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1989 68: 1075J DENT RESK.J. Anusavice and R.B. Lee

Effect of Firing Temperature and Water Exposure on Crack Propagation in Unglazed Porcelain  

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Effect of Firing Temperature and Water Exposureon Crack Propagation in Unglazed Porcelain

K.J. ANUSAVICE and R.B. LEE

Departm-nent of Dental Bionmaterials, College o(f Dentistry, University of Floiida, Gainesville, Florida 32610-0446

Static ftitigue of dental ceramics results from the interaction of resid-ialf tensile stress cand an aqueous environment. 'This phenomenllon isa potential cause of delayedl crack formation and propagation in ce-raimic or metal-ceranmic restorations. For dental ceramics, the influ-ence of microstructural dlejects such as porosity or fissures caused byincomplete sintering is riot kloown-. 7hle objectives of this study wereto characterize the influence ofporosity (produced by Lindeifirilng) onthe crack propagation resistance of two jeldspathic porcelains (111(d todetermine whether lower stress corrosion susceptibility or higher frac-tur-e toughness accountss for the superior thermal shock resistance of0oe of these ceramics.

Wae111underfired blas oeach porcelain, 25 mmn X 4 mnm X 4 mmt,by as nutch acs 84°C below their recommended firing temperatures.Atlter polishing the specimens through 0. 05 pm alumina, we inducedcracks in their sum faces wit/i a Vickers mnicrohardness indentert Semi-circular cracks, which were produced under an applied indenter loadof 19.6 N, grewt wit/i tine during storage in distilled watem- at 37C.Underfiring of bat/ ceramics caused a slight increase ini Jracturetoughness and a relatively .small change in pore volume fraction untilwe underfired the ceramics at 30'C or more. The crack propagationdata indicate that the higher them-al shock resistance of one of theceramics-as measured previously by a water-quench technique-maybe due to its greater resistance to sti-ess cotr-asion at the initial stageofcrack proptigation.

.1 Dent Res 68(6):1075-1081, June, 1989

Introduction.

The resistance of brittle materials to crack propagation is pri-marily influenced by surface flaws and microstructural fea-tures. Dental ceramics are also susceptible to crack growthenhanced by moisture. The presence of subcritical tensile stressesin porcelain due to thermal contraction incompatibility maycause delayed failure of metal-ceramic restorations. Such re-sidual tensile stresses, even if relatively low in magnitude, maycause existing flaws to grow with time in the presence of mois-ture. This reduction in strength, which is caused by a chemicalreaction at crack tips, is known as static fatigue. Jones et al.(1972) have shown that long-term loading of dental porcelainsat subcritical stress levels will produce failure over periods ofhours or months. Sherrill and O'Brien (1974) have shown thataluminous and feldspathic porcelains, when tested in water,exhibit a decrease in flexure strength of approximately 27%.

The presence of porosity in dental porcelain is generallybelieved to reduce its strength. However, Jones and Wilson(1975) have shown that there was no significant differencebetween the transverse strength of air-fired porcelains with apore volume of 5.6% and vacuum-fired porcelains with a porevolume of 0.56%. These authors suggested that the predomi-nantly spherical voids found in feldspathic porcelains did notreduce strength as markedly as did irregularly shaped voids.

Irregular, nonspherical voids might facilitate crack initiation

when subjected to transient incompatibility stress below thethreshold stress that produces failure in the presence of spher-ical voids. Such irregular voids may develop when porcelainis underfired.

The interactive effects of porosity, residual stress, and waterexposure on crack propagation in dental feldspathic porcelainshave not been previously investigated. Of primary concern isthe effect of residual tensile stress in porcelain veneers of metal-ceramic restorations caused by thermal contraction mismatchon crack initiation and propagation. Low-fusing porcelains formetal-ceramic restorations lack energy-dissipative microstruc-tural features that could significantly enhance the toughness ofthese materials. The glass matrix usually contains crystallineleucite which has the potential to inhibit crack initiation andpropagation. However, Morena et al. (1984) reported a higherfracture toughness value (1.27 MPa-m"'2) for a high-expansionexperimental frit than for a commercial porcelain (0.88MPa-m"2). The increased toughness of the experimental ma-

terials was attributed to the obstruction of cracks by leuciteparticles. The cracks were induced by a Vickers microhardnessindenter which was used for determination of fracture tough-ness.

