Microstructure evolution of ceramic materials during solidification from melts in high-gravity combustion synthesis

Article abstract of DOI:10.1016/j.materresbull.2011.11.038

A variety of bulk ceramic materials were prepared via melt-casting by high-gravity combustion synthesis, and the microstructure evolution during solidification was investigated. Most single-phase ceramics like Al ₂O₃, YAG, and MgAl₂O₄ showed a faceted crystal shape because they have higher melting entropies. The cooling condition had an evident influence on the solidification process, resulting in different microstructures. As an example, a strongly textured structure consisting of well-arranged faceted crystals was observed in the surface layer of Al₂O₃ samples. Compared with single-phase ceramic materials, multiphase eutectic ceramics have lower melting temperatures and can provide more opportunity for tailoring microstructures. From the experimental results, the microstructure of eutectic ceramics was related with the volume fraction of each component, and various eutectic structures with submicron interphase spacings were produced such as lamellae, fibers, and three-dimensional interpenetrating frameworks. No cracks were found at the interface between different phases in the eutectic ceramics, indicating an excellent interfacial bonding strength.

Full text of DOI:10.1016/j.materresbull.2011.11.038

Materials Research Bulletin 47 (2012) 222–234
Contents lists available at SciVerse ScienceDirect
Materials Research Bulletin
journal homepage: www.elsevier.com/locate/matresbu
Microstructure evolution of ceramic materials during solidification from melts
in high-gravity combustion synthesis
a,
a
b
a
c
d
Guanghua Liu *, Jiangtao Li , Kexin Chen , Yixiang Chen , Zhijun Jiang , Yuepeng Song
a
Technical Institute of Physics and Chemistry, Chinese Academy of Sciences, Beijing 100190, China
b
Department of Materials Science and Engineering, State Key Laboratory of New Ceramics and Fine Processing, Tsinghua University, Beijing 100084, China
High-Technology Research and Development Center, The Ministry of Science and Technology of the People’s Republic of China, Beijing 100044, China
Mechanical and Electronic Engineering College, Shandong Agricultural University, Shandong 271018, China
c
d
A R T I C L E I N F O
A B S T R A C T
Article history:
A variety of bulk ceramic materials were prepared via melt-casting by high-gravity combustion
synthesis, and the microstructure evolution during solidification was investigated. Most single-phase
Received 24 July 2011
Received in revised form 25 October 2011
Accepted 11 November 2011
Available online 22 November 2011
2 3 2 4
ceramics like Al O , YAG, and MgAl O showed a faceted crystal shape because they have higher melting
entropies. The cooling condition had an evident influence on the solidification process, resulting in
different microstructures. As an example, a strongly textured structure consisting of well-arranged
2 3
faceted crystals was observed in the surface layer of Al O samples. Compared with single-phase ceramic
Keywords:
materials, multiphase eutectic ceramics have lower melting temperatures and can provide more
opportunity for tailoring microstructures. From the experimental results, the microstructure of eutectic
ceramics was related with the volume fraction of each component, and various eutectic structures with
submicron interphase spacings were produced such as lamellae, fibers, and three-dimensional
interpenetrating frameworks. No cracks were found at the interface between different phases in the
eutectic ceramics, indicating an excellent interfacial bonding strength.
A. Ceramics
A. Composites
B. Chemical synthesis
D. Microstructure
ß 2011 Elsevier Ltd. All rights reserved.
1
. Introduction
In contrast to metals and alloys, which are generally produced by
size is limited only by crucibles. By this technique, ceramic melts
are contained in molybdenum, tungsten or iridium crucibles, and
unidirectional solidification is realized by slowly pulling the
crucibles off the hot region. The apparent thermal gradients in the
Bridgman technique are generally below 10 8C/mm, which results
casting from melts, polycrystalline ceramics are generally fabricated
bysinteringfromfinepowders. Forthemelt-castingofceramics,two
major technical obstacles are involved. The first one is the difficulty
in preparing stable and homogeneous ceramic melts because most
ceramic materials have extremely high melting points and some
decompose instead of melting at elevated temperatures. Therefore,
the melting of ceramic materials is oftenrealized by efficient heating
media such as high-frequency induction, lamp mirror furnaces and
lasers [1–5]. Another obstacle in the melt-casting of ceramics is
exaggerated grain growth during solidification from hot melts. A
large grain size is usually not desirable in ceramic materials because
it will impair the mechanical strength. In order to obtain fine
microstructure, multiphase eutectic ceramics are more frequently
prepared than single-phase ceramics [6–8].
in growth rates of <100 mm/h and interphase spacings of >10
mm.
Another approach is the Czochralski method, in which crucibles are
also required but direct contact between crucibles and growing
materials is avoided. In addition, the floating-zone method is often
used to prepare eutectic ceramics such as Al /YSZ (Y
2
O
3
2 3
O -
stabilized ZrO ), Al /YAG (Al 12 garnet) and Al O /YAG/YSZ
2
2
O
3
5
Y
3
O
2 3
[11–13]. This method does not require crucibles, and a relatively
small volume of sample is melted by high-frequency induction or
lasers. In the floating-zone method, larger thermal gradients up to
3
10 8C/mm can be attained, leading to higher growth rates and
smaller interphase spacings below 1
mm. Other methods for solid
pulling from a melt meniscus include edge-defined film-growth
and micro-pulling down [14–16], which can provide thermal
The preparation of bulk ceramics by direct solidification can be
carried out by unidirectional solidification in a container or pulling
of a solid from a melt meniscus [9,10]. The Bridgman method is a
suitable technique for preparing large samples, where the ingot
2
gradients of 10 8C/mm.
Despite the differences in technical details, all the above-
mentioned methods require an external heating source to achieve
high temperatures for melting ceramic materials. In contrast to
these methods, a furnace-free technique called high-gravity
combustion synthesis has been recently reported to prepare bulk
ceramics by direct solidification [17,18]. By this technique, highly
*
Corresponding author. Tel.: +86 10 82543695; fax: +86 10 82543695.
E-mail address: liugh02@mails.tsinghua.edu.cn (G. Liu).
0025-5408/$ – see front matter ß 2011 Elsevier Ltd. All rights reserved.
doi:10.1016/j.materresbull.2011.11.038

