Microstructure and Mechanical Properties of Porous Alumina Ceramics Fabricated by the Decomposition of Aluminum Hydroxide

Article abstract of DOI:10.1111/j.1151-2916.2001.tb01065.x

The mechanical properties of Al₂O,-based porous ceramics fabricated from pure Al₂O₃ powder and the mixtures with Al(OH)₃ were investigated. The fracture strength of the porous Al₂O₃ specime

Full text of DOI:10.1111/j.1151-2916.2001.tb01065.x

J. Am. Ceram. Soc., 84 [11] 2638–44 (2001)
journal
Microstructure and Mechanical Properties of Porous Alumina Ceramics
Fabricated by the Decomposition of Aluminum Hydroxide
Zhen-Yan Deng, Takayuki Fukasawa,* and Motohide Ando
Synergy Ceramics Laboratory, Fine Ceramics Research Association, Nagoya 463-8687, Japan
Guo-Jun Zhang* and Tatsuki Ohji*
Synergy Materials Research Center, National Institute of Advanced Industrial Science and Technology,
Nagoya 463-8687, Japan
The mechanical properties of Al O -based porous ceramics
Porous Al O ceramics have attracted considerable attention for
2 3
many years because of their good thermal stability at elevated
temperatures. Traditional methods for fabricating porous Al O
2 3
2
3
fabricated from pure Al O powder and the mixtures with
2
3
Al(OH) were investigated. The fracture strength of the porous
3
Al O specimens sintered from the mixture was substantially
ceramics for structural applications include partially sintering
2
3
3–7
higher than that of the pure Al O sintered specimens because
Al O powder
and forming the porous structures by adding
8,9 4,5
2
3
2 3
of strong grain bonding that resulted from the fine Al O
fugitives to the starting powder. Green and coworkers found
that before any densification occurs, the formation of necks
between touching particles by surface diffusion can increase the
elastic modulus to ϳ10% of the fully dense value. Ostrowski and
2
3
grains produced by the decomposition of Al(OH) . However,
3
the elastic modulus of the porous Al O specimens did not
2
3
increase with the incorporation of Al(OH) , so that the strain
3
7
to failure of the porous Al O ceramics increased considerably,
R o¨ del found that a small uniaxial pressure during densification
2
3
especially in the specimens with high porosity, because of the
unique pore structures related to the large original Al(OH)3
particles. Fracture toughness also increased with the addition
has almost no effect on the relationship between strength and
9
porosity in Al O ceramics. Lyckfeldt and Ferreira fabricated
2
3
porous Al O materials by a starch-consolidation method and
2
3
of Al(OH) in the specimens with higher porosity. However,
obtained high-porosity Al O . Recently, some new ways to fabri-
3
2 3
fracture toughness did not improve in the specimens with
lower porosity because of the fracture-mode transition from
intergranular, at higher porosity, to transgranular, at lower
porosity.
cate porous Al O ceramics, such as by pulse electric current
2
3
1
0
sintering (PECS)
and by the reaction bonding of Al O
2 3
1
1,12
(RBAO)
from an Al O /aluminum mixture have been devel-
2 3
10
oped. Oh et al. found that PECS could enhance grain neck
growth and increase the fracture strength of porous Al O and
2
3
1
1
Al O –SiC nanocomposites. Claussen et al. found that strong
2
3
grain bonding could be obtained by RBAO, because the strength of
porous Al O fabricated by RBAO was appreciably higher than
I. Introduction
2
3
OROUS ceramics are widely used as filters, sensors, and catalyst
supports, as well as lightweight structural components. Usu-
that fabricated by partial sintering.
