Enhancement of Properties of Zirconia-Based

Ceramic Composites by Controlling the

Characteristics of Reinforcement Species

Muhammad Abdullah

Department of Metallurgy and Materials Engineering

Pakistan Institute of Engineering and Applied Sciences

Nilore, Islamabad, Pakistan

September 2012

Enhancement of Properties of Zirconia-Based

Ceramic Composites by Controlling the

Characteristics of Reinforcement Species

Muhammad Abdullah

A thesis submitted for partial fulfillment of the requirement for the degree of

Doctor of Philosophy (Ph.D.) in the Department of Metallurgy and Materials Engineering

Pakistan Institute of Engineering and Applied Sciences

Nilore, Islamabad, Pakistan

September 2012

ii

Declaration

I declare that any material in this thesis, which is not my own work, has been

identified and that no material has previously been submitted and approved for the

award of a degree by this or any other university.

Signature: ________________________

Author’s Name: Muhammad Abdullah

It is certified that the work in this thesis is carried out and completed under my

supervision.

Supervisor:

Dr. Mirza Jamil Ahmad

Pakistan Institute of Engineering and Applied Sciences (PIEAS)

Nilore, Islamabad

Co-Supervisor:

Dr. Mazhar Mehmood

Pakistan Institute of Engineering and Applied Sciences (PIEAS)

Nilore, Islamabad

iii

Certificate

Certified that the work contained in this thesis, is carried out by Mr. Muhammad

Abdullah under my supervision and that, in my opinion, it is fully adequate, in scope

and quality, for the degree of Doctor of Philosophy.

Dr. Mirza Jamil Ahmad

Supervisor

Submitted through:

Dr. Mazhar Mehmood

Head

Department of Metallurgy & Materials Engineering

Pakistan Institute of Engineering & Applied Sciences

Islamabad-45650, Pakistan

iv

International Journal Publications

1. M Abdullah, J Ahmad, M Mehmood, H Mujtaba ul and H Maekawa. "Synthesis

of Al2O3 whisker-reinforced yttria-stabilized-zirconia (YSZ) nanocomposites

through in situ formation of alumina whiskers," Ceram Int, 37[7] 2621-2624

(2011).

2. M Abdullah, J Ahmad, M Mehmood, Waqas H and M Mujahid. ―Effect of

Deflocculants on Hardness and Densification of YSZ-Al2O3(Whiskers &

Particulates) composites‖ Compos Part B-Eng, 43[3] 1564-1569 (2012).

3. M Abdullah, J Ahmad, M Mehmood. ―Effect of sintering temperature on

densification, microstructure and hardness of Al2O3 whisker-reinforced yttria-

stabilized-zirconia (TZ-3Y) nanocomposites‖ Compos Part B-Eng, 43[4] 1785-

1790 (2012).

4. M Abdullah, M Mehmood and J Ahmad. " Single Step Hydrothermal Synthesis

of 3D Urchin like Structures of AACH and Aluminum Oxide with Thin Nano-

Spikes" Ceram Int, 38[5] 3741–3745 (2012).

5. M Abdullah, J Ahmad and M Mehmood. ―Influence of Al2O3 Whisker

Concentration on Flexural Strength of the Al2O3(w)-ZrO2 (TZ-3Y) Composite‖

Ceram Int, 38[8] 6517–6523 (2012).

6. M Abdullah, J Ahmad, M Mehmood and S A Akbar. ―Low Temperature

Degradation of Al2O3 Whiskers Reinforced 3mol%Y2O3 Stabilized Tetragonal

ZrO2 Composites‖ J Compos Part B-Eng, Under Review.

7. M Abdullah, J Ahmad, M Mehmood, M Mujahid, S. S. Z. Zaidi. ―Hydroxyapatite

Containing, Alu mina Whiskers Reinforced TZ-3Y Composite for Dental Material

Application‖ J Mater Sci: Mater Med Under Review

International Conference Papers

1. M Abdullah, J Ahmad, M Mehmood and M Mujahid. “Effect of CTAB on

improvement of properties of Al2O3(w)–Al2O3(n)–ZrO2(n) (3mole Yttria stabilized

zirconia) Nanocomposite‖ in Proceedings of the International Conference on

Knowledge and Education Technology, Chengdu, China.

v

Table of Contents

Acknowledgements ............................................................................................................................... xii

Abstract ................................................................................................................................................ xiii

1 Introduction ................................................................................................................................... 1

References ........................................................................................................................................... 5

2 Literature Survey .......................................................................................................................... 7

2.1 Ceramic Materials.......................................................................................................................... 7 2.2 Zirconia and its Polymorphs .......................................................................................................... 7

2.2.1 Stabilization of Zirconia and its Advantages ......................................................................... 8 2.3 Applications of Zirconia Ceramics .............................................................................................. 10

2.3.1 Zirconia as a Dental Material ............................................................................................. 10 2.3.2 Zirconia Based Refractory Materials .................................................................................. 11 2.3.3 Zirconia in Cutting Tools Industry....................................................................................... 12

2.4 Composite Materials .................................................................................................................... 12 2.5 Ceramic Matrix Composites (CMCs) .......................................................................................... 12 2.6 Zirconia-Alumina Composites ..................................................................................................... 13 2.7 Fiber Reinforced Ceramic Composites ........................................................................................ 14 2.8 Hydrothermal Synthesis .............................................................................................................. 15

2.8.1 Brief History ........................................................................................................................ 15 2.8.2 What is Hydrothermal Synthesis .......................................................................................... 16 2.8.3 The Merits of Hydrothermal Synthesis ................................................................................. 18 2.8.4 Synthesis of Alumina Whiskers ............................................................................................ 18

2.9 Deflocculant and Dispersion of Reinforcement in Composite .................................................... 19 2.10 Hydroxyapatite for Improving Biocompatibility ....................................................................... 20

2.10.1 Hydroxyapatite and its Preparation .................................................................................. 20 2.10.2 Hydroxyapatite for Bio-Composites ................................................................................... 21

2.11 Solid-State Sintering .................................................................................................................. 22 2.11.1 Sintering of Whisker Reinforced Composites ..................................................................... 23

2.12 Concluding Remarks and Objective of the Thesis ..................................................................... 24 References ......................................................................................................................................... 26

3 Characterization Techniques ..................................................................................................... 38

3.1 A Brief Outline of Synthesis Processes ....................................................................................... 38 3.2 Characterization Techniques ....................................................................................................... 38

3.2.1 X-Ray Diffraction ................................................................................................................. 39 3.2.2 Scanning Electron Microscopy ............................................................................................ 40 3.2.3 Fourier Transform Infrared Spectroscopy (FTIR) ............................................................... 41 3.2.4 Thermogravimetric Analysis (TGA) ..................................................................................... 41 3.2.5 Mechanical Properties ......................................................................................................... 42 3.2.6 In-Vitro Biocompatibility Study ........................................................................................... 46

References ......................................................................................................................................... 48

4 Hydrothermal Synthesis of AACH Whiskers ........................................................................... 50

4.1 Experimental Procedure .............................................................................................................. 51 4.1.1 Hydrothermal Synthesis of AACH Whiskers ........................................................................ 51

4.2 Results and Discussion ................................................................................................................ 53 4.2.1 Fourier Transform Infrared Spectroscopy ........................................................................... 60

4.3 Conclusions ................................................................................................................................. 62 References ......................................................................................................................................... 63

5 Effect of Deflocculants on the Properties of Composite ........................................................... 66

5.1 Experimental Procedures ............................................................................................................. 68 5.1.1 Preparation of AACH Whiskers ........................................................................................... 68 5.1.2 Synthesis & Characterization of TZ-3Y+Al2O3 Composites ................................................ 68

5.2 Results ......................................................................................................................................... 70

vi

5.2.1 AACH Whiskers, Alumina Powder & Tetragonal Zirconia ................................................. 70 5.2.2 Effect of Deflocculants on Densification and Hardness of Composites ............................... 71

5.3 Discussion.................................................................................................................................... 75 5.3.1 Role of CTAB in Dispersion of Alumina .............................................................................. 75 5.3.2 Role of PABA in Dispersion of Alumina .............................................................................. 79

5.4 Conclusions ................................................................................................................................. 81 References ......................................................................................................................................... 82

6 Effect of Whisker Addition on Mechanical Properties of Composite ..................................... 85

6.1 Experimental Procedures ............................................................................................................. 86 6.2 Results and Discussion ................................................................................................................ 87

6.2.1 Whisker Concentration Effect on Hardness ......................................................................... 90 6.2.2 Whisker Concentration Effect on Flexural Strength ............................................................ 92

6.3 Conclusions ................................................................................................................................. 96 References ......................................................................................................................................... 97

7 Effect of Sintering Temperature on Morphology of Whisker and Properties of Composite 99

7.1 Experimental Procedures ............................................................................................................. 99 7.2 Results and Discussion .............................................................................................................. 101

7.2.1 Microstructure Analysis ..................................................................................................... 101 7.2.2 Effect of Sintering Temperature and Whisker Concentration on Density .......................... 103 7.2.3 Effect of Sintering Temperature on Whisker Morphology ................................................. 105 7.2.4 Effect of Sintering Temperature and Whisker Concentration on Hardness ....................... 107

7.3 Conclusions ............................................................................................................................... 111 References ....................................................................................................................................... 112

8 Low Temperature Tetragonal Phase Stability of TZ-3Y Composite Reinforced with

Alumina Whiskers ..................................................................................................................... 114

8.1 Experimental Procedures ........................................................................................................... 115 8.2 Results and Discussion .............................................................................................................. 117

8.2.1 Effect of Sintering Temperature on LTD ............................................................................ 119 8.3 Conclusions ............................................................................................................................... 126 References ....................................................................................................................................... 127

9 Effect of HAp Addition on Biocompatibility of the Composite ............................................. 129

9.1 Experimental Procedure ............................................................................................................ 130 9.1.1 Preparation of AACH Whiskers ......................................................................................... 130 9.1.2 Preparation of HAp Nanoparticles .................................................................................... 130 9.1.3 Formation of TZ-3Y+Al2O3+HAp Nanocomposites .......................................................... 130 9.1.4 Cell Adhesion ..................................................................................................................... 132

9.2 Results and Discussion .............................................................................................................. 132 9.2.1 Surface Morphologies ........................................................................................................ 132 9.2.2 Fracture Mechanism .......................................................................................................... 134 9.2.3 Densification of Composite ................................................................................................ 135 9.2.4 Effect of Whisker Addition on Hardness and Flexural Strength of Composite .................. 137 9.2.5 X-Ray Diffraction Analysis ................................................................................................ 141 9.2.6 Cell Adhesion Biocompatibility Test .................................................................................. 144

9.3 Conclusions ............................................................................................................................... 147 References ....................................................................................................................................... 148

10 Conclusions ................................................................................................................................ 151

vii

List of Figures

Fig. 2.1(a-c): Illustration of three polymorphs of ZrO2; (a) cubic; (b) tetragonal and

(c) monoclinic .......................................................................................................................... 8 Fig. 2.2: Schematic description of transformation toughening at microstructural level. .........................10 Fig. 2.3: Micrograph showing neck formation during initial stage of the sintering. ................................23 Fig. 3.1: Schematic of Density measurement equipment through Archimedes principle. .......................43 Fig. 3.2: Ring on Ring test arrangement, for determining Bi-Axial flexural strength. ............................45 Fig. 4.1(a-c): (a) Laboratory scale high pressure reactor (Limbo 350) by Büchi Glasuster,

Switzerland used for hydrothermal synthesis; (b) Corrosion resistant silver plated

metallic seal O-ring; (c) Heating unit. ........................................................................................52 Fig. 4.2(a-d): SEM images of hydrothermally synthesized AACH urchins (a) U6 at low

magnification; (b) U6 at high magnification; (c) U8 at low magnification; (d) U8

at high magnification. ................................................................................................................53 Fig. 4.3 (a-f): (a) Low magnification image of U10 sample showing a number of urchins-

like particles; (b-e) High magnification images of various micro-urchins showing

size and morphology of individual urchins; (f) Higher magnification image of

individual thorn of a micro-urchin. ............................................................................................54 Fig. 4.4 (a-d): (a) Low magnification image of whiskers formed in U12 sample; (b-d)

High magnification images of sample U12 showing individual whiskers and their

agglomerates ..............................................................................................................................55 Fig. 4.5 (a-b): Showing octahedron of the Al-O; (a) Isometric view ......................................................57 Fig. 4.6 (a-d): Schematic illustration of AACH whiskers and micro-urchin formation in

samples containing different amounts of urea; (a) the U6 sample; (b) U8 sample;

(c) U10 sample and (d) U12 sample ..........................................................................................58 Fig. 4.7: Thermal gravimetric curve of U10 sample. ...............................................................................59 Fig. 4.8(a-b): SEM images of hydrothermally synthesized U10 AACH urchins calcined at

650ºC; (a) Low magnification; (b) High magnification .............................................................60 Fig. 4.9: FTIR spectra of U10 sample calcined at different temperatures; the spectrum of

as prepared sample is also given for comparison. ......................................................................61 Fig. 4.10: XRD patterns of U10 sample calcined at different temperatures; the spectrum of

as prepared sample is also given for comparison. ......................................................................62 Fig. 5.1: Morphology of as received α-Al2O3 particulate .........................................................................70 Fig. 5.2(a-c): X-Ray diffraction of (a) as received TZ-3Y, (b) sintered TZ-3Y and (c) TZ-

3Y+Al2O3(w) pellets. ...................................................................................................................71 Fig. 5.3(a-b): Hardness of pellets (a) [90%TZ-3Y+10% Al2O3]100-x+[CTAB]x (b)

[90%TZ-3Y+10% Al2O3]100-x+[PABA]x. ...................................................................................72 Fig. 5.4(a-b): Relative density of pellets (a) [90%TZ-3Y+10% Al2O3(w)]100-x+[CTAB]x

(b) [90%TZ-3Y+10% Al2O3(w)]100-x+[PABA]x. .........................................................................73 Fig. 5.5(a-d): SEM micrographs at low and high magnification of the sintered pellets of

(a&b) [90%TZ-3Y +10%Al2O3(p)]99+[CTAB]1 (c&d) [90%TZ-3Y

+10%Al2O3(w)]99+[CTAB]1. ........................................................................................................74 Fig. 5.6(a-b): Fractograph of composite (a) at lower magnification (b) at higher Magnified

identifying the presence of whisker in composite matrix. ..........................................................75 Fig. 5.7(a-b): (a) FTIR Spectrum for samples containing 20wt% alumina inform of

AACH, with and without 1wt% CTAB, before and after calcinations. (a-i) AACH

with CTAB calcined at 650ºC (a-ii) AACH without CTAB calcined at 400ºC

(a-iii) AACH with CTAB uncalcined (a-iv) AACH without CTAB uncalcined.

(b) magnification of the FTIR peak at wave number of 3437cm-1. ...........................................77 Fig. 5.8: Structure of PABA molecule .....................................................................................................78 Fig. 5.9(a-c): TGA of (a) pure AACH and [90%TZ-3Y+10%Al2O3(W)]75with (b) CTAB

and (c) PABA Composite. .........................................................................................................80 Fig. 6.1(a-b): (a) XRD patterns of sintered ZrO2+Al2O3(w) nanocomposites with various

amounts of Al2O3 whiskers; (b) Slow scan rate XRD pattern around (012)

reflection of α-Al2O3 for sample with 10 wt.% Al2O3 whiskers. ...............................................88 Fig. 6.2(a-f): SEM micrographs of: (a) & (b) as-fractured 90% ZrO2(3Y)-10% Al2O3(W)

pellet along with (c) EDS spectra on a 1.5µm × 1.5µm area in (a) and a

(d) whisker in (b); (e) & (f) polished and thermal etched surfaces of samples with

10 wt% and 50 wt% alumina whiskers, respectively. ................................................................89

viii

Fig. 6.3(a-b): (a) Relative densities and (b) Vicker hardness of sintered pellets having

various amounts of alumina whiskers, with & without CTAB. For comparison

hardness values from Nevarez-Rascon et al. are also given. ......................................................91 Fig. 6.4(a-b): Flexural strength of the TZ-3Y+Al2O3 composite with different weight

fraction of Al2O3 whiskers: (a) with CTAB and (b) without CTAB. .........................................93 Fig. 6.5: SEM micrograph of crack formed at the corner of the Vicker hardness indenter

on the sintered surface of the TZ-3Y+10wt%Al2O3 composite sample, showing

crack deflection and crack bridging effects. The morphology encircled on the

micrograph is a ligamentary bridge between two surfaces of the crack formed by

alumina whisker. ........................................................................................................................94 Fig. 6.6(a-b): Relative density of the TZ-3Y+Al2O3 composite with different weight

fraction of Al2O3 whiskers: (a) with CTAB and (b) without CTAB. .........................................95 Fig. 7.1(a-c): SEM micrographs showing size and morphology of the inter-granular pores

for 10wt% Al2O3(w)+90wt%TZ-3Y composite sample surface sintered at

temperature of (a) 1400ºC; (b) 1500ºC; (c) 1650ºC. ................................................................102 Fig. 7.2(a-b): Relative densities of sintered composite samples having varying amount of

alumina whiskers; (a) without CTAB; (b) with CTAB............................................................104 Fig. 7.3(a-f): SEM micrographs and EDS of sample surface sintered at 1650ºC;

(a) containing 5wt% Al2O3(w) and (b) 15wt% Al2O3(w); (c) Spot EDS of dark

grain; (d) Spot EDS of the white grain; (e) EDS spectra on aprox. 8µm×8µm area

(f) Spot EDS of less dark grain. ...............................................................................................106 Fig. 7.4(a-d): SEM micrographs of fractured composite samples containing 10wt%

Al2O3(w) and sintered at (a) 1400ºC;(b) 1500ºC; (c) 1650ºC (low magnification

image) (d) 1650ºC (high magnification image). ......................................................................107 Fig. 7.5(a-b): Vicker hardness of sintered pellets having varying amount of alumina

whiskers; (a) without CTAB; (b) with CTAB. ........................................................................108 Fig. 7.6(a-c): SEM micrographs of Vicker indentation showing the crack propagation path

at one arm of the indent impression for samples with 10wt% Al2O3(w) and sintered

at (a) 1400ºC; (b) 1500ºC; (c) 1650ºC. ....................................................................................110 Fig. 8.1(a-b): XRD patterns of composite samples containing pure TZ-3Y and 20wt%

Al2O3(w) (a) sintered at 1400, 1500 and 1650ºC after hydrothermal aging at135ºC

for 10 hr; (b) sintered at 1400ºC and without hydrothermal aging. .........................................118 Fig. 8.2: Amount of monoclinic phase produced due to hydrothermal treatment of samples

containing 0, 10 & 20 wt%Al2O3(w) in the matrix of TZ-3Y. ...................................................120 Fig. 8.3: Dilatometer curves of (10 and 20wt%) Al2O3(w) containing TZ-3Y samples. .........................121 Fig. 8.4(a-e): Micrographs of the hydrothermally synthesized AACH whiskers (a&b)

SEM micrographs at (a) lower and (b) higher magnifications; (c-e) TEM

micrographs at (c) lower and (d&e) high magnification. .........................................................122 Fig. 8.5: Backscattered image of hydrothermally aged TZ-3Y+20wt% Al2O3(w) composite

sample sintered at 1650ºC. .......................................................................................................123 Fig. 8.6: Density of the Al2O3 whisker reinforced TZ-3Y composite samples sintered at

different temperatures. .............................................................................................................124 Fig. 8.7(a-b): SEM images of the surface of TZ-3Y+20wr% Al2O3(w) composite sample

sintered at (a) 1400ºC; (b)1650ºC. ...........................................................................................125 Fig. 9.1(a-f): Micrograph of sintered surface of composited samples containing different

concentration of HAp (a) Pure HAp; (b) CT-ZA(w); (c) 10HA-CT-ZA(w);

(d) 20HA-CT-ZA(w); (e) 30HA-CT-ZA(w); (f) 40HA-CT-ZA(w). .............................................133 Fig. 9.2(a-h): SEM micrographs of fractured surface of composite samples with different

compositions (a) pure ZA(w); (b) 20HA-CT-ZA(w); (c) pure HAp; (d) areal EDS

of pure HAp; (e) 30HA-CT-ZA(w); (f) 40HA-CT-ZA(w). .........................................................135 Fig. 9.3(a-b): Relative density of sintered pellets of ZrO2+Al2O3 (whiskers or particulates)

having different amount of HAp (a) without CTAB. (b) with CTAB. ....................................136 Fig. 9.4(a-b): Vickers hardness of sintered pellets of ZrO2+Al2O3 (whiskers or

particulates) having different amount of HAp (a) with 1wt% CTAB. (b) without

CTAB .......................................................................................................................................138 Fig. 9.5(a-b): Flexural strength of sintered pellets of ZrO2+Al2O3 (whiskers or

particulates) having different amount of HAp (a) without CTAB. (b) with 1wt%

CTAB HAp. .............................................................................................................................139 Fig. 9.6(a-b): SEM images showing the path of crack propagation at corner of Vickers

hardness indent on sintered pellet of (a) pure ZA(w), (b) 30H-CT-ZA(w). ................................140

ix

Fig. 9.7(a-d): XRD patterns of samples; (a) pure ZA(w) composite sintered at 1400ºC;

(b) pure HAp sintered at 800 & 1400ºC; (c) 20, 30 & 40H-CT-ZA(w);

(d) magnified image of 40H-CT-ZA(w). ...................................................................................143 Fig. 9.8(a-e): Micrograph of surface of pellets after cell culturing (a) ZA(w); (b) 10H-CT-

ZA(w); (c) 20H-CT-ZA(w); (d) 30H-CT-ZA(w); (e) 40H-CT-ZA(w). ..........................................145 Fig. 9.9(a-c): Micrograph of surface of pellets after cell culturing (a&b) 40H-CT-ZA(w)

after further 48 hours; (c) 30H-CT-ZA(w) after further 48 hours. .............................................146

x

List of Tables

Table 4.1: Change in pH of solution with amount of urea added in the precursor ..................................55 Table 5.1: Types of samples prepared .....................................................................................................69 Table 7.1: Average grain size of samples sintered at different temperatures. .......................................101 Table 8.1: Reflections as used in Toraya’s method ...............................................................................117 Table 9.1: Sample designations for various compositions ....................................................................131

xi

Dedicated to

My Honorable Parents and sweet family

xii

Acknowledgements

First, I whole-heartedly thanks to ALLAH Almighty, Who gave me courage, health,

and energy to accomplish my dissertation. No doubt, without His help, this study,

which required untiring efforts, would have not been possible to complete. Motivation,

encouragement, guidance, corrections, advices and overall support are the key

elements required from the supervisor(s) to write and complete a thesis of good

standard and quality. It is a matter of utmost pleasure for me to extend my gratitude

and give due credit to my supervisor Dr. Mirza Jamil Ahmad whose support has

always been there in time of need. Acknowledgement would be incomplete without

extending my gratefulness to one of my most respected teacher Dr. Mazhar Mehmood,

Head of the Metallurgy and Material Engineer Department, PIEAS, whose technical

and moral support at every stage made this study possible.

I would also like to pay my gratitude to Dr. Sheikh Ali Akbar, Professor, The Ohio

State University (OSU), USA for his help and valuable guidance during my stay at

OSU.

My thanks are also due to Dr. Hassan Waqas, PINSTECH for his colossal contribution

in form of suggestions and encouragement to improve the quality of research and

thesis dissertation.

I am also thankful to Higher Education Commission (HEC), Government of Pakistan

for providing me the financial support to carry out this research activity.

Finally, yet importantly, I would like to extend many thanks to my entire family for

moral support and prays for my health and success in completing this study.

Muhammad Abdullah

August 2012

xiii

Abstract

The purpose of the present work was to develop high strength ZrO2 based composites

with the use of Al2O3 whiskers as reinforcement, which could be able to meet the

modern needs of load bearing structural and biomedical materials. For this purpose,

3mol% Y2O3 doped tetragonal (TZ-3Y) was used as the matrix material. The alumina

whiskers were formed in situ during calcination and sintering from Aluminum

Ammonium Carbonate Hydroxide (AACH) whiskers, which were in turn produced

by hydrothermal synthesis technique using urea and aluminum nitrate as precursor

materials. It was found that the morphology of AACH structures was dependent on

the urea content forming urchin like structures at lower urea concentration, which

transformed into whiskers as the urea content was increased. After the preparation of

alumina whisker reinforcement, the second most important study was the

optimization of alumina whisker content in the Al2O3(w)-TZ-3Y composites to

achieve best mechanical properties. With increasing the whisker concentration, the

hardness increased up to 10 wt% addition and then decreased. It was because the

higher concentrations of whiskers resisted particle rearrangement resulting in

introduction of porosity in sintered products. Consequently, 10wt% alumina whisker

content was taken as optimum. To improve the uniform distribution of alumina

whiskers in the matrix further, deflocculants were also employed and their optimized

concentrations were determined. 1.0 wt% cetyltrimethylammonium bromide (CTAB)

gave the best results as a dispersant. Optimization of sintering temperature was also

very important. It was observed that 1400oC was a too low temperature for complete

sintering process and resulted in poor mechanical properties of the sintered product,

which were attributed to incomplete removal of porosity. On the other hand,

temperatures higher than 1500oC were too high. At 1650

oC, the whiskers were

diffused and formed alumina rich grains losing their whisker-like morphology.

However, the best mechanical properties were observed for the sample sintered at

1500oC, which was decided as the optimized sintering temperature for Al2O3(w)-TZ-

3Y composite. Effect of sintering temperature on low temperature phase stability of

the composite was also studied. Higher sintering temperature was useful for the

hydrothermal stability of tetragonal phase in whisker-reinforced composites,

although such sintering temperatures were deleterious for tetragonal phase stability in

monolithic TZ-3Y. A composition 10 wt% Al2O3(w)-TZ-3Y with 1wt% CTAB,

xiv

sintered at 1500ºC was concluded as the best composite having better reinforcement

dispersion, high density, excellent mechanical properties and improved hydrothermal

stability of tetragonal phase.

Furthermore, to investigate the possibility of using the aforementioned optimized

composite as bioactive material, its biocompatibility was improved by incorporating

the Hydroxyapatite (HAp) nanoparticles, synthesized through precipitation technique,

into the composites. It was observed, that although the increase in the HAp content

resulted in improving the biocompatibility but it was deleterious for mechanical

properties of the composite. The optimized sintering temperature was also lowered due

to decomposition of HAp at higher temperatures. Consequently, a composite with

30wt% HAp sintered at 1400oC was assessed as the optimized composition for bio-

composite applications.

