

to improve the mechanical properties of the composite.

133

ZrB2

C

ZrC

B4C

A G 298

(Kcal/mol)

-76.049

0

-46.194

-9.151

A G IOOO

(Kcal/mol)

-73.225

0

-44.713

-8.782

A G 2000

(Kcal/mol)

-66.708

0

-42.028

-7.171

Table 5-1. Gibbs free energy of formation of ZrB2, C, ZrC, and B4C at 298K, 1000K,

and2000K[12].

134

So = 40mm

W = 4mm : a = 2 m m H

;0 3 - ^ 4mm Si = 20mm

B= 3mm

50mm

Figure 5-1. Detailed dimensions of SCS 9a fiber-reinforced composites for fracture

toughness.

135

Figure 5-2. FE-SEM images of SCS 9a fiber-reinforced composite; (a) Top surface

view, (b) Cross-sectional view.

136

S

10000

8000

6000

4000

2000

0

i

•

•

_

•

l Cross section

• ZrB2

• SiC

• •

1 1A 11 i * L 1 A

• i i i i i i

20 30 40 50 60 70 80 90 100

20

fr

s

=

12000

10000

8000

6000

4000

2000

0

-

-

1

1

•

- i

1

• A ...

i , i

1

• 1 1

i

Top surface

•ZrB2

• SiC

1?, i

•

1" •! 1 I i i i i i i i

20 30 40 90 50 60 70 80

20 Figure 5-3. Typical X-ray diffraction pattern of SCS 9a fiber-reinforced ZrB2-RB SiC

composite

137

I r-v h ing patte- n [*S/VJ ^* i

W^S ' $ & & & $

Figure 5-4. Typical indentation mark and cracking pattern of SCS 9a fiber composite.

138

* ^3L W;J*£=£Z3»'V%3BB£.

fes- "<SS. . ?S:

Figure 5-5. FE-SEM images of SCS 9a fiber-reinforced composite fracture.

139

-Pf» * * " * * * «

- ^ s ^ i f i S

Figure 5-6. Whole cross-section view of SCS 9a fiber-reinforced composite.

140

Figure 5-7. Typical FE-SEM images of cracks inside and outside the composite

141

Figure 5-8. FE-SEM Images of residual carbon from reaction bonding and EDS

result.

142

Aa • " ' *-5

Figure 5-9. The FE-SEM image of the interface between fiber and matrix.

143

Figure 5=10 Typical TEM microstructures and SAED (Selected Area Electron

Diffraction) of the composite

144

Figure 5-11. Typical TEM microstructural images and EDS results of the interface

between the fiber and the matrix; (a) Carbon coating layer of the fiber, (b) ZrB2 grain,

(c) ZrB2 and SiC agglomerate (d) ZrB2 and SiC grains.

145

References



Vol.90, Issue 5, pl347-1364

[2] Jiecai Han, Ping Hu, Xinghong Zhang, and Songhe Meng, (2007) Oxidation

behavior of zirconium diboride-silicon carbide at 1800 °C, Scripta Materialia 57,

p825-828.

[3] Ronald Loehman, Erica Corral, Hans Peter Dumm, Paul Kotula, and Raj an

Tandon, (2006) Ultra high temperature ceramics for hypersonic vehicle applications,

SAND 2006-2925

[4] M. M. Opeka, I. G. Talmy, and J. A. Zaykoski, (2004), Oxidation-based materials

selection for 2000 °C + hypersonic aerosurfaces: Theoretical considerations and

historical experience, J. Mater. Sci., 39, p5887-5904.

[5] Sufang Tang, Jingyi Deng, Shijun Wang, and Wenchuan Liu, (2007), Fabrication

and characterization of an ultra high temperature carbon fiber-reinforced ZrB2-SiC

matrix composite, J. Am. Ceram. Soc, 90, [10], p3320-3322.

[6] J. A. Dicarlo, H. M. Yun, and J. B. Hurst, (2004), Fracture mechanisms for SiC

fibers and SiC/SiC composites under stress rupture conditions at high temperatures,

Applied mathematics and computation, 152, p473-481.