P
In the present study, a mixture of ␣-Al
2
O
3
and Al(OH)
experiences a 60%
volume contraction during decomposition and produces fine Al
grains, porous Al ceramics with high porosity and strong grain
bonding should be obtainable. On the other hand, pore structures
are modified by the addition of Al(OH) when the partial-sintering
method is used, because the pore morphology is related in some
way to the shape of the original Al(OH) particles. The goal of this
research was to optimize pore structures and strengthen grain
3
was
ally, these porous materials are fragile, and their mechanical
properties are inferior to those of the corresponding dense mate-
rials. Improvement of the mechanical properties of porous ceram-
ics is an important issue for increasing their reliability in practical
applications.
used as the starting powder. Because Al(OH)
3
O
2 3
O
2 3
1
3
3
Usually, the mechanical properties of porous ceramics are
related to such factors as pore volume fraction, pore structures, and
grain bonding of the matrix grains. Optimizing pore structures
represents an effective way of improving the mechanical proper-
3
bonding to improve the mechanical properties of porous Al
ceramics.
O
2 3
1
ties of porous ceramics. For example, Shigegaki et al. used
tape-casting to fabricate porous Si N with an anisotropic config-
3
4
uration of grains and pores and found that the strain-to-failure of
II. Experimental Procedure
2
the Si N was appreciably increased. Liu used different-sized
3
4
poly(vinyl butyral) particles as pore-forming agents to fabricate
porous hydroxyapatite ceramics, and his results showed that
specimens with small macropores exhibited higher strength than
those with large macropores.
Highly pure ␣-Al O (99.99%, 0.21 ␮m, TM-DAR, Tamei
2
3
Chemical Co., Nagano, Japan) and Al(OH)3 (99.99%, 2.5 ␮m,
High Purity Chemical Co., Tokyo, Japan) were used in the present
experiment. Two types of starting powders were prepared: Al O
2
3
and Al O ϩ Al(OH) , referred to as A and AH, respectively.
2
3
3
The processing route for this work was the same as that in a
1
4
previous study. Green-body billets with a relative density of 53%
were prepared by cold-pressing and then sintered in a furnace with
an air atmosphere. The heating rate was set to 1°C/min at
J. R o¨ del—contributing editor
Ͻ1000°C, and 10°C/min at Ͼ1000°C for the Al(OH) -containing
3
billets, so that the decomposition of Al(OH) would be complete
3
before sintering began. A heating rate of 10°C/min was used for
the pure Al O billets. All specimens were held at the sintering
temperature for 30 min and then cooled to room temperature at a
rate of 10°C/min. Different sintering temperatures were used to get
Manuscript No. 188448. Received August 3, 2000; approved June 27, 2001.
Supported by METI, Japan, as part of the Synergy Ceramics Project. Supported in
part by NEDO. The authors are members of the Joint Research Consortium of
Synergy Ceramics.
2
3
*
Member, American Ceramic Society.
2
638

November 2001
Porous Alumina Ceramics Fabricated by the Decomposition of Al(OH)3
2639
different porosity specimens, as shown in Fig. 1, where the number
after AH represents the volume percentage of Al(OH)3 in the
starting mixtures; e.g., AH60 represents 60 vol% Al(OH) in the
3
mixture. As the decomposition of Al(OH) produces 60% volume
3
contraction, the porosity of the Al(OH) -containing billets is
3
greater than in the pure Al O billets before sintering, and the AH
2
3
billets required a higher sintering temperature to reach the same
1
5,16
density as that of the pure Al O billets.
2
3
The sintered specimens were cut into pieces measuring 3 mm ϫ
mm ϫ 40 mm, and then ground and beveled before strength and
4
toughness measurements. Strength was determined by three-point
bend tests, with a span of 30 mm and a crosshead speed of 0.5
mm/min. Fracture toughness was measured using the single-edge
1
7
notched-beam test, with a notch depth of ϳ2.0 mm and a notch
width of 0.1 mm. A span of 16 mm of the three-point bend test was
used for strength measurement of the toughness specimens. Six
specimens were used for each strength or toughness point.