G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
223
Table 1
Nominal chemical compositions and calculated phase compositions of the ceramic samples prepared by high-gravity combustion synthesis.
Sample
Al
Nominal chemical composition (mol%)
Al
T
m/e (8C)
Calculated phase composition (vol%)
Al
Tad (8C)
2
O
3
2
O
3
2054
1950
2135
1825
1825
1860
1715
2
O
3
2884
2593
2476
2859
2884
2678
2678
YAG
2
62.5 Al O
3
+ 37.5 Y
2
O
3
YAG
MgAl
2 3
59 YAG + 41 Al O
MgAl
YA-1
YA-2
AZ
2
O
4
50 Al
80 Al
95 Al
62 Al
65 Al
2
2
2
2
2
O
3
3
3
3
3
+50 MgO
2 4
O
O
O
O
O
+20 Y
+ 5 Y
+ 37 ZrO
+ 16 Y
2
O
3
2
O
3
84 Al
67 Al
2
O
O
3
3
+ 16 YAG
+ 33 ZrO
2
2
+ 1 Y
2
O
3
2
YAZ
2
O
3
+ 19 ZrO
2
2 3 2
49 YAG + 37 Al O + 14 ZrO
T
m/e: melting point for single-phase ceramics or eutectic points for multiphase eutectic ceramics; Tad: calculated adiabatic temperature.
exothermic aluminothermic reactions are utilized to realize high
temperature and produce ceramic melts. At the same time, a high-
gravity field is applied to accelerate the separation between
ceramic and metallic melts and removal of gas bubbles from the
melts. After cooling and solidification, bulk ceramics are obtained
together with metallic ingots. In high-gravity combustion synthe-
sis, the heat energy created by chemical reactions is used for
melting ceramic materials, and thus external heating source is not
necessary. At the same time, combustion reactions progress very
quickly and the whole synthesis process lasts only tens of seconds.
Accordingly, high-gravity combustion synthesis can offer a fast and
economical way to prepare bulk ceramics. Up to now, high-gravity
combustion synthesis has been applied to prepare a variety of
ceramic materials, including single-phase ceramics, multiphase
eutectic ceramics, glasses and glass–ceramics [19–25].
For preparing bulk ceramics by high-gravity combustion
synthesis, microstructural evolution during solidification plays
an important role in determining the structure and properties of
the solidified samples, and a systematic study on this problem is
necessary for the manipulation of microstructure and optimization
of processing parameters. This paper makes a detailed investiga-
tion on microstructure evolution during the solidification of
ceramic melts in high-gravity combustion synthesis. The effects of
phase composition and cooling conditions on the solidification
process and microstructural development are discussed for both
single-phase and multiphase eutectic ceramic materials.
particle size of 2–3
reagents. The Al and NiO powders were used to produce Al
the aluminothermic reaction of 2Al + 3NiO = Al + 3Ni. The
nominal chemical compositions of all the samples are listed in
Table 1.
The raw materials were mixed in ethanol and homogenized by
planetary ball milling for 1 h, and then dried at 80 8C for 8 h to
prepare a uniform mixture of reactant powders. The powder
mixture was cold-pressed into a round compact and loaded into a
graphite crucible. The crucible was coated by carbon felt and then
put into a steel vessel. The vessel was mounted on a Ni-based
superalloy rotator, which was placed in a closed reaction
chamber. A schematic illustration of the equipment is shown in
Fig. 1(a).
m
m, and the oxides were analytical-grade
2 3
O
by
2 3
O
The reaction chamber was evacuated to a vacuum of <100 Pa
and then the rotator was started. By centrifugation, a high-gravity
field was induced. The strength of the high-gravity field was
0
0
evaluated by a high-gravity factor of g /g, where g and g mean the
centrifugal acceleration and normal gravitational acceleration,
respectively. In this work, a high-gravity factor of 900 was applied.
The high-gravity field could be modulated by controlling the
rotation frequency of the rotator. When the designed high-gravity
factor was reached, the combustion reaction was triggered by
passing an electric current through a tungsten coil closely above
the sample. In the reaction, a large amount of heat energy was
created and the products were melted. Then, ceramic and metallic
melts were separated in a high-gravity field because of their
density difference. After cooling and solidification, bulk ceramic
samples and Ni ingots were obtained in graphite crucibles, as
sketched in Fig. 1(b). Fig. 1(c) shows the photographs of some as-
prepared ceramic samples as well as a Ni ingot.
2
. Experimental
2 3 2
Commercial powders of Al, NiO, Y O and ZrO were used as
raw materials. The Al powder had a purity of >99% and an average
Fig. 1. A schematic illustration of the equipment for high-gravity combustion synthesis and photographs of as-prepared samples: (a) a simple drawing of the equipment,
0
2
where the high-gravity field is induced by centrifugation and the high-gravity acceleration is calculated by g =
v
ꢀ(R
1 2
+ R )/2; (b) bulk ceramics and Ni metal ingot obtained in
the graphite crucible; (c) photographs of some ceramic and glass samples as well as the Ni ingot prepared by high-gravity combustion synthesis.

