Reactive Hot Pressing of Alumina-Silicon Carbide Nanocomposites

Article abstract of DOI:10.1111/j.1551-2916.2004.00299.x

A dense alumina-silicon carbide (Al₂O₃-SiC) nanocomposite was synthesized in situ from the reaction of mullite, aluminum, and carbon by reactive hot pressing (RHP). Transmission electron microscopy investigation showed that in situ-formed, nanometer-sized SiC particles were mainly entrapped in the matrix grains, whereas submicrometer-sized particles were located at the grain boundaries or triple points of the Al ₂O₃. In addition, no amorphous phase was observed at the interfaces of the Al₂O₃ and SiC grains, which indicated strong direct bonding. Fracture-surface analysis by scanning electron microscopy revealed an intrafracture mode. The bending strength of the nanocomposite RHP-treated at 1800°C was 795 ± 160 MPa, and the fracture toughness, measured by the indentation method, was 3.1 MPa·m^1/2.

Full text of DOI:10.1111/j.1551-2916.2004.00299.x

J. Am. Ceram. Soc., 87 [2] 299–301 (2004)
journal
Reactive Hot Pressing of Alumina-Silicon Carbide Nanocomposites
,
†
,‡
†
,‡
Guo-Jun Zhang,* Jian-Feng Yang,* Motohide Ando, and Tatsuki Ohji*
Synergy Ceramics Laboratory, Fine Ceramics Research Association (FCRA), Nagoya, Aichi 463-8687, Japan
Synergy Materials Research Center, National Institute of Advanced Industrial Science and Technology (AIST), Nagoya,
Aichi 463-8687, Japan
1
6
17,18
A dense alumina-silicon carbide (Al O –SiC) nanocomposite
composite powders. Borsa et al. and Amroune et al.
recently
2
3
was synthesized in situ from the reaction of mullite, aluminum,
and carbon by reactive hot pressing (RHP). Transmission
electron microscopy investigation showed that in situ-formed,
nanometer-sized SiC particles were mainly entrapped in the
matrix grains, whereas submicrometer-sized particles were
located at the grain boundaries or triple points of the Al O . In
studied the reaction conditions for obtaining SiC whiskers in an
Al O matrix from kaolin, andalusite, or kyanite (Al O ⅐SiO ) and
2
3
2
3
2
carbon in an argon atmosphere.
However, because CO gas is emitted during the above-
mentioned reaction process, only Al O –SiC composite powders
2
3
2
3
can be prepared; subsequent hot pressing or pressureless sintering
is needed to obtain dense composite materials.
addition, no amorphous phase was observed at the interfaces
of the Al O and SiC grains, which indicated strong direct
2
3
Very recently, Al O –SiC composites were prepared from the
2 3
bonding. Fracture-surface analysis by scanning electron mi-
croscopy revealed an intrafracture mode. The bending
strength of the nanocomposite RHP-treated at 1800°C was
reaction of SiO –Al-C by the self-propagating high-temperature
2
19,20
synthesis (SHS) technique.
The reaction synthesis of Al O –
2 3
SiC composites without CO release during the process was
accomplished by aluminum addition. Dense Al O –SiC compos-
7
95 ؎ 160 MPa, and the fracture toughness, measured by the
2
3
1/2
indentation method, was 3.1 MPa⅐m
.
ites also can be fabricated by subsequent hot pressing or pressure-
less sintering of the composite powders obtained from the SHS
technique.
I. Introduction
The use of aluminosilicates with a high Al O content is
2
3
beneficial for obtaining Al O –SiC composites with a relatively
2
3
EACTION synthesis is a promising process for directly fabricat-
low SiC content, which is especially common for nanocompos-
R
ing nanomaterials, which are difficult to obtain by the normal
sintering of nanometer-sized starting powder compacts because of
Oswald ripening at high temperatures. Alumina-silicon carbide
