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Synthesis and Characterization of Nitrogen-rich Calcium -Sialon Ceramics Yanbing Cai Department of Physical, Inorganic, and Structural Chemistry Stockholm University 2009
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Page 1: Synthesis and Characterization of Nitrogen-rich Calcium ...200496/FULLTEXT01.pdf · Synthesis and Characterization of Nitrogen-rich . ... based ceramics are densified via liquid phase

Synthesis and Characterization of Nitrogen-rich

Calcium -Sialon Ceramics

Yanbing Cai

蔡 雁 兵

Department of Physical, Inorganic, and Structural Chemistry

Stockholm University

2009

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Doctoral Thesis 2009

Department of Physical, Inorganic and Structural Chemistry

Stockholm University

10691 Stockholm

Sweden

Cover

Nitrogen-rich corner in the Ca-Si-Al-O-N system

Faculty opponent:

Associate Professor Elis Carlström

Department of Materials Applications

Swerea IVF

Mölndal, Sweden

Evaluation committee:

Professor Yvonne Andersson, Uppsala University

Researcher Leif Hermansson, Uppsala University

Associate Professor Mats Johnsson, Stockholm University

Substitute:

Professor Ulf Hålenius, Swedish Museum of Natural History

© Yanbing Cai

ISBN 978-91-7155-823-7, pp. 1-57 Printed in Sweden by PrintCenter US-AB, Stockholm 2009

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To my family, for their love and support!

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I

ABSTRACT

In this thesis, a synthesis concept has been developed, which uses nitrogen-rich liquid

phases for sintering of Ca--sialon ceramics. First, keeping the Si/Al ratios constant, the

effects of N/O ratio on the properties and microstructure were investigated through a liquid

phase sintering process. Second, nitrogen-rich Ca--sialon ceramics, with nominal

compositions: CaxSi12-2xAl2xN16, x < 2.0, were synthesized and characterized. Third,

mechanical and thermal properties of nitrogen-rich Ca--sialons were investigated in terms

of high temperature deformation resistance, reaction mechanism, phase stability and

oxidation resistance, and further correlated to their phase assemblage and microstructure

observation.

It has been found that increasing the N/O and Ca/Al ratio simultaneously in the materials

could result in development of a microstructure with well shaped, high-aspect-ratio Ca--

sialon grains, and an improvement in both toughness and hardness.

For the nitrogen-rich Ca--sialon, mono-phasic -sialon ceramics were obtained for 0.51 ≤

x ≤ 1.32. The obtained Ca--sialon ceramics with elongated-grain microstructures show a

combination of high hardness and high fracture toughness. Compared with the oxygen-rich

Ca--sialons, the nitrogen-rich Ca--sialons exhibited approximately 150 oC higher

deformation onset temperatures and decent properties even after the deformation.

The -sialon phase was first observed at 1400 oC, however the phase pure Ca--sialon

ceramics couldn’t be obtained until 1800 oC. The nitrogen-rich Ca--sialons were thermal

stable, no phase transformation observed in the temperatures range1400 – 1600 oC. In

general, mixed -sialon showed better oxidation resistance than pure -sialon in the low

temperature range (1250-1325 oC), while -sialons with compositions located at -sialon

border-line showed significant weight gains over the entire temperature range tested

(1250-1400 oC).

Key words: Nitrogen-rich, SiAlON, Sintering, Microstructure, Properties

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LIST OF PUBLICATION

I. Yanbing Cai, Zhijian Shen, Jekabs Grins, Saeid Esmaeilzadeh, and Thomas Höche

Self-Reinforced Nitrogen-Rich Calcium -Sialon Ceramics

J. Am. Ceram. Soc., 90 (2), 608-613, 2007

II. Yanbing Cai, Zhijian Shen, Jekabs Grins, and Saeid Esmaeilzadeh

Sialon Ceramics Prepared by Using CaH2 as a Sintering Additive

J. Am. Ceram. Soc., 91 (9), 2997-3004, 2008

III. Yanbing Cai, Zhijian Shen, Thomas Höche, Jekabs Grins, Saeid Esmaeilzadeh

Super Plastic Deformation of Nitrogen-Rich Ca--Sialon Ceramics

Mater. Sci. Eng. A, 475, 81-86, 2008

IV. Yanbing Cai, Mats Nygren, Zhijian Shen, Jekabs Grins, and Saeid Esmailzadeh

Thermal Properties of Nitrogen-Rich Ca--Sialons

J. Eur. Ceram. Soc., accepted, 2009

Papers not included in this thesis

V. Yanbing Cai, Jekabs Grins, Zhijian Shen, and Saeid Esmaeilzadeh

Nitrogen-Rich -Sialons Stabilized by Y, Yb, and Nb Cations

in manuscript

VI. Changming Xu, Yanbing Cai, Flodström Katarina, Kjell Jansson, Zheshen Li, Guoqiang

Zhu, and Saeid Esmaeilzadeh

Spark Plasma Sintering of B4C Ceramics: the Effects of Milling Medium and TiB2

Addition

J. Mater. Res., submitted, 2009

Paper I, II and III are reprinted with the permission from the publishers.

II

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TABLE OF CONTENTS

ABSTRACT ........................................................................................................................... I

LIST OF PUBLICATION.....................................................................................................II

TABLE OF CONTENTS .................................................................................................... III

1. INTRODUCTION ............................................................................................................. 1

1.1 Background................................................................................................................... 1

1.2 Silicon nitride and SiAlON ceramics ........................................................................... 1

1.3 Oxynitride glasses and liquid phase sintering .............................................................. 3

1.4 Properties and applications........................................................................................... 6

1.5 Sintering techniques ..................................................................................................... 8

1.6 Parameters influencing the properties .......................................................................... 9

1.7 Toughening Si N based ceramics through anisotropic grain growth3 4 ......................... 9

2. AIMS OF THE WORK ................................................................................................... 11

3. EXPERIMENTAL........................................................................................................... 12

3.1 Specifications of starting powders.............................................................................. 12

3.2 Composition design .................................................................................................... 12 3.2.1 Nitrogen-rich liquid phase sintering of Ca--sialon ............................................ 12 3.2.2 Synthesis of nitrogen-rich Ca--sialon ................................................................ 14

3.3 Sintering...................................................................................................................... 16

3.4 Mechanical property characterizations....................................................................... 16

3.5 Microstructure and chemical composition characterization....................................... 16

3.6 Phase analysis ............................................................................................................. 17

3.7 Compressive deformation tests................................................................................... 17

3.8 Thermal properties characterization ........................................................................... 18

4. RESULTS AND DISCUSSIONS ................................................................................... 19

4.1 Nitrogen-rich liquid phase sintering of Ca--sialon .................................................. 19

III

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4.1.1 Phase compositions............................................................................................... 19 4.1.2 Microstructures and mechanical properties.......................................................... 20 4.1.3 Compressive deformation..................................................................................... 23

4.2 Synthesis of nitrogen-rich Ca--sialon ...................................................................... 25 4.2.1 Phase compositions............................................................................................... 26 4.2.2 Solubility of calcium in -sialon, and lattice parameters..................................... 27 4.2.3 Microstructures and mechanical properties.......................................................... 29

4.3 Superplastic deformation of nitrogen-rich Ca--sialon ............................................. 31 4.3.1 Compressive deformation behavior...................................................................... 32 4.3.2 Microstructures and properties ............................................................................. 32

4.4 Thermal properties of nitrogen-rich Ca--sialon ....................................................... 35 4.4.1 Reaction sequence and phase evolution ............................................................... 36 4.4.2 Thermal stability................................................................................................... 38 4.4.4 Oxidation resistance ............................................................................................. 40

5. SUMMARY..................................................................................................................... 44

6. FUTURE WORK ............................................................................................................ 45

ACKNOWLEDGMENT ..................................................................................................... 48

REFERENCES .................................................................................................................... 50

PAPER I - IV

IV

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1

1. INTRODUCTION

1.1 Background

Silicon nitride and its solid solutions, SiAlON ceramics, exhibit excellent mechanical

properties, good oxidation and thermal shock resistance at both room and high temperature.

The high wear resistance and excellent mechanical properties make these ceramics

competitive for applications such as cutting tools, balls and rollers for bearings, valves for

automotive engines or as turbine components. Because of the highly covalent bonding nature,

and the volatility of Si3N4 at high temperatures, Si3N4 has to be densified with sintering

additives by pressureless sintering (PLS), gas pressure sintering (GPS), hot pressing (HT), hot

isostatic pressing (HIP) or spark plasma sintering (SPS). The properties of silicon nitride

based ceramics are strongly dependent on their microstructures and phase compositions, and

on the chemistry of the intergranular grain boundary phase.

1.2 Silicon nitride and SiAlON ceramics

Si3N4 exists in two structural modifications: trigonal Si3N4 and hexagonal -Si3N4, both

of which are built up of Si(N4) tetrahedral with space groups P31c and P63/m, and unit cell

parameters a = 7.7541(4) Å, c = 5.6217(4) Å for -Si3N4 (PDF 41-0360), and a = 7.6044(2)

Å, c = 2.9075(1) Å for -Si3N4 (PDF 33-160), respectively.

SiAlON is a solid solution of Si3N4. The discovery of the -sialon solid solutions in the 1970s

[1],[2], based on substitution of Al and O for Si and N in -Si3N4, increased the interest in and

widened the range of possible applications of Si3N4 based ceramics. The general formula for

-sialon is Si6-zAlzOzN8-z, where the z value is in the range of 0 ~ 4.2 for samples formed at

1760oC. [3],[4] Fig.1.1 shows the behaviour diagram for the Si-Al-O-N system at 1730oC. The

phase that has attracted most interest is the -sialon solid solution region; ceramics with

compositions located in this region exhibit good sinterability and improved mechanical

properties.

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Fig.1.1 Phase relationships in the SiAlON system at temperature of 1730oC.[5],[6]

-sialon, the solid solution of -Si3N4[7] , was found later, with a general formula expressed

as MexvSi12-(m+n)Alm+nOnN16-n, where x < 2, x = mv, and m (Al-N), n (Al-O) replace (m+n)

(Si-N) bonds.

The unit cell parameters change is very small in the substitution of Al-O bonds for Si-N bonds

in the -sialon structure, because the bond lengths of Si-N (1.74 Å) and Al-O (1.75 Å) are

very similar; whereas the substitution of Al-N and Al-O bonds for Si-N bonds in -sialon

results in a considerable unit cell size expansion, due to the bond length difference between

Al-N (1.87 Å) and Si-N (1.74 Å). The cell parameters of -Si3N4 are a = 7.7541 Å,

c = 5.6217 Å, and for -Si3N4 they are a = 7.6044 Å and c = 2.9075 Å. The changes of cell

dimensions due to bond substitutions in -sialon, and most of the rare earth cation-doped -

sialons, fit reasonably with the empirical relationships:

-sialon: [8]

za = (a - 7.603)/0.0297 (Å)

zc = (c - 2.907)/0.0255 (Å)

z = (za + zc)/2 (Å)

-sialon: [7, 9]

a = 7.706 + 0.0117 m + 0.0824 rM + 0.055 x (Å)

c = 5.578 + 0.0259 m + 0.0774 rM + 0.0l71 n (Å)

where rM is the radius of the M ion.