Young’s modulus of the porous Al O was measured by the
2
3
pulse-echo method (Model No. 5052 PRX50 Pulser Receiver,
Panametrics, Inc., Waltham, MA), according to JIS R1602. The
data for Young’s modulus represent an average of two or three
specimens.
The pore-size distribution of the sintered porous Al O speci-
2
3
mens was measured by mercury intrusion porosimetry (Model No.
Autopore 9220, Shimadzu Corp., Kyoto, Japan). The density of the
sintered specimens was measured by the Archimedes method. The
morphology and pore structures of the porous Al O were ob-
2
3
served, using the fresh-fractured surface of the sintered specimens,
by SEM.
III. Results and Discussion
(1) Fracture Strength
The relationship between the bending strength and the relative
Fig. 2. Dependence of bending strength on the relative density of porous
density of porous Al O ceramics prepared from different powders
2
3
Al O ceramics prepared from (a) A and AH60 powders and (b) AH60 and
2
3
is shown in Fig. 2. Adding an increasing amount of Al(OH) to the
3
AH90 powders for the sintered specimens with high porosity.
starting powder increased the bending strength of the porous
Al O appreciably, especially for the specimens with high poros-
2
3
ity, compared with those of the same porosity prepared from pure
Al O powder. For example, the specimens prepared from AH90
higher activity than ordinary Al
bonding; similar results have been found by Claussen et al.
Figure 3 shows the pore and grain morphologies of porous Al
specimens with higher porosity prepared from pure Al
Al(OH) -containing powders, indicating that the bonding area
between the Al grains in the AH specimens apparently was
O
grains and can produce good
2
3
2
3
1
1
powder had higher strength than those prepared from AH60, as
shown in Fig. 2(b). The maximum porosity of the porous Al O
O
3
2
3
2
treated by the present method was as high as 62%, and a bending
strength as high as 50 MPa at 50% porosity was obtained.
The reinforcing mechanism for increasing the bending strength
of porous Al O by adding Al(OH) to the starting powder is the
O and
2 3
3
O
2 3
larger than that in pure Al
porosity.
O sintered specimens with comparable
2 3
2
3
3
strong interface bonding between the grains produced by the
decomposition of Al(OH) , discussed later. In fact, the decompo-
3
1
4
sition of Al(OH) formed very fine Al O grains, which have a
3
2 3
(
2) Elastic Modulus
The dependence of Young’s modulus on the relative density of
porous Al O ceramics prepared from different powders is shown
2
3
3
in Fig. 4(a). For comparison, the results of Lam et al. are also
added, where ␳ ϭ 0.62 and ␳ ϭ 0.50 represent the initial relative
0
0
densities of the two types of Al O compacts used by those
2
3
researchers. Unlike the fracture strength, which was dependent on
the porosity, the elastic modulus of the porous Al O did not
2
3
increase with the addition of Al(OH)3 to the starting powder.
Because the fracture strength increased and the elastic modulus did
not increase, the strain to failure (bending strength divided by
Young’s modulus) increased with the addition of Al(OH) to the
3
starting powder, especially for the high-porosity specimens, as
compared with the results for the porous Al O prepared from pure
2
3
Al O powders, as shown in Fig. 4(b). High strain to failure
2
3
implies high reliability and good thermal-shock resistance in
1
application.
3
Lam et al. sintered porous Al O ceramics by using two types
2
3
of pure Al O compacts with different initial densities, and found
2
3
that the sintering temperature and fracture strength of porous
Al O increased as porosity increased in the initial compact, a
2
3
Fig. 1. Dependence of relative density on the sintering temperature for
the specimens prepared from different starting powders.
finding similar to our present results. However, the elastic modulus
also increased as the compact porosity increased (Fig. 4(a)), so that

2
640
Journal of the American Ceramic Society—Deng et al.
Vol. 84, No. 11
Fig. 3. SEM micrographs of the fracture surfaces of the porous Al O specimens prepared from (a) pure Al O powder and sintered at 1100°C (porosity
2
3
2 3
4
1.14%) and (b) AH60 powder and sintered at 1200°C (porosity 43.07%).
