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Accepted Manuscript
Investigation of the effect of Al2O3–Y2O3–CaO (AYC) additives
on sinterability, microstructure and mechanical properties of SiC
matrix composites: A review
Mahdi Khodaei, Omid Yaghobizadeh, Alireza Alipour Shahraki,
Sadeq Esmaeeli
PII:
DOI:
Reference:
S0263-4368(18)30449-9
doi:10.1016/j.ijrmhm.2018.08.008
RMHM 4774
To appear in:
International Journal of Refractory Metals and Hard Materials
Received date:
Revised date:
Accepted date:
9 July 2018
13 August 2018
18 August 2018
Please cite this article as: Mahdi Khodaei, Omid Yaghobizadeh, Alireza Alipour Shahraki,
Sadeq Esmaeeli , Investigation of the effect of Al2O3–Y2O3–CaO (AYC) additives
on sinterability, microstructure and mechanical properties of SiC matrix composites: A
review. Rmhm (2018), doi:10.1016/j.ijrmhm.2018.08.008
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Investigation of the effect of Al2O3–Y2O3–CaO (AYC) additives on
sinterability, microstructure and mechanical properties of SiC matrix
composites: A Review
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Mahdi Khodaeia, 1 , Omid Yaghobizadeh , Alireza Alipour Shahrakia, Sadeq Esmaeelic
Composite Materials & Technology Center, Malek Ashtar University of Technology , Tehran, Iran.
Department of Materials Engineering, Imam Khomeini International University (IKIU), Qazvin, Iran.
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Department of Ceramic, Shahreza Branch, Islamic Azad University, Shahreza, Iran.
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Abstract:
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Appropriate properties of SiC ceramic such as high hardness, low density, high melting point
and high elastic modulus make this material as a favorite candidate for different industrial
applications. Although some disadvantages including high sintering temperature, low
sinterability, and low fracture toughness have restricted the use of this material, previous studies
showed that using Al2 O3 -Y2 O3 additives plays an effective role in the improvement of
sinterability as well as the enhancement of the properties of these composites. Moreover, the
addition of CaO results in the acceleration of the formation of molten phase and the
improvement of sinterability. In addition, the use of these additives cause the formation of the
intermetallic phases of Al5 Y3 O12 (YAG) and CaY2 O4 and by activating the mechanisms of crack
deflection, crack bridging, phase transformation, strengthening the grain boundary and changing
the fracture mode from intergranular to transgranular results in improved mechanical properties.
This paper attempts to investigate the effect of using Al2 O3 –Y2 O3 –CaO (AYC) additives on
sinterability, microstructure, and mechanical properties of SiC matrix composites including the
composites reinforced with SiC fibers and SiC matrix nano-composites. Finally, the effect of the
post-sintering annealing process under two conditions i.e., with and without applying pressure
(pressureless sintering) on microstructure and mechanical properties has been studied.
Keywords: SiC-Al2 O3 –Y2 O3 –CaO; Fracture toughness; Flexural strength; Hardness; Microstructure;
Crack bridging; Crack deflection.
1. Introduction
Silicon carbide has been widely studied over the past few years and is a promising material for
high-temperature engineering applications because of its favorable properties such as high elastic
1
Corresponding author,
E-mail address: mahdi.khodaei01@gmail.co m
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modulus, superior hardness, good corrosion resistance, low thermal conductivity, low density
and low thermal expansion coefficient [1, 2]. It has found extensive applications including diesel
motor parts, gas turbines, industrial heat exchangers, high-temperature energy exchanger
systems, hot gas filters, anti-abrasion parts in various atmospheres, medical implants and optical
mirrors [3, 4].
This ceramic has two different crystalline structures including cubic β-SiC and α–SiC. The
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latter has a specific type of one-dimensional polymorphism called Polytypism. Polytypes are
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similar in two dimensions of the closed packed planes but differ in order of sequence in the
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dimension perpendicular to those planes. The polytype 3C (cubic) with a zinc-blend
crystallographic structure is known as β-SiC. All non-cubic polytypes are known as α–SiC. 6H,
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4H and 2H are the most common α–SiC polytypes with a hexagonal symmetry. Instable β–SiC
transforms to one of α–SiC polytypes at high temperatures and the β to α phase transformation is
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highly a function of thermal history, starting materials, impurities, atmosphere and pressure [511]. For example, the addition of a little amount of Al, B and C to β-SiC usually results in a
transformation to the 4H polytype while the addition of some B and C to β-SiC leads to a
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transformation to the 6H polytype. Moreover, β-SiC preferably transforms to the 2H polytype in
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the presence of AlN [12, 13]. The atmosphere of the sintering process also affects β to α phase
transformation and the grain growth behavior severely [7, 14, 15]. For instance, although argon
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atmosphere results in an improvement in the phase transformation and forms elongated grains,
nitrogen atmosphere inhibits β to α transformation and grain growth. Nader et al. [16] have also
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reported that the β to α transformation depends the initial amount of β and the sintering
atmosphere. The phase transformation rate decreases with an increase in the initial amount of β
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in the starting powder during sintering in the presence of nitrogen. Furthermore, it has been
demonstrated that the crystalline structure of the starting powder is not of particular importance
in terms of compaction and the β to α transformation does not improve the compaction of SiC
[17].
Nevertheless, the resultant microstructure depends on the type of starting powder. Equiaxed
grains can be obtained from the powders with a high portion of α–SiC while elongated grains can
be formed from β-SiC or β-SiC containing the nucleating particles of α-SiC [17, 18]. The
materials with a high fracture toughness occasionally show a microstructure with large elongated
grains [6, 19, 20]. Large elongated grains result in an increase in the fracture toughness in SiC
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through crack bridging [18, 21] or crack deflection [20, 22-24]. As it was mentioned earlier,
annealing process also plays an important role in controlling the phase transformation.
According to the reports, annealing along with applying a uniaxial compression (25 MP) greatly
delays the β to α transformation and prevents grain growth [25, 26]. Previous studies also
demonstrated that the contribution of β to α transformation is less than 30 % in the annealed
materials as a result of applying compression and the microstructure contains fine and equiaxed
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grains while the aforementioned contribution becomes more than 30 % in the samples annealed
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without applying compression and elongated grains form in the microstructure of these samples
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[27].
Sintering of SiC can be done using both liquid phase and solid state methods. Solid state
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sintering needs some additives to change the energy level of SiC in order to be able to become
compacted. Even with additives and at high temperatures it is still difficult to compact SiC
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because of its predominant covalent bonds and a low self-diffusion coefficient. Such
characteristics inhibit the mass transfer mechanisms which are responsible for compactability
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[28-30].
Carbon, boron and aluminum are the usual additives added to SiC during solid state sintering
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[1, 4]. Liquid phase sintering is the most common method for manufacturing SiC ceramics. The
presence of the liquid phase accelerates the mass transfer and decreases the sintering temperature
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and the sintering time and as a result, a fine-grained and more uniform microstructure with better
mechanical properties and particularly fracture toughness is obtained. Final microstructure
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depends on multiple factors including α-SiC and β-SiC proportion, amount and the type of
additives as well as the sintering temperature and time [19, 29, 31, 32]. Liquid phase sintering
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also depends on the composition of the starting materials, type and the amount of additives along
with the sintering temperature and time [19, 29, 31, 32]. Y2 O 3, Al2 O3 -Yb2 O 3 , Al2 O 3 -La2 O3 ,
Al2 O 3 -Dy2 O 3 and Al2 O3 -Y2 O3 -CaO are the most common additives used for this purpose [4].
The use of Alumina (Al2 O3 ), Yttrium oxide (Y2 O3 ) and Calcium oxide (CaO) as the additives in
the sintering of SiC is important in terms of lowering the sintering temperature, increasing the
toughness, helping the formation of liquid phase and improving the sinterability. These additives
increase the mechanical properties and improve the microstructure through forming the
intermetallic phases of Al5 Y3 O12 (YAG) and CaY2 O4 . In addition to the aforementioned
advantages, Al2 O3 , Y2 O3 and CaO improve the mechanical properties as well as the
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microstructure of the samples compared to the non-containing additive samples through the
mechanisms such as crack deflection, crack bridging, phase transformation, strengthening the
grain boundary and changing the fracture mode from intergranular to transgranular [3].
On the other hand, Al2 O3 and Y2 O3 are the most common additives that can react with the
SiO2 present on the surface of the SiC powders and increase the sinterability with the formation
of a eutectic composition. CaO is added with the purpose of lowering the vapor pressure of the
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growth rate and weight loss during the sintering process [33].