2

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In Fig.1.2, a schematic phase relationship of the Ca-Si-Al-O-N system is presented. The Ca-

-sialon phase formation region, compared to most of the rare earth doped RE--sialons, is

much wider due to the high solubility of Ca cation in the -sialon structure (0.3 ≤ x ≤1.4 with

compositions on the tie line Si3N4-CaO: 3AlN [10, 11], whereas solubility ranges of RE--

sialons are in the range of 0.33 < x < 1.0 along the tie line Si3N4-RE2O3:9AlN) [12], depending

on the ionic sizes of the cations.

Fig.1.2 Schematic phase behavior diagram of the Ca-Si-Al-O-N system (incorporating data

from Wang, Hewett et al. and present research results) [10, 13]

The O-sialon [14] (Si2-xAlxO1+xN2-x, 0.04<x<0.4 at 1760 oC) is a derivative of Si2N2O and is

characterized by a greater oxidation resistance than most of the Si3N4 based ceramics. Due to

its superior oxidation resistance, O-sialon based materials,[15] like O-sialon-ZrO2, have found

application in handling of molten metals. But single phase O-sialon is hard to synthesize

because of impurities in starting powders and a narrow phase formation region.

1.3 Oxynitride glasses and liquid phase sintering

Si3N4 based ceramics are densified via liquid phase sintering, and they contain amorphous

phases, which may be described as M-Al-Si-O-N oxynitride glasses, at triple grain junctions

and grain boundaries. This intergranular glass phase is prone to partial devitrification if post

3

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sintering annealing treatments are applied. The effect of nitrogen on properties is much

greater than that of cations.[16] For a constant cation composition, the viscosity increases by

more than 2 orders of magnitude as 18 eq % oxygen is replaced by nitrogen. For rare earth

SiAlON glasses of constant composition, viscosity, Young’s modulus (E), hardness (H) and

glass transition temperature (Tg) increase linearly with decreasing ionic radius. Studies by

Becher et al. [17] on Si-Al-RE-based oxynitride glasses indicated that increases in the N:O

ratios of these glasses substantially increase the E, H, and Tg values, which can be attributed

to changes in network strength and non-bridging anion contents, but there was no direct

evidence for this interpretation.[18]

The presence of residual grain-boundary phases in the sintered Si3N4 based ceramics may

result in deterioration of high temperature properties, especially when a glassy phase has

relatively low glass transition temperature and viscosity.[19-21] Trying to improve the high

temperature performance of the Si3N4 based ceramics, efforts were made to control grain-

boundary chemistry, either by reducing the impurity contents of the glass[19] or by adding

more refractory-glass-forming oxides, such as Y2O3[22, 23], Lu2O3

[24], Yb2O3[25], Sc2O3

[26] as

sintering additives, or by crystallizing the intergranular phases through post heat treatment.[27,

28] However, studies have indicated that even a small addition of oxide to the SiAlON will

greatly impair the high temperature properties. On the other hand, in the preparation of

oxynitride glasses, improvement in properties has been obtained through controlling the N/O

ratio.[17] The substitution of N for oxygen in the Si(O4) network of tetrahedra results in a

higher bonding density per unit volume of the glass, thus improving the desired mechanical

properties of the glasses.[21, 29, 30] Recent research on Ca-Si-O-N glasses has also indicated that

the hardness and glass transition temperature are improved dramatically with a nitrogen

content of 70 eq %.[31] Therefore, improvements in high-temperature properties are expected

if Si3N4 based ceramics can be fabricated by forming a nitrogen-rich liquid phase during the

sintering process.

The elongated -sialon grains were usually obtained via oxygen-rich liquid phase sintering,[32-

34] resulting in toughening of materials, however compromised by decreasing hardness and

declining high temperature properties because of the presence of a residual oxygen-rich glass

phase, or secondary grain boundary phases.[34-36]

4

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Traditionally, the aspect ratio of -Si3N4 grains is considered to be a function of the rare earth

ionic radius. In compositions of RE0.4Si9.6Al2.4O1.2N14.8, Re = Yb, Er, Y, Dy, Gd, Sm, and

Nd, the amount of elongated RE--sialons was found increased with increasing ionic radii of

rare earth cations.[37] Painter and Becher et al.[38-39] introduced a differential binding model to

explain the mechanisms of interfacial segregation of rare earth cations and chemical bonding

in regions of variable N/O content. The anisotropic grain growth is found to originate from the

site competition between rare earth cations and Si for bonding at N-rich -Si3N4 interfaces

and with the O-rich glasses. The differential binding energy of rare earths (DBE, relative to

that of Si) quantifies the bond preference of rare earths with O in the glass over N at the

interface. Rare earths for which DBE > 0, like Lu, reflect the preference for bonding with O in

glass (O-rich region). Rare earths for which DBE < 0, like La, effectively compete with Si for

N at the -Si3N4 grains interfaces (N-rich environments). Elements that segregate to the prism

planes of the -Si3N4 grains impede the attachment of Si-based silicon nitride growth units,

and the extent of this limitation reduces the growth in diameter and leads to anisotropic grain

growth along the c direction. The segregation of Ca cations at grain boundary films is also

observed in calcium doped Si3N4 ceramics. [40] The calcium segregation in grain boundary

films changes the film composition dramatically by substituting more N3- for O2- anions to

maintain the local stoichiometry.

While there is little evidence for the existence of Al-N bonds in the oxynitride glasses, [41] it

was found that Si-O, Si-N and Al-O are the major bonds in Si-Al-Y oxynitride glasses [42, 43]

and the glasses are built up of Si(O4), Si(O3N), Si(O2N2) and Al(O4) tetrahedral.[44] The

reason is that, theoretically, tetrahedra of Si(O4), Si(O3N), Si(N2O2) and Al(O4) are more

stable than the others. Research by Becher and Sun et al. [45-47] has indicated that interfacial

bonding strength, which dominates the debonding behaviour of the interface between

oxynitride glasses and Si3N4 grains, increases with increasing Al and O contents of the

epitaxial SiAlON layers formed on the grains. The decrease in binding energies with N

substitution for O has important impact on the interface strength, thus the fracture toughness,

since the binding energies of Si(N,O)4 and Al(N,O)4 tetrahedra systematically decrease as

nitrogen replaces the oxygen (Fig.1.3).[45, 46]

5

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0 1 2 3 40

2

4

6

8

10

12

14

16

18

20

22

24

SiN4

SiN3O

1

SiN2O

2

SiN1O

SiO4

AlO4

AlN1O

AlN2O

2

AlN3O

1

AlN4

AlN4-x

Ox

SiN4-x

Ox

Bo

nd

ing

En

erg

y (e

V)

x in AlN4-x

Ox or SiN

4-xO

x tetrahedra

Fig.1.3 Bonding energies of Al(N,O)4 and Si(N,O)4 tetrahedra.

1.4 Properties and applications

The properties of Si3N4 based materials are strongly related to the phase compositions and

morphology, as representatively shown in Table 1.1. The Vickers hardness of Si3N4 based

materials is in the range of 14–22 GPa, and toughness in the range of 3–10 MPa·m1/2.

Generally, the -phase Si3N4 and SiAlON (having equiaxed morphologies) are harder, but

lower in toughness than -phase Si3N4 and SiAlON with elongated grains. But this is not

always the case, since harder and tougher -sialon ceramics have been reported in the

literature.

Table 1.1 Properties of Si3N4 based ceramics [6, 48, 49].

Unit cell parameter (Å) Mechanical properties

a c Morphology

HV10 (GPa) KIC (MPa·m-1/2)

-Si3N4 7.7541(4) 5.6217(4) Equiaxed <20 3

-Si3N4 7.6044(2) 2.9075(1) Elongated <16 47

-sialon 7.8017.864 5.6795.720 Equiaxied

Elongated

16–22

16–22

3–4

5–10

-sialon 7.6107.716 2.9113.007 Elongated 14–18 5.0–7.0

6

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The unique combinations of properties such as good wear and creep resistance abilities, high

strength, toughness and hardness allow Si3N4 based ceramics to be candidates for a wide

variety of applications. A summary of general property requirements of Si3N4 based materials

for specific applications is listed in Table 1.2 [50, 53].

Table 1.2 Property requirements for specific applications

Application Key properties

Cutting tool inserts Hardness, toughness, thermal shock resistance, strength, thermal conductivity,

chemical stability

Bearings Toughness, strength, hardness, smoothness, porosity, grain size, wide range of

temperature

Wear parts Toughness, hardness, smoothness, coefficient of friction, generally room

temperature

Heat engines Reliability, high temperature corrosion resistance, toughness, strength, thermal

shock resistance, Weibull modulus, coefficient of friction

Other thermal applications Chemical resistance, thermal stress resistance, creep resistance

Table 1.3 Representative properties of Si3N4 based ceramics and other cutting tool materials.

Coefficient

of thermal

expansion,

Young’s

modulus

E

Bending

strength

Toughness

KIC

Hardness

HV10

Hot

hardness

HV1, 800oC Tool material

10-6 K-1 GPa MPa MPa·m1/2 GPa MPa

Ref.

-sialon 3.4 300 650 4–7 14–17.5 850–1350 52

-Si3N4 -

-

530 4.2

14.7

(HV1) - 53

Ce--sialon 0

.3

1-3.2 300 0–900

0-4.0 300

lon

–19

6.0

n (seeded) 860 0

% 4 5 0

na 7 0

3-SiC -

900

C 4.3 300–400 400–800 3 25 - 57

SiC-SiCw - 250–270 580 12–18 - - 57

-

-

72 4.3

18

(HV1) - 53

Sialon 3. 80 5.5–6.5 - - 54

Y--sialon 3. - 3.5–5.5 20–22 - 55

SiCw-Y--sia - - - 3.5–6.9 19–20 - 55

MoSi2-Y--sialon - - - 3.5–5.5 17 - 55

Ca--sialon - - - 3.0– 16–19 - 10

Ca--sialo - - 560– 5.2–8. - - 56

WC-Co 6wt 5. 630 2100 10 7. 65 52

Alumi 8. 400 600 3-5 17 80 52

Al2O 400 850 8.5 - - 57

ZrO2 - 205 600– 3–9 - - 57

Si

7

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The most widely used cutting tool materials are based on Al2O3 (pure or dispersed with TiC,

TiN, ZrO2, SiCw), cemented carbide (WC–Co), and Si3N4 (-Si3N4, -sialon). A

comparison of properties of these typical cutting tool materials is listed in Table 1.3. The high

thermal shock resistance of SiAlON ceramics, determined by the interaction of thermal

conductivity, tensile strength, thermal expansion coefficient, and Young’s modulus, make

them particularly suitable for heavy-feed-rate or interrupted-cutting machining operations.