Measurements of fracture toughness may provide insightinto analyses of thermal shock resistance of dental porcelain.In a previous study, Anusavice et al. (1981) measured themean thermal shock resistance (reported as AT, the differencebetween the annealing temperature and the quenching temper-ature at which cracking occurred) of metal-opaque porcelain-body porcelain crowns. For one porcelain (Vita body porce-lain, shade 540, Vident Corporation, Baldwin Park, CA) withfour types of alloys (Ni-Cr, Au-Pd, Au-Pd-Ag, and Pd-Ag),the AT value (1700C) was much greater than that (1450C) ofsimilar crowns fabricated with another commercial porcelain(Ceramco gingival porcelain, shade Al, Ceramco, Inc., EastWindsor, NJ), even though the transient tensile stress reportedby DeHoff et al. (1983) at the surface of body porcelain ineach case was comparable. Thus, it is possible that this dif-ference could be explained on the basis of a potentially greaterfracture toughness of Vita (V) body porcelain.

Recent investigations of the fracture behavior of brittle ce-ramics have focused mainly on the use of techniques employ-ing controlled surface flaws. These techniques-which includedouble torsion, double cantilever, and notched designs (Frei-man, 1980)-are not well-suited for materials that are availablein small quantities. For relative fracture toughness comparisonsof dental ceramics, the indentation method is useful (Marshall,1983; Morena et al., 1984). Another potential benefit of thismethod is the ability to produce residual stresses (within thecontrolled surface flaw area) which can serve as the drivingforce for crack propagation in environmentally sensitive ce-

ramics.The objectives of this study were to investigate the crack

propagation resistance of two body porcelains as a function ofincomplete sintering, and to characterize their susceptibility tostatic fatigue by means of the indentation technique. The hy-potheses tested were that incomplete sintering enhances crack

1075

Received for publication October 23, 1987Accepted for publication February 16, 1989This study was supported by NIDR Grant DE06672 and RCDA

DEOO 125.

1076 ANUSA VICE & LEE

Fig. 1-Schematic illustration of indentation crack pattern. The initialcrack size (c0) was determined immediately after removal of the micro-hardness indenter.

growth, and that the higher thermal shock resistance of Vitabody porcelain is a result of its greater fracture toughness.

Materials and methods.Two commercial body porcelain products, designated as C

and V, were selected for study because of their known thermalshock behavior. We prepared three bars of each porcelain, 25mm x 5.5 mm x 5.5 mm, for each set of experimentalconditions. To minimize the effect of residual surface stressresulting from the formation of a glaze layer and from rapidcooling, we polished each bar through 0.05 pLm alumina to thedimensions of 25 mm x 4 mm x 4 mm. We prepared 24specimens of each porcelain by firing three specimens each (ofporcelain C and porcelain V) in vacuum at 971'C and 920'C,respectively, and three specimens each in vacuum at a tem-perature 14'C above, and at temperatures 14AC, 280C, 420C,560C, 70'C, and 840C below these recommended tempera-tures. After we polished each bar through 0.05 awm aluminaabrasive, we indented them at six widely separated locations

(3 mm apart) with a load of 19.6 N in a microhardness tester(Tukon Microhardness Tester, Model MO, Page Wilson Cor-poration, Bridgeport, CT). Thus, 18 measurements were madefor each condition:

Typical crack patterns, which are shown in Fig. 1, wereproduced by the indentation process. We measured the initialcrack length (c0) within two minutes after indentation to min-imize errors associated with moisture-enhanced crack growth.We stored these specimens in distilled water at 370C. We madesubsequent measurements of crack growth at intervals of onemonth over a period of six months.We prepared another series of 48 specimens for each por-

