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Available online at www.sciencedirect.com

ScienceDirect

Journal of the European Ceramic Society 34 (2014) 2453–2463

3Y-TZP ceramics with improved hydrothermal degradation resistanceand 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

a Department of Metallurgy and Materials Engineering, KU Leuven, Kasteelpark Arenberg 44, B-3001 Leuven, Belgiumb KU Leuven BIOMAT, Department of Oral Health Sciences, KU Leuven & Dentistry, University Hospitals Leuven, Kapucijnenvoer 7 blok a, B-3000 Leuven,

Belgiumc 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 2014Available online 6 March 2014

bstract

ifferent factors such as the way of incorporating the Y2O3 stabilizer, alumina addition and sintering temperature were assessed with the goal tomprove the low temperature degradation (LTD) resistance of 3Y-TZP without compromising on the mechanical properties. The degradation ofydrothermally 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 withptimized mechanical properties. Alumina addition along with the use of Y2O3 stabilizer coated starting powder allowed to combine both anxcellent toughness and LTD resistance, as compared to alumina-free and stabilizer co-precipitated powder based equivalents. Transmission

3+

lectron microscopy revealed that the improved LTD resistance could be attributed to the segregation of Al at the grain boundary and theeterogeneously distributed Y3+ stabilizer.

2014 Elsevier Ltd. All rights reserved.

aodzbfase

p

eywords: Coated 3Y-TZP; Co-precipitated 3Y-TZP; Degradation; Alumina

. Introduction

Yttria stabilized tetragonal zirconia polycrystals (Y-TZP) arencreasingly used as dental restorative material in recent yearsecause they possess excellent mechanical properties, such asuperior strength and fracture toughness while they also exhibitiocompatibility and aesthetic potential.1–3 The high strengthnd toughness of zirconia ceramics are due to stress-inducedhase transformation toughening. Around a propagating crack,he metastable tetragonal grains can transform to the mono-linic structure, t–m transformation, and the associated volume

xpansion induces compressive stresses and eventually reducesr inhibits further crack propagation.2,4 However, the tetrago-al grains can also spontaneously transform to monoclinic in

∗ Corresponding author. Tel.: +32 16321192.E-mail address: kim.vanmeensel@mtm.kuleuven.be (K. Vanmeensel).

YtTtHuc

ttp://dx.doi.org/10.1016/j.jeurceramsoc.2014.02.026955-2219/© 2014 Elsevier Ltd. All rights reserved.

humid environment in the 20–300 ◦C range,5,6 like in theral cavity, a phenomenon being referred to as low temperatureegradation (LTD).5 The mechanism of water incorporating intoirconia to induce t–m transformation is not yet fully understood,ut it is well established that the degradation starts from the sur-ace and proceeds inwards. Surface uplift7,8 and micro-cracks9

re subsequently induced which can result in macrocracks9 andurface roughening.10 Consequently, the mechanical10–13 andven the aesthetic properties14 of Y-TZPs are affected.

The LTD susceptibility of Y-TZP is affected by differentarameters. The degradation can be retarded by increasing the2O3 content.5 For dental applications, 3 mol% yttria stabilized

etragonal zirconia polycrystals (3Y-TZP) are commonly used.15

he sintering condition is a predominant factor influencing bothhe mechanical properties and LTD behaviour of 3Y-TZPs.15–19

owever, the mechanism was not well studied and different man-facturers produce dental 3Y-TZP materials at different sinteringonditions.3

2 an C

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oLttm

tpaiiyac

2

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ad

3aplctoZ

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454 F. Zhang et al. / Journal of the Europe

It was reported that high purity ZrO2 starting powder waseeded to inhibit the transformation20 but it was also claimedhat the beneficial effect of additives like TiO2, Fe2O3 and Al2O3epends on their amount.20 A small amount of homogeneouslyistributed alumina (<0.5 wt.%) can decrease the LTD suscep-ibility of 3Y-TZP.21 Commercially available dental ZrO2 isherefore commonly doped with 0.25 wt.% alumina.

A homogeneous distribution of the Y2O3 stabilizer is gener-lly believed to increase the degradation resistance of TZPs22

ut different opinions are recently challenging this.5 Moreover,he mechanical properties of yttria-coated Y-TZP were reportedo be more attractive.17,23,24 Due to the inhomogeneous Y2O3istribution, coated Y-TZP had a higher toughness compared too-precipitated Y-TZP.23 Moreover, coated Y-TZP is expectedo have a higher LTD resistance.17 A simple nitrate processas been reported for the processing of stabilizer-coated ZrO2tarting powder, resulting in TZPs with excellent toughness andnhomogeneous stabilizer distribution.24,25 The LTD behaviourf yttria-coated powder based TZPs has however not been stud-ed.

Furthermore, studies on dental 3Y-TZP were focused eithernly on the mechanical properties17,18,24 or only on theTD behaviour.6,16,19,26,27 It is however important to preservehe excellent mechanical properties, especially the fractureoughness, when developing hydrothermally stable Y-TZP bio-aterials.The aim of this work was to investigate the influence of

he sintering condition, alumina addition and the way of incor-orating the Y2O3 stabilizer on the mechanical propertiesnd LTD behaviour of 3Y-TZPs. 3 mol% yttria stabilizer wasncorporated in the starting ZrO2 powder by a nitrate coat-ng technique as compared to commercially available 3 mol%ttria co-precipitated ZrO2 powder. Alumina-free and 0.25 wt.%lumina-doped yttria-coated and co-precipitated powder basederamics were investigated.

