mechanical properties and low-temperature aging of...

Mechanical properties and low-temperature aging oftetragonal zirconia polycrystals processed by hotisostatic pressing

J. Munoz-SaldanaTechnical University Hamburg-Harburg, Advanced Ceramics Group, D-21073, Hamburg, Germany

H. Balmori-Ramırez and D. Jaramillo-ViguerasDepartment of Metallurgical Engineering, Escuela Superior de Ingeniería Química e IndustriasExtractivas (ESIQIE), Instituto Politecnico Nacional, A.P. 75-872, Mexico City 07300, Mexico

T. IgaCeramic Science Department, National Industrial Research Institute of Nagoya, Nagoya 462, Japan

G.A. SchneiderTechnical University Hamburg-Harburg, Advanced Ceramics Group, D-21073, Hamburg, Germany

(Received 10 January 2003; accepted 11 July 2003)

The influence of grain size and density of yttria-tetragonal zirconia polycrystals(Y-TZPs) ceramics on mechanical properties and on low-temperature aging degradation(LTD) in air and in hot water was investigated. A TZP powder containing 3 mol%Y2O3 was consolidated by slip casting and densified by the sintering/hot isostaticpressing (HIP) method. Only the presintered samples that contained less than 0.15%open porosity reached near full density after HIP. The best conditions to reach fulldensity were found to be attained by presintering and HIP both at 1400 °C. At theseconditions, some of the best mechanical properties such as modulus of rupture andWeibull modulus reached 1397 ± 153 MPa and, 10.6, respectively. These values wereclearly higher than those obtained from sintered bodies and samples hot isostaticallypressed at 1600 °C. Aging degradation of 3Y-TZP materials can be avoided throughmicrostructural design. Fully dense materials with a critical grain size <0.36 �m didnot show any evidence of degradation after extreme aging conditions at pressurizedautoclaving in hot water at 100, 200, and 260 °C for 8 h. We propose a criterion topredict degradation in air as well as in hot water for the characterized materials basedon the microstructure and density control of the samples.

I. INTRODUCTIONTetragonal zirconia polycrystals (TZPs) have exten-

sively been studied during the past 15 years because oftheir excellent mechanical properties. They are based onmetastable tetragonal solid solutions of Y2O3 or CeO2 inZrO2.1 Their mechanical properties can be optimized bykeeping the amount of solute oxides as small as possibleand a grain size just below a critical size to avoid thespontaneous tetragonal (t) → monoclinic (m) transfor-mation during processing. Doing this, it is possible toactivate the transformation toughening mechanism.2 Inthe case of yttria-TZP (Y-TZP), optimum fracture tough-ness is obtained for materials with 2 mol% Y2O3 and agrain size around 0.3 �m. For materials having 3 mol%Y2O3, higher fracture strength is obtained, but the frac-ture toughness is lower. However, processing of 3 mol%Y-TZP is easier because its critical grain size increases toabout 1 �m.3

Despite Y-TZPs being one of the strongest structuralceramics, they experience a catastrophic degradationwhen exposed in air at temperatures between 100–400 °Cfor prolonged times, which is known as low-temperaturedegradation (LTD).4 This response is even faster inwater-rich environments.5 A spontaneous tetragonal (t)→ monoclinic (m) transformation at the surface underthese ambient conditions is responsible for the degrada-tion. The transformation produces severe surface crack-ing that drastically reduces the fracture strength of TZPbodies. Yoshimura6 and Lawson7 summarized the pos-sible mechanisms controlling this phenomenon, but anexact answer is not yet known. Sato et al.5 proposed thatthe mechanism responsible for the degradation is thechemisorption of water to form Zr(OH)4 at the surface,which results in the accumulation of strain energy andthus in a phase transformation tetragonal (t) → mono-clinic (m). A similar model6 was proposed, based on the

J. Mater. Res., Vol. 18, No. 10, Oct 2003 © 2003 Materials Research Society 2415

dissolution of Y2O3 in water or for the formation ofY(OH)3, which destabilizes the matrix, resulting in in-creasing residual stresses as the grain size increases.

LTD resistance has experimentally been shown to besuperior in materials with higher Y2O3 content, homo-geneously distributed stabilizer oxide, small grain size,and high density.6 For instance, Tsukuma et al.8 ob-served that 3 mol% Y-TZP with grain sizes greater than0.4 �m degraded only after 200 h at 400 °C. Masaki9

prepared TZP with different Y2O3 contents by conven-tional sintering in air, hot pressing, or hot isostatic press-ing (HIP). He found that TZP with more than 3 mol%Y2O3 and with grain sizes smaller than 0.4 �m processedby HIP did not degrade nor transform to monoclinic after2000 h in air at temperatures of 100–400 °C. His resultsalso suggest that there is a relationship between grain sizeand density of the ceramic to avoid degradation.