High room-temperature hardness and hot hardness give them excellent abrasion and

eformation resistance.

.5 Sintering techniques

to the structure to

and

often exhibit tailored microstructures and properties

ependent on the processing techniques.

d

1

To fabricate a useful Si3N4 based ceramic, the aim must be to provide a liquid at high

temperature that will allow densification through a solution–diffusion–reprecipitation process,

and then, by proper cooling or by heat treatment, to incorporate the liquid in

give a single-phase product, or to produce a crystalline intergranular phase.

Direct reaction sintering (RS) means the nitridation and sintering of Si particles in an N2

atmosphere. However, because of the energy costs of a slow, high-temperature process

the difficulties of achieving dense sintering, the RS process is not inexpensive in practice.

Dense Si3N4 and SiAlON ceramics as structural materials are usually densified by hot

pressing (HP, only for producing simple shape materials), or hot isostatic pressing (HIP,

capable of producing complex shape materials but requiring complicated equipment design),

although pressureless sintering (PLS, requiring a large amount of liquid phase) of pre-

compacted powders has been commercially applied. Recently, spark plasma sintering (SPS,

capable of densifying ceramics at low temperatures and in short time, but is also limited in

obtaining materials with simple shape) has been reported to densify Sialon ceramics within a

just few minutes. The resulting ceramics

d

8

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1.6 Parameters influencing the properties

There are a number of factors influencing the properties of Si3N4 based ceramics, such as

phase composition, sintering additives, processing methods and microstructure. Normally,

equiaxed -Si3N4 and -sialon are harder, while elongated -Si3N4 and -sialon are tougher.

For high strength materials a microstructure consisting of fine equiaxed grains is desired; a

microstructure consisting of moderately elongated grains with some equiaxed grains filling in

among them has been proved to produce ceramics with high toughness. Regarding high

temperature properties, the amount and properties of the grain boundary glass phase strongly

influences the high temperature strength, creep resistance, and therefore the hot hardness.

The intergranular grain boundary phases formed after liquid phase sintering of Si3N4 based

ceramics often cause serious degradation of high temperature properties. A number of

investigations have indicated the practical difficulties in the formation and densification of

pure, single phase -sialon using oxides as sintering additives, which usually cannot be fully

incorporated into the -sialon structure, but will be present as intergranular grain boundary

glass phases or secondary grain boundary phases. Use of non-oxide additives can significantly

reduce the amount of glassy phase and increase its softening temperature and viscosity in

comparison with Si3N4 sintered with oxides. Greskovich et al. [58] used Be3N2 or BeSiN2 as a

sintering aid and obtained -sialon with good creep resistance, high temperature bending

strength, but a relatively low toughness (~ 4 MPa·m1/2). Uchida et al. [59] found that VN, YN

and Mg3N2 were quite effective additives for consolidating Si3N4. Recently, Xie et al. [60]

prepared Ca--sialon ceramics using Ca3N2 as a precursor, but did not report about their

mechanical properties.

1.7 Toughening Si3N4 based ceramics through anisotropic grain growth

In the sintering of oxide ceramics, it is important to prevent abnormal grain growth so as to

obtain a microstructure with fine and equiaxed grains. In the preparation of Si3N4 based

ceramics for engineering applications, however, a bimodal microstructure consisting of a

small number of large elongated grains and a large number of small equiaxed grains is

9

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important to ensure high toughness as well as high strength. Toughening through insitu

formation of elongated grains, rather than addition of whiskers, has advantages over

processing. In addition, in whisker toughened Si3N4 based materials the matrix is frequently

not compatible chemically with the second phase. Reaction between them will not provide the

weak interface required for crack deflection or accommodation of vacancies. But a coherent

interface is required to maintain other properties like hardness and strength.

Self-reinforced Si3N4 and SiAlON materials have been obtained with toughness as high as 8–

10 MPa·m1/2. A most reasonable explanation for such high toughness is whisker like

reinforcement by the development of elongated -Si3N4, -sialon, or recently -sialon grains

in the microstructures. The elongated grains act as fibers in the sintered body. When a crack is

propagating, the surface energy to by-pass a grain and the friction to pull out the grain require

the applied stress to spend more energy to fracture the material. Thus, the reason for the high

toughness might be the shape of the grain rather than the intrinsic toughness of Si3N4 itself,

because the fracture toughness of the single crystal Si3N4 [61, 62], being 1.9-2.8MPa·m1/2, is

lower than that of sintered Si3N4 containing elongated -Si3N4 grains and grain boundary

phase, while the single crystal Si3N4 exhibits a higher Vickers hardness of 28–30 GPa.

10

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2. AIMS OF THE WORK

The successful development of -sialons with elongated grains has stimulated the research

activities on Si3N4 based ceramics. The particular properties of Ca--sialon ceramics: high

solubility of Ca2+ in the -sialon structure, high-temperature thermal stability, and feasibility

of forming a whisker like microstructure make the Ca--sialon an interesting system for

extensive investigation. Again, the linear increase of hardness, toughness, and glass transition

temperature with increasing N/O ratio found in the oxynitride glasses also suggests a possible

way of synthesizing -sialon ceramics with enhanced properties. For the Ca-Si-Al-O-N

system, a number of studies have been carried out with respect to solubility limits, phase

formation region, and properties. However, there is no systematic study on solubility range,

phase identification, microstructures, and properties of nitrogen-rich Ca--sialon ceramics.

Therefore, in this work, firstly effects of nitrogen-rich liquid phase sintering on the properties

and microstructures of Ca--sialon ceramics are studied. Then synthesizing and

characterization of nitrogen-rich Ca--sialon ceramics are investigated. Specifically the goals

are:

Examining the influence of glass compositions on the properties and microstructure of the

Ca--sialon ceramics with varying N/O but fixed Si/Al ratios, which can be designed

simply by adding CaH2 into -sialon (Si6-zAlzOzN8-z, with 0 < z < 4.2) based

compositions.

Synthesizing nitrogen-rich Ca--sialon ceramics with nominal compositions CaxSi12-

2xAl2xN16 extending along the Si3N4-1/2Ca3N2:3AlN tie line. Correlating phase

assemblages, lattice parameters, mechanical and thermal properties of the nitrogen-rich

Ca--sialon ceramics with their calcium contents and microstructures.

11

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3. EXPERIMENTAL

3.1 Specifications of starting powders

The starting powders (Table 3.1) are commercial -Si3N4 (UBE, SN-10E), -Si3N4 (Nobel

Grade P95 Size H), Al2O3 (Alcoa, SG16), AlN (HC Stark Grade A), CaH2 (AlfaAesar) and

CaCO3 (E. Merck Darmstadt). When calculating the overall compositions, corrections were

made for the small amounts of oxygen present in the Si3N4 and AlN raw mixtures.

Table 3.1 Starting powders used in present study

Chemicals Manufacturer Specifications

Si3N4 Nobel Grade P95 Nobel Industries Grade P95 Size H, >95%, O<2.0%

Si3N4 UBE SN-E10 UBE Europe GmbH BET:9~13m2/g,O<2.0%, <5wt%

AlN Grade A HC Stark O < 1.0 %, d50 7.0 - 11.0 μm, BET < 2.0 m2/g

Al2O3 ALCOA, SG16 Sandvik Coromant, 99.9%

CaH2 AlfaAesar -10 mesh, 98%, Mg < 1%

CaCO3 E. Merck Darmstadt 99.9%,

3.2 Composition design

3.2.1 Nitrogen-rich liquid phase sintering of Ca--sialon

In order to investigate the effects of N/O ratios on the microstructures and properties of

SiAlON materials, nitrogen-rich compositions were designed in present study by adding CaH2

into -sialon (Si6-zAlzOzN8-z, with 0 < z < 4.2) based compositions so as to obtain

compositions with fixed Si/Al but varying N/O ratios (Table 3.2).In order to enhance reaction

kinetics CaH2 instead of Ca3N2 is used as a precursor. CaH2 decomposes at relatively low

temperatures, at ca. 400oC, to fine-grained and nominally oxygen-free Ca metal. Previous

studies by us[31] have shown that alkali earth and rare earth metals exhibit comparatively high

reactivities towards Si3N4 and AlN in N2 gas atmosphere, most probably in the form of

freshly formed, and therefore reactive, nitrides. As illustrated in Fig.3.1, the resulting

12

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nitrogen-rich overall compositions, located on the tie line of -sialon-2Ca3N2, are above the

-sialon plane but not compatible with AlN’ polytypoids that might be observed in

compositions theoretically located on the -sialon plane, but actually below it due to the

ubiquitous existence oxides on the surfaces of the starting Si3N4 and Al2O3 particles. The

corresponding glass forming region of the nitrogen-rich compositions at sintering

temperatures, somewhere within the square-base pyramid of 2Ca3N2-4AlN-Si3N4-3SiO2-

2Al2O3, is remote from the region containing CaO.

Table 3.2 Starting compositions with addition of extra CaH2 into -sialon based compositions

CaH2 Si3N4 Al2O3 AlN Sample

Nominal

Compositions (in eq%) wt% wt% wt% wt%

A0 Ca0Si80.1Al19.9O22.2N77.8 0 73.69 26.31 0

A1 Ca0.6Si79.7Al19.7O22.1N77.9

1 72.95 26.05 0

A3 Ca1.8Si78.7Al19.5O21.8N78.2 3 71.48 25.53 0

A5 Ca3.1Si77.7Al19.2O21.5N78.5 5 70.01 25.00 0

A7 Ca4.4Si76.6Al19.0O21.2N78.8 7 68.53 24.47 0

A9 Ca5.7Si75.6Al18.7O21.0N79.0 9 67.06 23.95 0

A11 Ca7.0Si74.5Al18.5O20.7N79.3 11 65.58 23.42 0

B0 Ca0Si80.3Al19.7O14.1N85.9 0 75.43 15.64 8.93

B1 Ca0.6Si79.8Al19.6O14.0N86.0 1 74.68 15.49 8.84

B3 Ca1.8Si78.8Al19.3O13.9N86.1 3 73.17 15.17 8.66

B5 Ca3.0Si77.9Al19.1O13.7N86.3 5 71.66 14.86 8.49

B7 Ca4.3Si76.8Al18.9O13.5N86.5 7 70.15 14.54 8.31

B9 Ca5.6Si75.8Al18.6O13.3N86.7 9 68.64 14.23 8.13

B11 Ca6.9Si74.8Al18.3O13.2N86.8 11 67.13 13.92 7.95

B13 Ca8.2Si73.7Al18.1O13.0N87.0 13 65.62 13.61 7.77

B15 Ca9.5Si72.7Al17.8O12.8N87.2 15 64.12 13.29 7.59

C0 Ca0Si89.6Al10.4O9.7N90.3 0 86.45 9.67 3.87

C1 Ca0.6Si89.1Al10.3O9.6N90.4 1 85.58 9.57 3.85

C3 Ca1.76Si88.0Al10.2O9.5N90.5 3 83.85 9.38 3.77

C5 Ca3.0Si86.9Al10.1O9.4N90.6 5 82.12 9.18 3.69

C7 Ca4.2Si85.8Al10.0O9.3N90.7 7 80.39 8.99 3.61

C9 Ca5.4Si84.7Al9.8O9.2N90.8 9 78.67 8.80 3.54

C11 Ca6.7Si83.6Al9.7O9.0N91.0 11 76.94 8.60 3.46

C13 Ca8.0Si82.4Al9.6O8.9N91.1 13 75.21 8.41 3.38

C15 Ca9.3Si81.3Al9.4O8.8N91.2 15 73.48 8.22 3.30

13

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AlN polytypes

-sialon plane

4/3CaAlSiN3

Gehlenite

Anorthite

3SiO2

C B A overall composition

L2

L1

O-sialon-sialon

4/3(AlN·Al2O

3)

