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High Resolution Electron Microscopy Investigations
of Interface and Other Structure Defects in Some Ceramics
Shanghai Institute of Ceramics, Chinese Academy of Sciences, Shanghai, 200050, China
fullerene; sialon; silicon nitride; high Tc superconductors; ceramics; bioceramics;
human enamel; hydroxyapatite; electron microscopy; high-resolution electron
Interface, grain boundary, and other structure defects are the most important
structural factors to affect the properties of ceramics materials. The present paper shows the
relationship between the properties and those structure features such as grain boundaries, phase
boundaries, interfaces, twins, intergrowths, dislocations, point defect aggregates, order-disorder,
and other structure defects in different kinds of ceramics materials. At present this research covers:
C60, sialon-based ceramics (a-sialon/SiC(w) composite, Y-a-sialon/b-sialon composite), high Tc
superconductors (YBa2Cu3O7, YBa2Cu4O8, Bi2Sr2CaCu2O8, Bi2Sr2Ca2Cu3O10), and bioceramics
(hydroxyapatite, chlorapatite) and so on. The structure features mentioned above were characterized by high-resolution electron microscopy; so the structure details are at an atomic level and the
related physical, chemical, engineering, even biological phenomena can be understood at an atomic
and molecular level. Microsc. Res. Tech. 40:177–186, 1998. r 1998 Wiley-Liss, Inc.
Fullerenes with a soccerball-shaped C60 molecule
structure is a brand-new form of crystalline carbon in
addition to graphite and diamond, and it has been
widely investigated for its unique characteristics. Until
now, both scanning electron microscope and transmission electron microscope (TEM) have been effectively
used to study morphology, microstructure, crystal defects, and phase transformation (Banhart et al., 1992;
Dravid et al., 1991; Muto et al., 1993; Verheijen et al.,
1992; Wang and Buseck, 1991).
As to the crystal structure of solid C60, the molecules
can stack either in a hexagonal close-packed (hcp)
lattice with a 5 1.002 nm, c 5 1.6636 nm (Krätschmer
et al., 1990), or a face-centered-cubic (fcc) lattice with
a 5 1.4172 nm (Fleming et al., 1991). Fcc and hcp
structures are both stable at room temperature,
whereas, from the energetic point of view, the fcc
structure is much more stable. Moreover, the differences in the structure result from the different conditions under which crystals grow (Haluska et al., 1993;
Ven Tendeloo et al., 1991; Verheijen et al., 1993). To
some extent, (0001)hcp face can be considered as the
structure defect caused by the stacking faults of (111)fcc
face in fcc C60 crystals. We confirmed some structure
features mentioned above and made some new discoveries that we will discuss here.
Sialon-based ceramics have been recognized as
materials with a large potential in applications at
elevated temperatures where hardness, thermal shock
resistance, and high bending strength are important.
In order to quicken these applications, current research
on sialon-based ceramics is attempting to further improve their mechanical properties at either normal or
high temperature by focusing on two aspects (Braue,
1992; Liu et al., 1994): (1) correct selection of the
composition of interface (grain boundary phases) with
high refractoriness and good oxidation resistance that
can be crystallized thoroughly or partly during cooling
or heat-treatment processing to improve grain boundary strength; and (2) inducing in situ or second phase
(particle, whisker, fiber, and so on) reinforcement mechanisms in these ceramics to enhance mechanical properties, in which in situ reinforcement mechanism gives
the so-called mosaic texture effect caused by the different phases compounded microstructure. The second
phase reinforcement makes crack deflection, crack
bridging, debounding, microcracking, and/or whisker
pull-out possible, thereby increasing energy absorption
when cracks propagate through the composites.
At this point, the influence of these mechanisms
would depend strongly upon the nature of the interface
among different phases. It is, therefore, of interest to
find out different phase boundary configurations in
sialon-based ceramics. For years we have been working
on grain boundaries including phase boundary in Ln-asialon/b-sialon (Ln 5 Y, Nd, and Yb) multiphase ceramics (Liu et al., 1994, 1995a; Wen and Feng 1984, 1986)
as well as Y-a-sialon/SiC(w) composites (Nordberg et
al., 1993) and some important results will be presented
High Tc superconductors remain an attractive
topic of research although the materials have been
investigated for nearly 10 years since they were found.
This is related in part to finding prominent ways to
obtain superconductors with higher Jc for the applications and, in part, to understanding the unique characShulin Wen’s current address is Material Science Division, 212, Argonne
National Laboratory, 9700 S. Avenue, Argonne, IL 60943.
*Correspondence to: Quian Liu, Shanghai Institute of Ceramics, Chinese
Academy of Sciences, 1295 Dingxi Road, Shanghai, 200050, China.
Received 22 March 1996; accepted in revised form 8 April 1996
teristics of new kinds of superconductors, such as
mercury cuprates, C60-based as well as oxycarbonatescontaining materials.