Alternative approaches have been reported to increasethe LTD resistance of these materials, where additionaldopants such as silica glass phase10 or CeO2

11 were in-troduced to the Y-TZP. However, the critical factor con-trolling the LTD resistance in these compositionsremained the control of microstructure. Parallel to theseresults, it has been proposed that HIP is a good method toimprove the LTD resistance of Y-TZP7,11 due to thecontrol of the density and fine grain size that can beproduced by this method.

HIP of Y-TZP has been the subject of a number ofstudies.8,9,12–14 Tsukuma et al.8 prepared TZP materialsby HIP with a modulus of rupture (MOR) value of1400 MPa. Hogg et al.13 showed that the properties andthe microstructure of these materials depend on the HIPtemperature. Tsukuma et al.8 and Masaki9 studied theirlow-temperature aging degradation and as mentioned be-fore their results were obtained by aging in air. Aging inwater of hot isostatic pressed 3Y-TZP was reported byChevalier et al.15 However, their observations weremainly directed to explain the mechanism of degradationand not to study the effect of the HIP conditions on themicrostructure and LTD resistance.

To characterize the influence of various sintering pa-rameters involving HIP on the control of microstructuresto reach both high mechanical properties and LTD resis-tance (in hot water) of 3 mol% Y-TZP ceramics was theaim of the current study. The novelty in our work isthe proposed simple criteria to predict degradation in airor in hot water for the characterized materials based onthe microstructure (grain size and density) of the sinteredsamples.

II. EXPERIMENTAL

Commercial zirconia powder (Tosoh TZ-3YS) con-taining 3.1 mol% Y2O3 and having an average particlesize of 0.3 �m was used. This powder was mixed with

distilled water in a ratio of 30 vol% of solids and ball-milled for 24 h in a plastic jar. An ammonium polyacry-late-based dispersant (0.02 ml/g of powder, SerunaD-350, Chukyo Yushi, Japan) was used to disperse thepowders and stabilize the suspension. The suspension pHvaried between 7 and 9, but this had a negligible influ-ence in the results because the powder is well dispersedunder these conditions.16,17 The powder was consoli-dated into plates of 70 mm × 40 mm × 5 mm by pressurecasting in plaster-of-Paris molds.

The plates were sintered for 1 h between 1300 and1500 °C at intervals of 50 °C. The samples sintered at1300, 1350, or 1400 °C were subsequently hot isostati-cally pressed in argon at 1400 or 1600 °C for 1 h at apressure of 150 MPa. The plates sintered at 1500 °C werekept as a reference for comparison. A summary of thesample preparation and identification is presented inTable I. The density and porosity of the sintered sampleswere analyzed by the Archimedes method. Phases wereidentified by x-ray diffraction (XRD). At least one sur-face of some of the sintered samples was polished withdiamond paste of different particle sizes (9, 6, 3, 1, and0.5 �m) and finally with a silica suspension (Mastermet,Würtz-Buehler, Germany). To reveal the microstructure,the polished samples were thermally etched at 1250 °Cfor 2 h. Micrographs (at least five of every sample) weretaken via scanning electron microscopy (SEM) and viaatomic force microscopy (AFM) (Nanoscope D3000,Digital Instruments, Santa Barbara, CA). The grain sizewas determined by the linear intercept method from theSEM and AFM micrographs.

The hardness (H ) and Young’s modulus (E) weremeasured through nanoindentation experiments con-ducted with the AFM combined with a Hysitron system(Triboscope, Hysitron, Minneapolis, MN). This setup al-lows continuous loading of the sample and the simulta-neous measurement of the applied force F and theindenter displacement h. H and E are determined fromthe partial unload portions of the load–depth curve. Tech-nical details on measurement of H and E are describedelsewhere.18 Measurements were performed with a cube-corner diamond tip as an indenter to a maximum load

TABLE I. Sinter/HIP conditions and sample identification.

Sinteringtemperature

(°C)HIP temperature

(°C)Identification of

samplesa

1500 �� S15001300 1400 S1300/HIP14001350 1400 S1350/HIP14001400 1400 S1400/HIP14001300 1600 S1300/HIP16001350 1600 S1350/HIP1600

aS, sintering or presintering at a temperature of; HIP, hot isostatic pressingat a given temperature.

J. Munoz-Saldana et al.: Mechanical properties and low-temperature aging of TZPs processed by hot isostatic pressing

J. Mater. Res., Vol. 18, No. 10, Oct 20032416

Fmax between 500–5000 �N, which corresponds to theindentation depths of about 80–310 nm. The tip-shapefunction of the indenter (the relationship between thedisplacement h and the projected contact area A�) wasdetermined by reference indentations in fused silica.