. Experimental procedures

3 mol% Y2O3 coated ZrO2 powder was prepared accordingo the nitrate route.28,29 The proper amount of Y2O3 (99.9%,cros, Geel, Belgium) was dissolved in warm nitric acid (65%,igma–Aldrich, Bornem, Belgium) and mixed with pure mono-linic ZrO2 nanopowder (grade TZ-0, Tosoh, Japan) in ethanoln a polyethylene container on a multidirectional mixer (Tur-ula, Basel, Switzerland) for 24 h. An appropriate amount ofl(NO3)3·9H2O (Aldrich Chemical Company) was added to

he suspension to obtain 0.25 wt.% Al2O3 doped Y2O3-coatedrO2 powder. Y-TZP milling balls (grade TZ-3Y, Tosoh, Japan)ith a diameter of 5 mm were added to the suspension to break

he ZrO2 and Al2O3 agglomerates during mixing. After drying,he yttrium nitrate coated powder was calcined in air for 30 mint 800 ◦C. A second multidirectional mixing was performed inthanol for 48 h to break the hard agglomerates after calcina-

ion. Ceramic discs were obtained by cold isostatic pressingCIP) at 300 MPa and subsequent pressureless sintering in air for

or 4 h at 1350–1550 ◦C. Reference ceramics were processedrom 3 mol% Y2O3 co-precipitated ZrO2 powder (grade TZ-3Y,

4sps

eramic Society 34 (2014) 2453–2463

osoh, Japan) and the 0.25 wt.% Al2O3 containing equivalentgrade TZ-3Y-E, Tosoh, Japan).

The nomenclature of the experimental 3Y-coated ZrO2,.25 wt.% Al2O3 doped 3Y-coated ZrO2 and commercial pow-er based 3Y-coprecipitated ZrO2 and 0.25 wt.% Al2O3 dopedY-coprecipitated ZrO2 ceramics used throughout the text isZ-3Y, CZ-3Y-0.25Al, TZ-3Y and TZ-3Y-0.25Al, respectively.

The density of the sintered ceramics was measured in ethanolccording to the Archimedes principle, assuming a theoreticalensity of 6.05 g/cm3 to calculate the relative density.30

Phase identification was done by X-ray diffraction (XRD,003-TT, Seifert, Ahrensburg, Germany) using Cu K� radiationt 40 kV and 40 mA. Rietveld analysis of the XRD pattern waserformed with X’Pert Highscore Plus to quantitatively calcu-ate the amount of tetragonal (P42/nmc (1 3 7) space group) andubic ZrO2 phase (Fm-3m (2 2 5) space group). The Y2O3 con-ent in the remaining tetragonal ZrO2 phase was calculated basedn the calculated a and c unit cell parameters of the tetragonalrO2 phase according to the formula:31

O1.5 (mol%) = 1.02311 − (ct/√

2at)

0.001498(1)

The microstructure was analyzed by scanning elec-ron microscopy (SEM, XL-30FEG, FEI, Eindhoven, Theetherlands). Cross-sectioned ceramics were polished and ther-ally etched for 25 min at 1250 ◦C in air. SEM images were

sed to measure the average grain size of the sintered ceramicssing IMAGE-PRO software according to the linear interceptethod. The reported grain sizes are the as-measured values

btained from at least 1000 grains, without any correction.Scanning transmission electron microscopy (STEM) images,

lemental mapping and energy-dispersive spectroscopy (EDS)ere performed with a FEI Titan 60-300 “cubed” transmission

lectron microscope to examine the distribution of Y3+, Al3+ andr4+ around the grain boundaries. Electron transparent samplesere prepared by ion-milling with an Ion Slicer (EM-09100IS,

eol, Japan). The transmission electron microscope was operatedt 200 kV and a high magnification having a mapping resolutionelow 0.19 nm. 5–7 grain boundaries in each specimen werenalyzed. The quantitative elemental mapping was acquired toalculate the Y3+, Al3+ and Zr4+ concentration profile usingSPRIT 1.9 software.

The Vickers hardness was measured on a hardness testerModel FV-700, Future-Tech Corp., Tokyo, Japan) with a loadf 9.8 N and a dwell time of 10 s. The indentation toughnessas evaluated from the radial crack pattern accompanying theickers indentations and calculated according to the Anstisquation.32 The reported values are the mean and standard devi-tion of 10 indentations. An E-modulus value of 210 GPa wassed to calculate the fracture toughness of all ceramics.