However, there are still some problems that should
be solved before practical applications can be achieved.
One of these is to enhance the Jc by improvement of
microstructure features. In fact, the superconductivity
acts structurally at electronic and atomic levels and is
characterized by Tc. However, Jc acts at the micro-level
in the structure when defects are involved. From the
viewpoint of microstructure and superconducting physics, some of the microstructure features, such as dislocations, stacking faults, and structural transformation,
are very effective flux pinning sites for enhancing
critical current density in highly anisotropic oxide
superconductors according to theoretical and technical
At this moment ceramics-sintered materials only
give a Jc value as high as 103–104 A/cm2, which is not
suitable for practical usage of high Tc superconductors.
The main problem for the lower Jc is due to the weak
link between the superconducting grains. In oxide
superconductors, only Cooper-pairs can tunnel over a
nonsuperconducting region 0.5 nm (parallel to the
C-axis) and 2 nm (perpendicular to the C-axis), so the
grain boundaries with varied thickness would form a
weak link. At this point, the microcracking resulting
from anisotropic thermal expansion could form the
weak link. Non-stoichiometry near grain boundaries
could also form a weak link; especially high-angle grain
boundaries with anisotropy in conductivity in crystal
must form a weak link. On the other hand, some grain
boundaries serving as pinning centers can contribute a
potential effect for Jc as seen in the case of thin films
that can have a Jc as high as 107 A/cm2, as the pinning
centers are the regions of inhomogeneities in the lattice, resulting in a local change in Tc or in the correlation length and acting as strong flux pinning centers
(Van der Beek et al., 1995). We began this kind of
research early on (Wen et al., 1987, 1988) and some new
results of our progress will be presented here.
Bioceramics is a very important part of biomaterials, which is a new and very promising both for
research and development. The reason for even more
interest in bioceramics is its close relation with biomineralization occurring in the human body. Mineralization is the process in which living organisms convert
various ions from solution into biominerals. The investigation of interface structure in bioceramics is most
important due to its special relation with the abovementioned process of biomineralization in human body.
The structures of interfaces in bioceramics differ greatly
from those synthetic bioceramic materials due to cellular activities that create and promote necessary physiochemical conditions for the nucleation and growth of
bioceramics from solution. We have performed this kind
of research for years (Wen et al., 1989) and some recent
results will be described here.
Our TEM observations of crystalline morphology and
microstructure of C60 crystals, prepared from C60/
benzene solution, showed that the morphology is controlled by the growth rates of 51116, 51006, and 51106
Fig. 1.
HREM image showing C70/C60 intergrowth.
crystal faces of the crystal, based on the selection rules
of crystal planes. The crystal has an fcc structure with
some 51116 microtwinning as well as stacking faults, but
is easily damaged by the electron irradiation. These
kinds of defects and damage may cause the local
transformation from fcc to hcp structure (Liu et al.,
1993). Our HREM investigation has also shown that
C70 molecules intergrow with C60 molecules at the
content level of a few percent. Figure 1 is a HREM
micrograph indicating the intergrowth between C70 and
Compared with hcp graphite and cubic diamond,
allotropes of carbon, C60 is intriguing because of its
spherical shape, resulting in different electronic, chemical, physical, and crystallographic properties. Both
electron energy loss spectroscopy (EELS) and energy
loss near edge structure (ELNES) are essential methods for characterizing materials mostly by using a TEM
equipped with an electron energy analyzer. In our
present work, the EELS and ELNES were used to
characterize and compare two kinds of carbon, C60 and
The bulk plasmon excitation of the valence electrons
in C60 crystals has shown two characteristic lossenergies of 6.51 and 26.26 eV ([111]fcc), but 7.47 and
27.32 eV for graphite ([0001]hcp), which are consistent
with other authors’ reports (Egerton, 1986; Saito et al.,
1991a). The bulk plasmons are raised from two distinct
groups of valence electrons in C60 and graphite. For C60,
each atom has one p-electron and three s-electrons,
with s-orbital along the neighboring atoms and porbital extending outward and inward towalds the C60
cage. Furthermore, it is probable for the p-orbital to
extend more outside the cage than the inside. It is the
weakly bound p-electrons that are attributed to bulk
plasmon excitation at the lower energy level, whereas
all the valence electrons (p- and s-electron) are attributed to the higher energy level excitation. Carbon
atoms in graphite are bonded between the carbon layer
planes by 2pz-orbital, evenly distributing above and
below the planes. The p-electrons are free to move
throughout the planes of C6 rings. The strongly bound
s-electrons join each carbon atom with three others in
co-planar bonding. Similar to C60, the energy-loss spectrum of graphite at a lower level is identified to be due
to the weakly bound p-electrons, and the spectrum at a
higher level due to the p- and s-electrons.
