Synthesis of fully dense nanostabilized undoped tetragonal zirconia

Article abstract of DOI:10.1111/j.1551-2916.2010.03695.x

Zirconia-based materials, with very low or null doping contents (YO _1.5 between 0% and 4% mol), were prepared in the form of bulk nanostructured materials through a two-step process. Nanometric powders of size-stabilized tetragonal zirconia were synthesized by the Pechini method and subsequently densified using the high-pressure field-assisted rapid sintering method at temperatures around 850°-900°C for 5 min under a pressure of ～800 MPa. By this method, theoretical densities of ～98% were achieved while retaining the grain size in the final product between 40 and 50 nm. The transition to the thermodynamically stable monoclinic phase was consequently prevented, allowing for the first time the synthesis of size-stabilized zirconia in a high-density bulk form.

Full text of DOI:10.1111/j.1551-2916.2010.03695.x

J. Am. Ceram. Soc., 93 [7] 2092–2097 (2010)
DOI: 10.1111/j.1551-2916.2010.03695.x
r 2010 The American Ceramic Society
ournal
J
Synthesis of Fully Dense Nanostabilized Undoped Tetragonal Zirconia
Filippo Maglia,w,z Monica Dapiaggi,y Ilenia Tredici,z Beatrice Maroni,z and Umberto Anselmi-Tamburiniz
zDepartment of Physical Chemistry, University of Pavia, 27100, Pavia, Italy
yDepartment of Earth Science, University of Milano, 20133, Milano, Italy
Zirconia-based materials, with very low or null doping contents
(YO1.5 between 0% and 4% mol), were prepared in the form
of bulk nanostructured materials through a two-step process.
Nanometric powders of size-stabilized tetragonal zirconia were
synthesized by the Pechini method and subsequently densified
using the high-pressure field-assisted rapid sintering method at
temperatures around 8501–9001C for 5 min under a pressure of
B800 MPa. By this method, theoretical densities of B98%
were achieved while retaining the grain size in the final product
between 40 and 50 nm. The transition to the thermodynamically
stable monoclinic phase was consequently prevented, allowing
for the first time the synthesis of size-stabilized zirconia in a
high-density bulk form.
The size-induced stabilization has been exclusively explored on
nanopowders or in thin films.22 No attempt to retain such char-
acteristics in high-density bulk materials has ever been reported.
The possibility to obtain bulk tetragonal, or cubic, zirconia with
no dopant content would open the possibility to investigate the
functional and structural properties of zirconia in conditions
never explored before. The difficulty encountered in the prepa-
ration of bulk nanocrystalline materials, especially when a very
low level of porosity is required, so far prevented the possibility to
obtain dense zirconia samples with a grain size below the critical
value for size-stabilization of the tetragonal phase.
It has recently been shown that the high-pressure spark
plasma sintering (HP-SPS) allowing rapid sintering cycles at
high pressures (up to 1 GPa), can be used to prepare nearly fully
dense nanocrystalline materials with minimal grain growth.23–25
Through this approach, the feasibility of consolidation of yttria
fully stabilized zirconia (8% YSZ) and ceria samples to relative
densities exceeding 98% and with a grain size below 20 nm has
been demonstrated.
I. Introduction
IRCONIA at ambient pressure exists in three different poly-
Z
morphs: monoclinic (room temperature—11751C), tetrago-
nal (11751–23701C), and cubic (23701–26801C). Tetragonal and
cubic polymorphs are those relevant for technological applica-
tions, based on their oxygen conductivity as in solid oxide fuel
cells, sensors, and in catalysis. The tetragonal and cubic phases
are traditionally stabilized at room temperature by doping with
lower valence cations (e.g., Ca21, Mg21, Y31).1
The aim of this work was the application of a technique sim-
ilar to the HP-SPS, known as high-pressure field-assisted rapid
sintering method (HP-FARS) to the densification of zirconia
nanopowders with null or low dopant levels with the intent of
obtaining bulk nanostabilized tetragonal ZrO2 samples at do-
pant levels below the stabilization limit for the tetragonal phase
at room temperature.