For low temperature degradation (LTD) testing, acceleratedydrothermal degradation testing was used and mirror polishedpecimens were autoclaved in steam at 134 ◦C and 0.2 MPa up to

0 h. At predefined times, the test was interrupted to measure theurface monoclinic ZrO2 phase content by means of XRD. XRDatterns were recorded on both polished flat surfaces of eachpecimen 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 1Rietveld compositional analysis of 3Y-TZPs sintered for 2 h.

c-ZrO2 fraction (wt.%) Y2O3 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

1350 ◦C 0.9 0.1 0.3 0 2.64 3.06 2.72 2.981400 ◦C 2.7 0.2 5.1 0.1 2.89 3.04 2.91 2.951450 ◦C 18.7 1.8 17.9 20.8 2.69 2.82 2.52 2.611 ◦1

ap[

V

w(

puR5L1tsc

V

(ws

Sh

3

3

iasdTfppcra

ttd1nzt

3

a1tToaaritT0a

ccfiner grained matrix. The larger grains had a size of 400–650 nmfor the ceramic sintered at 1350 ◦C and up to 2 �m when sinteredat 1550 ◦C. The ZrO2 grains were more bi-modally distributed

500 C 20.8 18 20.2 21.9

550 ◦C 20.4 20.4 27.0 24.0

nd a scan size of 0.02◦. The volume fraction of monoclinichase was calculated according to the formula of Toraya et al.33]

m = 1.311(I−111m + I111

m )

1.311(I−111m + I111

m )I101t

(2)

ith I, the intensity of monoclinic (−111 and 111) and tetragonal1 0 1) phase peaks indicated by the subscripts m and t.

The in-depth t–m transformation profile was measured onolished cross-sections of 40 h hydrothermally treated specimensing micro-Raman spectroscopy (Senterra, Bruker, Germany).aman scattering was excited with a laser at a wavelength of32 nm through a 100× objective (lateral resolution ≤1 �m).ine scans perpendicular to the surface with an interval of

�m were applied. The Raman spectra were recorded from 45o 1500 cm−1 with 20 s integration time of 3 successive mea-urements per point. The monoclinic phase content (Vm) wasalculated according to the formula of Clarke and Adar:34

m

I179m + I190

m

(I179m + I190

m ) + 0.97(I142t + I256

t )(3)

The intensities (I) of the characteristic bands of the tetragonal142 and 256 cm−1) and monoclinic (179 and 190 cm−1) phaseere quantified with background subtracted spectra using OPUS

pectroscopy software.The depth of the transformation zone was also measured by

EM investigation of polished cross-sectional images of 40 hydrothermally treated specimens.

. Results

.1. Phase composition

The X-ray diffraction patterns of all investigated ZrO2 ceram-cs densified at different sintering temperatures indicated thatll ceramics were fully tetragonal zirconia (Y-TZP). However, amall amount of cubic phase was possibly not clearly revealedue to the peak overlap of the cubic and tetragonal ZrO2 phases.herefore, a Rietveld refinement of the XRD pattern was per-

ormed to quantitatively calculate the amount of cubic ZrO2hase and the Y2O3 content in the remaining tetragonal ZrO2

hase. The results of the Rietveld refined phase calculation of theeramics sintered for 2 h at different temperatures are summa-ized in Table 1. Cubic ZrO2 phase was formed in all 3Y-TZPst a sintering temperature of 1350–1550 ◦C. As the sintering

Ft

2.54 2.62 2.41 2.572.52 2.56 2.40 2.49

emperature increased, more cubic zirconia was formed andhe average yttria content in the remaining tetragonal phaseecreased, except for CZ-3Y and CZ-3Y-0.25Al sintered at350 ◦C in which the average yttria content of tetragonal zirco-ia phase was low. The average yttria content in the tetragonalirconia phase of CZ-3Y and CZ-3Y-0.25Al ceramics was lowerhan for TZ-3Y and TZ-3Y-0.25Al respectively.

.2. Microstructure and grain size measurements

The average grain size, graphically presented in Fig. 1, ofll TZPs increased with increasing sintering temperature from350 ◦C to 1550 ◦C. The evolution was similar after 2 and 4 h sin-ering, but the grain size after sintering for 4 h was clearly larger.he average grain size of CZ-3Y was slightly smaller than thatf TZ-3Y when sintered at 1350–1450 ◦C, and slightly largert 1450–1550 ◦C. The average grain sizes of CZ-3Y-0.25Alnd TZ-3Y-0.25Al were higher than that of CZ-3Y and TZ-3Y,espectively. Although the grain size of CZ-3Y was only slightlyncreased upon Al2O3 addition, the grain growth was substan-ially accelerated for TZ-3Y-0.25Al. The average grain size ofZ-3Y-0.25Al was considerably higher than that of CZ-3Y-.25Al and the largest of all investigated ceramic compositionst any sintering temperature and time.

Representative microstructures of all investigated 3Y-TZPeramics are shown in Fig. 2. In the CZ-3Y and CZ-3Y-0.25Aleramics, some substantially larger grains were embedded in a

ig. 1. Average grain size of 3Y-TZPs sintered for 2 h or 4 h at different sinteringemperature.

2456 F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463

Z-3Y

ipegg

m0cbw

ca

tAToa

FA

Fig. 2. SEM images of CZ-3Y (a), CZ-3Y-0.25Al (b), T

n the coated 3Y-TZPs (Fig. 2a and b) compared to the co-recipitated powder based ceramics (Fig. 2c and d). It can bexpected from the phase analysis (Table 1) that the larger ZrO2rains were due to the formation of larger cubic phase ZrO2rains.