With regard to the ELNES of C60 and graphite, as the
ELNES may reflect a density of states above the Fermi
level to some degree, K-shell excitation in C60 is very
different from that of graphite. Figure 2 shows the
energy-loss spectra in the carbon K-shell excitation of
the solid C60 and graphite (Saito et al., 1991b). There is
a sharp peak at 285 eV and a band ranging from 290 to
305 eV, in which the former is assigned to a transition to
antibonding p* orbital and the later to a transition to
antibonding s* orbital, respectively, the same assignment as graphite. The p* resonance at 285 eV is
considered to be unsaturated(sp or sp2) carbon bonds,
which implies sp2 hybridization in the C60 molecule.
A distinguishable ELNES feature of C60 crystal is a
shoulder observed at 289 6 1 eV (indicated by arrow A
in Fig. 2). According to total-energy calculation of solid
C60 (Saito et al., 1991), the dispersion of the energy
bands is very small due to weak interaction between
C60 molecules in the solid. Therefore, the 289 eV
shoulder corresponds to unoccupied p* states and
reflects the high degeneracy of energy levels in C60
owing to its extremely high symmetry. Since the interpretation of ELNES is based on the density of states in
the energy band diagram, which is related to a longrange order or disorder in a solid, ELNES can provide
information about the arrangement of atoms in the
local areas as well as extensive knowledge on the
short-range extended X-ray absorption fine structure
(EXAFS). Comelli et al. (1988) have reported that the
width of the s* resonance band in X-ray absorption
near edge structure (XANES) of amorphous carbon
increases toward that of graphite with the increase of
annealing temperature, i.e., the increase of the degrees
of order. Therefore, the s* band of solid C60 with an
intermediate value between those for graphite and
amorphous carbon suggests that the structural order in
the solid C60 is lower than that in graphite but higher
than that in amorphous carbon.
Furthermore, in our ELNES examination about the
effect of electron beam irradiation on the structural
stability of solid C60([111]fcc) and graphite([0001]hcp), it
has been found that with increasing irradiation time,
p* and s* bands of C60 and graphite all shift to the
lower energy-loss sides. For C60, the relative shifts of s*
band are 1.0, 2.8, and 5.6% after 10-, 20-, and 30-second
irradiation (10 PA/cm2), respectively, with 2.0, 5.0, and
10.0% shift in p* band at the same time. For graphite,
the relative shifts of s* band are 2.0 and 16.0% after 10and 20-second irradiation, respectively, with 8.0 and
19.0% shifts in p* band at the same time. These
phenomena suggest that with the increase of irradiation time, the order of the atomic arrangement in solid
C60 and graphite gradually decreases, or order-disorder
transition occurs. Comparatively, the p* bands shift to
the left more than s* bands, which means p* bonds
(weakly bound electrons) are easier to be broken after
Ln-a-sialon/bsialon(Ln 5 Y, Nd, and Yb) multiphase ceramics, the
mixture of powder of Si3N4, SiO2, Al2O3, AlN, Y2O3,
Nd2O3, and Yb2O3 was used accordingly as the starting
materials depending on the sintering additive (Y2O3,
Nd2O3, Yb2O3) selected.
Y-a-sialon/b-sialon ceramics were sintered by gas
pressure sintering (GPS) at 1900°C for 3 hours and
Fig. 2. Electron energy-loss spectra of graphite (a) and solid C60 (b)
in the region of K-shell excitation. Relative intensity in the ordinate
axis represents the electron intensity relative to that of the zero-loss
peak (the zero-loss peak height is normalized to unity).
then the as-sintered specimens were heat-treated at
1400°C for 24 hours. In this material system, the
a/b-sialon, as two phase composites, are compatible
with Y3Al5O12(YAG) as an interface phase, forming a
compatibility tetrahedron Si3N4-a-sialon-b-sialon-YAG.
Based on the phase diagram researches (Sun et al.,
1991; Yen and Sun, 1993), chemical compositions and
microstructures falling into this tetrahedron are expected to have excellent mechanical properties because
of the co-existence of the refractory grain boundary
phase YAG and the mosaic texture composed of equiaxed
a-sialon as well as elongated b-sialon grains. As a
consequence, the ceramics thus prepared can maintain
high bending strength at a temperature as high as
Nd- and Yb-a-sialon/b-sialon multiphase ceramics
were fabricated by hot-press sintering to evaluate the
improvement in microstructure and mechanical properties that can be achieved by from a8 to b8 transformation induced by heat treatment (1450°C, 96 hours). For
Nd-doped composites, the stability of the a-sialon phase
is very sensitive to thermal history, so that the a-sialon
phase transforms almost totally to the b-sialon phase
after undergoing a long heat treatment with Nd-M8
(aluminum-containing melilite, Nd2Si3O4N3) and
NdAlO3 left at pockets of grain junctions as grain
boundary crystalline phases. The high fracture tough-
ness (K1c) is characteristic of these materials derived
from more elongated b-sialon grains existing in the
specimens. In contrast, Yb2O3 is a kind of effective
additive for the a-sialon phase stability, i.e., a-sialon
can keep untransformed in Yb-system, with only a few
b-sialon phases in the composites plus Yb-garnet
(Yb3Al5O12) and Yb-J-phase (Yb4Si2O7N2) as grain
boundary crystalline phases.
