Sintering and grain growth in nanocrystalline ceramics

l

NanoSTRUCTUREDMATERIALS VOL. 1, PP. 173-178, 1992 ALL RIGHTSRESERVED

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SINTERING AND GRAIN GROWTH IN NANOCRYSTALLINE CERAMICS R.S. Averback, H.J. HSfler, H. Hahn and J.C. Logas Department of Materials Science and Engineering University of Illinois at Urbana-Champaign Urbana, IL. 61801

Nanocrystalline ceramics have been produced by the method of inert gas condensation of ultra-small particles and in situ compaction. Sintering and grain growth characteristics have been investigated by a variety of techniques, including: x-ray diffraction, gravimetry, BET, scanning electron microscopy and small-angle neutron scattering. It was observed that the green body densities of these materials vary greatly with compaction pressure and temperature. Densities > 90 % bulk. density could be achieved without grain growth in TiO2 using pressure assisted sintering at = 500 °C. Grain growth studies showed that the grain size varied as t 1/3. Moreover, the grain growth in dense materials was greatly accelerated compared to that in less dense material suggesting pore stabilization of grain size. Grain size could be partially stabilized by suitable doping or by pressure-assisted sintering. Correlations of grain size to various mechanical properties will be presented.

Introduction Scientific interest in ultra-fine grained powders for processing of ceramic components is motivated by the promise of improved sinterability, reduction in flaw sizes and low-temperature superplastic deformation. Recently, Gleiter and co-workers have developed a new method, which combines the method of inert gas condensation of small particles with in situ powder compaction, for synthesizing materials with grain sizes of 5 - 15 nm [1]. An attractive feature of this method is that the entire process is conducted in an ultra-high vacuum chamber so that the particles surface can be kept clean and the ceramics bodies kept pure. Although the initial particle size of the powders produced by nanophase processing is extremely small, the microstructural properties of the fully sintered materials depend on a variety of factors such as agglomeration, pore size distribution, sintering treatment and impurities [2]. Consequently, ceramic components which are formed from green bodies with very small particle sizes, can have very large final grain sizes and high porosities. These features are common when abnormal grain growth takes place. The present research was initiated to gain an understanding of sintering behavior in nanophase ceramics produced by the inert gas condensation method. We have done this by measuring the density and grain size of nanophase (n-) TiO2 and n-ZrO2 as functions of sintering temperature and sintering time. We have also examined the samples by scanning electron microscopy (SEM) for abnormal grain growth, by nitrogen absorption (BET) for the size distribution of the residual porosity and by small angle neutron scattering (SANS) for closed porosity. Preliminary attempts to control grain growth in n-

173

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TiO2 were made by using either pressure-assisted sintering or impurity doping. Finally, measurements of Vickers' microhardness were made to begin to develop relationships between grain size and mechanical properties. Experimental Methods The nanophase samples employed for this study were prepared by the Gleiter method of inert gas condensation and in-situ compaction [1 ]. Oxide powders are somewhat more difficult to prepare since they usually can not be evaporated directly. For Ti and Zr, however, it is possible to oxidize the metal nanophase powders and subsequently compact them. The base Ti powder was prepared by evaporation from a tungsten boat into a He gas atmosphere at " 4 mbar. Zirconium, which has a much higher cohesive energy than Ti, (6.32 eV versus 4.86 eV) was "vaporized" using a magnetron sputtering Source operated with a He-Ar gas mixture at an ambient pressure of ,- 0.7 mbar [3]. For both oxides, the initial particle size was 10 - 15 nm, and surprisingly, in light of the work of Granqvist and Buhrmann [4], it was rather insensitive to the evaporation or sputtering conditions. The n-TiO2 sample containing Y were prepared by first alloying the Y with the Ti and then vaporizing the alloy, analogous to the procedure for pure Ti. This method is acceptable for this alloy as Y is soluble in Ti, and in dilute solutions, the ratio of their vapor pressures is not far different from the ratio of their molar fractions (that of Y is somewhat higher). The stoichiometry of the ceramic oxides was determined by Rutherford backscattering spectrometry, although this method is generally useful only when deviations from stoichiometry are greater than = 5% in oxides. These measurements revealed that the as compacted n-TiO2 samples, which were black in color, have a stoichiometry of TiOI.8 whereas after sintering in oxygen, the sample turns white and the stoichiometry becomes close to TiO2 [5]. Diffraction patterns obtained from TEM indicated that while both samples were predominantly comprised of the ruffle phase, other unidentified phases were also present in the off-stoichiometric samples [6]. The density and porosity of the samples were determined from a combination of gravimetry, (BET) and (SANS) measurements. The BET measurements were employed on samples in their early stages of sintering when the porosity was open. The gravimetry measurements, using Archimedes' principle consistently yielded high densities on these low density samples. When the densities were ~- 85%, the two measurements came into agreement, presumably because the viscosities of our displacement media, carbon tetrachloride or de-ionized water, are greater than N2 gas. Samples with densities above 90% had closed porosity. The SANS studies were performed at ILL at Grenoble [7]. One of the difficulties of employing SANS on our nanophase ceramic samples is that the scattering from grain boundaries superposes with that from pores and the separation of these components is not straight forward. However, by compacting high density samples at 500 °C using high pressures, "~ 1 GPa, the scattering from grain boundaries could be calibrated as the contribution from pores to the scattering in such samples is small [7]. Results Densification and Grain Growth: Several sintering conditions were applied to the green body ceramics. The "standard" treatment was to oxidize the compacted powder at 350 °C under a small load, - 1 MPa, and then to compact it again at 150 °C at 1.4 GPa to maximize the green body density. The standard treatment led to fully oxidized specimens (ruffle for TiO2 and monoclinic for ZrO2) which had green body densities of 0.7 - 0.8 theoretical bulk density [5,8]. These samples were subsequently sintered for periods of - 24 hrs. at successively higher temperatures. The grain size and density of the sample was measured after each sintering step. Figs. 1 and 2 show representative data for TiO2 and ZrO2 respectively. Data for the densification of conventional