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liquid phase and decreasing the sintering temperature. This can lead to a decrease in the grain
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Due to the importance of using CaO additive along with Al2 O3 and Y2 O 3 , the most important
studies on the effect of CaO on the sinterability and the properties of SiC-matrix composites
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including the composites fabricated by pressureless sintering and hot pressing, the composites
reinforced with SiC fibers, and SiC matrix nano-composites have been reviewed. In the end, the
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papers on the effect of the post-sintering annealing process have been reviewed.
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2. Sintering methods of SiC-based composites
2.1. Pressureless sintering
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Eom et al [34]. studied the sinterability, mechanical and thermal properties of the sintered SiC
ceramics using the pressure-less method with the addition of CaO to a composition of Al2 O3 ,
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Y2 O 3 at different temperatures (1750-1900 °C).
They investigated the effect of sintering temperature on the thermal and mechanical properties
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of the SiC ceramics sintered with the additives of Al2 O3 , Y2 O 3 and CaO. The starting materials
used in this research included 91 wt. % β-SiC and 5.5, 3.1 and 0.4 wt. % the additives of Al2 O3 ,
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Y2 O 3 and CaO, respectively. According to the reported results, the SiC samples containing the
additive of Al2 O3 sintered in a temperature range between 1700 to 1850 °C achieved a theoretical
density of above 97%. The results also demonstrated that with an increase in the sintering
temperature to 1900 °C, the density values of the samples decrease again. The researchers
believe that the addition of CaO along with Al2 O3 and Y2 O3 decreases the eutectic temperature
which, in turn, improves the sinterability of the samples. On the other hand, a decrease in the
density of the samples with an increase in the temperature was attributed to a reaction between
SiC and the oxide additives and the production of volatile products such as Al2 O and CO. The
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microstructure of the sintered samples is shown in Fig 1. The grain size of theses samples at
1700, 1750, 1800, 1850 and 1950 °C is 0.8, 0.9, 1.5, 1.8 and 2.8 um, respectively.
As the sintering temperature increases from 1750 to 1800 °C, a tough microstructure
containing rather large elongated grains (flake-like grains in three dimensions) and small
equiaxed grains was achieved. According to their report, the samples with a high toughness are a
result of using additives.
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The results demonstrate that with a further increase in the temperature, the grain size increase
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and the contribution of the elongated grains with a higher aspect ratio increase in the
microstructure consisting of only elongated
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microstructure. As the temperature reaches 1900 °C, all the equiaxed grains disappear and a
grains remains. Microstructural developments
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occurred in this investigation are highly related to the phase transformation of 3C => 6H => 4H
which occurs at high temperatures. It's worth mentioning that the β to α phase transformation
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results in the growth of elongated grains with a high aspect ratio.
Based on the reported results, with the increase in the sintering temperature (from 1700 °C to
1850 °C) the fracture toughness has increased from 2.3 to 5.3 MPa.m1/2 . Nonetheless, with
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increasing the sintering temperature to 1900 °C, the toughness has decreased to 5.1 MPa.m1/2 .
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The increase in the fracture toughness of the samples can be attributed to the formation of
elongated grains and the activation of crack deflection and crack bridging mechanisms [19, 20,
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35]. A decrease in the toughness of the samples can also be ascribed to more fracture of large
SiC grains sintered at 1900 °C which is represented in Fig 2.
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According to the published report, the flexural strength has reached 633 MPa with increasing
the sintering temperature from 1700 °C to 1850 °C and decreased to 377 MP as the temperature
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reached 1900 °C. These researchers believe that the high fracture strength of the sintered samples
at 1800 °C and 1850 °C can be because of their high density and the high fracture energy of the
Al2 O 3 –Y2 O3 –CaO additive system compared with the other additive systems. On the other hand,
the porosity present in the sample, that an example of which is shown in Fig 3, is considered as
the main factor affecting the decrease in the flexural strength.
The hardness of the samples increased from 23.1 GP to 29.1 GPa with an increase in the
sintering temperature (1700 to 1850 °C) and as the temperature reached 1900 °C the hardness
decreased to 27.4 GPa. In the current study, the hardness of the SiC samples has been attributed
to the remaining porosity, the amount of the second phase, the reaction of the additives with the
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matrix and grain size. Of which the remaining porosity is considered as the most important
factor.
Furthermore, the results presented showed that the electrical conductivity of SiC ceramics
increases with increasing the sintering temperature as a result of a decrease in the oxygen content
in the network of SiC grains. In this investigation, the best results were achieved at 1850 °C. The
density, flexural strength, fracture toughness, hardness and thermal conductivity of the samples
, respectively when the structure contained 62.2 % 4H, 35.7 % 6H and 2.1 % 3C.
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at this temperature were reported to be 99%, 628 MPa, 5.3 MPa.m1/2 , 29.1 GPa, and 80 W.(m.K)-
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Zawrah et al [36], have also investigated the liquid phase sintering of α-SiC in the
presence of CaO using the pressure-less sintering method. They used more CaO and sintered two
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samples named S1 and S2 made of SiC containing a constant amount of the additives Al2 O3 and
Y2 O 3 (7 and 2 wt. %, respectively) and various amounts of CaO (1 and 3 wt. % for S1 and S2,
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respectively) at a temperature range between 1750 and 1900 °C until achieving a theoretical
density of above 95 percent. According to the data presented, the compaction process of the
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samples was carried out through the liquid phase produced by Al2 O3 , Y2 O3 and CaO.
The microstructural studies also showed that a microstructure containing elongated grains,
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equiaxed and small grains (with an average grain size around ≥10µm) was formed. The relative
densities of S1 and S2 samples at different sintering temperatures are depicted in Fig 4. As it can
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be seen, the highest density was achieved at 1800 °C which is close to YAG melting point.
According to the report, sample S1 has a higher density than sample S2. This behavior is
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attributed to the lower CaO content in S1 sample and it seems that the release of SiO, Al2 O and
CO gases has increased during the sintering process and led to a lower relative density.
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The X-ray diffraction patterns of S1 and S2 samples which were sintered at 1750 and 1800 °C
are presented in Fig 5. As can be seen, the Al5 Y3 O12 (YAG) phase is formed in both samples.
With an increase in CaO content, the oxide phase of CaY2 O4 was also formed. In this research,
the liquid phases of Al5 Y3 O12 and CaY2 O4 , improve the sinterability through activating the quick
diffusion paths in grain boundaries.
The fracture toughness, hardness and elastic modulus of the samples sintered at different
temperatures are presented in table 1. According to the table, these properties have improved up
to 1800 °C and then have decreased at 1900 °C. This trend is in accordance with the changes of
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the density of the samples. This shows that density plays a key role in determining the
mechanical properties in this study, too.
According to the SEM images of the fracture surfaces (fig 6), the fracture mode of S1 and S2
samples is a mixture of intergranular and transgranular fractures. The reason behind this
phenomenon is the weak interface developed a result of a mismatch between the thermal
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expansion of the liquid and matrix phases.
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2.2.Hot pressing method
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Yoshida et al. [37] and some other researchers. [38-41] investigated the effect of CaO on SiC
matrix composites reinforced with SiC fibers. Several researchers [42-47] have also studied the
effect of this material on the fabrication of SiC nano-composites by the hot pressing method. The
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findings of these researchers have been reported in the following.
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2.2.1. Hot pressing of SiC matrix composites reinforced with SiC fibers
Yoshida et al. [37] studied the properties of β-SiC matrix composite reinforced with SiC fibers
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under pressure at high temperature and room temperature. In this research, polycarbosilaneimpregnated 2D woven SiC fiber-reinforced SiC composites using sintering additives of an
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Al2 O 3 -Y2 O3 -CaO (20 wt. %) were fabricated by the hot pressing method. In this study, the
samples were sintered for 1 hour at 1650, 1700 and 1750 °C, under a pressure of 40 MPa. In Fig.
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7, the displacement-load curve of the SiC-SiCf composite obtained at room temperature is
presented. According to the results, the composite failure occurred in a non-brittle manner and
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this behavior increased with decreasing the sintering temperature.
The results showed that in this process, the penetration of polycarbosilane into a fabric
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made of Hi-Nicalon fibers was effective in the non-brittle failure. In Fig. 8, the strength
properties of the composite at room temperature and high temperature are shown.
As shown in Fig. 8, the SiC-SiCf composite hot pressed at 1700 °C has reached the highest
strength value. According to the results, the strengths of the composites hot pressed at 1700 °C
or 1750 °C did not decrease up to the testing temperature of 1200 °C.
The composite fracture energy has increased with increasing the testing temperature and has
reached its maximum at 1200 °C. The results show that an increase in the resistance to crack
propagation and the fiber pull-out dominance has occurred due to the softening of the grain
boundary phase in the composite matrix at the high temperature. In this research, the SiC
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condensation process has been accomplished by the additives of Al2 O3 -Y2 O3 -CaO during the
liquid phase synthesis mechanism, and the glassy phase of SiO 2 has been formed on the SiC
surface with the help of other oxide additives. In Fig. 9, the thermal conductivity of the
composite measured at room temperature is presented.