As a result, Yb-a-sialon/b-sialon ceramics possess
higher hardness (Hv) owing to equiaxed a-sialon grains
developing in the materials. With regard to the tendency of a8 to b8 transformation, besides the effect of
time and temperature on the treatment period, the
kind of additive is another important factor to produce
stable a-sialon phase structure. Since the transformation is regarded as a process of solution-diffusionprecipitation, i.e., solution of a-sialon, diffusion of atoms through a grain boundary oxynitride liquid,
followed by precipitation of b-sialon from the liquid, the
structure of a-sialon is partially or wholly dismembered
by solution and reconstructed by precipitation. During
the heat-treatment, Nd-a-sialon is unstable relative to
b-sialon and M8, probably because on cooling down, the
large interstices (one located around the origin with a
radius of 0.15 nm, the other with a dimension of 0.1 3
0.24 nm in the middle of the unit cell; Grun, 1979) in the
a structure are not big enough to accommodate the
relatively large Nd cation (RNd31 5 0.0995 nm). Therefore, the solubility of Nd cation at 1450°C might be zero
in equilibrium. But because Yb cation (RYb31 5 0.0886
nm) is smaller, it has quite a high solubility in the
structure at the same temperature to enter the channels to prevent a-sialon from being transformed. That is
the reason why Nd-a-sialon is very unstable and much
more Nd-containing grain boundary phases exist in the
The SiC whiskers incorporated Y a-sialon is well
known as SiC whiskers reinforced a-sialon composites,
usually prepared by hot pressing sintering at 1750–
1800°C. The indentation fracture toughness and hardness in the reinforced a-sialon/SiC(w) composites were
greatly increased, especially the toughness values perpendicular to the whisker orientation. From the observations of SEM and TEM, the dominating toughening
mechanisms are probably crack deflection and some
crack bridging; there is no extensive whisker pull-out
effect although a few protruding whiskers are involved.
This indicates a strong SiC(whisker)/sialon interfacial
bond. The absence of an amorphous interface layer
confirmed by the HREM study supports this hypothesis. Other reports have indicated that the presence of
thin amorphous layers between SiC whisker and sialon
matrix gives a weak-link boundary, but its absence
makes the bond strong (Braue, 1992; Braue et al., 1990;
Das Chowdhury et al., 1992). Our observations show
that the interfaces between SiC(whisker) and sialon in
this composite were mostly a direct and strongly bonded
contact; only a minor part of these interfaces consisted
of a 2–5-nm-thick amorphous layer. Most grain boundary phases accumulate at grain junctions. Figure 3 is
our HREM image showing the interface between sialon
and SiC, where the grain boundary phase with some
amorphous layer (see arrowheads) is located at the
pocket of grains, but the most interfaces between SiC
Fig. 3. HREM image of a-sialon/SiC(w) composites. Arrowheads
indicate the grain boundary phase located at the interface between
sialon and SiC whisker.
and sialon grains are ‘‘clean’’ only with a mismatched
atoms arrangement.
As for the structural ceramics, like sialon-based
ceramics, the grain boundary influences many processes (such as diffusion and mass-transfer) and properties (such as creep and corrosion). Generally speaking,
both high- and low-angle grain boundaries exist in
ceramics, and the structure of both kinds of grain
boundaries seems to be less compact owing to the
existence of electrostatic potential and instability of the
grain boundary. Furthermore, in the case of sialonbased ceramics, the grain boundary configuration and
composition are pertinent to their mechanical properties. As is well known, grain boundaries usually exhibit
disordered microstructure. From HREM micrographs,
grain boundary configurations can be briefly classified
into two types: dihedral angle 2QÞ 0 or 2Q 5 0. In the
light of thermodynamics analysis, dihedral angle depends on the relationship between the interfacial (gl)
and grain boundary (gb) energies. If the Q # 30°, the
grain boundary phase is stable at pockets between
grain junctions. During cooling or heat-treatment processing, some grain boundary phases can be crystallized, but some remain amorphous, depending on the
chemical composition of grain boundary phases and
other effects from the kinetic process.
Meanwhile, the size and shape of grain boundary
phases gradually change from the wetting liquid state
to the crystalline state or ‘‘frozen’’ glassy state. The
typical grain boundary phases are shown in Figure 4
from HREM observations, which shows that the movement of the grain boundary would be towards the
center of the curvature in the period of crystallization,
and some steps of crystal growth can be seen in Figure
4. Figure 5 shows the lattice image of two grains and
their boundary phase in Y-a-sialon/b-sialon ceramics,
indicating that the original amorphous phase has been
crystallized after heat treatment.
