Processing and properties of sol-gel-derived alumina/silicon carbide nanocomposites

Article abstract of DOI:10.1111/j.1551-2916.2004.01024.x

Dense alumina/5 vol% SiC nanocomposites were prepared by sol-gel processing using nanosized (180 nm) precoated SiC powders and a commercial boehmite sol. The SiC powder was precoated with boehmite by a controlled heterogeneous precipitation from an aluminum nitrate solution. The coated SiC powder was then dispersed in a boehmite sol, gelled, calcined, and densified by gas pressure sintering under argon atmosphere at 7-8 MPa pressure. The dependence of the calcination conditions on densification, the effect of seeding on the microstructural development, as well as the mechanical behavior of the sintered specimens, are presented and discussed.

Full text of DOI:10.1111/j.1551-2916.2004.01024.x

J. Am. Ceram. Soc., 87 [6] 1024–1030 (2004)
journal
Processing and Properties of Sol–Gel-Derived Alumina/Silicon
Carbide Nanocomposites
†
‡
,§
Unnikrishnan Nair Saraswathy Hareesh, Martin Sternitzke, Rolf Janssen,* and Nils Claussen*
Advanced Ceramics Group, Technische Universit a¨ t Hamburg-Harburg, D-21071 Hamburg, Germany
Krishna Gopakumar Warrier*
Ceramics Division, Regional Research Laboratory (CSIR), Trivandrum, India
1
1
Dense alumina/5 vol% SiC nanocomposites were prepared by
sol–gel processing using nanosized (180 nm) precoated SiC
powders and a commercial boehmite sol. The SiC powder was
precoated with boehmite by a controlled heterogeneous pre-
cipitation from an aluminum nitrate solution. The coated SiC
powder was then dispersed in a boehmite sol, gelled, calcined,
and densified by gas pressure sintering under argon atmo-
sphere at 7–8 MPa pressure. The dependence of the calcina-
tion conditions on densification, the effect of seeding on the
microstructural development, as well as the mechanical behav-
ior of the sintered specimens, are presented and discussed.
of the SiC particles in the alumina matrix. Messing et al.
conducted a more detailed study on the sol–gel preparation, which
was focused on the quantification of the ratio of inter/intragranular
SiC particles. The effect of ␣-alumina seeds on the microstructural
development was also investigated. Nevertheless, hot pressing was
used for the densification and the mechanical properties of the
composites were not presented. The use of an alumina sol as a
sintering additive to promote full densification in pressureless-
sintered alumina/SiC composites with up to 50 vol% SiC was
recently reported by Yang and Troczynski. The alumina/SiC
nanocomposite precursor contained a fraction of alumina together
with a hydrated alumina sol to achieve an enhanced dispersion of
SiC and the fine scale mixing of the constituent powders.
The use of precoated particles or whiskers was found to be very
effective in ceramic/ceramic composites for both the dispersion of
1
4
I. Introduction
LUMINA/SILICON CARBIDE nanocomposites, by virtue of their
the reinforcement phase and the pressureless sintering to nearly
A
improvements in strength, creep resistance, and wear resis-
tance over monolithic alumina, attracted much attention over the
1–3
15–17
full densities.
Generally, precoating reinforcement or matrix
powders with a precursor of the matrix phase was shown to
enhance densification and dispersion characteristics not only in
composites but also in monolithic systems like SiC, Si N , and
1
–14
past decade.
The concept was first introduced by Niihara
4
and reviewed recently by Sternitzke. It is believed that the
improvements in mechanical properties result from the incorpora-
tion of the inert nanosized SiC particles by reducing the flaw size
or by residual stress due to thermal expansion mismatch between
3
4
18–23
mullite.
The aim of the present work was to facilitate the
advantages of both processes (sol–gel processing using alumina
seeds, as well as precoating of the SiC particles) to develop
alumina/SiC nanocomposites with a uniform and fine microstruc-
ture. Additionally, gas-pressure sintering at low pressures of 7–8
MPa was used for densification. Besides details of the processing
route and the resulting microstructures, mechanical properties of
the materials such as strength, toughness, and hardness were also
investigated.
