ZrO2CeO2Y2O3: electrical and mechanical property relationships with the microstructure

Materials Science and Engineering, B8 ( 1991 ) 265- 271 265

ZrO2-CeO2-Y203: electrical and mechanical property relationships with the microstructure

Nasruilah Khan* and Brian C. H. Steele Department of Materials, Imperial College, London SW 7 2BP (U.K.)

(Received September 20, 1990; in revised form February 13, 1991 )

Abstract

ZrO2-CeO2-Y203 ternary ceramic alloys were studied to investigate their electrical and mechanical behaviour. The electrical conductivity, fracture strength and toughness of one composition from each tetragonal (t), cubic (c) and two-phase (t + c) region were measured. These properties were correlated with the microstructure. The highest conductivity was found in the c-phase samples, while the t-phase samples exhibited the best fracture properties. The two-phase specimens showed intermediate behaviour for both the electrical and mechanical properties.

1. Introduction

A variety of alloys of ZrO2 ceramics have been investigated over the last decade [1-3]. Most of the studies were devoted to binary alloys doped with oxides of calcium, magnesium or yttrium. These ceramics possess high fracture strength, high toughness, good chemical resistance, a thermal expansion coefficient comparable to those of steels and low thermal conductivity. These properties allow their application in harsh environments under severe loading conditions.

Applications of these materials include fabrication of hot extrusion dies and cutting of indus- trial materials, particularly magnetic tapes, plastic films, tough fibres, fabrics and cables. They are also being used in automotive engine parts as exhaust port liners [4].

Certain ZrOz alloys also demonstrate significant electrical conductivity, which is useful in applications such as solid electrolytes in fuel cells and gas sensors [2, 5]. In particular, ZrO2-Y203 ceramics have shown exceptional ionic conductivity [6]. Moreover, the fracture strength and fracture toughness in this type of ceramic can be modified by compositional and microstructural control. Y203-ZrO2 tetragonal polycrystalline (Y-TZP) ceramics possess the highest strength

*Present address: Atomic Energy Commission, PO Box 1509, Islamabad, Pakistan.

and toughness among the Y203-ZrO2 ceramic alloys, whereas the fully stabilized ZrO 2 (FSZ) system offers the highest ionic conductivity. FSZ materials are being extensively used as solid electrolytes, although they possess lower mechanical strength and thermal shock resistance.

The solubility of Y203 in ZrO 2 is known to be quite high. Yet, in order to maintain charge neutrality, defects are created as follows:

Y203 z~o~ / oo o 2Yzr + V o + 3 0

Initially, the concentration of oxygen vacancies increases with increasing dopant concentration. Since the conductivity is proportional to the oxygen vacancy concentration, the ionic conductivity also increases with dopant concentration. How- ever, as the oxygen vacancy concentration becomes large, the probability of defect inter- action increases. Defect interactions have the net effect of reducing the mobility of the oxygen vacancies and consequently decreasing the ionic conductivity [7].

It has been shown that solid solutions of ZrO2 and CeO2 can also be formed. However, the Zr 4 + and Ce 4+ ions are both quadravalent under normal conditions, and solid solutions formed by these two compounds will not generate any oxygen ion vacancies. The unit cell volume of pure ZrOz could be increased by substituting cerium ions for zirconium ions [8]. Since Ce 4+

0921-5107/91/$3.50 © Elsevier Sequoia/Printed in The Netherlands

ions are much larger than Zr 4 + ions, this lattice stretching is expected to result in a decrease in activation energy.

Y-TZP ceramics suffer the disadvantage of degradation in fracture strength and toughness in a hot humid atmosphere owing to destabilization of the tetragonal phase [9]. Tetragonal CeO2-ZrO 2 ceramics (Ce-TZP) possess a useful fracture strength and fracture toughness (10-12 MPa m ~/2) [10]. In addition, this system has been found to be more stable in aqueous environments at slightly elevated temperatures, in contrast to the considerable degradation in strength observed in Y-TZP materials [9]. These materials can also have a larger critical grain size compared to Y-TZE It has also been observed that more than 80% of the tetragonal grains transform to monoclinic grains under stress [10].