1
ites. In addition, composites with an even lower content and
homogeneous distribution of SiC in the Al O matrix are easily
2
3
obtainable by adding a small amount of Al O . This paper reports
2
3
(
Al O –SiC) composites have received significant attention in the
2 3
a method of preparing Al O –SiC nanocomposites based on the
2
3
past decades because of their outstanding properties, including
reaction of mullite, aluminum, and carbon, as follows:
1
–6
7–11
high bending strength,
excellent creep behavior,
and wear
1
2,13
resistance. Al O –SiC composites are produced convention-
ally by hot pressing mechanically mixed Al O and SiC powders
2
3
3͑3Al O ⅐2SiO ͒ ϩ 8Al ϩ 6C 3 13Al O ϩ 6SiC
(1)
2
3
2
2
3
2
3
(
or whiskers) in a nonactive atmosphere. Submicrometer- or
The calculated content of SiC phase in the obtained composite was
18.25 vol%. Reactive hot pressing (RHP) was used to fabricate the
material, and the microstructure and properties of the obtained
product were briefly characterized.
nanometer-sized SiC particles are needed to produce Al O –SiC
2
3
nanocomposites. Such SiC powders usually contain oxide scale on
the particle surfaces, because of the easy oxidation of fine
particles. This oxide scale is believed to degrade the high-
temperature mechanical properties of the obtained composites. In
1
4
the case of Al O –SiC composites, problems are toxicity
2
3
(w)
II. Experimental Procedures
during whisker handling and the high cost of whisker production.
On the other hand, in situ synthesis of composites can eliminate the
above-mentioned problems and has been attracting the attention of
investigators.
The preinvestigation revealed that starting materials with too
coarse a particle size, e.g., 50–100 ␮m aluminum powder or
1
0–15 ␮m graphite powder, were unfavorable for preparing dense
Al O –SiC composites were produced from mixtures of alumi-
2
3
and homogeneous composites. Accordingly, starting powders with
a fine particle size were selected for the present study. The starting
powders were mullite, aluminum, and carbon. The characteristics
of the raw powders, as provided by the suppliers, are listed in
Table I. The stoichiometric powders obtained according to reaction
(1) were mixed in ethanol with Al O balls for 72 h, in a plastic
1
5
nosilicates and carbon by Chaklader et al. Those researchers
reported that high temperature, Ͼ1500°C, was needed to complete
the reactions. Ultimately, a dense Al O –SiC composite with
nanometer-sized SiC particles homogeneously distributed in an
Al O matrix was obtained by hot pressing the synthesized
2
3
2
3
2
3
bottle, and then dried. The obtained mixed powders were RHP-
treated in a BN-coated graphite die under a pressure of 30 MPa for
6
0 min in an argon atmosphere. Material with a relative density of
D. J. Green—contributing editor
Ͼ95% of the theoretical density (TD) was impossible to obtain
when the RHP temperature was Ͻ1700°C. Accordingly, an RHP
temperature of 1800°C was used in the present study. The heating
rate was 10°C/min. At temperatures Ͻ500°C, a vacuum of ϳ1.3 ϫ
Manuscript No. 10047. Received December 16, 2002; approved July 3, 2003.
This work was supported by AIST, METI, Japan, as part of the Synergy Ceramics
Project. Part of the work was supported by NEDO.
Ϫ2
1
0
Pa was maintained in the furnace. Up to 1.3 atm of argon gas
was then added to the chamber, and the pressure used for hot
pressing was gradually increased to 30 MPa. The obtained product
measured 45 mm ϫ 42 mm ϫ 6 mm.
*
Member, American Ceramic Society.
†
‡
Synergy Ceramics Laboratory.
Synergy Materials Research Center.
2
99