4
the alumina matrix and the SiC inclusions. However, widespread
applications of such structural ceramic nanocomposites are largely
hindered by difficulties in processing. Conventional processing
methods involve solid-state mixing of the constituent powders,
followed by hot pressing or pressureless sintering and hot isostatic
1
–3,5,6
pressing under inert atmospheres.
alone often failed to achieve acceptable levels of densification
Pressureless sintering
7
,8
leading to insufficient mechanical properties.
Moreover, the
difficulty in achieving a homogeneous dispersion of the SiC
particles in the alumina matrix leads to agglomeration problems
having detrimental effects on densification and mechanical prop-
II. Experimental Procedure
(
1) Activation and Precoating of the SiC Powder
One hundred grams of ␣-SiC powder (average particle size 180
9
–11
erties. Alternative preparation techniques are based on sol–gel
1
2,13
nm, FCP 40, Norton, Lillesand, Norway) were mixed with 100 mL
of hydrofluoric acid (40% solution) and stirred for 5 min. The
mixture was then diluted with water to 500 mL and the suspension
was allowed to settle. The supernatant was decanted off and the
powders were repeatedly (five times) washed with distilled water
and dried at T Ͻ 70°C.
or polymer processing
tion problems.
which were used to avoid agglomera-
9
By using a sol–gel route as reported earlier by Xu et al. and
Haaland et al. it is possible to produce a homogeneous dispersion
1
0
For one batch, 4 g of the activated SiC powder was ultrasoni-
cally dispersed in an aqueous medium for 30 min followed by ball
milling for 12 h. The resulting SiC suspension was used for the
precoating step. The SiC powder dispersion filled up to 1 L with
distilled water was mixed with 0.025M aluminum nitrate solution
(AR, Riedel deHaen, Seelze, Germany). The amount of the
aluminum nitrate solution was adjusted to yield 15 wt% alumina
respective to the SiC powder. The mixture was heated to a
temperature of more than 90°C and stirred continuously. The
precipitation of boehmite was then initiated by the dropwise
addition of ammonia until the pH reached a value of 7.5. The
coated SiC particles were removed from the suspension by
centrifugation at 4000 rpm for 10 min. The coated SiC powders
B. Derby—contributing editor
Manuscript No. 187270. Received December 11, 2001; approved January 15,
2
004.
This project was partly supported by the CSIR-KFA/DLR Indo-German Co-
operative Science Program. U. N. S. Hareesh acknowledges the financial support
from the German Academic Exchange Service (DAAD) for a research fellowship.
*
Member, American Ceramic Society.
†
Presently at Ceramics Division, Institute for New Materials, Saarbruecken, 66123
Germany.
‡
Materials Physics and Technology Group, Technische Universit a¨ t Hamburg-
Harburg.
§
Author to whom correspondence should be addressed.
1
024

June 2004
Processing and Properties of Sol–Gel-Derived Nanocomposites
1025
were repeatedly washed with ethanol, dried at 75°C for 48 h, and
finally calcined at 500°C for 2 h at a heating rate of 3°C/min.
(2) Processing of Nanocomposites
For the processing of the alumina/5 vol% SiC nanocomposite
powders a batch size of 50 g was used. The starting precursor was
a boehmite sol that was obtained by dispersing a commercially
available powder (Disperal, Condea, Brunsbuettel, Germany) at
pH 2.5–3. An amount of 647.60 g of the boehmite sol yielding
4
6.970 g of alumina was mixed with 22.32 g of an ␣-alumina seed
suspension yielding 2 wt% seeds. The precoated SiC powder was
then dispersed in the sol and the mixture was flocculated at pH 7.5
by the dropwise addition of ammonia. The resulting viscous gel
was centrifuged, washed, and dried. The alumina/SiC nanocom-
posite precursor was then calcined at temperatures in the range of
7
00° to 1000°C for 2 h under N atmosphere. After calcination the
2
precursors were ground, attrition milled in ethanol, and dried.