12/11(CaO·3AlN)

1/2(Ca3N

2)·3AlN

2Ca3N

26CaO

2Al2O

3

4AlN

Si3N

4

Fig.3.1. Schematic illustration of the Jänecke prism of the Ca-Si-Al-O-N system. Overall

compositions located on the tie line Ca3N2–-sialon (with low z values in formula

Si6-zAlzOzN8-z), are expected to produce /-sialons equilibrated with a nitrogen-rich liquid

with increasing Ca3N2 content at elevated temperatures.

3.2.2 Synthesis of nitrogen-rich Ca--sialon

The second group is nitrogen-rich Ca--sialon, named CCH series, with starting powders of

Si3N4 (UBE, SN-E10), AlN (H. C. Starck, grade A) and CaH2 (99.8%, Johnson Matthey

Chemicals Ltd.), with general compositions designed according to the formula CaxSi12-

2xAl2xN16, 0.2 x 2.6 (see in Fig. 3.2). Oxygen contents in the starting powders are taken

into account when calculating the actually achievable compositions. Two oxygen-rich Ca--

sialons, Cam/2Si12-(m+n)Alm+nOnN16-n with n = m/2 and m equaling 2.8 and 3.2, labeled as

respectively CCO14 and CCO16, were also prepared under the same conditions, using CaCO3

(99.9%, E. Merck Darmstadt ), Si3N4, and AlN as precursors (Table 3.3).

14

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0.0 0.4 0.8 1.2 1.6 2.0 2.4 2.8

N-rich -sialon+AlN

Si 6-nAl n

O nN 8-n

-Si 3N 4

+

-sial

on+-sialon(O

-rich

)

.

.

.

O-rich -sialon-s

ialon(N

-rich

)

N-rich -sialon

Si3N

4

-sial

on 4/3(A

lN A

l 2O 3

)

1/2(Ca3N

2) 3AlN

CaO 3AlN

n=4

n=3

n=2

n=1

x value

Fig.3.2 Tentative phase plane of Ca-sialon

Table 3.3 Nominal compositions of starting mixtures.

CaH2 Si3N4 AlN CaCO3 Sample

labels Overall compositions

wt% wt% wt% wt%

CCH02 Ca0.2Si11.6Al0.4N15.53O0.70 1.47 95.64 2.88 0

CCH04 Ca0.4Si11.2Al0.8N15.54O0.69 2.91 91.38 5.71 0

CCH06 Ca0.6Si10.8Al1.2N15.54O0.68 4.32 87.20 8.47 0

CCH08 Ca0.8Si10.4Al1.6N15.55O0.67 5.71 83.11 11.18 0

CCH10 Ca1.0Si10.0Al2.0N15.56O0.67 7.06 79.10 13.84 0

CCH12 Ca1.2Si9.6Al2.4N15.56O0.66 8.39 75.18 16.44 0

CCH14 Ca1.4Si9.2Al2.8N15.57O0.65 9.69 71.34 18.99 0

CCH16 Ca1.6Si8.8Al3.2N15.57O0.64 10.96 67.55 21.48 0

CCH18 Ca1.8Si8.4Al3.6N15.58O0.63 12.21 63.85 23.93 0

CCH20 Ca2.0Si8.0Al4.0N15.58O0.63 13.44 60.22 26.34 0

CCH22 Ca2.2Si7.6Al4.4N15.59O0.62 14.64 56.67 28.69 0

CCH24 Ca2.4Si7.2Al4.8N15.59O0.61 15.82 53.17 31.00 0

CCH26 Ca2.6Si6.8Al5.2N15.6O0.60 16.98 49.75 33.27 0

CCO14 Ca1.4Si7.8Al4.2N14.6O1.4 0 53.87 25.43 20.70

CCO16 Ca1.6Si7.6Al4.8N14.4O1.6 0 48.54 28.37 23.09

15

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3.3 Sintering

The starting powder materials, in batches of 30–50 g, were planetary milled in a sealed tank

with hexane as medium for one hour, using Si3N4 balls. The powders were dried in a vacuum

furnace and quickly moved into an argon-filled, water-free glove box. Pellets of the mixtures

(4.5g), compacted using a steel die, were hot-pressed in BN-coated graphite dies in nitrogen

atmosphere in a graphite resistance furnace at 1800oC for 4 h under 35 MPa uniaxial pressure.

The samples were heated to 1500oC with a heating rate of 20oC /min, held there for 1 h, then

quickly heated with a heating rate of 40oC /min to 1800oC and held there for 4 h, if not

specified otherwise.

3.4 Mechanical property characterizations

The densities of the sintered specimens were measured in water according to Archimedes’

principle. Samples for physical characterization were ground and carefully polished, using

standard diamond polishing techniques, down to a 1 m surface finish. Hardness (Hv10) and

indentation fracture toughness (KIC) were determined with a Vickers diamond indenter and a

98 N load (8-10 tests each), according to the method of Anstis et. al. [63]

3.5 Microstructure and chemical composition characterization

Microstructure observations on polished and fracture surfaces were carried out with a

scanning electron microscope (SEM) (JEOL JSM 880, or JSM-7000F) equipped with an EDX

system (acceleration voltage 20 kV Link Isis, Oxford Instruments) and a transmission electron

microscope (TEM, Model JEOL JEM 4010, acceleration voltage 400 keV, or JEM-3010

operated at 300 kV). Prior to SEM investigation, the polished surfaces of the samples were

etched in a molten mixture of KOH and KNO3 for 1–3 min before carbon coating. On

average, 5 area and 15 point EDX analyses of the cation composition were made on both

polished only and polished and etched surfaces. The accuracy of the analyses was ensured by

using a single phase anorthite standard (CaAl2Si2O8). The nitrogen and oxygen contents were

16

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determined with a LECO TC-436 DR machine , by taking the average of three separate

measurements. To determine the aspect ratio, the size of particles (length and diameter) was

measured quantitatively, using an image analysis program (Image Tool, UTHSCSA). 250–300

particles per specimen condition were measured in back-scattered SEM images from polished

and chemically etched surfaces.

3.6 Phase analysis

XRPD data was collected using a Guinier-Hägg focusing camera with 50 mm radius, with

CuK1 radiation and Si as the internal standard. The films were scanned with a micro-

densitometer, and the transmission data were processed with the program SCANPI.[64] Unit

cell parameters, x values and phase fractions were determined by the Rietveld method, using

the refinement program Fullprof [65] and data up to 2 = 88. Approximately 20–25 parameters

were refined, and the atomic coordinates for α-sialon [66], -sialon [67], AlN [68] and CaAlSiN3

[69] held fixed. Obtained estimated standard deviations of refined parameters were multiplied

by ~ 2 in order to correct for serial correlation in the data. [70] Unit cell parameters were

obtained by multiplying the values from the refinement by the correct unit cell parameter for

Si [71] divided by its refined value. The values were in very good agreement with values

obtained by using the unit cell refinement program PIRUM. [72] The relative Bragg peak

intensities for the Ca--sialons were found to be highly sensitive to the x value. The latter can

therefore be determined to a relatively high precision by the Rietveld method.

3.7 Compressive deformation tests

The compressive deformation[73] of selected samples (CCH02, 04, 08,12,16,20 and CCO14,

16, corresponding compositions are listed in Table 3.2) were performed with an SPS

equipment (Dr. Sinter 2050, Sumitomo Coal Mining Co.) by applying a uniaxial compressive

stress on cylindrical hot pressed specimens, 12 mm in diameter and 6 mm in height, placed

within a graphite die with 20 mm inner diameter. A load corresponding to an initial stress of

40 MPa was applied and held constant during the entire deformation process, implying that

17

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the applied compressive stress decreased from 40 MPa to 20 MPa when reaching 50% strain.

The compressive strain was defined as –L/L0, with L and L0 the shrinkage of sample

height and original height of the sample before deformation, respectively. A constant heating

rate of 40oC/min was applied.

3.8 Thermal properties characterization

In order to reveal the reaction sequences, starting powders of CCH16 composition were hot

pressed in the temperature range 600 -1800 oC for 2 hours using a pressure of 35 MPa, followed

by an X-ray powder diffraction analysis. For the same composition, thermal gravity analysis was

also carried out in a TG unit (SETARAM TAG 24, Setaram, France) using a heating rate of 10

oC/min in nitrogen atmosphere, with a purpose of monitoring the decomposition and nitridation

processes supposed to happen with elevated temperature up till 1200 oC.

Oxidation experiments were performed in the same TG unit at 1250 oC, 1325 oC and 1400 oC.

The samples were heated to these temperatures with heating rate of 10 oC/min and the duration of

the oxidation experiment was 20 h in flowing oxygen. Prior to oxidation, samples of the

approximate size of 10×3×1 mm3 were cut, ground, and polished with diamond suspensions down

to 1 m. The samples were connected to the hang down wire of the TG-unit via a notch with

depth of 1 mm that was made at the end of sample bars by using a 0.5 mm thick diamond blade.

The drift of the set up was corrected via recording time versus weight loss curves of Al2O3

dummies at 1250 oC, 1325 oC and 1400 oC, respectively.