In this study, the bulk density values of the composites were obtained in the range 2.86-2.72
g/cm3 (90-95% of the relative density). Thermal conductivity has also increased with increasing
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the bulk density and reached 7-14 W/m.K at room temperature. These results are due to the fact
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that the SiC matrix used is in below the submicron range and furthermore, a phase with a low
conductivity of SiC fibers is also effective in this process.
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thermal conductivity has been formed in the grain boundaries. In addition, the low thermal
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Lee et al. [38] prepared SiC composites reinforced with short fibers by the penetration of
polycarbosilane (PCS). These researchers prepared the samples using SiC short fibers (SA
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Tyranno) coated with polycarbosilane and a mixture of sub-micron β-SiC powder and Al2 O3 Y2 O 3 -CaO (containing 14 wt. % Al2 O 3 , 4 wt. % Y2 O3 and 2 wt. % CaO) as additives by the tape
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casting method. In Fig. 10, the process for the preparation of SiC matrix composite reinforced
with short fibers is presented. These green sheets were heated at 1750 °C under the argon
along the tape casting direction.
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atmosphere and 40 MPa pressure for 1 hour. According to the report, the SiC fibers are oriented
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According to the results, the volume density of the composites reached 2.66 g/cm3 . Fig. 11,
shows the XRD patterns of the composites. β-SiC, α-SiC and Y3 Al5 O12 (YAG) phases are
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observed in the structure. It should be noted that the formation of the YAG phase in these
composites can improve the process of sinterability and the strength of the matrix/fiber interface
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[39, 40]. In Fig. 12, the flexural strength variations of the composites at room temperature and
high temperatures are shown.
The highest flexural strength was reported to be 260 MPa at room temperature and at 1000 °C.
The flexural strength of composite A is higher than that of composite B at 1400 ° C. The fracture
surfaces of the composites are shown in Fig. 13.
As shown in Fig. 13, minor short fibers pull-out is observed regardless of the synthesis
conditions. Therefore, the fracture behavior of the composites is a completely brittle failure at
room temperature and temperatures up to 1400 ° C.
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In another study, Lee et al. [41] examined the effect of the volume fraction of the fibers on the
thermal and mechanical properties of SiC composites reinforced with unidirectionally aligned
short fibers. These composites were prepared by hot pressing with the addition of Al2 O 3 -Y2 O3 CaO (14 wt. % Al2 O 3 , 4 wt. % Y2 O 3 and 2 wt. % CaO) and the effect of different parameters on
the thermal and mechanical properties of SiC/SiC composites containing Hi-Nicalon short fibers
coated with BN has been studied. In Table 2, the bulk density values and the mechanical
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properties of hot-pressed SiC/SiC composite synthesized in this study are presented. As can be
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seen, the bulk density values have increased with increasing the sintering temperature, and the
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composites synthesized at 1700 °C have reached a density close to the full density value.
In this study, the low eutectic temperature of the Si-Al-Y-Ca-O oxide compound can be
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considered as an effective factor to increase the density at low temperatures. According to the
results, the fracture toughness of the SiC/SiC composites with the addition of Al2 O3 -Y2 O3 -CaO
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additive is 1.6 times more than that of the monolithic SiC sintered by the hot pressing method at
1750 °C.
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In Fig. 14, the TEM images of the matrix/fiber interface coated with BN are presented in the
sample sintered at 1750°C are depicted. The results of this research have shown that the
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microstructure consists mainly of uniform distribution of fibers in the matrix, fine SiC grains
along with a grain boundary phase and a matrix/fiber interface. Due to the presence of a grain
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boundary phase, it is believed that the composite is densified through the liquid phase
mechanism.
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According to the results, the flexural strength has increased with increasing the sintering
temperature while decreased with increasing the testing temperature. The fracture behavior of the
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SiC/SiC composites is also quite brittle.
The results of the scanning electron microscope (SEM) show that the fibers are pulled out in
the tensile region at the fracture surface but the length of the pulled out fibers is low.
In the following and in Fig. 15, the thermal conductivity of the SiC/SiC composites containing
different amounts of the reinforcing fibers is shown.
According to the results, the thermal conductivities of the hot pressed samples at 1750°C
decreased to 28-26 W/m.K with increasing the volume fraction of the fibers, which is due to the
low-conductivity of the fibers (Thermal conductivity of Hi-Nicalon fiber is equivalent to 7.8
W/m.K at 25 °C and 10.1 W/m.K at 500 °C). Another reason is the low thermal conductivity of
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this composite can be due to the presence of a high amount of grain boundary oxide phase in the
structure [37].
In another report [43], the thermal conductivity of the SiC body, for the synthesis of which,
Al2 O 3 -Y2 O3 -CaO additives have been used, has been investigated. The thermal conductivity of
this sample which has been sintered at 1750 °C by the hot pressing method has reached 32
W/m.K at room temperature. This is much less than the typical thermal conductivity of the SiC
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ceramic. The reason for this difference in thermal conductivity can be attributed to the
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microstructure of the final sample and the matrix phases. Due to the fact that the glassy phases,
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oxides and their compounds, such as YAG, which have been obtained due to the presence of
sintering aids and these agents have low thermal conductivities, the thermal conductivity of the
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2.2.2. Hot pressing of SiC nano-composites
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sample has dropped.
Mitomo et al. [42] investigated the effect of different additives on the sinterability of SiC nano-
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ceramics. These researchers have compressed two types of SiC powder called F and S with a
particle size of 90 and 280 nm, respectively using the hot pressing method with adding 7 wt. %
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Al2 O 3 , 2 wt. % Y2 O 3 and 1 wt. % CaO at different temperatures of 1750 °C and 1900 °C,
respectively, at 20 MPa pressure and at different times of 15 and 30 minutes, respectively. The
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characteristics of the raw materials in Table 3 and the hot pressing conditions to achieve the
highest density are given in Table 4. As can be seen, the sintering temperature of the samples
micron powders.
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synthesized with very fine powders is about 150 °C lower than those synthesized with the sub-
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In addition, the nano-ceramics synthesized with an average grain size of 110 and 510 nm,
respectively, were obtained from F and S powders, respectively and their microstructures are
shown in Fig. 16.
In this study, the deformation behavior of nano-ceramics was compared with that of a ceramic
made with a sub-micron particle size. The relationship between the true strain and the true stress
at 1700 ° C is shown under a constant strain rate in Fig. 17.
As shown in Fig. 17, the nano-ceramics exhibit a higher degree of deformation. The strain rate of
the nano-ceramics is shown in Fig. 18, materials with a deformation rate of over 10 -4 1/s are
generally known as super-plastic ceramics. According to the results presented in this study, the
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deformation rate of 5.4 – 10.0
has been obtained at 1900 °C and under a pressure of
96 MPa.
Also, Kim et al. [44] investigated the sinterability of β-SiC nano-powders by the hot pressing
method. Their samples were sintered at the temperature range of 1700 to 1850°C, under the
argon atmosphere for 30 minutes and under the pressure of 25 MPa. The researchers studied the
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ability to sinterability of three types of commercial SiC nano-powders (powders A, B and C) and
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a SiC nano-powder (powder D). It should be noted that a treatment for the removal of freecarbon at 600 °C, dipping the powder in a HF mixture to remove SiO 2 were carried out on
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powder D. The combination of Al2 O3 -Y2 O3 -CaO was used as a sintering additive material. These
four types of powder contain 90 wt. % SiC, 7 wt. % Al2 O3 , 2 wt. % Y2 O3 and 1 wt. % CaO. The
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characteristics of the starting SiC powders are given in Table 5.
Fig 19, shows the variation of the density of the sintered samples at different temperatures. As
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shown in this figure, the highest density of the samples synthesized with powders A and D was
obtained at 1800 °C, and no change was observed afterwards. In contrast, for the samples
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synthesized with B and C, the highest density was obtained at 1850 °C, and also slightly
increased above 1750 °C. The researchers' findings suggest that the samples synthesized with
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commercially available SiC nano-particles hardly reach a relative density of over 95%. It is
claimed that the high amount of free carbon and the low green density of commercial SiC nano-
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powders are the main causes of weak sinterability of these samples. Reducing the amount of free
carbon by oxidation and controlling SiO 2 content by acid pickling are the appropriate solutions
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for increasing the sinterability of these samples. According to the results, the samples
synthesized with powder D reached a density of about 97% of the theoretical density. The
minutes.