Bars of dimensions 4.5 mm ϫ 6 mm ϫ 40 mm were compacted
by uniaxial pressing at 25 MPa, followed by cold isostatic pressing
at 500 MPa. The sintering was conducted in a gas-pressure
sintering furnace at temperatures ranging from 1550° to 1700°C,
with a soaking period of 1 h under an argon atmosphere at a gas
pressure of 80 bar. The whole process is illustrated by the
flowchart in Fig. 1 and the sintering schedule is presented in Fig.
Fig. 2. Sintering schedule for alumina/SiC nanocomposites.
microscopy. SEM specimens were polished and thermally etched
at 1450°C for 1 h in a static argon atmosphere. The grain size
analysis was done by the linear intercept method using a correction
factor of ␲/2. The densities of the sintered materials were mea-
sured on ground specimens by the Archimedes principle using
distilled water as the reference medium. The theoretical densities
2
1
. The sintered bars were ground on all four sides, first with a
51-␮m- and then with a 15-␮m-diamond-impregnated wheel. At
least 500 ␮m of material was removed from all sides.
Monolithic alumina samples were also prepared from sol–gel
boehmite seeded with 2 wt% ␣-alumina following the method
adopted for the processing of nanocomposites. The alumina
precursors were sintered in air at 1550°C for 2 h.
3
were calculated using the values of 3.98 g/cm for alumina and
3
3
.21 g/cm for SiC.
The flexural strength of the materials was measured in four-
point bending on bars with dimensions of 2 mm ϫ 2.5 mm ϫ 30
mm. The tensile face and the adjoining sides were polished down
to a 1-␮m surface finish and the tensile edges were chamfered. The
strength testing was performed on a universal testing machine
(3) Microstructural and Mechanical Investigations
The precoated SiC powders were characterized by transmission
(
Series 1600, ATS, Butler, PA) with a crosshead speed of 0.5
mm/min. The inner and outer spans of the loading rig were 12 and
4 mm, respectively. Fracture toughness (K ) was measured by
electron microscopy (2000 FX II, JEOL, Eiching, Germany). The
TEM specimens were prepared by dispersing the powder on a
C-coated copper grit. Zeta potentials of aqueous suspensions of the
calcined powders were measured at pH values between 2 and 12
2
Ic
the indentation strength method. For that purpose an indent was
created on the midpoint of the tensile face of a four-point bend
testing beam by applying a load of 98.1 N using a Vickers indentor
by using NaNO as an electrolyte (Zetamiser 4, Malvern, Cam-
3
bridge, U.K.). The phase content of the nanocomposite powder
precursors was examined by X-ray diffraction (PW 1710, Philips,
Eindhoven, The Netherlands). Dilatometer studies of the com-
pacted bars were conducted by using the same sintering schedule
presented in Fig. 2 (Baehr Geraetabau GmbH, Luellhorst, Ger-
many). The microstructures of the sintered materials were inves-
tigated by scanning electron microscopy (SEM, Gemini 1530,
LEO, Oberkochen, Germany), as well as by transmission electron
(
Zwick 3212, Zwick, Ulm, Germany). During a standard four-
point bending test, the weakened specimen broke at a lower load,
the “rest load,” ␴ , and the fracture toughness was calculated after
the following equation:
R
2
4
Ϫ1 0.125
0.33 0.75
K ϭ 0.059͑EH ͒ ͑␴ P ͒
R
v
(1)
Ic
where E is Young’s modulus, H is the hardness, and P is the
contact load.
v
Hardness (H) measurements were conducted at room temperature
using a Vickers diamond indentor (Zwick 3212) under a load of
9
8.1 N with a loading time of 20 s. An average hardness was then
calculated from the results of 10 indentations using the following
equation:
F
H ϭ
(2)
2
2
a
where a is the half length of the indent and F is the load.