2
24
G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
The prepared ceramic samples were machined and then
characterized. The phase assemblage was identified by X-ray
diffraction (XRD; D8 Focus, Bruker, Germany). The microstructure
was examined by scanning electron microscopy (SEM; S-3400,
Hitachi, Japan), and chemical analysis for selected areas was
carried out by energy dispersive spectroscopy (EDS; INCA, Oxford
Instrument, UK).
3
. Results and discussion
High-gravity combustion synthesis is a complex process
including combustion reaction, phase separation, and solidifica-
tion. Combustion reactions generally progress very quickly and last
no more than 10 s. During the reactions, much heat energy is
released, leading to an abrupt increase in temperature, and the
products occur in a molten state. Then, phase separation happens,
which involves the separation between ceramic and metallic melts
and removal of gas bubbles from the melts. Finally, solidification
takes place during a cooling and bulk ceramic samples are obtained
together with Ni ingots.
Fig. 2. The relationship between the adiabatic temperature and the proportion of
diluent for the combustion reaction of 2Al + 3NiO = Al + 3Ni.
2 3
O
To realize the melt-casting of ceramics by high-gravity combus-
tion synthesis, it is required that the reaction temperature exceeds
the melting or eutectic points of the ceramic materials. For the
materials studied in this work, the adiabatic temperatures are all
much higherthanthe meltingor eutecticpoints, asshowninTable 1.
Consequently, the ceramic samples are expected to be produced by
solidification from melts instead of powder metallurgy.
In this work, both single-phase and multiphase eutectic
ceramics were prepared by high-gravity combustion synthesis,
as shown in Table 1. In the samples, no Al or NiO was found,
indicating that the aluminothermic reaction of 2Al + 3NiO = A-
gradient in melts. By high-gravity combustion synthesis, the
cooling and solidification of ceramic melts occur in a graphite
crucible (Fig. 1(b)). With a good thermal conductivity, the graphite
crucible is a major source of heat dissipation. In this case, a thermal
gradient occurs in the radial direction in melts, where the
temperature decreases from the center to the side. With such a
thermal gradient, solidification will take place first at the side and
then develop into the center. This is verified by SEM observation, as
2 3
shown in Fig. 3, where the growth direction of Al O crystals is
nearly parallel to the radial direction.
l
2
O
3
+ 3Ni was almost complete. In the XRD patterns of the ceramic
Besides crucibles, another source of heat dissipation is the fast
gas flow above the surface of melts under the condition of high-
speed rotation. In this case, a thermal gradient also existed in the
axial direction in the melts. This thermal gradient caused a
directional solidification from the surface to the center, as shown
in Fig. 4. Some huge rod-like crystals up to 1 mm growing along the
axial direction were observed. The formation of these crystals can
be probably attributed to the coupling of combustion reaction with
the high-gravity field, because both the propagation of combustion
samples, diffraction peaks of Ni were not found. EDS analysis
revealed that the content of Ni in the samples was below 1.0 wt%.
From these results, the phase separation between Ni and ceramic
melts was almost finished.
The prepared ceramic samples showed a diversity of micro-
structures, including faceted crystals, skeletal crystals, and fibrous,
lamellar and three-dimensional interpenetrating eutectic struc-
tures. Associated with SEM observations, the microstructural
evolution during solidification for single-phase ceramics and
multiphase eutectic ceramics will be discussed respectively as
follows.
0
wave and the high-gravity force (G ) occur in the axial direction,
where enhanced mass transportation and heat transfer are
expected.
2 3
During solidification, Al O showed a strong tendency to
3
.1. Single-phase ceramics
develop into faceted crystals, even in large dendrites, as shown
in Fig. 5. This agrees well with the Jackson’s criterion on interface
structures and crystal growth modes in solidification of melts.
According to the Jackson’s criterion [26], the interface structure is
identified by an interface roughness parameter defined as
3
.1.1. Al
Single-phase Al
combustion synthesis from the aluminothermic reaction of
Al + 3NiO = Al + 3Ni. For this reaction, the adiabatic tempera-
2 3
O
2 3
O ceramics can be prepared by high-gravity
2
2
O
3
a
m m
= AꢀDS /R, where DS is the melting entropy, R is the gas
ture reaches the boiling point of Ni (2884 8C). The vaporization of
constant, and A is a crystallographic parameter equal to the ratio
Ni is not desirable as this will increase the porosity in the solidified
between the number of bonds parallel to the surface and the total
samples. In order to decrease the reaction temperature, Al
powder is often added as diluent in the thermite reaction. With
increasing content of Al diluent, the adiabatic temperature for
the reaction of 2Al + 3NiO + xAl = (1 + x)Al + 3Ni decreases,
2
O
3
number of bonds. For a solidification process, if
solid interface is atomically rough and non-faceted interface is
expected. If > 2, the interface is atomically smooth and faceted
interface is produced because the crystal growth is limited by
nucleation rate. The melting entropies of ceramic compounds
a < 2, the liquid/
2
O
3
a
2
O
3
2
O
3
as plotted in Fig. 2. When too much diluent is added (x = 1.6), the
adiabatic temperature decreases to below the melting point of
investigated here are listed in Table 2. For Al
A is estimated to be 0.5–1, giving > 2. Therefore, atomically
smooth interface structure should be involved during the
solidification of Al , resulting in faceted crystals.
For the solidification of ceramic materials like Al
melting entropies, the energy barrier for nucleation is high, and the
crystal growth is thought to take place by two-dimensional
nucleation and growth. With as-existing crystals as nucleation
sites, new nuclei form and grow up. By iterative nucleation and
2 3 m
O , DS /R = 5.74, and
Al
2
O
3
(2054 8C). Considering the fact that the actual reaction
a
temperature is lower than the adiabatic one because of heat
dissipation, an appropriate content of diluent is expected as
2 3
O
0
.2 ꢁ x ꢁ 0.6.
2 3
O with large
During the solidification of ceramic melts, the driving force for
liquid–solid transformation and growth rate of solid phases are
closely related to the under-cooling degree. For this reason, the
microstructure evolution during solidification depends on thermal