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sintering temperature of these samples was 1800 °C and the duration of the sintering was 30
Nagano et al. [45] examined the effect of argon and nitrogen atmospheres in the sintering stage
on the weight loss of the samples synthesized with β-SiC powder with an average grain size of
90 nm. The researchers performed a heat treatment on β-SiC with various additives containing 7
wt. % Al2 O3 , 2 wt. % Y2 O3 and 1 wt. % CaO in the temperature range of 1600-1800 °C and in
both atmospheres of nitrogen and argon. The applied pressure was equivalent to 30 MPa and the
pressure application time was reported to be 40 minutes. The relationship between weight loss
and temperature of the heat treatment in the Ar atmosphere is shown in Fig. 20. As seen in Fig.
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20, the weight loss increases with an increase in the amount of additive, and as the temperature
increased, the weight loss in the samples increased. According to the results, the heat treated
samples at 1800 °C and under the atmosphere of argon have had a greater weight loss.
The relationship between weight loss and heat treatment temperature in the Nr atmosphere is
shown in Fig. 21, As can be seen, under the Nr atmosphere, similar to the Ar atmosphere, weight
loss has increased once more additives are used. Also, the weight loss has increased with
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increasing the temperature of the heat treatment. However, the amount of the weight loss of the
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samples in the Nr atmosphere is about half to one-third less than that in the Ar atmosphere.
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According to the results obtained from the XRD patterns, by performing the heat treatment
under the Nr atmosphere, the evaporation and decomposition of the grain boundary phase are
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delayed, and the replacement of nitrogen has increased with increasing the temperature and the
time.
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In addition, according to the results presented, the density of the samples heat treated in the Nr
atmosphere is much higher than the samples heat treated in the Ar atmosphere. It has also been
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shown that the weight loss mainly depends on the reaction in the interface between SiC grains
and the grain boundary phase.
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Lee et al. [46] studied the dense nano-structured silicon carbide ceramics fabricated through a
two-stage sintering process. In this study, the SiC-β powders were milled with the additives
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containing 7 wt. % Al2 O3 , 2 wt. % Y2 O3 and 1 wt. % CaO for 6 hours.
According to the information provided, the average particle size after milling has reached 20
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nm. The samples were synthesized in the argon atmosphere under the pressure of 20 MPa under
the conditions given in Table 6.
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In Fig. 22, the grain size–density curves of the samples of this research group have been
shown. As can be seen, using a two-stage sintering, the SiC ceramics are fully condensed, and
nano-grains (sample T5) have been obtained.
Fig. 23, shows the microstructure of sample T5. The grain size distribution in this sample is
uniform and the average grain size is reported to be 43 nm. It should be noted that the grain size
of sample Z2 after sintering at 1750 °C is 38 nm and its relative density has been reported to be
85%. It has been observed that the density of sample T5 after the two-stage sintering at 1550 °C
increased from 85% to 99%, while no severe grain growth has occurred. The reported results
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indicate that the two-stage sintering developed by Chen and Wang is also used in the liquid
phase sintering.
Lee et al. [47] have investigated the mechanical properties and the contact damages of
nanostructured SiC ceramics. In this study, the fine-grained β-SiC powders and the additives
containing 90 wt. % SiC, 7 wt. % Al2 O3 , 2 wt. % Y2 O3 and 1 wt. % CaO were used. Green
samples were hot pressed at 1750 °C and pressure of 20 MPa under an argon atmosphere and
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then hot pressed again at 1550 °C with a pressure of 20 MPa. In this study, the hardness and the
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toughness of nanostructured SiC ceramics were reported to be in the range of 23 GPa, 3.2
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MPa.m½, respectively. In this study, due to the ultrafine grains of the structure, the strength
properties and fracture toughness are not affected by the grain boundary phase.
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The results of the strength testing of the samples as well as the higher strength of
nanostructures ceramic can be explained by the Hall-Petch equation. This equation suggests that
⁄
where d is the grain size. The flexural
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the strength of a polycrystal linearly increases with
strength obtained in this study was in the range of 400-700 MPa. According to the results
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obtained, the fracture toughness of the nanostructured SiC ceramics is lower than that of the
micron-sized structures [17]. In contrast, the hardness and the strength values of the
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nanostructured SiC ceramics are higher than those with a micron-sized structure. Therefore, it
can be claimed that the nanostructured SiC ceramics possess an improved wear resistance and a
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higher fatigue resistance compared with the micron-sized SiC ceramics.
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2.3. Investigation of annealing process in SiC matrix composites
Until now, the annealing process in SiC composites has been studied in two forms i.e. with
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pressure and without pressure at temperatures above the sintering temperature. In the following,
the most important studies on this subject are presented.
Zhan et al. [27] investigated the microstructural analysis of β-SiC sintered by the liquid phase
during annealing under uniaxial compression and without compression. In this study, initial βSiC powder (87 wt. %) with an average particle size of 90 nm with additives containing 7 wt. %
Al2 O 3 , 2 wt. % Y2 O 3 and 1 wt. % CaO were used. The samples were hot pressed at 1750 °C for
40 minutes under the pressure of 25 MPa, under argon atmosphere and then, were annealed for 4
hours at 1900 °C under the pressure of 0.1 MPa with and without using uniaxial compression. As
shown in Fig. 24, the annealing simultaneously with applying the pressure, delayed the phase
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transformation and grain growth, and the obtained samples contained fine grains with β-SiC as
the main phase. In contrast, the microstructure in the annealed materials without applying the
pressure exhibited α-SiC elongated grains.
The EDS analysis results show that there is no difference in the degree of segregation of
aluminum and oxygen atoms at the grain boundaries, but there is a significant difference in the
segregation of yttrium atoms along the SiC grains for both the annealed samples under
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compression (SC1) and the annealed sample without compression (SC2). According to these
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researchers, the intensified segregation of yttrium at the grain boundaries is due to the
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application of pressure, which can be considered as a reason to delay the phase transformation
and grain growth. In Fig. 25, a 1 nm thin secondary phase is observed in the interface for both
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materials. The presence of grains of α/β type is a common characteristic for both materials.
Kim et al. [48] investigated the microstructure of SiC matrix composites, including large
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grains, during the annealing process under compression. In this study, a fine-grained β-SiC
powder (0.1 μm) with 3.3 wt. % of large α-SiC or β-SiC grains as the primary material was used
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and additives containing a mixture of 7 wt. % Al2 O3 , 2 wt. % Y2 O3 and 1 wt. % CaO are added
to the composition. The samples were hot pressed at 1750 ° C for 15 minutes under the pressure
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of 20 MPa, under argon atmosphere and then, annealed at 1850 °C for 0.5, 1, 2, and 4 hours
under argon atmosphere, under compression. In Table7, the sample designing is presented.
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The results indicate that the entrance of larger particles into the β-SiC system accelerates the
growth of the elongated grains during the annealing and therefore, no significant phase transition
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of β→α has occurred. The growth of matrix grains in a sample containing β-SiC particles was
slower than the growth of the sample grains containing α-SiC particles. The sample containing β-
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SiC particles, which were annealed at 1850 for 4 hours, exhibited a bimodal microstructure of
fine-matrix grains and large elongated grains. In contrast, the sample containing α-SiC particles
that were annealed in similar conditions, had a uniform microstructure including elongated
grains. The fracture toughness results of the samples are presented in Table 8.
As can be seen, with increasing the annealing time, the fracture toughness has increased. Also,
the shape of the grains has shifted from the equiaxed grains (A0, B0) to the elongated grains (A4,
B4). Due to the SiC grain growth, the ratio of diameter to grain size has increased on average.
Despite the different microstructures of the A4 and B4 samples, the fracture toughness of sample
A4 is approximately the same as that of sample B4.
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The microstructure of sample B4 consists of large and elongated grains, but their volume
fraction in the microstructure is low. In contrast, the microstructure of sample A4, including
shorter grains with a larger volume fraction in the microstructure. According to the research
carried out in this field [49], the greater fracture toughness in the annealed materials is due to the
bridging by the large and elongated grains.
According to studies [50], the diameter of large grains and the volume fraction of the
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reinforcement phase affects the fracture toughness mechanism through the crack bridging. The
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fracture toughness of the samples is a function of the volume fraction of large grains, the
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diameter of the large grains, and the aspect ratio of large grains. Fig. 26, which is presented by
Kim et al shows that the mechanisms of bridging and deflection of cracks by elongated grains are
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the most important reasons for increasing the fracture toughness of the annealed materials [51].
In order to reinforce the SiC ceramic bodies, Zhan et al. [52] examined the microstructure
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through the addition of seed to the initial composition and performing a supplementary annealing
process. In this study, fine-grained SiC ceramics (less than 1 um) were fabricated. The raw
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materials included very fine β-SiC powder (approximately 90 nm) and additives containing 7 wt.