III. Results and Discussion
(1) Coating of the SiC Particles
Commercial SiC powders contain impurities such as carbon,
oxygen, and other contaminants. Therefore, a heat treatment at a
temperature of about 650°C is necessary to burn out the free
2
5
carbon. In an earlier work, Krull et al. showed that such a
thermal treatment leads to the formation of a thin layer of silica on
the SiC particles due to an oxidation reaction during the heat
treatment. If not removed, the oxide layer on the SiC particles
leads to a glassy grain boundary phase in the final nanocompos-
Fig. 1. Flowchart showing the sol–gel processing of alumina/SiC
nanocomposites.
2
5
ites. Hence, the thermally activated powders were washed with

1
026
Journal of the American Ceramic Society—Hareesh et al.
Vol. 87, No. 6
Fig. 3. TEM micrograph of precoated SiC particles.
Fig. 4. Zeta potential as a function of pH for alumina, SiC and precoated
SiC.
1
0% hydrofluoric acid (HF) to leach out the surface silica. The
heat treatment of the SiC powder followed by the HF leaching not
only removes carbon and silica impurities but also activates the
surface of the SiC particles.
␣-alumina (see Fig. 5) yielded the maximum densification during
sintering. The 700°C calcined precursor consisting of a mixture of
␥-alumina and ␦-alumina showed a lower density after sintering.
Furthermore, the samples cracked during sintering due to the large
weight loss of 7.31%. This is associated with an incomplete
dehydroxylation of the precursor during calcination at 700°C,
resulting in residual hydroxyl groups. During sintering at high
temperatures the hydroxyl groups undergo a chemical decompo-
sition leading to pore channels, which later grow to introduce
cracks in the sample. By increasing the calcination temperature to
800°C, the weight loss is reduced to 6.55% and the sintered
samples attained a theoretical density (TD) of 97.6%. Calcining
the precursor at 900°C leads to an even lower weight loss and a
higher densification. Increasing the calcination temperature leads
to a higher dehydroxylation and thus to a decreased amount of
residual OH groups, which prevents the formation of cracks in the
sintered samples. The highest density and the minimum weight
loss were observed for the precursor calcined at 1000°C, indicating
complete dehydroxylation.
The activated SiC particles were then coated with boehmite by
a controlled heterogeneous precipitation from an aluminum nitrate
salt solution at T Ͼ 90°C and pH 7.5. Figure 3 shows a TEM
micrograph of precoated SiC particles with 15 wt% alumina. One
can see in Fig. 3 that the coating consists not of a uniform layer but
an agglomeration of ultrafine poorly crystallized particles adhering
firmly to the SiC surfaces. The thickness of the coating varies but
the surface of the SiC particles is completely covered with the
alumina precipitations. This prevents the direct contact between
SiC particles. The effect of the alumina coatings on the SiC
powder was investigated by zeta potential measurements as pre-
sented in Fig. 4. The uncoated SiC powder exhibits an isoelectric
point at about pH 4 and the pure alumina phase at about pH 9.
Figure 4 demonstrates that there is a clear shift of the isoelectric
point of the coated SiC powder from pH 4 to pH 8, indicating a
change in the surface characteristics from SiC to alumina. Thus,
the coated SiC particles are effective dispersoids in the processing
of alumina/SiC nanocomposites.
The enhanced sintering characteristic of the precursor calcined
at 1000°C is assigned to its composition containing transition and
␣
-alumina phases as well as to the enhanced diffusion rate of the
26
Al species. Moreover, the change in the pore morphology from
(2) Fabrication of the Nanocomposites
slit shaped pores in ␥-alumina to shallow hexagonal pores in
2
7
For the fabrication of alumina/SiC nanocomposites the pre-
transition alumina phases during the topotactic transformation
coated particles were dispersed in a boehmite sol. The precoating
step prevents the segregation of SiC particles during the floccula-
tion of the gel due to differences in the surface charges. As already
demonstrated, the precoated SiC powder show surface character-
istics that are identical to respect to sol particles. The flocculation
of the mixture results in a gel network in which the SiC particles
are held firmly and dispersed uniformly preventing the settling of
SiC particles. Figure 5 contains XRD measurements of an alumi-
na/5 vol% SiC nanocomposite precursor gel seeded with 2%
␣
-alumina before and after calcination at different temperatures to
demonstrate the phase transformations. After drying at 75°C the
gel consists mainly of boehmite, whereas after calcination at
7
00°C the ␥-alumina phase becomes dominant. A further increase
of 100°C in the calcination temperature results in the transforma-
tion to the ␦-alumina phase. At 900°C ␪-alumina starts to appear
though ␦-alumina is still the major phase. At 1000°C a mixture of
␦
- and ␪-transitional alumina phases as well as ␣-alumina is
present.
