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d e n t a l m a t e r i a l s 2 6 ( 2 0 1 0 ) 807–820

avai lab le at www.sc iencedi rec t .com

journa l homepage: www. int l .e lsev ierhea l th .com/ journa ls /dema

eview

ow temperature degradation -aging- of zirconia: A criticaleview of the relevant aspects in dentistry

anni Lughi, Valter Sergo ∗

ENMAT, Materials and Natural Resources Dept., University of Trieste, via Valerio 6A, Trieste I-34136, Italy

r t i c l e i n f o

rticle history:

eceived 26 June 2009

eceived in revised form

3 November 2009

ccepted 12 April 2010

eywords:

irconia

ental ceramics

etragonal-to-monoclinic

ransformation

ow temperature

a b s t r a c t

This review presents a critical survey of all experimental data about the low temperature

degradation of zirconia (often referred to as “aging”) due to the tetragonal-to-monoclinic

transformation, which have been collected at temperatures of interest for dental application

(room temperature to about 100 ◦C). It is shown that the main factors affecting the aging phe-

nomenon are (i) the stabilizer type and content, (ii) the residual stress and (iii) the grain size.

It is also shown that extrapolating the low temperature degradation rate from accelerated

aging tests can lead to unacceptable conclusions about the lifetime of the zirconia-based

components. Finally, based on the experimental evidence, a set of engineering guidelines

for the use of zirconia in restorative and prosthetic dentistry is proposed.

© 2010 Academy of Dental Materials. Published by Elsevier Ltd. All rights reserved.

egradation—aging

tabilizer

esidual stress

rain size

ngineering guidelines

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 808. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 808. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 810

ontents

1. Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .2. Terminology and background . . . . . . . . . . . . . . . . . . . . . . . . . . . .3. LTD and stabilizers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

4. LTD and stress . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .5. LTD and grain size. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .6. Experimental methods. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

∗ Corresponding author. Tel.: +39 040 5583702; fax: +39 040 572044.E-mail address: sergo@units.it (V. Sergo).

109-5641/$ – see front matter © 2010 Academy of Dental Materials. Puoi:10.1016/j.dental.2010.04.006

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 812

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 813

. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 815

blished by Elsevier Ltd. All rights reserved.

808 d e n t a l m a t e r i a l s 2 6 ( 2 0 1 0 ) 807–820

7. Limits of lifetime predictions at low temperature derived from accelerated aging tests . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 816ircon. . . . .. . . . .

8. Conclusions and engineering guidelines for the use of zAcknowledgement . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1. Introduction

Typing the words “zirconia OR ZrO2” on any scientific searchengine provides close to 25,000 publications from 1969 to2008. Clearly also review papers centered on this materialare numerous and, hence, it is necessary to justify the sup-posed need for a new one. We believe that the need stemsfrom two concomitant aspects: (i) the world of dentistry hasbegun to deal with zirconium oxide only in the recent yearsand some aspects of the long-term behavior of this materialsare not sufficiently known among dentists and researchers inthe field [1–3], and (ii) more specifically, in spite of the abovementioned body of literature, there are actually few contribu-tions reporting experimental data, and not extrapolations, onthe long-term degradation (aging) of zirconia at temperatureslow enough to be of interest for the dentists. For example, thecomprehensive review of Lawson [4] on the degradation of zir-conia due to the environment reports literally 3 data pointsbelow 100 ◦C. On the other hand, the case of the unusuallylarge premature failures of ceramic heads in hip joint pros-theses is widely known: between 2000 and 2002 a consistentnumber of ceramic balls made of yttria-stabilized zirconia,produced by Saint Gobain Desmarquest and marketed withthe name Prozyr®, failed prematurely because of a change inthe processing procedure which resulted in increased mono-clinic content (the best clinical description of the episode canbe found in [5] whereas a sound scientific interpretation of theproblem has been proposed only very recently, in 2009 [6]).

Chevalier, a worldwide expert in the field of aging of zirco-nia, recently presented an informed opinion on this problem[7]; in the concluding remarks, he adds that “The use of zir-conia for dental implants is quite young and in developmentphase. The issue of aging is still not discussed for these appli-cations.” Another review on the use of zirconia in dentistryconcludes also very conservatively, warning that long-termstudies are badly needed in the field [8]. A recent reviewfocused on the bacteriological aspects of zirconia for dentalapplications concludes that “A need for references concerningresistance to failure in long-term clinical trials is of paramountimportance for such systems.” [3]. The cited work of Denry andKelly [2] finishes stating that “It seems wise to keep in mindthat some forms of zirconia are susceptible to aging and thatprocessing conditions can play a critical role in the low tem-perature degradation of Zirconia”. In yet another very recentreview [6], in the few lines dedicated to zirconia in dental appli-cations, the authors conclude that aging “is expected to beno less of an issue for their manufacturers” [6]. This shouldinduce the dental community to be very cautious and in wontof a clear knowledge, based on experimental data, about theperformances of zirconia at low temperatures and long terms.

Thus, the present contribution will focus on the experi-mental evidence gathered in the temperature range of interestfor implanted dental materials, that is between 0 and 100 ◦Con the problem of long-term degradation, or aging, of zirconia.

ia as dental materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 817. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 818. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 818

The paper is organised as follows: firstly, a background on zir-conia will be presented, then a paragraph will be dedicated toeach of the three main factor affecting LTD: (i) stabilizer typeand content, (ii) stress and (iii) grain size. Then the experimen-tal methods used to study the LTD will be discussed followedby a section dedicated to a critical evaluation of the lifetimepredictions of zirconia implants. Finally, based on a com-prehensive evaluation of the experimental evidence, someengineering guidelines for the use of zirconia in restorativeand prosthetic dentistry are proposed in the last section.