% Al2 O3 , 2 wt. % Y2 O3 and 1.785 wt. % CaCO 3 . The powder mixtures were hot pressed at 1750
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°C for 40 minutes under the pressure of 30 MPa, under argon atmosphere and then annealed for 4
hours at 1850 °C under argon atmosphere.
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The microstructures of hot pressed and annealed ceramics without α-SiC seeds, included fine,
uniform, and equiaxed grains. In contrast, the annealed samples containing α-SiC seeds had a
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uniform and isotropic microstructure. Due to the excessive growth of β-phase on α-seeds, the
elongated grains are also found in the structure (Fig. 27).
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According to the results, the Weibull modulus and the fracture toughness of fine-grained SiC
ceramics have increased with increasing the grain size to 1 um. Changes in the results of this
study showed that a small amount of grain growth was favorable for the mechanical properties.
Flexural strength, fracture toughness and the Weibull modulus of the annealed materials
containing seeds were reported to be 835 MPa and 4.3 MPa.m1/2 , respectively. (Table 9).
Kim et al. [53] studied the behavior of crack healing in SiC ceramics sintered by the liquid
phase as a function of the temperature of the heat treatment and crack size during annealing
under compression over a period of 4 hours. The mixture of the powders used in this study
contained 0.87 % wt. α-SiC, 86.64 wt. % β-SiC, 7.05 wt. % Al2 O3 , 4.08 wt. % Y2 O3 and 1.36 wt.
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% CaO. The samples were hot pressed at 1810 °C for 1 hour, under the pressure of 25 MPa and
under argon atmosphere. Then, they were annealed at 1910 °C for 4 hours at an atmospheric
pressure of -40 °C. The results indicate that the heat treatment in air has been able to significantly
increase the strength against the indenter. In this study, the temperature of the heat treatment has
a great influence on the amount of crack healing and the recovery degree. It has been stated that
the optimal temperature of the heat treatment depends on the softening temperature of the inter-
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grain phase in each material. According to the results, after performing the heat temperature at
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the optimum temperature, the crack has almost completely disappeared. The results of this group
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showed that in SiC ceramics, sintered by the liquid phase with Al2 O3 -Y2 O 3 -CaO, the heat
treatment at 1100 °C for 1 hour in the air improves the strength, so that the strength reached 1054
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MPa (about 150% of the amount of the strength without indenter). Crack closure and the crack
reconnection due to oxidation of cracked surfaces have been introduced as the dominant
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mechanism of crack healing in SiC ceramics sintered by the liquid phase. The theoretical density
of the samples obtained in this study was reported to be 98.5%.
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Lee et al. [54] investigated the relationship between the microstructure and the fracture
toughness of the SiC bodies after annealing under compression in a period of 4 to 8 hours. The
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properties of the primary SiC powder are given in Table 10. The combination of Al2 O 3 , Y2 O3
and CaO has been chosen as additives. According to Table 11, five powder mixtures containing
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90 wt. % SiC, 7.5 wt. % Al2 O3 , 3.3 wt. % Y2 O3 and 1 wt. % CaO were used. The samples were
hot pressed at 1850 °C for 1 hour and at the pressure of 25 MPa under argon atmosphere. In the
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following, the hot pressed samples in the temperature range of 1850 to 1950 °C were annealed
for 4 to 8 hours under compression and in argon atmosphere to improve the grain growth.
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The growth of the cracks created by the Vickers indentation has been studied to examine the
mechanism of increasing the fracture toughness. The evidence suggests the occurrence of the
crack deflection mechanisms, elastic bridging, and frictional grain bridges which are shown in
Fig. 28. The results have shown that the mechanisms of improving toughness have changed
depending on the grain morphology. Kleebe [55], has reported that the dominant mechanism to
increase the toughness in SiC with YAG as the secondary phase is the occurrence micro-crack
toughening. Therefore, the mechanism of micro-crack toughening in these ceramics cannot be
unlikely, although it has been accepted that crack bridging and crack deflection are the most
important mechanisms for increasing toughness in these ceramics.
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According to the microscopic observations in this study (Fig. 28), it has been shown that in the
samples with elongated α-SiC grains with an aspect ratio >4, a length greater than 2 um, and a
grain thickness (t) lower than 3 um, the dominant mechanism for the improvement of toughness,
is crack deflection. However, in the samples with the grain thickness of 1<t<3 um and a length
greater than 2 um, the dominant mechanism is crack bridging. In this research, the toughness
values varied in the range of 5.4 MPa.m1/2 to 8.7 MPa.m1/2 , due to the microstructural properties
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such as grain thickness, aspect ratio and the total grain volume fraction. It has been reported that
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the difference in fracture toughness is mainly attributed to the amount of grains participating in
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the fracture toughness process.
In Fig. 29, the fracture surfaces of samples 2, 5 and 9 are depicted. As can be seen, the SiC
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grains are not pulled out. The fracture surface in Fig. 29 (b) indicates an increase in the
contribution of the crack deflection. The rough and transgranular fracture surfaces in some grains
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(with ar>4 and l /2 μm) have been reported from the samples that contain grains that are involved
in the crack bridging mechanism (Fig. 29 (a)). On the other hand, the fracture surface of sample
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9 (Fig. 29(c)), which contains large (t >3µm) elongated grains featuring a bimodal grainthickness distribution, showed a transgranular fracture of large grains, which was consistent with
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the observation of an etched surface (Fig. 28(d)).
Kim et al. [56] investigated the microstructure of SiC bodies after annealing under uniaxial
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compression over a period of 1, 3, 4, and 8 hours. The researchers used β-SiC powders
containing 1.1 wt. % α-SiC particles. In this study, the samples were hot pressed at 1800 °C and
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then annealed at 2000 °C under uniaxial compression of 25 MPa to enhance the grain growth.
The additives contain 5.7 wt. % Al2 O3 , 3.3 wt. % Y2 O 3 and 1 wt. % CaO. The annealing
Table 12.
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conditions, relative density and microstructural properties of the annealed samples are given in
The results obtained by the researchers have shown that when the hot pressed SiC samples by
the liquid phase are annealed without compression, the relative density usually decreases slightly
[18, 57]. The comparison of the samples from the study performed by Kim et al [57] with the
other researchers [26] indicates that annealed samples with compression (the relative density
greater than 98%) have higher relative density values than non-compressed annealed samples
(the relative density about 97%). Therefore, it seems that annealing pressure leads to avoid
decreasing the density during the annealing process.
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In all the hot pressed samples which are then annealed, the relative density values greater than
98% have been reported. It has been mentioned that after annealing for 1 hour at 2000 °C, all the
β-SiC phases were changed to α-SiC.
The microstructures of the hot pressed and annealed samples at different times are shown in
Fig. 30. The microstructures of the samples annealed for 1 hour consists of fine and coarse
equiaxed grains and relatively large and elongated SiC grains that are produced due to the
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abnormal growth. In contrast, the annealed for 3, 4 and 8 hours samples under compression
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contain small and large elongated grains.
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Bimodal grain thickness distribution in the samples annealed under compression has occurred
due to the abnormal growth in some grains. According to the results, the in-situ toughened
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microstructure is developed after annealing for 3 hours. Also, the grain thickness and the aspect
ratio of large grains increase with increasing the annealing time, but grain growth was mainly
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due to an increase in the thickness after 3 hours of annealing, which is due to the collision of
large grain. The flexural strength and the fracture toughness of the SiC samples which are
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sintered and the annealed under compression versus different times of the annealing process are
depicted in Fig. 31. As can be seen, the maximum flexural strength (about 500 Mpa) and fracture
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toughness (7.5 MPa.m1/2 ) are obtained in the sample annealed for 4 hours.
Lee et al. [58] investigated the microstructural stability of fine-grained SiC ceramics during the
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annealing under compression for 6 and 12 hours. In this study, the fine-grained SiC ceramics
with an average grain size of about 140 nm or even finer were fabricated by hot pressing at low
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temperatures using very fine β-SiC powders and Al2 O 3 -Y2 O3 -CaO (AYC) additives. The
additives include 7 wt. % Al2 O3 , 2 wt. % Y2 O 3 and 1 wt. % CaO. The samples were sintered at
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2050 MPa under argon atmosphere at 1750 °C for 30 minutes, and then were annealed at 1850
⁰C for 6 and 12 hours under argon atmosphere. The microstructure of hot pressed sample which
has been annealed for different times is shown in Fig. 32. As can be seen, after 6 hours of
annealing, the abnormal growth occurred in a few number of grains. The microstructure of the
sample which has been annealed for 6 hours consists of flake-shaped grains and equiaxed grains,
and after 12 hours of annealing, significant growth occurred in the flake-shaped grains. The
researchers have reported that the phase transformation of β→α from SiC in SiC-AYC was
responsible for accelerating the abnormal growth of grains and converting them into flakeshaped grains.