1

Chapter 1

1 Introduction

Due to variety of applications in the fields like electronic and optoelectronic

industry (semiconductors, insulators, magnetism, superconductors, gas-sensors, fiber

optics, energy storage), bio-ceramics (orthopedic, dental implants), high strength

materials, refractory materials, automobiles and defense industry (radomes of missiles),

demand for ceramics has seen a tremendous growth. Among various conventional

ceramics, zirconia occupies a unique place due to its mechanical properties, chemical

inertness (stability in both oxidizing and reducing environments), low cost and easy

availability. However, large volume change (~ 4%) associated with its transformation

from low temperature monoclinic to high temperature tetragonal phase (during

processing) makes it susceptible to catastrophic failure. This phase transformation

precludes the use of unstabilized zirconia for many applications [1]. In 1975, Garvie

suggested that by stabilizing high temperature tetragonal phase of zirconia to room

temperature (with the help of some stabilizing agent like yttria, ceria, calcia), its

mechanical properties can be significantly improved. He used the term ―ceramic steel‖

to emphasize excellent properties of stabilized zirconia [2]. Stabilized zirconia lends its

highly improved mechanical properties from stress induced transformation toughening

phenomenon [3]. This stabilization also induces thermal stability to zirconia. Among

various stabilized zirconia compositions, 3mol% Y2O3 stabilized tetragonal ZrO2 (TZ-

3Y) has gained wide spread applications in solid oxide fuel cells [4-5], dental implants

[6-7] and refractories [8]. Although TZ-3Y has excellent mechanical properties, studies

have shown that on exposure to hydrothermal conditions (such as in human body), it

may undergo severe strength degradations (also termed as low temperature thermal

degradation of zirconia) [9-11]. One of the approaches adopted for preventing this

degradation are use of some stable surface coating or formation of zirconia-alumina

composite [12]. However, due to the problem of peeling off or disassociation of surface

coatings from base metal [13], suitable choice is formation of zirconia-alumina

composite. Furthermore, addition of alumina also enhances the mechanical properties of

TZ-3Y.

2

It is well known that fracture toughness or flexural strength of composites

strongly depend upon morphology of the reinforcement [14]. This is the reason that to

improve strength of ceramic materials there has been extensive studies on incorporation

of strong ceramic whiskers into ceramic matrix [15-17]. Due to their structural

perfection, ceramic whiskers have excellent strength and toughness. Whisker-

reinforcement also imparts special toughening mechanisms such as crack deflection,

crack bridging and crack bowing etc. [18] to help improve flexural strength and fracture

toughness of the formed composite [19]. Therefore, a better choice may be use of

alumina in form of whiskers for formation of zirconia-alumina composite.

In addition to the type of reinforcement, processing is also a significantly

important step towards achieving a composite with good mechanical properties. During

processing main problem is agglomeration of reinforcement species. The aggregates act

as preferential sites for cracks initiation [20]. For very small sized whiskers (e.g., a few

tens of nanometers), it will almost be impossible to avoid agglomeration. Relatively

larger sized whiskers (e.g., 120 nm or so) could be a better choice. That is why in the

present study, choice of 120 nm dia. whiskers is made in order to minimize the chances

of agglomeration. Secondly, use of some suitable low temperature volatile

deflocculating agents for dispersing the reinforcement [21] may also be helpful. Vishista

et al. [21] have checked number of dispersing agents for alumina dispersion and

concluded that para-aminobenzoic acid (PABA) is the best for breaking alumina

agglomerates. Similarly Liu et al. have shown that by using cetyltrimethylammonium

bromide (CTAB) as deflocculant for alumina dispersion, properties of the produced

composite improve significantly [22]. Moreover, both of these compounds are volatile

at a temperature of about 650ºC that is much lower than the composite sintering

temperature. Therefore, these deflocculants may be removed easily through calcination

before sintering. In the present study, both of these dispersing agents were investigated

for dispersion of alumina whiskers in TZ-3Y matrix and the amounts of alumina

whiskers and deflocculants were optimized to obtain best properties.

Among processing parameters, uniform mixing of ceramic powders, their

compaction and sintering are other important issues to be addressed. To obtain dense,

pore free and fine grained microstructure, selection of suitable sintering temperature is

very important [23]. Specially, for whisker reinforced composites it is reported that at

lower sintering temperature the retained porosity and at higher temperature the change

in whisker morphology or aspect ratio severely affect mechanical properties of the

3

composite [24]. Consequent on the importance of sintering temperature selection, in the

present study, different sintering temperatures were investigated and an optimized

sintering temperature was suggested.

As a part of the present study, possibility of using the alumina-whisker-

reinforced zirconia composite for formation of bioactive composite was also

investigated. For this purpose, hydroxyapatite (HAp) was incorporated to improve

Osseointegration properties of the composites. Bioactive ceramics like tricalcium

phosphate or hydroxyapatite have poor mechanical but much better bio-medical

properties. On the other hand, yttria stabilized zirconia and alumina are bio-inert with

excellent mechanical properties in composite form. None of the components in these

composites is bio-toxic. Consequently, a composite having yttria-stabilized zirconia as

matrix and alumina whiskers as reinforcement with hydroxyapatite added to it could be

an excellent choice for load bearing biomedical applications such as artificial orthopedic

and dental implants. While exploring the possibility of using synthesized composite for

biomaterial applications, response of human body to the introduced material should also

be investigated. For this purpose, biocompatibility tests were performed to observe

response of live cells to the composite when placed in cell culturing media. Since an

excess of hydroxyapatite might cause undue deterioration of mechanical properties of

the composite, amount of hydroxyapatite was optimized for a good compromise

between mechanical properties and biocompatibility.

What distinguishes this work from previous works on the same subject is the

introduction of alumina whiskers instead of particulates. For production of alumina

whiskers, an innovative and quite simple hydrothermal synthesis technique was

employed. Initially, hydrothermally produced Ammonium Aluminum Carbonate

Hydroxide (AACH) whiskers were mixed with other components (i.e. yttria stabilized

zirconia and hydroxyapatite). This mixed powder was then calcined, compacted and

sintered for composite formation. It was observed that added AACH whiskers were

transformed into α-alumina during sintering. It was also found that addition of alumina

in form of whiskers imparts better mechanical properties to the composite than its

addition in form of particulates. Moreover, use of deflocculating agent for improving

dispersion of reinforcement in ceramic matrix further enhanced the properties. Other

worth mentioning distinction of this work is to find effect of sintering temperature on

morphology of alumina whiskers and its correlation with change in mechanical

properties of the composite. Finally, the possibility of the use of developed composites

4

for biomedical applications through improvement in biocompatibility was also

investigated.

The thesis has been divided into ten chapters. Chapter 1 gives a brief

introduction of the thesis. In chapter 2, a literature survey covering historic background

of composites and ceramics, along with different synthesis techniques etc. has been

presented. Chapter 3 provides a general overview of the experimental techniques used in

this work. The synthesis techniques are not discussed at this stage and are left to be

described in corresponding chapters, while characterization techniques have been

introduced in this chapter. Details about production of Aluminum Ammonium

Carbonate Hydroxide (AACH) whiskers through hydrothermal synthesis, effect of

different concentrations of precursor materials on the produced nano-morphologies and

change in AACH nanowhisker morphology with increasing urea content of solution has

been given in Chapter 4. Chapter 5 is concerned with the study of optimizing Al2O3(w)

concentration in Al2O3(w)-TZ-3Y composites for excellent mechanical properties. In

Chapter 6, selection of suitable dispersing agent and its optimum concentration for

dispersing the reinforcement (i.e. Al2O3(w)) to achieve best composite properties have

been covered. Chapter 7 is related to the evaluation of sintering temperature effect on

morphology of alumina whiskers in TZ-3Y matrix and its correlation with composites’

mechanical properties. Effect of sintering temperature on low temperature degradation

of tetragonal zirconia in TZ-3Y+Al2O3(w) composite, under mist environment during

hydrothermal treatment (mimicking the human body environment according to ASTM

standard F2345-03 [25]) is presented in chapter 8. In chapter 9, synthesis of HAp

nanoparticles through precipitation technique and their mixing with Al2O3(w) and TZ-3Y

for forming composite is discussed. Also in this chapter, optimization of HAp

concentration for obtaining a biocompatible material with acceptable mechanical

properties was investigated. Finally, chapter 10 gives a brief review of main conclusions

drawn from the present work.

5

References

[1] Anthony AM, Yeroucha.D. New Ceramics Based on Monoclinic Zirconia. Phil

Trans R Soc A. 1967;261(1123):504-13.

[2] Garvie RC, Hannink RH, Pascoe RT. Ceramic steel. Nature. 1975;258(5537):703-4.

[3] Porter DL, Evans AG, Heuer AH. Transformation-toughening in partially-stabilized

zirconia (PSZ). Acta Metall. 1979;27(10):1649-54.

[4] Wei CC, Li K. Yttria-Stabilized Zirconia (YSZ)-Based Hollow Fiber Solid Oxide

Fuel Cells. Ind Eng Chem Res. 2008;47(5):1506-12.

[5] Brown M, Primdahl S, Mogensen M. Structure/Performance Relations for Ni/Yttria-

Stabilized Zirconia Anodes for Solid Oxide Fuel Cells. J Electrochem Soc.

2000;147(2):475-85.

[6] Oliva J, Oliva XV, Oliva JD. One-year follow-up of first consecutive 100 zirconia

dental implants in humans: A comparison of 2 different rough surfaces. Int J Oral

Maxillofac Implants. 2007;22(3):430-5.

[7] Manicone PF, Rossi Iommetti P, Raffaelli L. An overview of zirconia ceramics:

Basic properties and clinical applications. J DENT 2007;35(11):819-26.

[8] Andrievskaya ER. Phase equilibria in the refractory oxide systems of zirconia,

hafnia and yttria with rare-earth oxides. J Eur Ceram Soc. 2008;28(12):2363-88.

[9] Chevalier J, Gremillard L, Deville S. Low-Temperature Degradation of Zirconia

and Implications for Biomedical Implants. Annu Rev Mater Res. 2007;37(1):1-32.

[10] Yoshimura M, Noma T, Kawabata K, Somiya S. Role of H2O on the Degradation

Process of Y-TZP. J Mater Sci Lett. 1987;6(4):465-7.

[11] Lange FF, Dunlop GL, Davis BI. Degradation during Aging of Transformation-

Toughened ZrO2–Y2O3 Materials at 250°C. J Am Ceram Soc. 1986;69(3):237-40.

[12] Schneider J, Begand S, Kriegel R, Kaps C, Glien W, Oberbach T. Low-

Temperature Aging Behavior of Alumina-Toughened Zirconia. J Am Ceram Soc.

2008;91(11):3613-8.

[13] MacDonald DE, Betts F, Stranick M, Doty S, Boskey AL. Physicochemical study

of plasma-sprayed hydroxyapatite-coated implants in humans. J Biomed Mater Res.

2001;54(4):480-90.

[14] Hartmann O, Kemnitzer M, Biermann H. Influence of reinforcement morphology

and matrix strength of metal–matrix composites on the cyclic deformation and fatigue

behaviour. Int J Fatigue. 2002;24(2–4):215-21.

6

[15] Zhang X, Xu L, Du S, Han J, Hu P, Han W. Fabrication and mechanical properties

of ZrB2-SiCw ceramic matrix composite. Mater Lett. 2008;62(6-7):1058-60.

[16] Zhang X, Xu L, Han W, Weng L, Han J, Du S. Microstructure and properties of

silicon carbide whisker reinforced zirconium diboride ultra-high temperature ceramics.

Solid State Sci 2009;11(1):156-61.

[17] Zhang XH, Xu L, Du SY, Han WB, Han JC. Crack-healing behavior of zirconium

diboride composite reinforced with silicon carbide whiskers. Scr Mater.

2008;59(11):1222-5.

[18] Belitskus D. Fiber and whisker reinforced ceramics for structural applications: M.

Dekker; 1993.

[19] Xu HHK, Martin TA, Antonucci JM, Eichmiller FC. Ceramic whisker

reinforcement of dental resin composites. J Dent Res 1999;78(2):706-12.

[20] Kissinger-Kane MC. Investigation and characterization of the dispersion of

nanoparticles in a polymer matrix by scattering techniques [Dissertation]. Florida:

University of Florida; 2007.

[21] Vishista K, Gnanam FD. Effect of deflocculants on the densification and

mechanical properties of sol-gel alumina. Mater Lett. 2004;58(11):1665-70.

[22] Liu D, Yan Y, Lee K, Yu J. Effect of surfactant on the alumina dispersion and

corrosion behavior of electroless Ni-P-Al2O3 composite coatings. Mater Corros.

2009;60(9):690-4.

[23] Handwerker CA, Blendell JE, Kaisser W, Meeting ACS. Sintering of advanced

ceramics: American Ceramic Society; 1990.

[24] Zhu SJ, Iizuka T. Fatigue behavior of Al18B4O33 whisker-framework reinforced Al

matrix composites at high temperatures. Compos Sci Technol. 2003;63(2):265-71.

[25] ASTM.org. Standard Test Methods for Determination of Static and Cyclic Fatigue

Strength of Ceramic Modular Femoral Heads. West Conshohocken, PA2008.

7

Chapter 2

2 Literature Survey

2.1 Ceramic Materials

The origin of ceramic is derived from Greek word ―Keramos‖, used for clays [1].

History of ceramics can be traced back since 30,000 year ago. From that time, the

ceramic products like bricks, tiles and cements were invented and have played vital role

in the evaluation and development of human civilization.

Modern technical ceramics can be divided into two main categories, oxides (like

ZrO2, Al2O3, MgO etc) and non-oxide ceramics (SiC, ZnS, Si3N4 etc). Out of these two,

the former type is famous due to its good oxidation resistance, chemical inertness and

electrical insulating properties, and extensively used in many applications like

refractory, structural and abrasive materials, etc. In this class, the zirconia has unique

position owing to its relatively better toughness and other mechanical properties.

However, the volume change associated with its high temperature phase transformation

precludes its use in pure form. The detail of the polymorphic phases of zirconia and

techniques to stabilize these phases are discussed below.

2.2 Zirconia and its Polymorphs

At atmospheric pressure, pure ZrO2 exhibits three crystalline forms. Liquid zirconia

freezes into cubic fluorite type structure at 2680°C, a transition from cubic to tetragonal

structure takes place at 2370°C, while tetragonal zirconia transforms to monoclinic

structure at 1170°C. The unit cells of three crystal structures are shown in Fig. 2.1.

Zirconia with completely monoclinic structure is also known as unstabilized

zirconia and has very poor toughness. If high temperature cubic phase is stabilized at

room temperature, it is called fully stabilized zirconia and offers improved toughness. If

full stabilization is not achieved, the resultant structure may consist of a mixture of

cubic, tetragonal and monoclinic phases. This is known as partially stabilized zirconia

and tetragonal phase in it is responsible for transformation toughening mechanism. Pure

tetragonal phase may also be stabilized.

8

Fig. 2.1(a-c): Illustration of three polymorphs of ZrO2; (a) cubic; (b) tetragonal and (c) monoclinic [2]

2.2.1 Stabilization of Zirconia and its Advantages

The metastable tetragonal or cubic structure of zirconia can be stabilized at room

temperature by doping it with cubic oxides like magnesia (MgO), calcia (CaO) or rare-

earth oxides like yttria (Y2O3), ceria (CeO2), etc. The stabilization is mainly based on

the creation of oxygen vacancies due to substitution of dopant atoms for Zr in lattice.

However, other factors such as stress, surface energy and fine particle size have also

been shown to play their role in stabilization of high temperature phases at room

temperature. Oxygen vacancies are produced in order to compensate for charge

neutrality as dopants generally used for stabilization have valence lower than zirconium

(Zr4+

); for example, yttrium has +3 valence. Yttrium is most widely used stabilizer and

two yttrium atoms will be associated with one oxygen vacancy. Actually, in ZrO2

lattice, Zr-O bond has a strong ionic character and 8-fold coordination of a Zr4+

cation

by O2-

anions (required by tetragonal and cubic structures) is not favored at room

temperature due to smaller size of Zr4+

cation. Resultantly, it likes to assume a 7-fold

coordination offered by monoclinic structure; therefore, monoclinic phase is stable at

room temperature. When Y3+

cations occupy Zr4+

sites in the lattice and to balance the

charge oxygen vacancies are produced. Since, ionic size of Y3+

is larger than Zr4+

,

vacancies are more associated with Zr4+

cations resulting in reduction of effective

coordination number of Zr4+

below seven (near six). Consequently, ZrO2 lattice adopts a

crystal structure with 8-fold coordination (i.e., cubic or tetragonal), so that after

incorporation of oxygen vacancies, coordination number becomes near seven [3].

Therefore, it is creation and association of O2-

vacancies with Zr4+

cations which is

9

responsible for stabilization of high temperature phases at room temperature [4]. This is

also supported by the fact that high temperature tetragonal phase can be formed by

exposing monoclinic ZrO2 to low oxygen partial pressures [5]. At lower concentration

of the oxygen vacancies (associated with 2–3 mol% of Y2O3) tetragonal structure is

stabilized, which is a slightly distorted form of cubic structure. When concentration of

oxygen vacancies is increased further by doping with 6–10 mol% of Y2O3, cubic

structure becomes stable.

There are also other phenomena, through which the high temperature phases

(particularly tetragonal phase) can be stabilized at room temperature, with other

dopants, for example, tetravalent cations , Ce4+

, etc. [6].

Generally, dopants with oversized cations such as Y2O3 and CeO have higher

solubilities, while those with undersized cations, e.g., MgO and CaO have lower

solubilities [2]. Yttria stabilized zirconia is generally preferred due to the associated

advantages such as easy stabilization of complete tetragonal phase and low sintering

temperature etc. [7].

The monoclinic to tetragonal phase transformation is associated with a 4%

reduction in volume, which produces large stresses and cracking of pure zirconia

making it unreliable for fabricated components. Fully stabilized zirconia does not

undergo any phase transformations and is thus used for a large number of applications

as a hard and chemically inert ceramic as well as in thermal barrier coatings in gas

turbines, non-metallic knifes, etc. Due to its hardness and optical properties, it is also

used in jewelry.

2.2.1.1 Transformation Toughening

A dramatic increase in use of zirconia was, in fact, realized with the discovery of

transformation toughening mechanism [8]. Garvie et al. were first to report about this

mechanism in 1975 in their famous paper entitled ―Ceramic Steel?‖. The title was

purposely selected to emphasize the similarity of certain properties of toughened

zirconia with steels. For example, elastic modulus and thermal expansion coefficient of

both are similar, they exhibit martensitic transformation, and the fracture toughness and

strength of ZrO2 can be improved by controlling tetragonal to monoclinic phase

transformation, etc.

Transformation toughening is based on the formation of meta-stable tetragonal

phase and hindrance in its transformation to monoclinic phase due to volumetric

10

expansion associated with this transformation. When a crack is produced and it

propagates through metastable tetragonal phase, free space created behind crack tip

induces phase transformation that results in a blockage for the crack to grow further

(Fig. 2.2). Consequently, fracture toughness of ceramic material increases due to

transformation.

Fig. 2.2: Schematic description of transformation toughening at microstructural level [9].

It may be assumed that there are three requirements for transformation

toughening to take place: First, presence of a metastable phase and possibility of its

transformation into a more stable phase under applied stress in crack-tip field. Second,

the transformation to be instantaneous and not a sluggish time dependent process

requiring long-range diffusion. Third, an associated shape and / or volume change as a

result of phase transformation [10]. Some of the applications of zirconia and its

composites are given below.

2.3 Applications of Zirconia Ceramics

2.3.1 Zirconia as a Dental Material

Since very beginning, metallic materials (such as titanium, gold etc.) are being

used as dental materials [11]. The work on use of ceramics in dental applications has

been expedited since 1970s. As a result, the use of ceramics for oral implants can be

traced to 30-40 years back [12-13]. Recently their ceramic counterparts are extensively

substituting metals due to following reasons.

11

Esthetics: Unlike color of metals (e.g., gray for titanium), the ceramic materials (like

alumina, zirconia) are white in color, thus mimicking natural teeth [14].

Health Hazards of Metals: The release of titanium particles and its corrosion products

may cause health hazards. A high concentration of titanium has generally been observed

in the vicinity of titanium oral implants [15]. Other hazards include sensitization of

patients with titania [16] or titania allergy [17] etc. With biocompatible ceramics, such

hazards may be avoided.

Hygienic: The dental implants and crowns made of ceramic materials retain less plaque

and calculus than their metallic counter parts and promote gum health.

Mechanical properties: Zirconia implant have high flexural strength i.e. 1200 MPa

[18].

Biocompatible: The biocompatibility and osseointegration of zirconium implants is

much higher than their titanium counter parts.

2.3.2 Zirconia Based Refractory Materials

Traditionally the refractory materials should have higher thermal shock resistance

and high temperature mechanical strength. They should be able to sustain significant

amount of elastic energy during thermal shock and show resistance against enhanced

slag attack and erosion by flowing molten metal. The phase-stabilized zirconia

reinforced with alumina is reported to have improved density, excellent strength and

thermal shock resistance. The addition of alumina not only improves its mechanical

properties but also enhances the phase (tetragonal) stability of zirconia. This is the

reason that presently in market several zirconia-alumina refractory bricks are available

[19-20].

The nanocrystalline ceramic material has been proved a potential candidate for

thermal barrier coating. By limiting the thermal conductivity through incorporation of

grain boundary scattering as an extensive phonon-scattering phenomenon thermal

barrier properties can be attained in coated components. This has been proved as an

effective strategy for many materials, but the major concern with the long-term stability

of such coating is the high temperature coarsening. Due to grain boundary pinning effect

offered by the dispersed alumina nanoparticles in stabilized zirconia, the composite is

being investigated as a potential candidate for such thermal barrier coatings.

12

2.3.3 Zirconia in Cutting Tools Industry

Advancement in the ceramic processing techniques has resulted in production of

new generation of high performance ceramic cutting tools with improved properties.

This was possible due to invention of materials with better properties such as higher

wear resistance, thermal shock resistance, hardness and fracture toughness. Presently,

the ceramic tools are being used for machining of various cast iron, ferrous and non-

ferrous alloys and so on. The grains of zirconia-alumina composite with improved

strength due to phase transformation toughening are being used in the cutting tool

industry. Furthermore, the ceramic nozzles made of zirconia alumina composite are also

available in market for spraying the abrasive materials [21-22].

2.4 Composite Materials

Throughout the human history, the need of composite materials has remained the

same, like from pre-historic mud bricks with straw reinforcement for ancient civil

structures to the advanced composite parts for aircraft engine. Generally speaking,

modern composite materials are combinations of two or more components, which may

be for instance, in the form of a stronger and stiffer reinforcement phase (having

particles, fibers, whiskers etc morphologies) arranged or dispersed in a polymeric,

ceramic or metallic matrix. When reinforcement and matrix are put together to form a

composite, properties of resultant product are superior to those of the constituent

components, depending on the proper selection of the reinforcements, their sizes, shapes

and distribution. Due to versatility of benefits associated with them, composites have

found applications in almost all scientific fields including aerospace, automobiles,

chemical sensors, scientific equipment, corrosion resistant tanks and pipes, sports

equipment, earthquake resistant structures and synthetic biomaterials, etc.

2.5 Ceramic Matrix Composites (CMCs)

Due to excellent physicochemical properties (refractory characteristics, chemical and

bio-inertness, excellent mechanical properties etc.), ceramic oxide materials have found

wide spread applications in modern processes and technologies [23]. The use of

ceramics extends from biomedical (artificial dental and bone implants, etc.) to civil

structures, electronic devices (i.e. magnets, semiconductors, superconductors, insulators,

fiber optics), structural refractories, and so on. In spite of various advantages of

13

ceramics, they suffer from an inherent problem of lower fracture toughness under

mechanical or thermo-mechanical load, which is associated with their poor resistance to

crack initiation and propagation. This problem can be circumvented through the

formation of ceramic matrix composites (CMC).

With advances brought about in manufacturing techniques for dental and

orthopedic ceramics, demand for new types of ceramic materials capable of

withstanding the advanced specifications has increased [24]. In such applications,

ceramic materials also beat their metallic counterparts due to their chemical inertness,

biocompatibility and high esthetic quality with good mechanical properties.

2.6 Zirconia-Alumina Composites

Like zirconia, the alumina is also a biocompatible inert oxide. Its good strength,

biological and chemical inertness, fatigue and fracture resistance have made it (specially

α-alumina) an attractive material for many applications, for example in electrical and

medical fields as well as in structural ceramics and computer electronics, etc. [25].

Alumina is also being used in certain clinical applications like alveolar ridge and

maxillofacial reconstruction, ossicular bone substitutes, knee prostheses, bone & dental

screws, segmental bone replacements and corneal replacements.

Due to high mechanical strength of alumina, it is also being used as a reinforcing

material in bio-composites. Recently, development of zirconia-alumina composites has

attracted much attention of the researchers [26-28]. Both zirconia and alumina have

lower resistance to fracture [27, 29] than when they are combined to make composites.

Although yttria-stabilized tetragonal zirconia (TZ-3Y) has comparatively better

mechanical properties and fracture toughness [2, 30-31], these properties may

significantly be enhanced by composite formation. The major problem associated with

tetragonal zirconia is its phase degradation at low / body temperature [32]. The

formation of alumina-zirconia composites has been suggested as a solution to overcome

low temperature degradation problem as well as to further improve mechanical

properties of zirconia [33]. Consequently, zirconia-alumina composites are considered

to withstand dental forces over a large span of time without fracture [34].

Presently, in market different types of zirconia are being offered for dental

applications including onlay (VINCRON), veneers, crowns, inlay etc., they contain

either yttria-doped zirconia or magnesia-doped zirconia or alumina toughened zirconia.

Attempts have also been made to combine two or three toughening mechanisms (e.g.,

14

transformation toughening, alumina toughening, etc.) together to get better properties. In

zirconia composites, grain size has also been found to play an important role. For

example if grain size is less than 0.5 µm then sluggish tetragonal—monoclinic

transformation may occur, but when the grain size becomes smaller than 0.2 µm, the

tetragonal—monoclinic transformation is unlikely to occur [35-36].

For the sake of dental material applications requirement is strong and tough

ceramic materials [37], which can be fulfilled through development of the zirconia

composites with enhanced fracture toughness [38] by combining various toughening

mechanisms.

2.7 Fiber Reinforced Ceramic Composites

While making composites, one of the important considerations is the shape and size

of reinforcements. Particularly, for high load bearing applications, fiber-reinforced

ceramic composites are very important [39]. That is why, in previous few decades, a lot

of work has been devoted to development of glass, SiC or carbon fiber reinforced

ceramic composites [40-45]. The fracture toughness of composite also depends on

mechanical properties of the fibers used. In comparison with polycrystalline fibers,

ceramic whiskers offer higher strength levels along with significantly enhanced

resistance to fracture. For instance, the tensile strength of SiN whiskers is 50GPa, while

that of glass fibers is 3GPa and in case of bulk glass this strength decreases to value of

0.1GPa [46-49]. Consequently, work on the use of strong ceramic whiskers in ceramic

matrices for formation of whisker-reinforced ceramic matrix composites has gained a lot

of attention [50-52].

Whiskers are single crystals with high crystal perfection and thus superior

mechanical properties. The shape of a whisker is generally acicular or needle like with

different dimensions. When used in formation of composites, they enhance their

mechanical properties by mechanisms such as crack bridging, crack deflection etc. [46,

53]. Some studies have shown that fracture toughness and flexural strength of dental

composites can be doubled by using whiskers [54]. The fracture toughness of SiC

whiskers is about 3 MPam-1/2

, and by incorporating these SiC whiskers in fine grained

(2µm) alumina matrix, fracture resistance of alumina increases, e.g., up to three times at

a concentration of 20 vol% whiskers [55-56].

On the other hand, for composites containing dispersed zirconia as

reinforcement, advantages of transformation toughening and/or microcracking [57-58]

15

may be incorporated. For example, an addition of 20 vol% of ZrO2 particles to mullite -

20 vol% SiC(whisker) composites enhances fracture toughness from 4.7 MPam-1/2

to 7.7

MPam-1/2

[59]. In this case, microcrack toughening also plays its role along with

whisker toughening. Microcrack toughening is based on the absorption of propagation

energy of a main crack by microcracks through branching of main crack [60].