The MOR was measured by three-point bend testingfrom 25–30 bars (40 mm × 44 mm × 3 mm) cut from theplates and machined with diamond tools. The corners ofthe bars were chamfered at 45°. Test fixtures were tung-sten carbide with a point-load span of 25 mm. Niiharaet al. method19 for KIC determination was used. Five toseven indents at 10 or 20 kg were made on surfaces pol-ished with a diamond paste of 0.5-�m particle size. Simi-lar results were obtained at both loads. The samples forMOR and KIC were annealed at 900 °C for 1 h prior totesting.

LTD was evaluated at temperatures of 100, 200, 260,and 300 °C in air for times up to 2000 h. Samples of20 mm × 4 mm × 3 mm were used for these experiments.Density measurements and XRD patterns of everysample were taken at intervals of 100 h and then reintro-duced into the furnace until the aging time was com-pleted. Another set of samples was immersed for 8 h in a1-l Ti-autoclave with water at 100 °C/0.14 MPa, 200 °C/1.40 MPa, and 260 °C/4.14 MPa (temperature/pressure).Sixteen pieces of every sinter/HIP treatment were testedat each autoclave condition. At the end of the aging treat-ments, the samples were cut in the transverse direction andobserved in an optical microscope or via SEM or AFM.

III. RESULTS AND DISCUSSION

In the following sections, we report the results anddiscussion of microstructure and mechanical propertiesas well as the low-temperature degradation for the sin-tered samples. For microstructure, we emphasize density,porosity, phase identification, grain size, and size of de-fects observed on the surface of the sintered sampleswhile in the section on mechanical properties, H, E, andKIC are reported. In the third section, the low-temperaturedegradation of the characterized samples in air and inwater is discussed.

A. Microstructure of the samples after sinteringor HIP

1. Density and porosity

Starting with the as-cast samples, which had a relativegreen density of 47.55%, the sintering behavior is pre-sented in the form of a curve showing the variation of therelative density as a function of the sintering temperaturein Fig. 1. Also in this figure, the variation of the open andclosed porosity with sintering temperature is presented.The density increases continuously as the sintering tem-perature increases. At sintering temperatures higher than

1450 °C, the samples reached +99% of the theoreticaldensity (�t). Samples sintered at 1300 °C (S1300) pre-sented a wider range of scattering, which was reduced bysintering at higher temperatures. The closed porosity de-creases continuously from around 7% for samples S1300to less than 0.3% for samples S1500.

The variation of density and porosity of the sampleshot isostatically pressed at 1400 °C and 1600 °C as afunction of the presintering temperature is presented inTable II. Several of the samples presintered at 1300 °Cexploded during HIP at 1400 °C or 1600 °C, and thedensity of those that survived was less than 99% theo-retical density with a great degree of scattering. Most ofthe remaining porosity in these samples was closed, butmeasurable amounts of open porosity were still present.The density of samples presintered at 1350 °C or1400 °C increased to more than 99.5% after HIP. Neg-ligible amounts of open porosity remained in samplesS1350/HIP1400 and S1350/HIP1600. It is remarkablethat samples presintered at 1400 °C attained almost theo-retical density (>99.9%) after HIP at 1400 °C.

Further comments on the reduction in density or dam-age of the samples presintered at 1300 °C and hot iso-statically pressed at the given temperatures are asfollows. In samples having open porosity, the Ar gasused to pressurize during HIP penetrates the pores andimpedes densification. In the worst case, samples

FIG. 1. Variation of density and porosity of as-cast 3Y-TZP as afunction of sintering temperature (�, density; �, open porosity; �,closed porosity).

TABLE II. Density and porosity after HIP.

SampleDensity(%�t)

Porosity (%)

Open Closed

S1300/HIP1400 98.33 (±2.2) 0.06 (±0.05) 1.67 (±2.16)S1350/HIP1400 99.85 (±0.06) 0.06 (±0.07) 0.05 (±0.05)S1400/HIP1400 99.95 (±0.01) 0.01 (±0.01) 0.01 (±0.01)S1300/HIP1600 99.04 (±1.2) 0.06 (±0.04) 0.96 (±1.21)S1350/HIP1600 99.74 (±0.2) 0.02 (±0.01) 0.26 (±0.16)

�t: Relative density, taken the theoretical density as 6.09 g/cm3.


J. Mater. Res., Vol. 18, No. 10, Oct 2003 2417

“exploded” or the pores led to growing cracks duringHIP. Druschitz et al.14 also concluded that successfuldensification of TZP by the sinter/HIP method requirespresintered samples without open pores. Our conclusionfrom data in Fig. 1 and Table II is that, in order to attainfull density after the HIP treatment, it is necessary for theopen porosity of the presintered samples to be lessthan 0.15%.

