Available online at www.sciencedirect.com
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Journal of the European Ceramic Society 34 (2014) 2453–2463
3Y-TZP ceramics with improved hydrothermal degradation resistance
and fracture toughness
F. Zhang a , K. Vanmeensel a,∗ , M. Inokoshi b , M. Batuk c , J. Hadermann c ,
B. Van Meerbeek b , I. Naert b , J. Vleugels a
b
a Department of Metallurgy and Materials Engineering, KU Leuven, Kasteelpark Arenberg 44, B-3001 Leuven, Belgium
KU Leuven BIOMAT, Department of Oral Health Sciences, KU Leuven & Dentistry, University Hospitals Leuven, Kapucijnenvoer 7 blok a, B-3000 Leuven,
Belgium
c Electron Microscopy for Materials Research (EMAT), University of Antwerp, Groenenborgerlaan 171, B-2020 Antwerp, Belgium
Received 4 July 2013; received in revised form 23 January 2014; accepted 9 February 2014
Available online 6 March 2014
Abstract
Different factors such as the way of incorporating the Y2 O3 stabilizer, alumina addition and sintering temperature were assessed with the goal to
improve the low temperature degradation (LTD) resistance of 3Y-TZP without compromising on the mechanical properties. The degradation of
hydrothermally treated specimens was studied by X-ray diffraction, micro-Raman spectroscopy and scanning electron microscopy.
Decreasing the sintering temperature decreased the LTD susceptibility of 3Y-TZPs but did not allow to obtain a LTD resistant 3Y-TZP with
optimized mechanical properties. Alumina addition along with the use of Y2 O3 stabilizer coated starting powder allowed to combine both an
excellent toughness and LTD resistance, as compared to alumina-free and stabilizer co-precipitated powder based equivalents. Transmission
electron microscopy revealed that the improved LTD resistance could be attributed to the segregation of Al3+ at the grain boundary and the
heterogeneously distributed Y3+ stabilizer.
© 2014 Elsevier Ltd. All rights reserved.
Keywords: Coated 3Y-TZP; Co-precipitated 3Y-TZP; Degradation; Alumina
1. Introduction
Yttria stabilized tetragonal zirconia polycrystals (Y-TZP) are
increasingly used as dental restorative material in recent years
because they possess excellent mechanical properties, such as
superior strength and fracture toughness while they also exhibit
biocompatibility and aesthetic potential.1–3 The high strength
and toughness of zirconia ceramics are due to stress-induced
phase transformation toughening. Around a propagating crack,
the metastable tetragonal grains can transform to the monoclinic structure, t–m transformation, and the associated volume
expansion induces compressive stresses and eventually reduces
or inhibits further crack propagation.2,4 However, the tetragonal grains can also spontaneously transform to monoclinic in
∗
Corresponding author. Tel.: +32 16321192.
E-mail address: [email protected] (K. Vanmeensel).
http://dx.doi.org/10.1016/j.jeurceramsoc.2014.02.026
0955-2219/© 2014 Elsevier Ltd. All rights reserved.
a humid environment in the 20–300 ◦ C range,5,6 like in the
oral cavity, a phenomenon being referred to as low temperature
degradation (LTD).5 The mechanism of water incorporating into
zirconia to induce t–m transformation is not yet fully understood,
but it is well established that the degradation starts from the surface and proceeds inwards. Surface uplift7,8 and micro-cracks9
are subsequently induced which can result in macrocracks9 and
surface roughening.10 Consequently, the mechanical10–13 and
even the aesthetic properties14 of Y-TZPs are affected.
The LTD susceptibility of Y-TZP is affected by different
parameters. The degradation can be retarded by increasing the
Y2 O3 content.5 For dental applications, 3 mol% yttria stabilized
tetragonal zirconia polycrystals (3Y-TZP) are commonly used.15
The sintering condition is a predominant factor influencing both
the mechanical properties and LTD behaviour of 3Y-TZPs.15–19
However, the mechanism was not well studied and different manufacturers produce dental 3Y-TZP materials at different sintering
conditions.3
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F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
It was reported that high purity ZrO2 starting powder was
needed to inhibit the transformation20 but it was also claimed
that the beneficial effect of additives like TiO2 , Fe2 O3 and Al2 O3
depends on their amount.20 A small amount of homogeneously
distributed alumina (<0.5 wt.%) can decrease the LTD susceptibility of 3Y-TZP.21 Commercially available dental ZrO2 is
therefore commonly doped with 0.25 wt.% alumina.
A homogeneous distribution of the Y2 O3 stabilizer is generally believed to increase the degradation resistance of TZPs22
but different opinions are recently challenging this.5 Moreover,
the mechanical properties of yttria-coated Y-TZP were reported
to be more attractive.17,23,24 Due to the inhomogeneous Y2 O3
distribution, coated Y-TZP had a higher toughness compared to
co-precipitated Y-TZP.23 Moreover, coated Y-TZP is expected
to have a higher LTD resistance.17 A simple nitrate process
has been reported for the processing of stabilizer-coated ZrO2
starting powder, resulting in TZPs with excellent toughness and
inhomogeneous stabilizer distribution.24,25 The LTD behaviour
of yttria-coated powder based TZPs has however not been studied.