Recently, the stabilization of metastable phases induced by
the nanostructure has attracted considerable attention, not only
in zirconia, but in several other oxides, including Al2O3,2,3
TiO2,4 and perovskites.5–7 This effect is generally observed
when the dimension of the crystalline domains is below some
critical value. Since its discovery by Garvie,8,9 a large number of
studies have been devoted to the nanoinduced stabilization of
zirconia, especially in the form of nanopowder.10–18 Despite this,
the mechanism responsible for the nanoinduced stabilization has
not been completely clarified. While the differences in surface
energy between the polymorphs is by far the most popular ex-
planation, the influence of anionic impurities, lattice strain,
structural similarities between the precursor materials and te-
tragonal zirconia, influence of lattice defects, and/or water vapor
have also been proposed.10 On the other hand, it is generally
agreed that nanocrystalline tetragonal zirconia is not just kinet-
ically metastable, but can be truly thermodynamically more
stable than the monoclinic form, as long as coarsening is
precluded.10,13
II. Experimental Procedure
ZrO2 nanopowders with a dopant (YO1.5) content ranging be-
tween 0 and 4 at.% Y (i.e., from pure ZrO2 to Zr0.96Y0.04O2ꢀx/2
)
were obtained using a modified Pechini method.26 The stability
range of the pure monoclinic phase at room temperature is cov-
ered by the composition interval considered. Zirconyl (IV)
nitrate and citric acid were taken in a molar ratio of 1:1. To a
proper amount of a commercial aqueous solution (35% weight
in nitric acid) of Zirconyl (IV) nitrate, a saturated solution of
citric acid was added (typically, to 30 mL of Zirconyl solution,
12 mL of saturated citric acid solution was added). The solution
was then stirred on a hot plate at 801C, to obtain a solution of
increasing viscosity. After a time period that depends on the
amount of solution used (3 h for the volumes reported above),
the viscous mass turned into a colorless transparent glass. Fur-
ther heating (typically 5 h for the reported amounts) provided a
white solid, which was subsequently ground in a mortar and
calcined in a furnace at 6251C for 2.5 h; the calcination temper-
ature was selected as the lowest that leads to a complete removal
of the carbonaceous residuals. The white powders were succes-
sively densified using a home-made HP-FARS apparatus. This
apparatus presents the same general design of an HP-SPS, but
uses low-voltage AC high-intensity current instead of the pulsed
DC current generally used in the SPS. A similar apparatus has
been used previously by Gauthier et al.27 In each experiment,
0.15 g of powders were loaded in a double-stage die. The low-
pressure section of the die was made of high-density graphite,
while the high-pressure section was made of silicon carbide and
As summarized by Li et al.,12 several different values of crit-
ical size can be found in the literature, ranging between 3 and 30
nm, depending on the synthetic route and the size determination
technique. Some authors also reported the possibility to stabilize
the cubic polymorph at even smaller sizes than those required
for the tetragonal phase,19–21 although some controversy exists
on this point.
J. Hellmann—contributing editor
Manuscript No. 26807. Received September 18, 2009; approved January 28, 2010.
wAuthor to whom correspondence should be addressed. e-mail: filippo.maglia@unipv.it
2092

July 2010
Dense Nanostabilized Undoped Tetragonal Zirconia
2093
tungsten carbide.23 The die was loaded in the HP-FARS appa-
ratus, which was evacuated to a pressure of 10 Pa. A moderate
initial uniaxial pressure (150 MPa) was applied. The tempera-
ture was then increased with a heating rate of 2001C/min. Once
the sample attained the designated temperature, the pressure
was rapidly increased to the final value. The sample (5 mm in
diameter, B1 mm thick) was held under these conditions for
5 min and then the pressure was quickly released and the power
was turned off. Temperatures were measured using a shielded
K-type thermocouple inserted in the lateral wall of the die.
Microstructural characterization of the samples was made
on uncoated fracture surfaces, using a high-resolution SEM
(HRSEM, Philips XL30s, Eindhoven, the Netherlands). The
average grain size was determined measuring at least 100 grains
on each HRSEM image using the software AnalySIS (Soft
Imaging System Corp., Lakewood, CO). The density of the
samples was measured using the Archimedes method and from
geometric and gravimetric measurements.
The X-ray powder diffraction patterns were collected at the
beamline ID31 of the synchrotron source European Synchrot-
ron Radiation Facility (ESRF) in Grenoble, with a wavelength
Fig. 1. XRD patterns of the Zr1ꢀxYxO2ꢀx/2 (0oxo0.04) powder sam-
ples after calcination at 6251C.