The studies show that whiskers generally offer five types of toughening

mechanism i.e. crack bowing [61], microcracking [62], crack deflection [63-64],

whisker pullout [65] and crack bridging [66]. In crack bridging, the crack propagation is

stopped due to friction interlocking between opposing fracture surfaces offered by the

ligamentary bridges between these surfaces. This kind of phenomena also occurs in

various type of composite materials like rubber-polymers [67], ceramic-metal

composites [68] and core-grained-alumina [69].

2.8 Hydrothermal Synthesis

2.8.1 Brief History

Primarily hydrothermal synthesis was started by geologists to simulate in

laboratory the natural hydrothermal phenomena occurring in earth crust. The history of

hydrothermal synthesis can be traced back to the mid of 19th

century [70] when

Schafthaul prepared quartz fine powders of sub-microscopic and nano-size by using a

digester containing freshly precipitated silicic acid. This synthesis route has gained

much attention with development of pressure vessel engineering. From 1840s to 1900s,

most of the experiments, in which nano-sized products were obtained, were considered

as failure due to lack of sophisticated electron microscopic techniques to analyze them.

If particle size grew into millimeter size, it was also considered failure, although they

could be single crystals [71]. Most of the chemical analysis techniques available were

suitable for product identification only. A systematic understanding of nanosized

products began with invention of X-Ray diffraction. Even then the hydrothermal

synthesis of bulk crystals did not flourish too much because in majority of experiments

the product was very fine grained and available X-ray data could not be used to analyze

them [72]. The other problem associated with these experiments was long experiment

time. Morey quotes that earlier hydrothermal experiment was a horrible experience

because the experiment lasted even for several months and they ended up with very fine

grained products which were difficult to analyze and experiments were considered as

16

failure [73]. Up to 1950s great variety of oxides, carbonates, germinates, silicates,

phosphates, sulphates, etc., were prepared and they were analyzed as fine crystalline

materials and attention was only focused on phase equilibrium studies. However, no

emphasis was on growth of single crystals, because of the lack of knowledge on

hydrothermal solution chemistry. Attempts to understand hydrothermal solution

chemistry started from 1950. In 1980s, with the availability of high-resolution scanning

electron microscopes (SEM’s), researchers started observing fine structures obtained

through hydrothermal synthesis, which were previously discarded regarding as failure.

The work on hydrothermal synthesis of fine and ultrafine structures with controlled

sizes and morphology began in 1990s. In last two decades with introduction of a new

term ―nanotechnology‖, hydrothermally prepared ceramic materials with submicron to

nano-sized structures have revolutionized the world of science.

Today with advancement in knowledge of thermodynamic principles and

chemical energy, understanding of solution chemistry has increased. That is why these

days the hydrothermal reactions are performed at relatively lower pressure and

temperature conditions with comparatively faster processing times. Hydrothermal

synthesis is also considered as green technology with no or little waste, less energy

consumption, no hazardous processing materials etc. and has gained status of science of

sustainability [74].

2.8.2 What is Hydrothermal Synthesis

In hydrothermal synthesis, a single phase or multiple phases are produced at high

temperature and pressure conditions from aqueous solutions. Recently with in-depth,

study of thermodynamics and the kinetic, involved in this synthesis technique, reactions

are also performed at ambient conditions. However, in case of growth of single crystals,

additional pressure is required to control solubility and growth reactions [75-76]. The

reactants (called precursors) are generally administered in form of gel, suspension,

solution etc. Organic or inorganic additives are used to serve different purposes i.e. to

control pH, to promote particle dispersion, to control crystal morphology. The upper

pressure and temperature limits of hydrothermal synthesis are up to 500MPa and

1000°C respectively [77]. But for large scale material production the practical limit of

pressure and temperature are kept at 100MPa and 350°C, below this limit the reaction

conditions are considered as mild while above this the conditions are considered severe.

17

These conditions (i.e. mild & severe) are considered from viewpoint of corrosion and

strength of construction material of processing vessel/autoclave. Previously processing

time was up to several months but presently with better understanding to hydrothermal

chemistry, this reaction time has been optimized to days and hours [78].

The exact mechanisms for crystal formation under hydrothermal synthesis is not

very clear [79-80]. However, there are two famous proposed processes involved in

hydrothermal synthesis of crystal formation; these two mechanisms are in situ

transformation and dissolution-precipitation [81]. Which mechanism is actually

involved depends upon solubility of precursor in the solution. In case of in-situ

transformation, a heterogeneous nucleation type of process is proposed where ions

nucleate on some un-dissolved oxide and in this way result into formation of a crystal.

For example, Hu et al. [82] have reported hydrothermal formation of BaTiO3

nanostructures in which dissolved barium deposits on the surface of small TiO2 particles

during treatment. The additional barium continues to diffuse to the formed nuclei until

supply of TiO2 and barium is exhausted. In second mechanism as suggested by

Pinceloup et al. [83], TiO2 particles dissolve in solution and forms some complex like

hydroxyl-titanium. The subsequent reaction is between this complex and barium ions in

solution to form the BaTiO3 nanostructures. In case of hydrothermal synthesis of

aluminum ammonium carbonate hydroxide (AACH) whiskers, a modified form of

dissolution precipitation mechanism may be involved. At ambient conditions, urea and

aluminum nitrate (the precursors used for AACH whisker formation) are easily soluble

in distilled water. However, during hydrothermal process this urea decomposes and

produces ample amount of ammonia and bicarbonate ions according to following

chemical reactions.

CO (NH2)2 + 3H2O → 2NH4+ + 2OH

–+CO2 (2.1)

CO2 + H2O → HCO3- + OH

- (2.2)

The produced ions (NH4+, OH

– & HCO3

-) along with Al

3+ ions in the solution combine

to form AACH.

NH4+ + Al

3+ + HCO3

- + 3OH

- → NH4Al(OH)2CO3 + H2O (2.3)

According to Li et al. [84], initially, an amorphous phase is formed, which acts

as nucleation sites for whiskers. The whiskers then grow at the expense of this

18

amorphous phase. Ultimately, Ostwald ripening may also take place in which smaller

whiskers are dissolved allowing bigger ones to grow. The pH of solution also changes

due to release of decomposition products (especially OH–) of urea. This pH change

plays vital role in defining morphology of AACH product. Ma et al. [85] and the

authors have reported the effect of change of pH of the solution during hydrothermal

conditions on morphology of the formed structures.

2.8.3 The Merits of Hydrothermal Synthesis

Hydrothermal synthesis has many advantages over conventional and non-

conventional methods. Nearly all kinds of ceramics powders, fibers, whiskers, coatings

etc. can be synthesized with the help of hydrothermal synthesis. Due to low processing

temperature and less waste produced, this synthesis technique is environment friendly.

Ability to precipitate powders directly from solutions, uniformity of nucleation and

growth, better control of morphology, size and aggregation make this technique very

useful [86-87]. A special advantage of hydrothermal process, based on high pressure

and temperature, is the synthesis of some of the phases that almost impossible to

produce by other techniques. For example large quartz single crystals can only be

synthesized through hydrothermal technique [75].

Unlike other synthesis techniques that require elevated temperatures,

hydrothermal synthesis is performed at lower or moderate temperatures, which results in

very small defects in final product. For example tungstates of Sr, Ca and Ba synthesized

hydrothermally do not contain schottky defects [88], resulting in improved luminescent

properties. Similarly, other defects like barium ion vacancies in barium titanate due to

substitution of oxygen ions with hydroxyl are not formed. This kind of defect degrades

the dielectric properties [89-90].

2.8.4 Synthesis of Alumina Whiskers

The work on development of one-dimensional (1D) nanostructures like nanorods

[91-93] etc. was started with the discovery of carbon nanotubes [94]. The examples of

1D structures made of metal oxides include Ga2O3 [95], ZnO [96], SiO2 [97] etc. These

are synthesized through vapor phase transport [98], laser ablation [99] or arc discharge

[100]. Due to remarkable properties associated with these structures, they are being

widely investigated.

19

Recently, 1D single crystal whiskers have attracted much attention due to their

wide range of potential applications in industry. During the last few decades various

kinds of whiskers like AlN [101], MgO [102], Al18B4O33 [103], etc., have been reported.

Due to high elastic modulus, melting point, chemical and thermal stability of alumina

whiskers, they occupy a unique position [104]. Various techniques have been employed

for the synthesis of alumina whiskers. Webb et al. [105] first prepared alumina whiskers

by heating alumina at 1300-1450°C. Zhang [106] used Al & TiO2 powders for synthesis

of alumina whiskers. Similarly, Al and SiO2 powders were used to prepare alumina

whiskers via vapor liquid and solid deposition in Ar at a temperature of 1300-1500°C

[107-109]. Other researchers have used different metal oxides in Ar atmosphere at 850-

1100°C for synthesis of alumina whiskers [110-111]. Okada et al. [112] prepared

alumina whiskers by using aluminum bromate whiskers, while Yu and Du [113]

prepared them from aluminum ethoxide used as precursor. However, these techniques

require severe reaction conditions (such as high processing temperature) and often the

precursor materials are quite uncommon and expensive. On the other hand,

hydrothermal synthesis of alumina whiskers requires readily available raw materials and

is performed at lower temperatures. It also helps in preparation of whiskers with highly

reduced defects in the structure.

2.9 Deflocculant and Dispersion of Reinforcement in

Composite

It must be noted that most of the physical properties of a composite depend on state

of dispersion of reinforcement inside matrix. An improved dispersion of reinforcement

can be achieved through mechanical mixing and / or by addition of suitable dispersing

agents. The chances of agglomeration of reinforcement are dependent on its size and

shape. Smaller sizes and shapes with high aspect ratio are more vulnerable to

agglomeration. Mechanical mixing is not so effective in such cases and even it may

cause unnecessary morphological distortion of the reinforcement and incorporation of

impurities. Consequently, one has to rely on dispersing agents or deflocculants [114]. A

dispersing agent should be such that it may be removed by evaporation during

calcinations or sintering steps. Uniform dispersion ensures not only high green and

sintered densities, but also improved mechanical properties.

20

The selection of a suitable deflocculant and determination of its optimized

amount is necessary to obtain maximum advantage in terms of improved mechanical

properties. On the basis of already published data on the dispersion of alumina in

different matrices for the formation of composite, the cetyl trimethylammonium

bromide (CTAB) [115] and p-aminobenzoic acid (PABA) [114, 116] are reported to

give better results. Vishista et al. [114] investigated a series of deflocculating agents for

alumina dispersion and proposed PABA as a best choice to obtain improved density and

mechanical properties. However, they did not investigate the effect of CTAB. Although

Francisco et al. have used CTAB for dispersion of zirconia in alumina matrix [116], use

of CTAB for dispersion of alumina in zirconia matrix has not been investigated yet.

Therefore, it is important to study deflocculating effects of CTAB and PABA on

dispersion of the alumina whiskers in TZ-3Y matrix.

2.10 Hydroxyapatite for Improving Biocompatibility

2.10.1 Hydroxyapatite and its Preparation

Hydroxyapatite (HAp) is a synthetic ceramic material and is crystalline compound

of calcium and phosphorous, which is similar to mineral present in the bone. The natural

bone consists of organic (proteins and collagen) and inorganic (carbonated

hydroxyapatite) parts [117]. Approximately more than 50% of bone or teeth is made up

of modified form of hydroxyapatite (i.e. carbonated hydroxyapatite with lower

crystallinity) [118]. It is believed that synthetic hydroxyapatite, which may be used for

artificial bone and dental implants, is entirely compatible with the body [119]. When

exposed to body fluid this synthetic hydroxyapatite forms bonds with the body tissues.

Due to its composition and tissue bonding property, it is an important member of

bioactive ceramics.

HAp is a member of calcium phosphate family. Other members of this family in

the order of decreasing solubility in the body fluid are tetracalcium phosphate

(Ca4P2O9), amorphous calcium phosphate, tri-calcium phosphate (Ca3(PO4)2), and

hydroxyapatite (Ca10(PO4)6(OH)2) [120]. It is known that HAp does not break under

physiological conditions [121]. It is thermodynamically stable in pH of the body fluids

and actively takes part in bone bonding. HAp is thermally unstable and may decompose

at temperatures between 800-1200°C depending upon its stoichiometry [122]. If thermal

processing (like sintering etc) is required, control of stoichiometry becomes very

21

important. If molar ratio of calcium to phosphorus (Ca/P) deviates from 1.67 in HAp, it

can decompose into tricalcium phosphate and tetracalcium phosphate. With lowering of

Ca/P ratio and increase in acidity of the environment, chances of HAp to dissolve are

enhanced. It is observed that if Ca/P is less than unity then acidity and solubility both

are very high. These effects are minimized when ratio comes near 1.67 (the

stoichiometric Ca/P ratio of HAp) [123]. By using the HAp with stoichiometric Ca/P

ratio it can be sintered at temperature up to 1400ºC [124].

For bone mimetic implants, HAp is synthetically produced through various

routes [125-129] with varying morphologies and stoichiometry. The synthesis routes

can be divided into two main categories, i.e. solid state methods and wet methods [130].

The wet methods contain chemical precipitation, sol–gel synthesis, co-

precipitation, hydrothermal synthesis, etc. [131]. A well-crystallized and

stoichiometrically suitable product can be prepared through hydrothermal and

precipitation methods. However, hydrothermal method needs relatively high

temperature and long heat treatment time. The nanometer sized HAp with good

stoichiometry and various morphologies can be produced through precipitation method

[132]. A simple and most commonly employed precipitation method for HAp

production is by using ammonium phosphate and calcium nitrate in basic condition in

presence of ammonium hydroxide [133] according to the following chemical reaction.

10Ca(NO3)2·4H2O + 6(NH4)2HPO4 + 8NH4OH → Ca10(PO4)6(OH)2 + 20NH4NO3 + 46H2O

2.10.2 Hydroxyapatite for Bio-Composites

Due to its chemical and crystallographic resemblance with natural calcified tissues

of vertebrates, HAp has attracted much attention as a biomaterial with high bio-affinity

[134-136]. It can make bond directly with bones that is required for osseointegration of

the implants when used as dental, orthopedic and maxillofacial applications. It is very

useful material for the formation of bio-ceramic coatings and as a bone filler material.

When large sections of bones are removed or bone augmentation is required (e.g., in

dental or maxillofacial applications), HAp powders, porous beads or blocks are used to

fill bone defects or voids. This HAp based filler provides scaffold and encourages the

speedy natural bone formation. In spite of its excellent biocompatibility, poor

mechanical properties are the main limitation in its use in load-bearing parts. Therefore,

to be used for load bearing biological applications, its mechanical properties (i.e.

22

strength, hardness etc.) should be improved [137-140]. For obtaining this goal, there are

two possible approaches; first is to make a HAp coating on metallic, ceramic or

polymeric implant (like Ti/TiO2 alloy or steel, ZrO2 etc.) for biocompatible surface

properties. In this way, body is willing to accept the implant. The main limitation of this

approach is problem of de-bonding from metallic substrate [141]. The second approach

may be through the formation of composite with some bio-inert material. This can

further be divided into two cases: one HAp added composite, second HAp based

composite. In the latter case, HAp is reinforced with morphologies like platelets, fibers

or powder. With the help of these reinforcements mechanical properties of HAp based

composites can be enhanced up to 3 times [139]. In a study on screw-shaped dental

implants, it has been reported that HAp based composites are better than pure Ti [142].

However, for load bearing applications further improvement in mechanical properties is

required. This can be achieved in composites in which biocompatibility of various

strong materials is increased by incorporating HAp. As HAp is a ceramic material, the

use of some ceramic material with comparatively better mechanical properties may be a

better choice. The zirconia-alumina system, a bio-inert composite [134], is well-known

for its higher strength values [143]. Therefore, proving to be one of the most suitable

materials for HAp added high strength composites.

2.11 Solid-State Sintering

The basic purpose of sintering is to develop interparticle bonding and increased

density. During sintering process, mass transport takes place between neighboring

particles inside the compact, initially resulting in the formation of a neck. The neck

grows as sintering proceeds and particles coalesce, small pores collapse forming the

larger pores. Then these pores also tend to diminish. The dimensions of the sample are

reduced due to an overall decrease in porosity [144-145].

There are several proposed solid-state sintering mechanisms, which explain

densification process and movement of atoms or vacancies to boundary and formation

of neck between neighboring powder particles. These mechanisms can be grouped into

two categories: bulk and surface transport mechanisms. The surface transport includes

evaporation / condensation and surface diffusion. Due to surface transport, centre of

mass of the particles does not move and coarsening takes place instead of densification.

But bulk transport includes volume diffusion, grain boundary diffusion and plastic flow

[146] which results in more densification.

23

One can identify three stages of sintering which may be called as initial,

intermediate and final stages. During initial stage, adjacent particles that are already in

contact form necks as shown in Fig. 2.3. Additional necks are formed as the particles

rotate and move closer together.

Fig. 2.3: Micrograph showing neck formation during initial stage of the sintering [147].

During intermediate stage, the necks continue to grow through surface and bulk

transport phenomena (possibly including evaporation and condensation, plastic flow,

viscous flow, volume diffusion, grain boundary diffusion etc.). The necks begin to come

in contact with each other and pores start to collapse, but still there are definite grain

boundaries between original powder particles. During this stage strength improves

through mass transport and Ostwald ripening phenomena [148]. In Ostwald ripening

small pores combine to form large pores and these pores further, reduce through mass

transport phenomena. During final stage of sintering, grain growth takes place and grain

boundaries start to disappear.

2.11.1 Sintering of Whisker Reinforced Composites

A very small work is reported on the effect of sintering temperature on properties

of whisker-reinforced composites. In a study on SiC (whiskers) reinforced Al2O3

composite, Lio et al. [149] found some dramatic decrease in fracture toughness with

increase in sintering temperature. They attributed this decrease in fracture toughness to

high temperature agglomeration of whiskers. They suggested that agglomeration may

24

change the stress field around Al2O3/SiC(w) interface, which affected toughening

mechanism. However, they did not present a solid justification for the decrease in

mechanical properties with increase in the sintering temperature. Feng et al. [150] have

reported an increase in whisker diameter with increasing sintering temperature of

TiB(whiskers)+Ti metal matrix composite. The increase in aluminum borate whisker’s

diameter from 0.2 to 2 μm by increasing sintering temperature from 1150 to 1350 °C

was also reported by Peng et al. [151]. Since changing the physical size or shape of

whisker has a significant effect on mechanical properties of composite [152], dramatic

change in mechanical properties of whisker-reinforced composites may be attributed to

the change in morphology or diameter of the whiskers.

2.12 Concluding Remarks and Objective of the Thesis

As discussed, the use of whiskers significantly improves the stress-bearing

capabilities of dental composites. Yttria stabilized zirconia composites offer

multifunctional properties (both mechanical and biocompatible) [153] and are thus

considered excellent materials for dental applications. During the last few decades,

many studies have been reported on use of SiC whiskers to reinforce zirconia for the

formation of composites [154-156]. But non-oxide reinforcements like SiC may result

into high temperature degradation or reaction between reinforcement and matrix [157].

On the other hand, oxide materials such as alumina have good compatibility with

zirconia matrix. Alumina is also a biocompatible material and its whiskers can be made

through a less expensive and environment friendly hydrothermal process. Therefore,

alumina whiskers reinforced yttria stabilized tetragonal zirconia composites may be an

excellent development to solve material-related problems of dental industry. Very

limited work has so far been reported on zirconia composites reinforced with alumina

whiskers. The only worth mentioning report appears to be that of Nevarez-Rascon et al.

[158]. However, they used whiskers of relatively smaller diameter i.e. 2-4 nm, which

were severely agglomerated and properties of composite could not be improved

significantly.

The prime objective of the present study is to develop high strength and

toughened zirconia based composites for various load bearing structural and biomaterial

applications. For this purpose, out of different polymorphs of zirconia, the tetragonal

zirconia was selected due to possessing transformation toughening characteristics. To

25

improve the mechanical properties of zirconia based composite, the alumina whiskers

were used as reinforcement material. To get uniform dispersion of alumina whiskers in

the matrix of tetragonal zirconia, low temperature volatile deflocculant (CTAB) was

added. Sintering temperature of the synthesized composite was also optimized. At the

end, the best zirconia-alumina composition bearing excellent mechanical properties was

further investigated to enhance the biocompatibility by adding appropriate concentration

of HAp.

26

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38

Chapter 3

3 Characterization Techniques

Present chapter focuses on various techniques used for the characterization of the

developed composites. Most of the experimental procedures adopted for synthesis of

composites under various conditions will be addressed in connection with the

experimental results in the respective chapters, although a brief outline of these

procedures will be given in section 3.1.

3.1 A Brief Outline of Synthesis Processes

1. Ammonium aluminum carbonate hydroxide (AACH) nanostructures were

synthesized by hydrothermal technique (chapter 4).

2. Composites with varying alumina whisker concentrations were prepared to study

the effect of whisker concentration on their mechanical properties (chapter 5).

3. Composites were synthesized with varying concentration and type of

deflocculants, in order to assess their effects on the properties of alumina whisker-

reinforced composites (chapter 6).

4. After optimizing the concentration of whiskers and the amount of deflocculants on

the properties of resultant composite, effect of sintering temperature on whisker

morphology and properties of composite was studied (chapter 7).

5. Low temperature tetragonal phase stability of TZ-3Y (3mol% Y2O3 stabilized

ZrO2) in the synthesized composite was accessed through hydrothermal technique

(Chapter 8).

6. Finally, composites having various amounts of hydroxyapatite (HAp) were

synthesized to study the effect of HAp on biocompatibility of these composites

(chapter 9).

3.2 Characterization Techniques

Techniques employed for characterization of composites prepared under different

conditions included X-ray diffraction (XRD), scanning electron microscopy (SEM),

Fourier transform infra-red (FTIR) spectroscopy, thermal gravimetric analysis (TGA),

39

hardness measurement, flexural strength measurement and biocompatibility tests.

Details of the characterization techniques are given below:

3.2.1 X-Ray Diffraction

X-ray diffraction (XRD) is a widely used non-destructive analytical technique in

which a monochromated X-ray beam is focused on a sample. X-rays are scattered by

electrons around atoms of the material, and may combine either constructively or

destructively depending upon arrangement of the atoms in the crystalline material [1].

As a result, high intensity diffracted beams are observed in specific directions, which

are fingerprints of the arrangement of atoms in that crystalline material. Bragg’s law

[2] provides a simple relationship between diffraction angle and interplanar distance of

atomic planes in the crystal lattice.

2 d sinθ = n λ (3.1)

Where

λ = wavelength of X-rays used,

d = interplanar distance,

n = order of diffraction, and

θ = diffraction angle.

Various kinds of useful information can be extracted from XRD data, i.e., type

of crystal structure, number of phases, degree of crystallinity, lattice parameters,

crystallite size, texture and even quantitative analysis [3].

In present work, X-ray diffraction was employed to investigate phases present

in the AACH, alumina, hydroxyapatite and synthesized composite samples. Similarly,

sintering and calcination-induced solid-state phase transformations in bulk and

whisker samples were also investigated.

The X-ray diffraction study was performed on Bruker D8 Discover XRD

machine, equipped with Cu Kα radiation (λ= 0.15418 nm) source at 0.020° intervals in

10° ≤ 2θ ≤ 90° range mostly at a rate of 2 degree per minute. Diffraction patterns were

assigned by comparing them with standard d-values from JCPDS (Joint Committee on

Powder Diffraction Standards) [4] files of different constituting phases.

40

3.2.2 Scanning Electron Microscopy

Scanning electron microscopy (SEM) is a very useful technique for imaging

surface morphology of a sample. A sharp beam of electrons scans the surface of

sample in a raster scan mode to produce an enlarged image of surface features.

Various kinds of signals can be obtained because of interaction of this electron beam

with sample such as secondary electrons, backscattered electrons and characteristic X-

rays, etc. Secondary electrons are generally used for imaging of topography, while

backscattered electrons allow imaging with information about kind of atoms

distributed on the surface. Characteristic X-rays are used for chemical analysis such as

energy dispersive X-ray spectroscopy (EDS). SEM images provide useful information

about the surface feature sizes and their morphology.

In SEM, non-metallic samples are difficult to observe due to problem of

charging of the surface by electrons [5-7]. On the other hand, metallic or electronic

conductor materials can be easily examined. Therefore, in case of non-metallic

samples, it is common practice to coat a very thin layer of gold or other conducting

material on the surface to allow electronic flow and avoid charging.

For topographic images, the JEOL-JSM5910 scanning electron microscope

was used at a voltage of 20 kV. As samples were oxides, i.e., non-conductors, they

were sputter-coated with high purity (99.999%) gold under vacuum using SPI-Module

sputter coater in order to prevent charging of the surface on exposure to electron beam.

Gold-coated sample was placed on a copper stud, which was electrically connected

with SEM stage through carbon tape.

Morphology of reinforcements and, porosity, grain size and phase distribution

of the sintered pellets were examined by SEM, mostly at a voltage of 20 kV. For the

biocompatibility study of the samples, adhesion of grown cultured cells on sample

surface was also confirmed through SEM at lower accelerating voltage. For EDS

analysis again the lower accelerating voltage (7~10kV) was used to reduce the

interaction volume of electrons. It is reported that depth resolution of EDS in SEM can

be increased with decreasing the accelerating voltage because the volume for

generation of characteristic X-rays become smaller at lower accelerating voltage [8].

The selection of lower accelerating voltage for EDS was also inspired from already

reported study in which high resolution EDS analysis was performed on zirconium

based composite at 7 kV [9].

41

3.2.3 Fourier Transform Infrared Spectroscopy (FTIR)

In FTIR, electromagnetic radiations in the infrared range are allowed to pass

through a sample; radiations with frequencies characteristic of various chemical bonds

in molecules in the sample are absorbed. Different combinations of radiation

frequencies are passed and fraction of total radiations absorbed by sample is recorded.

After a number of cycles with different frequency combinations have been recorded

through interferometer, data is analyzed to make an interferogram and Fourier

transformed to give absorbance or transmittance spectrum plotted against

wavenumbers [10]. Energy absorbed in infrared range of light, changes vibrational

state of the bond from ground to excited level. Energy absorbed for this transition is

unique for each bond, which can be used as fingerprint of specie and chemical bond.

The fingerprint of each molecule corresponds to certain wavenumber. The location of

IR bands by variety of bonds depends upon its chemical environment. Therefore, the

shift in wave numbers may occur due to interaction of neighboring atoms (dipole-

dipole interaction).

Consequently, FTIR is a very powerful tool for studying the chemical bonds

(functional groups) between elements in a compound [11]. In the present study, FTIR

was used for determination of various phases present at different stages of

decomposition of AACH whiskers as well as to study the effects of CTAB and PABA

in promoting dispersion of AACH whiskers in composites. For this purpose,

NICOLET 6700 FTIR system of Thermo Electron Co. USA, was used at room

temperature in the range of 400-4000 cm-1

(wavenumber).

3.2.4 Thermogravimetric Analysis (TGA)

Decomposition behavior of AACH whiskers and conditions for removal of

deflocculating agents by thermal treatment were studied through Thermogravimetric

analyzer. Thermogravimetric analyzer measures weight loss as the temperature of

sample is increased at a specified rate and gives information about thermal

decomposition behavior and thermal stability etc.