3
the strain to failure of the porous Al O fabricated by Lam et al.
in the AH specimens are in some way related to the original shape
of the Al(OH) particles, because the morphology of the Al(OH)
2
3
exhibited no improvement (Fig. 4(b)). The above results indicate
that the effect of initial compact porosity before sintering on the
elastic modulus of porous Al O ceramics differs for pure Al O
3
3
particle is retained after the Al(OH) particles have decomposed
3
1
4
into Al O grains.
2
3
2
3
2 3
and Al(OH) -containing billets.
Figure 5 is a schematic representation of the pore morphologies
of the pure Al O and the Al(OH) -containing billets during
3
The low elastic modulus in porous Al O ceramics sintered
2
3
2
3
3
from Al(OH) -containing powder originates from their unique
sintering. A spherical shape is assumed for the Al(OH) particles,
3
14
3
pore morphology and microstructure, because different grain and
pore arrangement results in a different relationship of elastic
because no special shape was observed. Because the average
Al(OH) particle (2.5 ␮m) was much larger than the Al O grains
3
2 3
1
8–21
modulus to porosity.
Actually, the pore shape and structures
produced by Al(OH) decomposition and the added Al O grains
3 2 3
(0.21 ␮m), the pore structures in the AH specimens can be viewed
as the stacking of large solid spheres that, in turn, are constructed
Fig. 5. Two-dimensional schematic representation of (a) general pore
Fig. 4. Dependence of (a) Young’s modulus and (b) strain to failure on
morphology, (b) the main pore channel related to Al(OH) particles in the
3
the relative density of porous Al O ceramics prepared from different
Al(OH) -containing billets, (c) general pore morphology, and (d) the main
2
3
3
powders.
pore channel in the pure Al O prepared billets during sintering.
2
3

November 2001
Porous Alumina Ceramics Fabricated by the Decomposition of Al(OH)3
2641
of another stacking of small Al O grains (Fig. 5(a)). Such pore
and numerous models and formulas have been proposed to explain
and fit the experimental data. The simplest empirical equation
is
2
3
morphology was evident in the AH specimens with high porosity,
as shown in Fig. 3(b). Of course, the small initial Al O grains
3
2,33
2
3
would fill part of the space of the large pores between Al(OH)3
particles, depending on the volume percentage of Al O grains in
E ϭ E exp͑ Ϫ bP͒
(1)
0
2
3
1
3,22
the initial mixtures,
but the basic pore structures should be
where E is the zero-porosity Young’s modulus, P the porosity,
0
similar to those in Fig. 5(a). The bimodal pore structure in the AH
specimens was also demonstrated by their pore size distributions
and b a characteristic constant. Although Eq. (1) is entirely
empirical, not based on theory, it is the best fit to the experimental
(see Fig. 6 of Ref. 14).
3
4
data. In fact, the characteristic constant b is related to the particle
The pore morphology of the porous Al O prepared from pure
2
3
19
20
stacking and pore shape. Rice used a minimum solid area
MSA) model to calculate the relationship of the elastic modulus to
Al O powders differs from that of the AH specimens. Usually,
2
3
(
some volume percentage of large spherical pores can exist in the
initial pure Al O green compacts, depending on the initial
porosity for different particle and pore stacking, and then fitted the
theory data by Eq. (1). He found that different particle and pore
stacking have different b values (Fig. 4 of Ref. 20); e.g., solid
spheres in cubic stacking (SSCS) correspond to b Ϸ 5, and
spherical pores in cubic stacking (SPCS) have a value of b Ϸ 3.