The main disadvantage of the Z r O 2 - C e O 2 ceramics is associated with the low value of ionic conductivity and the observation that these materials can suffer partial electronic conductivity in a reducing atmosphere [11]. The other drawback lies in the sintering difficulty.

An attempt has therefore been made to over- come the disadvantages in binary systems by adopting the ternary Z r O 2 - C e O E - Y 2 0 3 system. It was considered useful to study the effects of C e O 2 and Y203 addition on the electrical and mechanical properties. It was expected that Y203 would increase the ionic conductivity via formation of oxygen ion vacancies and that C e O 2 would promote the fracture toughness and thermal and chemical stability of the ternary system. It would also be interesting to study the effects of CeO 2 on the grain size and grain boundary conductivity in the Z r O 2 - C e O 2 - Y 2 0 3 system.

2. Experimental procedure

2'. 1. Sample preparation A range of samples having four different com-

positions were prepared by mechanical mixing of the appropriate quantities of commercial powders ZrO2-CeO2 and ZrOE-Y203 (TOSO Inc. Manufacturing Corporation, Japan). The powders were uniaxiaUy pressed to form ceramic pellets and then sintered at 1550 °C for the microstructure and impedance measurements.

For the mechanical properties, uniaxially pressed and sealed green compacts (bars) were isostatieally pressed at 200 MPa and sintered at 1580°C in air for 10 h. The densities of these

specimens were 99.4% of the theoretical values for tetragonal (t), 98% for tetragonal plus cubic (t + c) and 95% for cubic (c) phase samples. The theoretical densities were determined by the X-ray diffraction method (described in Section 2.2). The sintered test bars were ground and polished to minimize surface defects.

2.2. X-ray diffraction and microstructure Crystallographic measurements were carried

out to determine the crystal structure, phase content and theoretical density. A Philips diffrac- tometer was used with Cu Ka radiation with a 2 0 step of 0.01° and a count time of 1 s per step over the 20 ranges 26o-33 ° (to resolve the monoclinic (m) phase) and 700-76 °. This second range covers the cubic-related (400)c peaks conventionally chosen to resolve the cubic and tetragonal peaks. The X-ray data were processed on a BBC micro- computer using an in-house peak-stripping pro- gramme. Quantification of the phases present was done by peak area measurement using the cali- bration of Toraya et al. [12] for ZrO2 ceramics. The lattice parameters were determined by peak position and the corresponding unit cell volume was calculated. The theoretical density was obtained from the unit cell weight and volume.

The microstructure of the specimens was examined using a JEOL JSM (T-220 model) scanning electron microscope. The grain size (D = 2r) was determined by measuring the aver- age intercept length (L) and applying the formula

4~rr L -

3

Grains were considered to be spherical in shape and no correction was applied for grain shape.

2. 3. Mechanical measurements The fracture tests were carried out in a univer-

sal testing machine (Instron Model 1342). Three- point bend and single-edge notch beam (SENB) tests were applied in order to determine the fracture strength (Of) and fracture toughness (Klc) respectively. The results were obtained by using the relationships [13]

3FL of = 2bd 2

K~c= YOfC 1/2

where F is the load in newtons, L is the knife

edge span, b is the width, d is the depth, Y is a geometric factor and c is the notch depth (0.5 ram). The test bars' dimensions were 3 mm × 4 mm x 40 mm, the knife edge span was 30 mm and the cross-head speed was 0.2 mm min-~ for these measurements.

2.4. Electrical measurements Impedance spectroscopy analysis was per-

formed on disk samples 12.5 mm in diameter by 1 mm thick sputtered with platinum electrodes. A computer-controlled frequency response analyser (Hewlett-Packard Model HP4192A LF) was used for the electrical measurements. The measurements were conducted over a frequency range from 5 Hz to 13 MHz with an applied potential of 0.05 V and over a temperature range from 225 to 725 °C in air. Within this temperature range the frequency response of the equipment covers the grain lattice (bulk), the grain boundaries and the electrode arcs of the impedance plot.