18

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4. RESULTS AND DISCUSSIONS

4.1 Nitrogen-rich liquid phase sintering of Ca--sialon

In developing elongated -sialon grains, a nitrogen-rich liquid phase sintering method was

introduced by using CaH2 as a sintering aid, so as to vary the N/O ratio of the liquid phase

formed in the sintering process while keeping the Si/Al ratios constant ( see Fig.3.1). The

purpose is to obtain a material with a combination of high toughness and hardness through the

formation of elongated -sialon grains with a nitrogen-rich residual grain boundary glass

phase. The results indicate that:

With increasing CaH2 addition the phase contents changed from single -sialon to dual

-sialons and to single Ca--sialon. At low N/O ratios the microstructures contained

mainly equi-axed -sialon grains, and at high N/O ratios well faceted elongated Ca--

sialon grains. The improved toughness (KIC = 7.8 MPa·m½) and hardness (HV10 = 17.5

GPa) properties can be attributed to the formation of interlocked microstructures.

High-temperature compressive deformation tests indicated that the deformation onset

temperature is determined mainly by the Si/Al and N/O ratios, whereas the

deformation rate is affected by the microstructure, i.e. the morphology and amounts of

elongated -sialon grains and residual glass phase, especially for the sialons with low

N/O ratios.

4.1.1 Phase compositions

Single-phase -sialon materials were obtained without addition of CaH2 in all three series A,

B and C. Upon increasing the amount of CaH2 addition, both -, -sialon phases were found,

and single-phase Ca--sialon materials formed at high CaH2 contents (above 9 wt% for A and

above 7 wt% for B and C, see in Fig.4.1). There was no observation of AlN polytypoids or

calcium aluminum silicates, which are usually present as secondary phases in Ca--sialon

materials when using CaO as a precursor, implying a complete transformation of -Si3N4 to

19

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-or-sialon.

0 2 4 6 8 10 12 14 16

0

20

40

60

80

100

, %

Addition of CaH2, wt%

A series B series

0 1 2 3 4 5 6 7 8 90.4

0.6

0.8

1.0

1.2

1.4

1.6

1.8

zA

zB

zC

Cal

cula

ted

z, Å

Addition of CaH2, wt%

C series

Fig.4.1 ratio as a function of CaH2 addition. Pure Ca--sialon was obtained in A

series, B series and C series when CaH2 additions were above 9, 7 and 5 wt%, respectively.

Based on the refined lattice parameters, the estimates of z values for -sialon and x values for

-sialon (shown in insert of Fig. 4.1 ) were obtained by using the relationships between lattice

expansion and the extent of Al–O and Al–N bonds substitution for Si–N bonds provided by

Ekström et al. [8] and Wang et al., [10] respectively. The calculated z values (mean values of za

= (a - 7.603)/0.0297 Å, and zc = (c - 2.907) / 0.0255 Å) generally decrease with increasing

CaH2 content, whereas the calculated x values (mean values of xa = (a - 7.749) / 0.156 Å, and

xc = (c - 5.632) / 0.115 Å) show a reverse trend, corresponding to a unit cell expansion of the

Ca--sialon structure with increasing Ca2+ content.

4.1.2 Microstructures and mechanical properties

It was found that the addition of CaH is effective in controlling the fraction and size of

elongated Ca--sialon grains. The formation of a nitrogen saturated liquid, and the Ca2+

cations serving as glass network modifiers at elevated temperatures, kinematically facilitates

the development of elongated Ca--sialon grains. Statistic grain size distribution analysis,

used as a qualitative tool to show the aspect ratio development with respect to CaH2 addition,

20

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indicates that the number and aspect ratio of the formed elongated Ca--sialon grains strongly

depend on the amount of liquid phase during sintering. As shown in Fig.4.2, a finer-grained

microstructure was observed for the samples with low CaH2 addition, compared to the

specimens with high amounts of CaH2 addition. Increasing the N/O and Ca/Al ratio

simultaneously in the materials could result in the development of a microstructure with well

shaped, high aspect ratio Ca--sialon grains (Fig.4.3).

02

46

810

1214

16

C11C5

C0

C15

B15B11

B5B0

A11A9

A5A0

a) frequency count of grain aspect ratio%

0.0

0.5

1.0

1.5

2.0

2.5

C15C11

C5C0

B15B11

B5B0A11

A9A5

A0

b) frequency count of grain diameter

m

0.00.51.01.52.02.5

c) frequency count of grain diameter, B series

B15B13

B11B9

B7B5

B3B1

B0 m

Fig. 4.2. Statistic analysis of grain aspect ratio and diameter distribution shows the changes

with addition of CaH2 contents in the sialon ceramics: (a) frequency count of grain aspect

ratio, (b) frequency count of grain diameter, and (c) typical frequency count of grain diameter

of B series as a function of CaH2 addition.

21

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Fig.4.3 Representative back-scattered SEM images of SiAlON ceramics show the grain sizes

and morphologies without CaH2 addition: (a) A0, (c) B0 and (e) C0; with 11 wt% of CaH2

addition: (b) A11, (d) B11 and (f) C11.

As expected, both density and hardness increase with increasing N/O ratio (Fig.4.4). All

samples hot pressed at 1800oC are found to be fully densified, and the densities increase with

increasing amount of CaH2. About 5 wt% of CaH2 addition is enough to increase the

toughness from 3 – 4 MPa·m1/2 to 6–7MPa·m1/2. The highest indentation toughness is

22

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observed in sample B15 with 15 wt% of CaH2 addition (KIC = 7.8 MPa·m½, HV10 = 17.5

GPa).

0 2 4 6 8 10 12 14 160

2

4

6

8

10

12

14

16

18

20

22

0

2

4

6

8

10

12

14

Hardness A series B series C series

To

ug

hen

ess,

KIC

, MP

a.m

1/2

Har

dn

ess,

GP

a

Addition of CaH2, wt%

Toughness A series B series C series

Fig.4.4 Indentation hardness and toughness of the Ca--sialon ceramics as a function of CaH2

addition.

4.1.3 Compressive deformation

Results of high-temperature compressive deformation tests are shown in Fig.4.5. General

observations are that deformation onset temperatures increase with increasing Si/Al and N/O

ratios of starting compositions. Both for the -sialons and the -sialons, the increase of the

N/O ratio when going from compositions A (Si/Al = 3.03, N/O = 2.33) to B (Si/Al = 3.06,

N/O = 4.05) and finally C (Si/Al = 6.46, N/O =6.22 ) results in an increase of the deformation

onset temperature (Tonset, as representively shown in the insert of Fig.4.5 ) of 30–50oC.

However, the rate of deformation decreases upon an increase in the N/O ratio at a fixed Si/Al

ratio, as indicated by the curves for A0 and A11, or B0 and B11 in Fig.4.5, but there is no

significant accompanying effect on the deformation onset temperatures, especially not for the

samples in series C.

23

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1200 1300 1400 1500 1600 1700

0.0

0.1

0.2

0.3

0.4

0.5

0.6

1100 1200 1300 1400 1500 1600

0.000

0.001

0.002

0.003

ACH11

Tonset

1375 oC

- d

(

L /

L0)

/ d t

Temperature, oC

Temperature, OC

Sample Tonset

oC A0 1380 A11 1375 B0 1400 B11 1415 C0 1450 C11 1450

-

L /

L0

Fig.4.5 SPS compressive deformation behaviors of the SiAlON ceramics obtained under

constant load, indicating different deformation onset temperatures and rates. Representatively,

the deformation onset temperature of sample ACH11 is shown in the insert.

The Ca--sialon materials with high aspect ratio grains can be regarded as whisker-reinforced

ceramics. Increasing toughness with increasing N/O ratio at similar Si/Al levels, implies

formation of less strong Si-N and Al-N bonds in the forms of Si(N,O)4 or Al(N,O)4 tetrahedra

in the interfacial amorphous glass film, enhancing the interfacial debonding and thus

improving toughness. As showed in Fig.4.6, observations of interlocked elongated grains,

large elongated imprints, and protruding grains on the fracture surface suggest interfacial

debonding and grain pull-out from the matrix. Unlike the case of -sialon ceramics sintered

with extra oxides, the conservation and even slight increase in hardness with increasing CaH2

addition in present study, can be related to the presence of a nitrogen-rich grain boundary

glass phase, in which the substitution of nitrogen for oxygen, probably although not definitely

proven, leads to the increase of cross-linking that results in a more rigid glass network.

24

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Fig.4.6 Representative SEM images taken of nitrogen-rich Ca--sialon samples: (a) polished

surface of B13, and (b) fracture surface of B15, showing the interlocked microstructure and

grain pull-out effects toughening the Ca--sialon materials.

4.2 Synthesis of nitrogen-rich Ca--sialon

The intrinsic toughness of Si3N4 single crystals[61] is relatively low, about 1.9–2.8 MPa·m1/2,

compared with typical sintered Si3N4 ceramics, showing high toughness, about 7 MPa·m1/2 or

more,[74-76] which is attributed to the presence of elongated grains and grain boundary glass

phase favoring interfacial debonding of the reinforcing grains and the matrix. This indicates

that the observed high toughness of the Si3N4 based ceramics is not determined by Si3N4

itself but by the elongated grains and properties of the grain boundary glass phase.

The hardness of Si3N4 single crystals[61, 62] varies from 28 GPa to 35 GPa (HV0.3), depending

on the crystal orientations. The measured hardness of Si3N4 based ceramics is lower than

these values, ranging from 15 to 22 GPa, which indicates that, besides the phase composition

( phase or phase), density, amounts and properties of grain boundary glass phase and

secondary crystalline phases influence the hardness of Si3N4 based ceramics.

Thus, Si3N4 based ceramics that have a combination of hard phase with microstructure

consisting of elongated grains and very small amounts of nitrogen-rich grain boundary glass

phase can be expected to exhibit high enough high hardness and toughness to be suited for

most engineering applications. Starting from this observation, nitrogen-rich Ca--sialon

25

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ceramics were prepared and characterized with compositions (CaxSi12-2xAl2xN16, 0.2 x

2.6) along the Si3N4-1/2Ca3N2: 3AlN tie line. It was found that:

Pure and N-rich Ca--sialon ceramics were obtained for compositions with nominal

calcium contents 0.5 ≤ x ≤ 1.4.

Ca--sialon forms continuously in the compositional range 0 ≤ x ≤ 1.82 in the nitrogen-

rich region (CaxSi12-2xAl2xN16).

Self-reinforced microstructures were obtained for the nitrogen-rich Ca--sialon ceramics,

yielding a combination of high hardness (21 GPa) and fracture toughness

(~5.5 MPa·m½).

4.2.1 Phase compositions

Fig.4.7 shows the X-ray diffraction analysis results of the nitrogen-rich Ca--sialon ceramics

hot pressed at 1800oC, using Si3N4, AlN and CaH2 as precursors. Single phase nitrogen-rich

Ca--sialon ceramics are obtained for compositions with nominal calcium contents

0.5 ≤ x ≤ 1.4, whereas AlN and an additional phase, CaSiAlN3, are observed for higher x

values, implying that oxygen is preferentially incorporated in the -sialon phase, and the grain

boundary glass phase is rich in nitrogen.