[7] Nils Calussen, Suxing Wu, and Dietmar Holtz, (1994) Reaction bonding of

aluminum oxide (RBAO) composites: Processing, reaction mechanisms and

properties, J. of Euro. Ceram. Soc, 14, p97-109.

[8] Stanley R. Levine, Elizabeth J. Opila, Michael C. Halbig, James D. Kiser,

Mrityunjay Singh, and Jonathan A. Salem, (2002), Evaluation of ultra high

temperature ceramics for aeropropulsion use, J of Euro. Ceram. Soc, 22,

p2757-2767.

[9] R. Griffiths and C. Radford, (1965), Calculations in ceramics, Maclaren and sons

146

LTD, London, England, Chapter 4, Porous solids.

[10] ASTM C 1421-0lb, (1999), Standard test method for the determination of

fracture toughness of advanced ceramics at ambient temperatures, Annual book of

ASTM standards, V. 15. 01, American Society for Testing and Materials, West

Conshohocken, PA.

[11] J. Kovacik, (1999), Correlation between Young's modulus and porosity in porous

materials, J. Master. Sci., 18, pi007-1010.

[12] M. W. Chase, (1974), JANAF thermochemical tables

147

Chapter 6

Conclusions

In this dissertation, the investigation into ZrB2-based composites for

ultra-high-temperature applications is summarized in this chapter with the

conclusions of each previous research. It can be concluded that ZrE$2-based

composites are the most promising candidate materials for ultra-high-temperature

applications.

Chapter 1

The ultra-high-temperature ceramics was defined and brief development history

was explained. In addition, the applications for ultra-high-temperature ceramics were

introduced and then, the requirements for each application were discussed.

Chapter 2

An extensive and thorough literature survey of ultra-high-temperature ceramics

were made in chapter two. At first, how ZrB2-SiC ceramics were selected for

ultra-high-temperature ceramics was explained. Two important properties of

ZrB2-SiC ceramics, oxidation property and mechanical property, were explained with

the limitation of current research from other studies. The spark plasma sintering was

introduced with several advantages for nano and high temperature ceramics. Lastly,

liquid phase sintering mechanism was explained.

148

Chapter 3

ZrE$2 nano-sized SiC ceramics were investigated in chapter three. Spark plasma

sintering and liquid phase sintering effect enabled ZrB2 nano-sized SiC ceramics to

be fully dense. Incorporation of nano-sized SiC effectively hindered the grain growth

of ZrB2, but better dispersion process was needed to avoid the SiC agglomerates in

the ceramics. The hardness of the ceramics was dependent on composition, but the

fracture toughness was not related to the compositions. Each phase ZrB2 and SiC

grain was confirmed and the second phase, Zr(0, B)x, was observed by the TEM

microstructural analysis. In addition, the produced dislocations and stacking faults

were observed by the micropillar compression tests.

Chapter 4

ZrC-ZrB2-SiC ceramics were investigated in chapter four. The fully densed

ceramics were prepared by the spark plasma sintering and liquid phase sintering

effect. Elastic modulus, hardness, fracture toughness, thermal conductivity, and

electrical conductivity of ZrC-ZrB2-SiC ceramics decreased with increase of ZrC

content in the studied composition range. The fracture toughness of ZrC-ZrB2-SiC

ceramics was comparable with that of ZrB2 ceramics and that of ZrB2-SiC ceramics.

In addition, micropillar compression tests gave information about typical longitudinal

149

cracking behavior and generation of stacking faults.

Chapter 5

SCS 9a fiber-reinforced ZrB2-RB SiC composite was investigated in chapter five.

The composite was consolidated with conventional hot pressing with reaction

bonding of SiC. The prepared composite had high porosity and large amount of

cracks and these results affected to the mechanical properties of composite. However,

the fracture toughness of composite itself was four times higher than that of matrix

materials. The chemical reaction between fibers and matrix materials was not

observed based on thermochemical and TEM microstructural analysis. In addition,

reaction-bonded SiC was observed to complex structures.

150

abd2

Documents

fingers copper grid

eden prairie mn

transmission electron microscopy

edax energy dispersive

gibbs free energy

tungsten probe tips

point bending test

mesh screen size