G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
225
2 3
Fig. 3. SEM images showing varying microstructures in the radial direction in Al O samples.
2 3
Fig. 4. SEM images showing varying microstructures in the axial direction in Al O samples.
2 3
Fig. 5. SEM images and EDS results of Al O dendrites.
growth, large faceted crystals will be produced with a sufficient
material supply from the melt at the early stage of solidification.
When the solidification is close to the end, however, solid phase is
predominant and the melt becomes isolated. In this wise, material
supply for solidification is limited and incomplete growth of
faceted crystals takes place, resulting in a terraced morphology, as
shown in Fig. 6.
Table 2
Melting entropies of ceramic materials involved in this work.
ꢂ1 ꢂ1
Compound
Al O
D
S
m
(J mol
K
)
D
S
m
/R
2
3
47.72
122.38
81.64
5.74
14.72
9.82
At the surface of some Al
2
O
3
samples, bright layers with a
YAG
MgAl
ZrO
2 4
O
thickness of nearly 0.5 mm were produced. These layers were not
entirely smooth but consisted of small flakes with different
2
29.51
3.55

2
26
G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
2 3
Fig. 6. Terraced structures observed in as-prepared Al O samples.
orientations. The width of the flakes varied in the range of 3–5 mm,
and each flake was composed of orderly arranged faceted Al
of faceted crystals. The skeletal crystals are possibly induced by
fast cooling at melt surface. On one hand, fast cooling often results
in large under-cooling degree and high growth rate of crystals. On
the other hand, fast cooling means a short lifetime of melt, which
can lead to the lack of material supply. Without sufficient material
supply, the imperfections created during fast crystal growth
cannot be timely repaired, leading to the occurrence of skeletal
crystals.
2 3
O
crystals, as shown in Fig. 7. Such strongly textured structure was
also verified by XRD analysis. As shown in Fig. 8, an evident
difference in relative intensities of the diffraction peaks was
observed between the XRD patterns of the flakes and that of a
powder sample.
The formation of the strongly textured structure can be
attributed to two reasons. One is the intrinsic faceted growth
In contrast to the surface, in the center part of samples, rather
dense microstructure was obtained. In the fracture surface of the
center part, typical cleavage planes are visible, as shown in Fig. 10.
2 3
character of Al O , and the other is the particular solidification
condition at the melt surface, which is characterized by fast cooling
and free growth. Under high-frequency rotation, the gas flow
sweeping the melt surface causes fast cooling and induces rapid
solidification as well as crystal growth. On the other hand, crystal
growth at the melt surface is not constrained by crucible, and thus
2 3
Few secondary particle inclusions were found inside the Al O
crystals. From EDS analysis for six individual crystals with different
size and morphology (e.g., Figs. 5 and 6), the atomic ratios of Al and
O are 59.5–60.5% and 39.5–40.5%, respectively, showing a good
the Al
2
O
3
crystals can grow freely. It is thought that the strongly
2 3
consistency with the nominal composition of Al O .
textured structure has been produced by parallel continuous
growth. At first, individual primary nuclei are precipitated in the
melt and with different orientations. Then, based on each primary
nucleus, more nuclei occur and grow with an orientation identical
to the basis. Finally, by continuous nucleation and growth, an array
of crystals with the same orientation is produced, which is
recognized as a flake on a macroscopic scale. By the assembly of
such flakes, a strongly texture structure is obtained.
3.1.2. YAG
Besides Al
prepared by high-gravity combustion synthesis. YAG can be
prepared from the reaction of 10Al + 15NiO + 3Y = 2Y A-
12 + 15Ni. For this reaction, the calculated adiabatic tempera-
2 3
O , YAG is another example of single-phase ceramics
2
O
3
3
5
l O
ture is 2593 8C, which is much higher than the melting point of YAG
(1950 8C). Therefore, it is feasible to carry out the melt-casting of
YAG ceramics by high-gravity combustion synthesis.
2 3
In the prepared Al O samples, skeletal crystals were also
observed, as shown in Fig. 9. To some extent, the skeletal crystals
can be regarded as intermediate products from incomplete growth
Fig. 11 shows SEM images of as-prepared YAG ceramics,
without or with glass additives. In pure YAG ceramics without

