synthesis and characterization of nitrogen-rich - diva portal
TRANSCRIPT
Synthesis and Characterization of Nitrogen-rich
Calcium -Sialon Ceramics
Yanbing Cai
蔡 雁 兵
Department of Physical, Inorganic, and Structural Chemistry
Stockholm University
2009
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
To my family, for their love and support!
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
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
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
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
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.
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
-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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
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
~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
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
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
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
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
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
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
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
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
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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