2. Terminology and background

Zirconium oxide, ZrO2, is chemically an oxide and technolog-ically a ceramic material. It is basically insoluble in water. Itcan be dissolved in H2SO4 and in HF. In nature it is relativelyabundant (about 0.02% of the earth crust) [9]. Large depositsare present in Brazil as baddeleyte (monoclinic zirconia) andin Australia and India as zircon (ZrSiO4) sands. Pure zirconiapresents the phenomenon of allotropy, that is same chemi-cal composition but different atomic arrangement, among thefollowing crystallographic structures [10]:

Orthorombic↔monoclinic1170 ◦C←→ tetragonal

2370 ◦C←→ cubic2680 ◦C←→ liquid

The cubic structure is of the fluorite type, with oxygen ionsoccupying a simple cubic lattice and the zirconium ions occu-pying the center of half of the anionic cubic cells [9]. Examinedupon cooling, the transformation from cubic to tetragonal (c–t)and from tetragonal to monoclinic (t–m) is athermal and dif-fusionless (hence the term “martensitic” used to describe thistransformation, in analogy to what happens in steel). Further-more the t–m transformation occurs with a volume expansion(when unconstrained) of about 5 vol.% [9]. Till the late twen-ties, this enormous volume expansion prevented a large scaleuse of zirconia as refractory, because, upon cooling, it inducedsevere cracking of the bricks often leading to catastrophic fail-ure. Then Passerini [11] and Ruff et al. [12,13], independently,discovered that the tetragonal, or even the cubic form could beretained metastably at room temperatures by alloying zirco-nia with other cubic oxides, thus preventing the catastrophicfailure of pure zirconia. It becomes now clear why such oxideshave been termed “stabilizers”. To date there are reports ofbinary systems where zirconia is alloyed and stabilized withCao, MgO, Y2O3, CeO2, Er2O3, Eu2O3, Gd2O3, Sc2O3, La2O3 andYb2O3. By and large, the most studied stabilizers for bioma-terials applications are CaO [14], MgO [15], Y2O3 [16–19] andCeO2 [20,21], but only ZrO2–Y2O3 reached the actual statusof having a dedicated ISO standard for surgical application[22].

In 1976, in what is arguably the seminal paper concerning

zirconia, Hannink et al. [23], proposed that the t–m transfor-mation with its ensuing volume expansion, could be used toenhance the fracture toughness of zirconia-based materials.The fracture toughness is defined as the capability of a mate-

d e n t a l m a t e r i a l s 2 6

Fig. 1 – Schematic representing how the t–m transformationof ZrO2 increases fracture toughness. When a partcontaining metastable t-ZrO2 is subject to a remotemacroscopic tensile stress, the stress intensification due tothe presence of a crack tip is sufficient to transform somet-zirconia grains to the monoclinic form. Since thistransformation entails a volumetric expansion which isconstrained by the surrounding materials, the net result isa compressive stress acting on the surfaces of the crack,w

rG

K

wstwtnttmttb

hose propagation is thus hindered.

ial to stand a load in presence of a pre-existing flaw. Using theriffith approach [24] the fracture toughness, KIc, is defined as

Ic = Y�f

√c (1)

here Y is a geometrical adimensional parameter, �f is thetress at failure (sometimes referred to as resistance, or frac-ure resistance) and c is the initial length of the defect athich failure started. The schematic presented in Fig. 1 helps

o visualize how zirconia can improve the fracture tough-ess; the mechanism described in Fig. 1 is usually referredo as transformation toughening or, more completely, phaseransformation toughening (PTT) [25–27]. Depending on the

aterial surrounding the zirconia grains susceptible to the t–mransformation, different types of materials can be obtained;he most important ones are the following (the acronyms inrackets are the most widely used in the scientific literature):

(i) Partially stabilized zirconia (PSZ)—a matrix of cubic zirco-nia embedding transformable t-zirconia grains.

(ii) Zirconia toughened composites (ZTC)—a matrix with

high elastic modulus embedding transformable t-zirconiagrains. The most used matrix, with relevant biomedicalapplication in prostheses, is alumina and the material isknown as zirconia-toughened alumina (ZTA) [28].

( 2 0 1 0 ) 807–820 809

(iii) Tetragonal zirconia polycrystals (TZP)—the whole mate-rial is constituted by transformable t-zirconia grains.Differently from the two previous cases, which are two-phases materials, TZP is single phase.

Table 2 of Ref. [1] reports the commercially available mate-rials for dental use divided in the three aforementionedcategories.

The basic, indispensable requirement for phase PTT tooccur is that the t-ZrO2 grains be in thermodynamic metasta-bility at the conditions in which the crack is trying topropagate. It is exactly this thermodynamic metastabilitywhich lies at the base of the low temperature degradation,LTD, or aging, of zirconia: LTD (or aging) is defined as thespontaneous t–m transformation occurring over time at lowtemperatures, when the t–m transformation is not triggeredby the local stress produced at the tip of an advancing crack.Apart from other drawbacks of the t–m transformation occur-ring under these conditions (enhanced wear rates [29,30] withrelease of small zirconia grains in the surrounding environ-ment, roughening of the surface finish with both mechanicaland aesthetic worsening [31]) it is clear that, once alreadytransformed to the m-polymorph, zirconia cannot exhibit PTT,just like an used match cannot be lit again. Overall PPT and LTDare based upon the same phenomenon, the t–m transforma-tion of zirconia, and, to date, we cannot exploit the remarkablepossibility of the former without exposing ourselves to therisks of the latter.

The best way to understand the role of all factors involvedin the t–m transformation is to consider it a chemical reactiongoverned by the pertinent free energy change:

t-ZrO2 ⇒ m-ZrO2 �Gt−m (2)

The free energy term, �Gt−m, takes into account all contribu-tions relevant for the t–m transformation and can be split asfollows:

�Gt−m = �GCHEMt−m +�GSTRAIN

t−m +�GSURFt−m (3)

�GCHEMt−m is the traditional term obtained from the chemical

potentials: it takes into account the fact, imposed by nature,that at temperatures of interest for dental applications, theenergy of the monoclinic form is smaller than that of tetrag-onal. If the two other terms are vanishingly small, at thosetemperatures zirconia will be monoclinic. This term can bemade less negative, thus hindering the t–m transformation,by using the aforementioned stabilizers.

�GSTRAINt−m is the contribution to the total free energy change

due to strain and this term is also known as the strain energyterm. Stresses play a critical role in the t–m transformation, asdemonstrated by an important study of Schmauder and Schu-bert [32]: they showed that the t–m transformation does notoccur in stress-free grains of zirconia, even if these containan amount of stabilizer far below the stability of the phasediagram.

�GSURFt−m takes into account the size of the particles involved

in the t–m transformation and it is the contribution to the totalfree energy due to the presence of crystal surface of a givenareal extension. It is important to note that, as �GCHEM

t−m and

l s 2

810 d e n t a l m a t e r i a

�GSTRAINt−m scale with the cube of the crystal size while �GSURF

t−m

with the square of it, this last term becomes important onlywhen the size of the grains is small, tens of nanometers or so.