G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
227
2 3
Fig. 7. SEM images showing the strongly textured structure at the surface of Al O samples.
additives, coarse grains were produced and some of them were
larger than 500 m. At the boundary areas of the coarse grains,
large cavities up to 100 m were present. In the samples with
wt% SiO or 10 wt% Si–Al–Ca–O glass additives, both the porosity
and grain size were greatly reduced. The reduction of porosity and
grain size contributed to a higher hardness. For example, the
Vickers hardness of the YAG sample with 10 wt% Si–Al–Ca–O glass
additive was 8.7 ꢃ 0.2 GPa, which was 21% higher than that
(7.2 ꢃ 0.4 GPa) of pure YAG samples with prepared by high-gravity
combustion synthesis.
m
m
5
2
The above results show that, during the melt-casting of YAG
ceramics, proper additives are helpful to reduce the porosity and
limit grain growth. In the solidified YAG samples, the porosity
comes from two sources. One is the shrinkage cavities created
during fast solidification of the melt, and the other is the gas
bubbles not removed from the melt. The densities of YAG solid and
3
melt are 4.552 and 3.688 g/cm , respectively [27]. In this case, the
solidification of YAG melt is accompanied by a volume shrinkage
of 19%. Under a fast cooling rate and with a short lifetime of melt,
the volume shrinkage is difficult to be fully compensated, which
causes shrinkage cavities. Such shrinkage cavities can be filled by
additives with lower melting temperatures than YAG. When the
solidification of YAG is over, the additives still keep a liquid state
and can feed into the shrinkage cavities. Additionally, with a
longer lifetime than YAG melt, the additive melt can accommo-
date gas bubbles and reduce the porosity of solidified YAG
samples.
The effect of additives in limiting grain growth of YAG can be
attributed to depressed grain boundary migration. In high-gravity
combustion synthesis, YAG melt is intersected by the additive melt
and divided into smaller sections. Among these sections, mass
transportation and heat transfer are greatly blocked. In this way,
the grain growth of YAG is restricted into a smaller volume,
reducing the occurrence of huge grains. Moreover, in each small
section, YAG grains are separated by the additive melt and grain
2 3
Fig. 8. XRD patterns of as-prepared Al O samples: (a) pulverized powder; (b–d)
flakes in the surface layer.

2
28
G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
2 3
Fig. 9. SEM images of skeletal crystals in as-prepared Al O samples.
boundary migration is pinned, which is helpful to limit grain
growth.
thick dendrites occurred, as shown in Fig. 12. The smaller grains
did not fully develop or show regular faceted morphology. The
Similar to Al
2
O
3
, YAG also has a high melting entropy of
D
S
m
/
dendrites were composed of ultrafine crystallites below 1 mm, and
R = 14.72 (Table 2) and thus is expected to undergo faceted crystal
growth. This has been verified by SEM observation showing faceted
polyhedron YAG crystals (Fig. 11). Especially, a regular polyhedral
morphology was observed for the crystals growing in cavities with
no spatial constraint.
some crystallites were only ꢄ100 nm. The exact mechanism for the
formation of the dendrites is not clear yet. Perhaps they have
developed from isolated fractions of melt under a large under-
cooling degree.
Besides polyhedral faceted crystals, other crystallization shapes
were also found in the prepared YAG samples. At the boundary
areas or the surface of some large faceted crystals, smaller grains or
3.1.3. MgAl
2 4
O
Other than Al
2 3 2 4
O and YAG, MgAl O is also an important single-
phase ceramic material prepared by high-gravity combustion
Fig. 10. SEM images of fracture surface of Al
2
O
3
samples showing typical feature of cleavage planes.