For 2Q 5 0, this kind of configuration has high
energy due to its large interface area. Unless there is
enough space for the grain boundary phase to be
crystallized as illustrated in Figure 5, the grain boundary will remain amorphous since the stress field at the
interface, caused by the transition from the ordered
Fig. 4. HREM image of the outline of the grain boundary phase
located at the pocket of three grains in Y-a-sialon/b-sialon ceramics,
where some steps of crystal growth can be seen.
Fig. 5. Lattice image of two grains and their boundary phase in
Y-a-sialon/b-sialon ceramics, showing that the original amorphous
phase has been crystallized after heat treatment.
structure in the crystal to the disordered structure in
the glass, can theoretically be explained as a kind of
balancing surface force to stabilize the grain boundary
configuration (Clarke, 1987, 1989).
It is known that sialon-based ceramics show a high
bend strength at 1300°C (about 600 MPa). Therefore,
deformation behaviour of these materials at high temperatures relies not only on the properties of the
individual phase, but also on the relative distribution of
the phases under the stress or temperature. There are
at least two kinds of kinetic processes involved in the
formation and propagation of a crack at high temperatures under stress (Tsai and Raj, 1980). They are: (1)
grain boundary plastic sliding inducing a stress concentration at the pocket between grain junctions; and (2)
crack nucleation and propagation at grain boundary
and the pocket.
Therefore, the resistance of sliding increases with an
increase in 2Q value, so that the parallel grain boundary is more beneficial to plastic sliding. In this sense, it
is possible to precisely control the amount, distribution,
and configuration of grain boundary phases for improvement of mechanical properties at high temperature.
The grain boundaries in YBCO superconductor systems can be briefly classified into three types: (1) twin
boundary, (2) coherent boundary, and (3) incoherent
boundary with second phase (Wen et al., 1987). Twinning occurs commonly in the orthorhombic structure.
There are many kinds of twin boundaries in YBCO
ceramics, where the 110 twin is the most popular from
our HRTEM observations (Wen et al., 1988) in which
every O1-Cu1-O1 chain can keep its continuity through
110 twin boundaries and without interruption, rendering high Jc values. Furthermore, a kind of tilt boundary
with an angle of 52° along (023) plane was found in
YBCO, in which the perovskite layers could be maintained from one grain to another, but the periodic unit
in the grain boundary is just twice that for the perovskite layer. Moreover, the structure in the grain boundary area would be more open than that in the normal
area, so that the structure symmetry becomes lower
and the unit cell in the boundary area becomes twice as
large. In spite of the differences in the structure and
thickness for a variety of twin boundaries, they are able
to keep O1-Cu1-O1 chains to maintain continuity and
provide effective pinning centers due to the domain size
of 10 to 100 nm.
As the coherent boundaries directly connect two
grains without any grain boundary phases in between,
this kind of grain boundary can keep the continuity of
O1-Cu1-O1 chain that is believed to contribute to high
Jc. As to the incoherent boundaries, for example in
RBa2Cu3O7 ceramics, they have second phases, either
crystalline second phases or an amorphous layer. Obviously, this kind of grain boundary is an obstacle for high
Jc values as the O1-Cu1-O1 chains could not keep their
continuities at these boundaries.
Viewed from another angle, the grain boundary is a
kind of disordered structure, so the presence of dislocations (Kramers et al., 1990) is unavoidable in these
areas. They are mostly [010], [001], and 1/2(a1b)[001]
edge dislocations. For example, the (001) twin boundaries often terminated at (001) dislocation arrays. Some
dislocations could be potential pinning centers as reported in a few high Jc materials with a high density of
dislocations (Nakahara et al., 1989). Generally, the
superconducting current preferentially flows in the
Cu-O layer, so the potential pinning center, i.e., dislocation, should provide a potential well for the flux lines in
the Cu-O layer.
Apart from the above situation, grain boundaries can
also function as reaction site in preparation of oxide
superconductors. It was discovered in our research that
the reaction of Y-123 (YBa2Cu3O7) 1 CuO = Y-124
(YBa2Cu4O8) occurs more easily in the area near the
interface between a Y-123 thin film and a SrTiO3
substrate even with no extra CuO supplied during
Y-123 film formation as shown in Figure 6. It could be
seen that the area of the Y-123 film (T) near substrate
(S) consists of Y-124 phase. It is obvious that the
structure defects such as dislocations and lattice mismatch between the substrate and films have induced
the Y-123 = Y-124 reaction, implying that the interface
between Y-123 and SrTiO3 plays a role as starting
reaction sites.
Fig. 7. HREM image of a grain boundary acting as a reaction site
during the solid state reaction Y-123 1 CuO = Y-124.
Fig. 6. HREM image of interface between the YBCO thin film and
SrTiO3 substrate showing the Y-124 phase produced at the area near
the interface due to the structure defects formed there.