˚
of 0.399946 (3) A, calibrated against NIST 640c standard silicon
at room temperature. The beam was monochromated by a cryo-
genically cooled Si 111 double-crystal monochromator and a
bank of nine detectors was vertically scanned to measure the
diffracted intensity as a function of 2y. Each detector was pre-
ceded by a Si 111 analyzer crystal. The powders were contained
in a boron glass capillary (0.6 mm diameter), while for densified
samples, a fragment of the pellet was placed in a larger capillary
(1.5 mm diameter) and fixed with wax to avoid moving the
sample during spinning. The capillaries were spun for a better
counting statistics.
The cell refinements and the quantitative analyses were
performed using the Rietveld method, using the software
GSAS1EXPGUI.28,29 The crystallite size and the root mean
square (RMS) microstrains were evaluated using the software
MAUD30: the standard used for the evaluation of the instrumen-
tal broadening was LaB6 (NBS 660). The model used was the
isotropic one, in which the microstructural evaluation was made
over the whole diffraction pattern, assuming that the diffracting
domains were isotropic in shape. The crystallite size and the RMS
microstrain were evaluated for both the zirconia polymorphs
when they were present in adequate quantities. When the mono-
clinic phase was o10% wt., the whole peak broadening was as-
sumed to be originating from small crystallite size only.
corresponding structural and microstructural characteristics,
obtained by the refinement of the XRD data, are reported in
Table I. The diffraction peaks are significantly broadened, indi-
cating the formation of zirconia nanocrystals. The calculated
grain size of the tetragonal phase is around 15 nm with varia-
tions among the different compositions within the uncertainty
typical of the synthetic method (Table I). The grain size is
mainly controlled by the calcination conditions and smaller
grains could be obtained by lowering the calcinations tempera-
ture. Although full crystallization of the zirconia powder is
achieved at temperatures close to 4001C, the elimination of the
carbonaceous residuals requires temperature above 6001C. The
thorough removal of all these impurities is essential in order to
avoid their incorporation in the sintered samples.
All the X-ray reflections of Fig. 1 can be indexed by the
tetragonal and monoclinic polymorphs. In agreement with the
literature,8–18 the high-temperature tetragonal phase is the major
component of the product thanks to the stabilizing effect of the
nanometric particle size. As expected, there is a synergic stabi-
lizing effect played by both the reduced crystallite dimension
and the doping amount; consequently, the fraction of the mono-
clinic phase decreases as the amount of the stabilizing Y31 is
increased. The amount of monoclinic phase, calculated by
Rietveld refinement of the XRD data, decreases from 17.1% wt.
in the pure ZrO2 sample to below the detection limit in the
Zr0.96Y0.04O2ꢀx/2 one (Table I). The concomitant presence of the
monoclinic and tetragonal phases in pure (or low dopant con-
tent) zirconia was reported by most previous works and their
III. Results
XRD patterns of the Zr1ꢀxYxO2ꢀx/2 (0oxo0.04) powder
samples after calcination at 6251C are shown in Fig 1; the
Table I. Structural and Microstructural Characteristics of the Starting Powder
Tetragonal
Y at.%
3 w
˚
V (A )
RMS strain (ꢁ10ꢀ3
)
˚
a (A)
˚
c (A)
˚
c/a
oD4 (A)
0
0.5
2
5.0878 (1)
5.08841 (9)
5.09110 (9)
5.09989 (9)
5.1791 (3)
5.1792 (3)
5.1802 (2)
5.1672 (4)
67.03
67.05
67.13
67.20
1.0179
1.0178
1.0174
1.0132
148 (1)
151 (1)
161 (1)
125 (1)
3.21 (3)
3.31 (1)
3.42 (2)
5.31 (5)
4
Monoclinic
Y at.%
RMS strain (ꢁ10ꢀ3
)
˚
oD4 (A)
0
0.5
2
198 (4)
164 (3)
120 (4)
ND
2.31 (5)
1.78 (8)
2.3 (4)
ND
4
wPlease note that the volume was calculated using the original primitive cell used in the refinements, and not with the centered cell usually used in the literature, which is
reported for an easier comparison with the literature data (e.g., Garvie and colleagues8–18). RMS, root mean square.

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Journal of the American Ceramic Society—Maglia et al.