Furthermore, studies on dental 3Y-TZP were focused either
only on the mechanical properties17,18,24 or only on the
LTD behaviour.6,16,19,26,27 It is however important to preserve
the excellent mechanical properties, especially the fracture
toughness, when developing hydrothermally stable Y-TZP biomaterials.
The aim of this work was to investigate the influence of
the sintering condition, alumina addition and the way of incorporating the Y2 O3 stabilizer on the mechanical properties
and LTD behaviour of 3Y-TZPs. 3 mol% yttria stabilizer was
incorporated in the starting ZrO2 powder by a nitrate coating technique as compared to commercially available 3 mol%
yttria co-precipitated ZrO2 powder. Alumina-free and 0.25 wt.%
alumina-doped yttria-coated and co-precipitated powder based
ceramics were investigated.
2. Experimental procedures
3 mol% Y2 O3 coated ZrO2 powder was prepared according
to the nitrate route.28,29 The proper amount of Y2 O3 (99.9%,
Acros, Geel, Belgium) was dissolved in warm nitric acid (65%,
Sigma–Aldrich, Bornem, Belgium) and mixed with pure monoclinic ZrO2 nanopowder (grade TZ-0, Tosoh, Japan) in ethanol
in a polyethylene container on a multidirectional mixer (Turbula, Basel, Switzerland) for 24 h. An appropriate amount of
Al(NO3 )3 ·9H2 O (Aldrich Chemical Company) was added to
the suspension to obtain 0.25 wt.% Al2 O3 doped Y2 O3 -coated
ZrO2 powder. Y-TZP milling balls (grade TZ-3Y, Tosoh, Japan)
with a diameter of 5 mm were added to the suspension to break
the ZrO2 and Al2 O3 agglomerates during mixing. After drying,
the yttrium nitrate coated powder was calcined in air for 30 min
at 800 ◦ C. A second multidirectional mixing was performed in
ethanol for 48 h to break the hard agglomerates after calcination. Ceramic discs were obtained by cold isostatic pressing
(CIP) at 300 MPa and subsequent pressureless sintering in air for
2 or 4 h at 1350–1550 ◦ C. Reference ceramics were processed
from 3 mol% Y2 O3 co-precipitated ZrO2 powder (grade TZ-3Y,
Tosoh, Japan) and the 0.25 wt.% Al2 O3 containing equivalent
(grade TZ-3Y-E, Tosoh, Japan).
The nomenclature of the experimental 3Y-coated ZrO2 ,
0.25 wt.% Al2 O3 doped 3Y-coated ZrO2 and commercial powder based 3Y-coprecipitated ZrO2 and 0.25 wt.% Al2 O3 doped
3Y-coprecipitated ZrO2 ceramics used throughout the text is
CZ-3Y, CZ-3Y-0.25Al, TZ-3Y and TZ-3Y-0.25Al, respectively.
The density of the sintered ceramics was measured in ethanol
according to the Archimedes principle, assuming a theoretical
density of 6.05 g/cm3 to calculate the relative density.30
Phase identification was done by X-ray diffraction (XRD,
3003-TT, Seifert, Ahrensburg, Germany) using Cu K␣ radiation
at 40 kV and 40 mA. Rietveld analysis of the XRD pattern was
performed with X’Pert Highscore Plus to quantitatively calculate the amount of tetragonal (P42/nmc (1 3 7) space group) and
cubic ZrO2 phase (Fm-3m (2 2 5) space group). The Y2 O3 content in the remaining tetragonal ZrO2 phase was calculated based
on the calculated a and c unit cell parameters of the tetragonal
ZrO2 phase according to the formula:31
√
1.02311 − (ct / 2at )
YO1.5 (mol%) =
(1)
0.001498
The microstructure was analyzed by scanning electron microscopy (SEM, XL-30FEG, FEI, Eindhoven, The
Netherlands). Cross-sectioned ceramics were polished and thermally etched for 25 min at 1250 ◦ C in air. SEM images were
used to measure the average grain size of the sintered ceramics
using IMAGE-PRO software according to the linear intercept
method. The reported grain sizes are the as-measured values
obtained from at least 1000 grains, without any correction.
Scanning transmission electron microscopy (STEM) images,
elemental mapping and energy-dispersive spectroscopy (EDS)
were performed with a FEI Titan 60-300 “cubed” transmission
electron microscope to examine the distribution of Y3+ , Al3+ and
Zr4+ around the grain boundaries. Electron transparent samples
were prepared by ion-milling with an Ion Slicer (EM-09100IS,
Jeol, Japan). The transmission electron microscope was operated
at 200 kV and a high magnification having a mapping resolution
below 0.19 nm. 5–7 grain boundaries in each specimen were
analyzed. The quantitative elemental mapping was acquired to
calculate the Y3+ , Al3+ and Zr4+ concentration profile using
ESPRIT 1.9 software.
The Vickers hardness was measured on a hardness tester
(Model FV-700, Future-Tech Corp., Tokyo, Japan) with a load
of 9.8 N and a dwell time of 10 s. The indentation toughness
was evaluated from the radial crack pattern accompanying the
Vickers indentations and calculated according to the Anstis
equation.32 The reported values are the mean and standard deviation of 10 indentations. An E-modulus value of 210 GPa was
used to calculate the fracture toughness of all ceramics.