In the present study, different types of deflocculants were used for enhancing

the dispersion of reinforcement in composite. To prevent the samples from

disintegration during high temperature sintering the removal of these deflocculating

42

agents was necessary. TGA was used to study the thermal conditions for removal of

these deflocculants.

Mettler Toledo, TGA/SDTA851 thermogravimetric analyzer was used.

Measurement temperature was up to 1000ºC with heating rate of 20ºC/min in air. A

typical analysis procedure consisted of measuring 10mg of sample in an alumina

crucible. Crucible was then placed inside the analyzer and temperature was controlled

through a computer program from 40 to 1000°C at 20ºC/min. During heating the

weight loss of sample was measured through installed ultra-microgram resolution

high-precision balance over whole measurement range.

3.2.5 Mechanical Properties

3.2.5.1 Density Measurement

Density (i.e., mass per unit volume) of a component prepared through

compaction and sintering route is very important. It not only gives information about

success of processing parameters [12] but also is useful to explain the mechanical

properties of resultant part.

Density can be calculated straightforward if mass and volume can be

measured. Mass can be measured with high accuracy by using scientific balances

available today. However, for sintered components, it is not always possible to

measure volume exactly if some shape distortion due to shrinkage is involved. In such

cases, volume is generally measured by Archimedes principle. In this method, buoyant

force on the submerged object is used for volume measurement. This buoyant force

experienced by a submerged object is equal to weight of the fluid displaced by the

object and can be determined through the equation given below:

FB = Vp ρf g (3.2)

Where, FB =is buoyancy force, Vp is object volume, ρf is density of the fluid, g is

acceleration due to gravity. Since buoyancy force is weight of fluid displaced, Eq.

(3.2) can be written in terms of mass as given below:

∆m = Vp ρf (3.3)

43

Where ∆m is mass of the fluid displaced and is equal to the weight loss of sample

when submerged in fluid. It can simply be calculated from difference of weight of

sample in air and in the fluid. The volume of the sample (and fluid displaced) is

obtained using the following equation:

∆m / ρf = ma / ρo (3.4)

Where ma is mass of sample in air and ρo is density of the sample.

A schematic of apparatus and its arrangement for measurement of density through

Archimedes principle is shown in Fig.3.1.

Fig. 3.1: Schematic of Density measurement equipment through Archimedes principle.

The apparatus consists of a small perforated steel plate for sample placement

hanged freely with steel wire and placed on a high accuracy weighing balance. This

sample barrier plate is hanged freely in a plastic container containing a fluid (high

purity ethanol). The container is mounted on a support that is not on balance. Weight

of the sample is measured when submerged in fluid. All equipment is placed inside an

44

air tight enclosure so that the air does not disturbs the balance reading. Ultimately,

with the help of sample mass when submerged and in air, density is calculated by

using Eq. (3.4).

3.2.5.2 Hardness Measurement

Hardness measurements were performed on microhardness tester of Wolpert

Measuring Equipment Co. Ltd. (Model 401MVD). Before measurement, surface of

each sample was ground on a grinding wheel up to 1200 grit SiC paper followed by

polishing with diamond paste of 1 μm particle size. Polished samples were washed

with distilled water at room temperature and then rinsed with ethanol followed by

drying in oven at 250ºC for 4 h. For the hardness measurement, three samples for each

composition prepared under similar conditions were selected with approximately 10-

15 indents per sample and subsequently average of these values was calculated. The

hardness was tested according to ASTM standard C1327-99 [13] which is related with

Vickers indentation hardness for advance ceramics. Formed indentation size was

observed with the help of installed optical microscope and its size and hardness values

were measured through the software automatically. In general, a load of 1 kg was

employed for 15s except where stated otherwise.

3.2.5.3 Measurement Of Flexural Strength

Reliable information about tensile strength of a material is basic requirement for

its structural applications. Tensile test is easy to perform with ductile materials.

However, technical problem of aligning and gripping the test specimen is often quoted

as the reason for not performing tensile test for brittle materials [12-13]. For

measuring the strength of brittle materials like ceramics, flexural strength test is used.

Indirect tensile strength can be obtained in the form of flexural strength [14-15]. There

are different ways of measuring flexural strength like three point and four point

bending test for uniaxial geometry, ring on ring test for bi-axial geometry etc.

However, uniaxial strength test (i.e., four or three point bending test) is complicated

and requires expertise. Moreover stress state in most applications (specially for fiber-

reinforced ceramics) is not uniaxial but rather bi-axial [16]. Therefore, in present study

bi-axial flexural strength testing (usually meant for fiber reinforced composite) was

performed by using ring on ring test (at cross-head speed of 0.1 mm/min) as per

ASTM F394–78 standard [17]. For this purpose, universal testing machine made of

45

SANS, China was used. The test set up [17-18] was designed and got fabricated as

shown in Fig. 3.2.

For flexural strength measurement of the composite samples their surface

flaws were removed through polishing [19]. For this purpose, they were ground on

successively finer silicon carbide abrasive paper down to 1200 grit and polished with

1 micrometer diamond paste. These samples were then dried at 250°C for 4 h for

removal of any residual water vapors.

Fig. 3.2: Ring on Ring test arrangement, for determining Bi-Axial flexural strength.

In this arrangement thin disk shaped specimen was placed between two

metallic rings (lower ring being large and hollow at center & used for sample support

while upper ring being small and solid & used for loading the specimen).

Flexural strength was determined by using the equation given below:

(3.5)

46

S=biaxial flexural strength (MPa); P= load at fracture (N); d=specimen thickness at

fracture origin (mm); where the value of X & Y can be determined from equations

given below:

(3.6)

(3.7)

where ν = Poisson’s coefficient (ceramic = 0.25 [20], ISO 6872); A = radius of support

circle (mm) = 3.5mm in our case; B=radius of loaded area (mm) = 0.7mm in our case;

C=radius of specimen disk (mm).

3.2.6 In-Vitro Biocompatibility Study

3.2.6.1 Cleaning And Sterilization Of Samples

Before cell adhesion biocompatibility study, cleaning of the sample surface is an

important step because surface of the samples may have been contaminated by human

contact. For this purpose, samples were soaked in ethanol-acetone mixture in 1:1 v/v

ratio for one hour. (Acetone and ethanol are well-known cleaning agents for removing

grease and dust particles.) After that, the samples were washed with distilled water in a

flask. These washed samples were then dried at 80oC in a closed flask overnight. For

removing the any microorganisms, the dried samples were sterilized in an autoclave at

120°C, 15 psi for 40 minutes, followed by exposure to UV light. Throughput, the

samples were handled with sterilized tweezers.

3.2.6.2 Cell Culturing

In-vitro study was performed using RD (Rhabdomyosarcoma) cell line [21]

under scientifically controlled conditions. Cell culturing medium contained all the

necessary ingredients required for keeping the cells contamination free, for their

survival and growth (like serum, antibiotic cocktail etc). Samples of different

compositions were placed in a 96-well cell culturing plate. Equal amount of fluid

containing cell culturing media and RD cells was poured on each sample with the help

of sterilized pipettes. Samples were then placed inside the oven at 37.48ºC (human

body temperature) for incubation. After 72 h of incubation, the culturing media was

47

removed with the help of sterilized pipette and the samples were dried at the same

temperature for further 36 h. Afterwards, response of cells to the sample surface was

observed under the SEM at very low accelerating voltage of less than 3kV [22]. Prior

to SEM, the samples were gold coated through sputtering for avoiding the charging

effects. Subsequently to confirm that layer formed on the sample surface was that of

cells, the samples in which continuous well adhered surface was observed, were left in

the oven for further 48 h and were again studied through SEM.

48

References

[1] Smith BC. Fundamentals of Fourier transform infrared spectroscopy: CRC Press;

1996.

[2] Callister WD. Materials science and engineering: an introduction: John Wiley &

Sons; 2007.

[3] Cullity BD. Elements of x-ray diffraction: Addison-Wesley Pub. Co.; 1978.

[4] JCPDS–ICDD: PCPDFWIN, Version 2.0, August 1998.

[5] Markovic V, Marsh H. Microscopic techniques to examine structure in anisotropic

cokes. Journal of Microscopy. 1983;132(3):345-52.

[6] Marsh H, Crawford D. Structure in Graphitizable Carbon from Coal-Tar Pitch

HTT 750-1148-K - Studied Using High-Resolution Electron-Microscopy. Carbon.

1984;22(4-5):413-22.

[7] Bourrat X, Oberlin A, Escalier JC. Microtexture and Structure of Semi-Cokes and

Cokes. Fuel. 1986;65(11):1490-500.

[8] Sakurada T, Hashimoto S, Tsuchiya Y, Tachibana S, Suzuki M, Shimizu K.

Lateral Resolution of EDX Analysis with Ultra Low Acceleration Voltage SEM. J

Surf Anal. 2005;12(2):118-21.

[9] Friel JJ, Greenhut VA. Novel Technology for X-ray Mapping of Ceramic

Microstructures. J Am Ceram Soc. 1997;80(12):3205-8.

[10] Stuart BH. Infrared spectroscopy: fundamentals and applications: J. Wiley; 2004.

[11] Griffiths PR, Haseth JAD. Fourier transform infrared spectrometry: Wiley-

Interscience; 2007.

[12] Ring TA. Fundamentals of ceramic powder processing and synthesis: Academic

Press; 1996.

[13] ASTM.org. Standard test method for vickers indentation hardness of advanced

ceramics Philadelphia, PA,: West Conshohocken, PA; 1991.

[14] PrÖBster L, Geis-Gerstorfer J, Kirchner E, Kanjantra P. In vitro evaluation of a

glass–ceramic restorative material. J Oral Rehabil. 1997;24(9):636-45.

[15] Perez-Gonzalez A, Iserte-Vilar JL, Gonzalez-Lluch C. Interpreting finite element

results for brittle materials in endodontic restorations. Biomed Eng Online.

2011;10:44.

[16] Donald FA. Equibiaxial flexural strength testing of fiber-reinforced ceramics

2008.

49

[17] ASTM.org. Standard test method for biaxial flexure strength modulus of rupture

of ceramic substrates. Philadelphia, PA,1991.

[18] Chandrasekhar Kothapalli, Montgomery T. Shaw, Wei M. Strength distributions

of sintered hydroxyapatite nanoparticles. In: Proceedings of 11Ith

International

Conference on Fracture (ICF-XI). Conference, Conference 2005. p. 4811-7.

[19] Giordano R, Cima M, Pober R. Effect of surface finish on the flexural strength of

feldspathic and aluminous dental ceramics. Int J Prosthodont. 1995;8(4):311-9.

[20] Cattell MJ, Chadwick TC, Knowles JC, Clarke RL, Lynch E. Flexural strength

optimisation of a leucite reinforced glass ceramic. Dent Mater J. 2001;17(1):21-33.

[21] Elizabeth S. In vitro cytotoxicity of the epothilone analog, BMS 247550, in

pediatric malignacies. Centaurus. 2002;10, 31-37.

[22] Subramanian A, et al. Fabrication of uniaxially aligned 3D electrospun scaffolds

for neural regeneration. Biomed Mater. 2011;6(2):025004.

50

Chapter 4

4 Hydrothermal Synthesis of AACH Whiskers

Metal oxides with controlled nanostructured morphologies are of considerable

importance due to their unique properties enabling their fine-tuning for specific

purposes. Consequently, they find wide spread applications in the fields of gas-

sensing [1], photocatalysis [2], magnetism, energy storage [3-4], and high strength

composites , etc. There are various techniques employed for synthesis of these

nanostructures with controlled morphologies including vapor–liquid–solid (VLS)

methods [5], thermal evaporation [6], thermal decomposition [7], chemical vapor

deposition (CVD) [8], template-assisted methods [9] and solution-phase synthesis

[10]. These techniques have their own merits and demerits, but solution phase

synthesis offers special advantage of simplicity, cost-effectiveness and less damage to

environment. Various morphologies (nanorods, wires, sheets, flakes, networks etc.)

may easily be synthesized through control of solution chemistry, solution pH,

temperature, pressure, additives, etc. [11]. Among metal oxides, alumina occupies a

unique position due to its use as catalyst and membrane material [12]. Furthermore,

nanostructured aluminum hydroxide is quite useful for its absorption-desorption and

flame retarding characteristics [13-14].

Alumina and aluminum hydroxide have been prepared in different

morphologies including nano-whiskers [15], nano-flakes [16], nano-wires [17], nano-

fibers [18], nano-rods and nano-tubes [19] etc. They may be categorized as one- and

two- dimensional (1D & 2D) nanostructures. It is also possible to prepare three-

dimensional (3D) nano-architectures through solution-phase synthesis route [20].

These 3D nano-architectures offer their own benefits due to three dimensionally

connected nanostructures. Therefore, they are regarded as the next generation

nanostructures [21]. These 3D structures have potential applications in the areas of

electronics and optoelectronics [20]. In this chapter, formation and characterization of

3D urchin like aluminum ammonium carbonate hydroxide (AACH) and aluminum

oxide nanostructures have been presented. To the best of author’s knowledge, any

other group has not reported formation of such structures of aluminum oxides

51

synthesized through solution-phase (hydrothermal) route. These urchin-like structures

were composed of fine 60 nm thorns. A possible mechanism for formation of such

nano-architectures has also been suggested

4.1 Experimental Procedure

4.1.1 Hydrothermal Synthesis of AACH Whiskers

The AACH whiskers were synthesized by hydrothermal technique. For this

purpose a high-pressure reactor system (Limbo 350) of Büchi Glasuster AG,

Swizerland was employed. A bit detailed description of Limbo 350 system is given

below:

4.1.1.1 Reaction Autoclave:

A high-pressure autoclave made of stainless steel and Hastelloy (Fig. 4.1a) and

having PTFE crucible for sample placement was used as a reaction vessel for

hydtothermal synthesis of AACH whiskers. After pouring the precursor solution in

PTFE crucible, the crucible was placed in the autoclave. Autoclave was then covered

by a plate and screwed in position. Cover plate had welded ¼″ Swagelok connections

for various accessories like thermocouple, pressure gauge etc. A metallic O-ring with

silver coating (Fig. 4.1b) placed between cover plate and vessel ensured leak tight

sealing. Solution was heated by heating the outer jacket (Fig. 4.1c) surrounding the

solution container. Temperature of the solution was controlled with the help of a PID

controller. Equilibrium pressure was noted from pressure gauge.

4.1.1.2 Synthesis of AACH Whiskers

Analytical grade aluminum nitrate (Al(NO3)3.9H2O) and urea were used as

precursor materials. AACH whiskers were prepared from solutions containing 25 g of

aluminum nitrate with varying amounts of urea (6, 8, 10, 12 and 15 g) samples

prepared from these solutions were designated as U6, U8, U10, U12 and U15,

respectively. Above mentioned amounts of urea and aluminium nitrate were mixed in

70 ml of distilled water with the help of a magnetic stirrer. After 30 minutes of

stirring, solution was trasnfered into PTFE crucible. Crucible was placed inside

autoclave followed by its sealing. Temperature was set at 120ºC for 24 hrs. After

completing hydrothermal cycle, autoclave was cooled to room temperature. The

52

sample was then collected on a filter paper and washed repeatedly with distilled water

followed by drying in an oven for 24 hrs at 80ºC to 85ºC in air. Samples were

characterized through SEM, TGA and FTIR. The dried AACH whiskers were later

trasnsformed into alumina whiskers during high temperature calination or sintering.

Fig. 4.1(a-c): (a) Laboratory scale high pressure reactor (Limbo 350) by Büchi Glasuster, Switzerland

used for hydrothermal synthesis; (b) Corrosion resistant silver plated metallic seal O-ring; (c) Heating

unit.

(a)

(c) (b)

Cyclone magnetic drive

PID controller

Cover Plate with welded connections

Cyclone rotation-speed

controller

53

4.2 Results and Discussion

Morphologies of nanostructures prepared from solutions containing 6 and 8 g of

urea and designated as U6 & U8, respectively are shown in fig. 4.2 (a-d). SEM images

of U6 (Fig. 4.2 (a & b)) show cluster-like particles with 300-500 nm diameter and

having lengths up to 3 µm. They appear to have formed by joining of some

longitudinal features together. By increasing the amount of urea up to 8g (Fig. 4.2 (c &

d)), these features become well-defined with separation of their surfaces at interfaces.

Fig. 4.2(a-d): SEM images of hydrothermally synthesized AACH urchins (a) U6 at low magnification; (b) U6

at high magnification; (c) U8 at low magnification; (d) U8 at high magnification.

A further increase of urea to 10g (Fig. 4.3 (a & b)) results in formation of

urchin like morphology. The rods and clusters with thorn / spike like features may also

be seen scattered around these urchins. The thorns emanating from urchins are shown

in high magnification images of Fig. 4.3 (c-e). Diameters of individual thorns are

about 60 nm (Fig. 4.3 (f)).

(d) (c)

(a) (b)

54

Fig. 4.3 (a-f): (a) Low magnification image of U10 sample showing a number of urchins-like particles;

(b-e) High magnification images of various micro-urchins showing size and morphology of individual

urchins; (f) Higher magnification image of individual thorn of a micro-urchin.

Increasing urea to 12 g (U12 sample) results in the separation of the spikes of

these urchins like structures into individual AACH whiskers as shown in Fig. 4.4(a).

Various individual whiskers are focused in high magnification images of Fig. 4.4 (b-

d). The diameter of individual whiskers is approximately 120 nm. However, in most of

the cases, they are still agglomerated by joining together with longitudinal directions

parallel to one another. A further increase in urea content results in improving

separation of individual whiskers.

(a)

(f)

(d)

(e)

(c)

(b)

55

Fig. 4.4 (a-d): (a) Low magnification image of whiskers formed in U12 sample; (b-d) High magnification images of

sample U12 showing individual whiskers and their agglomerates

Consequently, urea is found to play a vital role in defining morphology of

produced AACH nanostructures. It is an organic compound that decomposes thermally

into ammonia and carbon dioxide at 70°–100°C [22] .

CO(NH2)2 + 3H2O → 2NH4+ + 2OH

–+CO2 (4.1)

The ammonia then reacts with aluminum ions (produced due to hydrolysis of

aluminum nitrate) in the presence of water vapors and forms AACH

(NH4Al(OH)2CO3) according to following chemical reaction:

CO2 + NH4+ + Al

3+ + 4OH

– → NH4Al(OH)2CO3 + H2O (4.2)

Table 4.1: Change in pH of solution with amount of urea added in the precursor

(b) (a)

(c) (d)

56

Sr. No. Sample

Designation

Urea Amount

(gm)

pH of filtrate after hydrothermal

treatment

1. U6 6 7.4

2. U8 8 7.9

3. U10 10 8.3

4. U12 12 9

5. U15 15 9.5

Increasing the amount of urea results in an increase in the number of OH- ions

in solution and thus a rise in pH of solution [23], as shown in Table 4.1.

This increase in pH also indicates provision of ample amount of reactants (OH-

and NH4+) and according to Von Weimarn's theory, particle size decreases as

concentration of precipitating agent increases [24]. Therefore, by increasing the

concentration of urea, size of nanostructure features is reduced [25], because energy

required for the formation of new surfaces becomes available. This results in

producing a structure in which rod-like structures start separating from each other. At

even higher urea concentrations, a complete separation of whiskers may be obtained.

Ammonia may also play an important role in controlling the morphology of

produced nanostructures. It is known that in AACH crystals Al and O atoms forms

octahedrons [26], as shown in Fig. 4.5. These octahedrons are connected with each

other through strong covalent bonds between oxygen atoms making long chains. The

radicals NH4+ and CO3

2- are connected through weaker bonds among these chains

[15]. When AACH decomposes, these weakly connected radicals are removed in the

form of NH3, CO2 and H2O and leave behind the alumina. One can break up reaction

(4.1) as follows:

CO(NH2)2 + H2O → 2NH3 + CO2 ↑ (4.3)

Al3+

+ 3NH3 + 3H2O → Al(OH)3 +3NH4+ (4.4)

Al(OH)3 + H2O → Al(OH)4− + H

+ (4.5)

57

Fig. 4.5 (a-b): Showing octahedron of the Al-O; (a) Isometric view [27]; (b) Top view.

During hydrothermal synthesis, decomposition of urea produces NH3, which

ultimately forms Al(OH)4. It is known that at lower pH electrostatic charge of

aluminum hydroxide is negative (its isoelectric point is attained at a pH of 9.2 [28]).

Due to electrostatic force, NH4+ can be absorbed on surface of aluminum hydroxide.

When amount of urea and consequently of NH4+ and HCO3

1- is less, chains of Al-O

octahedra are cross linked with each other, which results in formation of micron sized

particles as shown schematically in Fig. 4.6 (a). With increase in urea content, the

amount of produced/absorbed NH4+ and HCO3

1- is increased which causes separation

of chains and thus controls morphology of resultant AACH structures as sketched in

Fig. 4.6 (b to d).

O Al

58

Fig. 4.6 (a-d): Schematic illustration of AACH whiskers and micro-urchin formation in samples containing different amounts of urea; (a) the U6 sample; (b) U8 sample; (c) U10

sample and (d) U12 sample [29].

(b)

(c) (d)

(a)

HCO3-

HCO3-

HCO3-

HCO 3-

HCO3

-

HCO3

-

HCO3

-

HCO3-

HCO3-

HC

O 3-

HCO

3 -

HCO3-

HCO 3-

HCO3-

NH4+

NH4+

NH4+

NH 4+

NH4+

NH4+

NH4+

NH4+

NH4+

NH4+

NH 4

+

NH4+

NH4+

NH4+

Al OH O C NH4+

59

Fig. 4.7 shows thermal gravimetric curve of AACH sample. It shows that

decomposition of AACH to alumina and removal of volatile species like carbon

dioxide, ammonia and water vapors is completed as temperature is increased to about

550ºC. The total weight loss is about 63% which is in good agreement with theoretical

calculations [30].

2 NH4Al(OH)2CO3 → Al2O3+ 2NH3(12.23 %wt loss)+ 2CO2(31.65 %wt loss)+ 3H2O(19.42 %wt loss)

It shows that remaining 37% mass is that of aluminum oxide. Figure 4.8 (a &

b) shows the SEM images of whiskers after calcinations at 650ºC. In spite of removal

of volatile ingredients (NH3, H2O and CO2), size of individual spikes is not affected

significantly indicating that 63% weight loss just causes formation of porous

nanostructured spikes [31]. This is very important property for applications like

catalysis and gas-sensing etc.

Fig. 4.7: Thermal gravimetric curve of AACH sample.

60

Fig. 4.8(a-b): SEM images of hydrothermally synthesized U10 AACH urchins calcined at 650ºC; (a)

Low magnification; (b) High magnification

4.2.1 Fourier Transform Infrared Spectroscopy

Figure 4.9 shows FTIR results of AACH sample. As prepared AACH sample

shows infra-red absorptions due to aluminium, hydroxide, ammonia and carbonate

groups. Peaks at 3446 and 978 cm-1

are due to stretching and bending of hydroxyl

group in AACH [32]. Similarly, bands at 3000, 3173 cm-1

are due to symmetric and at

1385, 1721 cm-1

are due to asymmetric stretching of NH4+. The bands at 468, 611, 733

and 848 cm-1

are ascribed to vibrational modes of Al-O, while strong band at 1105 cm-

1 is due to Al-O-Al symmetric vibrations. The peaks at 1444 & 1539 cm

-1 are due to

asymmetric stretching of CO3 bonds.

Calcination at 250ºC results in a decrease in the intensity of almost all the

peaks, while no functional group is completely removed at this temperature. Sharpness

of absorption peaks is further decreased by calcining the sample at 300ºC. Individual

peaks corresponding to various groups appear to diffuse together to form wider bands.

The intensities of these broad bands highly decrease when calcination is performed at

400oC except that a broad shoulder formed by combining the peaks corresponding to

vibrational modes of Al-O (400 to 1000 cm−1

). Increasing calcinations temperature

further to 650oC causes the broad shoulder of Al-O bond to develop to maximum,

while all other groups disappear. Therefore, formation of pure alumina is ensured after

calcinations at 650oC, although appearance of broad shoulder suggests the absence of

crystallinity and formation of amorphous alumina phase [33-34]. It must be

emphasized that the amorphous porous nanostructured alumina will be of high value

for catalytic applications.

(a) (b)

61

Fig. 4.9: FTIR spectra of AACH sample calcined at different temperatures; the spectrum of as prepared

sample is also given for comparison.

Amorphous phase of calcined AACH was also confirmed by XRD analysis.

Figure 4.10 shows XRD patterns of U10 sample calcined at 400oC and 1200

oC. The

XRD pattern of as prepared AACH sample is also presented for comparison. It is

observed that structure of as-prepared sample is that of AACH (JCPDS Card No. 42-

0250). After calcining at 400ºC, almost all the volatile ingredients are removed and

structure becomes amorphous that is predominantly alumina. Further raising the

calcination temperature to 1200ºC results in the formation of pure α-Alumina phase

(JCPDS Card No. 01-010-0173).

62

Fig. 4.10: XRD patterns of AACH sample calcined at different temperatures; the spectrum of as prepared

sample is also given for comparison.

4.3 Conclusions

Ammonium aluminum carbonate hydroxide (AACH) was prepared by

hydrothermal route using urea and aluminum nitrate as precursor materials. It was

observed that the amount of urea is very important to control the morphology of

resultant nanostructured AACH. At lower urea contents, relatively larger particles are

formed, which transform to urchin like structures and then to AACH whiskers as urea

content is increased. Size of these urchin like nano-architectured particles was up to 4

µm. Diameter of the individual spikes of urchins was about 60 nm. However, fully

developed (separated) AACH whiskers formed at higher urea contents had diameters

of about 120 nm, although they were mostly agglomerated. TGA results showed that

AACH whiskers could be completely converted to alumina whiskers through

calcination at a temperature about 650oC. Complete removal of volatile ingredients of

AACH at 650oC was also verified by FTIR. FTIR and XRD results showed that low

temperature calcination gave amorphous alumina, which could be converted to -

alumina at higher temperatures such as at 1200oC.

63

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[17] Zhang J, Wei S, Lin J, Luo J, Liu S, Song H, et al. Template-Free Preparation of

Bunches of Aligned Boehmite Nanowires. J Phys Chem B. 2006;110(43):21680-3.

[18] Zhu HY, Gao XP, Song DY, Bai YQ, Ringer SP, Gao Z, et al. Growth of

Boehmite Nanofibers by Assembling Nanoparticles with Surfactant Micelles. J Phys

Chem B. 2004;108(14):4245-7.

[19] Hou HW, Xie Y, Yang Q, Guo QX, Tan CR. Preparation and characterization of

gamma-AlOOH nanotubes and nanorods. Nanotech. 2005;16(6):741-5.

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flowerlike Cu2O nanoarchitectures by a solution phase route. Cryst Growth Des.

2007;7(1):87-92.

[21] Zhou W, Wang ZL. Three-Dimensional Nanoarchitectures: Designing Next-

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Precipitation of Aluminum Basic Sulfate. J Am Ceram Soc. 1998;81(5):1377-9.

[23] Kara F, Sahin G. Hydrated aluminium sulfate precipitation by enzyme-catalysed

urea decomposition. J Eur Ceram Soc. 2000;20(6):689-94.