Vol. 93, No. 7
Table II. Sintering Conditions and Product Density
richer in yttrium in order to achieve a theoretical density similar
to the one obtained for lower contents (B95%–96%). As a
consequence, a small but consistent increase in the grain size was
observed as the doping content is increased (Table III). The
temperature range is anyway narrow as it varies between 8501C
for pure ZrO2 and 9001C for 4 at.% Y. It must be stressed that
the conditions reported in Table II are quite strict and the qual-
ity of the samples decreases dramatically either by reducing the
pressure and/or the temperature (insufficient density) or increas-
ing them (grain size exponentially increases with temperature).
The temperatures used in this study are considerably lower than
those required for the sintering of zirconia by traditional meth-
ods.1 At such low temperatures and short sintering times, the
sintering mechanism is most probably based on mechanisms
different from the thermal-activated diffusion of the atomic spe-
cies. The only fast-diffusing species in zirconia at low tempera-
tures is oxygen, but sintering and densification of these materials
require the diffusion of cations, which at these temperatures is
very limited.32 Furthermore, the rapid application of pressure
produces very high strain rates. Under these conditions, plastic
(or superplastic) deformation may play a very important role in
supporting a densification mechanism based on grain-boundary
sliding and grain rotation.
Y at.%
Heating rate (1C/min)
Temperature (1C)
Pressure (MPa)
0
0.5
2
200
200
200
200
850
875
875
900
790
800
825
825
4
relative amount appears to depend not only on the particle size
but also on the synthetic route used.8–18 Zhang et al.18 observed
the coexistence of both monoclinic and tetragonal polymorphs
when the average particle size was between 14 and 31 nm, while
only the tetragonal phase was observed below 14 nm, and only
the monoclinic one above 31 nm. The coexistence of the two
polymorphs in the intermediate grain size range is suggested to
result from the size distribution; because there is a certain grain
size distribution in one batch of samples, this implies that the
grain sizes of the monoclinic phase are larger than those of the
tetragonal one. Our results are not in agreement with this con-
clusion because, with the exception of the dopant-free sample
that shows a slightly higher value (19.8 nm) of the grain size of
monoclinic phase, we did not detect a significant difference be-
tween the crystallite size of two polymorphs (Table I).
Figure 2 shows the fracture surface of a sample with 2 at.% Y
sintered under the experimental conditions reported in Table II.
The fracture surface allows to appreciate the small grain size, the
good grain size homogeneity, and the absence of appreciable
residual porosity. Similar to what was observed on other mate-
rials such as fully stabilized YSZ,24,31 pure23 and Sm-doped
ceria,25 LSGM33 etc., the HP-FARS technique allowed the pro-
duction of high-density materials with a limited (typically two to
three times the initial value) grain growth even starting, as in this
case, from the badly agglomerated powder, which generally
produces sintered samples characterized by large mesoporosity.
Figure 3 shows the XRD patterns collected on the densified
samples. The most important result is indeed represented by the
phase composition, which is summarized in Fig. 4 (upper part)
as a function of the Yttrium content for both the starting pow-
der and the sintered samples. Although the amount of the
monoclinic phase is higher in the densified samples, due to
Ostwald ripening during the sintering process, the tetragonal
phase remains in all cases the major component; even in the pure
bulk ZrO2 sample, the amount of the monoclinic phase is only
around 25% in weight in the sintered sample. Figure 4 (lower
part) shows the variation of the grain size and stress content
determined by the MAUD program for both powder and dense
samples as a function of Y31 content. It should be observed that
the crystallite size of Fig. 4 is in good agreement with the particle
The calcined precursor nanopowders were densified by the
HP-FARS method. As already mentioned, the main aim was to
produce a high-density material with minimal grain growth
to prevent, as much as possible, the tetragonal to monoclinic
transition. Different sintering conditions were explored for this
purpose, with particular focus on the maximum sintering tem-
perature and uniaxial applied pressure. Table II summarizes the
experimental parameters that produced the best results in terms
of density and crystallite size. A very short hold time (5 min) was
used in all experiments as this parameter was shown to have a
limited influence on sample density, because the major portion
of the overall densification occurs before reaching the final tem-
perature.31 In a study on the densification via SPS of a com-
mercial nanometric powder of fully stabilized zirconia (Tosoh,
16 at.% Y), we showed that the temperature required to achieve
the 95% of the theoretical density decreases linearly with the
logarithm of the applied pressure.31 As a consequence, it was
possible to obtain a more than exponential decrease in the final
grain size by increasing the applied pressure. In agreement with
the above, in the present study also, the minimum grain growth
and highest density were obtained by using the maximum pres-
sure allowed by the instrument, which is around 800 MPa.