G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
229
2
Fig. 11. SEM images of as-prepared YAG samples: (a) with no additives; (b) with 5 wt% SiO ; (c–f) with 10 wt% Si–Al–Ca–O glass.
synthesis. MgAl
Al + 3NiO + MgO = MgAl
ed adiabatic temperature is 2476 8C and higher than the melting
point of MgAl (2135 8C). In this way, the melt-casting of bulk
MgAl ceramics can be realized by high-gravity combustion
synthesis.
2
O
4
can be synthesized from the reaction of
Typical SEM images of as-prepared MgAl
2
O
4
ceramics are
2
2
O
4
+ 3Ni. For this reaction, the calculat-
shown in Fig. 13. At the sample surface, a porous structure was
produced, consisting of equiaxed crystals > 100 m and large
cavities. In the center, a relatively dense structure was obtained
with plate-like crystals with a width > 10 m. At the boundaries of
the plate-like crystals, smaller grains and pores existed. The
m
2 4
O
2
O
4
m
Fig. 12. Different crystallization morphologies observed in YAG samples: (a and b) mixed microstructure with fine crystallites distributed at the boundary regions of the large
crystals; (c–f) thick dendrites produced at the surface of large crystals.

2
30
G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
Fig. 13. SEM images of as-prepared MgAl
2
O
4
samples: (a and b) porous structure with coarse crystals observed at the surface of samples; (c) dense structure with plate-like
2 4
crystals obtained in the center of samples; (d) faceted MgAl O crystals with an octahedral shape.
different structures observed at the surface and in the center of
samples are related to different solidification conditions, including
under-cooling degree and spatial constraint for crystal growth.
proportions of the components often induces a transition of
microstructure.
In comparison with single-phase ceramic materials, multi-phase
eutectic ceramics will show more complicated behaviors in
microstructure evolution during solidification. Here, both binary
and ternary eutectic ceramics have been prepared by high-gravity
combustion synthesis, and their microstructure evolution char-
acteristics during solidification from melts are discussed as follows.
MgAl
is between those of Al
MgAl should also undergo faceted crystal growth. This is
2
O
4
has a melting entropy of
m
DS /R = 9.82 (Table 2), which
2 3
O
and YAG. From the Jackson criterion,
2 4
O
confirmed by SEM observation showing faceted MgAl
2 4
O crystals
with an octahedral shape, as shown in Fig. 13(d).
3.2. Multiphase eutectic ceramics
In addition to single-phase ceramics, multiphase eutectic
ceramics can also be prepared by high-gravity combustion
synthesis. Compared with single-phase ceramics, multiphase
eutectic ceramics have two advantages. At first, the eutectic point
of two or more phases is usually lower than the melting point of
each individual phase. That is to say, eutectic systems have lower
melting temperatures than single-phase systems. Eutectic melts
have a longer lifetime, which is desirable for phase separation and
removal of gas bubbles. On the other hand, multiphase eutectic
ceramics can offer more opportunity for tailoring the microstruc-
ture. In single-phase ceramics, only one phase is involved and the
microstructure is characterized by grain size and morphology of
that phase. In multiphase eutectic ceramics, however, two or more
phases co-exist, and the microstructure is affected by more
factors, including the proportion, size and morphology of each
phase. For example, in eutectic ceramic materials, a change in
2 3
Fig. 14. XRD pattern of the YAG–Al O eutectic ceramic sample (YA-1).

G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
231
3.2.1. YAG–Al
2 3
O
The first example of binary eutectic ceramics is the YAG–Al
2 3
O
system. In this system, two compositions were studied, including a
eutectic one (YA-1) and a hypoeutectic one (YA-2), as shown in
Table 1. For both the compositions, the calculated adiabatic
temperatures are much higher than the eutectic temperature as
well as the melting points of Al
Fig. 14 shows the XRD pattern of the sample YA-1, which was
composed of major YAG and minor Al . This phase composition
2 3
O and YAG.
2 3
O
is consistent with that expected from the binary phase diagram
shown in Fig. 15. By calculation from the nominal chemical
composition, the volume fractions of YAG and Al
1%, respectively. The sample YA-2 also consisted of YAG and
Al , except that the major phase was Al . By calculation, the
volume fractions of Al and YAG in YA-2 are 84% and 16%.
2 3
O are 59% and
4
2
O
3
2 3
O
2 3
O
In the samples YA-1 and YA-2, different microstructures were
obtained, as shown in Fig. 16. YA-1 showed typical fine eutectic
2 3 2 3
Fig. 15. A simplified phase diagram of the Al O –Y O binary system.
structures with interphase spacings < 1
mm, but in YA-2 large
2 3
Al O crystals were produced and eutectics occurred only at the
Fig. 16. SEM images of YAG–Al
with large primary Al crystals and fine eutectics observed in YA-2 with a hypoeutectic composition; (f and g) EDS spectra for Al
and ‘‘2’’, respectively in (e).
2
O
3
eutectic ceramic samples: (a and b) uniform and fine eutectic structure observed in YA-1 with a eutectic composition; (c–e) mixed structure
2
O
3
2 3
O
and YAG crystallites, as marked with ‘‘1’’