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[25] Madhusudan P, Ran J, Zhang J, Yu J, Liu G. Novel urea assisted hydrothermal

synthesis of hierarchical BiVO4/Bi2O2CO3 nanocomposites with enhanced visible-

light photocatalytic activity. Appl Catal B-Environ. 2011;110(0):286-95.

[26] Mitsui I, Shima T, Nishiyama F, Komatsu T. Physical properties of aluminum

compound, NH4AlCO3(OH)2. Nippon Denshi Zairyo Gijutsu Kyokai Shuki Koen

Taikai Koen Gaiyoshu. 1977;14:5-8.

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[27] Johnson-Maynard J. http://soils.cals.uidaho.edu/soil205-90/lecture%207/index.h.

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[29] Li G-C, Liu Y-Q, Guan L-L, Hu X-F, Liu C-G. Meso/macroporous γ-Al2O3

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[30] Wu ZS, Shen YD, Dong Y, Jiang JQ. Study on the morphology of alpha-Al2O3

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66

Chapter 5

5 Effect of Deflocculants on the Properties of

Composite

Zirconia (ZrO2) nanocomposites offer excellent mechanical properties making

them highly attractive for a wide range of applications as structural & bio-ceramics

[1]. Due to high corrosion and fracture resistance of ZrO2 it is being used as dental

implants [2]. In comparison with other allotropic forms of zirconia, tetragonal

zirconia is more famous owing to the transformation-toughening phenomenon, which

gives good toughness along with high strength [3-5].

Although 3mol% Y2O3 stabilized tetragonal ZrO2 (TZ-3Y) presents good

toughness and excellent corrosion resistance, its hardness is less than many other

ceramics such as Al2O3, SiC, Si3N4 etc. [6]. The hardness of conventionally available

TZ-3Y is 12GPa [7] while that of Al2O3 is higher than 15 GPa [8]. Any compositional

modifications to improve hardness are accompanied by deterioration of toughening.

Composite formation is an easy route to improve strength or hardness of TZ-

3Y without affecting its other properties. For biological structural applications, a

biocompatible reinforcement is required. In this regard, Al2O3 is an excellent choice.

For last few decades, a lot of work has been devoted to synthesis and characterization

of zirconia-alumina composites in order to obtain better mechanical properties, which

excel both of the composite constituents [9].

Among various morphologies of the reinforcement constituent, whiskers with

relatively high aspect ratio provide best results in terms of strength enhancement.

Such morphologies are also useful for further improvement in toughness. Different

techniques for production of alumina whiskers have been reported like hydrothermal

[10], carburization, pyrolysis [11] and vapor phase growth [12] etc. Among these,

hydrothermal technique as reported by J. Li et al [10] is unique due to its low

temperature synthesis, cost effectiveness, less time consumption and use of common

precursor salts. We therefore prepared ammonium aluminum carbonate hydroxide

(AACH) precursor whiskers through hydrothermal technique, which were

subsequently transformed into alumina whiskers in-situ during sintering.

67

There have been various reports on improvement of mechanical properties by

whisker-reinforcement of ceramic composites [13-15]; however, alumina whiskers

reinforced tetragonal zirconia composites have been reported very rarely. Only worth

mentioning work is that by Nevarez-Rascon et al. [16]. The whiskers used by them

were of very high aspect ratio (1400 approx.), which resulted into severe

agglomeration. The presence of large agglomerates generally leads to low sintered

density [17] and consequently poor mechanical properties of structural ceramics. In

the present work, emphasis has been laid on prevention of agglomeration.

It must be emphasized that the control of reinforcement geometry such as the

aspect ratio of whiskers as well as its uniform dispersion in the matrix is very

important to obtain real advantage for composite formation. We know that

intermolecular bonding like van der waals and hydrogen bonding are very weak

interactions, but their influence increases if they become very large in number (like in

nanomaterials with very small size and large surface area). For example in DNA

(having 2nm size), the two helixes are held together by numerous hydrogen bonds. In

nano-whiskers with large surface area, this fact becomes particularly relevant and

hydrogen bonding induced agglomeration is really a big issue. Composites formed

without good dispersion of reinforcement showed insignificant deterioration of

properties. For increasing, the dispersion of reinforcement through use of some

suitable dispersing agent / deflocculant is a widely accepted approach for composites

[18-21].

In chapter 4, formation of alumina whiskers has been successfully achieved

using Ammonium aluminum carbonate hydroxide (AACH) precursor whiskers by

thermal decomposition at 1200oC. The precursor whiskers were prepared by

hydrothermal synthesis using urea and aluminum nitrate. The control of the size and

morphology was successfully achieved. An addition of optimum amount of 12 g urea

in 70 ml of water containing 25 g of aluminum nitrate, led to the formation of well-

established morphology of isolated whiskers with a diameter of 120 nm and a length

of about 5 µm.

AACH is more wettable in comparison with alumina. If AACH whiskers are

successfully dispersed in suspended particulate ZrO2, decomposition of AACH; then

subsequent calcination could possibly result in the formation of Al2O3 whiskers well-

distributed in the particulate ZrO2 powder. Among various techniques for dispersion

of reinforcement in the powder of matrix phase, use of deflocculants in the

68

suspensions seems attractive. Some work related to the use of deflocculating agents

CTAB [22] and PABA [23-24] on the successful suspension of alumina particles in

has already been reported for electrodeposited metal-matrix composites. The same

technique may possibly be used for the dispersion of Ammonium aluminum carbonate

hydroxide (AACH) whiskers in particulate zirconia.

Chapter 5 describes the work done in present study on improvement of

dispersion of alumina whiskers in TZ-3Y matrix. Effect of addition of deflocculating

agents like CTAB and PABA [22, 24] for successful mixing of precursor powders has

been investigated on the density and hardness of sintered zirconia-alumina

composites.

5.1 Experimental Procedures

5.1.1 Preparation of AACH Whiskers

Alumina whiskers were produced in-situ from transformation of Ammonium

Aluminum Carbonate Hydroxide (AACH) whiskers during sintering. The AACH

whiskers were produced by hydrothermal technique (detail given in chapter 4).

5.1.2 Synthesis & Characterization of TZ-3Y+Al2O3 Composites

The various compositions of TZ-3Y-Al2O3 composite samples prepared for

study on the effect of deflocculants are given in Table 5.1. Four types of samples were

prepared: two with CTAB and two with PABA as deflocculant additions, while the

two types having same deflocculant differed in geometry of reinforcement, i.e.,

whisker or particulate. In each type amount of deflocculant was varied as shown in

Table 5.1.

For preparation of these samples measured amounts of deflocculates were

dissolved in 99.9% analytical grade ethanol. The required amounts of reinforcement,

AACH whiskers (preparation method explained in section 4.2.1.2) or high-purity α-

Al2O3 particulates (Almatis, Germany; purity >99.99%; average particle size

~0.2µm), were dispersed in this solution separately. While the required amounts of

ZrO2 + 3 mol. % Y2O3 (TZ-3Y) were dispersed in separate beakers. Sonication for

each separate suspension was performed for 30 minutes followed by mixing of the

reinforcement and matrix material suspensions. After mixing, the composite

suspension was magnetically stirred overnight at 40oC in order for the ethanol to

69

evaporate. Total volume of the ethanol-deflocculant solution was fixed for each

sample. The AACH whiskers were added on the basis of required amount of alumina

formed as a result of its decomposition.

Table 5.1: Types of samples prepared

[90%TZ-3Y+10% Al2O3(w)]100-x+[CTAB]x Where;

x = 0, 1, 5, 10, 15

w = whisker

p = particulate

[90%TZ-3Y+10% Al2O3(p)]100-x+[CTAB]x

[90%TZ-3Y+10% Al2O3(w)]100-x+[PABA]x

[90%TZ-3Y+10% Al2O3(p)]100-x+[PABA]x

Dry mixture was thermally characterized by using Mettle Toledo,

TGA/SDTA851 up to the temperature of 1000ºC with heating rate of 20ºC/min in air.

On basis of TGA results, the powder was calcined at 650ºC to remove organic

deflocculants and other volatile species such as NH4+

, OH-, HCO

3- [10] and residual

water vapors. The calcined mixture was ground for ½h and pressed uniaxially at a

load of 500MPa in a steel disk die to make pellets of 10 mm diameter. The prepared

green pellets were sintered in a high temperature (Carbolite HTF-18/8) furnace at a

temperature of 1450ºC in air for 2 hours at heating and cooling rate of 2ºC/min to get

TZ-3Y+Al2O3 sintered composites.

The density of sintered pellets was measured in ethanol through Archimedes

method using high accuracy (10-6

gm) weighing balance. The detail of this

measurement is given in the experimental procedures chapter (section 3.2.5.1).

The Vickers hardness of sintered pellets was measured by using Vickers

micro-hardness tester (Model 401MVD WOLPERT) at room temperature. Four

samples per composition and approximately 10-15 indents per measurement were

made and the average hardness was quoted. The separation between neighboring

indents was kept more than 4 diagonal lengths of indentation impression as per

ASTM standard of ASTM C 1327-99 [25]. For hardness testing, a load of 1kg was

applied for 15 seconds.

Phase analysis of the samples at different processing steps was performed by

X-ray diffraction using Bruker D8 X-ray diffractometer with Cu-Kα radiation in θ-2θ

mode. The microstructural study was carried out using JEOL-JSM5910 scanning

electron microscope (SEM) at 20 kV. Before electron microscopy, samples were

70

gold-coated using SPI-Module sputter coater in order to avoid charging during

exposure to electron beam.

5.2 Results

5.2.1 AACH Whiskers, Alumina Powder & Tetragonal Zirconia

The AACH whiskers used for this study were prepared by using higher urea

content (synthesis of these whiskers has been described in chapter 4), which result in

formation of separated fully developed whiskers. These AACH whiskers transforms

in situ to alumina during sintering as verified through X-Ray diffraction shown and

discussed in upcoming paragraphs. While SEM image of as received α-Al2O3

particulates cluster with rough spherical morphology, that were used in this study is

shown in Fig. 5.1.

Fig. 5.1: Morphology of as received α-Al2O3 particulate

Figure 5.2(a-c), showing that after sintering of the compacts of TZ-3Y+AACH

at higher temperature (1450°C), transformation of mixed phase to tetragonal zirconia

and the conversion of AACH to α-Al2O3 takes place (Figure 5.2(c)). From the figure

5.2(b & c) it is also observed that the peak around 2θ value of 35º is present in both

zirconia and alumina patterns.

71

Fig. 5.2(a-c): X-Ray diffraction of (a) as received TZ-3Y, (b) sintered TZ-3Y and (c) TZ-3Y+Al2O3(w)

pellets.

5.2.2 Effect of Deflocculants on Densification and Hardness of

Composites

Figure 5.3(a-b) shows the effect of adding different concentrations of CTAB or

PABA on hardness of 90%TZ-3Y+10%Al2O3 composites. Average Vickers hardness

for samples containing alumina whiskers as reinforcement is 14.44GPa (maximum)

while in case of alumina particulates it is 13.72GPa at a concentration of 1 wt%

CTAB (Fig. 5.3(a)). At higher concentrations of CTAB, the average hardness

decreases slightly. PABA also shows somewhat similar behavior. The values of

average hardness are 14.34 GPa and 13.65 GPa for alumina whiskers and particulates

reinforcements, respectively, at a concentration of 1 wt% PABA. These values are

higher as compared to previously reported [16] for 10 wt% alumina whiskers and pure

TZ-3Y. It is also revealed that the average hardness decreases as the amount of

deflocculating agents is increased beyond 1 wt%.

72

Figure 5.4(a-b) shows the relative densities of 90%TZ-3Y+10%Al2O3

composite samples containing CTAB and PABA. The density increases as

deflocculating agent was added. At a concentration of 1 wt% CTAB, the sample

containing alumina whiskers shows a relative density of 93% while sample containing

alumina particulates shows relative density of 95.5%. Density also decreases with

increase in CTAB beyond 1 wt%.

Fig. 5.3(a-b): Hardness of pellets (a) [90%TZ-3Y+10% Al2O3]100-x+[CTAB]x (b) [90%TZ-3Y+10%

Al2O3]100-x+[PABA]x.

(a)

(b)

73

Fig. 5.4(a-b): Relative density of pellets (a) [90%TZ-3Y+10% Al2O3(w)]100-x+[CTAB]x (b)

[90%TZ-3Y+10% Al2O3(w)]100-x+[PABA]x.

(b)

(a)

74

The change in relative density with addition of PABA also shows a similar

behavior and a maximum in density is observed at 1 wt% PABA, beyond which the

density decreases. At 1 wt% PABA, densities of 91% & 95% are observed for

samples containing whiskers and particulates, respectively.

Figure 5.5 shows SEM images of surfaces of sintered pellets containing Al2O3

particulates (Fig. 5.5 (a-b)) and Al2O3 whiskers (Fig. 5.5 (c-d)) in the matrix of TZ-3Y

with 1 wt% CTAB. Low magnification images (Fig. 5.5 (a & c)) show that surface is

crack free, and high magnification images (Fig. 5.5 (b & d)) reveal that samples are

quite compact as little porosity is visible. The average grain size in both cases is

approximately 0.3µm.

Fig. 5.5(a-d): SEM micrographs at low and high magnification of the sintered pellets of (a&b) [90%TZ-3Y

+10%Al2O3(p)]99+[CTAB]1 (c&d) [90%TZ-3Y +10%Al2O3(w)]99+[CTAB]1.

To examine morphology and to determine the elemental composition of

whiskers, the pellets were fractured and observed under SEM. The fractographs at low

and higher magnifications are shown in Fig. 5.6. From the micrographs, it is revealed

that the whiskers have retained their morphology and in the chapter 6 on mechanical

properties of these composite, it is shown through EDS analysis that these whisker

(a) (b)

(c) (d)

75

morphologies have relatively higher aluminum content than matrix confirming it to be

alumina whiskers.

Fig. 5.6(a-b): Fractograph of composite (a) at lower magnification (b) at higher Magnified identifying the

presence of whisker in composite matrix.

5.3 Discussion

The difference in density of whiskers and particulates containing composite (Fig.

5.4(a-b)) suggests that whiskers resist complete shrinkage during sintering and very

small voids or porosity remain inside the composite which decreases its density.

The increase in density and hardness, Fig. 5.3(a-b), with addition of CTAB

and PABA is due to uniform dispersion of whiskers and particles. The density of

compacts depends upon many factors like applied load; morphology, size & shape of

the reinforcement; state of die wall lubrication; agglomeration etc [23]. The density

can also be improved due to better dispersion of reinforcement. If dispersion of

reinforcement increases then sample will have uniform shrinkage during sintering and

this result in decreasing the porosity and improving density as observed in Fig. 5.4(a-

b).

5.3.1 Role of CTAB in Dispersion of Alumina

There are two reasons behind better dispersion of Al2O3. First and main reason

as already mentioned, is due to addition of deflocculating agents and other is the use

of AACH and its subsequent transformation into Al2O3 in-situ during sintering.

CTAB is a well-known cationic surfactant. Addition of CTAB increases zeta potential

of Al2O3 [26]. In general, electrostatic repulsive force between particles increases

(a) (b)

76

with the increase in zeta potential. It is assumed that when alumina whiskers /

particulates are added into the CTAB containing solution then CTAB adsorbs on

surface of the alumina and produces a positive charge layer on whiskers / particles

surface. During mixing when secondary particles break then this surfactant

immediately coat the newly formed surfaces and these particles then repel each other.

In this way, agglomeration of particles is prevented and dispersion is improved. In

fact, increase in electrostatic repulsive force is responsible for the dispersion of

alumina (Al2O3).

77

Fig. 5.7(a-b): (a) FTIR Spectrum for samples containing 20wt% alumina inform of AACH, with and

without 1wt% CTAB, before and after calcinations. (a-i) AACH with CTAB calcined at 650ºC (a-ii)

AACH without CTAB calcined at 400ºC (a-iii) AACH with CTAB uncalcined (a-iv) AACH without

CTAB uncalcined. (b) magnification of the FTIR peak at wave number of 3437cm-1.

78

To further investigate the role of CTAB in dispersing alumina structures the

help of FTIR was sought. Fig. 5.7 shows Fourier Transform Infrared Spectroscopy of

samples containing 20wt% Alumina in form of AACH (NH4Al(OH)2CO3) before and

after calcination. Most of the peaks correspond to AACH [27]. Peaks in the region

3173cm-1

and 1388 cm-1

are due to symmetric and asymmetric bond stretching of NH4

of AACH. While the bands at 1382, 1444 and 1540cm-1

corresponds to asymmetric

stretching modes of CO3-2

in AACH. However, after calcinations these peaks of NH4

& CO3-2

are eliminated this means that volatile ingredients of AACH have been

removed as a result of calcination.

From this figure, it is observed that visually, the peaks of samples with and

without CTAB have no clear difference. However, on magnification a slight shift can

be observed in some of the peaks of two types of samples. For example, the peak

corresponding to O-H stretching appears at 3437cm-1

wavenumber. However, with

addition of 1wt% CTAB peak shifts slightly to higher wavenumber of 3441cm-1

fig.

5.7(b). This indicates that with addition of surfactant or deflocculating agent (CTAB)

hydrophobicity of the resultant materials increases [28]. Alternatively, we can say that

CTAB has attached with particles or whiskers by replacing water molecules. As the

CTAB is, a cationic surfactant so it has produced the net positive charge on particles

and this results in increasing inter-particle repulsion and does not let the particles to

agglomerate so dispersion of the reinforcement enhances. This is in good agreement

with already published work reporting that with addition of CTAB zeta potential of

Al2O3 increases [26]. In general, electrostatic repulsive force between particles

increases with the increase of zeta potential.

Fig. 5.8: Structure of PABA molecule

79

5.3.2 Role of PABA in Dispersion of Alumina

Similarly, addition of PABA also results in improving density and hardness of

the composite. PABA an organic compound (with formula (H2NC6H4CO2H) and is

also used as dispersing agent [29]. The structure of para-aminobenzoic acid (PABA:

Formula C7H7NO2) is shown in Fig. 5.8. This figure shows that PABA has acidic

(COOH-) and basic (NH2

+) groups at para sites of a single molecule [30]. These NH2

or COOH derivates of PABA are playing their role in dispersing the AACH in matrix

of TZ-3Y during mixing. It is speculated that there is a possibility of intramolecular

ionization, after which the carboxyl group of PABA specifically absorbs on surface of

AACH whisker while the ionized amine group which has positive charge, is towards

solution and this give alumina a net positive charge [31].

The other factor contributing towards improvement in properties is use of

AACH instead of Al2O3, and during sintering this AACH converts into Al2O3

whiskers. The Al2O3 whiskers formed from AACH have larger specific surface area

and have better sorption capacity and supported its dispersion in the matrix [32].

Other important observation that can be made from Fig. 5.3 (a-b) and Fig. 5.4

(a-b), is decrease in density and hardness with increasing the amount of CTAB or

PABA above 1 wt% concentration. When deflocculating agent is added above a

certain amount then it produces small amount of porosity which may results in

deteriorating properties of the composite [22]. This can be confirmed from

thermogravimetry (TG) analysis curve at Fig. 5.9 for samples containing 25%

deflocculating agent. TG thermogram of pure AACH in Fig. 5.9 shows total weight

loss of 63%. This is in good agreement with literature [10] that AACH

(NH4Al(OH)2CO3) decomposes to Al2O3 according to following reaction (as already

given in section 4.3),

2 NH4Al(OH)2CO3 → Al2O3(37%)+ 2NH3(12.23%)+ 2CO2(31.65%)+ 3H2O(19.42%)

For samples with [90%TZ-3Y+10% Al2O3(w)]75+[CTAB/PABA]25 amount of

precursor materials used for making 3gm mixture were AACH=0.818gm (which

converts into 0.3 gm Al2O3(w)); TZ-3Y=2.7gm & CTAB/PABA=0.88 gm. Therefore,

out of total weight (of mixture i.e. 4.398gm) remaining weight should be 3gm (i.e.

68.213% of total weight) but this comes out to be 69.7% and 71.3% for this specific

80

samples containing 25wt% CTAB and PABA respectively. This means that out of the

added 25%, approximately 1.5% and 3% weight of CTAB and PABA respectively is

not removed and remains within the composite even at high temperature.

Approximately, 3~4% combustion products of CTAB remains even at high

temperatures [33]. Therefore, when relatively large amount of deflocculating agent

was added then the remaining reaction products produce low density regions and

results in deterioration of density and hardness. Similarly, maximum hardness

achieved for CTAB is higher than the value achieved from PABA. The reason is

obvious from Fig. 5.9 that as relatively more mass of PABA as compare to CTAB

remains within the samples after heating, CTAB produces better results than PABA.

Fig. 5.9(a-c): TGA of (a) pure AACH and [90%TZ-3Y+10%Al2O3(W)]75with (b) CTAB and (c) PABA

Composite.

It is observed that hardness of the samples with Al2O3 whiskers is higher than

the samples with equal amount of Al2O3 particulates. This improvement in

mechanical properties of whiskers reinforced composites is due to crack deflection

and crack bridging effects [34] offered by the added whiskers the detail is described in

chapter 6.

81

5.4 Conclusions

Mechanical properties of TZ-3Y+Al2O3 nanocomposite can be increased by careful

selection of deflocculant. Addition of CTAB has slightly better effect on improving

the density and hardness of the composite as compare to PABA addition. The

maximum hardness for CTAB (1 wt%) was 14.44GPa at a composition of [90%TZ-

3Y+10% Al2O3(w)]99+[CTAB]1 while maximum hardness of 14.34GPa was obtained

for [90%TZ-3Y+10%Al2O3(w)]99+[PABA]1. The density of Al2O3 particulates

containing samples was higher than samples with Al2O3 whiskers, while hardness of

Al2O3 whiskers containing samples was more than Al2O3 particulates containing

samples.

82

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[17] Ozkan N, Briscoe BJ. Characterization of die-pressed green compacts. J Eur

Ceram Soc. 1997;17(5):697-711.

[18] Bondeson D, Oksman K. Polylactic acid/cellulose whisker nanocomposites

modified by polyvinyl alcohol. Compos Part A-Appl S. 2007;38(12):2486-92.

[19] Bondeson D, Oksman K. Dispersion and characteristics of surfactant modified

cellulose whiskers nanocomposites. Compos Interfaces. 2007;14(7-9):617-30.

[20] Petersson L, Kvien I, Oksman K. Structure and thermal properties of poly(lactic

acid)/cellulose whiskers nanocomposite materials. Compos Sci Technol. 2007;67(11–

12):2535-44.

[21] Jiang L, Morelius E, Zhang JW, Wolcott M, Holbery J. Study of the Poly(3-

hydroxybutyrate-co-3-hydroxyvalerate)/Cellulose Nanowhisker Composites Prepared

by Solution Casting and Melt Processing. J Compos Mater. 2008;42(24):2629-45.

[22] Liu D, Yan Y, Lee K, Yu J. Effect of surfactant on the alumina dispersion and

corrosion behavior of electroless Ni-P-Al2O3 composite coatings. Mater Corros.

2009;60(9):690-4.

[23] Vishista K, Gnanam FD. Effect of deflocculants on the densification and

mechanical properties of sol-gel alumina. Mater Lett. 2004;58(11):1665-70.

[24] Guimarães FAT, Silva KL, Trombini V, Pierri JJ, Rodrigues JA, Tomasi R, et al.

Correlation between microstructure and mechanical properties of Al2O3/ZrO2

nanocomposites. Ceram Int 2009;35(2):741-5.

[25] ASTM.org. Standard test method for vickers indentation hardness of advanced

ceramics Philadelphia, PA,: West Conshohocken, PA; 1991.

[26] Chen L, Wang L, Zeng Z, Zhang J. Effect of surfactant on the electrodeposition

and wear resistance of Ni-Al2O3 composite coatings. Mater Sci Eng: A. 2006;434(1-

2):319-25.

[27] Zhu Z, Sun H, Liu H, Yang D. PEG-directed hydrothermal synthesis of alumina

nanorods with mesoporous structure via AACH nanorod precursors. J Mater Sci.

2010;45(1):46-50.

[28] Xue W, He H, Zhu J, Yuan P. FTIR investigation of CTAB-Al-montmorillonite

complexes. Spectrochim Acta A. 2007;67(3-4):1030-6.

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[29] Sumita S, Rhine WE, Bowen HK. Effects of Organic Dispersants on the

Dispersion, Packing, and Sintering of Alumina. J Am Ceram Soc. 1991;74(9):2189-

96.

[30] Galaction AI, Camarut M, Cascaval D. Separation of p-aminobenzoic acid by

reactive extraction: 1: Mechanism and influencing factors. Chem Ind Chem Eng Q.

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[31] Kitayama M, Pask JA. Formation and control of agglomerates in alumina

powder. J Am Ceram Soc. 1996;79(8):2003-11.

[32] Fu G-f, Zhang J-x, Liu J, Wang Z-f. Wear behavior of Al-Si alloy matrix

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whisker-reinforced MoSi2. J Am Ceram Soc. 1987;70(4):79-81.

85

Chapter 6

6 Effect of Whisker Addition on Mechanical

Properties of Composite

Among various ceramic composite systems investigated so far, ZrO2-Al2O3(w)

system has been proved to be a potential candidate for various load bearing

applications [1-2]. It is expected, that the reinforcement of the zirconia with alumina

whiskers may result in improvement of the mechanical properties of the composite.

The knowledge of both these properties is very important from application point of

view.

As discussed in chapter 4, AACH whiskers are formed using aluminum nitrate

and urea. The concentration of the latter proves to be a predominant factor in

determining the morphology of the AACH whiskers. AACH has been shown to

decompose to form alumina.

In chapter 5, it was shown that alumina whiskers could be successfully

dispersed in the particulate zirconia by first suspension mixing of AACH whiskers are

zirconia followed by calcinations. This has been shown to facilitate the formation of

ZrO2-Al2O3(w) composites by sintering at 1450°C.

It has been shown in chapter 5 that the mechanical properties of TZ-3Y+Al2O3

nanocomposite can be improved by careful selection of deflocculant. Addition of

CTAB is slightly better in improving the density and hardness of the composite as

compared to PABA. The maximum hardness for CTAB (1wt%) was 14.44 GPa at a

composition of [90%TZ-3Y+10% Al2O3(w)]99+[CTAB]1 while maximum hardness of

14.34GPa was obtained for [90%TZ-3Y+10%Al2O3(w)]99+[PABA]1. Although the

density of Al2O3 particulates containing samples was higher than the samples with

Al2O3 whiskers, the hardness of Al2O3 whiskers containing samples was

comparatively higher.

Hardness of ceramic materials is one of the most frequently measured

properties. Its value helps designers to characterize the ceramics resistance to

deformation and fracture. Determination of hardness value of a ceramic material is

86

very crucial for approximately all load bearing application. It is crucial for prosthetic

hip-joint balls and sockets, molds and dies, wear and abrasion-resistant parts, ballistic

armor, valves, seals and cutting tools. Approximately all ceramics applications lists

minimum hardness requirements. For example, Vickers hardness requirement for

zirconia based surgical implants is equal to or greater than 11.8 GPa (ASTM F 1873-

98) [3]. More than 60% worldwide published hardness values for ceramic materials

are Vickers hardness [4]. Due to its vital importance, in present study hardness of the

synthesized composite is also measured by using vickers micro-hardness test.