Because the pore structures in porous Al O ceramics prepared
2
3
green-compact density, because of the nonuniform particle ar-
rangement. From the viewpoint of mechanical stability, the spher-
ical pores are more stable than the large pores of other shapes
during cold pressing. This fact is shown in Fig. 5(c), and can be
observed in the pure Al O sintered specimen with high porosity
2
3
2
3
from pure Al O and Al(OH) -containing powders can be viewed
2 3 3
(Fig. 3(a)). Of course, the higher the porosity of the green compact,
as one type of pore or particle stacking (PS1), with the matrix
constructed by the other particle stacking (PS2), then,
the greater the volume percentage of large spherical pores in the
pure Al O sintered porous specimen.
2
3
The unique pore structures in the AH specimens were further
confirmed by the change in their peak pore size, which represents
the diameter of the pore channels between the large solid
E ϭ EЈ exp͑Ϫb P ͒
0
1 1
b P
2 2
1
4
ϭ E
0
exp
ͩ
ͩ
Ϫ
ͪ
exp͑Ϫb P ͒
1 1
spheres. Figure 6 shows that the peak pore size of the AH60
specimens increased during densification. However, the peak pore
size of the pure Al O sintered specimens decreased as the
1 Ϫ P1
b2P2
2
3
ϭ E exp
Ϫ b P Ϫ
ͪ
(2)
densification progressed.
According to the sintering theory,
0
1 1
1
Ϫ P1
2
3,24
there is a critical coor-
dinate number, N , of grains around a pore during sintering. When
c
where P ϭ V /V and P ϭ V /V are the porosity in PS1 and PS2,
1
1
t
2
2
t
the coordinate number, N, is ϾN , the pore is stable and grows, but
c
respectively, with V and V the pore volume in PS1 and PS2,
1 2
if N Ͻ N , the pore contracts. The increase in peak pore size during
c
respectively, and V the total specimen volume; b and b are the
t 1 2
densification for the AH60 specimens is related to the large
coordinate number of the grains around the pores between the
large solid spheres. Figure 5(b) shows that a boundary movement
of the large pores in the AH specimens also accompanies the
change of pore shape and size, which indicates the difficulty of
increasing the necking area between the large solid spheres. The
characteristic constants of PS1 and PS2, respectively. The effec-
tive characteristic constant, b, of the porous Al O ceramic can be
2
3
obtained as follows:
b P ϩ b P /͑1 Ϫ P ͒
1
1
2
2
1
b ϭ
(3)
P
1
8–20
small solid contact area means a small elastic modulus.
The
situation of the pure Al O sintered specimens is different, because
where P ϭ P ϩ P is the total porosity of the specimen. If the
2
3
1
2
the coordinate number of the grains around the pores between the
Al O grains is small, so that the size of the main pore channels
characteristic constants of PS1 and PS2 are equal, b ϭ b ϭ b ,
then
1
2
0
2
3
always decreases during densification, as shown in Fig. 5(d). This
situation also implies that the large spherical pores in pure Al O
b P P
2
0
1
2
3
b ϭ b ϩ
Ͼ b0
(4)
0
sintered specimens are isolated.
The relationship between the porosity and the elastic modulus of
͑1 Ϫ P ͒͑P ϩ P ͒
1 1 2
3–5,18–20,25–31
porous ceramics has been studied by many researchers,
This equation indicates that the elastic modulus of porous ceram-
ics, which consists of one large agglomerate or pore stacking with
the matrix constructed by the same type of particle or pore
stacking, is always lower than that of porous ceramics with only
single-modal particle or pore stacking. In fact, the above conclu-
sions are drawn based on the minimum solid area, because Eq. (1)
is a good approximation of minimum solid area in porous
2
0,34
ceramics.
It is believed that Eq. (4) reflects the internal
physical properties of porous ceramics with bimodal pore struc-
tures.
If the above situation is true for the AH specimens, it is easy to
understand why the elastic modulus of those specimens is not
larger than that of pure Al O sintered specimens with a similar
2
3
initial green density (Fig. 4(a)). The dependence of b on the
porosity ratio of two stackings is shown in Fig. 7. If two stackings
of PS1 and PS2 are the same, the effective constant is always
larger than that of single-modal stacking, depending on the total
porosity. If two stackings are different, the effective constant
usually is located between the value of b and b , depending on the
1
2
porosity ratio and total porosity.