2. Phase identification

From the XRD characterization, it was found that allthe samples sintered and hot isostatically pressed at dif-ferent conditions have a tetragonal structure. Addition-ally, in the hipped samples a small shoulder at the left ofthe (111)-tetragonal peak, at 2� � 29.2°, was identified.Masaki et al.20 assigned this shoulder to a rombohedralphase of ZrO2. It was not possible to identify the pres-ence of the cubic phase in the 2� range of 72° to 76°, sothat the amount of cubic phase could not be determined.However, its presence is predicted by the ZrO2–Y2O3

phase diagram at ambient pressure.21 Applying the levelrule, at 1600 °C a ZrO2–3 mol% Y2O3 system in equi-librium would be composed of approximately 29% ofcubic ZrO2 containing 6.3 mol% Y2O3 and 71% of atetragonal phase with 1.7 mol% Y2O3.

3. Grain size and surface defects

Representative microstructures taken by AFM ofthe samples sintered and hot isostatically pressed are pre-sented in Fig. 2. The variation of grain size for differentsinter/HIP treatments is presented in Table III. In thistable, two values of grain size are reported for everysample: one for the average grain size and another for thebiggest grain size observed in the AFM and SEM micro-graphs. The average grain size increases as the sinteringor HIP temperature increases. The samples hot isostati-cally pressed at 1400 °C retained a very fine and homo-geneous grain size. The samples sintered at 1500 °C orhot isostatically pressed at 1600 °C have a higher grainsize and show a bimodal tendency. For instance, the av-erage grain size of the samples hot isostatically pressed at1600 °C is around 0.55 �m, but it is possible to seegrains as big as 1 �m. These grains grew at the expenseof smaller ones during HIP. Some researchers12 haveproposed that the bigger grains are cubic, but it is nec-essary to employ transmission electron microscopy to char-acterize every grain individually. The average grain size ofall samples prepared in the framework of this investiga-tion attained the nanometric size (smaller than 600 nm)and is comparable with grain size results obtained fromsintered 3Y-TZP samples using colloidal processed zir-conia nanopowder as reported by Vasylkiv et al.22

An image analysis of the recorded surfaces revealed avariety of intergranular defects only from samples S1500

and S1300/HIP1400, which is consistent with the resultsof residual porosity reported in the previous section. Spe-cifically in the samples S1500 the pores were as large as10 �m, with some of the pore clusters up to 100 �m.Pore clusters of smaller size were also found in thesamples presintered at 1300 °C and post-hot isostaticallypressed at the given temperatures (S1300/HIP1400,S1300/HIP 1600). Defects like these were not foundin any of the other samples, which suggests differences inthe flaw populations between S1500 and S1300/HIP1400and the rest of the samples. Some Ca-rich inclusions(not shown) were occasionally observed, which defi-nitely came from the plaster molds. These observationscoincide with the results of the density measurementsand prove the efficiency of HIP to close the majorpore defects introduced during the consolidation step ofthe powder.

B. Mechanical properties

An effective elastic modulus of 201 ± 20 GPa (con-sidering a Poisson ratio of 0.25)12 and a hardness of13.7 ± 2.2 GPa were measured through the nanoindenta-tion experiments. These average values were computedover 50 experimental data points. The former wasused for the determination of the fracture toughness ofthe fabricated samples. MOR results are reported as afailure probability, and the degree of transformationtoughening is reported as a KIC value as a function of themicrostructure are discussed in the next sections.

1. Failure probability

The average MOR and the corresponding standard de-viation as a function of the processing parameters arelisted in Table IV. Additionally, two-parameter Weibulldistribution plots23 of all experimental data for the dif-ferent sinter/HIP conditions are presented in Fig. 3.These results are represented in two graphics separatingthe plots and showing a nearly linear survival probabilitydistribution [Fig. 3(a)] from those that show high “scat-tering” behavior [Fig. 3(b)].

The former were calculated according to the standardWeibull equation,23

PF = 1 − exp�−kV� �

�o�m� . (1)

Here, PF is the failure probability of a sample volumeunder stress (�) considering two parameters: the Wei-bull modulus (m) and a normalizing constant (�o). Adimensionless “load factor” (kV) is also considered inEq. (1), which describes the fraction of volume which iseffectively loaded at the maximum stress. For the three-point bending test, kV is given as:

kV =1

2�m + 1�2 . (2)



The failure probability of survival for the ith element wascalculated with the following equation

PF =�i − 0.5�

n, (3)

where n is the number of samples tested. For every proc-essing condition, 21 to 27 samples were tested.

From these plots showing a linear survival probabilitydistribution, the following observations can be drawn.First, the sinter/HIP combinations that show this behavior

are S1350/HIP1400, S1400/HIP1400, and S1500. Oneshould also note that the maximum stress, �, of eachspecimen was used in the regression analysis despite theknown fact that fracture in three-point bend specimensvirtually never originates exactly at the line of maximumtensile stress. One should also note that the samples pre-pared by S1350/HIP 1400 and S1400/HIP1400 show thehighest density and the finest and more homogeneousmicrostructure. However, samples S1500 also display alinear survival probability distribution as shown in

FIG. 2. Microstructure AFM of samples sintered at different sinter/HIP combinations: (a) S1300/HIP1400, (b) S1350/HIP1400, (c) S1400/HIP1400, (d) S1500, (e) 1300/HIP1600, and (f) S1350/HIP1600.