For low temperature degradation (LTD) testing, accelerated
hydrothermal degradation testing was used and mirror polished
specimens were autoclaved in steam at 134 ◦ C and 0.2 MPa up to
40 h. At predefined times, the test was interrupted to measure the
surface monoclinic ZrO2 phase content by means of XRD. XRD
patterns were recorded on both polished flat surfaces of each
specimen in the 27◦ ≤ 2θ ≤ 33◦ range with a scan step of 2 s/step
F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
2455
Table 1
Rietveld compositional analysis of 3Y-TZPs sintered for 2 h.
c-ZrO2 fraction (wt.%)
1350 ◦ C
1400 ◦ C
1450 ◦ C
1500 ◦ C
1550 ◦ C
Y2 O3 content in t-ZrO2 (mol%)
CZ-3Y
TZ-3Y
CZ-3Y-0.25Al
TZ-3Y-0.25Al
CZ-3Y
TZ-3Y
CZ-3Y-0.25Al
TZ-3Y-0.25Al
0.9
2.7
18.7
20.8
20.4
0.1
0.2
1.8
18
20.4
0.3
5.1
17.9
20.2
27.0
0
0.1
20.8
21.9
24.0
2.64
2.89
2.69
2.54
2.52
3.06
3.04
2.82
2.62
2.56
2.72
2.91
2.52
2.41
2.40
2.98
2.95
2.61
2.57
2.49
and a scan size of 0.02◦ . The volume fraction of monoclinic
phase was calculated according to the formula of Toraya et al.
[33]
Vm =
−111 + I 111 )
1.311(Im
m
−111 + I 111 )I 101
1.311(Im
t
m
(2)
with I, the intensity of monoclinic (−111 and 111) and tetragonal
(1 0 1) phase peaks indicated by the subscripts m and t.
The in-depth t–m transformation profile was measured on
polished cross-sections of 40 h hydrothermally treated specimen
using micro-Raman spectroscopy (Senterra, Bruker, Germany).
Raman scattering was excited with a laser at a wavelength of
532 nm through a 100× objective (lateral resolution ≤1 ␮m).
Line scans perpendicular to the surface with an interval of
1 ␮m were applied. The Raman spectra were recorded from 45
to 1500 cm−1 with 20 s integration time of 3 successive measurements per point. The monoclinic phase content (Vm ) was
calculated according to the formula of Clarke and Adar:34
Vm
179
(Im
179 + I 190
Im
m
190
+ Im ) + 0.97(It142
+ It256 )
(3)
The intensities (I) of the characteristic bands of the tetragonal
(142 and 256 cm−1 ) and monoclinic (179 and 190 cm−1 ) phase
were quantified with background subtracted spectra using OPUS
spectroscopy software.
The depth of the transformation zone was also measured by
SEM investigation of polished cross-sectional images of 40 h
hydrothermally treated specimens.
temperature increased, more cubic zirconia was formed and
the average yttria content in the remaining tetragonal phase
decreased, except for CZ-3Y and CZ-3Y-0.25Al sintered at
1350 ◦ C in which the average yttria content of tetragonal zirconia phase was low. The average yttria content in the tetragonal
zirconia phase of CZ-3Y and CZ-3Y-0.25Al ceramics was lower
than for TZ-3Y and TZ-3Y-0.25Al respectively.
3.2. Microstructure and grain size measurements
The average grain size, graphically presented in Fig. 1, of
all TZPs increased with increasing sintering temperature from
1350 ◦ C to 1550 ◦ C. The evolution was similar after 2 and 4 h sintering, but the grain size after sintering for 4 h was clearly larger.
The average grain size of CZ-3Y was slightly smaller than that
of TZ-3Y when sintered at 1350–1450 ◦ C, and slightly larger
at 1450–1550 ◦ C. The average grain sizes of CZ-3Y-0.25Al
and TZ-3Y-0.25Al were higher than that of CZ-3Y and TZ-3Y,
respectively. Although the grain size of CZ-3Y was only slightly
increased upon Al2 O3 addition, the grain growth was substantially accelerated for TZ-3Y-0.25Al. The average grain size of
TZ-3Y-0.25Al was considerably higher than that of CZ-3Y0.25Al and the largest of all investigated ceramic compositions
at any sintering temperature and time.
Representative microstructures of all investigated 3Y-TZP
ceramics are shown in Fig. 2. In the CZ-3Y and CZ-3Y-0.25Al
ceramics, some substantially larger grains were embedded in a
finer grained matrix. The larger grains had a size of 400–650 nm
for the ceramic sintered at 1350 ◦ C and up to 2 ␮m when sintered
at 1550 ◦ C. The ZrO2 grains were more bi-modally distributed
3. Results
3.1. Phase composition
The X-ray diffraction patterns of all investigated ZrO2 ceramics densified at different sintering temperatures indicated that
all ceramics were fully tetragonal zirconia (Y-TZP). However, a
small amount of cubic phase was possibly not clearly revealed
due to the peak overlap of the cubic and tetragonal ZrO2 phases.
Therefore, a Rietveld refinement of the XRD pattern was performed to quantitatively calculate the amount of cubic ZrO2
phase and the Y2 O3 content in the remaining tetragonal ZrO2
phase. The results of the Rietveld refined phase calculation of the
ceramics sintered for 2 h at different temperatures are summarized in Table 1. Cubic ZrO2 phase was formed in all 3Y-TZPs
at a sintering temperature of 1350–1550 ◦ C. As the sintering
Fig. 1. Average grain size of 3Y-TZPs sintered for 2 h or 4 h at different sintering
temperature.