26

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26 28 30 32 34 36 38 40 42 44

Si

AlN

CaSiAlN3

-Si3N

4

-sialon

2, degree

CC26

CC24

CC22

CC20

CC18

CC16

CC14

CC12

CC10

CC08

CC06

CC04

CC02

Fig.4.7 XRD patterns of CaxSi12-2xAl2xN16 (0.2 ≤ x ≤ 2.6) SiAlON ceramics with increasing

calcium contents

4.2.2 Solubility of calcium in -sialon, and lattice parameters

The unit cell parameters of the Ca--sialons increase with increasing x, reflecting the extent

to which longer Al-N bonds replace Si-N bonds, and also the amount of Ca incorporated in

the structure. As seen in Fig. 4.8, the x values obtained from EDX analyses and XRPD data

are in very good agreement for xnom up to 2.2. The results show that the Ca contents are less

than nominal. Approximately 15% of the added Ca is present in the grain boundary glass and

secondary phases like CaAlSiN3. The difference between nominal and determined x values

increases linearly with x up to xnom ca. 2.2. This implies that, to the extent that the fraction of

glassy phase can be approximated to be constant for the materials, the concentration of Ca in

the glass is proportional to the concentration of Ca, x, in the -sialon.

27

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0.0 0.4 0.8 1.2 1.6 2.0 2.4 2.8

0.0

0.4

0.8

1.2

1.6

2.0

+ from EDS measurement from XRPD

Nominal x

x fr

om

XR

PD

Fig.4.8 x values determined by EDX analysis and from XRPD data, plotted vs. the nominal x-

value.

As illustrated in Fig. 4.9, single-phase Ca--sialon ceramics are obtained for 0.50 ≤ x ≤ 1.38.

The upper solubility limit of x = 1.82 is attained for xnom = 2.2, beyond which the unit cell

parameters a and c remain comparatively constant at values of ca. 7.95 Å and 5.77 Å,

respectively. The variation of the lattice parameters in dependence on x, determined from

Rietveld refinement, is shown in Fig. 4.9. The lattice parameters are well described by the

linear relationships:

a = 7.7541(4) + 0.1114(9)x (Å) (1)

c = 5.6217(4) + 0.0859(9)x (Å) (2)

28

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0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.05.60

5.65

5.70

5.75

5.80

5.85

7.65

7.70

7.75

7.80

7.85

7.90

7.95

8.00

x from XRPD

a=7.7541(4)+0.1114(9)x

-Si3N

4

Ca1.8

Si8.2

Al3.7

N16

CaxSi

12-2xAl

2xN

16

c (Å

)

a (Å

)

c=5.6217(4)+0.0859(9)x

-Si3N

4

Ca1.8

Si8.2

Al3.7

N16

CaxSi

12-2xAl

2xN

16

Fig.4.9 Lattice parameters of nitrogen-rich -sialon as a function of actual x-value.

4.2.3 Microstructures and mechanical properties

Self-reinforced microstructures containing elongated grains were obtained for the nitrogen-

rich Ca--sialon ceramics, yielding a combination of high hardness (21 GPa) and high

fracture toughness (~5.5 MPa·m½), see Fig.4.10.

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 2.2 2.40

2

4

6

8

10

12

14

16

18

20

22

24

3.2

3.6

4.0

4.4

4.8

5.2

5.6

6.0

6.4

6.8

7.2

7.6

8.0

KIC

(M

Pa·

m1/

2 )

Hv1

0 (

GP

a)

X of Ca

Fig.4.10 Indentaiton toughness and hardness as a function of Ca content in the nitrogen rich

Ca--sialon ceramics.

Microstructure observations show that coarser grains are formed with increasing Ca content.

In all cases inter-granular fracture is predominant, with pull-out of elongated grains. The dual-

29

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phase sample CCH02, containing Ca-sialon and -sialon, exhibits a bimodal microstructure

with elongated -sialon grains embedded in a matrix of equiaxed smaller -sialon grains (Fig.

4.11a). Single-phase, or almost single-phase, -sialon ceramics with low calcium contents,

e.g. CCH04, consist of very fine equiaxed grains (Fig. 4.11b). The amount and aspect ratio of

elongated -sialon grains increase with increasing x value. Grains with aspect ratios of six to

eight are observed in samples with high calcium contents.

Fig.4.11 SEM secondary electron images of fracture surfaces of samples (a) CCH02, (b)

CCH04, (c) CCH14, (d) CCH20,.

The TEM image of CCH20 shown in Fig. 4.12a reveals the presence of a very limited amount

of grain-boundary phase. The elongated grain morphology, for almost every grain, is well

illustrated by the corresponding SEM image, Fig. 4.12b, of a polished and subsequently

chemically etched surface.

30

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Fig.4.12 TEM image (a) and SEM image (b) of a polished and chemically etched surface of

sample CCH20, showing the formation of elongated grains and very limited amount of grain

boundary glassy phase.

4.3 Superplastic deformation of nitrogen-rich Ca--sialon

The high-temperature properties of Si3N4 based ceramics are largely dependent on the

amounts and properties of the secondary crystalline phase and the residual grain-boundary

phase that form during the sintering process. The presence of residual grain-boundary phases

may result in a deterioration of high temperature properties, especially when the glassy phase

has a relatively low glass transition temperature and a low viscosity. On the other hand, to

preserve the properties after deformation, or to prevent grain-boundary cracking during

superplastic deformation, a high cohesive strength at grain boundaries is crucial, and dynamic

grain-growth during large strain deformation should be avoided. In the present compressive

deformation, nitrogen-rich Ca--sialon ceramics show characters like:

Increased onset deformation temperatures, about 150oC higher than that of O-rich Ca-

-sialon (~ 1370oC);

Enhanced toughness (KIC) and preserved high hardness even after the deformation;

Crack-free superplastic deformation behavior at a strain rate of 2×10-3 s-1 accompanied

by considerable grain growth.

31

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4.3.1 Compressive deformation behavior

As shown in Fig.4.13, the deformation onset temperatures of the nitrogen-rich Ca--sialon

ceramics (~1520oC) are ca. 150oC higher than those of the oxygen-rich Ca--sialon ceramics

(~1370oC), and also higher than those of (Y,Yb)--sialon (1350–1375oC) with nano sized

microstructures[73] and Li--sialon (1300–1380oC)[77] with fine-grained microstructures. The

maximum compressive deformation rate of the nitrogen-rich Ca--sialon ceramics, 2.010-3 s-

1, is slightly lower than that of the oxygen-rich ones, 3.210-3 s-1.

1300 1400 1500 1600 1700 1800 1900

0.0

0.1

0.2

0.3

0.4

0.5

0.6 CCO14 CCO16 CCH04 CCH08 CCH12 CCH16 CCH20

Temperature,oC

-

L/L

0

N-rich Ca--sialonO-rich Ca--sialon

Fig.4.13 Compressive deformation behaviors of the Ca--sialon ceramics conducted in SPS at

heating rate of 40oC /min, and a constant load corresponding to an initial pressure of 40 MPa.

4.3.2 Microstructures and properties

A nitrogen-rich liquid phase can be expected to wet the -sialon grains better than one with

high oxygen content, since the former has a composition, especially a nitrogen content, close

to the compositions of the -sialon grains. Upon the formation of a nitrogen-rich liquid during

deformation, the good wetting of the -sialon grains guarantees fast deformation without

cracking via grain sliding or rotation. The glassy grain boundary phase with high nitrogen

content allows the mechanical properties of the nitrogen-rich Ca--sialon ceramics to be

32

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maintained at high temperatures, and has made it possible to make a material combining

excellent properties (HV10 ~21 GPa, KIC ~ 7 MPa·m1/2) and complex shapes that can be

obtained after proper compressive deformation.

Appreciable grain growth was found to have occurred during SPS deformation, resulting in an

amplified anisotropic microstructure (Figs. 4.14b, d and f) and a development of coarser and

more elongated grains. Compared with isothermal deformation, high hardness (HV10 = 18–

20 GPa) and toughness (KIC = 4–7 MPa·m1/2) of the nitrogen-rich Ca--sialon are maintained

after the isostatic pressure deformation. The anisotropic distribution of hardness and

toughness (Fig.4.15) is well explained by the observation of microstructures showing

anisotropic grain growth along the direction normal to the applied pressure (Fig.4.16 ).

33

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Fig.4.14 Representative SEM micrographs of the prepared ceramics before and after the

compressive deformation: CCH04 (a, HP1800oC×4h, b, SPS 1750oC×0min), CCH16 (c,

HP1800oC×4h, d, SPS 1750oC×0min) and CCO14 (e, HP1800oC×4h, f, SPS 1600oC×0min).

3

4

5

6

7

8

9

10

11

12

0.0 0.4 0.8 1.2 1.6 2.00

4

8

12

16

20

24

28

x of CaxSi

12-2xAl

2xN

16

KIC

(M

Pa·

m1/

2 )

HV

10 (G

Pa)

before after H

V10

KIC

Fig.4.15 Indentation hardness and toughness of the nitrogen-rich Ca--sialon ceramics

measured before and after compressive deformation, with the inset showing light-optical

micrographs of the cross sections of selected Ca--sialon ceramics.

Fig.4.16 SEM micrographs of specimen CCH16 after the compressive deformation (SPS at

1750oC×0min): (a) normal to the direction of pressure, (b) parallel to the direction of pressure.

It can be concluded that high nitrogen content as well as a reduced amount of residual glass

phase contribute to the increase of deformation onset temperature. N/O analysis and EDS

analysis indicate that the nitrogen-rich Ca--sialon has less than 2.2 at% overall oxygen

34

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content, and more than 85% of the nominal calcium content is incorporated into the -sialon

structure. TEM observations on the nitrogen-rich Ca--sialon with high x value reveal that the

glass phase is predominantly present as a thin film surrounding the grains (Fig.4.17).

Corresponding EELS analysis at triple-grain pockets indicates that the glassy phase has a high

N/O ratio, which furthermore increases with increasing x value.

Fig.4.17 TEM micrograph of the nitrogen-rich Ca--sialon CCH20 (x = 2.0, hot pressed at

1800oC for 4h), showing the presence of residual glassy phase at some triple grain pockets.

Corresponding EELS analyses indicate very low oxygen content in the grains but relatively

high oxygen and calcium contents at triple-grain junctions (arrowed).

4.4 Thermal properties of nitrogen-rich Ca--sialon

As mentioned above, the nitrogen-rich Ca--sialons (CaxSi12-2xAl2xN16 ) ceramics exhibit

improved toughness that in turn can be attributed to the formation of elongated -sialon grains. In

addition, increased resistance to high temperature deformation was also observed in nitrogen-rich

Ca--sialons, i. e. nitrogen-rich Ca--sialons exhibited about 150 oC higher deformation onset

temperature than those of their oxygen-rich counterparts. In this part, thermal properties of

nitrogen-rich C--sialons are investigated with respect to reaction mechanism, phase stability

and oxidation resistance with the aim to elucidate their high temperature properties. The results

indicated that:

35

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The decomposition of CaH2 takes place in the temperature region 200-300 oC and is

followed by a significant nitridation within the temperature range 600-800 oC. -sialon

phase is first observed at 1400 oC and monophasic Ca--sialon ceramics was prepared at

1800 oC.