From all the above it is clear that

(i) stabilizer type and content modify the first term, appliedand residual stresses the second and grain size, GS, thethird term of Eq. (3); each of the following paragraphs willdeal with one of the three terms.

(ii) the relative contribution of the different terms ultimatelydetermines the possibility for the transformation to occurat an arbitrarily considered initial time.

(iii) If at the considered initial time the transformation is notoccurring, aging takes place if and only if, over the elapsedtime range, some kinetically driven phenomena occur,thus modifying one or more of the terms above.

The treatment above is thermodynamic, time-independentin nature, and it allows to determine the boundaries of thethermodynamic stability of t- and m-phase. Before discussingthe details of each term of Eq. (3), it is necessary to introducethe approach used to rationalize the t–m transformation intime-dependent, kinetic terms. Zirconia ceramics can exist ina metastable t-form even if the sum of conditions expressedby the three terms of Eq. (3) indicated that m-zirconia is thestable form; this is due to the fact that a triggering mech-

anism is needed for the t–m transformation to occur—justas water will need to be cooled well under 0 ◦C in order tofreeze. Like many other phase transitions, the transforma-tion of (meta)stabilized zirconia from t- to m-polymorph is

Fig. 2 – Y2O3–ZrO2 phase diagram. The shaded areaindicates the region of metastable tetragonal phase,separated from the monoclinic and the metastablepseudocubic region by the T0(t/m) and T0(c/t) dashed lines,respectively. According to the engineering guidelinespresented in the last paragraph, it is suggested that Y-TZPfor dental applications should have compositions andoperating conditions included within the dark-shadedregion.

6 ( 2 0 1 0 ) 807–820

governed the by nucleation of the new phase followed by itsgrowth. The starting point for a kinetic analysis can be con-veniently assumed to be the Mehl–Avrami–Johnson, MAJ [33]equation, often used for time-dependent, isothermal phasetransitions. The typical form is

fm = 1− exp[−(btn)] (4)

where fm is the monoclinic volume fraction, t is the time, b isan activation energy term of the usual form (b = b exp(−Q/RT))and n is an exponent related to the nucleation and growthconditions. Different kinetics of the t–m transformation willbe a consequence of different process conditions, which affectthe activation energy and the transformation driving force (i.e.the free energy difference between the metastable and sta-ble state), as well as of different operating conditions, whichdefine the trigger mechanisms involved and their efficiency.The shape of the MAJ equation is sigmoidal and, in the fieldof the t–m transformation and aging, it has had a consid-erable success since it allows an extrapolation of the timesrequired for aging at low temperatures from data collected athigh temperatures, which are much faster to acquire. As itwill be shown later, this approach has some severe drawbacks,to the point of being almost useless in predicting lifetime oft-zirconia components.

3. LTD and stabilizers

Though, in principle, all stabilizers could be used to obtainboth PSZ and TZP, practically alkaline earth oxides are usedfor commercially available PSZ and Y2O3 and CeO2 for TZP. Asall other stabilizers, they are characterized by a large solubil-ity in zirconia and by the ability to form fluorite-type phasesover large temperature and compositional ranges. By defini-tion PSZ is a two-phase material, with the tetragonal phaseprecipitating in the cubic matrix [34]: the precipitation is arather delicate processing step since it is influenced by severalparameters: initial GS of the powder, stabilizer type and con-tent, temperature and time for nucleation and growth of thetetragonal precipitates [35–37]. It is no surprise that one singleproduct of this type has been available for the dental market(Denzir-M®, produced by Dentronic AB) which is a MgO-PSZ[38,39].

Y2O3 is the most widely used stabilizer in high-tech appli-cation of zirconia; because of charge balance reasons, theintroduction of Y2O3 in the ZrO2 lattice gives origin to oxygenvacancies which are one of the mechanisms proposed for thestabilization effect [40]. The portion of phase diagram reportedin Fig. 2 indicates the most agreed-upon stability and metasta-bility regions. Typically 3 mol% and 8 mol% of Y2O3 in ZrO2

are used to stabilize down to room temperature the tetrago-nal (Y-TZP) and the cubic form, respectively. From an aestheticpoint of view Y-TZP is white and can be rather easily colored byaddition of traces of rare earth elements, for example addingsoluble nitrites or nitrates of these elements to the Y-TZP pow-

der; subsequent sintering gets rid of the nitrogen-containinganions leaving the cations in the lattice of Y-TZP, which ulti-mately confers coloration to the part. Since long it has beenknown that Y-TZP is prone to LTD in presence of water [41–44].

2 6 ( 2 0 1 0 ) 807–820 811

TtbZi1tii[rdimettmttt

pa

-

-

-

taa(foois

m[tnl[sd

Fig. 3 – Results of hydrothermal aging tests at 150 ◦C forY-TZP and Ce-TZP. Ce-TZP is much more resistant to LTDthan Y-TZP, unless Y2O3 content reaches 4 mol%, which liesin the non-transformable Y-TZP region shown in Fig. 2.

d e n t a l m a t e r i a l s

here is clear evidence that few hours at 100 ◦C in water leadso more than 40% of monoclinic phase in 2Y-TZP (the num-er in front of Y indicates the percentage in mol of Y2O3 inrO2) [45]. An yttria content of 5 mol%, thus with a large start-

ng cubic content, diminishes the monoclinic value down to5%. Sato and Shimada found that annealing provoked no t–mransformation at 62 ◦C, whereas an increase to 82 ◦C resultedn about 1 and 22 mol% of m-zirconia for annealing conductedn air and water, respectively (in this study the GS was 1.1 �m)43]. In another study [46] aging of 3Y-TZP conducted at 70 ◦Cesulted in more than 50% m-phase after 5000 h (less than 200ays), with an average GS less than 0.5 �m. At the same time,

t is widely demonstrated that extensive transformation to the-polymorph causes a deterioration of the mechanical prop-

rties, especially the bending strength. Interestingly enough,here is some evidence that, for aging performed above 300 ◦C,he mechanical properties are retained, whereas there is a

inimum for temperatures around 200 ◦C, indicating that athis temperature the thermodynamic contribution in favor ofhe m-phase conjures up with the kinetic acceleration to yieldhe largest m-content [47].