The b value of the porous Al O ceramics sintered from
2
3
different powders has been obtained by fitting the experimental
data in Fig. 4(a) with Eq. (1) and is given in Table I, where the
3
Fig. 6. Dependence of peak pore size on the relative density of the porous
Young’s modulus of dense Al O is E ϭ 400 GPa. Because the
2
3
0
Al O ceramics prepared from A and AH60 powders.
pore structures in the AH specimens can be viewed as a large
2
3

2
642
Journal of the American Ceramic Society—Deng et al.
Vol. 84, No. 11
difficult to evaluate quantitatively the exact pore volume ratio of
different stackings in porous Al O specimens.
2
3
(3) Fracture Toughness
The dependence of fracture toughness (K ) on the relative
IC
density of porous Al O ceramics prepared from different powders
2
3
is shown in Fig. 8. The fracture toughness of porous Al O
2
3
increases with the addition of Al(OH) to the starting powder only
3
for specimens with higher porosity, compared with the pure Al O
2
3
sintered specimens. Also, the fracture toughness increases as the
volume percentage of Al(OH) in the starting powder increases
3
(Fig. 8(b)). At lower porosity, the fracture toughness does not
improve with the addition of Al(OH) . This phenomenon could be
3
attributed to fracture-mode transition from intergranular, for the
higher-porosity AH specimens, to transgranular, for the lower-
porosity specimens, as shown in Figs. 3(b) and 9(b). For pure
Al O sintered specimens, the fracture mode is always intergranu-
2
3
lar, independent of the porosity (Figs. 3(a) and 9(a)). Relative to
intergranular fracture, transgranular fracture in the low-porosity
AH specimens probably reduces the crack deflection and decreases
the critical strain energy release rate.
Transgranular fracture in the AH specimens at lower porosity
probably results from the defects formed in the Al O grains
2
3
during sintering, because the bimodal stacking in the AH billets
would result in nonuniform sintering driving stress and a resultant
nonuniform shrinkage strain inside and near the surface of the
large solid spheres related to the original Al(OH) parti-
3
2
2,23,36–38
cles.
(4) Reinforcing Mechanism
The elastic properties of porous ceramics, such as Young’s
modulus, are proportional to the minimum solid contact area
Fig. 7. Dependence of the effective characteristic constant b on the
porosity ratio of two particle or pore stackings (a) b ϭ b ϭ 3.56 (pure
1
2
Al O sintered ceramics with initial green density ␳ ϭ 0.53) and (b) b ϭ
2
3
o
1
3
.0 (SPCS) and b ϭ 5.0 (SSCS) in porous ceramics with different total
2
porosity.
sphere stacking, with the matrix constructed of a similar small-
grain stacking, and the characteristic constant of AH60 is larger
than that of a pure Al O sintered specimen with a similar green
2
3
density (␳ ϭ 0.53), in accord with the prediction in Fig. 7(a).
0
On the other hand, decreasing the initial green density would
increase the volume percentage of large pores in the green
3
5
compacts, because the porosity for a definite particle stacking is
a constant and the matrix particle stacking without large pores can
be viewed as a stable array; therefore, the volume percentage of
large pores in the pure Al O sintered specimens increases and the
2
3
porosity ratio of the large pore stacking to the Al O grain stacking
2
3
would increase. If the large pore stacking is viewed as SPCS and
the Al O grain stacking as SSCS, the characteristic constant of the
2
3
pure Al O sintered specimens would decrease as the green
2
3
density decreased, as listed in Table I. This suggestion is in good
agreement with the theoretical prediction in Fig. 7(b). Because the
pore shape and particle stacking will change during sintering, it is
Table I. Characteristic Constant of
Porous Al O Ceramics Prepared from
2
3
Different Powders
Starting powder
Characteristic constant, b
␳0
ϭ 0.62
0
ϭ 0.50
4.06
3.56
3.24
3.90
ϳ5
A, ␳ ϭ 0.53
␳
0
AH60
SSCS
SPCS
Fig. 8. Dependence of fracture toughness (KIC) on the relative density of
porous Al O ceramics prepared from (a) A and AH60 powders and (b)
ϳ3
2 3
AH60 and AH90 powders for the sintered specimens with high porosity.