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F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
Fig. 2. SEM images of CZ-3Y (a), CZ-3Y-0.25Al (b), TZ-3Y (c) and TZ-3Y-0.25Al (d) sintered for 2 h at 1450 ◦ C.
in the coated 3Y-TZPs (Fig. 2a and b) compared to the coprecipitated powder based ceramics (Fig. 2c and d). It can be
expected from the phase analysis (Table 1) that the larger ZrO2
grains were due to the formation of larger cubic phase ZrO2
grains.
STEM images and corresponding Y, Zr and Al elemental
mappings obtained by EDS-STEM measurements on CZ-3Y0.25Al and TZ-3Y-0.25Al are presented in Fig. 3. The Y3+
content around the grain boundary was higher compared to the
bulk of the grains in both CZ-3Y-0.25Al and TZ-3Y-0.25Al. Zr4+
was slightly depleted around the grain boundary, whereas Al3+
clearly segregated at the grain boundary of both CZ-3Y-0.25Al
and TZ-3Y-0.25Al.
The quantitative Y- and Al-concentration profiles along with
the Y/Zr ratio across the grain boundary are shown in Fig. 4. The
Al-concentration profiles were similar for CZ-3Y-0.25Al and
TZ-3Y-0.25Al (Fig. 4c). Al3+ segregated at the grain boundaries
over a width of 5 nm in both CZ-3Y-0.25Al and TZ-3Y-0.25Al,
and the concentration of segregated Al3+ was also similar. However, the Y3+ distribution was different. In CZ-3Y-0.25Al, Y3+
was more enriched at the edge of the grain, whereas the distribution of Y3+ was more homogeneous inside the grains of
Fig. 3. The STEM images for the grain boundary of CZ-3Y-0.25Al and TZ-3Y-0.25Al sintered for 2 h at 1550 ◦ C, and the corresponding elemental mapping of Y,
Al and Zr elements.
F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
2457
Fig. 4. Y-concentration (a), Y/Zr ratio (b) and Al-concentration (c) profiles across the grain boundaries in CZ-3Y-0.25Al and TZ-3Y-0.25Al.
TZ-3Y-0.25Al, which is clearly illustrated in the elemental mapping in Fig. 3. A clear peak was observed in the Y/Zr profile for
CZ-3Y-0.25Al, whereas the Y/Zr profile for TZ-3Y-0.25Al was
more spread (Fig. 4b). The Y/Zr ratio at the grain boundary
of CZ-3Y-0.25Al was high and higher than for TZ-3Y-0.25Al,
while the Y-concentration and the Y/Zr ratio in the core of the
grains in CZ-3Y-0.25Al were slightly lower than for TZ-3Y0.25Al.
Fig. 6a shows the evolution of fracture toughness as a function of the sintering temperature (2 h). The toughness of the TZPs
first gradually decreased with increasing sintering temperature
up to 1450 ◦ C and slightly increased with further increasing sintering temperature. The fracture toughness of the CZ-3Y and
CZ-3Y-0.25Al grades was higher than for the TZ-3Y and TZ3Y-0.25Al ceramics, independent on whether Al2 O3 was added
or not, implying that the fracture toughness of coated 3Y-TZPs
was higher than for co-precipitated 3Y-TZPs.
3.3. Mechanical properties
The evolution of the relative density and hardness as a function of the sintering temperature were similar after 2 h and 4 h
sintering, so only the evolution for ceramics sintered for 2 h
are shown in Fig. 5. Comparing Fig. 5a and b shows that the
hardness and density evolved in a similar way as a function of
the sintering temperature, initially increasing with increasing
sintering temperature up to 1450 ◦ C and slightly decreasing at
higher temperatures. The relative density and hardness of CZ3Y-0.25Al and TZ-3Y-0.25Al were higher than that of CZ-3Y
and TZ-3Y respectively, especially at sintering temperatures
below 1450 ◦ C, implying that the 0.25 wt.% Al2 O3 addition
substantially enhanced the densification for both coated and
co-precipitated powder based Y-TZPs. Furthermore, the relative density and hardness of 3Y coated TZPs were lower than
that of 3Y co-precipitated TZPs, independent on the alumina
addition.
Fig. 5. Relative density (a) and Vickers hardness (b) of the 3Y-TZPs as a function
of the sintering temperature.
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F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
Fig. 6. Fracture toughness of all 3Y-TZPs as a function of the sintering temperature (a) and average grain size (b).
Plotting the measured toughness as a function of the average grain size (Fig. 6b) revealed that the alumina addition did
not influence the fracture toughness of the co-precipitated and
coated 3Y-TZPs. A clear minimum in fracture toughness was
observed around an average ZrO2 grain size of 0.20–0.25 ␮m.
At an average grain size above 0.25 ␮m, the toughness of the
yttria-coated powder based ceramics was slightly higher than
for the co-precipitated powder based equivalents.
3.4. Low temperature degradation
3.4.1. Surface t–m transformation induced by
hydrothermal degradation
The monoclinic ZrO2 phase content, as measured by XRD, as
a function of the increasing hydrothermal treatment time on the
surface of CZ-3Y and CZ-3Y-0.25Al sintered at 1350–1550 ◦ C
for 2 and 4 h is presented in Fig. 7.