To date, there is no accepted mechanism to explain thehenomenon, but only informed speculations. The soundestre, in our opinion, the following.

Lange et al. [48] propose, based on some TEM observations,that water reacts with Y2O3 to form clusters rich in Y(OH)3;this leads to a depletion of the stabilizer in the surroundingzirconia grains which are then free to transform to mono-clinic.According to Yoshimura et al. [44], water vapor attacks theZr–O bond, breaking it and leading to a stress accumula-tion due to movement of –OH; this in turn generates latticedefects acting as nucleating agents for the subsequent t–mtransformation.Chevalier et al. [6] propose that O2− originating from thedissociation of water, and not OH−, is responsible for thefilling of oxygen vacancies which is believed to be one ofthe causes of destabilization and, hence, of LTD and for thelong diffusional path.

Irrespective of the mechanism, it is well established thathe t–m transformation (i) starts from the surface of the samplend then proceeds inward; (ii) it causes a surface uplift [49,50]nd microcracks [51] with ensuing aesthetic degradation [31],iii) it open the possibility for water to penetrate below the sur-ace, thus propagating the t–m transformation to the interiorf the sample [45], and, finally, (iv) it leads to the developmentf major cracks [51]. It is in response to these problems affect-

ng Y-TZP, that ceria, CeO2 has attracted a lot of interest astabilizer.

Ceria is needed in larger amount than yttria, in order toaintain the same regions of stability on the phase diagram

52]. A typical Ce-TZP composition, capable of t–m transforma-ion, contains 8 mol% of ceria. Above 12 mol% the system ison-transformable. Ce-TZP with 12 and 14 mol% showed neg-

igible degradation even after 500 h at 150 ◦C in water vapor53]; after 360 h in water at 80 ◦C Ce-TZP with 8 mol% ceriahowed less than 10% monoclinic on the surface [54]. Theifference with Y-TZP is striking (Fig. 3). Maybe even more

(Redrawn from Fig. 5 of Ref. [53] with permission fromAmerican Ceramic Society, Wiley-Blackwell).

important, the bending strength of these Ce-TZP samples isbasically unaffected by the aging cycles in water vapor [53]and even studies more focused on the dental applications ofCe-TZP confirm this finding [55]. Another interesting featureof the ceria–zirconia system, is that it shows the most pro-nounced pseudo-plastic behavior among ceramic materials: toa certain extent, Ce-TZP can bend before fracturing [56]. Actu-ally Ce-TZP can exhibit hysteresis loops on the stress-strainplane and even Shape Memory Effect, that is shape recoveryafter heating [57,58]. Chemically, Ce-TZP present the problemthat Ce4+ can be relatively easily reduced to Ce3+, which doesnot have the same stabilizing ability toward t-zirconia [59–61].Even sintering in air can lead to a reduction, and there isevidence that the chemical reduction can be triggered alsoat room temperature by the stress locally generated by thet–m transformation [62]. From an aesthetic point of view, Ce-TZP poses more problems than Y-TZP to the dentists: CeO2 isyellow and products based on commercially available Ce-TZPpowders range from light yellow to almost brownish. A furthercomplication is due to the fact that when Ce4+ is reduced toCe3+, due to the high concentration of oxygen vacancies Ce-TZP tends to became dark grey [61,62]. Since some foods havereducing capabilities (glucose and lactose, for example), thepossibility cannot be excluded that, on the long range, Ce4+

could be reduced in the oral cavity with consequent darken-ing of the zirconia-based part. We have not been able to finda single commercially available product for the dental marketmade of Ce-TZP, but there are studies being conducted on mix-tures between alumina, Al2O3, and Ce-TZP for applications indentistry [54,63–65]. It is also worth mentioning that there is an

important number of publications devoted to the study of zir-conia stabilized with mixtures of CeO2 and Y2O3, which showvery good mechanical properties and also a good resistance toLTD [66,67].

812 d e n t a l m a t e r i a l s 2

Fig. 4 – Schematic of a dental bridge subject to amastication load. In the point indicated by the arrow, thetensile stress can easily reach values of the order of severalhundreds MPa, often large enough to provoke locally the

t–m transformation.

4. LTD and stress

As mentioned in paragraph II, stress can affect the t–m trans-formation and the behavior of zirconia with respect to stresscan be easily rationalized in terms of Le Chatelier princi-ple: “when perturbed, a system in equilibrium tends to reactin order to minimize the perturbation”. When t-zirconia isput in tension, it minimizes this perturbation by expand-ing its dimensions and transforming to m-zirconia, becausethis expansion decreases the initially applied tensile stress(the same principle holds true for m-zirconia transformingto t-zirconia upon compression). Stresses can be applied orresidual and zirconia-based materials are subject to both indental applications. Applied stresses are immediate to visu-alize intuitively, but more difficult to compute precisely whenit comes to dental prostheses, given the complicated geome-tries of the part and of the loading conditions; a situation likethe one depicted in Fig. 4 can easily result in tensile stressesof the order of several hundreds of MPa. Given that stressesare defined as force/area, it is then possible to understandwhy manufacturers suggest lower bounds to cross-sectionalareas in prostheses: areas smaller than suggested increase thelocal stress possibly pushing it close to the limit of failure ofthe material. This aspect is made more complicated by theconcentration of stresses occurring at edges, discontinuities,crack tips, etc.

Residual stresses arise every time zirconia is fired at hightemperature and then cooled down to room temperature for

the practical use; fast cooling rates and/or coupling with othermaterials having a different coefficient of thermal expansion(CTE) lead to larger residual stresses. This is certainly the caseof zirconia in dentistry, where a firing step is always involved

6 ( 2 0 1 0 ) 807–820

in the preparation of the dental part. The general formulato calculate residual stresses arising between two dissimilarmaterials is the following:

�1 = E1

1− �1(˛1 − ˛2)(Tfir − TR) (5)

where �1 is the stress, E1 is the Young modulus, �1 is the Pois-son ratio, ˛ is the CTE, Tfir is the firing temperature and TR

is room temperature (in our case the body temperature); suf-fices 1 and 2 refer to the different materials. Eq. (5) predictsthat, when implanted in the oral cavity zirconia coupled withalumina will be subject to an approximate tensile stress of1 GPa: this value is sufficient to trigger the t–m transformationin several cases [61,68] (obviously depending on stabilizers,GS, etc.), and this t–m transformation can be a slow process[69], not necessarily a burst-like effect observed immediatelyupon reaching a certain tensile load [68]. Actually, the sameapproach to the determination of residual stresses is applica-ble if zirconia is coupled as a veneer/coating or as a laminateto another material, though in these latter cases the peculiarnature of stresses at and near the surface is not intuitive [70]and can lead to surface cracks [71].