It was also found in our experiment (Wen et al., 1994)
that the solid state reaction of Y-123 1 CuO = Y-124
occurred along grain boundaries. The samples for this
research were made from mixed powders consisting of
CuO and Y-123 phase with a long time heat-treatment.
Our HREM observations showed that the reaction
mostly starts from the structure defects at grain boundaries. This is probably due to the fact that the atoms
near the grain boundary are much easier to be activated at reaction temperature than those inside of
grains. Figure 7 shows at an atomic level how the
reactive CuO diffused into Y-123 structure along the a-b
plane of grain boundary forming Y-124, even Y-125
structure, replacing single O1-Cu1-O1 chain on the a-b
plane in Y-123 by double and even triple Cu-O chains.
Consequently, the lattice distortion accordingly occurred at the growth front of Y-124 and Y-125 phases.
In comparison with the stoichiometry Y-123, the
Y-deficient ones also have an orthorhombic structure
and the Y-123 main crystalline phase remains, but with
the shortage of Y, Y atoms should come out from some
unit cells in the lattice and extremely fine-scale defects
are anticipated to be possible pinning sites, so as to
increase the bulk pinning force (Fp) and Jc . The improvement of Jc and Fp should be related to the microstructural defects as flux pinning centers within Y-deficient
specimens. HREM observation has revealed that the
main microstructural feature is dominated by Y-123
layer-like assemblages with several kinds of defects in
the specimens: (1) highly dense, fine-scale, fault-like
defects along the (001) basal plane with a thickness of
several atomic layers, shown in Figure 8, and (2)
Fig. 8. HREM image of high-dense, fine-scale, and fault-like
defects in the Y-deficient samples.
localized superstructures, especially existing at grain
boundaries, shown in Figure 9. The fault-like defects
are considered to be effective pinning centers as they
have proved to have useful pinning roles in Y-123 bulk
materials with high Jc and Fp, which were prepared by
phase decomposition and atomic substitution routines.
The localized superstructure is related to oxygen nonstoichiometry, especially at grain boundaries as grain
boundaries play an important role in gas diffusion.
These localized superstructures imply the existence
of disordered areas with some small defects or local
variations of oxygen content. Generally, the principal
directions of the oxygen ordering resulting in superstructures are [001]p, [110]p, [210]p, and [310]p (Raveau et al.,
1991). In fact, the superstructures only exist in very
tiny areas due to local fluctuation of oxygen, forming
local new periodicity, as shown in Figure 9.
The inclusions in the Y-123 phase material are usually not considered to be beneficial to the superconduc-
Fig. 9. HREM image of a local superstructure at the grain boundary of the Y-deficient samples.
tivity because of their large size. However, some inclusions, like fine Y-211(Y2BaCuO51x) particles in the
Y-123 matrix could serve as a pinning center (Murakami, 1990), since these inclusions can either induce
a great amount of dislocations, stacking faults, and
other fine-scale defects along the inclusions or have
some special direction.
We have observed some inclusions with a diameter of
30–50Å within the grains (Wen et al., 1988). They have
very special direction; their direction may be easier to
be normal to the Cu-O layer forming a rather suitable
potential well for flux line in the Cu-O layer.
The remarkable feature of the microstructure in
oxide superconductors, not only in the old species such
as BSCCO, YBCO, but also in copper oxycarbonates
families as well as C60-based ones, would be the intergrowth of multiphases. The difficulty for obtaining pure
phase or removing intergrowth from structure comes
from the very small difference in thermodynamic stability between homologous series.
The intergrowth in (BiO)2Sr2Can21CunO2n12 with n 5
1,2,3,4,5,7,9 make it difficult to prepare the pure single
phase (Losch et al., 1991; Wen, 1989, 1991). At this time
only the phases with n 5 1 (2201 phase), and n 5 2
(2212 phase) have been obtained without additional
oxides. The phase with n 5 3 (2223 phase) needs a
partial substitution of Bi/Ca by Pb to get a pure phase,
whereas the phases with n 5 4,5,7,9 have not yet been
isolated as one single phase. The intergrowth of lamella
of various BSCCO phases always has coherent (001)
interfaces but without diffusion in (001) direction,
obtaining larger critical current and giving stronger
flux line pinning when magnetic fields are applied due
to the reduction of the activation energy.
The intergrowth in the YBCO system involves the
compounds among the Y2Ba4Cu61nO141n families with
n 5 0,1,2,3,4. Both theoretical and experimental investigations showed that very fine low Tc phase particles
(like Y-124 or Y-247 phases) dispersed in the high Tc
Y-123 matrix may create good pinning sites due to the
favorable intergrowth interface of the (010) or (100)
planes with the formation of stacking defects at a
suitable size scale of 1–3 nm in thickness and 10–100
nm in width in most cases.