Fig. 3(a), even when the density and grain size combi-nation from these samples are clearly worse. This, in factis reflected in the difference of the calculated Weibullmodulus m of 5.7 when comparing with the modulus10.6 of the S1400/HIP1400 samples. One can concludethat a post-HIP treatment increases strongly the fracturestrength and reliability. As can be seen in Table II, thetotal porosity of these hot isostatically pressed sampleswas nearly zero. This difference is consistent with theaverage MOR of the samples presintered at 1300 or1350 °C and hot isostatically pressed at 1400 or 1600 °C,which is clearly higher than for samples S1500.

In the case of the plots showing the “scattering” be-havior of Fig. 3(b), we observe that the samples presin-tered at 1300 °C and hot isostatically pressed at the giventemperatures as well as the one presintered at 1350 °C andhot isostatically pressed at 1600 °C show a “scattered”plot, which is typical of samples having multiple flawdistributions.

For these samples, the standard Weibull distributionfunction was extended to a failure probability functionfor concurrent flaw distributions according to Johnson.24

This extension of the Weibull statistics considers an ad-ditional component, which has two independent means offailure, A and B, with associated probabilities of failurePFA and PFB (Eq. 4). Thus, the survival probability Ps �(1 − PF) of a component showing this behavior is then theproduct of the two survival probabilities.

PF � 1 − (1 − PFA) (1 − PFB) . (4)

This equation is valid regardless of the form of thedistribution functions describing PFA and PFB. Finally,considering that the two-parameter volume Weibull dis-tribution is valid for both flaw populations, Equations 1and 4 yield.

PF = 1 − exp�−kv� �

�oA�mA

− kv� �

�oB�mB� . (5)

This failure probability function contains four adjust-able parameters: mA, �oA, mB, and �oB. The completedata set of both Weibull parameters for samples at eachprocessing condition are shown in Fig. 3.

FIG. 3. Fracture statistics (Weibull plots) based on three-point-bendMOR as a function of the sinter/HIP combinations. Data are repre-sented in (a) standard and (b) concurrent flaw distribution two-parameter Weibull distributions.

TABLE III. Summary of microstructure and observations of low-temperature aging degradation.

Sample

Relativedensity(%�t)

Grain size(�m)

average/maximum

Agingin air

Aging inwater

S1500 99.5 0.37/0.57 Minimal CatastrophicS1300/HIP1400 94.2–100 0.32/0.42 No MinimalS1350/HIP1400 100 0.31/0.48 No MinimalS1400/HIP1400 100 0.32/0.51 No NoS1300/HIP1600 99.4–100 0.54/0.85 Minimal CatastrophicS1350/HIP1600 100 0.57/0.98 Minimal Catastrophic

TABLE IV. Mechanical properties of the prepared Y-TZP samples asa function of the processing parameters.

SampleModulus of

rupturea (MPa)KIC (10 kg)(MPa m1/2)

KIC (20 kg)(MPa m1/2)

S1500 816 (165) 4.9 5.2S1300/HIP1400 1113 (467) 5.1 5.5S1350/HIP1400 1346 (190) 5.0 5.4S1400/HIP1400 1397 (153) 5.1 5.4S1300/HIP1600 1076 (419) 5.0 5.2S1350/HIP1600 1280 (247) 4.9 5.2

aStandard deviation in parenthesis



Independent of the “scattering” behavior, it is interest-ing to observe that the data for every HIP temperatureconverge to a single line at the high-strength region of theWeibull plot, irrespective of the presintering tempera-ture, giving in all cases a Weibull modulus higher than 8.Moreover, this convergence occurs at higher MOR val-ues for the samples hot isostatically pressed at 1400 °C thanfor the samples hot isostatically pressed at 1600 °C.

Now focusing on the Weibull data, which points out asecond flaw population, the samples presintered at1300 °C and hot isostatically pressed at 1400 or 1600 °Cgive exactly the same modulus of about 1.5. This value iseven clearly smaller than the modulus of 2.5 from thesamples S1350/HIP1600. We explain the behavior of twoflaw distributions present in the samples presintered at“low” temperature as a consequence of the variability ofthe density and, more specifically, of the residual openporosity left after HIP (Figs. 1 and 2, Table III). It isworth remembering that some of the samples presinteredat 1300 °C exploded during HIP. From these results, onecan conclude that the presintering temperature has astrong influence on the variability of the MOR from thehot isostatically pressed samples. Furthermore, the sinter/HIP method is a good method to increase the fractureresistance of Y-TZP samples because the MOR of presin-tered samples with less than 0.15% open porosity wasincreased by HIP on the strict condition that nearly zeroopen porosity before HIP is satisfied.