We will now address the important issue connected withthe consequence of the residual stresses generated by aninitial t–m transformation on the prosecution of furthert–m transformation, that is the “autocatalitic effect”. Thisis explained with the help of Fig. 5. Basically, when somet-zirconia transforms to m-zirconia, we have a two-phasematerial with m-zirconia in compression and the remnantt-zirconia in tension, as predicted theoretically [72] anddemonstrated experimentally [73]; as a rough estimate, forevery 10 vol.% of t–m transformation, the tensile stress inthe remnant t-zirconia increases by 250 MPa. Because of thistensile stress, t-zirconia will be more prone to further transfor-mation to m-zirconia. This is the so-called autocatalitic effectwhich can be immediate at the moment of the crack propa-gation [68], but which can also be diffusion-controlled withmuch longer characteristic times [62,69]. In this latter casealumina/Ce-TZP samples were subject to bending and, at thebeginning of the experiments, no m-zirconia could be detectedon the samples. The bending induced an initial, burst-like,t–m transformation and, at some locations, the m-zirconia was20% of the total zirconia amount. This initial transformationinduced a tensile stress in the remnant t-zirconia of the orderof 700 MPa. After 100 days at room temperature, because ofthis tensile stress, further t–m transformation was observed,with the monoclinic content increasing from 20 to 35% [69,73].The exact nature of the change in the sample, which provokedthe slow t–m increase was not known, but it was clearly adiffusional-type process [69].

In some alumina–zirconia laminate composites [74] wehave found that tensile stresses as low as 400 MPa could lead toaging of 3Y-TZP after 5 years at room temperature and ambienthumidity. Actually the data are even more surprising: layers ofpure zirconia in macroscopic tension did not show aging after

4 years at room temperature and ambient humidity, whereaslayers of intimately mixed alumina/zirconia under macro-scopic tension aged considerably, up to 25 vol.%, and the samelayers under macroscopic compression still showed aging, up

d e n t a l m a t e r i a l s 2 6 ( 2 0 1 0 ) 807–820 813

Fig. 5 – Schematic representation of how a partialtransformation of zirconia grains to the m-polymorphincreases the tensile stress in the remnant t-zirconiagrains. (a) The whole microstructure of the sample,extending also around the thick broken line, is composedby t-ZrO2 with the top surface exposed to water (liquid orvapor). (b) One grain transform to monoclinic; on the topsurface, it is free to expand, provoking an uplift of thesurface itself. On the left, right and below, the newly formedm-grain is constrained by the surrounding untransformedmaterial. Since this grain is limited in its expansion, itimplies that it is in a large compression, −�1, whereas itspush against the surrounding materials gives origin to asmall tensile stress, +�1. (c) When more grains transform tomonoclinic, they are less constrained and the compressivestress diminishes to −�2, whereas they put in larger tensilestress, +�2, the remaining t-ZrO2. The length of the thickblack arrows in (b) and (c) is qualitatively proportional to themagnitude of the stress (arrows pointing outwards indicatetensile stress and vice versa). Microcracks originating fromthe t–m transformation facilitate the inward diffusion ofwater. The inset above right reports qualitatively the effectof the m-content on the residual stress in the remnant m-and t-ZrO2. Roughly, an increase of 10% of m-ZrO2

increases the tensile stress in t-ZrO2 by 250 MPa [73].

Fig. 6 – Histogram reporting the amount of m-ZrO2 presentin alumina/zirconia samples aged for 8 years at roomtemperature and ambient humidity. The bars in the figurerefer to scatter among samples subject to different coolingrates and with different grain size but same aluminacontent, and, hence, they do not represent experimentalerrors. The GS of ZrO2 in these samples is about 300 nmand the stabilizer content is 2.5 mol% Y2O3. These datawere collected in 2008; at the beginning of the experiments,

in 2000, no trace of m-ZrO2 could be detected on any of thesample [75].

to 5 vol.%. (For a correct definition of macro- and micro-scopicstresses see ref [70,74].)

Even more surprisingly, in a series of experiments con-ducted on alumina–zirconia samples (10:90 in volume) after7 years at room temperature and ambient humidity, we havefound evidence of spontaneous t–m: some preliminary dataare reported in Fig. 6 and they will be the subject of a dedi-cated paper. The samples belonged to a batch prepared in 2001to study the effect of GS and cooling rate on the stress accu-mulated at room temperature [75]. In 2001 no m-zirconia couldbe detected and the t-zirconia (2.5Y-TZP) was subject to tensilestresses in the range 50–200 MPa, with a GS of 300 nm. As canbe seen from Fig. 6, in 2008 the average m-content among allsamples was more than 3 vol.%, with peaks of 6 vol.%. To ourknowledge these are the only available data on LTD obtainedafter such long times and at room temperature. It is interest-ing to note that, according to the phase composition/stabilityphase diagram developed for Y-TZP with accelerated agingtests [76], 2.5Y-TZP with a GS of 300 nm should be well withinthe t-stability region: our data demonstrate that submicromet-ric GS and an yttria content of 2.5 mol% are not sufficient toprevent LTD over some years in presence of a relatively lowtensile stress.

5. LTD and grain size

Numerous researchers have reported that reducing the aver-age GS in zirconia-based ceramics has a beneficial effect on

814 d e n t a l m a t e r i a l s 2 6 ( 2 0 1 0 ) 807–820

Table 1 – Critical grain sizes for LTD-resistant 3Y-TZP ceramics. Rows 1 through 4 report experimental data whereretention of the tetragonal phase is also dependent on kinetic factors; Row 5 shows a theoretical prediction of the grainsize that guarantees thermodynamically stable tetragonal zirconia.