Graphical estimations of the bulk and grain boundary resistivities were made according to a recent book [14] which includes summaries of experimental details.

The impedance results were graphically presented by the computer in terms of the complex resistivity (Pc), which is defined as

Z A lOc-

d

where Z is the real part of the complex impedance plot (ohms), A is the cross-sectional area (cm 2) and d is the thickness (centimetres) of the specimen.

3. Results

3.1. Mechanical behaviour The results for the mechanical properties of

the various phases are presented in Table 1. The

267

highest fracture strength and fracture toughness can be seen for the t-phase samples, the lowest for the c-phase samples. The two-phase samples exhibited intermediate behaviour. In general, 12 bar specimens from each composition were prepared to determine the fracture properties and a variation of +5% was noted in the results obtained.

A typical microstructure of a t-phase sample (ZCY1) with a grain size of 1.0 ,urn is shown in Fig. 1 (a). The figure indicates an increase in critical grain size as a result of CeO 2 addition com-

Fig. 1. Microstructures of (a) 3Y203-2CeO2-95ZrO 2 (TZP) and (b) 8Y:O3-2CeO2-90ZrO~ " (FSZ) sintered at 1580 °C for 10h.

TABLE 1

Fracture strength and fracture toughness of selected ZrO2-CeO2-Y203 (ZCY) compositions

Sample Phase CeO 2 Y203 (mole%) (mole%)

of Notch width c Kit (MPa) (nun) (MPa m 1/2)

ZCY1 t 2 3 ZCY1 t 2 3 ZCY2 t + c 2 6 ZCY4 c 2 8

850 0.6 13 850 0.1 7 225 0.1 3 147 0.1 2

268

5 1.8XI0

E u

E

6 X 10 4

E

N

o o o 05

o " • % 5 • 0 3 0 0 i °

t bulk 40 0 • 02 0 5 ° 0 • cO

"4 3 electrode ~'~..'~

. . . . . . . ~x~° s 2~,0 s - ~o s

Z reo l ( o h m - cm)

5 XIO

Fig. 3. Impedance plots of 2%CeO2-ZrO 2 with different amounts of Y203 at 300 °C: o, 3% Y203 o, 8% Y 2 0 3 . Num- bers are log (frequency).

Fig. 2. Scanning electron micrographs of fracture surfaces of (a) TZP and (b) FSZ ceramics.

pared with a critical size of less than 1 /~m for binary 3mole%Y203-ZrO2 [15]. The large grains of a c-phase sample are shown in Fig. 1 (b).

When the notch width was kept at 0.6 mm, the scatter in Ktc data was not significant for a particular composition. Sample ZCY1 was selected to investigate the influence of notch width on Klc. The effect of notch width was found to be pro- nounced in this study (Table 1).

Fracture surfaces of samples ZCY 1 and ZCY3 are shown in Fig. 2. The figure shows that the fracture in the t-phase composition was typified by an intergranular mode. The two-phase and c-phase samples were less dense and a microstructure similar to fine-grained TZP samples could not be obtained. In the (t + c)- or c-phase specimens the fracture might occur transgranu- larly as shown in Fig. 2(b). No evidence of t--" m transformation in (t + c)-material was found when the fracture faces of these samples were subjected to X-ray diffraction for phase analysis. Therefore no enhancement in toughness was recorded.

3.2. Electrical conductivity The conductivity results of the four composi-

tions were compared. The impedance plots for the t- and c-phase samples at 300 °C are shown in Fig. 3. It can be seen that the bulk resistivity (Pgi) is higher than the grain boundary resistivity (pg~) in both the t- and c-phase samples. The figure also illustrates the slightly lower bulk resistivity in the t-phase compared with c-phase at low temperature (300 °C). However, the t-phase samples exhibited higher strain boundary resistivity at this temperature. The measured conductivity results at selected temperatures (285 and 685°C) are listed in Table 2.