The densification under the maximum applied pressure re-
quired slightly higher sintering temperatures for the composition
Table III. Structural and Microstructural Characteristics of the Sintered Samples
Tetragonal
Y at.%
3 w
˚
V (A )
RMS strain (ꢁ10ꢀ3
)
˚
a (A)
˚
c (A)
˚
c/a
oD4 (A)
0
0.5
2
5.08381 (8)
5.08449 (7)
5.08738 (7)
5.09336 (9)
5.1869 (2)
5.1860 (1)
5.1836 (2)
5.1769 (4)
67.02
67.04
67.08
67.15
1.0203
1.0199
1.0189
1.0164
412 (3)
452 (5)
473 (6)
472 (9)
1.56 (3)
1.39 (2)
1.31 (3)
1.35 (6)
4
Monoclinic
Y at.%
RMS strain (ꢁ10ꢀ3
)
˚
oD4 (A)
0
0.5
2
491 (12)
643 (23)
805 (13)
ND
3.3 (2)
4.3 (2)
2.9 (4)
ND
4
wPlease note that the volume was calculated using the original primitive cell used in the refinements, and not with the centered cell usually used in the literature, which is
reported for an easier comparison with the literature data (e.g., Garvie and colleagues8–18). RMS, root mean square.

July 2010
Dense Nanostabilized Undoped Tetragonal Zirconia
2095
Fig. 2. Fracture surface of a sample with 2 at.% Y sintered under the
experimental conditions of Table II.
size observed for the HRSEM images of the fracture surfaces
(Fig. 2).
Figure 5 shows the comparison of the weight percentage of
the monoclinic phase as a function of the crystallite size of the
tetragonal phase for both the powders and the densified sam-
ples. The amount of zirconia that is transformed into the mono-
clinic polymorph during sintering is much larger as the dopant is
reduced. As expected, there is a combined effect of the grain size
reduction and the amount of doping in the stabilization of the
tetragonal phase. It is reasonable to suggest, in agreement with
the data of Fig. 4, that the effect of introducing Y31 in the
nanostructured tetragonal zirconia shifts the critical value for
the transition to the monoclinic phase towards larger grain sizes.
Under the sintering condition adopted in this study, the grain
size increases from around 15 nm in the starting powder to 43–
50 nm in the densified samples. It could appear surprising that
such a low monoclinic content is present in the sintered samples,
especially for the lowest doping content. When compared with
the available literature data, concerning pure zirconia in the
powder form, the grain size exhibited by our sintered samples
would, in fact, appear hardly compatible with the presence of
large amounts of size-stabilized tetragonal phase. This point has
been recently discussed by Li et al.12 In their theoretical work,
the authors suggested the possibility of a larger critical grain size
(37 nm) in nanostructured bulk materials than in powders
(10 nm) at room temperature. This result is justified by a con-
siderably larger energy cost in the transition from the tetragonal
to monoclinic phase when the free surface contribution
(Dg5 0.18 J/m) is replaced by the grain-boundary contribution
Fig. 4. Upper part: variation of the amount of tetragonal phase as a
function of the Y content (black symbols: powder; white symbols: sin-
tered samples). Lower part: variation of the amount of crystallite size of
the tetragonal phase as a function of the Y content (black symbols:
powder; white symbols: sintered samples).
(Ds5 1.24 J/m). Further increase in the critical grain size, pos-
sibly above 40 nm, can occur, according to the same authors,
thanks to other factors such as the existence of internal strain
and in particular the presence of a surrounding rigid matrix (as
in our bulk samples) that prevents the volume expansion ac-
companying the tetragonal to monoclinic transition.
In our densified samples, the stabilization of the tetragonal
phase could be also due to some structural alteration introduced
by the fast sintering cycle that may create a quenched-in meta-
stable state inside the material. The structural parameters of the
tetragonal phase in the powder and in the sintered material were,
therefore, analyzed and compared. Figure 6 shows the varia-
tions of cell volume (upper part) and of the individual cell
parameters (lower part) with YO1.5 content. As expected, due to
the larger ionic size of the Y31 ion, cell volume linearly increases
with an increasing Y31 content even though an opposite trend
can be observed for the individual cell parameters (lower part of
Fig. 6). The c/a ratio (Fig. 7) is generally used as the parameter
accounting for the ‘‘tetragonality’’ degree of the structure, in-
tended as a deviation from the ideal cubic one (c/a 5 1). As ex-
pected, the tetragonality decreases as the amount of yttrium is
Fig. 3. XRD patterns of the densified samples with a nominal compo-
sition Zr1ꢀxYxO2ꢀx/2 (0oxo0.04).