Ref. Reported critical grain size (�m) Notes

[42] 0.3 Aging 1000 h at 300 ◦C in air[92] 0.36 Aging 8 h at 100 to 260 ◦C in water

ging 1 ◦

ging 1edict

[93] 0.37 A[106] 0.52 AThis work ∼0.05 Pr

the stability of the tetragonal phase [42,48,77–79], and there-fore on LTD. Some confusion seems to arise in the literatureon this subject, much of it generated by the ambiguous defi-nition of the critical GS. Some researchers, who have focusedon a purely thermodynamic approach to the problem, iden-tify it as the maximum GS for which the tetragonal phase isthermodynamically stable, i.e. it will never transform. On theother hand, there is a number of contributions where the criti-cal GS is simply associated to the maximum size for which thetetragonal phase is observed under given conditions–therebyimplying (i) metastability of such phase rather than thermo-dynamic stability, and (ii) the need to account for kineticconsiderations. This ambiguity leads to reports of the criticalGS that spread over a broad range.

In the following, in order to frame the role of GS on thephase stability of zirconia and, in turn, on LTD, we will reviewthe thermodynamic approach – which sets the intrinsic limitsof this complex phenomenon – and then briefly mention themain factors that influence the thermodynamic equilibria andthe kinetics by which transformation occur. We will then pro-vide some of the key data (Table 1) selected among the mostsignificant when keeping in mind LTD, thus laying the basisfor a discussion and for drawing some engineering guidelinesin the following section.

Though here we shall discuss the specific role of GS as aseparate matter, as mentioned, chemical stabilizers, stress,and GS are actually interlinked and affect one another in acomplex way. For example, an increase in the content of sta-bilizer induces a reduction of the grain size [80], while a largergrain size leads to higher local stress [81].

As for a purely thermodynamic approach, one importantimplication of Eq. (3) is that, for small crystal sizes, the over-all free energy is heavily influenced by the term associated tothe crystal surface. By developing the �GCHEM

t−m and the �GSURFt−m

terms, one can see (details can be found for example in Refs.[77,78]) that for a given temperature, even below the transfor-mation temperature of an infinitely large solid (∼1170 ◦C, asseen above), �Gt–m can be negative if the radius is less thana critical value. The approach outlined here also leads to theprediction that the critical size depends on the temperatureT as 1/rc = A−BT where A and B are constants depending onthe surface energy and the transformation enthalpy [82]. Prac-tically, the lower the temperature, the smaller is the criticalsize.

The key implication for the purpose of the present review

is that there exists a critical radius rc such that for r < rc a zir-conia crystal can be thermodynamically stable in the t-formeven at low temperature. In other words, a reduction of thegrain size below a certain critical value has the potential of

00 h at 100 to 500 C in water00 h at 100 to 500 ◦C in air

ion based on the thermodynamic approach proposed in Refs. [80,82]

fully inhibiting LTD. This simple thermodynamic approachsuccessfully predicts the critical size at various temperaturesfor powders of both pure and stabilized zirconia. Although thevast majority of the available data refer to measurements con-ducted above 100 ◦C, a number of reports suggest that at roomtemperature for pure zirconia powders the critical size is inthe range of 5–10 nm [77,83,84]; similar values can be calcu-lated following a general framework proposed by Mayo andco-workers and well tested for temperatures above 300 ◦C [85].This approach also enables the calculation of the critical sizefor Y-TZP powders, which is about 90 nm for 1.5Y-TZP.

While powder data provide a useful framework for reason-ing on design guidelines (for example about the choice of thestarting GS), they alone are of little practical interest; it is moreimportant to focus on the relationship between GS and LTD inzirconia-based densified components, a problem with impor-tant consequences but much more complex to tackle. Thecomplexity arises from the fact that crystals in a solid (grains)are constrained by each other and are not free to expand orcontract as they are in a powder. This leads to stress genera-tion due to changes of volume/shape associated with thermalexpansion or phase transformations; moreover, in a solid body,stresses can be concentrated by local defects. The role of stressin LTD has already been described; here we will only remarkthat, generally, an applied hydrostatic pressure has the effectof increasing the critical size, while non-hydrostatic stressesintroduce a broadening of the transition between the m- andt-phase—i.e. the coexistence of the two phases over a rangeof temperatures [78]. All in all, the t-form can be stable inzirconia solids even if the grains are much larger than the crit-ical GS identified above for powders. By including the effectof strain energy on the formation of microcracks and twins[81,86], or the effect of the stabilizers [82,85], accurate predic-tions of the critical GS in densified solids as a function of thetemperature in Y-TZP ceramics have been obtained at tem-peratures above 300 ◦C. At room temperature, a critical GS of31 nm has been calculated by Garvie for pure zirconia ceram-ics [78], while Mayo and co-workers derive the critical GS atvarious temperatures and for various contents of stabilizer: forexample, from their work it is possible to extrapolate a value ofabout 150 nm for zirconia stabilized with 1.5 mol% Y2O3 (onceagain, in dense ceramics). Table 1 presents experimental data(thus affected by kinetic issues) about the critical GS of 3Y-TZPdense ceramics, together with the theoretical prediction of thethermodynamic critical GS.

As mentioned above, much of the confusion that seems toemerge from the literature about the critical GS arises fromthe fact that it is customary to indicate by this term the max-imum size for which a t-form can be observed, regardless for

2 6

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d e n t a l m a t e r i a l s

ts stable or metastable nature. This distinction, however, is ofrimary importance when the role of the GS on LTD is con-idered: while a thermodynamically stable tetragonal zirconiaill never transform, a metastable phase is prone to transfor-ation upon even slight variations of the system conditions

nd these variations are likely to occur if given enough time74,87–90].

From an engineering point of view, in the design processf a zirconia ceramic product it is of little importance if theetragonal phase is stable and will never degrade, or if it isetastable and will degrade in a time that is longer than the

ifetime of the product. In the latter case, however, it is criticalo know the transformation kinetics of the material.