3
00
Communications of the American Ceramic Society
Vol. 87, No. 2
Table I. Characteristics of the Starting Powders
Starting
powder
Mean
particle size
Chemical composition
Manufacturer
Mullite
0.74 ␮m
KM102 grade, SiO 28.1 wt%, Al O 71.4 wt%, impurities
KCM Corp. (Nagoya, Japan)
2
2 3
include (wt%) Fe O 0.017, TiO 0.003, CaO 0.011, MgO
2
3
2
0
.013, Na O 0.01, K O 0.003, and ZrO 0.167
2 2 2
99.9% pure; impurities include Fe 0.1%, Si 60 ppm, Cu 20 ppm
Aluminum
Carbon
3 ␮m
High Purity Chemicals Laboratory
(Saitama, Japan)
13 nm
2600# grade, volatile 1.8%, impurities include Al Ͻ 100 ppm;
Na and Fe Ͻ 30 ppm; Si Ͻ 10 ppm, Ca, Ni and Mo Ͻ 3 ppm
Mitsubishi Chemical (Tokyo,
Japan)
Specimens for property evaluation were cut from the RHP
product and then ground with a 600-grit diamond wheel. The
density was measured by the water-displacement method. TD
process will occur for reaction (1). Thus, a normal heating program
was adopted for RHP treatment of the system in the present
investigation.
3
3
values of 3.97 g/cm for Al O and 3.22 g/cm for SiC were used
2
3
for calculating the TD of the composite. The phase composition
was determined by X-ray diffractometry (XRD), using CuK␣
radiation. The three-point bend strength was measured on bars
measuring 3 mm ϫ 4 mm ϫ 42 mm after the four edges had been
beveled using 1000-grit SiC abrasive paper; the span was 30 mm,
and the crosshead speed was 0.5 mm/min. The strength data were
based on an average of five measurements. The fracture toughness
was measured by the indentation method, using a 5 kg load, as in
(
2) Characterization of the Composite
The XRD results for mixed-powder compacts heat-treated at
different temperatures showed that the ␤-SiC phase appeared after
ϳ1200°C and that the reactant mullite disappeared at ϳ1600°C.
The XRD patterns of the obtained composites RHP-treated at
1
600°, 1700°, and 1800°C indicated that the detectable phases
were Al O and ␤-SiC. However, if the RHP temperature was
2
3
Ͻ1700°C, composites with Ͼ95% TD could not be obtained. The
properties of the composite RHP-treated at 1800°C are listed in
Table II. This table shows that a fully densified material was
produced.
2
1
previous studies.
The equation used for calculating fracture toughness, K , was
IC
0
.5 0.5
Ϫ1.5
KIC ϭ 0.026E P aC
(1)
The mean bending strength of the composites was 795 MPa,
although the standard deviation was large. The highest value of
bending strength was 1004 MPa. These results indicate that better
mechanical properties probably are obtainable by more carefully
machining and polishing the specimens and by optimizing such
manufacturing parameters as the particle size of the raw materials,
the mixing method, and the RHP program. On the other hand,
the fracture toughness, measured by the indentation method,
was as low as that of monolithic Al O and SiC, which suggests
where P is the indentation load, a the half-length of the indent, C
the half-length of the crack, and E the elastic modulus of the
composite. The value of E is calculated from the elastic moduli of
the components (EAl2O3 ϭ 380 GPa and ESiC ϭ 414 GPa),
according to the rule of mixtures for particulate composite with
random phase distribution (E ϭ ⌺E V ϭ 386 GPa, here E and V
i
i
i
i
stand, respectively, for the elastic modulus and the volume fraction
of the components Al O and SiC). The data for hardness and
2
3
2 3
toughness were based on an average of 10 measurements. Scan-
ning electron microscopy (SEM; Model JSM-5600, JEOL, Tokyo,
Japan) observation of the fracture surfaces was conducted at 20
kV, and transmission electron microscopy (TEM; Model
H-9000UHR III, Hitachi Ltd., Tokyo, Japan) analysis of the
microstructure was conducted at 300 kV.
that the toughening effect caused by the in situ-formed SiC was
insignificant.
Figure 1 shows a TEM photograph of the obtained nanocom-
posite. The grain size of the matrix Al O ranges from 0.5 to 3 ␮m
2
3
in this image. The in situ-formed, nanometer-sized SiC particles
are mainly entrapped inside the Al O grains, whereas the
2
3
submicrometer-sized particles are located at the grain boundaries
or triple points of the Al O . This different location of the SiC
2
3
III. Results and Discussion
particles of different sizes agrees with the grain growth dynamics
2
4
in ceramic materials. Such a type of SiC particle distribution is
(
1) Thermodynamic Considerations
According to the thermodynamic data, we can calculate the
enthalpy of reaction (1) under standard conditions to be ⌬H°298
2
2
ϭ
Ϫ1871 kJ, indicating that the reaction is exothermic. The free
enthalpy function, ⌬G°(T), in the temperature range of 298-2000 K
is as follows:
Ϫ3
⌬
GЊ͑T͒ ϭ ⌬HЊ298 Ϫ T⌬SЊ ϭ Ϫ1871 ϩ 336 ϫ 10 T͑kJ͒ (2)
The value of ⌬G° is highly negative in the experimental temper-
ature range, demonstrating the thermodynamic possibility of reac-
tion. On the other hand, the calculation method used in the
2
3
theoretical analysis of the SHS process gives an adiabatic
temperature (T ) of 1070°C for this exothermic reaction if the
ad
reaction is conducted at 25°C. Empirically, a T Ն 1800°C is
ad
2
3
necessary for an SHS process. This means that no SHS
Table II. Characteristics of the Al O –SiC Nanocomposite
2
3
Reactive Hot Pressed at 1800°C
Relative
density
(%TD)
Bending
strength
(MPa)
Vickers
hardness
(GPa)
Fracture
Fig. 1. TEM photograph of the obtained nanocomposite by reactive hot
pressing at 1800°C showing the distribution of the in situ formed SiC
particles, i.e., nanometer-sized ones entrapped into the Al O grains and
Phase
composition
toughness
1
/2
(MPa⅐m
)
2
3
the submicrometer-sized ones located at the grain boundaries or triple
points of Al O .
Al O , ␤-SiC
99.9
795 Ϯ 160
17.7
3.1
2
3
2
3

February 2004
Communications of the American Ceramic Society
301
reactive hot pressing. In situ-formed nanometer-sized SiC particles
were mainly entrapped in the Al O matrix grains, whereas
2
3
submicrometer-sized particles were located at the grain boundaries
or triple points of the Al O . The grain boundaries between Al O
2
3
2 3
and SiC bonded directly, with no amorphous phase. The intrafrac-
ture mode dominated the fracture process, because of the strongly
bonded interfaces and residual tensile stress in the Al O grains.
2
3
The bending strength of the nanocomposite RHP-treated at 1800°C
was 795 Ϯ 160 MPa, and the fracture toughness, measured by the
1
/2
indentation method, was 3.1 MPa⅐m
.
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7
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(
8
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1
0
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2
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[
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(
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Ⅺ