2. Transformation toughening

The fracture toughness KIC of studied specimens wasnot affected by the processing parameters (Table IV).Values around 5 MPa m1/2 were measured for all thesamples. These values of toughness do not reveal a sig-nificant reinforcement through phase transformation dur-ing mechanical loading. A relatively straight crackintergranular path was observed. The amount of mono-clinic and tetragonal phases on the fracture surface ofsome of the MOR bars was analyzed by XRD. Less than5% monoclinic zirconia was always detected. Based onthese results, it can be concluded that transformationtoughening of these samples did not occur during theVickers indentation event. We explain this behaviorbased on the grain size of the samples as follows. Mostof the samples have grain sizes either much smaller(0.3–0.5 �m) than the critical grain size for thetetragonal → monoclinic transformation in 3 mol% Y-TZP3 or the bigger grains (>1 �m) were of cubic phase.This assumption needs to be further investigated toclarify the absence of transformation toughening.

Based on the MOR and KIC values for the samplesprepared in this work, we can summarize the observa-tions from the characterization of the mechanical prop-erties as follows. The MOR increased after HIP by

eliminating the biggest defects and refining the micro-structure with the consequence that no transformationtoughening could be activated under these conditions. Asmentioned before, the typical pore clusters up to 100 �mof the samples sintered at 1500 °C and the one hot iso-statically pressed at 1600 °C in comparison with sampleshot isostatically pressed at 1400 °C explains why theformer has lower MOR values. The mechanical proper-ties reported here are in close agreement with the resultsof Tsukuma et al.8

C. Low-temperature degradation

The low-temperature degradation of TZP in air andin hot water was followed by measuring the changeof density and the tetragonal to monoclinic phase trans-formation of the samples aged at different temperatures.The tetragonal to monoclinic phase transformation thatoccurs during LTD draws a change of density becausethe monoclinic phase has a smaller density than the te-tragonal phase. Such transformation can also involve sur-face microcracking, which destroys the ceramic bodyand contributes to the reduction of the sample’s bulkdensity.

1. Low temperature aging in air

The change of density with aging temperature for2000 h at 100, 200, 260, and 300 °C in air is presented inFig. 4(a) for the different sinter/HIP treatments. The den-sity of all the samples hot isostatically pressed at 1400 °Cremained constant after aging at different aging tempera-tures. This was the case even for the sample S1300/HIP1400, whose density was only 98 ± 2% of theoreticaldensity. On the other hand, a small decrease in the den-sity of samples S1500 and for samples presintered andhot isostatically pressed at 1600 °C after aging at 100 °Cand higher temperatures was observed. A density reduc-tion can be produced by surface microcracking or by theappearance of the monoclinic phase at the surface, be-cause it is less dense than the tetragonal phase. It was notpossible to detect by XRD the presence of monoclinicZrO2 at the surface of the tested bars that presented agingin air. Nevertheless, the aging products were observedusing scanning force microscopy. In the AFM micro-graphs of the aged samples (not shown), the t → m trans-formation was distinguished as a distortion of the typicaltopography of the machining surface lines of samplesS1500 and S1300/HIP1600. Such distortions were notobserved at the surface of the samples S1400/HIP1400,confirming that the samples processed at these condi-tions resisted aging in air at 260 °C for 2000 h. The sameobservations were made for the samples for which ag-ing was identified as “minimal” (Table III), even afteraging in hot water.



2. Low temperature aging in water

The change of density with aging temperature at 100,200, and 260 °C in the pressurized autoclave is presentedin Fig. 4(b) for the different sinter/HIP treatments. Forthese tests, samples were taken from different plates pre-pared at every sinter/HIP combination. Every point inthis graphic represents an average of values from 16different samples.

The highlight of these results is the tendency to mini-mal and in some cases zero density variation of thesamples hot isostatically pressed at 1400 °C, without anyinfluence of the presintering temperature after the ex-treme aggressive aging conditions in pressurized hot wa-ter. Aging in the autoclave at 200 °C and 260 °C had aminor effect in the density even of samples S1300/HIP1400 and S1350/HIP1400, which resisted aging inair. It is worth remarking that the density of samplesS1400/HIP1400 did not change after the aging treatmentsbed in either air or hot water.

Additionally to this, samples S1500, S1300/HIP1600,and S1350/HIP1600 display a different behavior than thesamples hot isostatically pressed at 1400 °C. The densityof these samples decreases rapidly as the aging tempera-ture increases, although a few samples survived aging at100 °C. The values for 260 °C are not reported for thesesamples because the tested bars broke into small frag-ments at this aging condition. With these results, one cansee that samples S1500, S1300/HIP1600, and S1350/HIP1600 did not resist aging in hot water.