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Journal of the American Ceramic Society—Maglia et al.
Vol. 93, No. 7
Fig. 5. Variation of the amount of monoclinic phase as a function of
the crystallite size of the tetragonal phase for the 2 at.% Y composition
(circles), 0.5 at.% Y (squares), and pure zirconia (triangles). Black sym-
bols, powder; open symbols, sintered samples.
increased as, according to the phase diagram, the structure
should turn to the cubic polymorph at Y at.%5 16.
A marked difference could, on the other hand, be found in the
structural parameters of the powder and sintered samples of
equal composition. The sintered samples always show a smaller
/aS, a larger /cS, and a larger c/a ratio for all the explored
compositions. This variation of the unit cell parameters, that
ultimately leads to a smaller cell volume in the sintered samples,
can in principle be explained on the basis of the following hy-
potheses: (1) a permanent deformation of the unit cell due to the
stress applied during the densification, (2) the loss of larger Y31
ions from the tetragonal structure, or (3) the creation of excess
oxygen vacancies. The first hypothesis does not appear realistic,
because the applied pressure is not sufficient to introduce such
permanent modifications in the crystal structure. The absence of
any trace of Y2O3 appears to exclude the second hypothesis also.
The third hypothesis, i.e. the formation of non equilibrium ox-
ygen vacancy distribution, appears on the other hand more re-
alistic, due to the relatively fast heating/cooling cycle used and
the reducing environment (due to the graphite dies) typical of
the FARS technique. Even though the white color displayed by
the sintered samples appears to exclude the presence of a large
oxygen substoichiometry in these samples, a small variation of
their number and distribution cannot be a priori excluded.
Another obvious alteration introduced by sintering is an in-
crease in the grain size that on average is tripled in the densified
samples. The cell parameter could therefore vary as an effect of
the change in grain size. A comparison with the literature data
concerning this point was possible only for the pure zirconia
(Fig. 8). The variation of the lattice parameters and of the c/a
ratio observed in our samples is in good agreement with the
trend observed by Lamas et al.34 who evaluated these parame-
ters in powders of particle size between 5 and 11 nm prepared by
different synthetic routes. The strong decrease in the c/a ratio is
taken by the authors as a possible indication of a transition
toward a cubic structure (c/a 5 1) at even smaller grain sizes,
while the increase in the c/a parameters as the grain size in-
creases is expected to lead to the formation of the monoclinic
polymorph. Even though a similar trend was observed by other
authors also,11 a clear explanation on this point is not at hand.
The data shown in Fig. 8 indicate anyway as the sintered
samples prepared by HP-FARS well fit within the trend exhib-
ited by zirconia nanopowder and no alteration due to the high
Fig. 6. Upper part: Variation of the cell volume of the tetragonal phase
as a function of the Y content for the starting powder and the sintered
samples. Lower part: Variation of the unit cell parameters of the tetrag-
onal phase as a function of the Y content for the starting powder and the
sintered samples.
Fig. 7. Variation of the c/a ratio of tetragonal phase as a function
of the Y content. Black symbols, powders; open symbols, densified
samples.

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Dense Nanostabilized Undoped Tetragonal Zirconia
2097
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Fig. 8. Upper part: variation of the /aS (open symbols) and /cS
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c/a ratio for the tetragonal phase as a function of grain size; circles, this
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In the present work, we have carried out the synthesis of pure
and low doping (o4 at.% YO1.5) zirconia in the form of bulk
nanostructured samples. The starting powder, prepared by the
Pechini method, was densified using the High-Pressure Field-
Activated Rapid Sintering method. High relative density
(498%) samples could be obtained at temperatures as low as
8501–9001C and for sintering times of only 5 min. As a conse-
quence of the mild sintering conditions, a limited grain growth
was observed; the sintered samples showed grain sizes around
45–50 nm for all the compositions investigated.
Thanks to the reduced dimension of the grain size, it was
possible, for the first time, to retain the nanostabilization of the
tetragonal phase even in bulk high-density samples. Only a very
limited, or null, transition toward the thermodynamically stable
monoclinic phase was detected during the sintering procedure.
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Acknowledgments
This work has been supported by Fondazione Cariplo.
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