Besides, the high Tc copper oxides, alkali-metaldoped fullerides(MxC60) are another kind of interesting
superconductor. In the case of RbxC60 (x 5 3,4,6), we
have found that only Rb3C60 is a superconductor with
rubidium ions in both octahedral and tetrahedral interstices of FCC packing C60s. Consequently, the intergrowth between superconducting phase Rb3C60 and
non-superconducting phases RbxC60(x 5 4,6) occurs,
where lattices of RbxC60 (x 5 4,6) appear in some
domains surrounded by Rb3C60, but there are no boundaries evident between these phases because of the very
close d values of some lattice planes in these phases
(Wen et al., 1993).
In the most recently discovered ceramic superconductor HgBa2Ca2Cu3O81y, there are several members of
the homologous series of HgBa2Can21O2n121y, with n 5
0,1,2,3, . . ., where two types of intergrowth structures
between Hg-1212 and Hg-1223 as well as Hg-1223 and
Hg-1234 have been observed. The congruence relationships for these two types of intergrowth structure are
50016Hg-1212 2 50016Hg-1223 ,
,110.Hg-1212 2 ,110.Hg-1223 ,
50016Hg-1223 2 50016Hg-1234 ,
,110.Hg-1223 2 ,110.Hg-1234 .
Due to the mismatch of the atomic arrangement, a
strain field region is formed at the conjunction of the
intergrowth structure. An appropriate combination of
intergrowth structures may lead to the formation of
double immobile dislocations. From the configuration of
the atomic arrangement on the crystal plane of (110),
the mismatch region is stretched along the direction of
the nearest Ba-Ba bonding between the inter-crystal
unit cells since this is the greatest deviation in the
atomic radii (Lam et al., 1995).
As for the new large family copper oxycarbonates,
where most of them are high Tc superconductors
(Raveau et al., 1993), a great number of compounds
exhibit structures composed of single octahedral perovskite layers. The single octahedral perovskite layers are
interconnected either through layers of carbonate groups
or through mixed layers; these mixed layers contain
both carbonate groups and [AO]` ribbons with the
rock-salt structure (A 5 Tl, Bi, Hg, Pb). Most of these
compounds are closely related to the Sr2CuO2(CO3)(S2CC) structure and to the (Tl0.5Pb0.5Sr2CuO5) (1201)
structure of the thallium and bismuth cuprates
(Armstrong and Edwards, 1992; Parkin et al., 1988).
Especially, the critical temperature of these materials
is higher than that of the parent structures, S2CC and
1201 phases. Otherwise, the flexibility of their structure can accommodate species with different dimension
and coordinations such as carbonate groups and thallium or mercury cations between the octahedral copper
layers, so as to play a role in adjusting the superconductivity. There are three recently discovered new
superconductors: a 76K superconductor Hg0.6V0.4Sr4Cu2 (CO3 )O61z that exhibits a normal intergrowth
‘‘S2CC-1201’’ structure, and a 77K superconductor
Tl2/3Cr1/3Sr4Cu2(CO3)O7 and a 50K superconductor
Tl0.8V0.2Sr4Cu2(CO3)O7; the latter two exhibit a (110)collapsed ‘‘S2 CC-1201’’ structure (Maignan et al., 1995).
The high Tc superconductor is a large family and has
very complex microstructure characteristics; besides
the microstructural defects mentioned above, there are
many other kinds of imperfect structures that may
influence superconductivity. The intergrowth and superstructure incommensurate modulation are very important features in the BSCCO system (Horiuchi et al.,
1988; Matsui et al., 1988) where the extra oxygen
interstitial leads to a hole-doping, atomic displacement
and structural distortion. The structural transformation between the superconductor (SC) and the antiferromagnetic semiconductor (AF) in the BSCYCO (Bi2
Sr2Ca12xYxCu2O 81y) system is accompanied by a sharp
drop in Tc around x 5 0.4, where a structural transformation from a tetragonal to an orthorhombic phase is
produced due to a certain amount of substitution of
Ca21 for Y31 (Wen et al., 1993).
Human enamel is a very important bioceramic and
much attention has been paid to research for dental
decay repair (Featherstone et al., 1987; Takuma, 1980;
Wen et al., 1985).
The main inorganic constituent both in human bone
and teeth is hydroxyapatite, Ca10(PO4)6(OH)2, which is
in the form of an ordered prism with a hexagonal
structure oriented perpendicular to the surface in
enamel. These ordered prisms are the hardest tissue
found in the human body. However, these prisms are
easily attacked by the organic acids which are byproducts of oral bacterial metabolism.
We have shown that the dental decay process appears
to start at some defects. These defects are mostly point
defect aggregates and dislocations, stacking faults, and
grain boundaries located both inside crystals and at
grain boundaries. We also have shown that the dental
decay process, in fact, is a process of demineralization
and has three steps according to high-resolution electron microscopy observations (Wen, 1989). In order to
control the demineralization, remineralization as a
reversible process has been investigated and may be
considered to be promising for clinical application.