celain by firing six specimens at each of the eight temperatures.We polished the specimens through 0.05 pm alumina, an-nealed them for 100 h (3520C for porcelain V and 3860C forporcelain C), and indented them with a load of 19.6 N. Weindented one group of 24 specimens in air and stored it inwater at 370C. We indented another group with a drop ofparaffin oil (to minimize moisture-enhanced crack growth) andthen stored it at 370C in oil. We prepared an additional 15specimens for each porcelain by firing at the recommendedmaturing temperature. We annealed the bars to reduce surfacestresses caused by rapid cooling or grinding. The annealingtemperatures were two-thirds of the glass transition tempera-ture values reported for these porcelains. We avoided higherannealing temperatures to minimize the risk of microstructuralchanges associated with the leucite crystalline phase.To determine the range of indentation loads in which frac-

ture toughness equations are valid, we indented a group ofthree specimens for each porcelain fired at the recommendedtemperature at 18 locations with loads (P) of 3.9 N, 7.9 N,11.8 N, 15.7 N. and 19.6 N. For fracture toughness relationsto be valid, values of P/c 3/2 should be invariant with load(Marshall, 1983). As for all other specimens, indenter loadswere maintained for 18 sec prior to removal. These specimenswere also stored in test tubes containing water at 370C.As a reference standard, we used soda-lime-silica glass plates

(microscope slides) measuring 12 mm x 12 mm x 0.9 mmthick. We embedded the specimens in metallographic mount-ing resin after annealing them at 3860C for 100 h. We made30 indentations in each of three specimens at a load of 19.6N. In one specimen, we made the indentations with a drop ofparaffin oil to minimize slow crack growth resulting from con-tact of the cracks with atmospheric moisture.We made all measurements of crack lengths using the mi-

croscope of the hardness testing machine. We verified cracksize measurements on specimens with stabilized cracks by meansof a measuring microscope (Unitron, Model BTMD, UnitronInstruments, Inc., Plainview, NY).We calculated fracture toughness (K1c) using the relations

developed by Lawn and Fuller (1975) for the residual stressintensity factor,

Kr = XrP/c3/2 (1)

where c is the crack length, and Xr is defined (Lawn et al.,1980) by

Xr - § (E/H) (2)

for radial cracks produced by Vickers indentations. If one as-sumes that the crack system is in equilibrium during and afterthe indentation procedure, radial cracks remain stable (c = c,,and K, = K1c). The applied load (P), modulus of elasticity(E), hardness (H), and initial crack size (c0) are required forthe calculation of Kc. The calibration constant (§) for Vickersindentations is generally taken as 0.016 [according to the cal-culations of Anstis et al. (1981)] for a variety of reference

J Dent Res June 1989

-.*-a No c0

CRACK PROPAGATION IN PORCELAIN

ceramics. We calculated hardness values for each indentationsite using the expression,

H = 0.47P/a2 (3)For soda-lime-silica glass and all porcelain specimens, we

used a value of 73.4 GPa (Wiederhorn, 1969) for the elasticmodulus (E). Although porosity is known to reduce the elasticmodulus of these porcelains, no correction was made, sincedata on the variation of E with porosity were not available.We determined the volume fraction of porosity in each speci-men on the indented surfaces by means of quantitative mi-croscopy. We superimposed a 10 x 10 grid on magnifiedimages of the porcelain surfaces, and performed point-countingto determine the volume fraction of porosity (Vv), in accord-ance with procedures described by Underwood (1973). Wedivided the number of grid intersection points lying withinpores by the total number of points lying in the grid sectionthat was superimposed on each microstructure. We made 15measurements for each group of three specimens.

Results.For analyses of fracture toughness, it is essential that the

surface contain no significant residual stresses prior to inden-tation. If such surface stresses were present, a systematic var-iation in P/c,,3/2 with load would result. To determine whetherresidual stresses were present within the porcelain surfaces, weinitially made plots of P/c03/2 vs. P for each of the two com-mercial porcelains. The mean values and 95% confidence in-tervals for each mean for five loads are shown in Fig. 2. Ascan be seen in this Fig., the ratio P/c,3'2 was independent ofload. Therefore, we selected the maximum load, 19.6 N, forall measurements in this study, to minimize measurement in-accuracies associated with smaller cracks.To ensure that the magnitudes of fracture toughness values

were valid, we made calculations for the soda-lime-silica glassspecimens and compared these values with published valuesobtained by conventional test fhethods. Values of K1c and in-itial crack size are given in Table 1. Data are included for as-received specimens and for specimens that were polished through0.05 Am alumina. All specimens had been annealed at 386TCfor 100 h. The 95% confidence interval of the mean and coef-ficient of variation values are also given. The values of K1c(Table 1) were in excellent agreement with the value of 0.75MPa-m12 reported by Wiederhorn (1969) for soda-lime-silicaglass.The static fatigue behavior or the relative susceptibility of

the two body porcelains to water-enhanced crack growth in thepresence of residual stress can be determined from plots of logc vs. log t. Gupta and Jubb (1981) introduced a relativelysimple method for measuring the stress corrosion susceptibility