The resistivity results from plots over a range of temperature (200-725 °C) such as given in Fig. 3 were used to draw Arrhenius plots (log(oT) vs. log(l/T)). The Arrhenius plots for the bulk and grain boundary phases of these compositions are shown in Figs. 4(a) and 4(b) respectively.

These plots were used to obtain the activation e n e r g i e s (Ea) and pre-exponential factors (o0) given in Table 2. The errors in E a and o0 were less than 0.5% and 10% respectively in this work for the bulk conductivity. The error in o0 was much higher in the grain boundary results. The conductivities for both bulk and grain boundaries at 800 °C were extrapolated from the activation energies and pre-exponential factors.

The activation energy (Ea) and pre-exponential factor (%) values obtained from Figs. 4(a) and 4(b) are lower in the t-phase compared with the c-phase (Table 2). As a result of a large o0 value, the c-phase attains a conductivity higher than that of the t-phase at 800 °C.

The conductivity of all the phase regions is plotted in Fig. 5 against the Y203 composition at 285, 485 and 800 °C. The plots demonstrate that

TABLE 2 Conductivity, activation energy and pre-exponential factor for t-, (t + c)- and c-phase samples

269

Sample % x 106 % °C Ea O ° at 285 °C at 685 (eV mole-i) (S cm-i) (Scm -I) (Scm -])

% at 800 °C (Scm -I)

Ogb

at 800 °C (S cm -1)

ZCYI (t) 3.87 0.006 0.91 330 0.019 0.26 ZCY2(t + c) 3.20 0.011 1.00 2200 0.041 0.78 ZCY3(t + c) a 3.83 0.007 0.94 590 0.023 0.53 ZCY4(c) 2.69 0.016 1.06 5050 0.054 1.35

aSample ZCY3 contains 4 mole% Y203 .

2

1

0

p,. b -2 .Q

(a)

ZCY1 =

0 8 ' i o " 1'.2 ' 14 1'6 ' l:a " 2 0

1000/T'(1/K)

ZCYI o

i ° p-

b -2

-Q -3

I 0 1.2 1,L 1.6 18 2 0

1000/T (l/K) (b) Fig. 4. (a) Bulk and (b) grain boundary conductivity com- parisons o~ low CeO 2 ternary compositions.

the conductivity rises with Y203 addition in the t-phase region, decreases in the c-phase region and shows an intermediate behaviour in the two- phase region at 285 °C. At 485 °C the c-phase attains a conductivity a factor of 2 greater than that of the t-phase composition. An increase in Og i can be seen from 4 to 6 mole% Y203 at this temperature.

At 800 °C the c-phase achieves a large differ- ence in the conductivity value. The experimental results in this investigation confirmed that the t-phase compositions are the best bulk ionic conductors at temperatures lower than 400 °C and that the c-phase exhibits the best ionic conductivity at fuel cell application temperatures (800 °C or above). The two-phase compositions (ZCY2 and

_~ ×

3

x

6 i ( c } a t 800C

/,

2

I 8

6

2 5

( a l a t 285C

3.

2-

J

2 /, 6 8 10

Mole% y~.tiia

Fig. 5. Conductivity behaviour in three phase regions for 2% CeO2-ZrO 2.

ZCY3) have an intermediate conductivity which depends upon the phase ratio.

4 . D i s c u s s i o n

4.1. Mechanical behaviour The mechanism responsible for the superior

mechanical properties in t-ZrO 2 is the t--,m phase transformation which occurs in the stress field of the propagating crack [10]. Therefore t-phase with high transformation potential is required to achieve such superior mechanical properties.

The critical grain size of the ternary system was more than 1 / tm for such compositions compared with less than 0.2 gm for binary 2mole%Y2Oa-ZrO2 [16]. The critical grain size of the ternary system was increased as a result of

27O

CeO 2 addition, which allows the fabrication of a ceramic with larger t-grains which can transform to a greater extent compared to binary Y-TZP materials. Hence an improvement in toughness can be expected from ternary systems. The results for the t-phase suggest that transformation toughening is responsible for the high values of strength and toughness.