The first step of mineralization starts at the area less
than a nanometer inside crystals or sometimes at the
grain boundary with point defect aggregates and dislocations involved in most cases. At this stage of mineralization, only a small amount of calcium vacancies are
induced in the lattice and the whole structural frame
seems to be perfect. As for the second stage, the
demineralization area becomes large by several nanometers with a serious lattice distortion in the structure.
The chemical composition in these areas changes to less
calcium than normal. The third stage of mineralization
makes the area as large as more than a dozen nanometers with whole hydroxyapatite collapsed structurally
and both calcium and phosphorus being lost chemically.
Fig. 10. HREM of a grain boundary in synthetic chlorapatite,
showing the grains (1,2, and 3) and their very ‘‘clean’’ boundaries
(indicated by the arrow).
The research of interface in bioceramics is not only
related to the dental decay process as mentioned above
but is also related to a better understanding of the
complex mechanisms that determine and control the
process of biomineralization in the human body and aid
in the design of new materials for use in the biomedical,
microelectronic, and microengineering fields.
We started this interface research from the comparison of human enamel with synthetic bioceramics such
as synthetic hydroxyapatite and synthetic chlorapatite
(Krajewski et al., 1992). We characterized synthetic
hydroxyapatite and chlorapatite structurally by X-ray
diffraction and electron microscopy, and found that the
microstructure features between hydroxyapatite and
chlorapatite are almost the same: the grains are quite
large with a diameter of 2–3 µm on average depending
on temperature and the period of time during the
process, and the grain boundaries are very ‘‘clean’’ as
seen in Figure 10.
It is quite reasonable to observe very ‘‘clean’’ grain
boundaries in synthetic bioceramics due to the very
pure raw materials used and the rather high temperature and long period of time applied in the processing.
Further observations show that the grain boundaries,
some with small angles and other with large angles, are
distributed randomly. This is a typical microstructure
feature of synthetic ceramics.
However, in human enamel the grains are extremely
small and on a nanometer scale with a preferable
orientation along the [001] zone axis as seen in Figure
11. The interface or grain boundary phase in human
enamel has very complicated structures in comparison
with synthetic hydroxyapatite and chlorapatite. The
interface has three kinds of structural features.
The first kind is most simply like the case of synthetic
hydroxyapatite as shown at the grain boundary between grain A and B (or grain A and E) in Figure 11.
Here the grain boundaries are ‘‘clean’’ and coherent
with certain angles. They are: S31(18°), S14(22°),
S13(28°), S26(32°), S7(36°), S19(46°), S37(50°), and
Fig. 11. HREM (dark field) of human enamel imaged by an electron
beam being perpendicular to a tooth surface, showing its microstructure feature with preferable [001] direction, very small grain size, and
a complicated interface including the amorphous phase (arrows).
copy observations indicate that the area with the size of
20–30 nm is almost amorphous but some very tiny
crystallites are dispersed there. The enamel contains
1% protein (Takuma et al., 1975), which may distribute
this kind of pocket since these pockets along the [001]
direction are usually as long as 500 nm although they
have the width of only 20–30 nm according to our
From the above comparison between human enamel
and synthetic hydroxyapatite, the true interface is the
third kind of grain boundaries in human enamel and
can be considered as a true interface by which living
organisms (protein) convert the ions of Ca21 and (PO4)32
into hydroxyapatite prism.
In summary, there are three kinds of grain boundaries in human enemal. The microstructures of the first
and second kind of boundaries have been characterized
by the present study and it was found that they have a
very similar microstructural feature to synthetic nanostructured ceramic materials. However, the third kind
of grain boundaries is greatly different from the former
two due to a living organism, protein, involved. The
third kind of boundaries is even more complicated
because of a mass exchange between inorganic ions and
protein moleculars through these boundaries. The research for this kind of boundaries has been made in
progress and submitted to J. Appl. Phys.
Fig. 12. HREM of human enamel showing the grain boundaries
between the grains A, B, and C.
The second kind of grain boundaries has a more
complicated structure as shown at the grain boundary
between grain C and D. The most featured grain
boundaries of this kind are shown at the area between
grain B and C (indicated by arrowhead) in Figure 12.
The second kind of grain boundaries phase has a
similar structure with inside grains but with serious
lattice distortion judged by a Moiré pattern. The angles
of grain boundaries can be other values rather than
certain values as in the first kind of grain boundary
mentioned above. This may be due to 3% of other
constituents (such as F, Cl, Mg, Zn, Sr, Na, Pb, Al, Ba,
and carbonate) existing in grain boundaries, and substitution produced that causes point defects and lattice
The third kind of grain boundaries is at the pocket of
grains (arrowhead in Fig. 12) where 5 grains (A,B,C,D,
and E) are involved. High-resolution electron micros-
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