Like hardness, determination of strength of a material is also very important

for its application, design data, quality assurance and predicting its performance over

service life. Ceramic materials are brittle in nature and determination of their tensile

strength is very difficult due to three main reasons. First, it is difficult to prepare

specimens of specific geometry required for tensile tests. Second, it is also difficult to

grip them in tensile test machine jaws without fracturing. Third, for tensile test they

are required to be kept perfectly aligned because a small misalignment may result in

failure. In view of these difficulties, flexural strength testing is widely recommended

for ceramics [5]. Moreover, fracture toughness of ceramic is directly proportional to

its flexural strength [6]. Therefore, flexural strength becomes one of the most

demanding physical properties of ceramic materials used under stress environment.

According to best of author’s knowledge, no study on flexural strength of alumina

whisker reinforced TZ-3Y composite has been reported.

In Chapter 6, the effect of whisker concentration on the hardness and flexural

strength of zirconia-alumina(w) composites has been studied. It is expected that

current findings will help enhance the operational life of bio-ceramics for dental and

orthopedic applications.

6.1 Experimental Procedures

Ammonium Aluminum Carbonate Hydroxide (AACH) whiskers were prepared

from precursor aqueous solution of aluminum nitrate and urea at 120oC. The detail

given in chapter 4.

AACH whiskers and high purity (99.99%) tetragonal ZrO2-3Y powder were

separately suspended in ethanol using ultrasonic bath (Elma E 30 H, 37 kHz) with or

without 1% CTAB. Both suspensions were then mixed together, kept at 60°C while

87

stirring until nearly all the ethanol was evaporated. Samples with varying amounts of

AACH corresponding to 5, 10, 15, 20, 30 & 50 % Al2O3 were prepared.

Samples without CTAB were calcined at 400°C, while CTAB containing

samples were calcined at 650ºC (based on TGA studies presented in chapter 5) for

complete removal of CTAB and volatile ingredients of AACH. Pellets of about 1.5 g

with a diameter of 10 mm were then prepared at a load of 500MPa. Green pellets

were sintered for two hours at 1450°C in a muffle furnace (Carbolite HTF-18/8).

Heating and cooling rates were maintained as 1°C/min.

Density of the sintered pellets was measured by geometric method, while

hardness was determined by using Vicker Microhardness Tester (KARL FRANK,

Germany) at a load of 10Kgs.

Flexural strength of the sintered composite samples was measured by using

universal testing machine made by SANS, China. Due to simplicity of method and

ease of sample preparation ring on ring flexural strength test method [7] was adopted

as per ASTM F394–78 [8] standard with a cross head speed of 0.1 mm/min. For each

composition, measurement was performed on at least six samples and data was

reported as an average along with standard deviation showing how much variation or

"dispersion" exists from average.

6.2 Results and Discussion

Figure 6.1 shows typical XRD patterns of sintered pellets with varying amounts of

alumina prepared from AACH whiskers dispersed in zirconia powder, as detailed in

the experimental methods. Formation of -alumina in the matrix of tetragonal

zirconia is confirmed. Weak reflections of -alumina in XRD patterns for 10wt%

alumina (pattern ii) and 20wt% alumina (iii) may be due to low concentration, and

low X-ray scattering cross-section of aluminum as compared to zirconium.

Nevertheless, existence of X-ray reflections was confirmed at slow scans as typically

shown in Fig. 6.1(b) for sample with 10 wt% alumina. XRD patterns of the sintered

pellets were not affected by the addition of CTAB in the precursor suspensions.

88

Fig. 6.1(a-b): (a) XRD patterns of sintered ZrO2+Al2O3(w) nanocomposites with various amounts of Al2O3

whiskers; (b) Slow scan rate XRD pattern around (012) reflection of α-Al2O3 for sample with 10 wt.% Al2O3

whiskers.

Figure 6.2(a) shows a typical fracture surface of sample with 10wt% alumina,

with corresponding EDS spectrum (Fig. 6.2(c)) that confirms 10wt% alumina in the

sample. A corresponding high magnification image is shown in Fig. 6.2(b), along with

EDS spectrum obtained from a typical whisker (Fig. 6.2(d)) confirming it to be

alumina; the counts from zirconium seem to come from matrix. It is clear that

whiskers retain their shape during their transformation from AACH to alumina and

subsequent sintering. Whiskers are also seen to be well surrounded by zirconia matrix.

SEM images were also obtained after mechanical polishing with subsequent thermal

etching, as typically shown in Fig. 6.2(e) and 6.2(f). Relatively compact structure is

obtained at 10 wt% alumina, while high porosity regions are numerously seen in

sample with 50 wt% Al2O3, the poor ability to sinter at higher concentrations may

possibly be related with agglomeration.

For obtaining improved mechanical properties of ceramics, porosity is very

deleterious due to two main reasons. First, pores reduce cross-sectional area on which

load is applied. Second, they act as stress concentrators [9] which amplifies the

applied stress. The well-known empirical relationship between flexural strength and

porosity suggests that flexural strength decreases exponentially with increase in the

(b)

(a)

89

volume fraction of porosity [10]. Therefore, to obtain material with good hardness and

flexural strength pore free microstructure is necessary.

0.00 1.00 2.00 3.00 4.00 5.00 6.00 7.00 8.00 9.00 10.00

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40

80

120

160

200

240

280

320

360

400

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a

0.00 1.00 2.00 3.00 4.00 5.00 6.00 7.00 8.00 9.00 10.00

keV

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0

15

30

45

60

75

90

105

120

135

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Fig. 6.2(a-f): SEM micrographs of: (a) & (b) as-fractured 90% ZrO2(3Y)-10% Al2O3(W) pellet along with

(c) EDS spectra on a 1.5µm × 1.5µm area in (a) and a (d) whisker in (b); (e) & (f) polished and thermal

etched surfaces of samples with 10 wt% and 50 wt% alumina whiskers, respectively.

As shown in Fig. 6.3(a), relative density of sintered pellets decreases with

alumina content. Nevertheless, addition of CTAB results in a slightly higher density.

This strengthens our view that improved dispersion of alumina whiskers in zirconia

may enhance ability to form relatively compact composites.

(a)

(c)

(b)

(d)

(e) (f)

90

6.2.1 Whisker Concentration Effect on Hardness

Figure 6.3(b) shows hardness of sintered ZrO2-Al2O3 pellets as a function of

alumina content, with or without addition of CTAB in the precursor suspensions. A

maximum hardness of about 14 GPa is observed for 10wt% of Al2O3 whiskers. For all

the cases, addition of CTAB results in a relatively higher hardness as related with

improved compactness of composites (Fig. 6.3(a)). Possibly due to similar reasons

(i.e., increased porosity) hardness decreases significantly with addition of 20 wt%

alumina whiskers and above. For comparison, hardness of composites formed by

Nevarez-Rascon et al. [11], employing high aspect ratio alumina whiskers has also

been shown in the figure. High achievable hardness in case of present study seems to

be related with improved dispersion, and compactness. Improved dispersion may also

partly be related with ease in dispersion of precursor AACH instead of alumina

whiskers.

An important factor in improvement of dispersion of alumina whiskers may be

a drastic volume change associated with transformation of AACH whiskers into

alumina whiskers. This results into production of mesoporous alumina nanostructure

[12]. This mesoporous structure at higher energy state seems also to assist in

improving sinterability of the composite. Since ultrasonic bath with a frequency of 37

kHz was used to disperse AACH whiskers and zirconia powder in alcohol and it is

known that minimum agglomerate size in ultrasonic bath is dependent on the

frequency [13], an average agglomerate size between 100 to 200 nm would be

expected [14]. A subsequent drastic decrease in size during calcination results in very

small alumina agglomerate size at any localized region provided that sufficient

zirconia particles are present around to surround it. However, as volume percentage of

AACH whiskers is increased, sufficient zirconia particles may not be available to

surround the dispersed whiskers and re-agglomeration of whiskers may take place. A

further improvement in dispersion has been obtained due to deflocculating effect of

cationic surfactant CTAB [15].

91

Fig. 6.3(a-b): (a) Relative densities and (b) Vicker hardness of sintered pellets having various

amounts of alumina whiskers, with & without CTAB. For comparison hardness values from Nevarez-

Rascon et al. are also given.

(a)

(b)

92

6.2.2 Whisker Concentration Effect on Flexural Strength

On the basis of above mentioned studies on the hardness of the composite, only

samples containing up to 20 wt% Al2O3(w) were characterized for flexural strength

measurement. Figure 6.4(a-b) shows the effect of addition of alumina whiskers on

flexural strength of the composites. For pure TZ-3Y, the value of flexural strength is

945±35 MPa (measured through ring on ring test method). This flexural strength

value for TZ-3Y is a little bit lower than its previously reported value of 1000 MPa

[16] measured through three or four point flexural test technique. The flexural

strength is generally lowered due to presence of cracks or flaws. The measurement of

flexural strength through ring on ring technique is more authentic because, in

three/four point test methods where the specimens used are in form of bars, cracks

that are parallel to longitudinal direction of the bar do not participate in lowering the

flexural strength. While in the ring on ring technique, the crack that is present in any

direction in sample lowers the flexural strength [17] and depicts the real strength

value. This is the reason that strength value varies from one technique to the other

[18].

The strength data in Fig. 6.4 show some scattering. In case of ceramics, the

failure starts from a flaw that may be introduced during synthesis. These flaws may be

in the form of cracks, pores, inclusions or surface scratches. Due to statistical

distribution in size, location and geometry of the flaws, strength also shows some

scattering [19].

The value of flexural strength (Fig. 6.4(a&b)) increases with addition of

whiskers and reaches maximum values of 1325±65 MPa and 1212±60 MPa at 10

wt% whisker concentration, for the samples with and without 1 wt% CTAB,

respectively. These strength values are much higher than that of TZ-3Y and real tooth

[20]. However, further addition of whiskers results in a decrease in the flexural

strength.

93

Fig. 6.4(a-b): Flexural strength of the TZ-3Y+Al2O3 composite with different weight fraction of Al2O3

whiskers: (a) with CTAB and (b) without CTAB.

6.2.2.1 Flexural Strength Increase with Whiskers from 0-10 Wt%

To investigate the reason for increase in composite strength with increasing

whisker concentration (from 0-10 wt% whiskers), refer back to Fig. 6.2(a). It shows

that alumina whiskers have retained their morphology; therefore, these whiskers have

produced the crack deflection and bridging effects, which result in improving the

mechanical properties of the composite. The phenomenon of crack deflection can also

be verified from Fig. 6.5; a scanning electron micrograph showing the zigzag path of

Vickers indentation-induced crack. Crack deflection is the operative process reported

for these whisker reinforced composites. The morphology encircled on the

micrograph in Fig. 6.5 appears like a ligamentary bridge (formed by the whisker)

between opposite surfaces of the crack. We know that whiskers are randomly oriented

in the TZ-3Y matrix. Under applied indenter load a crack is produced

which encounters whiskers as it propagates. This encounter can happen near whisker

edge or its centre. If the encounter is near the whisker edge, the crack deflects around

it due to crack bowing and follows a zigzag path [21]. But if this encounter is near the

94

whisker centre then the crack tip appears to arrest as it encounters the interface and

then reappears in the matrix ahead of the whisker, leaving whisker intact and

functioning as a ligamentary bridge [22]. This intact type bridging by the whisker is

termed as ―frictional bridging‖ [23-25] and results in improved strength of the

composite.

Fig. 6.5: SEM micrograph of crack formed at the corner of the Vicker hardness indenter on the sintered

surface of the TZ-3Y+10wt%Al2O3 composite sample, showing crack deflection and crack bridging

effects. The morphology encircled on the micrograph is a ligamentary bridge between two surfaces of

the crack formed by alumina whisker.

6.2.2.2 Flexural Strength Decrease with Whiskers above 10 Wt%

From Fig. 6.4(a-b) we can conclude that whisker toughening effect is prominent

up to the concentration of 10 wt% Al2O3(w), beyond which the toughness decreases. It

must be emphasized that two opposing factors are operative in the whisker-reinforced

composites. One is the toughening effect and the other is agglomeration and a

consequent decrease in density. The toughening effect plays role through various

processes to hinder crack propagation such as crack deflection, crack bowing and

crack bridging, etc. Agglomeration of whiskers is not so prominent at lower

95

concentrations and therefore with increasing whisker content, the flexural strength

and toughness increase. However, as the concentration of whiskers goes beyond a

critical value (10 wt % in the present case), the agglomeration effect starts

dominating. The agglomerates create low density regions and result in lowering of

strength. As the whisker content is increased further, more low density regions are

created and more decrease in flexural strength is observed. This conclusion is also

supported by the density versus whisker content data shown in Fig. 6.6 (a, b). Up to

10 wt% whisker content, the density remains almost constant and it starts decreasing

as the whisker content is increased further. This is true for both kinds of composites,

i.e., with and without CTAB.

Fig. 6.6(a-b): Relative density of the TZ-3Y+Al2O3 composite with different weight fraction of Al2O3

whiskers: (a) with CTAB and (b) without CTAB.

The low density regions are formed due to formation of pores / flaws around

or inside agglomerates, thus deteriorating the strength of composites [26]. According

to some reports, incompatibility stresses between whiskers and matrix are responsible

for agglomeration of the whiskers [27]. These stresses arise with increase in whisker

concentration [28]. Therefore, when the whisker concentration is lower than 10 wt%

then these stresses are small, and the strengthening mechanism is whisker toughening.

96

On the other hand, when the whisker content increases above 10 wt%, harmful effect

of incompatibility stresses becomes dominant and this results in whisker

agglomeration and decrease in flexural strength.

From Fig. 6.4(a), it is observed that the flexural strength of the samples

containing 1 wt% CTAB [15] is higher than samples without CTAB. The CTAB

increases the dispersion of the whiskers [29], which results in improving strength of

the composite.

From above discussion we can conclude that by optimizing the alumina

whisker loading (10 wt%) in TZ-3Y matrix, the composite with excellent flexural

strength (i.e. 1325±65 MPa) superior to that of real tooth (225 MPa) [20] can be

synthesized. The reason for high strength may be the retained morphology of the

whiskers resulting in crack bowing, crack deflection and frictional bridging type of

effects. The energy absorbed in these crack whisker interactions results in improving

the composite strength. With improved strength, it is expected that these composites

can be good candidates for dental implants in future.

6.3 Conclusions

A new facile technique has been developed for improvement of dispersion of

alumina whiskers in ZrO2 matrix, as reported in previous chapters. In the present

chapter, it has been seen that optimum amount of whisker addition lies around 10%,

with hardness value of as high as 14.4 GPa, and the flexural strength of ~1325MPa.

SEM images of fractured surface of the composite reveal that the whiskers have

produced the crack deflection and bridging effects which contributes towards

improving the mechanical properties of the composite. Above 10%, excessive

increase in porosity possibly related with agglomeration and poor sinterability, the

mechanical properties start deteriorating.

97

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99

Chapter 7

7 Effect of Sintering Temperature on Morphology

of Whisker and Properties of Composite

In chapter 6, it was concluded that best mechanical properties for alumina whisker

reinforced zirconia composite could be obtained at 10wt% alumina whisker loading.

Above this concentration, owing to the agglomeration of whisker and retained porosity

deteriorates the mechanical properties.

For better mechanical properties of composite, the parameters such as density,

grain size, and porosity should be controlled. Optimum sintering temperature is of

great importance as sintering influences both density and grain growth. As far as the

sintering of zirconia-alumina composite is concerned, it is a relatively well developed

process and different reports mention sintering at 1200 to 1650ºC [1-2]. For TZ-

3Y+20wt% α-Al2O3 (particulate) composite somewhat higher sintering temperature (>

1400ºC) is required for good mechanical properties [2]. However, according to best of

the author’s knowledge no study is available on sintering temperature effect on

mechanical properties and alumina whisker morphology in TZ-3Y-Al2O3(w)

composites.

In the present chapter the effect of different sintering temperatures properties

of in-situ transformed alumina whisker reinforced tetragonal zirconia ceramic

composite has been studied. Change in hardness and density of the composite was then

correlated with change in whisker morphology through microscopic evaluation.

Moreover, beneficial use of CTAB on dispersion of alumina reinforcement in these

composites has also been investigated. Finally, based on the results obtained, an

optimum sintering temperature for this type of composite has been proposed.

7.1 Experimental Procedures

For present study, alumina whiskers were transformed in-situ from Ammonium

Aluminum Carbonate Hydroxide (AACH) whiskers, which were synthesized through

hydrothermal technique as reported by J. Li et al. [3] by using high pressure reactor

100

(Limbo-350). The precursor materials were aluminum nitrate and urea and autogenous

hydrothermal was performed at 120oC for 24hrs.

Hydrothermally synthesized AACH whiskers, 3 mol% Y2O3 stabilized ZrO2

and Cetyl trimethylammonium bromide (CTAB) were used as starting materials. The

TZ-3Y and AACH whiskers were dispersed separately in analytical grade ethanol

(with or without 1 wt% CTAB) and sonicated for 30 min for breakage of agglomerates.

Both suspensions were mixed together and kept at 60°C while stirring until nearly all

the ethanol was evaporated. Samples with varying amounts of AACH corresponding to

5, 10, 15 and 20 wt% Al2O3 were prepared. Composites with higher concentration of

whisker (i.e. more than 20wt%) were not included in this study because higher

concentration produced agglomeration, which lead to deterioration of the composite

properties as mentioned in chapter 6 [4].

To remove volatile products of AACH and CTAB the samples without CTAB

were calcined at 400°C while the samples containing CTAB were calcined at 650°C.

Pellets of diameter 10 mm were prepared by using uniaxial hydraulic press at a load of

500MPa for 2 minutes. Green pellets were sintered at three different temperatures, i.e.

1400, 1500 and 1650°C in a muffle type (Carbolite HTF-18/8) furnace at heating and

cooling rates of 5°C/min.

Density of sintered pellets was measured in high purity ethanol through

Archimedes principle by using high accuracy weighing balance. Morphology of

whisker and surface of the sintered/fractured pellets was observed by using scanning

electron microscope (JEOL JSM-6490LA) at an accelerating voltage of 2-20KV. The

average grain size was measured through linear intercept method. Chemical

composition and phase purity were determined through Energy Dispersive X-Ray

Spectroscopy (EDS or EDX) attached with SEM.

Vicker hardness of sintered pellets was measured with the help of

microhardness tester (Series 401MVD WOLPERT). For each composition, four

samples were used while approximately 10-15 indents were made for each

measurement and average hardness was calculated. Separation between neighboring

indents was kept more than 4 diagonal lengths of indented impression as per ASTM C

1327-99 standard [5]. For vicker micro-hardness test, a load of 1Kg was applied for 15

seconds and dimension of formed indent was observed by using installed optical

microscope.

101

7.2 Results and Discussion

7.2.1 Microstructure Analysis

Figure 7.1(a-c) shows SEM micrographs of the surface of alumina whisker

reinforced TZ-3Y composite samples sintered at three different temperatures (1400,

1500, 1650°C). Average grain sizes calculated from these micrographs through linear

intercept method are given in Table 7.1. As expected with increase in sintering

temperature the grain size increases.

Table 7.1: Average grain size of samples sintered at different temperatures.

Sintering Temperature (°C) Average grain size (μm)

1400 0.3

1500 0.5

1650 1.5

Unlike the samples sintered at 1400°C composite samples sintered at 1500°C

and 1650°C showed equiaxed grains with well-developed grain boundaries and

minimum residual porosity. Reason for grain growth was the high sintering

temperature, which provided activation energy required for grain boundary migration

through enhanced high temperature diffusion phenomenon [6].

102

Fig. 7.1(a-c): SEM micrographs showing size and morphology of the inter-granular pores for 10wt%

Al2O3(w)+90wt%TZ-3Y composite sample surface sintered at temperature of (a) 1400ºC; (b) 1500ºC; (c)

1650ºC.

(a)

(b)

(c)

103

7.2.2 Effect of Sintering Temperature and Whisker

Concentration on Density

Figure 7.2 (a-b) shows the effect of alumina whiskers concentration (5, 10, 15

and 20wt%) and sintering temperature on relative density of the composite. Three main

observations can be inferred from these graphs. First, with increase of whisker content

the sintered density decreases. It is because at higher whiskers concentration whiskers

form larger agglomerates, which resist complete shrinkage of the composite during

sintering and cause porosity and consequent decrease in sintered density. However,

decrease in density with increase in alumina whiskers content is not as significant at

higher temperatures as it is at lower temperatures.

Secondly, it is observed that with the addition of deflocculating agent (i.e., 1

wt% CTAB), density improves for all compositions of the composite due to better

dispersion of alumina whiskers in matrix material. It is because a poor dispersion and

consequent formation of larger agglomerates in the absence of CTAB results in leaving

larger voids during sintering due to difference in shrinkage rate of agglomerate and

matrix [7].

A third observation is about increase in density with an increase in sintering

temperature, for the same whisker concentration. It can be correlated with micrographs

of sintered surface of the composite as shown in Fig. 7.1 (a-c); which shows that due to

decreased diffusion at lower temperatures, comparatively larger pores (encircled in

Fig. 7.1 (a) ) are left in the sintered samples resulting in poor density. Moreover,

numbers of triple junction pores (indicated by arrows in Fig. 7.1 (b, c)) are more in the

samples sintered at higher temperature (i.e. 1650°C). Reason is, the quadruple junction

pores are comparatively unstable configuration [8] and with an increase in sintering

temperature grains grow, decreasing the size of pores and transforming them into triple

junction pores, which are much smaller in size and thus the density of composite

increases.

104

Fig. 7.2(a-b): Relative densities of sintered composite samples having varying amount of alumina

whiskers; (a) without CTAB; (b) with CTAB.

(a)

(b)

105

7.2.3 Effect of Sintering Temperature on Whisker Morphology

Figure 7.3(a-f) shows SEM micrographs and EDS spectra on surface of a sample

sintered at 1650°C. Two types of grains (dark and bright) are clearly visible in these

SEM image of sample with 15 wt% Al2O3(w) (Fig. 7.3(b)). The number of dark grains

is, however, much less in case of sample with 5 wt% Al2O3(w) (Fig. 7.3(a)). The EDS

on an area of 8 m × 8 m (Fig. 7.3(e)) shows peaks from both aluminum (Al) and

zirconium (Zr). The intensity of peak from Zr is higher than that from Al showing its

composition to be near about 15%Al2O3. The EDS on individual dark and bright grains

is shown in Fig. 7.3(c-f). In case of bright grain, Zr peak is high and Al peak is

diminished (Fig. 7.3(d)), showing it to be zirconia rich. Just an opposite case is

observed for a fully dark grain (Fig. 7.3(c)) showing it to be alumina rich. However,

EDS on a relatively less dark grain (Fig. 7.3(f)) shows that it contains both Al and Zr,

while the amount of Al is higher than Zr. The appearance of alumina and Zirconia in

form of dark and bright grains (characterized through EDS) is typically found in

zironia reinforced with alumina particulates [9].

The possible explanation for formation of aluminum-rich and lean grains may

be solid-state diffusion process during high temperature sintering. It is known that at

lower temperatures alumina has lower solubility in zirconia and solubility increases

with an increase in temperature [10-11]. Jerebtsov et al. [12] estimated that the

maximum solubility of alumina in tetragonal zirconia (at eutectic temperature of

1866ºC) is about 8±2 wt%. Therefore, aluminum lean grains consist of solid solution

of alumina in zirconia are formed. The left behind alumina (whiskers) then transforms

to alumina rich grains due to instability of highly mesoporous structure of alumina

whiskers at higher temperatures. Since optimum sintering temperature of zirconia is

lower than alumina [13-14], zirconia rich grains may form easily at lower temperatures

or at earlier stages during sintering at higher temperatures. While mesoporous alumina

structure retains its morphology in case of sintering at lower temperatures and loses it

in case of sintering at higher temperatures.

106

Fig. 7.3(a-f): SEM micrographs and EDS of sample surface sintered at 1650ºC; (a) containing 5wt%

Al2O3(w) and (b) 15wt% Al2O3(w); (c) Spot EDS of dark grain; (d) Spot EDS of the white grain; (e) EDS

spectra on aprox. 8µm×8µm area (f) Spot EDS of less dark grain.

In order to confirm retention of alumina whiskers morphology at lower temperatures,

the fractured surfaces of the samples were observed by using SEM. It was revealed that

by sintering at 1500°C, alumina retained its whisker morphology (encircled in

Fig.7.4(a-b)). However, diameter of whiskers in samples sintered at 1500°C was higher

than that in samples sintered at 1400°C due to high temperature spheroidization of the

whiskers [15]. But for samples sintered at 1650°C, instead of whiskers, typical grain-

like morphology was observed (Fig. 7.4(c-d)) due to accelerated diffusion process

among the neighboring whiskers having mesoporous nature [16] and larger surface

area [17].

Hence, change in morphology of the whiskers at higher sintering temperature

is responsible for decrease in density with drop in Al2O3(w) concentration (Fig. 7.2(a-

b)).

107

Fig. 7.4(a-d): SEM micrographs of fractured composite samples containing 10wt% Al2O3(w) and sintered

at (a) 1400ºC;(b) 1500ºC; (c) 1650ºC (low magnification image) (d) 1650ºC (high magnification image).

7.2.4 Effect of Sintering Temperature and Whisker

Concentration on Hardness

Figure 7.5(a-b) shows the effect of sintering temperature on hardness of the

composites. For all sintering temperatures, hardness increases with increasing the

alumina whisker content, reaching up to a maximum at 10 wt% Al2O3(w) content and

then decreases. These maximum hardness values are 13.37, 14.47 and 14.10 GPa for

samples sintered at 1400, 1500 and 1650°C, respectively. A further improvement in

hardness to 14, 15.11 and 14.76 GPa, respectively, was obtained through addition of

deflocculating agent (i.e. 1 wt% CTAB). Among the samples having 10 wt% Al2O3(w)

sintered at different temperatures, a maximum hardness was obtained for a sintering

temperature of 1500°C. This value of hardness is higher than already reported

maximum value of 14.4 GPa for same kind of material sintered at 1450°C [4].

108

Fig. 7.5(a-b): Vicker hardness of sintered pellets having varying amount of alumina whiskers; (a)

without CTAB; (b) with CTAB.

(a)

(b)

109

The maximum hardness at 1500°C is associated with combined effect of three

factors: 1) retention of whiskers morphology (which is absent in case of sintering at

1650oC), 2) formation of highly dense structure (which lacks in case of sintering at

1400oC) and 3) comparatively fine grain structure (Hall-Petch relationship) in

comparison with that at 1650oC. Whisker morphology deflects the propagating crack

and in this way causes hindrance in its propagation by absorbing crack propagation

energy. This can be seen from crack propagation path in Fig. 7.6(a-c) emanating from

a corner of a Vickers hardness indent, which shows a zigzag shape due to crack

deflection. Consequently, hardness of the composite increases [18]. The alumina

whisker acts as a ligamentary bridge between two separating surfaces of the crack and

offered hindrance to crack propagation [19]. This hindrance result in improving the

hardness and strength of the composite.

A drop in hardness with increase in alumina whisker content above 10 wt%

may be explained on the basis of comparatively sharp decrease in density of the

composites above 10 wt% of alumina whisker content (Fig. 7.2(a-b)). At lower

concentration (i.e. less than 10 wt% whisker) disperse well in matrix and cause crack

deflection effect. However, at higher concentrations, agglomeration effect

predominates resulting in fall in hardness [20-21].