Post heat treatment of the nitrogen-rich Ca--sialons in the temperature range1400 –

1600 oC revealed that these SiAlONs are stable, i. e. no →-sialon transformation was

fond.

The nitrogen-rich Ca--sialons are less resistant to oxidation, while the mixed -

sialon (low Ca-content) shows better oxidation resistance than pure -sialon at low

temperatures (1250-1325 oC).

4.4.1 Reaction sequence and phase evolution

Representatively shown in the TG-curve of ball-milled powders of the composition CCH16 in

Fig.4.18, most of the weight loss occurred in the temperature region 200-300 oC, which is in

agreement with the observed decomposition temperature of CaH2 in various CaH2 mixtures [78 -

80]. The weight loss is almost completed at 300 oC. The observed weight loss (0.38 % for CCH16)

is less than the calculated one (0.53%), indicating that CaH2 is partly decomposed in connection

with the ball milling procedure and/or when the mixed powders were vacuum- dried. Starting

from 600 oC, a significant weight gain was observed, and ended at T> 800 oC. Assuming that the

weight gain can be ascribed to the reaction:

3Ca(s) + N2(g) = Ca3N2(s)

The theoretical weight gain amounts to 1.9 wt% compared with the observed 1.5wt%. The

difference between observed and calculated weight gain can be ascribed to the observation

that Ca also forms other compounds, see in Fig.4.19, in which, crystalline phases like CaSiN2,

and CaSiN4 phases are observed below 1450 oC; and monophasic Ca--sialon ceramics are

prepared at 1800 oC. But there is no observation of intermediate phases like gehlenite

(CaAl2SiO7) and/or AlN polytypiods which appeared in preparation of Ca--sialon using

CaO as a starting powder.[81]

36

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0 200 400 600 800 1000 1200-0.50

-0.25

0.00

0.25

0.50

0.75

1.00

1.25

1.50

1.75

Temperature, oC

wei

gh

t g

ain

, wt

%

Fig.4.18. TG curve of the powder mixture CCH16 recorded in flowing N2. Weight loss occur

around 200 oC, and a significant weight gains are observed between temperatures 600-800 oC.

20 22 24 26 28 30 32 34 36 38 40 42 44

: C

aSiN2

aaa

aabbb

aa

600oC

1800oC

1500oC

1450oC

1400oC

1300oC

1200oC

1000oC

800oC

Two theta, degree

bb

a

a

aa

a

aaa

aa

aa

bbb

aa

aa

bbb

a a a a

a

a a aa

aa

a aa aa

a

a

a

aa

a

b

x

xxx

aa

aa

aaa

a

aaa a

a

x

aa

A

A

AA A

A

A

A

A

A A

A

A

AAA

Aaa

A

aa

A

a AA

unknown phasea4SiN

4

Si

aa

aaaaa a a

aa

aaa

a

aa AA

A: AlN

a

-sialona: -Si3N

4

b: -Si3N

4

Si

A

xxx

x:: C

Fig.4.19 XRPD patterns of the sample CCH16 hot pressed at different temperatures.

37

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4.4.2 Thermal stability

It has been shown that the thermal stability of Ca--sialons prepared with oxides additives is

superior to the ones of rear earth stabilized SiAlONs[82,83] and so are rear earth stabilized SiAlONs

co-doped with Ca.[84]In this part, the thermal stability of the nitrogen-rich Ca--sialons has been

further studied. Hot pressed compacts of the compositions CCH02 (an -sialon), CCH04 (an -

sialon containing traces of -sialon) and CCH16 (monophasic -sialon) have been post-heat

treated in the temperature region 1400-1600 oC for 24 hours. The X-ray powder patterns of the

post heat treated samples are given in Fig.4.20 below. Although very slight variation on lattice

parameters were observed, but it can be concluded that there is no significant difference between

the X-ray powder patterns of the sintered and post heat treated samples, i. e. no to phase

transformation takes place in these samples. In contrast, most -sialons stabilized by rear earth

cations are prone to transform into -sialon and grain boundary phases in the temperature region

1300-1600 oC.[85-88] This finding, together with the finding by Hewett[82] and Mandal et al.,[83] that

oxygen-rich Ca--sialons are stable even when sintered with excess glass phase, indicate that

calcium stabilized -sialons is probably the most thermal stable sialon phase over a wide

composition and temperature range. Low oxygen contents of starting compositions often yield

less amounts of grain boundary glass phase in the sintered body, and the formed grain boundary

phase has due to its higher nitrogen content higher viscosity than an oxygen rich grain boundary

phase. In the former case the to phase transformation is expected to be kinetically retarded.

38

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28 30 32 34 36 38

a -sialon

-sialon

Si

1600 oC

1550 oC

1450 oC

1400 oC

2 Theta,degree

CCH02 HPed

28 30 32 34 36 38

b

-sialon-sialon

Si

CCH04 HPed

1600 oC

1550 oC

1450 oC

1400 oC

2 Theta,degree

28 30 32 34 36 38

c -sialon

Si

1600 oC

1550 oC

1450 oC

1400 oC

CCH16 HPed

2 Theta,degree

Fig.4.20 X-ray powder diffraction analysis of samples CCH02 (an -sialon), CCH04 (an -

sialon containing traces of -sialon) and CCH16 (an -sialon) post heat treated for 24 h at

temperatures ranging from 1400 to 1600 oC for 24.

Low and high resolution transmission electron microscopy studies focused on grain boundaries

and triple-grain pockets provides us with some additional information about the glassy grain

boundary phase. Thus direct bonding of grains is observed in sample CCH04, which has a starting

compositions located very near to the -sialon phase border, as seen in Fig.4.21a and Fig.4.21b,

but grains separated by a thin film seems also to be present, see Fig.4.21a. In sample CCH16

(Fig.4.21c and Fig.4.21d), whose composition is close to the other end of Ca--sialon phase

region, a thin glass film with thickness about 0.9 nm is observed, but it is hard to estimate the

dimension of glass phase at the triple-grain pockets. The observations seems to indicate that these

samples contain only a small amount of grain boundary phase which in turn might explain the no

→-sialon transformation takes place in these samples.

39

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~0.9 nm

Fig.4.21. TEM and corresponding HRTEM images for nitrogen-rich Ca--sialons hot pressed at

1800 oC for 4 hours: CCH04 (a) and (b); CCH16 (c) and (d).

4.4.4 Oxidation resistance

In principle, the nitrogen-rich -sialon compositions are thermally less stable in presence of

oxygen, i. e. more prone to be oxidized, than their oxygen rich counterparts. Initially a thin

oxygen rich layer is formed and as the oxidation proceeds, the thickness of the product layer

increases, and the diffusion of the reactants and product gases through the product layer is

retarded. The shape of the recorded oxidations curves at 1250 and 1325 oC (see Fig.4.22) are

generally speaking of the parabolic type while at 1400 oC deviation from the parabolic type of

occurs, seemingly a consequence of formation of bubbles in larger quantities, located between the

sialon matrix and the outer part of the oxide scale, see in Fig. 4.23.

40

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0 20000 40000 60000 800000.00

0.02

0.04

0.06

0.08

0.10a1250 oC

CCH02 CCH04 CCH08 CCH12

W

/A0,

mg

.mm

-2

t, s0 20000 40000 60000 80000

0.00

0.02

0.04

0.06

0.08

0.10

b1325 oC

CCH02 CCH04 CCH08 CCH12

W

/A0,

mg

.mm

-2

t, s

0 20000 40000 60000 800000.00

0.02

0.04

0.06

0.08

0.10C1400 oC

CCH02 CCH04 CCH08 CCH12

W/A

0, m

g.m

m-2

t, s

Fig.4.22 Weight gains as a function of time of the nitrogen-rich Ca--sialons oxidized at 1250 oC

(a), 1325 oC (b) and 1400 oC (c) for 20 hours in flowing oxygen.

Fig.4.23 An SEM image shows the formation of bubbles beneath the glass film in sample CCH08

oxidized at 1250 oC.

41

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The oxidation products, identified via their X-ray powder patterns, are listed in Table 4.1

Cristobalite (SiO2) and wollastonite (CaSiO3) are the main crystalline phases detected in the

surface oxide scale. Very small amount of anorthite (CaAl2Si2O8) phase was detected in samples

oxidized at 1250 oC.

Table 4.1 Phases in oxidation scales identified via their X-ray powder patterns¶

sample CCH02 CCH04 CCH08 CCH12

As sintered at

1800 oC

-sialon (62 mol/%

-sialon (38 mol/%)

-sialon (97 mol/%)

-sialon (3 mol/%) -sialon (100 mol/%) -sialon (100 mol/%)

Oxidized at

1250 oC

SiO2, s

CaSiO3, s

CaAl2Si2O8,w

SiO2, s

CaSiO3,w

CaAl2Si2O8,w

SiO2, s

CaSiO3,w

CaAl2Si2O8,vw

SiO2,m

CaSiO3,w

Oxidized at

1325 oC

SiO2, s

CaSiO3, m

SiO2, s

CaSiO3, vw SiO2, w SiO2,w

Oxidized at

1400 oC

SiO2, s

CaSiO3, w SiO2, s SiO2,vw Amorphous

¶ Strength of reflections in XRPD pattern: s-strong; m- medium; w- weak; vw- very weak

Crystalline CaSiO3 and SiO2 phase is present in the oxide scale for all samples oxidized at

lower temperatures, but in decreasing amount with increasing Ca content. A Ca-Si-Al-O glass is

formed at all temperatures and the amount of crystalline phases in the oxide scale decreases with

increasing Ca-content. With increasing temperature and calcium content, both the amounts of

glass phase and number of bubbles increased.

The oxidation of Ca-sialon ceramics involves concurrently ongoing inward diffusion of oxygen

and outward diffusion of metal cations and nitrogen product, resulting in compositional gradients.

This is evident when EDS mapping technique is applied to cross section region obtained by ion-

polishing techniques. As shown in Fig.4.24, calcium and oxygen are enriched in the surface area,

and gradually decrease inward, while silicon and nitrogen exhibit a positive elemental

distribution, together with a homogenouse distribution of aluminium over the whole area. The

gradient distribution of elements at the reaction region suggest that the diffusion of calcium and

oxygen through the oxide layer and glass film is a rate controlling step in the oxidation of

nitrogen-rich Ca-sialon ceramics.