TABLE 1INITIAL CRACK SIZE AND CALCULATED Kic VALUES FOR THE

SODA-LIME-SILICA GLASS REFERENCE STANDARD

c, (>Lm) CV (%) Kc (MPa m) CV (%)As-received 130.7 4 4.0* 7.8 0.73 + 0.03: 10.5

(indented in air)As-received 132.2 - 5.1 10.4 0.71 + 0.03 10.5

(indented in oil)Polished 130.1 + 5.7 11.9 0.74 + 0.05 19.0

(indented in air)*Values preceded by + correspond to the 95% confidence interval for

the mean.

coefficient (n). The velocity of slow crack growth is usuallydescribed by

dc 1K1 \nv = dt V K / (4)

where vo is the critical velocity of crack growth at failure, K,is the stress-intenstty factor, and KIc is the fracture toughnessmeasured from the indentation method. After indentation, K,is equal to the right side of Eq. (1). For v,,t/c0 > > 1, Guptaand Jubb showed that

In c = (322) In t + I

I( 2n In + (3n+) In[( 2 VoXr3n+2 in[3n+2 2 Kr(,

(5)

(6)

where I is the intercept on the In c axis. It follows from Eq.(5) that the slope of the log c vs. log t plot is equal to 2/(3n + 2).Shown in Fig. 3 is a plot of log c vs. log t for annealed

soda-lime-silica glass specimens indented in air or in paraffinoil and aged in distilled water and oil, respectively, at 37TCfor three months. [The data points obtained at t = 2 min arenot shown because of the extended time scale.] However, themean values fit their respective linear plots very well. "Good-ness of fit" values were greater than 97%. Values of stresscorrosion susceptibility coefficient (n) are given in Table 2 forboth as-received and polished specimens at the initial one-month period and for the period between one and three months.Values of n were determined from linear regression analyses.Higher values of n indicate a lower stress corrosion suscepti-bility.The value of n found for soda-lime-silica glass specimens

aged in water during the first month (26.5) was higher thanthat (14.8) reported by Gupta and Jubb (1981) for the sameindentation load. These authors reported that values of n rangedbetween 14.7 and 19.2 for specimens aged in water and be-tween 21.8 and 24.8 for specimens aged in air at room tem-perature. Gupta and Jubb obtained these data at indentation

cam

ECu

a-cm

to0

0~

0uu

Porcelain C

I0 I100

Porcelain V

I0,15 10 15 20 25

LOAD (N)Fig. 2-Load-crack size parameter, P/c,, as a function of load.

Vol. 68 No. 6 1077

1078 ANUSA VICE & LEE

TIME (sec).6

2.3 F

2.2

2.16.( 6.2 6.4 6.6 6.8

.4

1.2

CLJ

cm

E

0b0

.0

0.8

7.0 7.2

LOG TIME (sec)Fig. 3-Log c vs. log t for annealed soda-lime-silica glass specimens

stored in distilled water and oil at 370C.

loads of 17.7 N, 19.6 N, 24.5 N. and 29.4 N. The only ap-parent difference in the two studies is that we used a higheraging temperature (37TC) in the present investigation, althoughit is possible that differences in the compositions of the soda-lime-silica glass specimens may account for this disparity.