STEM images and corresponding Y, Zr and Al elementalappings obtained by EDS-STEM measurements on CZ-3Y-

.25Al and TZ-3Y-0.25Al are presented in Fig. 3. The Y3+

ontent around the grain boundary was higher compared to theulk of the grains in both CZ-3Y-0.25Al and TZ-3Y-0.25Al. Zr4+

as slightly depleted around the grain boundary, whereas Al3+

ewt

ig. 3. The STEM images for the grain boundary of CZ-3Y-0.25Al and TZ-3Y-0.25l and Zr elements.

(c) and TZ-3Y-0.25Al (d) sintered for 2 h at 1450 ◦C.

learly segregated at the grain boundary of both CZ-3Y-0.25Alnd TZ-3Y-0.25Al.

The quantitative Y- and Al-concentration profiles along withhe Y/Zr ratio across the grain boundary are shown in Fig. 4. Thel-concentration profiles were similar for CZ-3Y-0.25Al andZ-3Y-0.25Al (Fig. 4c). Al3+ segregated at the grain boundariesver a width of 5 nm in both CZ-3Y-0.25Al and TZ-3Y-0.25Al,nd the concentration of segregated Al3+ was also similar. How-

ver, the Y3+ distribution was different. In CZ-3Y-0.25Al, Y3+

as more enriched at the edge of the grain, whereas the dis-ribution of Y3+ was more homogeneous inside the grains of

Al sintered for 2 h at 1550 ◦C, and the corresponding elemental mapping of Y,

F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463 2457

profile

TpCmowg0

3

tsahtsh3absctta

tfiutC3owas higher than for co-precipitated 3Y-TZPs.

Fig. 4. Y-concentration (a), Y/Zr ratio (b) and Al-concentration (c)

Z-3Y-0.25Al, which is clearly illustrated in the elemental map-ing in Fig. 3. A clear peak was observed in the Y/Zr profile forZ-3Y-0.25Al, whereas the Y/Zr profile for TZ-3Y-0.25Al wasore spread (Fig. 4b). The Y/Zr ratio at the grain boundary

f CZ-3Y-0.25Al was high and higher than for TZ-3Y-0.25Al,hile the Y-concentration and the Y/Zr ratio in the core of therains in CZ-3Y-0.25Al were slightly lower than for TZ-3Y-.25Al.

.3. Mechanical properties

The evolution of the relative density and hardness as a func-ion of the sintering temperature were similar after 2 h and 4 hintering, so only the evolution for ceramics sintered for 2 hre shown in Fig. 5. Comparing Fig. 5a and b shows that theardness and density evolved in a similar way as a function ofhe sintering temperature, initially increasing with increasingintering temperature up to 1450 ◦C and slightly decreasing atigher temperatures. The relative density and hardness of CZ-Y-0.25Al and TZ-3Y-0.25Al were higher than that of CZ-3Ynd TZ-3Y respectively, especially at sintering temperatureselow 1450 ◦C, implying that the 0.25 wt.% Al2O3 additionubstantially enhanced the densification for both coated and

o-precipitated powder based Y-TZPs. Furthermore, the rela-ive density and hardness of 3Y coated TZPs were lower thanhat of 3Y co-precipitated TZPs, independent on the aluminaddition.

Fo

s across the grain boundaries in CZ-3Y-0.25Al and TZ-3Y-0.25Al.

Fig. 6a shows the evolution of fracture toughness as a func-ion of the sintering temperature (2 h). The toughness of the TZPsrst gradually decreased with increasing sintering temperaturep to 1450 ◦C and slightly increased with further increasing sin-ering temperature. The fracture toughness of the CZ-3Y andZ-3Y-0.25Al grades was higher than for the TZ-3Y and TZ-Y-0.25Al ceramics, independent on whether Al2O3 was addedr not, implying that the fracture toughness of coated 3Y-TZPs

ig. 5. Relative density (a) and Vickers hardness (b) of the 3Y-TZPs as a functionf the sintering temperature.

2458 F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463

Fa

ancoAyf

3

3h

asf

mtwd4ih

Fd

Fa

u1cdfrseo

t30d

Tw4a

ig. 6. Fracture toughness of all 3Y-TZPs as a function of the sintering temper-ture (a) and average grain size (b).

Plotting the measured toughness as a function of the aver-ge grain size (Fig. 6b) revealed that the alumina addition didot influence the fracture toughness of the co-precipitated andoated 3Y-TZPs. A clear minimum in fracture toughness wasbserved around an average ZrO2 grain size of 0.20–0.25 �m.t an average grain size above 0.25 �m, the toughness of thettria-coated powder based ceramics was slightly higher thanor the co-precipitated powder based equivalents.

.4. Low temperature degradation

.4.1. Surface t–m transformation induced byydrothermal degradation

The monoclinic ZrO2 phase content, as measured by XRD, as function of the increasing hydrothermal treatment time on theurface of CZ-3Y and CZ-3Y-0.25Al sintered at 1350–1550 ◦Cor 2 and 4 h is presented in Fig. 7.