It is observed that by increasing the sintering temperature to 1650°C, hardness

of the composite decreases to lie in between that at 1400°C and 1500°C (Fig. 7.5(a-b)).

This is true up to 15 wt% Al2O3(w) content. However, at a concentration of 20 wt%

alumina whisker, the maximum hardness is obtained at 1650°C. It is because at this

concentration of alumina whisker, dispersion is not so good and whisker morphology

causes poor density. On the other hand, change in morphology to rounded grains gives

better density and thus improvement in hardness [22].

110

Fig. 7.6(a-c): SEM micrographs of Vicker indentation showing the crack propagation path at one arm of

the indent impression for samples with 10wt% Al2O3(w) and sintered at (a) 1400ºC; (b) 1500ºC; (c)

1650ºC.

Additionally, for all the composite samples containing different concentrations

of alumina whiskers, addition of CTAB (1 wt%) resulted in an increase in hardness of

the composite due to better dispersion of whiskers [4]. Addition of CTAB increases the

zeta potential of the alumina particles / whiskers [23], because of which the repulsive

force between them increases resulting in improvement in dispersion and consequent

enhancement of the mechanical properties.

111

7.3 Conclusions

Effect of sintering temperature on density and hardness of the alumina whisker

reinforced composites has been studied. Based on the present work, optimum-

sintering temperature was concluded to be 1500°C as maximum hardness was

obtained at this temperature. An optimum combined effect of whisker-like

morphology, smaller grain size and higher density seems to be responsible for higher

hardness in this case. Composites sintered at 1400°C had lower density and higher

retained porosity, while samples sintered at 1650°C showed excessive grain growth

and change in morphology of the whiskers to rounded grains. These factors resulted in

comparatively lower hardness of composites sintered at 1400°C and 1650°C.

Moreover, addition of 1 wt% CTAB increased density and hardness of all the

composite samples because of better dispersion of alumina whiskers.

112

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incorporated ZrO2-20 wt% Al2O3 by Spark Plasma Sintering and Characterization. J

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[2] Den Exter P, Winnubst AJA, Burggraaf AJ. The preparation and characterization

of Y-TZP/20 wt% alumina. J Eur Ceram Soc. 1993;11(6):497-507.

[3] Li J, Li W, Nai X, Bian S, Liu X, Wei M. Synthesis and formation of alumina

whiskers from hydrothermal solution. J Mater Sci. 2010;45(1):177-81.

[4] Abdullah M, Ahmad J, Mehmood M, Mujtaba ul H, Maekawa H. Synthesis of

Al2O3 whisker-reinforced yttria-stabilized-zirconia (YSZ) nanocomposites through in-

situ formation of alumina whiskers. Ceram Int. 2011;In Press, Accepted Manuscript.

[5] ASTM.org. Standard test method for vickers indentation hardness of advanced

ceramics Philadelphia, PA,: West Conshohocken, PA; 1991.

[6] Su H, Tang X, Zhang H, Zhong Z, Shen J. Sintering dense NiZn ferrite by two-

step sintering process. J Appl Phys. 2011;109(7):07A501-07A-3.

[7] Vishista K, Gnanam FD. Effect of deflocculants on the densification and

mechanical properties of sol-gel alumina. Mater Lett. 2004;58(11):1665-70.

[8] Tikare V, Braginsky M, Olevsky E, Johnson DL. Numerical Simulation of

Anisotropic Shrinkage in a 2D Compact of Elongated Particles. J Am Ceram Soc.

2005;88(1):59-65.

[9] N. P. Bansal, Choi SR. Processing of Alumina-Toughened Zirconia Composites.

Cleveland, Ohio: National Aeronautics and Space Administration, Glenn Research

Center; 2003.

[10] Guo X, Tang C-Q, Yuan R-Z. Grain boundary ionic conduction in zirconia-based

solid electrolyte with alumina addition. J Eur Ceram Soc. 1995;15(1):25-32.

[11] Bernard H. Commissariat a` l' Energie,. Atomique, CEN-Saclay, France. 1981. p.

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[12] Jerebtsov DA, Mikhailov GG, Sverdina SV. Phase diagram of the system:

Al2O3–ZrO2. Ceram Int. 2000;26(8):821-3.

[13] Sun Y-h, Zhang Y-f, Guo J-k. Microstructure and bending strength of 3Y-TZP

ceramics by liquid-phase sintering with CAS addition. Ceram Int. 2003;29(2):229-32.

[14] Liu F, Zhi WC, Jin YN. Microstructure and Luminance Properties of Integrated

Translucent Alumina Tube Used in Metal Halide Lamps. Adv Mat Res.

2011;177:272-6.

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[15] Krishnarao RV, Mahajan YR. Formation of SiC whiskers from raw rice husks in

argon atmosphere. Ceram Int. 1996;22(5):353-8.

[16] Zhu Z, Sun H, Liu H, Yang D. PEG-directed hydrothermal synthesis of alumina

nanorods with mesoporous structure via AACH nanorod precursors. J Mater Sci.

2010;45(1):46-50.

[17] Bandopadhyay AK. Nano Materials: New Age International (P) Ltd.; 2007.

[18] Liu H, Weisskopf K-L, Petzow G. Crack Deflection Process for Hot-Pressed

Whisker-Reinforced Ceramic Composites. J Am Ceram Soc. 1989;72(4):559-63.

[19] De Muynck W, De Belie N, Verstraete W. Effectiveness of admixtures, surface

treatments and antimicrobial compounds against biogenic sulfuric acid corrosion of

concrete. Cement Concrete Comp. 2009;31(3):163-70.

[20] Krylov AV, Barinov SM, Ivanov DA, Mindlina NA, Parilák L, Dusza J, et al.

Influence of SiC whisker size on mechanical properties of reinforced alumina. J

Matter Sci Lett. 1993;12(12):904-6.

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1996;22(3):199-205.

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114

Chapter 8

8 Low Temperature Tetragonal Phase Stability of

TZ-3Y Composite Reinforced with Alumina

Whiskers

Due to its excellent strength, toughness and hardness properties, 3 mol% yttria-

stabilized tetragonal zirconia (TZ-3Y) is a well known ceramic material. Unlike other

ceramic materials that have lower fracture toughness, TZ-3Y, due to its transformation

toughening mechanism, is a unique system for many load bearing applications [1].

Since its discovery, TZ-3Y has seen a tremendous growth in its applications, with

thermal barrier coatings, wear resistant applications and biomaterial applications, to

name a few. Zirconia, particularly stabilized tetragonal zirconia, is being used for

framework of fixed partial dentures and synthetic dental implants because of its

excellent mechanical properties and superior esthetics and bio-inertness [2].

However, along with its excellent properties and varied applications, stabilized

tetragonal zirconia (TZ-3Y) also has a major drawback. It undergoes tetragonal to

monoclinic phase transformation in presence of water or water vapor in a temperature

range between room temperature and approximately 400ºC depending upon the used

stabilizer, its concentration as well as grain size of the ceramic. This monoclinic phase

has 4% higher volume than the tetragonal phase, which may result in degradation or

catastrophic failure of the composite part. Micro-cracking and then loss of strength can

result from this phase transformation. This is a concern for biomaterial application

such as in a humid atmosphere of body fluids at body temperature over a long period

of time or when it is autoclaved for sterilization. Kobayashi et al. [3] were the first to

report on the aging transformation of zirconia and its deleterious effect on components

especially in biomedical industry. Subsequently, Haraguchi et al. [4] reported the

deleterious effect of transformation on formal head implants during their use in total

hip arthroplasty. In response to this, in 2001, the US Food and Drugs Administration

(FDA) warned against use of stabilized zirconia formal heads and recalled all

manufactured products by the Saint Gobain, a French company [5]. Because of its

importance, this area attracted due attention for scientific research [6-7].

115

Where addition of alumina to TZ-3Y to form composites is useful to improve

the mechanical properties of zirconia [8], it is also useful to enhance thermal stability

of the composites. It has been proposed that addition of alumina helps to suppress

propagation of monoclinic phase transformation in the bulk and results in improving

hydrothermal phase stability [9-10]. Ban et al. in their study of the effect of alumina on

hydrothermal phase stability observed that fracture toughness of alumina added

zirconia could be retained after autoclaving [11].

In the present study, alumina has been added to the composites in the form of whiskers

in order to improve their mechanical properties. Due to non-spherical morphology of

the added alumina whiskers, it is expected that they would resist shrinkage during

sintering and may result in enhancement of porosity in the sample [12]. Since porosity

also plays an important role in deteriorating the hydrothermal phase stability of

zirconia and zirconia composites, the study of the effect of sintering conditions on

shrinkage and consequently on hydrothermal phase stability for alumina whisker

reinforced composites will be important. For particulate alumina reinforced TZ-3Y

composites, the sintering behavior may be found in the book by Basu et al. [13].

The present chapter presents our results on hydrothermal phase stability as

well as on dilatometric behavior of alumina whisker reinforced TZ-3Y composites in

order to investigate proper compositions, which show best properties for biomedical

applications.

8.1 Experimental Procedures

The alumina whiskers were prepared through in situ transformation of ammonium

aluminum carbonate hydroxide (AACH) whiskers during sintering. Method for

reparation of AACH whiskers has been explained in chapter 4.

The AACH whiskers were mixed with 3 mol% Y2O3-stabilized ZrO2

(abbreviated as TZ-3Y) powder in analytical grade ethanol containing 1 wt% CTAB

(for improving the whisker dispersion) and mixed using magnetic stirrer on electric hot

plate for 12 hrs at temperature of 50ºC so that all the ethanol was evaporated. Samples

with varying amounts of AACH corresponding to 0, 10 and 20 wt % Al2O3 were

prepared. The dried mixture was calcined at 650ºC for removal of volatile ingredients

of AACH and CTAB. It was then dry pressed at a load of 500MPa and sintered at three

temperatures i.e. 1400, 1500 and 1650ºC for 2 hrs at a heating rate of 5ºC/min in a

Lindberg/Blue furnace followed by cooling inside the furnace.

116

The sintered samples were transferred into a 50 ml capacity Teflon-lined

autoclave (Parr bomb), filled with distilled water up to 75% of the total volume, sealed

and hydrothermally treated at 135ºC for 10 hr. This hydrothermal time was selected as

per ASTM standard F2345-03 [14], which describes that 10 hr hydrothermal treatment

is equivalent to 20 years in-vivo aging in human body environment [15].

The phase analysis of hydrothermally treated samples was performed by XRD

in θ-2θ mode for a 2θ range of 24°-34° with scan rate of 0.5°/min. For XRD

measurements, Bruker D8 Discover X-ray diffractometer with Cu-K radiation was

used. The volume fraction of monoclinic phase, Vm, in these samples was determined

by using equation (8.1) as per the Toraya’s method [16]. The integrated intensity ratio

Xm was calculated according to equation (8.2) below.

m

mm

X

XV

311.11

311.1

(8.1)

)101()111()111(

)111()111(

tmm

mmm

III

IIX

(8.2)

The subscripts m and t are for the monoclinic and tetragonal phases, respectively. For

calculating the integrated intensity ratio, the intensity was determined through the

Jade-5 XRD Pattern Processing software. The intensities of tetragonal and monoclinic

phases of zirconia were found from the peaks at the 2 reflections given in Table 8.1

[16].

Moreover, the thermal shrinkage due to sintering was determined with the help

of a high temperature dilatometer (Orton Dilatometer Model 1600D). The study was

performed in air from 25 to 1600ºC at a heating rate of 3ºC/min.

For microscopic analysis, the high-resolution field-emission (FEG) scanning

electron microscope (SEM) and CM-200 transmission electron microscope (TEM)

were used. The relative density of samples was measured by Archimedes method by

using high purity ethanol.

117

Table 8.1: Reflections as used in Toraya’s method

Reflection Phase 2-Theta Value

)111(

Monoclinic 28.2º

(111) Monoclinic 31.5º

(101) Tetragonal 30.2º

8.2 Results and Discussion

Figure 8.1 (a) shows typical XRD patterns of the composite samples sintered at

three different temperatures (1400, 1500 and 1650ºC) with subsequent hydrothermal

treatment at 135ºC for 10h. For reference, XRD patterns of typical as-sintered samples

have been shown in Fig. 8.1 (b). (Detailed structural study of as-sintered samples has

been shown in previous chapters.) The samples contained varying amounts of Al2O3

(whiskers) in matrix of TZ-3Y. Before hydrothermal aging, no XRD peaks could be

assigned to monoclinic phase; while the peaks of monoclinic phase appeared after

hydrothermal aging as can be seen in Fig. 8.1 (a).

118

Fig. 8.1(a-b): XRD patterns of composite samples containing pure TZ-3Y and 20wt% Al2O3(w) (a)

sintered at 1400, 1500 and 1650ºC after hydrothermal aging at135ºC for 10 hr; (b) sintered at 1400ºC

and without hydrothermal aging.

(a)

(b)

119

8.2.1 Effect of Sintering Temperature on LTD

The volume percent of the produced monoclinic phase (Vm) due to hydrothermal

aging was quantified following the procedure mentioned in experimental section above

and plotted in Fig. 8.2 against sintering temperature. The TZ-3Y samples with and

without alumina whiskers showed different behavior, as discussed below:

For the monolithic TZ-3Y, volume fraction of monoclinic zircon (Vm)

increased with increase in sintering temperature. This is in agreement with earlier

reports [15, 17]. The effect of grain size on the low temperature degradation is

complicated. However, it is confirmed that the decrease in grain size has beneficial

effect on the stability of tetragonal phase of zirconia [18-19]. There are two proposed

reasons for this positive effect of grain size on tetragonal phase of zirconia. The first

one is by Hallmann et al. [15], they has mentioned that during hydrothermal treatment

at lower temperature local stresses arises from the thermal mismatch between the

adjacent grains due to slight anisotropy of the coefficient of thermal expansion along

different crystallographic axes in zirconia. They suggested that for the small equiaxed

grains this stress distribution is uniform. However, the large grains produced due to

sintering at higher temperature show inhomogeneity in the distribution of this stress.

During hydrothermal aging, this inhomogeneous stress results in enhancing the LTD.

The other approach is related with tendency of yttria to diffuse preferably towards

grain boundaries. The diffusion of yttria and heterogeneity in its distribution is

appreciable at temperatures ≥1500ºC [20]. At lower sintering temperature, yttria does

not diffuse appreciably to produce any heterogeneity, as result of which the samples

exhibit improved stability against LTD.

For TZ-3Y+Al2O3(w) composites, Fig. 8.2 shows that monoclinic phase

fraction, Vm, decreases with an increase in sintering temperature, which is just an

opposite behavior to that of monolithic TZ-3Y. At lower sintering temperatures

(1400ºC), monoclinic phase fraction is even higher than that of monolithic TZ-3Y. By

increasing the sintering temperature, the monoclinic phase fraction decreases becoming

lower than that of monolithic TZ-3Y at 1650oC. The phenomenon behind this behavior

may be associated with change in relative density with rise in the sintering

temperature. To investigate it, dilatometric and relative density analyses of the sintered

composites were performed.

120

Fig. 8.2: Amount of monoclinic phase produced due to hydrothermal treatment of samples containing

0, 10 & 20 wt%Al2O3(w) in the matrix of TZ-3Y.

8.2.2.1 Dialtometry of TZ-3Y+ Al2O3(W) Composites

Figure 8.3 shows the dilatometric curve of TZ-3Y samples reinforced with 10

and 20wt% alumina whiskers up to 1600ºC. In both the samples, the densification

starts (i.e. 0.5% of length contraction) above 970ºC. The shrinkage rate remains

dramatically high from 1050ºC to 1250ºC and is attributed to the shrinkage of TZ-3Y

[21]. Above 1250ºC the shrinkage rate becomes slow and ultimately this gives a total

linear shrinkage of 25% (for 20 wt% Al2O3(w)+TZ-3Y sample) at 1600ºC.

121

For TZ-3Y composites reinforced with 20 wt% Al2O3 particulates, it has been

reported that the total dilatometric densification at 1600ºC is 21% [22]. The difference

in total densification may be associated with the difference in reinforcement

morphology. The whisker morphology is expected to resist densification of the

composite. The SEM images in Fig. 8.4(a & b) show the AACH whiskers prepared by

hydrothermal synthesis. The whiskers have long aspect ratio. However, the TEM

images (Fig. 8.4(c, d & e)) show that the sides of the whiskers are porous. A typical

diameter of a whisker is approximately 100 nm, while the pore diameter varies from 2

to 20 nm. These pores were produced at very start of the TEM imaging due to

exposure of whiskers to high energy electrons. Higher densification is expected when

such a porous structure will collapse at higher sintering temperatures.

Fig. 8.3: Dilatometer curves of (10 and 20wt%) Al2O3(w) containing TZ-3Y samples.

122

In chapter 7 on study of sintering behavior of TZ-3Y+Al2O3(w) composites, it

has been shown through fractured surface micrography that in these composites the

whisker-like elongated morphology is retained at lower sintering temperatures (i.e., at

1400ºC), which means low density regions are present in this case. Such low density

regions are responsible for enhanced fraction of monoclinic phase after hydrothermal

treatment. However, in case of higher sintering temperatures (1650ºC), zirconia having

higher diffusion rates forms zirconia rich grains leaving behind porous alumina

structure which later collapse to from alumina rich grains. The loss of elongated porous

whiskers results in enhanced total densification (i.e., 25%) at 1600ºC (Fig. 8.3). The

formation of alumina rich and lean grains may also be confirmed from the

backscattered SEM image in Fig. 8.5 that shows topographic grain contrast on the

surface of a sample containing 20 wt% Al2O3(w). The zirconia rich grains appear bright

due to higher atomic number of zirconium than aluminum resulting in stronger

backscattering effect. The alumina concentration in a grain can roughly be assessed by

the visual darkness of the grains.

Fig. 8.4(a-e): Micrographs of the hydrothermally synthesized AACH whiskers (a&b) SEM micrographs at (a) lower

and (b) higher magnifications; (c-e) TEM micrographs at (c) lower and (d&e) high magnification.

123

Fig. 8.5: Backscattered image of hydrothermally aged TZ-3Y+20wt% Al2O3(w) composite sample

sintered at 1650ºC.

8.2.2.2 Relative Density

The increase in densification with increase in the sintering temperature can also

be verified from the relative density change with sintering temperature (Fig. 8.6). From

this Fig, 8.6, it is observed that at the same sintering temperature by increasing the

whisker concentration the relative density decreases due to the presence of elongated

porous alumina whiskers which resist densification. However, at higher sintering

temperature (i.e., 1650ºC) the difference in density for different whisker contents

decreases owing to the loss of elongated morphology of the whiskers. For the same

composition, the relative density increases with an increase in sintering temperature,

which is due to enhanced diffusion as well as the collapse of whisker morphology.

124

Fig. 8.6: Density of the Al2O3 whisker reinforced TZ-3Y composite samples sintered at different

temperatures.

Therefore, one reason for an improvement in LTD with increase in sintering

temperature is the increase in density. At lower sintering temperature, the density is

low and the water vapors go inside the pores of composite increasing the monoclinic

phase fraction. Such pores can also be seen in the SEM micrographs of the samples

sintered at 1400ºC as shown in Fig. 8.7(a). However, at higher sintering temperatures,

the density is sufficiently improved and the porosity in the microstructure vanishes

(Fig. 8.7(b)). The exposure area to water vapors decreases resulting in decreased

monoclinic phase formation during hydrothermal aging. Another reason for

improvement of LTD in alumina-reinforced samples at higher sintering temperature

may be the formation of alumina-zirconia solid solution.

125

Fig. 8.7(a-b): SEM images of the surface of TZ-3Y+20wr% Al2O3(w) composite sample sintered at (a)

1400ºC; (b)1650ºC.

8.2.2.3 Mechanisms of LTD and Role of Alumina

Although there is no consensus on the mechanism of LTD in TZ-3Y, two

mechanisms are generally proposed. According to Sato et al. [23], the reaction

between ZrO2 and water at low temperature is responsible for the formation of

monoclinic phase. According to them, the water makes way into the flaws / pores and

forms Zr-OH bonds releasing the strain responsible for the stability of the tetragonal

phase. It results in tetragonal to monoclinic phase transformation. On the other hand,

Lange et al. [18] suggest that water vapors at low temperature react with Y2O3 on the

surface of zirconia or inside the pores and form Y(OH)3. This results in depletion of

stabilizing agent in ZrO2 grains and promotes the formation of monoclinic phase.

It is known that alumina has low solubility in zirconia that increases with an

increase in temperature [24-25]. Jerebtsov et al. [26] estimated the maximum solubility

of alumina in tetragonal zirconia (at eutectic temperature of 1866ºC) to be about 8±2

wt%. Like yttrium, aluminum also has trivalent ionic state and it forms substitutional

solid solution in zirconia. In zirconia, alumina produces the same number of oxygen

vacancies as that of yttira. These oxygen vacancies are responsible for stabilizing the

metastable t-phase of ZrO2. The chemistry of this substitution of alumina is similar to

that of yttria in zirconia as illustrated in following chemical reactions (8.3 & 8.4) using

Kröger–Vink notation.

Y2O3 2Y′Zr + Vo

•• + 3O

xo (8.3)

Al2O3 2Al′Zr + Vo

••+ 3O

xo (8.4)

(a) (b)

126

Consequently, enhanced dissolution of substantial alumina in the tetragonal

zirconia solid solution at higher temperatures may also be a dominant factor to

stabilize tetragonal phase and decrease LTD during hydrothermal ageing.

It may be noticed from Fig. 8.2 that hydrothermal phase stability of the sample

with 10 wt% Al2O3(w) is much better than the sample with 20 wt% Al2O3(w) . In

previous chapters, it has been shown that the sample with 10 wt% Al2O3(w) shows the

highest hardness and toughness values. While the density of this composite is also

better than that of 20 wt% Al2O3(w) as shown in Fig. 8.6. Therefore, it may be

concluded with greater confidence that this composite (TZ-3Y) with 10 wt% Al2O3(w)

is quite suitable for biomedical applications. This is due to reasonable density, greater

hardness, comparatively lower porosity, reinforcement effects as well as enhanced

stability of tetragonal zirconia related with partial dissolution of alumina into

tetragonal zirconia in the form of solid solution. It can also be suggested here that

alumina is less prone to precipitate to form its hydroxide, in comparison with yttria.

For the 20 wt% Al2O3(w) containing sample, increase in porosity allows a

greater water ingression and thus a faster degradation of tetragonal zirconia phase.

8.3 Conclusions

In this chapter, the phase stability of TZ-3Y with and without alumina whisker

reinforcement under hydrothermal conditions was studied. The hydrothermal

treatment at 135oC resulted in the formation of monoclinic phase. The higher alumina

whisker contents and lower sintering temperatures enhanced the destabilization of

tetragonal phase. It was attributed to lowering of density of the composites with

increase in whisker content and decrease in sintering temperature. Higher porosity in

composites provides more surface area for water to react with matrix phase and cause

degradation of the stabilized phase. Higher sintering temperatures enhanced the

stability of tetragonal phase, which could be associated with the formation of

spherical alumina particles from whiskers and resultant high density composites. The

composite with 10wt% Al2O3 whiskers was better as it showed greater tetragonal

phase stability.

127

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[4] Haraguchi K, Sugano N, Nishii T, Miki H, Oka K, Yoshikawa H. Phase

transformation of a zirconia ceramic head after total hip arthroplasty. J Bone Joint

Surg B. 2001;83B(7):996-1000.

[5] Administration UFaD. St. Gobain Desmarquest Zirconia Ceramic Femoral Head.

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[6] Chevalier J. What future for zirconia as a biomaterial? Biomaterials.

2006;27(4):535-43.

[7] Lughi V, Sergo V. Low temperature degradation -aging- of zirconia: A critical

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[8] Nevarez-Rascon A, Aguilar-Elguezabal A, Orrantia E, Bocanegra-Bernal MH.

Al2O3(w)-Al2O3(n)-ZrO2 (TZ-3Y)n multi-scale nanocomposite: An alternative for

different dental applications? Acta Biomater. 2010;6(2):563-70.

[9] Schneider J, Begand S, Kriegel R, Kaps C, Glien W, Oberbach T. Low-

Temperature Aging Behavior of Alumina-Toughened Zirconia. J Am Ceram Soc.

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[10] Kawai Y, Uo M, Wang Y, Kono S, Ohnuki S, Watari F. Phase transformation of

zirconia ceramics by hydrothermal degradation. Dent Mater J. 2011;30(3):286-92.

[11] Ban S, Suehiro Y, Nakanishi H, Nawa M. Fracture toughness of dental zirconia

before and after autoclaving. J Ceram Soc Jpn. 2010;118(1378):406-9.

[12] Tiegs TN, Becher PF. Sintered Al2O3-SiC-Whisker Composites. Am Ceram Soc

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[13] Basu B, Katti DS, Kumar A. Advanced Biomaterials: Fundamentals, Processing,

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[14] Chang JD, Billau K. Bioceramics and alternative bearings in joint arthroplasty:

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proceedings: Steinkopff; 2007.

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[15] Hallmann L, Mehl A, Ulmer P, Reusser E, Stadler J, Zenobi R, et al. The

influence of grain size on low-temperature degradation of dental zirconia. J Biomed

Mater Res B Appl Biomater. 2012;100B(2):447-56.

[16] Toraya H, Yoshimura M, Somiya S. Calibration Curve for Quantitative Analysis

of the Monoclinic-Tetragonal ZrO2 System by X-Ray Diffraction. J Am Ceram Soc.

1984;67(6):C-119-C-21.

[17] Campbell FC. Elements of metallurgy and engineering alloys: ASM

International; 2008.

[18] Lange FF, Dunlop GL, Davis BI. Degradation During Aging of Transformation-

Toughened ZrO2-Y2O3 Materials at 250°C. J Am Ceram Soc. 1986;69(3):237-40.

[19] Garvie RC. The Occurrence of Metastable Tetragonal Zirconia as a Crystallite

Size Effect. J Phys Chem. 1965;69(4):1238-43.

[20] Matsui K, Ohmichi N, Ohgai M, Yoshida H, Ikuhara Y. Effect of alumina-

doping on grain boundary segregation-induced phase transformation in yttria-

stabilized tetragonal zirconia polycrystal. J Mater Res. 2006;21(9).

[21] Mazaheri M, Simchi A, Dourandish M, Golestani-Fard F. Master sintering

curves of a nanoscale 3Y-TZP powder compacts. Ceram Int. 2009;35(2):547-54.

[22] Moritz T, Eiselt W, Moritz K. Electrophoretic deposition applied to ceramic

dental crowns and bridges. J Mater Sci. 2006;41(24):8123-9.

[23] Sato T, Shimada M. Transformation of Yttria-Doped Tetragonal ZrO2

Polycrystals by Annealing in Water. J Am Ceram Soc. 1985;68(6):356-.

[24] Guo X, Tang C-Q, Yuan R-Z. Grain boundary ionic conduction in zirconia-based

solid electrolyte with alumina addition. J Eur Ceram Soc. 1995;15(1):25-32.

[25] Bernard H. Commissariat a` l' Energie,. Atomique, CEN-Saclay, France. 1981. p.

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Al2O3–ZrO2. Ceram Int. 2000;26(8):821-3.

129

Chapter 9

9 Effect of HAp Addition on Biocompatibility of

the Composite

Zirconia is popular due to its better mechanical properties and bio-inertness [1].