Fig. 3 shows that crack growth continued beyond the initialmeasurement in the presence of air and paraffin oil for theannealed glass specimens. The stress corrosion susceptibilityfor the three-month period was greater (smaller n value of20.7) for specimens stored in water than for those stored in oil(n = 31.5). As-received glass specimens indented in air andstored in water at 370C exhibited mean crack sizes (c) of 130.7+ 4.0 pLm, 172.8 + 2.2 plm, 176.6 + 2.1 pm, and 178.8+ 2.2 Axm at periods of two minutes, one month, two months,and three months, respectively. Thus, crack growth was neg-ligibly small for the glass specimens during the later period ofaging.The effect of firing temperature on fracture toughness is

shown in Fig. 4 for each porcelain (nonannealed condition).The 95% confidence interval is represented by the dotted lines.The recommended firing temperatures are 971'C for porcelainC and 920'C for porcelain V. A slight increase in K1c resultedfrom a reduction in firing temperature for each porcelain. Thetoughness of porcelain C relative to that of porcelain V is

TABLE 2STRESS CORROSION SUSCEPTIBILITY COEFFICIENT (n) FORANNEALED SODA-LIME-SILICA GLASS AGED IN OIL AND

WATER AT 370CPolished As-received

Time (mo) Oil Water Oil Water0 to 1 105.9 25.2 42.2 26.5ito 3 20.1 15.7 31.5 20.7

940 920 900 880 860 840

T (0C)Fig. 4-Effect of firing temperature on fracture toughness of non-

annealed porcelain C and porcelain V specimens.

slightly greater and is less sensitive to a reduction in firingtemperature.

For the nonannealed specimens in Fig. 4, analysis of Kcvalues by the Tukey multiple range test revealed no significantdifferences (p>0.05) between the values of specimens under-fired by 14'C. For porcelain C, K1c was significantly greaterthan that of porcelain V at increments 14AC above (p<0.001),and 420C (p<0.0002) and 70'C (p<0.04) below their rec-ommended firing temperatures. For the specimens annealed100 h subsequent to firing, the Tukey multiple range test analy-sis revealed that the fracture toughness of porcelain C wassignificantly greater than that of porcelain V at the recom-mended firing temperature (p<0.001), and also when under-fired by 14AC (p<0.0001), 420C (p<0.0004), 700C (p<0.0001),and 840C (p<0.03).

Fig. 5 shows initial crack size (co) values (produced by an

indenter load of 19.6 N) for nonannealed specimens as a func-tion of firing temperature. The mean values of co ranged from87.1 Am (9290C) to 98.5 ,um (971'C) for porcelain C andfrom 87.6 ,um (8920C) to 104.0 Aim (934TC) for porcelain V.Corresponding mean values for annealed specimens ranged from92.1 rm (9010C) to 106.8 Vtm (915'C) for porcelain C and99.5 ptm (836TC) to 120.6 VAm (920TC) for porcelain V.

Values of crack size (c) for nonannealed porcelain speci-mens are plotted in Fig. 6 as a function of firing temperatureand time of exposure in distilled water at 370C. In this case,the anomalous values corresponding to the firing temperatureof 8920C (see Fig. 5) were excluded for clarity, and to illustratethe comparative crack growth with increased time. Comparedwith the previous Fig. for initial crack length, this plot of cracklengths for different exposure times resulted in a slightly steeper

E

w

N

CllUen

CD0

I I I I I 1 I1 1

P=19.6 N

n2=0.7- water

- oil n=31.5

Annealed Soda-Lime-Slica Glass

I I I I I I I I I I II I I I I A I a I I

J Dent Res June 1989

:6 I 07

.l11

CRACK PROPAGATION IN PORCELAIN

1 10

100

E00

iiNco

C.)

-J

z

901

01 I I I I I I I

Porcelain V

100 ---..

II LI I I

940 920 900 880 860 840T (0C)

Fig. 5-Initial crack size, c,, vs. firing temperature for nonannealedporcelain specimens.

120

I,110

LL

C)

00o

0

90h

940 920 900 880 860 840 820 800T (0C)

Fig. 6-Crack size of porcelain V specimens as a function of firingtemperature and time of exposure to water at 37TC.

slope and a narrowing of the 95% confidence interval. Forporcelain V, crack growth was progressively suppressed as thefiring temperature was reduced. This relative suppression wasnot observed for porcelain C.Shown in Fig. 7 is a plot of log c as a function of log t for

properly fired (A\T = 0C), nonannealed specimens of por-celains C and V stored in water at 370C. Although the stresscorrosion susceptibility coefficient is higher for porcelain C inthis case, we found no consistent difference between the twoporcelains for either the annealed specimens (Table 3) or non-annealed specimens (Table 4). A 19.6 N indentation load wasused for each case.