From the nanocomposite precursors calcined at different tem-
peratures, green bodies were compacted and sintered as described
in the experimental section. Table I summarizes the phase content
of the nanocomposite precursor calcined at various temperatures as
well as their densification behavior. The precursor calcined at
Fig. 5. X-ray powder diffraction pattern of sol–gel alumina/SiC nano-
1
000°C having a mixture of transitional alumina phases along with
composite precursors calcined at different temperatures.

June 2004
Processing and Properties of Sol–Gel-Derived Nanocomposites
1027
Table I. Phase Formation and Densification Behavior of the Seeded Nanocomposite Precursors Calcined at
Different Temperatures and Sintered at 1700°C/90 min
Calcination
temperature (°C)
Green density
(% TD)
Sintered density
(% TD)
Weight loss during
sintering (%)
Alumina phases
7
8
9
00
00
00
␥-Alumina ϩ ␦-alumina
46.25
48.37
49.77
48.23
Cracked
97.6
7.31
6.55
6.18
5.15
Transitional alumina phases ϩ traces of ␥-alumina
Transitional alumina phases
97.7
1
000
Transitional alumina phases ϩ ␣-alumina
99.2
may also contribute to the enhanced densification characteristics of
In our studies we found a porosity of more than 3% in the case of
the unseeded nanocomposites even after sintering at 1700°C for 90
min, which is indicative of a coalescence of pores.
1
000°C calcined precursor.
Dilatometer studies carried on these samples (Fig. 6) show a
two-step shrinkage curve with a total linear shrinkage rate of 22%.
The first step corresponds to the transformation of transitional
alumina phases to ␣-alumina, which is associated with about 9%
linear shrinkage. The shrinkage associated with the densification
process starts at temperatures above 1300°C and extends to over
(
4) Microstructural Development
The effect of ␣-alumina seedings on the microstructural devel-
opment is presented by SEM investigations on polished and
thermally etched specimens (Fig. 8). At 1550°C (Figs. 8(a) and
(b)) both seeded and unseeded samples possess a large amount of
closed porosity. The average matrix grain size of the seeded
samples is less than 1 ␮m with grains having an equiaxed
morphology. The pores are around 0.5 ␮m in size though bigger
pores were also observed. However, the unseeded sample (Fig.
8(b)) shows elongated grains with an average grain size slightly
above 1 ␮m. The elongated morphology of the unseeded sample is
a result of the distinctive vermicular grain growth arising from the
redistribution of porosity during the formation of ␣-alumina from
the transition alumina matrix.
After sintering at 1650°C for 1 h under 80 bar argon pressure,
the seeded samples achieved 98% TD and the microstructure (Fig.
8(c)) shows fine grains with a uniform distribution of SiC particles.
It appears from the SEM micrographs that the SiC particles are
mostly located at the grain boundaries and fewer particles are also
1
700°C. At 1300°C it can be assumed that the samples have no
more open porosity as the densities at this point reach 90%
theoretical.
(3) Effect of the ␣-Alumina Seeding
The densification behavior of the sol–gel composite precursor
with and without ␣-alumina seeds on sintering in the temperature
range of 1550° to 1700°C are presented in Fig. 7. One can see that
the seeded samples attain higher densities than the unseeded
counterparts at a given sintering temperature. Moreover, densities
greater than 96% TD could not be achieved for samples without
seeds even during sintering at 1700°C for 90 min. The better
sintering behavior due to the ␣-alumina seeding has also been
observed earlier in sol–gel-derived alumina, and is attributed to the
28
2
8
finer grain size of the alumina matrix. Kumagai and Messing
suggested that the seeding particles control the growth of
-alumina into the matrix, which consists of transition alumina
9
seen within the matrix grains. Messing et al. demonstrated that
␣
the entrapment of SiC particles takes place during vermicular
growth of the ␣-alumina phase into the transition alumina matrix.