Transformation toughening mechanism [2] of 3mol% yttria stabilized tetragonal

zirconia (TZ-3Y) also makes it ideal for load bearing applications. Based on our

results presented in earlier chapters, it has been identified that 10wt% addition of

Al2O3(w) in TZ-3Y significantly improves its mechanical properties. In addition, it also

slows down low temperature hydrothermal degradation of TZ-3Y for an increased

service life [3].

Hydroxyapatite (HAp; Ca10(PO4)6(OH)) is famous for its unique resemblance

with bone & dental material. It can make bond with bone very easily, which is the

requirement for osseointegration of implants [4-5] as well as for its use to repair the

damaged bones and teeth. However, its poor mechanical properties limit its

applications in pure form [6].

For obtaining a demanded composite material with optimum strength and

good biocompatibility, it seems interesting to make an addition of HAp in the

Al2O3(w)+TZ-3Y composite. The former may facilitate bioactivity and the latter may

be required for its excellent mechanical properties. Similar combinations have also

been reported by other authors [7-9].

HAp exhibits poor mechanical properties in comparison with tetragonal zircon

and alumina. Therefore, it seems necessary to optimize its content for a certain

required minimum level of mechanical properties along with sufficient

biocompatibility. Therefore, it seems necessary to study the effect of HAp addition on

the mechanical properties and biocompatibility.

The present chapter focuses on the addition of HAp in the TZ-3Y-Al2O3(w)

composites their mechanical properties, cyto-toxicity (and bioactivity). The later was

studied through in-vitro cell adhesion biocompatibility tests.

130

9.1 Experimental Procedure

9.1.1 Preparation of AACH Whiskers

The procedure for preparation of AACH whiskers has been described in Chapter 4.

9.1.2 Preparation of HAp Nanoparticles

Hydroxyapatite nanoparticles were synthesized through solution-precipitation

technique [10]. For this purpose, 23.6g of calcium nitrate was dissolved in 350ml of

distilled water and vigorously mixed for 2h with magnetic stirrer. The solution of

diammonium hydrogen phosphate ((NH4)2HPO4) was formed separately by dissolving

7.92g of diammonium hydrogen phosphate in 250ml of distilled water. The latter

solution was added drop wise into calcium nitrate solution under agitation, which

resulted in precipitation of HAp. During the experiment, the pH of the solution was

maintained at 11 with the help of ammonium hydroxide (NH4OH) addition. After the

precipitation the slurry was continued to be mixed for 12-15h for deceasing the grain

size of the HAp powder as described elsewhere [11]. The precipitates of HAp were

filtered and dried at a temperature of 80°C for 24 hours. After drying, the

agglomerates/lumps of HAp were de-agglomerated by grinding.

9.1.3 Formation of TZ-3Y+Al2O3+HAp Nanocomposites

The starting materials for formation of these composites were the

aforementioned synthesized HAp nanoparticles, AACH whiskers, high-purity α-Al2O3

particulates (Make: Almatis, Germany; purity >99.99%; average particle size ~0.2µm)

and ZrO2 + 3 mol% Y2O3 (thereafter abbreviated as TZ-3Y). The measured amounts

of these particulates or whiskers were separately suspended in ethanol using

ultrasonic bath (Elma E 30 H, 37 kHz) with or without 1wt% CTAB. The three

suspensions, i.e., of TZ-3Y, AACH/Al2O3 and HAp were mixed together, kept at

60°C while stirring until nearly all the ethanol was evaporated.

The dried mixtures with and without CTAB were calcined in air at a

temperature of 650°C and 450ºC respectively for removal of volatile ingredients of

AACH and HAp, residual water vapors and CTAB.

The calcined mixture was slightly ground and pressed uniaxially under a load

of 500MPa in a disk steel die of diameter 10mm. The pressed green pellets were

131

sintered at a temperature of 1400ºC in air for 2hours using high temperature

(Carbolite HTF-18/8) furnace at heating and cooling rates of 2ºC/min. Through XRD,

it was confirmed that added AACH was transformed into α-Al2O3 in sintered

composite samples.

For identification purpose the composite samples containing 90%TZ-

3Y+10%Al2O3(w) were designated as ZA(w), while those with 90%TZ-

3Y+10%Al2O3(p) were represented by ZA(p). For samples with various amounts of

HAp, ZA(w), ZA(p) and CTAB, designations as shown in Table 9.1 were used.

The density of sintered pellets was determined by Archimedes method in

ethanol using high accuracy weighing balance. While for theoretical densities (TD)

measurement the rule of mixtures was employed.

Table 9.1: Sample designations for various compositions

Composition Designation for samples

without CTAB

Designation for samples

with 1wt% CTAB

90% (ZA(w) / ZA(p)) + 10 % HAp 10H-(ZA(w) / ZA(p)) 10H-CT-(ZA(w) /ZA(p))

80% (ZA(w) / ZA(p)) + 20 % HAp 20H-(ZA(w) / ZA(p)) 20H- CT-(ZA(w) /ZA(p))

70% (ZA(w) / ZA(p)) + 30 % HAp 30H-(ZA(w) / ZA(p)) 30H- CT-(ZA(w) /ZA(p))

60% (ZA(w) / ZA(p)) + 40 % HAp 40H-(ZA(w) / ZA(p)) 40H- CT-(ZA(w) / ZA(p))

For flexural strength measurement samples surface was polished [12]. For this

purpose, surface was first ground on silicon carbide papers up to 1200 grit followed

by polishing with 6- and 1-micrometer diamond paste. The samples were then dried at

250°C for 4h to remove water vapors.

The flexural strength of polished samples was determined through ring on ring

test method according to the ISO 6474 standard [13]. For this purpose a universal

testing machine SANS CMT4304 (Shenzhen SANS Testing Machine Co., China)

with cross head speed of 0.1 mm/min, was used. For each composition, measurement

was performed on at least four samples and data was reported as an average.

Vickers hardness on polished surfaces was measured by using a Vickers

microhardness tester (Series 401MVD WOLPERT Group) at room temperature using

a load of 1kg for 15 seconds. Four samples per composition were prepared, and

approximately 10-15 indents per sample were made and the average hardness was

132

calculated. Separation between neighboring indents was kept more than 4-diagonal

lengths of indented impressions as per ASTM standard of ASTM C 1327-99 [14].

For X-ray diffraction analysis high resolution XRD (D8 discover XRD)

machine was employed. Morphological and microstructural investigations were

performed by scanning electron microscope (JEOL, JSM-6490A, Japan).

9.1.4 Cell Adhesion

Cell adhesion is the primary single most important test to assess the cell

response to designed biomaterial. In present study this test was performed on sintered

pellets by using Rhabdomyosarcoma (RD) [15] cell line under scientifically

controlled conditions. First of all the samples were sterilized in an autoclave at 15 psi

(1.034 bar) and 120°C for 40 minutes. Sterilized samples of 10mm diameter were

placed in a 48-well cell culture plate. Fixed amount of cell-containing media (serum

and antibiotic cocktail) was poured in each well with sample using sterilized pipette.

The cell culturing was performed in three stages. First, samples were incubated in an

oven for 72 h at 37°C. Second, after 72hrs of cell culturing, culture media was

aspirated and samples were allowed to dry in the same oven at temperature of 37°C

for further 36hrs. The growth of cultured cell line was then viewed by SEM. In third

part of the experiment, the samples containing 30 & 40 wt% HAp were kept dry in the

same oven for further 48 h to examine the response of fungus and other contaminants

to the cultured cell line. This was to verify that the formed / cultured layer on the

sample surface was really from the cell line.

9.2 Results and Discussion

9.2.1 Surface Morphologies

Figure 9.1(a-e) shows SEM micrographs of sintered surface of the composite

samples with different compositions i.e. pure HAp (Fig. 9.1(a)), pure ZA(w) (Fig.

9.1(b)) and combination of both HAp and ZA(w) (Fig. 9.1(c-e)).

133

Fig. 9.1(a-f): Micrograph of sintered surface of composited samples containing different concentration

of HAp (a) Pure HAp; (b) CT-ZA(w); (c) 10HA-CT-ZA(w); (d) 20HA-CT-ZA(w); (e) 30HA-CT-ZA(w); (f)

40HA-CT-ZA(w).

From these micrographs, it is observed that pure HAp sample (Fig. 9.1(a))

consists of large grains with average size of 8µm, while pure ZA(w) sample (Fig.

9.1(b)) shows fine grain structure with an average size of 250nm. For all other

samples (Fig. 9.1(c-e)) containing different weight fractions of HAp in ZA, coarse

grains of HAp embedded in the matrix of ZA(w) are observed. The reason for higher

grain growth of HAp is its lower sintering temperature than ZA [16]. On the surface

of sintered composite samples, few whiskers like morphologies are also evident.

134

9.2.2 Fracture Mechanism

Figure 9.2(a-h) shows SEM micrographs and EDS of the fractured surface of

composite. The roughness of fractured surface for pure ZA(w) composite (Fig. 9.2(a))

suggests that fracture mode is dominantly inter-granular. While smooth surface of

pure HAp sample (Fig. 9.2(b)) is because of trans-granular mode of fracture. However

in samples containing different concentrations of HAp in matrix of ZA(w) composite

mixed mode of fracture (combination of inter & trans-granular) can be observed in

Fig. 9.2(d-f).

The HAp sintered at lower temperatures (i.e. 1200°C) has been reported to

undergo inter-granular type of failure due to retained porosity at the grain boundary

that seems to facilitate intergranular crack initiation and growth [17]. While for HAp

sintered at a higher temperature of 1400°C exhibits trans-granular mode of fracture

(Fig. 9.2(b), which is also in agreement with reported literature [18].

From areal EDS analysis of fractured surface of pure HAp sample (Fig.

9.2(c)), Ca/P ratio is calculated and it comes out to be 1.62 [19] which confirms the

stability of HAp at sintering temperature of 1400°C.

In the sintered composite samples, whiskers have retained their morphology,

which may, for instance, be seen encircled in Fig. 9.2(f).

It is also important to note from Fig. 9.2(d-f) that the fracture does not appear

to take place at interphase boundary between ZA and HAp. This suggests that

bonding strength between ZA and HAp is reasonable [20]. This strong interfacial

bond provides better stress distribution and transfer of load to strong ZA composite

during loading resulting in enhanced mechanical properties of the composites.

135

0.00 1.00 2.00 3.00 4.00 5.00 6.00 7.00 8.00 9.00 10.00

keV

0

400

800

1200

1600

2000

2400

2800

3200

3600

Co

unts

OK

a

PK

a

CaK

a

CaK

b

Fig. 9.2(a-h): SEM micrographs of fractured surface of composite samples with different compositions (a)

pure ZA(w); (b) 20HA-CT-ZA(w); (c) pure HAp; (d) areal EDS of pure HAp; (e) 30HA-CT-ZA(w); (f) 40HA-

CT-ZA(w).

9.2.3 Densification of Composite

Relative density of sintered composite pellets without and with CTAB is shown

in Fig. 9.3 (a) and (b), respectively. The relative density of pure HAp pellet was 96%.

However, in composite samples, the sintered density decreased with an increase in

HAp content. The relative density was even lower than that of pure HAp samples

prepared under similar conditions.

(a) (b)

(c) (d)

(e) (f)

136

Fig. 9.3(a-b): Relative density of sintered pellets of ZrO2+Al2O3 (whiskers or particulates) having

different amount of HAp (a) without CTAB. (b) with CTAB.

Although it has already been established for TZ-3Y+Al2O3 composites

(chapter 5) that alumina particulate reinforced composites show higher relative

densities than alumina whisker reinforced composites, for HAp containing composites

(a)

(b)

137

the effect of reinforcement morphology was also studied. The results were similar to

those observed for TZ-3Y+Al2O3 composites as shown in Fig. 5.4(a) and (b).

Similarly, the improvement in relative density with the addition of 1 wt%

CTAB was also quite in accordance with our previous observations (chapter 5).

9.2.4 Effect of Whisker Addition on Hardness and Flexural

Strength of Composite

Figure 9.4(a) and (b) show the Vickers hardness of the alumina

particulates/whiskers reinforced composites with and without addition of 1wt%

CTAB respectively. Hardness of pure ZA(p) and ZA(w) composite samples were

13.2GPa & 13.68GPa, respectively. These hardness values were decreased with

addition of HAp. This seems partly due to relatively softer nature of HAp itself, apart

from increase in porosity manifested by decrease in relative density with the addition

of HAp. It may be worth-mentioning that the hardness of these composites remains

higher than that of pure HAp samples in spite of significant fraction of retained

porosity (Fig. 9.3).

It is also observed in Fig. 9.4 that alumina whisker reinforcement gives better

hardness as compared to alumina particulate reinforcement in spite of lower density

(Fig. 9.3), in agreement with TZ-3Y+Al2O3 composites (chapter 5). Reason for this

hardness improvement is the elongated morphology of reinforcement (i.e. alumina

whisker) contributing to strength through whisker toughening effects.

Figure 9.5(a-b) shows flexural strength of composite samples reinforced with

alumina particulates and whiskers. Like hardness, flexural strength also decreases

with increasing HAp content, although it remains higher than that of pure HAp

sample.

The addition of CTAB has been shown to provide with better mechanical

properties because of improved dispersion of alumina (AACH) whiskers in the

particulate mixture of HAp and zirconia, in agreement with Chapter 5.

138

Fig. 9.4(a-b): Vickers hardness of sintered pellets of ZrO2+Al2O3 (whiskers or particulates) having

different amount of HAp (a) with 1wt% CTAB. (b) without CTAB

(a)

(b)

139

Fig. 9.5(a-b): Flexural strength of sintered pellets of ZrO2+Al2O3 (whiskers or particulates) having

different amount of HAp (a) without CTAB. (b) with 1wt% CTAB HAp.

The reason behind decrease in strength with increase in HAp content may be

understood more clearly by looking at SEM images given in Fig. 9.6. The figure

(a)

(b)

140

shows crack propagation path extending from corner of a Vickers hardness indent.

Figure 9.6 (a) shows extensive crack deflection in a pure ZA(w) sample, while the

crack path is intergranular. It shows that grain interior are relatively too strong to

allow transgranular crack propagation. As it has been shown in section 6.3.1.2 of

chapter 6, whiskers result in bridge formation between separating grains and thus

increase the fracture toughness [21]. Whiskers also play role in crack deflection [22].

Due to these crack-whisker interactions, the fractured surface of sample ZA(w) was

rough as shown in Fig. 9.2(a). On the other hand, transgranular crack propagation

through HAp grains may be seen in 30H-CT-ZA(w) sample in Fig. 9.6 (b) due to its

lower strength. Therefore, as the amount of HAp in composites increases, crack finds

more paths through softer phase and fracture toughness and flexural strength decrease.

Fig. 9.6(a-b): SEM images showing the path of crack propagation at corner of Vickers hardness indent on

sintered pellet of (a) pure ZA(w), (b) 30H-CT-ZA(w).

It is important to note that the values of flexural strength for whisker

reinforced HAp containing ZA(w) composites is being reported for the first time.

However, strength of particulate reinforced composite (H-ZA(p)) has already been

reported in literature [9]. The flexural strength in previous studies was measured by

four-point flexural test method. In present work, the strength is determined through

ring on ring method [23]. A comparison between flexural strength reported in

previous studies [9] and that determined in present work for HAp containing ZA(p)

composites reveals that the present value is slightly lower than the previously reported

one. It may be due to difference in the method of measurement. Determination of

flexural strength through ring on ring test method is better representative of the

material [24], as it involves cracks / defects in all different directions. On the other

hand, in four-point method, the cracks parallel to tensile direction of bar shaped

141

specimen do not affect the determination of flexural strength. Also flexural strength of

a ceramic material depends not only on composition (i.e. type of material) but also on

other factors such as defect size and its distribution, specimen size, testing method

and synthesis route etc. [25].

9.2.5 X-Ray Diffraction Analysis

For phase identification of samples, the X-ray diffraction (XRD) analysis was

performed as shown in Fig. 9.7(a-d). Figure 9.7(a) shows XRD pattern of ZA(w)

samples sintered at 1400ºC. Only peaks of tetragonal zirconia and -alumina are

visible.

XRD pattern of pure HAp sintered at 1400oC is shown in Fig. 9.7(b). All the

peaks can be assigned to HAp, without any indication of decomposition products

appear. When ZA(w) is present along with HAp (Fig. 9.7 (a) & (b)), small peaks of

tricalcium phosphate (TCP; Ca3(PO4)2) also appear in the pattern. This TCP seems to

form due to reaction between HAp & ZrO2 at high temperature [9].

142

(a)

(b)

143

Fig. 9.7(a-d): XRD patterns of samples; (a) pure ZA(w) composite sintered at 1400ºC; (b) pure HAp

sintered at 800 & 1400ºC; (c) 20, 30 & 40H-CT-ZA(w); (d) magnified image of 40H-CT-ZA(w).

(c)

(d)

144

9.2.6 Cell Adhesion Biocompatibility Test

Figures 9.8 (a-e) show the SEM images of samples containing 0-40wt% HAp in

ZA(w) composite kept for 108 hours (72 hours with cell culturing media and 36 hours

without cell culturing media as described in section 9.2.4) in controlled atmosphere.

While Fig. 9.9 (a-c) shows composites, containing 30 & 40 wt% HAp kept for further

48 hours in oven without cell culturing media.

Figure 9.8 (a-c) which is for the samples containing 0%, 10% and 20% HAp

shows cell line growth although the cells do not properly adhere to the sample

surface. With increasing the amount of HAp, bonding of cells with underneath

composite surface increased, and well developed, adherent layer of the cells formed

on the surfaces of the samples containing 30% & 40% HAp, as shown in Fig. 9.8 (d)

& 9.8 (e), respectively). Close observation of the cultured cell line revealed that there

was little / no intercellular spacing. The cells were closely packed and outline of the

individual cells could not be distinguished. It is because by giving considerable time

for growth, the cell line grows to the point that a layer with no intercellular spacing is

formed [26]. It happens when RD cell line is clean and free from any contact

inhibition proteins [27].

In order to confirm that the surface layers observed in Fig. 9.8 are those of

cells and not some kind of artifact, the samples were kept in oven for further 48 hours

without cell culturing media to avoid cell growth and allow growth of contaminants

generally adhering to cells. As shown in Fig. 9.9 (a-c), some biological contamination

(e.g. bacteria, fungi, mycoplasma/virus etc) were grown on sample surface after 48

hours, which is a clear indication of the formation of cells layer on the samples that

allowed biological contamination from environment [28].

145

Fig. 9.8(a-e): Micrograph of surface of pellets after cell culturing (a) ZA(w); (b) 10H-CT-ZA(w); (c) 20H-CT-

ZA(w); (d) 30H-CT-ZA(w); (e) 40H-CT-ZA(w).

It was further observed that cell layers had formed on all the samples with /

without HAp. This is because TZ-3Y & alumina are bio-inert materials [29] and do

not cause any bio-toxicity. However, the adhesion of cells to the sample surface was

poor in the absence of HAp and improved as the amount of HAp was increased.

Below 30 wt% of HAp (Fig. 9.8(a-c)), the adhesion was not proper to keep the full

grown layer of cells. However, by increasing HAp concentration to ≥ 30 wt% (Fig.

146

9.8(d & e)) the adhesion improved sufficiently to keep full layer of cells adhered to it

due to presence of sufficient amount of HAp. Moreover, the TCP phase produced due

to reaction between ZrO2 and HAp further enhances biocompatibility because it

dissolves much faster in body fluid than HAp and acts as seed for formation of new

bones and is a source of Ca and P for growth of bones [4].

Fig. 9.9(a-c): Micrograph of surface of pellets after cell culturing (a&b) 40H-CT-ZA(w) after further 48

hours; (c) 30H-CT-ZA(w) after further 48 hours.

Consequently, on the basis of mechanical and biological properties discussed

above, a TZ-3Y+Al2O3 composite with 30wt% HAp and 1wt% CTAB may be

regarded as optimum composition for load bearing biomedical applications such as

dental restorations etc. It is the composition with sufficient biocompatibility (keeps

grown cell layer) as well as proper mechanical properties, i.e., hardness of 7.47GPa

and flexural strength of 684MPa. The mechanical properties for this composition are

even better than natural tooth which has hardness of 3.53GPa [30] and flexural

strength of 225 MPa [31].

147

9.3 Conclusions

A novel composite, (TZ-3Y+10wt% Al2O3(w))+30wt% HAp, with excellent

optimum cell adhesion and mechanical properties was synthesized for use in various

biomedical applications. A summary of the study performed in this regard is given

below:

TZ-3Y+10wt% Al2O3(w) composite, which has been proved in previous chapters to be

an excellent composition for load bearing biomedical applications was used as base

composition to enhance its bioactivity. The biocompatibility of the composite was

improved through the addition of hydroxyapatite (HAp). The addition of HAp was

found to decrease the hardness and flexural strength of the composites. It was

attributed to lower hardness and flexural strength of HAp, which allows transgranular

crack propagation through HAp grains. Although the pure HAp was stable during

sintering at 1400oC, its addition to TZ-3Y+10wt% Al2O3(w) composite resulted in its

partial decomposition to tricalcium phosphate. It was considered to be associated with

its reaction with ZrO2. However, the biocompatibility, as measured through cell

adhesion test, was found to improve with the addition of HAp. Below 30wt% HAp,

the adhesion was not so good, while at 30 wt% or above the adhesion was quite well.

Since increase in HAp content was not good for mechanical properties, therefore, the

minimum HAp content (i.e.,30wt%) which gave acceptable biocompatibility was

chosen as an optimum composition. The hardness and flexural strength for this

composition are 7.47 GPa and 684 MPa, respectively.

148

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151

Chapter 10

10 Conclusions

The purpose of the present work was the development of high strength zirconia based

composites through incorporation of alumina whiskers for advanced structural and

biomedical applications. In this regard, various synthesizing and processing aspects

were covered like preparation of alumina whiskers, improvement of their dispersion

in the matrix, optimization of sintering temperature and composition for best

combination of mechanical properties and hydrothermal phase stability. Effect of the

addition of HAp on processing parameters as well as on improvement in

biocompatibility of the composites was also studied. The main points, which may be

concluded at the end of this work, are as given below:

1. Alumina whiskers were introduced through decomposition of ammonium

aluminum carbonate hydroxide (AACH) in situ during sintering of the

composites. For this purpose, AACH whiskers were prepared by hydrothermal

route using urea and aluminum nitrate as precursor materials. The role of

concentration of precursor materials was very important for controlling the

morphology of nanostructures of AACH. At lower urea contents, relatively

larger particles were formed, which transformed to urchin-like structures and

then to AACH whiskers as urea content was increased. The size of urchin-like

nano-architectured particles was up to 4 µm while the diameter of individual

spikes on urchins was about 60 nm. However, the fully developed (separated)

AACH whiskers formed at higher urea contents had approximate dimensions

of 120 nm diameter and 5 µm length. The whiskers were subsequently used for

formation of composites.

2. TGA results revealed that volatile ingredients of AACH whiskers could be

completely removed by calcining the samples at temperatures higher than

650oC. Similarly, FTIR showed that complete decomposition of AACH takes

place at 650oC or above resulting in the formation of amorphous alumina as

152

revealed by XRD measurements. This amorphous alumina transforms to -

alumina at temperatures higher than about 1200oC.

3. The 3 mol% yttria stabilized tetragonal zirconia (TZ-3Y) was used as matrix

material due to its transformation toughening behavior, in which AACH

whiskers were incorporated. The AACH transformed into alumina whiskers

during sintering. For comparison, composites reinforced with particulate

alumina were also prepared. To achieve better dispersion of reinforcement in

the matrix phase, effect of CTAB and PABA as deflocculants was also studied.

It was found that each defloculant gave best results at 1 wt% concentration.

Between the two deflocculants, CTAB was relatively better than PABA in

improving the density and hardness of the composites. CTAB was completely

volatilized up to 650oC, while PABA or its combustion products were expected

to remain until sintering temperature and thus offer some detrimental effect.

4. Comparison between alumina whisker and alumina particulate reinforced

composites revealed that particulates were better to get higher densities of the

composites, while the hardness of particulate-reinforced composites was poorer

than that of whisker-reinforced composites. As far as the reinforcement

concentration is concerned, maximum hardness (i.e., 14.41 GPa) was obtained

at 10wt% of alumina whiskers, when CTAB was used as deflocculant. At

higher whisker contents, severe agglomeration deteriorated the mechanical

properties.

5. Like hardness, the flexural strength also increased with increasing the whisker

concentration approaching ~1325MPa at 10 wt% of alumina whiskers. Above

this concentration, the flexural strength decreased owing to severe

agglomeration of alumina whiskers.

6. In order to optimize sintering temperature, sintering of composite reinforced

with 10 wt% alumina whiskers was performed at three different temperatures,

i.e., 1400oC, 1500

oC and 1650

oC. Composites sintered at 1400°C had lower

density and higher retained porosity, while samples sintered at 1650°C showed

excessive grain growth and change in morphology of whiskers to rounded

grains. These factors resulted in poor hardness of composites sintered at

1400°C and 1650°C. Maximum hardness was obtained at 1500°C, which was

153

selected as the optimized sintering temperature. Combined effect of whisker-

like morphology, optimum grain size and higher density seem to be responsible

for higher hardness in this case.

7. To study the tetragonal phase stability of the developed composite in low

temperature moist atmosphere, the hydrothermal treatment at 135oC for 10 h

was employed. In monolithic tetragonal zirconia (TZ-3Y) the volume fraction

of monoclinic zirconia after hydrothermal treatment increased with increase in

sintering temperature. The addition of alumina whiskers enhanced the

destability of tetragonal phase significantly at samples sintered at lower

temperature. However, this destability decreased with increase in sintering

temperature, a behavior quite opposite to monolithic zirconia case. The

addition of alumina whiskers resulted in the formation of low-density regions

or pores into the composite. These pores act as points where water can absorb

and react with zirconia to form monoclinic phase. With increase in sintering

temperature, the porosity decreased and so did the destabilization of tetragonal

phase.

8. It was also observed that the composite with 10wt% Al2O3 whiskers after

sintering at 1650oC showed tetragonal phase stability even better than

monolithic zirconia. It was attributed to the formation a compact composite

due to transformation of whisker-like alumina to particulate alumina as well as

dissolution of alumina into zirconia forming tetragonal-phase-stabilizing solid

solution.

9. To improve the bioactiveness of the composition concluded as having best

mechanical properties (TZ-3Y+10wt%Al2O3(w)) the hydroxyapatite (HAp) was

incorporated. As HAp decomposes at higher temperatures, a lower sintering

temperature, i.e., 1400oC was used. The effect of various amounts of HAp on

the composite properties was studied. The density, hardness and flexural

strength decreased with an increase in HAp content. It was attributed to the

lower hardness and strength of HAp, which was present in the form of

relatively larger grains in the composite and allowed transgranular crack

propagation through its grains.

154

10. Although pure HAp did not decompose at 1400oC, it partially decomposed

when added to the composite. Reaction between zirconia and HAp was

attributed for this decomposition resulting in the formation of tricalcium

phosphate.

11. The biocompatibility of the composites was determined through in-vitro cell

adhesion test. The composites with less than 30 wt% HAp showed poor

adhesion of cell layer, while at 30 wt% or above the adhesion was very good.

The hardness and flexural strength for this composition are 7.47 GPa and 684

MPa, respectively, which are much better than natural tooth. Therefore, it was

regarded as a novel bioactive composite.