Shown in Fig. 8 is a plot of P/C3/2 vs. time for nonannealedspecimens which were fired at the recommended temperatures,

lbalr

150

140F

130__

:x

30 LNVoC.)

cc0

120_

1OF

o0-

106 107TIME AFTER INDENTATION (sec)

Fig. 7-Log c vs. log t for nonannealed porcelain specimens aged indistilled water at 370C.

indented, and aged in water at 370C for periods of one to sixmonths. The parameter (P/c3/2) is proportional to K, as definedin Eq. 1. For the specimens shown in Fig. 8, the residual stressfor crack growth in water has been reduced to an equilibriumstate for porcelain C, while for porcelain V, crack growth isnearly complete.The pore volume for each porcelain as a function of the

temperature difference between the actual and recommendedfiring temperatures is shown in Fig. 9 for body porcelains Cand V. As shown in this Fig., porcelain C was more sensitiveto void formation when underfired by 40'C or more below therecommended firing temperature. Pore volumes ranged from1.6 + 0.2% to 9.4 + 0.7% for porcelain C and from 1.6 -+0.2% to 5.7 + 0.9% for porcelain V between the recom-mended firing temperature and a firing temperature 840C belowthat recommended.

Discussion.Underfiring slightly increased the 'apparent' fracture tough-

ness of the two body porcelains studied, as shown in Fig. 4.Porosity can affect the calculated value of K1c, in several ways.First, pores may act to arrest cracks by reducing localizedstress. They can also serve as crack precursors if they are smallor irregular in shape. During the indentation process, poresmay contribute to structural densification and thereby reducethe driving force for crack growth as the indenter is removed.One would expect a reduction in the elastic modulus and pos-sibly the hardness with an increase in pore volume (Coble andKingery, 1956). Thus, the value of Xr would change and therebyaffect values of Kic. Since the initial size of semicircular cracksis reduced, an increase in K1c would result. Significant changesin K1c may not occur until the pore volume increases above athreshold value, although the effect would vary depending uponthe size and distribution of porosity.The role of leucite particles in inhibiting crack growth is not

known. Morena et al. (1984) observed that leucite particles do

Porcelain CP=19.6 N

a n - - _~~~~~~~~~~--t,~~~~~~~~ ~, -', ,-,-, s ,

980 960 940 920 900 81

NonannealedP=19.6 N

n=8.9

I I I l I I I

~~~ <~~~Porcelain V

~~~~~~~~~,~~ ~ ~ ~ ~ 1lamosIP

--I~~~~~~~~~~~~~t 2 min

li

I 11U.

90 , I I I I I I I I I I

7;; .

H .)..

Vol. 68 No. 6 1079

-10

I1

au,

TABLE 3STRESS CORROSION SUSCEPTIBILITY COEFFICIENT (n) FOR ANNEALED SPECIMENS AS A FUNCTION OF FIRING TEMPERATURE

Porcelain C Porcelain Vt< lmo t = lto3mo to lmo t = lto3mo

T (C) Oil H20 Oil H20 T (0C) Oil H20 Oil H20985 87.8 59.0 38.7 6.8 934 47.0 69.7 12.0 18.0971* 55.4 52.0 27.5 9.7 920* 48.4 107.9 20.4 7.3957 63.8 54.7 26.8 5.9 906 63.5 63.3 12.3 27.3943 75.1 50.5 12.5 7.0 892 55.9 59.5 16.6 22.2929 75.5 78.2 24.8 8.8 878 52.2 61.8 10.8 13.4915 46.4 51.2 16.3 11.7 864 61.9 92.0 11.6 8.5901 47.1 48.0 58.4 12.0 850 39.8 58.9 20.2 15.5887 54.0 49.7 21.3 9.5 836 62.6 71.2 59.1 17.4*Recommended firing temperature.