With increasing sintering temperature and time, the transfor-ation was faster. A clear distinction could be made between

he TZPs sintered for 2 or 4 h ≤ 1400 ◦C or for 2 h at 1450 ◦C,hich were not or only modestly influenced by hydrothermalegradation, and those sintered for 4 h at 1450 ◦C or for 2 or

h ≥ 1500 ◦C that were more susceptible to hydrothermal age-ng. The monoclinic phase content of TZPs sintered at 1350 ◦Cardly changed with prolonged hydrothermal degradation time

ig. 7. Surface monoclinic ZrO2 phase content as a function of hydrothermalegradation time for CZ-3Y and CZ-3Y-0.25Al.

atwtc(

wto

fowfi

ig. 8. Fitted b parameter of the JAMK equation for all TZPs as function of theverage grain size.

p to 40 h and remained below 5 vol%. Materials sintered at400 ◦C or at 1450 ◦C for 2 h started to transform and the mono-linic phase content slowly increased with longer hydrothermalegradation. The monoclinic phase content of the TZPs sinteredor 4 h at 1450 ◦C and 2 or 4 h at 1500 ◦C and 1550 ◦C quicklyose during the first 20 h of hydrothermal testing. A monoclinicaturation level, which was lower than 100% due to the pres-nce of non-transformable cubic ZrO2, was reached after 30 hf testing in CZ-3Y ceramics.

Comparing Fig. 7 for CZ-3Y and CZ-3Y-0.25Al showedhat CZ-3Y-0.25Al transformed at much slower rates than CZ-Y, especially when sintered at more elevated temperature..25 wt.% alumina addition considerably increased the degra-ation resistance of CZ-3Y ceramics.

For the commercial co-precipitated powder based TZ-3Y andZ-3Y-0.25Al ceramics, a substantially enhanced degradationas also observed when increasing the sintering time from 2 to

h at 1450 ◦C, and TZ-3Y-0.25Al had a lower susceptibility togeing than the TZ-3Y.

The surface t–m transformation curves for all TZPs followed sigmoidal shape as a function of degradation time, implyinghe surface degradation of yttria-coated powder based 3Y-TZPsas determined by a nucleation and growth process,26 similar

o co-precipitated powder based 3Y-TZPs. The transformationurves were fitted by the Johson–Mehl–Avrami–KolmogorowJMAK) equation:26

Vm

Vms

= 1 − exp(−(bt)n) (4)

ith Vms, the saturation level; b and n are parameters describinghe rate of the nucleation and growth, and spatial characteristicsf the crystallization process respectively.27,35

The kinetic b parameters for all grade TZPs sintered at dif-erent temperature–time combinations are plotted as a functionf the average ZrO2 grain size in Fig. 8. The b value increased

ith increasing sintering temperature. It was negligible small

or all ceramics with a grain size ≤0.21 �m, but substantiallyncreased 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)a on tim ◦

aA

ast(lTshdtscaTc

3c

fdattdaso

csmXz

tt

fdT1bftTpbsmCttdo

o2ttdtdasoXm1w

nd the surface monoclinic ZrO2 phase content profile as a function of degradati

bove 0.21 �m, the b parameter for ceramics with 0.25 wt.%l2O3 was significantly lower than for ceramics without Al2O3.The b value of CZ-3Y and CZ-3Y-0.25Al was either as low

s that of TZ-3Y and TZ-3Y-0.25Al at smaller ZrO2 grainize, i.e. at low sintering temperature, or significantly lowerhan 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 pro-onged sintering time. For Al2O3-free TZPs, the b value ofZ-3Y grade was much larger than that of CZ-3Y at grainizes above 0.25 �m. Below the critical grain size of 0.25 �m,owever, the b values were nearly independent on the pow-er preparation method. For 0.25 wt.% alumina doped TZPs,he b value of co-precipitated 3Y-TZPs (TZ-3Y-0.25Al) wastill larger than that of coated 3Y-TZPs (CZ-3Y-0.25Al) but theritical grain size shifted to 0.30 �m due to the grain growthcceleration effect of alumina addition. Therefore, coated 3Y-ZPs showed a higher surface degradation resistance thano-precipitated 3Y-TZPs.

.4.2. Depth of transformation in coated ando-precipitated 3Y-TZPs

The value of the kinetic b parameter from the surface trans-ormation curves revealed that coated 3Y-TZPs had a higheregradation resistance than co-precipitated 3Y-TZPs, especiallyt high sintering temperature. However, only the transforma-ion on the top surface (<10 �m) was measured by XRD andhe transformation propagation into the bulk material, whichetermines the final deterioration of Y-TZP material, was notssessed. Therefore, the transformation propagation was mea-ured by Raman spectroscopy to better compare the degradationf coated and co-precipitated 3Y-TZPs.

Fig. 9a shows the depth transformation profiles acquired onross-sectioned 40 h hydrothermally treated 3Y-TZP specimens

◦

intered for 2 h at 1550 C. For comparison, the surface transfor-ation profiles as a function of degradation time, as obtained byRD, are plotted in Fig. 9b. Fig. 9a shows that the monoclinic

irconia content decreased in a non-linear way and dropped from

ldi

e as measured by XRD (b) for 4 different 3Y-TZPs sintered for 2 h at 1550 C.

he saturation level to zero within a very short distance of lesshan 5 �m.