The XRD patterns of the samples aged in hot water at100 °C did not show any difference with respect to thoseof the samples after sinter/HIP. Figures 5(a) and 5(b)show the XRD patterns obtained from the surface of thesamples after being aged at 200 °C and 260 °C. Most ofthe surface of samples S1500, S1300/HIP1600, andS1350/HIP1600 transformed to monoclinic ZrO2. TheXRD for the samples aged at 260 °C were taken from amixture of the fragments collected at the end of the ag-ing. Samples S1300/HIP1400, S1350/HIP1400, andS1400/HIP1400 only display the tetragonal pattern.

The surface of samples S1500, S1300/HIP1600, andS1350/HIP1600 was severely damaged after aging at200 °C. A monoclinic (m) surface layer was observedin a zoom-in of the cross section of the bars (Fig. 6),whose depth varies between 200 and 250 �m. Insample S1500, the layer presents a distinct “wavy” pat-tern [Figs. 6(a) and 6(b)]. In samples hot isostaticallypressed at 1600 °C, the phase transformation layerwas more planar (Figs. 6(c) and 6(d)]. The layer wasseverely cracked and separated from the tetragonal (t)core at several locations and more frequently at thecorners.

In SEM (Fig. 7), the monoclinic layer appeared veryporous. In sample S1500, the layer propagates throughpore clusters [Fig. 7(a)]. In these samples, many visiblecracks were generated. The phase transformation layergrows in a perpendicular direction to the cracks gener-ated during aging [Fig. 7(b)], leading to the “wave” ge-ometry. At some locations, untransformed tetragonalregions are surrounded by a monoclinic layer. In someof these samples, a monoclinic band formed far awayfrom the surface was observed [Fig. 7(d)]. Figure 7(c)shows the interface between the transformed layer andthe rest of the sample after the detachment of the trans-formed material.

As mentioned above, samples S1500, S1300/HIP1600,and S1350/HIP1600 aged at 260 °C broke into frag-ments. Diagonal cracks emerging from the cornerscaused the separation of the sample into four fragments(Fig. 8). The XRD pattern of the outer and inner surfaceslooked exactly as the one reported in Fig. 8(b) of thesamples sintered at temperatures higher than 1500 °C.These results confirm the presence of at least 85% ofmonoclinic phase in the fragments. The perfect geometry

FIG. 4. Variation of density after aging in (a) air for 2000 h and (b)hot water for 8 h at different temperatures and for different sinter/HIPcombinations.



of the broken fragments can be interpreted as a result ofa very fast aging process. Moreover, these results indi-cate that the tetragonal → monoclinic transformation in-duced by low-temperature aging in pressurized hot water

can quickly penetrate any Y-TZP sample to the core andproduce a catastrophic failure if critical microstructureand conditions for aging are fixed.

3. LTD resistance through microstructural design

A summary of the starting microstructure and agingobservations of the characterized samples is presented inTable III. In this table, the term “minimal” indicates thateven when there was no microscopic (optical and SEM)nor XRD evidence of monoclinic ZrO2 at the surface ofthe aged samples, the transformation was detected in aslight decrease in their bulk density and in a high-resolution AFM analysis of the surface topography (notshown). The sample S1400/HIP1400 did not show anysurface distortion under the characterization with AFMafter any of the aging treatments.

The aging behavior in air can be resumed as follows.Samples with relative densities greater than 94.5% andgrain sizes smaller than 0.6 �m can resist aging in airfor 2000 h.

The relationship between grain size, density, and LTDresistance in hot water (at 260 °C) of the tested sam-ples, irrespective of the processing route followed foreach sample, is presented in Fig. 9. Some results fromMasaki9 (aging in pressurized air at T < 800 °C) andChevalier et al.15 (aging in pressurized hot water up to0.2 MPa between 70 and 130 °C) are also shown. All thesamples with an average grain size greater than 0.36 �mand a heterogeneous distribution of grain sizes failedcatastrophically in hot water and presented some evi-dence of deterioration in air. A sharp boundary is ob-served at this value, which we propose to be consideredas a critical grain size for LTD resistance. From theseresults, it is clear that the grain size plays a crucial role inthe resistance to low-temperature degradation in pressur-ized hot water of Y2O3–TZP. The density (also dependenton the grain size) plays an important factor in the descrip-tion of the resist–degrade boundary. For instance, sampleswith the smallest density prepared in the current work(S1300/HIP1400) survived the severe attack of waterwhile the denser samples S1500, S1300/HIP1600, andS1350/HIP1600 failed catastrophically at the same agingconditions. This behavior is a consequence of the dif-ferent grain size ranges for the two sets of samples be-cause the former has a finer grain size. Masaki9 observedthat samples with a grain size of 0.2 �m and a low den-sity (around 5.45 g/cm3) failed after aging in air for2000 h. On the other hand, samples with high densityand small grain sizes resisted the mentioned aging con-ditions. Moreover, Masaki9 proposed a criterion to pre-dict aging behavior in the form of a resist–degradeboundary (dotted line in Fig. 9), which in general sug-gests equivalent effects of grain size and density on LTDresistance.