TABLE 4STRESS CORROSION SUSCEPTIBILITY COEFFICIENT (n) FOR NONANNEALED SPECIMENS AGED IN WATER AT 370C

Porcelain C Porcelain VT (0C) t < 1mo t= 1 to 6mo T (C) t < 1mo t- 1 to 6mo985 57.5 10.5 934 15.0 4.8971 68.5 8.9 920 103.0 4.0957 57.4 6.8 906 79.1 7.6943 80.3 5.7 892 59.6 3.6929 51.3 6.1 878 198.0 7.2915 82.6 7.0 864 105.0 4.2901 54.5 3.7 850 123.0 5.3887 51.0 10.6 836 110.0 6.2

22 I

20

18_

16k

14

10k

8

w

-i

cr0

0.a.

0 1 2 3 4 5 6TIME AFTER INDENTATION (mo)

Fig. 8-Plot of P/C3/2 vs. time determined from nonlinear regressionanalysis.

not arrest crack growth in incisal porcelain (C) specimens. Thedeflection of cracks away from these particles was attributedto a radial tensile stress field around each particle.

In our study, we observed cracks passing through both thecenters of pores and along the perimeters of other pores (Fig.10). Nadeau and Bennett (1978) observed that cracks wereattracted to cylindrical holes (0.5 mm in diameter) with an areafraction of 0.2 because of the concentrated stress on the dia-metral plane that is coplanar with the crack. They also con-cluded that pores can arrest cracks by attracting the crack andlowering the localized stress. Furthermore, extra energy is ex-pended when the crack curves away from the main crack plane.

6

4

2

+20 0 -20 -40 -60 -80AT (0C)

Fig. 9-Pore volume as a function of difference between the actualfiring temperature and the recommended firing temperature.

These mechanisms could explain the crack growth suppressionin porcelain V specimens (Fig. 6) with high void volumes.

For the recommended firing cycle, nonannealed porcelain C

CY

E

0-

0.

CL

NonannealedP= 19.6 N

0

C

I4 I I I

-~ ~ ~ocli

Porcelain C- Porcelain V

II I~~~~~~~~~~~~

1-At&1.w01

Al

LI'

ANUSA VICE & LEE J Dent Res June 19891080

12 _

nI

nL

J'6 V 6 (RCX4K FRUPZ&IalON JA IPORCE1LAIN 1081

for slurries of groundI glass Ileached in water, where AHis therectom ntalp). in thi sudy, the inerasd scptibilitf

dental nore lains with aging time nay be icusd hy an in_erea I hydroxyl ion concentration at the ira k tip.

4 ~~~~~~~~~~~~~~~~~~~~~~~~~~Itis appAent from Fables 3 and 4 that paraffin oil did notad quately pr )tet the crack tips against exposure to trace

qunii s f W tr. This finding uprtedbyt bcrfti3n 31u0p't a id Jnbb (198. Othr x igto hase hadlittle success in eliminating th nluneo moistr re whenusing tofu ne, mineralI oil, immersion oil, or kerosene IThpost-inldentation growth Of era ks in oil indic ites a high sus-ceptibility of denta porcelains to mhoi'ture-ass ce ra' rwh

Becauseo thi senstivity to moistur, meansurementsoshould be made in dry nittrogen or immediately after ~indenta-ton in order for acr esvli of Ki to be obt mned.

The absence of a clear trenidtoward stres corrosion suscep-~tibility when firn ~ te uperature is decersed snggets th t po-rosity h is little effect on moilture-asited crack growth, Possiblybecause the pores an ctious and becarl tecon, Am

r that Cannot rdiybe displaced by water near the rkFig0SF~~~timgeolpokir r on por hr V ~~~~~ tip. lbeseresnts sugg st that the enhanced thirmal shock icmnw~~~ldwundefir d by 56 C~~~~~~~~~sistance of 'p ctala V may more likelyba islt of its~lower

TABLE.5 ~~~~~~~~~~~~~snseeptibility to s ress corrosion at the initial stage of crackRELA lVE CRACK GROWTH IN WATER AT 37 C propagation.

min m, ' MO 3 mo 4MOm MO 6 fnO RELFLRVNCLSPorclain C 1.00 1.12 1.14 1.18 1.,23 1.25 12 N S R( NIU ANBKrMRHlF0B11