42

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Fig.4.24. EDS mapping analysis focused on cross section region in sample CCH02 oxidized at

1250 oC. Note the depletion of calcium, and bubbles in the oxides scale (arrowed), but enrichment

of calcium and oxygen at the surface.

43

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5. SUMMARY

The nitrogen-rich liquid phase sintering concept is introduced by using CaH2 as a sintering

aid, so as to vary the N/O ratio of the liquid phase formed in the sintering process while

keeping the Si/Al ratios constant. With increasing addition the phase contents change from

single -sialon to dual -sialons and to single Ca--sialon. The microstructures contain

mainly equi-axed -sialon grains at low N/O ratios, and well faceted elongated Ca--sialon

grains at high N/O ratios. The improved toughness (KIC = 7.8 MPa·m½) and hardness (HV10 =

17.5 GPa) properties can be attributed to the formation of interlocked microstructures. High-

temperature compressive deformation tests indicate that the deformation onset temperature is

determined mainly by the Si/Al and N/O ratios, whereas the deformation rate is affected by

the microstructure.

Nitrogen-rich Ca--sialon ceramics, with compositions CaxSi12-2xAl2xN16, (0.2 x 2.6)

along the Si3N4-1/2Ca3N2:3AlN tie line, form continuously within the compositional range

x = 0 to at least x = 1.82. An empirical relationship is proposed, relating Ca cation solubility

to the unit cell expansion in the -sialon structure. The obtained materials demonstrate a

combination of excellent mechanical properties with high deformation onset temperatures,

due to the formation of self-reinforced microstructures and the nature of the low amount of

nitrogen-rich grain boundary glass phase.

The decomposition of CaH2 takes place in the temperature range 200-300 oC and is followed by a

significant nitridation within the temperature range 600-800 oC. -sialon phase is first observed at

1400 oC and monophasic Ca--sialon ceramics is prepared at 1800 oC. Post heat treatment of the

nitrogen-rich Ca--sialons in the temperature range1400-1600 oC reveals that these SiAlONs are

stable. The nitrogen-rich Ca--sialons are less resistant to oxidation, while the mixed -sialon

(low Ca-content) shows better oxidation resistance than pure -sialon at low temperatures (1250-

1325 oC).

44

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6. FUTURE WORK

This is a rarely touched corner in the Si3N4 system. It will be interesting to explore the family

of nitrogen-rich SiAlONs doped with rear earth elements, for example nitrogen-rich -sialons

stabilized by Y, Yb, and Nb. It will be interesting also to apply the concept of nitrogen-rich

liquid phase sintering to other systems such as AlN, BN, or SiC that are chemically and

thermal dynamically compatible with nitrogen.

Most of the past research on SiAlONs has focused on dense materials potentially for high

temperature applications. But for nitrogen-rich SiAlONs, the poor oxidation resistance might

be an obstacle. However, synthesis of nitrogen-rich powders, co-doped with functional

elements might be of interest for low temperature, functional applications.

To correlate the properties to microstructures, techniques such as SPS deformation, HREM

observations, EDS and EELS analysis, and overall N/O ratio assessment by combustion

method, have been applied. However, it should be pointed out that, it is still an issue regarding

the nature of grain boundary glass phase, especially its N/O ratio and elemental distribution

within grain boundaries in the nitrogen-rich Sialon ceramics. In cooperation with Institut de

Physique de la Matière Complexe, Switzerland, mechanical loss measurement is on going to

try to reveal the effects of grain boundary chemistry and N/O ratio on high temperature creep

behaviors of the nitrogen-rich Ca--sialon, and RE--sialons ( RE=Y and Yb). Here we just

present the most recent results for the nitrogen-rich Ca--sialon measured by mechanical

spectroscopy.

45

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Mechanical spectroscopy[89, 90] studies the absorption spectra of mechanical energy under the

conditions of applied periodic external mechanical field. In this method, the internal friction

Q-1, Q-1 = Tan (δ), where δ is the phase angle between the stress and the strain, can be used to

access the deformation mechanisms, since the appearance of characteristic internal friction

peaks, as a function of temperature or frequency, is connected to the anelastic movement of

defects in materials. Damping backgrounds and peak Q-1 are correlated to the refractoriness of

the glass phase or the softness of the materials. High values of background and large Q-1

values at specific temperatures reflect the creep resistance of studied materials

In present study, the mechanical spectroscopy is used to associate the deformation behavior of

Ca--sialon ceramics to their microstructure changes after thermo treatment, particularly to

link the grain boundary sliding with intergranular grain films or amorphous at triple-grain

pockets.

0 200 400 600 800 1000 1200 1400

0.80

0.84

0.88

0.92

0.96

1.00

1.04

0 200 400 600 800 1000 1200 1400

0.00

0.01

0.02

0.03

0.04

0.05

0.80

0.84

0.88

0.92

0.96

1.00

1.04

909 oC E (

T)

/ E (

25oC

)

heating IF

Inte

rnal

Fri

ctio

n, Q

-1

Temperature, oC

845 oC

1096 oC

1010 oC

heating stifness

cooling IF

cooling stifness

Fig. 6.1 Typical internal friction (IF) and shear modulus spectrum for nitrogen-rich Ca--

sialon. Note the disappearance of IF peak and decrease of relaxation background on cooling

procedure.

The deformation behaviours of the nitrogen-rich Ca--sialons are quite similar. Fig. 6.1

shows a typical creep behavior for the nitrogen-rich Ca--sialon (CCH02) measured by

mechanical spectroscopy. During the heating cycle, the characteristic IF peak appears around

909 oC, accompained by a decrease in stiffness (shear modulus). At 1330 oC, a relative

stiffness above 94% is still maintained, implying that the nitrogen-rich Ca--sialon is

46

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resistance to creep deformation. The IF peak can be attributed to the glass transition of the

amorphous phase located at triple-grains pockets (see in Fig. 6.2a). In which, the -sialons

grains are surrounded by a very thin glass film (with thickness less then 1 nm).

During cooling cycle, an irreversible disappearance of internal friction peak is observed and

the background is low. This implies that the amorphous phase in the nitrogen-rich Ca--sialon

is not stable, and tends to be crystallized upon heat treatment. As clearly shown in HREM

image of sample CCH04 obtained after SPS deformation (Fig. 6.2b), the re-crystallization of

amorphous phase resulted in modulus increase and direct bonding between grains, resulting in

a glass free Ca--sialon.

Fig. 6.2 HREM image of sample CCH04: (a) before SPS deformation, (b) after SPS

deformation. Note the grain boundary glass at triple-grain pocket in the hot pressed sample,

but the direct bonding of grains in sample after SPS deformation.

Effects of N/O ratio and amounts of grain boundary glass phase, as well as grains size on

creep behaviour can also be drawn from mechanical spectroscopy. Details of this part,

together with spectra for nitrogen-rich RE-sialons, will be addressed in a separated paper.

47

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ACKNOWLEDGMENT

With the finish of the main text, I am very grateful to all the people who have helped in the

past years. Sincere thanks go to:

Dr. Saeid Esmaeilzadeh, my supervisor, young and creative, who has given me the

opportunity to make some contribution to this fundamental area. Without his help and support,

this thesis would not exist.

Dr. James Shen, my co-supervisor, for his recommendation, encouragement, detailed

discussions, constructive suggestions, and providing academic inspiration as well as

enthusiasm for doing this work.

Dr. Jekabs Grin, for his stream-like knowledge flow. If you like, you can always get right

response and learn something new from a discussion with him. He was not my listed

supervisor, but actually he has been.

Dr. Kjell Jansson, for his detailed technical instructions and patient demonstrations whenever

I encountered problems in the microstructure observation.

Dr. Zhe Zhao, for fruitful discussions, suggestions, sharing of experimental staffs and skills,

also sharing the great time out of work.

Professor Mats Nygren, for great help in thermal properties analysis and instruction in SPS

techniques.

Dr. Thomas Höche of Leibniz-Institut für Oberflächenmodifizierung e. V. for help with TEM

observation and EELS analysis, and hard efforts paid on sample preparation.

Dr. Daniele Mari of Institut de Physique de la Matière Complexe,EPFL, Switzerland for help

in high temperature mechanical loss measurement.

Thanks to all staffs and researchers at the department for helping and supporting me:

Hans-Erik Ekström and Per Jansson for constructing, service and reparation of scientific

equipments. Specials thanks go to Hans-Erik for reparation of hydraulic pump and assembling

of HP furnace. Jaja Östberg for carbon coating of SEM specimens and kindly sharing of her

TEM specimen box. Lars Göthe, for recording and developing lots of Guinier Hägg films.

48

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Per-Erik Persson and Rolf Eriksson for help with computer and network problems. Ann-Britt

Rönell, Eva Pettersson, Hillevi Isaksson,Agnes Laurin for always being so helpful.

Thanks to Professor Sven Lidin and Professor Lennart Bergström for project consultation and

creating such a wonderful working environment. Thanks to Professors Xiaodong Zou, Sven

Hovmöller, Margareta Sundberg and Osamu Terasaki, associate professor Mats Johnsson,

Lars Eriksson, Dr.Yasuhiro Sakamoto for the interesting courses you give.

Thanks to Ashkan Pouya, Katarina Flodstrom, Bahman Etemad of Diamorph AB for the

fruitful discussions and great cooperation.

Thanks to my colleagues Mirva Eriksson ( my first aid on puzzles about interesting

documents, which unfortunately were written in Swedish ), Ali Sharafat and Abbas Haakem

(我们的巴基斯坦兄弟, Assalam-o-alaikum!),Dr. Daniel Grüner (我们的德国兄弟) Jovice

BoonSing Ng, Linnea Andersson, Ehsan Jalilian, Petr Vasiliev, and Bertrand Faure for being

around when needed and being good lunch-companions.

Special thanks to Tuping Zhou and your family for kind help. You are welcome anytime!

Thanks also go to Daliang Zhang for the great help in TEM and HRTEM observations. Wish

success to you and your family!

Thanks to all my Chinese friends. I have been here with you all! I must list here for sake of

remember and memory: Hong Peng, Liqiu Tang, Jing Liu, Shuying Piao, Juanfang Ruan,

Nanjiang Shu, Changming Xu, Junliang Sun, Mingrun Li, Huijuan Yue, Shiliang Huang,

Changhong Xiao, Yao Cheng, Yan Xiong, Cong Lin, Guanghua Liu, Dong Zhang, Guoying

Zhao, Jie Xiao, Yihong Liu, Shuai Li, Liang Ran, Bing Tang, Yi Sun et al. It’s my pleasure to

have been accompanied with you all!

Thanks also go to badminton team members, for the pleasure time we have been together!

Finally, I would like thank my wife Qian Dong, my parents, relatives and friends in China, for

their endless support, encouragement and love. And, to my son Yingnan, who has created so

much funny time and so many precious memories for us.

Wish you all good luck!

49

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