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3. Summary:
Silicon carbide (SiC) has been paid attention by researchers owing to its special applications in
the fabrication of cutting tools and high-temperature applications. Its low fracture toughness, the
presence of covalent bonds and its low self-diffusivity challenge obtaining a body with a high
relative density and proper toughness. The researchers showed that using Al2 O3 and Y2 O3
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additives results in improved mechanical properties and microstructure. According to literatures,
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CaO is another additive that can remove the weak points along with Al2 O3 and Y2 O3 . The
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addition of CaO decreases the vapor pressure of the liquid phase and as a result, the sintering
temperature decreases, the formation of the liquid phase is accelerated, and the sinterability is
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improved. This can lead to a decrease in the grain growth rate and weight loss during the
sintering process. Moreover, investigations showed that using this additive along with Al2 O3 and
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Y2 O 3 leads to the formation of intermetallic compounds of Al5 Y3 O12 (YAG) and CaY2 O4 . The
mechanical properties of the composite are improved by activating the mechanisms of crack
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deflection, crack bridging, phase transformation, strengthening the grain boundary and changing
the fracture mode.
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In this class of composites, an excessive increase in the sintering temperature decreases the
density of samples, this phenomenon is attributed to a reaction between SiC and the oxide
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additives and the production of volatile products such as Al2 O and CO.
In the samples fabricated by the pressureless sintering method, it was observed that an
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excessive increase in temperature caused an increase in the grain size as well as an increase in
the contribution of elongated grains with a high aspect ratio. Even in some cases, as the
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temperature reaches 1900 °C, all equiaxed grains disappeared and the structure consisted only of
elongated grains. In these cases, it was observed that the fracture toughness dropped which is due
to the fracture of large SiC grains. The results of the research show that at high temperatures,
increased resistance to crack propagation and the occurrence of fiber pull-out occur due to the
softening of grain boundaries in the composite matrix at high temperatures.
Studies showed that the high density and high fracture energy of the Al2 O 3 -Y2 O3 -CaO additive
system are the most important reasons for improving the flexural strength of these composites
compared with other additive systems. On the other hand, porosity is one of the most important
reasons for the loss of this property.
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As expected, the hardness of these composites depends on the factors such as microstructure,
porosity, the distribution of additives, the amount of additives, the reaction of the additives with
the matrix, grain size, and finally, the sintering temperature and time.
In some cases, an increase in temperature initially caused an improved hardness, followed by a
slight decline due to the grain growth and an increase in porosity due to evaporation. According
to studies, the most important and most effective reason for reducing hardness has been reported
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to be the residual porosity in the composite structure. In general, it was found that among all the
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factors, density is one of the most important factors affecting mechanical properties.
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According to studies, several factors affect the sinterability of these composites. One of these
factors is the size of the primary powders. It has been reported that the use of ultra-fine powders
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has reduced the sintering temperature up to 150⁰C compared with those fabricated with submicron powders. It has also been found that in some cases the high amount of free carbon and
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the excessive thickness of the SiO 2 layer formed on the surface of the particles are of the reasons
for reduced sinterability. Reducing the amount of free carbon by oxidation and controlling SiO 2
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content by acid pickling are the appropriate solutions for increasing the sinterability of these
samples.
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Studies showed that the type of sintering atmosphere is effective in weight loss during heat
treatment so that the weight loss of the sample sintered under N is less than that of the sample
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sintered under argon. According to the phase analyses, by performing the heat treatment under
the N atmosphere, the evaporation and decomposition of the grain boundary phase are delayed,
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and the replacement of nitrogen has increased with increasing the temperature and the time. In
addition, the density of the sample sintered under N is higher than that sintered under Ar. It can
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be claimed that weight loss mainly depends on the reaction in the interface between SiC grains
and the grain boundary phase.
Investigations showed that the atmosphere of the sintering process is highly effective on the β
to α phase transformation. Argon generally accelerates the transformation and the formation of
elongated grains. In contrast, nitrogen prevents the transformation and grain growth. It was also
found that the β to α phase transformation depends on the amount of initial β-SiC. The studies
showed that the transformation rate decreases with increasing the β contribution in the presence
of nitrogen. According to studies, the type of starting material is also effective in the
microstructure. Equiaxed grains can be obtained from the powders with a higher contribution of
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α-SiC, while elongated grains can be obtained from β-SiC or β-SiC containing the nucleating
particles of α-SiC.
Another factor that plays an important role in the transformation of β to α is the post-sintering
annealing process. According to the reports, the annealing under compression is effective in
delaying the transformation and preventing the grain growth. It was observed that under these
conditions, the microstructure included fine grains and β-SiC as the main phase. In contrast, the
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microstructure of the samples annealed without compression included elongated α-SiC grains.
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causes a delay in the phase transformation and grain growth.
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The application of the pressure causes severe yttrium segregation in the grain boundaries which
As mentioned before, the presence of elongated grains increases the fracture toughness of SiC
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samples by activating the crack bridging and crack deflection mechanisms. Studies showed that
increasing the annealing time generally increases the fracture toughness of the samples, and also
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the morphology of the grains shifts from equiaxed to elongated. In the sample that underwent
post-sintering annealing, the fracture toughness also depends on the diameter of large grains, the
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volume fraction of the reinforcement phase, and the aspect ratio of large grains and it can be
stated that bridging and deflection of cracks by these elongated grains are the most important
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reasons for increasing the fracture toughness of the annealed composites and the results showed,
the mechanisms of improving fracture toughness have changed depending on the grain
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morphology.
It should be noted, however, that prolonged annealing can cause another decrease in the
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fracture toughness. It seems that the formation of cavities in the structure as a result of overannealing and the occurrence of inter-granular fracture is the most important reasons for the
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reduction of mechanical properties.
A summary of the sintering behavior of the SiC matrix composites which have optimum
properties containing the Al2 O3 -Y2 O3 -CaO (AYC) additives fabricated by various methods are
presented in Table 13.
Declarations of interest: none.
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T
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Carbide Ceramics, J. Am. Ceram. Soc. 86 (2003) 465–470.
[54] S.G. Lee, Y.W. Kim, M. Mitomo, Relationship between microstructure and fracture
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Figure Captions:
Fig. 1. the microstructures of the SiC ceramics sintered at different temperatures for 2 hours under argon
atmosphere: (a) 1750 ⁰C (b) 8
⁰C (c) 8
⁰C and (d) 9
⁰C [34].
⁰C [34].
IP
9
T
Fig. 2. The SEM images of the crack propagation path in the samples sintered at (a) 1850 ⁰C and (b)
Fig. 3. The origin of fracture of the SiC ceramic sintered at 1850 ⁰C [34].
CR
Fig 4. The relative density of the SiC ceramics of S1 and S2 sintered at different sintering temperatures
[36].
US
Fig. 5. The XRD patterns of S1 and S2 samples sintered at 1750 and 1800 ⁰C [36].
Fig. 6. The SEM images of the fracture surfaces of the samples sintered at 1800 ⁰C [36].
AN
Fig. 7. The displacement-load curve of the SiC-SiCf composite hot pressed at different sintering
temperatures obtained by testing at room temperature [37].
Fig. 8. The strength properties of the SiC-SiCf composite as a function of testing temperature [37].
M
Fig. 9. The thermal conductivity properties of the hot pressed SiC-SiCf composite measured at room
temperature as a function of bulk density (the symbols indicate different sintering temperatures) [37].
ED
Fig. 10. The process for the preparation of SiC matrix composite reinforced with short fibers using
polycarbosilane (PCS) and hot pressing [38].
(b) composite B [38].
b=
–SiC
CE
a= Y3 Al5 O12
PT
Fig. 11. The XRD patterns of SiC/SiC composites sintered at 1750⁰C for 1 hour (a) composite A and
c=
–SiC
Fig. 12. The flexural strength curve of the SiC matrix composite reinforced with short fibers sintered at
AC
1750 ⁰C as a function of the testing temperature [38].
Fig. 13. the room-temperature fracture surfaces of the SiC/SiC composites sintered at 1750 ⁰C ,(a)
composite A and (b) composite B [38].
Fig. 14. The TEM images of matrix/fiber interface of the SiC/SiC composite containing 30 vol. % fibers
sintered at 1750 ⁰C [41].
Fig. 15. Thermal conductivity and thermal diffusivity of the SiC/SiC composite sintered at 1750 ⁰C as a
function of the volume fraction of fibers [41].
Fig. 16. The microstructures of the materials sintered from a) F and b) S powders [42].
Fig. 17. The stress-strain relationship compared with the reference ceramic synthesized with a sub-micron
powder [42].