ACriti vrrrtrr o iiniiafio Tec iue fo Meauinrg~ciurPorcelaiii V 1.0 1.10 1. 3 1.22 1.321 37~ 1,39 es. Mcasremn§isICin ) 33ANUSAVIC ii K J RINGt E RD OS -\~iiS, nd

Values snirirg a Commo ndeln wer rt significantly diffrn.KING Di (1)51 A I in. 51 1oTetfrP rcath Mea Svit Jrs.Dct Re i6018( i69i

specimens showed the largest crack growth in water during the eCOBtE RlKt a KiNGERY w-!-0 196) Litc I Porcwiiwy oi Phyicafirst month; porcelain V speciiiimens exhibited greater craek Si AIm J ciS 9373DEHOLv'I.PH. ANUSAVicr K.~an BOYCE K 1. 1983) Anilysis ofgrowth from the second through the shixt months. We can also i~ Sr i kIrM1S iI)lt6 9see this effect in the plot Of Pc/2~vM. time in Fig 8. the 9paaeter, is proportional to the s~tressrintenisity factor. R MNSW 1811i M f sI(ts cec nSnethe initial crack sizes were, not significantly different ~ echiiniigy, DRK uhimin d NJ K id ss, N' wYork Academic(p0.05), values ot c/caar given in TAbl 5 to illustratth iiI'K ttnuN iSwCxs Pi

relative growth rates. The horizontal bars indicate values which C~6si 6~ss Ai CemS10 64-C '1 14aentsignificantly, different (p>.05). The stress orso JONES W TONiES P.A. in WItLSON, 11L1 117) Iiei:tini

susceptibility coefficients wereii consistent with the crack growth thii iid T1 i M ii~abieta)i rmbehavior for these specimen. As shown in Tabl 4,n values N5 SNi Sit Ii 13iiCiiwere 68.5 and 103.0 for porcelains C an V rspctveyin! Ra in 203t6.~the fist Month, while the values decicrased to 8,.9 arid 4.0, LANnviv\Sx aMA.Ixi (I 9811 ii i p1slerespectively, in the second through sixth months. I rDr.iCriiI iiiK11iinSy i i/AnThe values of n determined in the first month fore the an Ci S

LAWN FBIR FriFLP II F17)Eqifir Fe -ikvS .sealed speiens (fable 3) are, similar to te valued of 39 o Fra jr Alaieratamied by Somna et al. (1980)~for porcelain consisting of a glass MARSHALLFDu ei1otrl~ flaws ci rrsip~matrix and crstalline particlesof mulie qatz, and corunm Knp n Vi e ii taioiii/ CcrSo i6612dum. This susceptibility to stress corrosion could be3 a result M Ik V1(/K1,)PA L~ ~RiOf thieincrease in hydroxyl, ion Concentration at the tip of the C (4tlu uhsicC Mr~il riT1 aI Po I JDeno R~6 234 Ahradial crack frontts. Wiederhorin (1972) hy~pothesized that hy- NADEAI J.S HnENNLETT, KC. (I97)1 5 Mertsu t Dise isedl Piaedroxyltons attack- siloxane bonds Rand that silanolate group r ii vior ii,s Frctr MhaIc (PCramnics,chan be hydrolyzed bywaeWtIorilnl rup4ndh Pr 'cig atriernai Syr M~anccldroxyl ions throughh the reaction: C l IH 'ai t c woinu

Prss n0 961-97oH K iU i R K ISiuH R SiDt (7) §AEILLi C BL W (In37)TI verseSi gn t ItniiAuu

Vadredpti PFocetiir! ~JcR' 536 6)3~0K SiO ±H20 ~R SiOH ± OH () soM T.F MAITSUL 151i YMM~ENi8Apiiiti

He supported this hypothesis with data which showN~ed anL increase Cr F iDoiLirPIi iPr1r ACa5 (30:66-i 69Iin crack velocity, with increasing hydr~oxyl ion concentration. UNtDiRWOO0 E-Et 173).A I c ii1 Oalliv54rit iph In.

T1he aging temperatures should also affect the stress corsio Metals Hiandbouk, V61 Meto 'r ~y Sfr ie cld P i Diiia Kng