One key factor affecting the kinetics in a number of wayss once again GS. First, local stress arises from the thermal

ismatch between grains, due to the slight anisotropy of theTE along the different crystallographic axes in zirconia, and

his stress is higher for larger grains [81]. GS also affects theinetics of transformation in metastable t-zirconia due to the

nfluence it has on the nucleation process. It has been men-ioned that the metastable t-form transforms to monoclinicia thermally activated nucleation and growth processes. Theinetics, therefore, depends on the number of nuclei thatan form. In a number of ways, including one described inhe following, grain boundaries can be considered as pref-rential nucleation sites, and this leads to the expectationhat the transformation becomes kinetically favored when

large number of grain boundaries is available—i.e. whenhe GS is small. In summary, although reducing the GS of

etastable t-zirconia ceramics has a beneficial effect as iteduces the thermodynamic driving force for transformation,rom a kinetic point of view a small GS might favor a fasterransformation through efficient nucleation of the m-phase athe grain boundaries. This effect was in fact experimentallybserved in experiments conducted on 2 and 3Y-TZP morehan 98% dense ceramics: after exposure in water at 80 and20 ◦C, samples with grain size of less than 0.6 �m showedore m-content than those with grain size ranging between

.6 and 0.8 �m [91]. This evidence is somewhat in contradic-ion with other experimental observations indicating that arain size of ∼0.35 �m or smaller hinders degradation [92,93].he guidelines presented in the last paragraph are based on

he fact that even a small amount of porosity contributes tonhancing the nucleation effect observed in Ref. [91]: conse-uently, the increase of transformation rate associated withhe reduction of grain size might be avoided in ceramics closeo the theoretical density [92,93].

. Experimental methods

n this section we briefly present the experimental methodshich can provide evidence for the occurrence of the aginghenomenon. A more extensive review can be found in Ref.

18], with the exception of the stress measurements, whichre not treated at all in that work.

Low temperature aging of zirconia is usually conducted inutoclave or steam chambers, where the pressure of waterapor, the temperature and the elapsed time are the controlledxperimental variables. In some cases the ceramic samples

( 2 0 1 0 ) 807–820 815

are simply immersed in water. The observed experimentalvariable is the m-zirconia content.

By definition, the difference between the t- and m-polymorphs is crystallographic and, hence, the sovereigntechnique to investigate crystallographic differences is X-raydiffraction. The most famous contribution on the use of XRDto distinguish between m- and t-zirconia is due to Garvie andNicholson [94]. In this contribution it was shown how the mon-oclinic content Xm, relative to the zirconia phase (and not tothe overall material), could be determine by the following for-mula:

Xm = I(1 1 1)m + I

(1 1 1)m

I(1 1 1)m + I

(1 1 1)m + I

(1 0 1)t

(6)

where I represents the intensity of the XRD peaks and thenumbers in brackets are the miller indices of the pertinentcrystallographic planes. Practically, a XRD scan between 25and 30◦ of 2theta covers the range necessary to observe allpeaks used in the formula above. Though well establishedand widely known, XRD is somewhat limited in the lateralresolution on most of the commonly available instruments;the typical spot is of the order of some millimeters and thesample should be flat. Another technique which can yield amuch higher lateral resolution is based on the Raman effect,that is the inelastic scattering of light due to the interactionwith vibrating matter [95]. This forms the basis of Raman spec-troscopy and, when implemented with an optical microscopein a back-scattered configuration, the same objective lens usedto observe the sample is then used to deliver a laser light onthe region of interest of the sample [96]. The light rejected bythe sample is dispersed/analyzed in the frequency domain,providing vibrational information (similar in nature to thoseprovided by Infra Red spectroscopy) on the sample under anal-ysis. Zirconia has a very strong Raman signal both in the t- andm-variants and the two vibrational spectra are very much dif-ferent. Using Raman spectroscopy, the amount of m-zirconiawith respect to the total amount of zirconia can be determinedaccording to [97]:

Xm = I(1 8 1)m + I

(1 9 2)m

I(1 8 1)m + I

(1 9 2)m + 0.97(I(1 4 8)

t + I(2 6 4)t )

(7)

where now the apices indicate the position of the pertinentRaman bands expressed in wavenumbers (cm−1). For the spe-cific issue of t–m transformation, Raman spectroscopy is afar more convenient method than XRD, as recognized also byChevalier et al., in their review on LTD [98]: beside the bet-ter lateral resolution (which can be as good as some �m),on typical spectroscopes a good Raman spectrum of t- andm-zirconia can be acquired in few seconds, as opposed tothe tens of minutes necessary for XRD. A further specificadvantage of spectroscopic methods is that they allow todetermine mechanical stresses experimentally, since mostRaman and luminescence bands change their position with

stress [99]: practically, it is possible to determine stresses fromthe shift of the spectroscopic bands, measured with respect toa reference stress-free state. This method is known as piezo-spectroscopy [99,100] and it has provided a unique insight

816 d e n t a l m a t e r i a l s 2

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6 ( 2 0 1 0 ) 807–820

about the interplay between t–m transformation, aging andstress [69,73,74,101]. For the materials of interest here (Y-TZP,Ce-TZP and alumina) piezo-spectroscopy is the only viablemethod to determine stresses experimentally and it has theadvantage that the same spectroscopic data provide infor-mation also on the phase composition of the material underinvestigation.

Thermal methods (differential thermal analysis, DTA; dif-ferential scanning calorimetry, DSC; thermal dilatometry, TD)can also provide evidence of the t–m transformation, but theyare limited to the transformation occurring during the experi-ment and, hence, they are practically useless for aging studies.

Methods sensitive to the surface topography (optical inter-ference, OI; scanning electron microscopy, SEM and atomicforce microscopy, AFM) can provide indirect evidence of thet–m transformation in the form of uplift from flat surfaces. Ifthe same sample is observed after time they can also provideindirect evidence of aging, but they cannot be used for bulkquantitative assessment.

7. Limits of lifetime predictions at lowtemperature derived from accelerated agingtests

Degradation rates at room or body temperature of Y-TZPceramics are currently not available, and accelerated testsat intermediate temperature (100–300 ◦C) are the only basisfor extrapolating an estimate of the transformation rate and,hence, of the product lifetime. This approach relies on theassumption that the transformation rate follows the sameArrhenius-like trend down to room/body temperature. Unfor-tunately, such extrapolation could lead to a significant error inestimating room/body temperature lifetimes.

Firstly, the uncertainty associated with the determinationof the activation energy is generally high. In Table 2 we reportthe estimates for the amount of m-phase formed after 10 years(a reasonable lifetime for dental applications), as well as thetime it takes for 25% of monoclinic to develop (i.e. the maxi-mum acceptable amount, based on current international ISOstandards [22]). The effect of the activation energy uncertaintyis clearly dramatic, so that no sound lifetime predictions canbe made on the reliability of these ceramics at low tempera-ture.