With increasing sintering temperature and time, the transformation was faster. A clear distinction could be made between
the TZPs sintered for 2 or 4 h ≤ 1400 ◦ C or for 2 h at 1450 ◦ C,
which were not or only modestly influenced by hydrothermal
degradation, and those sintered for 4 h at 1450 ◦ C or for 2 or
4 h ≥ 1500 ◦ C that were more susceptible to hydrothermal ageing. The monoclinic phase content of TZPs sintered at 1350 ◦ C
hardly changed with prolonged hydrothermal degradation time
Fig. 8. Fitted b parameter of the JAMK equation for all TZPs as function of the
average grain size.
up to 40 h and remained below 5 vol%. Materials sintered at
1400 ◦ C or at 1450 ◦ C for 2 h started to transform and the monoclinic phase content slowly increased with longer hydrothermal
degradation. The monoclinic phase content of the TZPs sintered
for 4 h at 1450 ◦ C and 2 or 4 h at 1500 ◦ C and 1550 ◦ C quickly
rose during the first 20 h of hydrothermal testing. A monoclinic
saturation level, which was lower than 100% due to the presence of non-transformable cubic ZrO2 , was reached after 30 h
of testing in CZ-3Y ceramics.
Comparing Fig. 7 for CZ-3Y and CZ-3Y-0.25Al showed
that CZ-3Y-0.25Al transformed at much slower rates than CZ3Y, especially when sintered at more elevated temperature.
0.25 wt.% alumina addition considerably increased the degradation resistance of CZ-3Y ceramics.
For the commercial co-precipitated powder based TZ-3Y and
TZ-3Y-0.25Al ceramics, a substantially enhanced degradation
was also observed when increasing the sintering time from 2 to
4 h at 1450 ◦ C, and TZ-3Y-0.25Al had a lower susceptibility to
ageing than the TZ-3Y.
The surface t–m transformation curves for all TZPs followed
a sigmoidal shape as a function of degradation time, implying
the surface degradation of yttria-coated powder based 3Y-TZPs
was determined by a nucleation and growth process,26 similar
to co-precipitated powder based 3Y-TZPs. The transformation
curves were fitted by the Johson–Mehl–Avrami–Kolmogorow
(JMAK) equation:26
Vm
= 1 − exp(−(bt)n )
Vms
Fig. 7. Surface monoclinic ZrO2 phase content as a function of hydrothermal
degradation time for CZ-3Y and CZ-3Y-0.25Al.
(4)
with Vms , the saturation level; b and n are parameters describing
the rate of the nucleation and growth, and spatial characteristics
of the crystallization process respectively.27,35
The kinetic b parameters for all grade TZPs sintered at different temperature–time combinations are plotted as a function
of the average ZrO2 grain size in Fig. 8. The b value increased
with increasing sintering temperature. It was negligible small
for all ceramics with a grain size ≤0.21 ␮m, but substantially
increased at grain sizes above 0.25 ␮m. Moreover, at grain sizes
F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
2459
Fig. 9. Monoclinic ZrO2 phase content as a function of the depth below the surface, as measured by Raman spectroscopy, after 40 h hydrothermal degradation (a)
and the surface monoclinic ZrO2 phase content profile as a function of degradation time as measured by XRD (b) for 4 different 3Y-TZPs sintered for 2 h at 1550 ◦ C.
above 0.21 ␮m, the b parameter for ceramics with 0.25 wt.%
Al2 O3 was significantly lower than for ceramics without Al2 O3 .
The b value of CZ-3Y and CZ-3Y-0.25Al was either as low
as that of TZ-3Y and TZ-3Y-0.25Al at smaller ZrO2 grain
size, i.e. at low sintering temperature, or significantly lower
than for TZ-3Y and TZ-3Y-0.25Al at larger ZrO2 grain size
(≥0.25/0.30 ␮m), i.e. at higher sintering temperature and/or prolonged sintering time. For Al2 O3 -free TZPs, the b value of
TZ-3Y grade was much larger than that of CZ-3Y at grain
sizes above 0.25 ␮m. Below the critical grain size of 0.25 ␮m,
however, the b values were nearly independent on the powder preparation method. For 0.25 wt.% alumina doped TZPs,
the b value of co-precipitated 3Y-TZPs (TZ-3Y-0.25Al) was
still larger than that of coated 3Y-TZPs (CZ-3Y-0.25Al) but the
critical grain size shifted to 0.30 ␮m due to the grain growth
acceleration effect of alumina addition. Therefore, coated 3YTZPs showed a higher surface degradation resistance than
co-precipitated 3Y-TZPs.
3.4.2. Depth of transformation in coated and
co-precipitated 3Y-TZPs
The value of the kinetic b parameter from the surface transformation curves revealed that coated 3Y-TZPs had a higher
degradation resistance than co-precipitated 3Y-TZPs, especially
at high sintering temperature. However, only the transformation on the top surface (<10 ␮m) was measured by XRD and
the transformation propagation into the bulk material, which
determines the final deterioration of Y-TZP material, was not
assessed. Therefore, the transformation propagation was measured by Raman spectroscopy to better compare the degradation
of coated and co-precipitated 3Y-TZPs.