It is apparent that the fracture toughness is dependent on the notch width and the results vary between 7 and 13 MPa m ~/2 for samples of the same composition. The phenomenon behind the notch sensitivity, however, is not clear.

The strength and toughness of the (t + c)- and c-phase samples were lower than for the t-phase because of the larger Y203 content. The strength and toughness decrease with addition of Y203. The transformation potential decreases with decrease in t-phase, therefore no enhancement in strength and toughness can be expected as a result of transformation. The lower values than those expected may be interpreted as a result of lower density. Therefore microstructure develop- ment with experimental density close to that of theoretical density is required. This may be achieved by preparing ceramic powders of small particle size with uniform size distribution or by hot isostatic pressing.

4. 2. Electrical conductivity Investigations on the relationship between the

microstructure and the impedance characteristics have been made. The ionic conductivity results follow the Arrhenius law in the measurement temperature range.

The conductivity results in Fig. 5 show that the c-rich phase is more conducting above 485 °C when compared to the t~phase composition. The low temperature behaviour indicates a decrease in conductivity from t- to c-phase. This can be attributed to the formation of defect interactions among the vacancies and the y3 + ions as a result of electrostatic attractive forces between oppo- sitely charged carriers. This effect increases with addition of Y203 and the effective charge carrier concentration decreases.

As the temperature increases, it gives rise to conductivity of the c-phase composition and the Arrhenius plots of the c-phase intersects those of the t-phase as shown in Fig. 4(a). This behaviour in the cubic ternary phases demonstrates that as the temperature rises, the concentration of free vacancies increases.

The activation energies are lower in compositions with t-rich phases compared to those with c-rich phases, which indicates a larger number of defect associates in the c-phases. These associates have defect pair energies (Ha) which can change the total activation energy (E a) according to

E~=Hm+H~

where H m is the enthalpy for mobility of oxygen ion vacancies. An increase in either H m or H a can cause a reduction in conductivity.

The influence of CeO2 in ZrO 2 solid solution is different from that of Y203 addition. Cerium, having a + 4 valency similar to zirconium, does not introduce vacancies in the oxygen sublattice to maintain charge neutrality. The ionic radius of Ce 4+ is 0.97 A compared with 0.78 A for Zr 4+ and 1.019 A for y3+ [8]. It was therefore expected that addition of CeO 2 would reduce the activation energy as compared with the binary Y203-ZrO2 as a result of introducing more open paths for vacancy mobility which could be associated with the observed increase in lattice parameter values [ 17].

The a0 values in the c-phase were an order of magnitude higher than those in the t-phase, with an intermediate behaviour in the two-phase region (Section 3). The increase in a 0 was found to be proportional to the vacancy concentration introduced by Y203 incorporation. The values of a0 give an indication of the vacancy concentration when the results for the t- and c-phases are compared. The % value for binary ZrO2-CeO 2 is only 13 S cm- 7, which provides a good argument for the dependence of a~, on the vacancy concentration [17].

The activation energy behaviour of the t-phase can be explained via the defect model for the dilute CeO2-Y203 system [18]. This model assumed an ordered array of defect associates (2Y/ce-V~') and suggested that the conductivity was proportional to the third root of the dopant concentration.

4.3. Grain boundary conductivity Since ZrOz-based ceramics are polycrystalline

materials, it is essential to consider the influence of grain boundaries. A grain boundary layer of non-conducting impurities such as SiO2 may be responsible for grain boundary resistance. Segre- gation of dopant on the grain boundaries and triple point can also be responsible for reduction in conductivity at low temperature as a result of

complex formation. The low conductivity of small grain tetragonal material compared with cubic samples can be attributed to the larger surface area in the t-ZrO2 ceramics for a similar impurity level (SiO2, etc). The grain boundary arc disappeared above 700 °C. Therefore in ZrO2 ceramics grain boundary resistance is not important above 700 °C.