2
32
G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
2 3
final structure of YA-2 is made up of primary Al O crystals and
YAG–Al eutectics at the boundaries of the primary crystals.
2 3
O
3
.2.2. Al
Another example for binary eutectic ceramics prepared by
high-gravity combustion synthesis is the Al –ZrO system. From
the Al –ZrO phase diagram, the eutectic composition was
selected, as shown in Table 1, where 1 mol% Y was added to
stabilize the tetragonal structure of ZrO . For the eutectic
2 3 2
O –ZrO
2
O
3
2
2
O
3
2
2 3
O
2
composition, the calculated adiabatic temperature is 2678 8C
and higher than the eutectic point of 1860 8C.
Fig. 17 shows XRD pattern of the as-prepared Al
eutectic ceramic sample (AZ). From the pattern, the sample was
composed of Al and ZrO . In the sample, both polymorphs of
ZrO were observed, with either tetragonal or monoclinic lattice
structure. The occurrence of monoclinic ZrO can be attributed to
the relatively low concentration of Y . From the nominal
chemical composition of the sample AZ, the volume fractions of
Al and ZrO are calculated to be 67% and 33%, respectively.
The microstructure features of the Al –ZrO eutectic ceramic
2 3 2
O –ZrO
2
O
3
2
2
2
2 3
O
2 3 2
Fig. 17. XRD pattern of the Al O –ZrO eutectic ceramic sample (AZ).
2
O
3
2
2
O
3
2
boundaries of the large crystals. The different microstructures
observed in YA-1 and YA-2 should be attributed to the different
sample existed on different size scales, as shown in Fig. 18. On a
larger scale, the sample consisted of colonies with an average size
volume fractions of Al
phase diagram, for YA-1 with a eutectic composition, the eutectic
’’ takes place during the whole
solidification process, leading to a homogeneous eutectic struc-
ture. For YA-2 with a hypoeutectic composition, during solidifica-
2
O
3
and YAG in them. According to the binary
of >10
mm. In the center of each colony, the eutectic structure was
unaffected, but at the boundaries the eutectic structure became
irregular. The interphase spacing in the center (<1 mm) was much
smaller than that at the boundaries.
During the solidification of eutectics, the interphase spacing is
determined by the balance of two opposite energy contributions.
On one hand, transverse diffusion of solutes must occur over
greater distance for thicker lamellae or fibers, therefore, favoring
smaller interphase spacings. On the other hand, the interfacial
energy associated to interphase boundaries will increase as the
width of lamellae or fibers decreases, thus favoring larger
reaction ‘‘liquid ! YAG + Al
2
O
3
9
tion primary Al
precipitation of Al
2
O
3
crystals will be precipitated first. The
causes a segregation of Y in the liquid,
2
O
3
2
O
3
whose composition gradually changes along the liquidus line (AE
line in the phase diagram in Fig. 15). When the composition of the
liquid reaches the eutectic point, eutectic reaction occurs with
further cooling and continued solidification. Consequently, the
2 3 2
Fig. 18. SEM images of Al O –ZrO eutectic ceramic samples: (a and b) an overview of colonies; (c) and (d) fine eutectic structures in a colony.

G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
233
interphase spacings. In general, the interphase spacing is related
with under-cooling degree and growth rate, which can be
expressed as:
2
lDT ¼ C
where
1
and V
l
¼ C
2
;
l
is the interphase spacing,
DT is the under-cooling degree,
V is the growth rate, C
formulae, a greater under-cooling degree is expected in the center
1
and C are the constants [28]. From the
2
than that at the boundaries of the colonies in the Al
2 3 2
O –ZrO
eutectic ceramics.
According to the model proposed by Hunt and Jackson [28], the
microstructure of binary eutectics depends on the melting
entropies of the components. If the large-volume component
has a high melting entropy (
melting entropy ( < 2), a ‘‘complex regular’’ microstructure will
be produced, which shows many features of the lamellar or fibrous
eutectic microstructure. In this work, the Al –ZrO eutectic
sample possesses a phase composition of 67%Al + 33%ZrO in
volume fraction. The large-volume component of Al
a > 2) and the minor one has a small
a
2 3 2
Fig. 19. XRD pattern of the YAG–Al O –ZrO eutectic ceramic sample (YAZ).
2
O
3
2
2
O
3
2
2
O
3
has a
has a
higher melting entropy of
lower melting entropy of
D
S
m
/R = 5.74, and the minor ZrO
2
the volume fraction of ꢄ30%. In the Al
sample prepared in this work, the volume fraction of the minor
component (ZrO ) is 33%, which is almost the critical point for the
2 3 2
O –ZrO eutectic ceramic
m
DS /R = 3.55. In this case, the eutectic
sample showed a complex microstructure, where both fibrous and
lamellar features were observed, as shown in Fig. 18(c and d). In
2
transition between the two structures. As a result, both fibrous and
lamellar structures were obtained in the sample.
the fibrous structure, fine ZrO
were embedded in a continuous Al
2
fibers with a width of 100–200 nm
matrix. The lamellar
O
3
Another factor affecting the microstructure of eutectics is the
growth rate. During the solidification of eutectic ceramics,
microstructure evolution is generally related with the thermal
gradient, concentration gradient, and growth rate. For given
thermal and concentration gradient values, the solidification
mechanism depends on growth rate. When the growth rate is
below a critical value, planar growth front is stable, and with
increasing growth rate shallow cells will occur as a secondary
phase. The microstructure transition from coupled to cellular and
then to shallow cells with increasing growth rate has been
2
structure was somewhat incomplete, which was more exactly a
mixture of lamellar and fibrous features.
The microstructure of binary eutectics is usually connected
with the volume fractions of the two components. It was reported
that [29], if the surface energies are isotropic or with no significant
anisotropy, the fibrous structure is stable when the volume
fraction of the minor component is below 32%, and when the
fraction is >32% a lamellar structure will be produced. This
prediction has been proved by experimental results on several
metal eutectics, where a lamella-to-fiber transition takes place at
2 3 2
observed in Al O –ZrO system [11]. In this work, the discrepancy
2 3 2
Fig. 20. SEM images showing the eutectic structures in YAG–Al O –ZrO ternary eutectic ceramics.