Fig. 9a shows the depth transformation profiles acquired on
cross-sectioned 40 h hydrothermally treated 3Y-TZP specimens
sintered for 2 h at 1550 ◦ C. For comparison, the surface transformation profiles as a function of degradation time, as obtained by
XRD, are plotted in Fig. 9b. Fig. 9a shows that the monoclinic
zirconia content decreased in a non-linear way and dropped from
the saturation level to zero within a very short distance of less
than 5 ␮m.
A faster surface t–m transformation resulted in a deeper transformation propagation inside the 3Y-TZP specimens. TZ-3Y
degraded faster than CZ-3Y since the transformed depth for
TZ-3Y (about 40 ␮m) was much higher than for CZ-3Y (about
15 ␮m). Upon adding 0.25 wt.% alumina, the transformed zone
became much thinner for both CZ-3Y and TZ-3Y, and the transformation front was observed at about 2 ␮m and 7 ␮m below
the surface for CZ-3Y-0.25Al and TZ-3Y-0.25Al, respectively.
The degradation retarding effect of 0.25 wt.% alumina was more
pronounced for CZ-3Y ceramics, which resulted in the most stable CZ-3Y-0.25Al grade. XRD results (Fig. 9b) showed that the
surface of TZ-3Y, CY-3Y and TZ-3Y-0.25Al was saturated with
monoclinic ZrO2 after 40 h degradation, whereas the surface of
CZ-3Y-0.25Al was not yet saturated. No transformation saturation plateau was observed by micro-Raman measurement inside
the CZ-3Y-0.25Al (Fig. 9a), since the monoclinic ZrO2 content
decreased from about 25 vol% at the surface to 0 within a depth
of 3 ␮m.
Fig. 10 shows the corresponding cross-sectional SEM images
of 40 h hydrothermally treated 3Y-TZP specimens sintered for
2 h at 1550 ◦ C. Due to the extensive intergranular fracture and
the internal stresses from the volume expansion associated with
the degradation induced t–m transformation,36,37 a distinct border between degraded and pristine material was visible. The
transformed grains in the degraded layer were easily pulled out
during sample polishing and the transformed zone appeared
as a roughened layer, whereas the pristine bulk material was
smoothly polished and free of porosity. The thickness of the
observed degradation layer confirmed the results obtained by
XRD and Raman spectroscopy. After 40 h of hydrothermal treatment, the thickness of the degradation layer for CZ-3Y (about
12 ␮m) was much smaller than for TZ-3Y (about 36 ␮m), and
was reduced upon adding 0.25 wt.% alumina. Only about 3 ␮m
layer appeared to be transformed in TZ-3Y-0.25Al, and the
degradation was only slightly visible at a depth below 1 ␮m
in CZ-3Y-0.25Al.
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F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
Fig. 10. SEM images of the degraded zone from the cross-sections of 3Y-TZPs sintered for 2 h at 1550 ◦ C after 40 h hydrothermal degradation.
4. Discussion
The strategy of increasing the stabilizer content to improve
the LTD resistance of TZP materials reduced their potential for
stress-induced transformation. However, the excellent fracture
toughness attributed to the stress-induced transformation toughening is one of the important reasons for Y-TZP ceramics to be
used in prosthetic dentistry. Therefore, it is important to develop
hydrothermally stable Y-TZP materials without compromising
on the fracture toughness.
In this work, 3 methods were observed to be able to improve
the LTD resistance of 3Y-TZPs, i.e. decreasing the sintering
temperature, adding 0.25 wt.% alumina, and incorporating the
stabilizer by yttria coating of the starting ZrO2 powder. In the discussion below, the influence of these 3 factors on the mechanical
properties and LTD behaviour of 3Y-TZPs are highlighted.
4.1. The influence of sintering temperature
The sintering temperature considerably influenced the
mechanical properties (Figs. 5 and 6) and hydrothermal degradation (Figs. 7 and 8) of all investigated 3Y-TZP grades in
a similar way. The degradation susceptibility of all 3Y-TZPs
significantly increased with increasing sintering temperature,
which can be attributed to an increased tetragonal ZrO2 grain
size (Figs. 1 and 8), a larger fraction of cubic zirconia and
a decreased average yttria stabilizer content in the remaining
tetragonal grains (Table 1). This result is in agreement with an
earlier report claiming that the presence of cubic grains had a
harmful impact on the LTD resistance of Y-TZPs, and cubic
grains were enriched with yttrium concomitantly resulting in a
decreased yttrium content in the remaining tetragonal grains.38
Therefore, the increased LTD resistance of 3Y-TZPs obtained
by decreasing the sintering temperature is actually achieved by
increasing the average yttria stabilizer content in the tetragonal zirconia phase. Despite the increased LTD resistance, it
can be expected that decreasing the sintering temperature would
decrease the driving force for transformation toughening, compromising on the fracture toughness of the 3Y-TZP. This was
indeed confirmed in Fig. 6, illustrating that the toughness of
all 3Y-TZPs decreased with decreasing sintering temperature or
decreasing grain size at a sintering temperature above 1450 ◦ C or
at a ZrO2 grain size above 0.20 ␮m. Earlier studies have clearly
reported that decreasing the sintering temperature, i.e. decreasing the grain size of Y-TZP ceramics, significantly decreased
their transformability and concomitant fracture toughness.18,39
Although it is shown in Fig. 6 that the fracture toughness
of 3Y-TZPs increased again upon further decreasing the sintering temperature from 1450 ◦ C to 1350 ◦ C, the sintering of
3Y-TZPs and especially alumina-free TZPs at a temperature
below 1450 ◦ C resulted in residual porosity that compromised
on the hardness. The increased fracture toughness could only
be attributed to an increased amount of closed pores. Fig. 5
shows that 1450 ◦ C was the optimum sintering temperature for
all investigated Y-TZPs based on density and hardness. The
critical grain size between 0.21 and 0.25 ␮m above which the
hydrothermal degradation is dramatically enhanced (Fig. 8) was
reached for all Y-TZPs when sintered for 2 and 4 h at 1450 ◦ C
(Fig. 7).