A faster surface t–m transformation resulted in a deeper trans-ormation propagation inside the 3Y-TZP specimens. TZ-3Yegraded faster than CZ-3Y since the transformed depth forZ-3Y (about 40 �m) was much higher than for CZ-3Y (about5 �m). Upon adding 0.25 wt.% alumina, the transformed zoneecame much thinner for both CZ-3Y and TZ-3Y, and the trans-ormation front was observed at about 2 �m and 7 �m belowhe surface for CZ-3Y-0.25Al and TZ-3Y-0.25Al, respectively.he degradation retarding effect of 0.25 wt.% alumina was moreronounced for CZ-3Y ceramics, which resulted in the most sta-le CZ-3Y-0.25Al grade. XRD results (Fig. 9b) showed that theurface of TZ-3Y, CY-3Y and TZ-3Y-0.25Al was saturated withonoclinic ZrO2 after 40 h degradation, whereas the surface ofZ-3Y-0.25Al was not yet saturated. No transformation satura-

ion plateau was observed by micro-Raman measurement insidehe CZ-3Y-0.25Al (Fig. 9a), since the monoclinic ZrO2 contentecreased from about 25 vol% at the surface to 0 within a depthf 3 �m.

Fig. 10 shows the corresponding cross-sectional SEM imagesf 40 h hydrothermally treated 3Y-TZP specimens sintered for

h at 1550 ◦C. Due to the extensive intergranular fracture andhe internal stresses from the volume expansion associated withhe degradation induced t–m transformation,36,37 a distinct bor-er between degraded and pristine material was visible. Theransformed grains in the degraded layer were easily pulled outuring sample polishing and the transformed zone appeareds a roughened layer, whereas the pristine bulk material wasmoothly polished and free of porosity. The thickness of thebserved degradation layer confirmed the results obtained byRD and Raman spectroscopy. After 40 h of hydrothermal treat-ent, the thickness of the degradation layer for CZ-3Y (about

2 �m) was much smaller than for TZ-3Y (about 36 �m), andas reduced upon adding 0.25 wt.% alumina. Only about 3 �m

ayer appeared to be transformed in TZ-3Y-0.25Al, and theegradation was only slightly visible at a depth below 1 �mn CZ-3Y-0.25Al.

2460 F. Zhang et al. / Journal of the European Ceramic Society 34 (2014) 2453–2463

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. Discussion

The strategy of increasing the stabilizer content to improvehe LTD resistance of TZP materials reduced their potential fortress-induced transformation. However, the excellent fractureoughness attributed to the stress-induced transformation tough-ning is one of the important reasons for Y-TZP ceramics to besed in prosthetic dentistry. Therefore, it is important to developydrothermally stable Y-TZP materials without compromisingn the fracture toughness.

In this work, 3 methods were observed to be able to improvehe LTD resistance of 3Y-TZPs, i.e. decreasing the sinteringemperature, adding 0.25 wt.% alumina, and incorporating thetabilizer by yttria coating of the starting ZrO2 powder. In the dis-ussion below, the influence of these 3 factors on the mechanicalroperties and LTD behaviour of 3Y-TZPs are highlighted.

.1. The influence of sintering temperature

The sintering temperature considerably influenced theechanical properties (Figs. 5 and 6) and hydrothermal degra-

ation (Figs. 7 and 8) of all investigated 3Y-TZP grades in similar way. The degradation susceptibility of all 3Y-TZPsignificantly increased with increasing sintering temperature,hich can be attributed to an increased tetragonal ZrO2 grain

ize (Figs. 1 and 8), a larger fraction of cubic zirconia and decreased average yttria stabilizer content in the remainingetragonal grains (Table 1). This result is in agreement with an

arlier report claiming that the presence of cubic grains had aarmful impact on the LTD resistance of Y-TZPs, and cubicrains were enriched with yttrium concomitantly resulting in aecreased yttrium content in the remaining tetragonal grains.38

h3s

ZPs sintered for 2 h at 1550 ◦C after 40 h hydrothermal degradation.

herefore, the increased LTD resistance of 3Y-TZPs obtainedy decreasing the sintering temperature is actually achieved byncreasing the average yttria stabilizer content in the tetrag-nal zirconia phase. Despite the increased LTD resistance, itan be expected that decreasing the sintering temperature wouldecrease the driving force for transformation toughening, com-romising on the fracture toughness of the 3Y-TZP. This wasndeed confirmed in Fig. 6, illustrating that the toughness ofll 3Y-TZPs decreased with decreasing sintering temperature orecreasing grain size at a sintering temperature above 1450 ◦C ort a ZrO2 grain size above 0.20 �m. Earlier studies have clearlyeported that decreasing the sintering temperature, i.e. decreas-ng the grain size of Y-TZP ceramics, significantly decreasedheir transformability and concomitant fracture toughness.18,39

Although it is shown in Fig. 6 that the fracture toughnessf 3Y-TZPs increased again upon further decreasing the sin-ering temperature from 1450 ◦C to 1350 ◦C, the sintering ofY-TZPs and especially alumina-free TZPs at a temperatureelow 1450 ◦C resulted in residual porosity that compromisedn the hardness. The increased fracture toughness could onlye attributed to an increased amount of closed pores. Fig. 5hows that 1450 ◦C was the optimum sintering temperature forll investigated Y-TZPs based on density and hardness. Theritical grain size between 0.21 and 0.25 �m above which theydrothermal degradation is dramatically enhanced (Fig. 8) waseached for all Y-TZPs when sintered for 2 and 4 h at 1450 ◦CFig. 7).