5. Conclusions

The c-phase composition exhibits lowest fracture strength and fracture toughness; on the other hand, it attains highest conductivity at fuel cell application temperatures. The increase in E a in the c-phase has led to the recognition of the importance of defect association in this ternary system.

The grain boundary conductivity is more important in t-phase compositions compared to the c-phase, particularly at temperatures below 500 °C. The grain boundary conductivity of t-, (t+c)- and c-phase systems was improved by CeO2 addition as a result of an increase in grain size.

Since the critical grain size of ternary ceramics increased with increase in CeO2 addition, we can fabricate ceramics with relatively large grain size to cause a reduction in grain boundary resistance and improved degradation of the t ~ m transformation.

The two-phase system could be useful for those situations where both the mechanical and electrical properties are required to be optimized.

Acknowledgment

One of the authors (N.K.) would like to thank Dr. K. A. Qureshi, Technical Manager, Pakistan Atomic Energy Commission for improving the

271

manuscript and providing the facilities to prepare this presentation.

References

1 A.H. Heuer and L. W. Hobbs (eds.), Advances in Ceram- ics, Vol. 3, Science and Technology of Zirconia 1, The American Ceramic Society, OH, 1981.

2 N. Claussen, M. Ruhle and A. H. Heuer (eds.), Advances in Ceramics, Vol. 12, Science and Technology of Zirconia 11, The American Ceramic Society, OH, 1984.

3 N. Yamamoto and H. Yanagida (eds.), Advances in Ceramics, Vol. 24B, Science and Technology of Zirconia 111, The American Ceramic Society, OH, 1988.

4 M. R. Pascucci, in H. M. Schwartz (ed.), Engineering Applications of Ceramic Materials, American Society for Metals, Metals Park, OH, 1985.

5 B.C.H. Steele, Chem. Ind., 6 (1986)651-656. 6 N. Bonanos, J. Drennan, R. K. Slotwinski, B. C. H. Steele

and E. P. Butler, Silicate Ind., (September-October 1985) 127-132.

7 J. E Baumard and P. Abelard, in N. Claussen, M. Ruhle and A. H. Heuer (eds.), Advances in Ceramics, Vol. 12, Science and Technology of Zirconia 11, The American Ceramic Society, OH, 1984, pp. 555-571.

8 R.D. Shannon, Acta Crystallogr. A, 32 (1976) 751. 9 H. Schubert and G. Petzow, in N. Yamamoto and H.

Yanigida (eds.), Advances in Ceramics, Vol. 24B, Science and Technology of Zirconia Ill, The American Ceramic Society, OH, 1988, p. 21.

10 K. Tsukuma and M. Shimada, J. Mater Sci., 20 (1985) 1178-1184.

11 E. P. Baker, M. Iqbal and B. E. Knox, J. Mater. Sci., 12 (1977) 305-310.

12 H. Toraya, M. Yoshimura and S. Somiya, J. Am. Ceram. Soc., 7(1984) C119-C121.

13 R. W. Davidge, Mechanical Behaviour of Ceramics, Cam- bridge University Press, Cambridge, 1979, p. 165.

14 J. R. MacDonald (ed.), Impedance Spectroscopy, Wiley, 1987.

15 E E Lange and D. J. Green, Adv. Ceram., 3 (1984) 217. 16 M. Watanabe, S. rio and I, Fukuura, in N. Claussen,

M. Ruhle and A. H. Heuer (eds.), Advances in Ceramics, Vol. 12, Science and Technology of Zirconia H, The American Ceramic Society, OH, 1984, pp. 391-398.

17 N. Khan, Ph.D. Thesis, University of London, 1990. 18 B. C. H. Steele, in T. Takahashi (ed.), High Conductivity

Solid Ionic Conductors, Recent Trends and Applications, World Scientific, London, 1989, pp. 403-446.

Documents

ZrO2CeO2Y2O3: electrical and mechanical property relationships with the microstructure