In summary, 3Y-TZPs and especially alumina-free TZPs
should be sintered ≥1450 ◦ C to obtain full densification and
high hardness. From the LTD resistance point of view however,
3Y-TZPs have to be sintered ≤1450 ◦ C. When 3Y-TZPs were
sintered at 1450 ◦ C, their fracture toughness reached a minimum
F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
at an average ZrO2 grain size of 0.20–0.25 ␮m (Fig. 6). Therefore, fully dense 3Y-TZP ceramics with high LTD resistance
and high toughness cannot be obtained by only adjusting the
sintering condition.
4.2. The influence of 0.25 wt.% alumina addition
The addition of 0.25 wt.% Al2 O3 accelerated the densification and increased the hardness of both CZ-3Y and TZ-3Y
materials when sintered at 1350–1400 ◦ C (Fig. 5). In addition,
Al2 O3 did not influence the fracture toughness of 3Y-TZPs
(Fig. 6b). Therefore, upon adding 0.25 wt.% Al2 O3 , fully dense
3Y-TZPs could be sintered at lower temperature to obtain a
higher LTD resistance without compromising on the hardness
and fracture toughness.
Moreover, the alumina addition itself decreased the susceptibility of both coated and co-precipitated 3Y-TZPs to
hydrothermal degradation, which was clearly shown by the fact
that the kinetic parameter b value in the JMAK equation for
both CZ-3Y and TZ-3Y decreased upon adding 0.25 wt.% Al2 O3
(Fig. 8). The transformation propagation inside CZ-3Y and TZ3Y was also significantly retarded (Figs. 9 and 10). Although it is
commonly assumed that the addition of Al2 O3 decreases the susceptibility of Y-TZPs to LTD due to a decreased grain size,40 the
addition of 0.25 wt.% Al2 O3 in this work was found to increase
the average ZrO2 grain size of both CZ-3Y and TZ-3Y ceramics
(Fig. 2). Therefore, the LTD retarding effect of Al2 O3 addition cannot be attributed to a grain size reduction.TEM results
(Figs. 3 and 4) clearly showed that Al3+ segregated at the ZrO2
grain boundaries over a width of 5 nm in both CZ-3Y-0.25Al
and TZ-3Y-0.25Al. This is in agreement with literature reports
claiming that a small (<0.3 wt.%) amount of Al2 O3 can dissolve
in the zirconia grains during sintering, resulting in the segregation of Al3+ ions at the grain boundaries and Al3+ was reported to
increase the Y3+ segregation at the grain boundary in Y-TZP.41
The grain boundary region is crucial in the degradation of YTZPs because it is believed to be the nucleation site for the
t–m transformation and the path for transformation to propagate
during degradation.42–44
In summary, a grain boundary with a small amount of segregated aluminium in solid solution enhances the hydrothermal
degradation resistance of 3Y-TZP ceramics without compromising on the fracture toughness.
4.3. The influence of incorporating the stabilizer by Y2 O3
coating the starting ZrO2 powder
Fig. 8 clearly shows that coated 3Y-TZPs had a higher
surface degradation resistance than co-precipitated 3Y-TZPs,
the transformation propagation profiles (Fig. 9) and images
(Fig. 10) confirmed the higher degradation resistance of coated
3Y-TZPs than that of co-precipitated 3Y-TZPs, independent
on the alumina addition. Furthermore, the fracture toughness
of yttria-coated powder based Y-TZPs was higher than that of
the yttria co-precipitated powder based ceramics, independent
on the Al2 O3 addition (Fig. 6). Therefore, coated 3Y-TZPs
2461
combine a higher LTD resistance and a higher fracture
toughness compared to co-precipitated 3Y-TZPs.
The advantage of incorporating the stabilizer by yttria coating of the starting ZrO2 powder is believed to be related to the
heterogeneously distributed Y3+ at the grain boundary. TEM
investigation (Figs. 4 and 5) showed that Y3+ was more enriched
at the edge of CZ-3Y-0.25Al grains, whereas the Y3+ distribution was more homogeneous in the TZ-3Y-0.25Al grains. A clear
peak in the Y/Zr ratio was observed at the CZ-3Y-0.25Al grain
boundaries, which was not the case at the TZ-3Y-0.25Al grain
boundaries. This was definitely due to the radically different
locations of the yttria stabilizer in the starting ZrO2 powder.
Y2 O3 stabilizer was already inside the ZrO2 grain for yttria
co-precipitated powders, so Y3+ segregated towards the grain
boundary from the bulk of TZ-3Y-0.25Al, whereas for CZ3Y-0.25Al the coated Y2 O3 layers had to dissolve and diffuse
into the ZrO2 grain from the surface during the sintering process, which might also contribute to the low yttria content in
the tetragonal zirconia phase of CZ-3Y and CZ-3Y-0.25Al sintered at 1350 ◦ C. Due to the slow diffusion of Y in zirconia at
1350 ◦ C,45,46 it is possible that only part of Y was dissolved and
diffused into the zirconia phase and the rest of Y2 O3 still located
at the surface of zirconia.