In summary, 3Y-TZPs and especially alumina-free TZPshould be sintered ≥1450 ◦C to obtain full densification and

igh hardness. From the LTD resistance point of view however,Y-TZPs have to be sintered ≤1450 ◦C. When 3Y-TZPs wereintered at 1450 ◦C, their fracture toughness reached a minimum

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F. Zhang et al. / Journal of the Europe

t an average ZrO2 grain size of 0.20–0.25 �m (Fig. 6). There-ore, fully dense 3Y-TZP ceramics with high LTD resistancend high toughness cannot be obtained by only adjusting theintering condition.

.2. The influence of 0.25 wt.% alumina addition

The addition of 0.25 wt.% Al2O3 accelerated the densifi-ation and increased the hardness of both CZ-3Y and TZ-3Yaterials when sintered at 1350–1400 ◦C (Fig. 5). In addition,l2O3 did not influence the fracture toughness of 3Y-TZPs

Fig. 6b). Therefore, upon adding 0.25 wt.% Al2O3, fully denseY-TZPs could be sintered at lower temperature to obtain aigher LTD resistance without compromising on the hardnessnd fracture toughness.

Moreover, the alumina addition itself decreased the sus-eptibility of both coated and co-precipitated 3Y-TZPs toydrothermal degradation, which was clearly shown by the facthat the kinetic parameter b value in the JMAK equation foroth CZ-3Y and TZ-3Y decreased upon adding 0.25 wt.% Al2O3Fig. 8). The transformation propagation inside CZ-3Y and TZ-Y was also significantly retarded (Figs. 9 and 10). Although it isommonly assumed that the addition of Al2O3 decreases the sus-eptibility of Y-TZPs to LTD due to a decreased grain size,40 theddition of 0.25 wt.% Al2O3 in this work was found to increasehe average ZrO2 grain size of both CZ-3Y and TZ-3Y ceramicsFig. 2). Therefore, the LTD retarding effect of Al2O3 addi-ion cannot be attributed to a grain size reduction.TEM resultsFigs. 3 and 4) clearly showed that Al3+ segregated at the ZrO2rain boundaries over a width of 5 nm in both CZ-3Y-0.25Alnd TZ-3Y-0.25Al. This is in agreement with literature reportslaiming that a small (<0.3 wt.%) amount of Al2O3 can dissolven the zirconia grains during sintering, resulting in the segrega-ion of Al3+ ions at the grain boundaries and Al3+ was reported toncrease the Y3+ segregation at the grain boundary in Y-TZP.41

he grain boundary region is crucial in the degradation of Y-ZPs because it is believed to be the nucleation site for the

–m transformation and the path for transformation to propagateuring degradation.42–44

In summary, a grain boundary with a small amount of seg-egated aluminium in solid solution enhances the hydrothermalegradation resistance of 3Y-TZP ceramics without compromis-ng on the fracture toughness.

.3. The influence of incorporating the stabilizer by Y2O3

oating the starting ZrO2 powder

Fig. 8 clearly shows that coated 3Y-TZPs had a higherurface degradation resistance than co-precipitated 3Y-TZPs,he transformation propagation profiles (Fig. 9) and imagesFig. 10) confirmed the higher degradation resistance of coatedY-TZPs than that of co-precipitated 3Y-TZPs, independent

n the alumina addition. Furthermore, the fracture toughnessf yttria-coated powder based Y-TZPs was higher than that ofhe yttria co-precipitated powder based ceramics, independentn the Al2O3 addition (Fig. 6). Therefore, coated 3Y-TZPs

tyba

eramic Society 34 (2014) 2453–2463 2461

ombine a higher LTD resistance and a higher fractureoughness compared to co-precipitated 3Y-TZPs.

The advantage of incorporating the stabilizer by yttria coat-ng of the starting ZrO2 powder is believed to be related to theeterogeneously distributed Y3+ at the grain boundary. TEMnvestigation (Figs. 4 and 5) showed that Y3+ was more enrichedt the edge of CZ-3Y-0.25Al grains, whereas the Y3+ distribu-ion was more homogeneous in the TZ-3Y-0.25Al grains. A cleareak in the Y/Zr ratio was observed at the CZ-3Y-0.25Al grainoundaries, which was not the case at the TZ-3Y-0.25Al grainoundaries. This was definitely due to the radically differentocations of the yttria stabilizer in the starting ZrO2 powder.

2O3 stabilizer was already inside the ZrO2 grain for yttriao-precipitated powders, so Y3+ segregated towards the grainoundary from the bulk of TZ-3Y-0.25Al, whereas for CZ-Y-0.25Al the coated Y2O3 layers had to dissolve and diffusento the ZrO2 grain from the surface during the sintering pro-ess, which might also contribute to the low yttria content inhe tetragonal zirconia phase of CZ-3Y and CZ-3Y-0.25Al sin-ered at 1350 ◦C. Due to the slow diffusion of Y in zirconia at350 ◦C,45,46 it is possible that only part of Y was dissolved andiffused into the zirconia phase and the rest of Y2O3 still locatedt the surface of zirconia.