In the case of the seeded precursors, this mechanism is retarded as
the seeds provide nucleation sites for the crystallization of
␣-alumina. Hence, it is reasonable to have reduced intragranular
SiC in the case of seeded samples.
Figure 8 demonstrates a marked difference in the dispersion of
the SiC particles between the seeded and unseeded specimens. In
the case of the unseeded samples (Figs. 8(b), (d), and (f)), the SiC
particles are accumulating at the triple grain junctions together
with a nonuniform dispersion. In this case, SiC particles cannot be
observed within the alumina matrix grains. This is in contrary to an
phases. In this case an aggregate free structure is formed.
In the absence of seeding particles the transformation from the
transition alumina phases into ␣-alumina requires first the forma-
tion of ␣-alumina nuclei involving a higher activation energy.
Furthermore, the subsequent growth of the ␣-alumina colony
results in an aggregated microstructure. Alternatively, Roy and
2
9
Yarbrough presented an explanation on the basis of the concept
of epitaxial nucleation in accordance with the classical sintering
theory. Pores are trapped in the interstices of grains as a result of
the vermicular growth. Thus, it is difficult to eliminate these pores
and they can coalesce during grain growth. On the other hand, if
pores are associated with high angle grain boundaries they can
easily be eliminated by vacancy diffusion along grain boundaries.
9
earlier report by Xu et al. where a model based on the growth of
␣-alumina into the transition alumina was proposed for the
Fig. 6. Shrinkage curve of a sol–gel nanocomposite precursor calcined at
Fig. 7. Densification behavior of seeded and unseeded sol–gel nanocom-
1
000°C.
posite precursors as a function of sintering temperature.

1
028
Journal of the American Ceramic Society—Hareesh et al.
Vol. 87, No. 6
Fig. 8. SEM micrographs of seeded nanocomposites sintered at (a) 1550°, (c) 1650°, and (e) 1700°C, as well as of unseeded nanocomposites sintered at
b) 1550°, (d) 1650°, and (f) 1700°C.
(
9
development of intragranular composites. However, in this work
dispersion of SiC particles and hardly any agglomerates are
observed. At sintering temperatures of 1650° and 1700°C a
significant coarsening and agglomeration of SiC particles are
observed. This suggests a difference in grain boundary mobility of
the matrix alumina grains in the presence of seeds during the final
stages of sintering. In the case of the seeded samples the micro-
structures consist of both uniformly sized alumina grains and
homogeneously distributed SiC particles. In the unseeded counter-
part, the SiC particles are getting dragged by grain boundaries. The
mechanism remains unclear since SiC particles are believed to be
inert at the sintering temperatures.
TEM micrographs of seeded nanocomposites sintered at
1650°C and 1700°C are presented in Figs. 9(a) and (b), respec-
tively. Again, the increase in alumina matrix size with increasing
sintering temperature is obvious. In contradiction to the SEM
micrographs in Fig. 8, the TEM investigations reveal a uniform
dispersion of the SiC particles at the grain boundaries and within
the alumina matrix grains. It is unclear why the SEM sample
preparation technique used leads to the overestimation of the
it was also mentioned that the coarsening and agglomeration of
SiC particles at the grain boundary retard the grain boundary
breakaway during grain growth and decrease the tendency for
entrapping the SiC particles. It can be seen in Fig. 8(d) that SiC
particles have undergone a significant clustering and coalescence
during grain growth, thereby reducing the grain boundary break-
away. Most of the SiC particles are swept to the grain boundaries
or to triple grain junctions.