Due to the higher amount of yttria located at the grain boundary in coated 3Y-TZPs, the grain boundary of yttria coated
TZPs was more stable than that of yttria co-precipitated TZPs.
It therefore reduces the susceptibility of the coated powder
based ceramic towards hydrothermal degradation. As explained
before, the grain boundary stability is very important for the stability of the complete TZP material because grain boundaries are
believed to be the starting point for transformation and also to
act as the preferred path for water radicals to propagate into the
material.42–44 A stable ZrO2 grain boundary with high Y3+ , i.e.
a high Y/Zr at the grain boundary, is beneficial for increasing the
degradation resistance of TZPs, and it is therefore not necessary
to homogeneously increase the transformation resistance of the
bulk of the grains by increasing the overall yttria content. Fig. 9a
and b clearly shows that the surface transformation in the initial stage of degradation and also the depth transformation after
reaching the saturation point were slower for coated 3Y-TZPs
compared to co-precipitated 3Y-TZPs. The higher degradation
resistance of CZ-3Y-0.25Al could was also enhanced by the finer
microstructure compared to TZ-3Y-0.25Al (Fig. 1).
The higher fracture toughness of coated 3Y-TZPs was
also attributed to the non-homogeneously distributed yttria.
In coated 3Y-TZPs, more yttria was located at the grain
boundary, which in turn lead to a lower amount of yttria in
the core of the grains (Fig. 4b and Table 1). The grain core
was responsible for the higher toughness due to an enhanced
transformation toughening contribution.24,29,47 Moreover, the
low yttria content core is not located at the critical point of
the degradation, i.e. grain boundary, and it is protected by the
more stable grain boundary. Thus, the high transformability of
the core did not enhance the LTD susceptibility. In addition,
yttria-coated starting powder based Y-TZP was reported to
be able to have a core–shell grain structure,23 i.e. grains with
an yttria-enriched tetragonal shell and a lower yttria content
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F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463
core, which is sensitive to transformation and responsible for
a higher toughness.48 The higher fracture toughness of coated
3Y-TZPs was also revealed from the results of a Rietveld
refinement, indicating that the average yttria content in the
tetragonal zirconia phase of the coated Y-TZPs was lower than
for co-precipitated Y-TZPs (Table 1). A similar inhomogeneous
yttria and ceria distribution was reported for stabilizer coated
starting powder based TZPs.25 The yttria content distribution in
coated TZPs sintered from yttrium nitrate coated ZrO2 powders
was measured to vary from 2 to 8 mol%, which was broader and
more inhomogeneous than for co-precipitated powder based
3Y-TZP sintered under exactly the same conditions.24
Therefore, incorporating stabilizer by coating yttria in the
starting powder can be a new way to optimize the degradation
resistance and also the fracture toughness of 3Y-TZP within the
same material.
In summary, in order to co-optimize the LTD resistance and
mechanical properties, it was essential to add 0.25 wt.% Al2 O3
and incorporate the stabilizer by an alternative Y2 O3 coating
route, especially when the average ZrO2 grain size was above
0.25 ␮m.
5. Conclusions
Increasing the sintering temperature significantly enhanced
the hydrothermal degradation of 3Y-TZP ceramics due to an
increased tetragonal ZrO2 grain size, a larger fraction of cubic
zirconia and a decreased average yttria stabilizer content in
the remaining tetragonal grains. Only limited degradation was
observed at ZrO2 grain sizes below 0.20 and 0.25 ␮m for the
Al2 O3 -free and 0.25 wt.% Al2 O3 -doped ceramics, respectively.
When sintered for 4 h at 1450 ◦ C or higher, corresponding to an
average grain size above 0.25 ␮m, degradation of 3Y-TZPs was
considerably increased. However, decreasing the sintering temperature did not allow obtaining a degradation resistant 3Y-TZP
with simultaneously optimized mechanical properties.
Al2 O3 addition and incorporating the stabilizer by yttria coating of the ZrO2 starting powder had a pronounced effect on
retarding the degradation without compromising on the transformation induced fracture toughness. This could be attributed to
the segregation of Al3+ and the heterogeneously distributed Y3+
at the grain boundary, respectively. 3Y-TZPs made from yttriacoated ZrO2 starting powder had a high Y/Zr ratio at the grain
boundary and lower yttria content in the core of the grain, thereby
combining a higher LTD resistance and higher fracture toughness compared to yttria co-precipitated ZrO2 starting powder
based ceramics. The cumulative positive effect of the addition
of 0.25 wt.% Al2 O3 as well as the use of yttria-coated starting
powder resulted in an enhanced resistance against low temperature degradation of the CZ-3Y-0.25Al ceramic with optimized
fracture toughness.
Acknowledgements
This work was performed within the framework of the
Research Fund of KU Leuven under project 0T/10/052 and
the Fund for Scientific Research Flanders (FWO) under grant
G.0431.10N. K. Vanmeensel thanks the Fund for Scientific
Research Flanders (FWO) for his postdoctoral fellowship.
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