Due to the higher amount of yttria located at the grain bound-ry in coated 3Y-TZPs, the grain boundary of yttria coatedZPs was more stable than that of yttria co-precipitated TZPs.

t therefore reduces the susceptibility of the coated powderased ceramic towards hydrothermal degradation. As explainedefore, the grain boundary stability is very important for the sta-ility of the complete TZP material because grain boundaries areelieved to be the starting point for transformation and also toct as the preferred path for water radicals to propagate into theaterial.42–44 A stable ZrO2 grain boundary with high Y3+, i.e.

high Y/Zr at the grain boundary, is beneficial for increasing theegradation resistance of TZPs, and it is therefore not necessaryo homogeneously increase the transformation resistance of theulk of the grains by increasing the overall yttria content. Fig. 9and b clearly shows that the surface transformation in the ini-ial stage of degradation and also the depth transformation aftereaching the saturation point were slower for coated 3Y-TZPsompared to co-precipitated 3Y-TZPs. The higher degradationesistance of CZ-3Y-0.25Al could was also enhanced by the finericrostructure compared to TZ-3Y-0.25Al (Fig. 1).The higher fracture toughness of coated 3Y-TZPs was

lso attributed to the non-homogeneously distributed yttria.n coated 3Y-TZPs, more yttria was located at the grainoundary, which in turn lead to a lower amount of yttria inhe core of the grains (Fig. 4b and Table 1). The grain coreas responsible for the higher toughness due to an enhanced

ransformation toughening contribution.24,29,47 Moreover, theow yttria content core is not located at the critical point ofhe degradation, i.e. grain boundary, and it is protected by the

ore stable grain boundary. Thus, the high transformability of

he core did not enhance the LTD susceptibility. In addition,ttria-coated starting powder based Y-TZP was reported toe able to have a core–shell grain structure,23 i.e. grains withn yttria-enriched tetragonal shell and a lower yttria content

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462 F. Zhang et al. / Journal of the Europe

ore, which is sensitive to transformation and responsible for higher toughness.48 The higher fracture toughness of coatedY-TZPs was also revealed from the results of a Rietveldefinement, indicating that the average yttria content in theetragonal zirconia phase of the coated Y-TZPs was lower thanor co-precipitated Y-TZPs (Table 1). A similar inhomogeneousttria and ceria distribution was reported for stabilizer coatedtarting powder based TZPs.25 The yttria content distribution inoated TZPs sintered from yttrium nitrate coated ZrO2 powdersas measured to vary from 2 to 8 mol%, which was broader andore inhomogeneous than for co-precipitated powder based

Y-TZP sintered under exactly the same conditions.24

Therefore, incorporating stabilizer by coating yttria in thetarting powder can be a new way to optimize the degradationesistance and also the fracture toughness of 3Y-TZP within theame material.

In summary, in order to co-optimize the LTD resistance andechanical properties, it was essential to add 0.25 wt.% Al2O3

nd incorporate the stabilizer by an alternative Y2O3 coatingoute, especially when the average ZrO2 grain size was above.25 �m.

. Conclusions

Increasing the sintering temperature significantly enhancedhe hydrothermal degradation of 3Y-TZP ceramics due to anncreased tetragonal ZrO2 grain size, a larger fraction of cubicirconia and a decreased average yttria stabilizer content inhe remaining tetragonal grains. Only limited degradation wasbserved at ZrO2 grain sizes below 0.20 and 0.25 �m for thel2O3-free and 0.25 wt.% Al2O3-doped ceramics, respectively.hen sintered for 4 h at 1450 ◦C or higher, corresponding to an

verage grain size above 0.25 �m, degradation of 3Y-TZPs wasonsiderably increased. However, decreasing the sintering tem-erature did not allow obtaining a degradation resistant 3Y-TZPith simultaneously optimized mechanical properties.Al2O3 addition and incorporating the stabilizer by yttria coat-

ng of the ZrO2 starting powder had a pronounced effect onetarding the degradation without compromising on the transfor-ation induced fracture toughness. This could be attributed to

he segregation of Al3+ and the heterogeneously distributed Y3+

t the grain boundary, respectively. 3Y-TZPs made from yttria-oated ZrO2 starting powder had a high Y/Zr ratio at the grainoundary and lower yttria content in the core of the grain, therebyombining a higher LTD resistance and higher fracture tough-ess compared to yttria co-precipitated ZrO2 starting powderased ceramics. The cumulative positive effect of the additionf 0.25 wt.% Al2O3 as well as the use of yttria-coated startingowder resulted in an enhanced resistance against low temper-ture degradation of the CZ-3Y-0.25Al ceramic with optimizedracture toughness.

cknowledgements

This work was performed within the framework of theesearch Fund of KU Leuven under project 0T/10/052 and

he Fund for Scientific Research Flanders (FWO) under grant

2

eramic Society 34 (2014) 2453–2463

.0431.10N. K. Vanmeensel thanks the Fund for Scientificesearch Flanders (FWO) for his postdoctoral fellowship.

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