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Fig. 18. The strain rate of nano-ceramics in comparison with the other published data. In this figure, C
and T indicate the under-tension and under-compression data, respectively [42].
Fig 19. The density values of the SiC synthesized with four different starting materials (according to
Table 5) [44].
Fig. 20. The relationship between the weight loss and the heat treatment temperature under the Ar
atmosphere [45].
T
Fig. 21. The relationship between the weight loss and the heat treatment temperature under the Nr
IP
atmosphere [45].
Fig. 22. The grain size-density curves of the samples sintered by different methods (according to Table 6)
CR
[46].
Fig. 23. The microstructure of sample T5 after sintering at 1750 ⁰C followed by annealing at 1550 ⁰C for
US
8 hours [46].
Fig. 24. The microstructural images of the polished samples: a) the annealed material under compression
(SC1) and b) the annealed material without compression (SC2) [27].
AN
Fig. 25. The HRTEM images of the amorphous grain boundaries: a) between two β-SiC (from <110>
direction) in SC1 (the thickness of amorphous phase is 1 nm) and b) between α-6H (from <1120>
M
direction) and β-3C (from <110>) in SC2 (the thickness of amorphous phase is about 1 nm) [23].
Fig 26. The relationship between the fracture toughness and (a) the volume fraction of large grains (b),
ED
the second root of the grain diameter (c) and the aspect ratio of large grains. Materials A ( ) and B ( ) are
depicted [48].
PT
Fig. 27. The SEM images of the polished samples: (a) the hot pressed sample, (b) the seedless annealed
sample and (c) the annealed seed containing sample [53].
CE
Fig. 28. SEM micrographs of crack profiles in annealed samples: (a) sample 2, (b) sample 3, (c) sample 5,
and (d) sample 9 (refer to Table 11). Elastic bridging (sample 2), elastic bridging and frictional grain
AC
bridges (sample 3), crack deflections (sample 5), and cut elongated grains (sample 9) were observed [55].
Fig. 29. Typical fracture surfaces of annealed samples: (a) sample 2, (b) sample 5, and (c) sample 9 (refer
to Table 11) [55].
Fig. 30. The microstructures of the samples annealed for (a) 1 hour, (b) 3 hours, (c) 4 hours, and (d) 8
hours [57].
Fig. 31. The flexural strength and the fracture toughness of the SiC samples which are sintered and
annealed under compression versus different times of the annealing process [57].
Fig. 32. The microstructure of hot pressed samples which have been annealed for different times: (a) the
hot pressed sample, (b) the sample annealed for 6 hours, and (c) the sample annealed for 12 hours [59].
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Table 1. The fracture toughness, hardness and elastic modulus of the samples sintered at different
temperatures [36].
Sintering temperature (°C)
1800
1750
1900
S2
S1
S2
S1
S2
KIC (MPa.m1/2 )
4.1
3.9
5.7
4.2
4.5
4.0
HV (GPa)
20.7
19.2
23.2
22.12
22.6
20.0
E (GPa)
290
280
410
380
330
310
CR
US
AN
M
ED
PT
CE
AC
IP
S1
T
Properties
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Table 2. The bulk density values and the mechanical properties of the hot pressed SiC/SiC composites
[41].
AC
CE
PT
ED
M
AN
US
30 vol. %
Flexural strength
at RT/ MPa
291
373
330
237
372
367
269
287
311
Fracture toughness at
RT/MPa.m1/2
2.57
3.12
3.23
2.63
3.23
3.24
2.94
3.17
3.30
T
20 vol. %
Bulk density /
Mg/m3
2.96
3.10
3.09
2.80
3.06
3.12
2.96
3.13
3.13
IP
10 vol. %
Sintering temp
°C
1650
1700
1750
1650
1700
1750
1650
1700
1750
CR
Fiber volume
fraction
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F
S
Average particle size (nm)
90
280
Content (wt. %)β
98
97
Oxygen content (wt. %)
1.00
0.18
Free carbon (wt. %)
1.88
0.16
AC
CE
PT
ED
M
AN
US
CR
IP
Powder
T
Table 3. Characteristics of the raw materials [42].
ACCEPTED MANUSCRIPT
F
S
1750
1900
15
97.2
30
99
110
510
70
200
AC
CE
PT
ED
M
AN
US
CR
IP
Powder
Hot pressing
Temperature (°C)
Time (min)
Relative density (%)
Average grain size (nm)
Area base (ds)
Number base (dn )
T
Table 4. The synthesis conditions and the properties of the sintered materials [42].
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Table 5. The characteristics of the starting SiC powders [44].
Specific
surface
area (m2 /g)
Oxygen
Free Carbon
A
48
0.45
B
71
C
D
Impurity (mass %)
manufacture
3.73
β
1.30
6.92
β
94
1.08
6
52
1.83
1.75
Sumitomo-Osaka Cement
Co. Japan
Materials Institute Tech.
Inc., Richmond, CA
Nanostructured and
Amorphous materials, Inc.,
Los Alamos, NM
Oxidized and acid-treated
powder A
IP
CR
β
US
AN
M
ED
T
phase
PT
CE
AC
Designation
Average
particle size
(nm)
β
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Table 6. The sintering schedule in this research [46].
Specimens
Conventional Sintering
C1
C2
C3
C4
Sintering conditions
(temperature/holding time)
1500°C/8h
1550°C/8h
1750°C/5 min
1750°C/30 min
Zero-time sintering
Z1
Z2
Z3
Z4
1700°C/0h
1750°C/0h
1800°C/0h
1850°C/0h
Tow-Step sintering
T1
T2
T3
T4
T5
1650°C/0h 1500°C/8h
1700°C/0h 1500°C/8h
1700°C/0h 1550°C/2h
1700/0h 1550°C/8h
1750°C/0h 1550°C/8h
AC
CE
PT
ED
M
AN
US
CR
IP
T
Sintering method
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Table 7. The sample designing [48]
Sample Designation
Material with α-SiC seeds
Material with β-SiC seeds
A0
B0
A1
B1
A2
B2
A3
B3
A4
B4
AC
CE
PT
ED
M
AN
US
CR
IP
T
Annealing time
at 1850 °C (h)
Hot-pressed
0.5
1
2
4
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Table 8. Polytype and fracture toughness of the hot pressed and annealed samples [48].
A0
A2
A4
B0
B2
B4
Hot pressed
1
4
Hot pressed
1
4
Relative
density
(%)
98.1
97.7
97.2
98.3
97.6
97
Polytype content (vol. %)
2H
3C
2
92
80
65
96
85
75
4
4H
6H
5
10
6
15
25
15
20
5
T
Annealing at 1850 °C (h)
AC
CE
PT
ED
M
AN
US
CR
IP
Material
Fracture toughness
(MPa.m1/2 )
2.0
3.7
5.5
2.0
2.9
5.4
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Table 9. The mechanical properties of fine-grained SiC materials [53].
Materials
Flexural Strength
(MPa)
569±169
632±69
835±72
Weibull modulus
m
5
11
14
AC
CE
PT
ED
M
AN
US
CR
IP
T
As-Hot-pressed
Annealed Seedless
Annealed Seeded
Fracture toughness
(MPa.m1/2 )
1.9±0.1
3.5±0.4
0.2±4.3
ACCEPTED MANUSCRIPT
Table 10. The properties of primary SiC powders [55].
Average particle size(µm)
0.32
0.45
0.50
0.70
Specific surface area(m2 /g)
20
18
15
10
Crystalline phase impurity
β
α
α
α
Oxygen (wt. %)
0.22
0.37
0.40
<0.70
Free carbon (wt. %)
0.87
1.08
0.20
<0.60
Manufacturer
Ibiden Co, Gifu,
Japan
Showa Denko,
Tokyo, Japan
Designation
β-SiC
Norton AS,
Lillesand,
Norway
NT
Lonza-Werke,
WaldshutTiengen
LZ
IP
CR
AN
SD
M
ED
PT
CE
AC
T
β-SiC
(ultrafine)
US
Type of Powder
α-SiC
α-SiC
(A-1)
(FCP-15e)
Characteristic
α-SiC
(UF-10)
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Table 11. The composition, annealing conditions, microstructural characteristics and fracture toughness of
the samples [55].