FIG. 5. XRD patterns of the surface of the samples after aging inwater at (a) 200 °C and (b) 260 °C.



We observe that instead of considering the influenceof density, it is probably more appropriate to consider therole that open porosity plays in the resistance to low-temperature degradation. The microstructural evidence(Fig. 7) shows that the tetragonal → monoclinic trans-formation occurs preferentially near pores and cracks.

These defects permit the access of water to the interior ofthe samples and are effective stress concentrators, wherelocal microcracking can start.

In materials with nearly theoretical density and biggrain sizes (samples hot isostatically pressed at 1600 °C),the aging mechanism is related with the critical grain size

FIG. 6. Optical micrographs of the transverse section of samples after aging in water at 200 °C: (a and b) sample S1500, (c) samples S1300/HIP1600, (d) samples S1350/HIP1600.

FIG. 7. SEM micrographs of the surface layer of samples S1500 after aging in water at 200 °C.



and the heterogeneity of the Y2O3 in the matrix. Thiscriteria based on the critical grain size for LTD resistanceconfirms the results obtained by Tsukuma et al.8 andMasaki9 for samples aged in static air. This simulta-neously gives credence to the theory that increasinginternal stresses affects the response of TZP to low-temperature degradation.25

As mentioned before, the other reason for failure ofmaterials with big grains is related to the heterogeneousdistribution of Y2O3 in the matrix. The samples with alarger grain size were sintered or hot isostatically pressedat temperatures above 1500 °C. During sintering or HIP,the Y2O3 partitions into the cubic and tetragonalphases.21 The equilibrium amount of the cubic phaseincreases at higher temperatures. Simultaneously, andprobably more important, the Y2O3 content of the te-tragonal phase is reduced to a value near 1.7 mol% at1600 °C. Whether or not the materials in the currentwork attained equilibrium is not known, but the forma-tion of some cubic phase in 3 mol% Y-TZP is well docu-mented.12,26 In previous studies about the influenceof Y2O3 content on the low-temperature degradation ofTZP, ceramics with <2 mol% Y2O3 always failed.8,9

Then, it is reasonable to assume that the tetragonal grainsformed at high temperatures are less resistant to agingand fail because they contain a smaller amount of stabi-lizer oxide. It is interesting to observe that the fragmentsof samples S1500, S1300/HIP1600, and S1350/HIP1600aged at 260 °C are composed of approximately 85% ofmonoclinic zirconia, which formed from the former te-tragonal grains. This amount is in close approximation tothe amount of 70% of tetragonal phase calculated byapplying the lever rule to the phase diagram of Scott.21

IV. CONCLUSIONS

Fully dense 3Y-TZP samples could be obtained byHIP at 1400 °C of slip casting specimens followed by asinter treatment at temperatures higher than 1400 °C,where open porosity was reduced to less than 0.15% foroptimal post-HIP treatment. Concerning the mechanicalproperties, the fracture resistance and the Weibull modu-lus varied dramatically with the sintering and HIP tem-peratures. The highest values of fracture resistance (of

FIG. 8. Photographs of typical broken fragments from samples thatfailed catastrophically at 260 °C in the autoclave.

FIG. 9. Density versus grain size relations for samples after aging inhot water at 260 °C. Results from Masaki9 (aging in air) and fromChevalier et al.15 are also presented.



the order 1600 MPa) with a Weibull modulus of 10.6were measured for fully dense materials (S1400/HIP1400). Moreover, the MOR decreases with increas-ing the HIP temperature. The LTD resistance of 3Y-TZPin air (100–130 °C) and in pressurized water (100–260 °C) depends critically on the grain size. Sampleswith densities higher than 94.5% of the theoretical den-sity and grain sizes smaller than 0.6 �m resisted aging inair for 2000 h. Only fully dense samples with a grain sizesmaller than 0.36 �m resist degradation in the more ag-gressive, pressurized hot-water environments.

ACKNOWLEDGMENTS

The authors thank Consejo Nacional de Ciencia y Tec-nología (CONACYT), and coordinación general de pos-grado e investigación-Instituto Politécnico Nacional(CGPI-IPN), for financial support of this project. Theassistance of M. Machida from National Industrial Re-search Institute of Nagoya (NIRIN) to perform the hotisostatic pressing (HIP) experiments, T. Scholz for thenanoindentation experiments, and of N. Claussen and M.Swain for helpful discussion is greatly appreciated.

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