On sintering the nanocomposite precursors at 1700°C the
samples attained an excellent densification of up to 99.2% TD in
the case of the seeded composition. The resulting microstructure
(
Fig. 8(e)) shows a homogeneous distribution of the SiC particles.
There is a considerable grain growth compared with the sample
sintered at 1650°C and the average matrix grain size is on the order
of 4 ␮m.
In the case of the unseeded composites the change in micro-
structures with increasing sintering temperatures from 1550° to
1
700°C is noteworthy. The microstructure at 1550°C shows a good

June 2004
Processing and Properties of Sol–Gel-Derived Nanocomposites
1029
Fig. 9. TEM micrographs of nanocomposites sintered at (a) 1650° and (b) at 1700°C.
intergranular SiC particles. However, the smaller SiC particles
seem to be entrapped within the alumina grains and the larger SiC
particles seem to remain at the grain boundaries. The curved grain
boundaries formed as a result of pinning by SiC particles are
characteristic of the nanocomposite microstructure. By TEM it was
observed that some triple grain junctions contain zirconia resulting
from debris during the attrition milling step. However, zirconia
segregations at grain boundaries were not detected.
II compares the mechanical properties of the alumina reference
material to the nanocomposite sintered at 1700°C using the seeded
sol–gel precursor. The strength of the nanocomposite is improved
by 25% (504 to 630 MPa) compared with the monolithic alumina,
but the fracture toughness is reduced by 6% (4.32 to 4.06
1
/2
MPa⅐m ). The hardness shows an increase of 12%, from 17.2
GPa for monolithic alumina to 19.3 GPa in the case of the
nanocomposites. Considering the general proposition that process-
ing flaws present in alumina are either large volume pores or large
columnar grains formed as a result of discontinuous grain
(5) Mechanical Properties
3
0,31
growth,
the observed strength improvement in the case of the
Figure 10 presents the dependency of the four-point bending
nanocomposites can be explained on the basis that such processing
flaws are absent in alumina/SiC nanocomposites. Common pro-
cessing flaws observed in nanocomposites are associated with SiC
agglomerates formed during the cold compaction followed by
strength of the seeded alumina/5 vol% SiC nanocomposites with
the sintering temperature. A maximum value of 630 MPa is
achieved for the samples sintered at 1700°C having a TD of
9
9.2%. The increase in fracture strength with increasing sintering
1
2
differential shrinkage during sintering.
temperature can be attributed to the improved densification. Table
IV. Conclusions
The present study has indicated the possibility of processing
alumina/5 vol% SiC nanocomposites from sol–gel-derived precur-
sors involving a boehmite sol, as well as a precoated SiC powder.
The flocculation of the boehmite sol containing precoated SiC
particles resulted in the formation of a gel network in which the
SiC particles were uniformly distributed due to their similar
surface charges. The nanocomposite precursor derived after calci-
nation of the gel at 1000°C containing seeds consists of a mixture
of transition alumina phases along with ␣-alumina and SiC. Such
precursors could be densified to Ͼ99%TD by gas pressure
sintering at 1700°C, resulting in microstructures with an excellent
dispersion of SiC particles. The sintered nanocomposites show a
2
5% increase in strength over monolithic alumina.
By utilization of sol–gel processing for alumina/SiC nanocom-
posites one can avoid two costly steps, namely, the use of ultra fine
and pure alumina powders and hot pressing, which are commonly
used in the literature. Furthermore, in the present paper the crucial
steps for obtaining a homogeneous nanocomposite microstructure
by sol–gel processing was pointed out and solutions such as the
precoating of the SiC powder and seeding were presented.
Fig. 10. Four-point bending strength values of seeded composites as a
function of sintering temperature.
Table II. Comparison of Mechanical Data for Alumina
(
Sintered at 1550°C) and the Nanocomposite Sintered at
Acknowledgments
1
700°C
We thank Professor H. Balmori Ramirez, Department of Metallurgical Engineer-
ing, ESIQIE, Mexico, for the zeta potential measurements.
Property
Alumina
Nanocomposites
4
-point bending strength (MPa)
504
630
Fracture toughness (indentation
4.32
4.06
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