90
90
89
89
89
85
1
1
1
5
85
85
5
5
85
LZ
α-SiC
Average
aspect ratio
5
1900/4
1950/4
1850/4
1900/4
1950/4
1900/4
1900/4
1950/4
1950/8
0.91
1.25
0.82
1.04
1.02
1.13
1.64
2.81
3.84
2.82
2.61
2.80
2.97
3.36
2.57
1.95
2.48
1.98
T
1
2
3
4
5
6
7
8
9
NT
α-SiC
Key grain content ¤ (l>2µm)
ar>4
Key∑
l<t<3µm
t<3µm
grain
(VKG2)
(VKG1 )
(VT)
17.2
37.7
51.3
9.6
60.9
66.2
14.0
33.0
46.9
20.1
45.6
60.2
33.8
46.9
68.4
7.4
51.3
58.1
2.7
59.7
60.9
12.5
50.8
51.4
3.1
35.8
36.3
CR
β-SiC
SD
α-SiC
Mean grain
thickness
(µm)
US
Sample
Annealing
condition
(°C/h)
IP
SiC composition (wt. %)*
Fracture
toughness
(MPa.m1/2 )
7.1±0.3
8.4±0.4
6.2±0.3
7.0±0.4
8.7±0.2
6.6±0.2
6.8±0.3
7.2±0.5
5.4±0.4
*For a batch composition of 90 wt. % SiC, 5.7 wt. % Al2 O3 , 3.3 wt. % Y2 O3 , and 1 wt. % CaO. See Table
I for explanation of the SiC designations. ¤ The abbreviations ‘’ ar’’, ‘’t’’, and ‘’l’’ represent aspect ratio, grain
common grains of KG1 and KG2
AC
CE
PT
ED
M
VT=VKG1 +VKG2 -V
AN
thickness, and grain length, respectively. The total volume fraction of key grains is given as
ACCEPTED MANUSCRIPT
Table 12. The annealing conditions, relative density and the microstructural characteristics of the
annealed samples [57].
Microstructural parameters
Matrix grains
2000°C/1 h
98.0
0.99
25 MPa
2000°C/3 h
3-h annealed
98.2
2.07
25 MPa
2000°C/4 h
4-h annealed
98.3
2.19
25 MPa
2000°C/8 h
8-h annealed
98.1
2.63
25 MPa
a
Hot – pressing conditions:1800°C/1 h/25 MPa
b
Theoretical density = 3.284 g/cm3
c
Total volume fraction of key grains (VT-VKG1 +VKG2 -V
1-h annealed
Thickness
(µm)
Aspect
ratio
2.07
3.37
1.46
40.7
0
3.11
5.43
2.82
58.5
0
3.15
5.60
3.36
62.3
0
3.16
8.66
3.15
50.1
0
AN
US
12
common grains of KG1 and KG2)
M
ED
PT
CE
AC
Key grains c
(vol. %)
Matrix
Large
grains
grains
Aspect
ratio
T
Thickness
(µm)
Large grains
IP
Annealing
conditions
CR
Sample
a
Relative
density b
(%)
ACCEPTED MANUSCRIPT
Table 13. A summary of the sintering behavior of the SiC matrix composites which have optimum
properties containing the Al2 O3 -Y2 O3 -CaO (AYC) additives fabricated by various methods.
Pressureless
1850
Pressureless
1800
AC
Fracture
Density
Hardness Toughness Strength
(% or
(GPa)
(MPa)
gr/cm3 )
(
)
29.1
ref
5.3
628
[34]
US
CR
IP
99
T
Sintering and
Annealing
temperature
( )
23.2
5.7
-
[36]
1750
3.12
-
3.30
370
[41]
hot press
1900
99
-
-
-
[42]
hot press
1750
>95
-
-
-
[44]
ED
M
AN
>95
hot press
CE
β–SiC:91
wt.%
Al2 O3 :5.5
β–SiC,
wt.%
Al2 O3 , Y2 O3 ,
Y2 O3 :3.1
CaO
wt.%
CaO:0.4
wt.%
α–SiC: 90
wt.%
Al2 O3 :7
α–SiC
wt.%
Al2 O3 , Y2 O3 ,
Y2 O3 :2
CaO
wt.%
CaO:1
wt.%
β–SiC
Al2 O3 :14
wt.%
β–SiC,
Y2 O3 : 4
Al2 O3 ,Y2 O3 ,CaO,
wt.%
SiC fiber(Hi–
CaO :2
Nicalon)
wt.%
SiC
fiber:20
vol.%
β–SiC:90
wt.%
Al2 O3 :7
β–SiC (nano),
wt.%
Al2 O3 ,Y2 O3 , CaO
Y2 O3 :2
wt.%
CaO:1
wt.%
β–SiC:90
wt.%
Al2 O3 :7
SiC (nano),
wt.%
Al2 O3 , Y2 O3 ,
Y2 O3 :2
CaO
wt.%
CaO:1
wt.%
Method
sintering
PT
Powders starting
Optimum
percentage
β–SiC (nano),
Al2 O3 :Y2 O3 ,
CaO
Two–step
Sintering
1750/0 h
1550/8 h
99
-
hot press
and Anneal
HOT
press:1750
Annealing:1550
hot press
and Anneal
-
-
[45]
-
-
[46]
US
CR
IP
-
23
3.2
600
[47]
hot press:1750
Annealing:1850
97.2
-
5.5
-
[48]
hot press
and Anneal
HOT
press:1750
Annealing:1850
>97
-
4.3
835±72
[53]
hot press
and Anneal
Hot press:1810
Annealing:1910
98.5
-
7.0±0.47
721±35
[54]
M
AN
97
ED
α –SiC
Al2 O3 , Y2 O3 ,CaO
3.15
PT
β–SiC (nano)
Al2 O3 , Y2 O3 ,CaO
1750
CE
β–SiC (nano)
Al2 O3 , Y2 O3 ,
CaO
hot press
AC
SiC (nano),
Al2 O3 , Y2 O3 ,
CaO
β–SiC:90
wt.%
Al2 O3 :7
wt.%
Y2 O3 :2
wt.%
CaO:1
wt.%
β–SiC:90
wt.%
Al2 O3 :7
wt.%
Y2 O3 :2
wt.%
CaO:1
wt.%
β–SiC:90
wt.%
Al2 O3 :7
wt.%
Y2 O3 :2
wt.%
CaO:1
wt.%
α–SiC:90
wt.%
Al2 O3 :7
wt.%
Y2 O3 :2
wt.%
CaO:1
wt.%
β–SiC
Al2 O3 :7
wt.%
Y2 O3 :2
wt.%
1.785 wt.
% CaCO3
T
ACCEPTED MANUSCRIPT
β–
SiC:86.64
wt.%
α–
α–SiC, β–SiC,
SiC:0.87
AlN, Er2 O3 ,
wt.%
Y3 Al5 O12 , Al2 O3 ,
Al2 O3 :7.05
Y2 O3 , CaO
wt.%
Y2 O3 :4.08
wt.%
CaO:1.36
ACCEPTED MANUSCRIPT
hot press
and Anneal
HOT
prees:1800
Annealing:2000
>99
-
8.7±0.2
-
[55]
7.5
500
[57]
IP
HOT
press:1800
Annealing:1950
PT
ED
M
AN
US
CR
hot press
and Anneal
CE
β–SiC,α–SiC,
Al2 O3 ,Y2 O3 , CaO
β–SiC:89
wt.%
α–SiC:1
wt.%
Al2 O3 :5.7
wt.%
Y2 O3 :3.3
wt.%
CaO:1
wt.%
AC
β–SiC,α–SiC,
Al2 O3 ,Y2 O3 , CaO
β–SiC:89
wt.%
α–SiC:1
wt.%
Al2 O3 :5.7
wt.%
Y2 O3 :3.3
wt.%
CaO:1
wt.%
T
wt.%
98.3
-
ACCEPTED MANUSCRIPT
Highlights:
The addition of CaO decreases the vapor pressure of the liquid phase and as a result, the
sintering temperature decreases.

An excessive increase in the sintering temperature decreases the density of samples,
increases the grain size as and contribution of elongated grains with a high aspect ratio.

The hardness of SiC-AYC composites depends on the factors such as microstructure,
porosity, the distribution of additives, the amount of additives, the reaction of the
additives with the matrix, grain size, and finally, the sintering temperature and time.

Sintering in argon atmosphere accelerates the transformation and the formation of
elongated grains. In contrast, nitrogen prevents the transformation of β to α and grain
growth.

Elongated grains increases the fracture toughness of SiC samples by activating the crack
bridging and crack deflection mechanisms.

The formation of cavities in the structure as a result of over-annealing and the occurrence
of inter-granular fracture are the most important reasons for the reduction of mechanical
properties.
AC
CE
PT
ED
M
AN
US
CR
IP
T

Figure 1
Figure 2
Figure 3
Figure 4
Figure 5
Figure 6
Figure 7
Figure 8
Figure 9
Figure 10
Figure 11
Figure 12
Figure 13
Figure 14
Figure 15
Figure 16
Figure 17
Figure 18
Figure 19
Figure 20
Figure 21
Figure 22
Figure 23
Figure 24
Figure 25
Figure 26
Figure 27
Figure 28
Figure 29
Figure 30
Figure 31
Figure 32
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