Secondly, the only available data for LTD at room temper-ature indicate faster transformation rates than those derivedby extrapolating data obtained at medium-low temperature.For example, Lughi and Clarke [88–90] showed that in zirco-nia thermal barrier coatings (3Y-TZP and 4.5Y-TZP) obtainedby electron-beam deposition, significant amounts of mono-clinic (up to 100%, depending on the processing) appear afteronly few months in laboratory air at room temperature. Also,Sergo [74] showed that up to 25% monoclinic can form at roomtemperature after only 6 years in 3Y-TZP zirconia/aluminalaminates. It should be noted, however, that coatings and lam-inates are not expected to behave exactly like Y-TZP ceramics:

the microstructure of the coatings is columnar and presentssome inter-columnar open porosity; also, in zirconia–aluminacomposites, zirconia is under significant residual stress. Thisresidual stress, in particular, is expected to be a factor in accel-

2 6

essEst(btfzsmin

oTaeHaraoebablteco

8f

Acuasf

d e n t a l m a t e r i a l s

rating the transformation rate at room/body temperature,ince lower temperatures imply higher residual stress: con-idering the example of alumina–zirconia composites, fromq. (5) one can estimate that in the zirconia phase the ten-ile residual stress is 40–50 MPa higher at room temperaturehan at temperatures commonly used for the accelerated testse.g. 134 ◦C as in the ISO Standards [22]). This difference mightecome significant in accelerating LTD. Analogous considera-ions hold even in pure Y-TZP, where a thermal stress arisesrom the anisotropy of the thermal expansion coefficient inirconia. Although in this latter case much lower residualtress is expected at room temperature with respect to inter-ediate temperature (<10 MPa), such temperature difference

s expected to enhance local stress concentration as well, withegative effects on LTD.

Lastly, one critical point in the current design approachf TZP is specific to the International ISO Standards for Y-ZP-based implants for surgery. The standard imposes thatmaximum of 25 wt.% of monoclinic is present after an accel-rated aging test conducted for 5 h at 134 ◦C and 2 bar [22].owever, this requirement does not provide any informationbout the actual lifetime at near-room temperature. The maineason, apart from the spread in predicted lifetimes associ-ted to the abovementioned uncertainty in the determinationf the activation energy, arises from the fact that the activationnergy for the t–m transformation in a given material cannote derived from such a “single datapoint” procedure (indeed,t least two accelerated tests at different temperatures woulde the minimum requirement for a rough estimate). Therefore,

ifetime prediction at room or body temperature according tohe ISO standards relies on the assumption that the activationnergy is approximately the same for all stabilized zirconiaeramics—a risky assumption, considering the wide spreadf values reported in the literature (Table 2).

. Conclusions and engineering guidelinesor the use of zirconia as dental materials

fter presenting the most relevant experimental data, we pro-eed now to formulate some possible guidelines for a practical

se of zirconia as dental materials. First of all the issue must beddressed whether the aging of t-zirconia to the m-polymorphhould be avoided or can be lived with, also in light of theact that there is some evidence that a limited and superfi-

Table 3 – Engineering guidelines for use of zirconia-based mate

Property

Density % ofStabilizer content: Y2O3 mol%Grain size �mResidual stress MPa

Monoclinic content FractInitialAfter aging test (134 ◦C, 2 bar H2O, 1 h)

Strength in 4 point bending MPaInitialAfter aging test

( 2 0 1 0 ) 807–820 817

cial t–m transformation can causes an increase of the strengthof the material [102]. The reason is that the t–m transforma-tion induces a local overall compressive stress on the surface,which tends to close a potential advancing crack (clearly thesuperficial compression must be accompanied by an equili-brating sub-surface tension, otherwise the part would not bein mechanical equilibrium). The data presented, however, arelimited to strength values collected right after sintering andthe possible effect of aging has not been examined. In prin-ciple, a surface engineering approach could be imaginable,whereby a tight control of the superficial transformation inthe samples is adopted, but this is totally impracticable in astandard dentistry laboratory.

Some studies indicate that alumina/zirconia mixtures areless prone to LTD than pure zirconia [103,104], but the datapresented here induce us to suggest that at or near room tem-perature and everything else been equal (GS of t-zirconia andstabilizer type and content), ZTA is probably more sensitive toageing than pure t-zirconia because of the microscopic tensilestress to which t-zirconia is subject when mixed to alumina[74].

In order to establish viable engineering guidelines, wefirstly need to assess the design constraints as they can begathered from the experimental evidence presented in theprevious sections and substantially summarised in Table 2:

- LTD occurs in pure zirconia samples stabilized with as muchas 2.5 mol% Y2O3;

- LTD occurs in alumina/zirconia samples where zirconia isstabilized with 2.5 mol% Y2O3 and 8.5 mol% CeO2;

- tensile stresses of the order of some hundreds MPa are suf-ficient to promote long-term LTD;

- intimately mixed alumina/zirconia (where zirconia is inmicroscopic tension) samples are more prone to LTD thanlayers of pure zirconia subject to macroscopic tensilestresses;

- zirconia samples with GS of the order of 300 nm can exhibitLTD over some years;

- All experimental evidence points to the fact that the t–mtransformation occurs initially on the samples surface, andthence penetrate underneath;

- water vapor possibly causes and certainly accelerates LTD.

The guidelines that we are about to propose are basedon the previous conclusions and on the following hypoth-

rials in restorative and prosthetic dentistry.

Unit Requirement

Theoretical density >99%3.5 < x < 8GS < 0.3� < 300 (tensile)

ion wrt to total zirconiaNegligibleNegligible

600600

l s 2

r

818 d e n t a l m a t e r i a

esis: the key properties of zirconia (high chemical stability,biocompatibility, low Young’s modulus, large natural avail-ability, consolidated industrial know-how) justify its use evenin absence of phase transformation toughening, because toimplement PPT on zirconia-based dental components implies,to date, exposing them to LTD on a time scale not acceptable inrestorative and prosthetic dentistry. Simply stated, t-zirconiashould be prepared and implanted in a way to make it non-transformable. This will reduce the strength by some amount,but will also make the material completely unaffected by LTD[47].

Based on all the above, the engineering guidelines for theuse of zirconia as dental materials are presented following thescheme of ISO standards [22] in Table 3.

Acknowledgement

FONDO TRIESTE is acknowledged for partial funding.

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