2
34
G. Liu et al. / Materials Research Bulletin 47 (2012) 222–234
in the growth rates of different parts is probably one of the reasons
responsible for the diversity of microstructures observed in the
For most single-phase ceramic materials like Al
2 3
O , YAG and
MgAl with high melting entropies, a faceted crystal growth
2 4
O
Al
2
O
3
–ZrO
2
eutectic ceramic sample.
mode was operative in solidification, which is consistent with
Jackson’s criterion on interface roughness. The solidification
process strongly depended on cooling conditions, and the
discrepancy in cooling conditions for different parts of samples
resulted in a diversity of microstructures. As an example, a strongly
textured layer composed of orderly arranged faceted crystals was
3.2.3. YAG–Al
2 3 2
O –ZrO
Besides binary eutectic systems, YAG–Al
2 3 2
O –ZrO ternary
eutectic ceramics have also been prepared by high-gravity
combustion synthesis. The eutectic composition was determined
according to the Al
2
O
3
–Y
2
O
3
–ZrO
2
phase diagram and reported in
2 3
observed at the surface of Al O samples. In the solidification of
Table 1. For the eutectic composition, the calculated adiabatic
temperature is 2678 8C, which is much higher than the eutectic
point of 1715 8C.
YAG ceramics, the incorporation of proper glass additives was
helpful to decrease the porosity and reduce grain growth.
Compared with single-phase ceramics, multiphase eutectic
ceramics have lower melting temperatures, which is desirable for
phase separation and removal of gas bubbles. In addition, eutectic
ceramics consist of multiple components and can provide more
opportunity for tailoring microstructures. From the experimental
results, the microstructure of eutectic ceramics was related with
the volume fractions of components, and various eutectic
structures were produced such as lamellae, fibers, and three-
dimensional interpenetrating frameworks. The eutectic struc-
tures showed interphase spacings on a submicron scale. In the
eutectic ceramics, no cracks were found at the interface between
different phases, indicating an excellent interfacial bonding
strength.
The XRD pattern of as-prepared YAG–Al
sample (YAZ) is shown in Fig. 19. The sample consisted of YAG,
Al and tetragonal ZrO , with no occurrence of monoclinic ZrO
By calculation from the nominal chemical composition, in the
sample YAZ, the volume fractions of YAG, Al and ZrO are 49%,
7% and 14%, respectively.
In the YAG–Al –ZrO
structure was observed, as shown in Fig. 20. Most colonies had a
size of >30 m and some were larger than 100 m. In the center
2 3 2
O –ZrO eutectic
2
O
3
2
2
.
2
O
3
2
3
O
2 3
2
eutectic sample, a typical colony
m
m
parts of the colonies, two microstructure features were found,
including lamellar and three-dimensional interpenetrating struc-
tures. The eutectic structures were composed of YAG and Al
crystallites with interphase spacings of <1 m, and ultrafine ZrO
particles were distributed at the interface between YAG and Al
2
O
3
m
2
2
O
3
.
Acknowledgements
At the boundaries of the colonies, larger interphase spacings were
observed.
This work was supported by National Natural Science Founda-
tion of China (Grant Nos. 50932006, 51002163, and 51001111) and
Beijing Natural Science Foundation (Grant No. 2112043).
In the YAG–Al
have higher melting entropies and ZrO
sample of YAZ, YAG was the major phase with a volume fraction of
2 3
O
–ZrO
2
ternary eutectic system, YAG and Al
2 3
O
2
has a lower one. In the
4
1
Al
9%, and the minor Al
4%. According to the model proposed by Hunt and Jackson [29],
and ZrO will form lamellar and fibrous structures,
2 3 2
O and ZrO had volume fractions of 37% and
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