November 2001
Porous Alumina Ceramics Fabricated by the Decomposition of Al(OH)3
2643
Fig. 9. SEM micrographs of the fracture surfaces of the porous Al O specimens prepared from (a) pure Al O powder and sintered at 1300°C (porosity
2
3
2 3
7
.5%) and (b) AH60 powder and sintered at 1450°C (porosity 10.7%).
2
0,34,35,39
between the ceramic particles.
If the spherical particles
IV. Conclusions
The mechanical properties of porous Al O ceramics prepared
are considered and the particle stacking is a cubic array, the
minimum solid area is the necking area between the particles, i.e.,
the grain bonding area. If the particle stacking isn’t the cubic array,
the minimum solid area is a representation of the grain bonding
area. Because Young’s modulus is an internal physical property of
ceramic materials, it is independent of surface defects.
2
3
from pure Al O powder and the mixtures with Al(OH) particles
2
3
3
were investigated, with the following results:
1) The fracture strength of porous Al O ceramics increased
(
2
3
with the addition of Al(OH) to the starting powder because of
3
strong grain bonding resulting from the fine Al O grains pro-
The strength of porous ceramics is also related to the minimum
2 3
2
0,34,35
duced by the decomposition of Al(OH)3.
solid area.
affect the strength, such as surface defects. Recently, Flinn et al.
studied the evolution of defect size and strength of porous Al O
However, there are other possible factors that
4
0
(2) Unlike the fracture strength, which was dependent on
porosity, the elastic modulus of porous Al O ceramics did not
2
3
2
3
increase with the addition of Al(OH) , so that the strain to failure
during sintering and found that the size of surface defects has only
a negligible influence on the strength of porous ceramics, provided
the defect size is large compared with the microstructural features.
Those experimental results showed that porous Al O ceramics
with different-sized surface defects and different matrix grain sizes
follow the same strength-porosity relationship. In fact, the defect-
stress concentration effects on mechanical properties of porous
3
of the porous Al O ceramics increased considerably. The de-
2
3
crease in the elastic modulus with the addition of Al(OH) was a
3
result of the unique pore microstructures related to the large
original Al(OH) particles in the porous Al O .
2
3
3
2 3
(3) The fracture toughness also increased with the addition of
Al(OH)3 in the higher-porosity Al O ceramics. However, the
2
3
4
1
addition of Al(OH)
caused no improvement in the fracture
3
ceramics are limited because of the pore-stress interactions,
toughness of porous Al O ceramics with lower porosity, because
2
3
especially for the porous materials with high porosity, and by
bending tests. Therefore, if there are no other reinforcing mecha-
nisms, the strength of porous ceramics is also determined by the
minimum solid area. The pore structures in the present Al O
of the fracture-mode transition from intergranular, at higher
porosity, to transgranular, at lower porosity.
2
3
specimens are very fine, similar to those fabricated by Flinn et
al.; the surface defect effects on the bending strength could be
neglected.
4
0
Acknowledgment
The authors are grateful to Dr. You Zhou for invaluable discussion and Dr.
Jian-Feng Yang for experimental assistance.
Because the fracture mode was intergranular for the high-
porosity Al O specimens, the bonding interface should be the
2
3
most stress-concentrated place, and its area could be viewed as a
representation of the minimum solid area. The nominal interface
bonding strength can be evaluated by
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