SINTERING AND GRAIN GROWTH

175

micron-sized powders are also shown in Fig. 1. These data clearly show the enhanced sinterability of nanophase oxide ceramics, with densification occurring several hundred K lower than in conventional powders. 150 DENSITY O n-TiO2 p---OGPa n-TiO2 p=l GPa A commercialTiO2 GRAIN SIZE D n-TiO^ !o=0OPa r~-'ri(~2 !0,=1GPa

L

n" LU

I

GRAIN

I

I

I

SIZE

,-7.

100

< 100

O

nO

I---

Z n" (5

o

z_

I

!

I/J.m at IO00*C

o

z w

Iii

(5

50 a

< 5o

z hi

kU > <

W

I

500 I000 SINTERING TEMPERATURE (*C)

0

FIG. I. Density and grain size o f nanophase TiO2 as a function o f sintering temperature [ 5 ].

I

t

1

I

0

200 400 600 800 1000 1200 SINTERING TEMPERATURE[°C]

FIG. 2. Density and grain size o f nanophase ZrO2 as a function o f slntering temperature [ 8 ].

Although sinterability appears to be improved in nanophase ceramics, it is also observed in Figs. 1 and 2 that substantial grain growth takes place before densification is complete. This observation is shown more clearly in Fig. 3 where grain size is plotted as a function of density. Initially, the grain growth is rather small as the ceramic body densities, but once the sintering body attains a density of = 90% bulk density, the grain growth becomes very rapid. 1000 I I I I io-S = as prepared 800

• 2h/550*C • 15hi 7OO*C

---I-Z~32]

t6 9

i°

\

i tO~o

200 Io-" 85 90 95 1~ DENSITY[%1 FIG. 3. Grain size versus density m a p s for n-TiO 2 and n-ZrO 2 [8]. 70

75

80

FIG. 4.

i

i

,

,,,,,I

i

I0

I

,

,,IL

I00 pore diameter (nrn)

i

,

i

,,I,,

I000

Pore size distributions in n-TiO2 as a function o f temperature obtained by BET [5].

To help understand the sintering behavior, BET measurements were employed to examine the evolution of the pore structure. These results are shown in Fig. 4 for n-TiO2. It is observed that

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the initial pore size distribution is rather broad with many large pores. These results are in good quantitative agreement with the SANS results, although the neutron scattering shows a significant number of very small pores not seen by the BET measurements [7]. We tentatively ascribe this extra porosity to closed porosity within small agglomerates. As expected, the pore size distribution shifts to larger sizes with increased densification, making further densification increasingly more difficult. The porosity becomes closed when the density reaches ,, 90%, about the same density where grain growth becomes rapid. Presumably, the open pore structure acts as a strong impediment to grain boundary migration. The SEM micrographs shown in Fig. 5 indicate that the grains were equiaxed, rather uniform in size, and showing no indication of abnormal grain growth. It appears, therefore, that the pores are dragged along with the grain boundaries, causing their coalescence and growth, and at the same time, slowing grain boundary migration and preventing abnormal grain growth. Grain growth during isothermal annealing was also studied in nanophase n-TiO2 samples that had been pre-densified using pressure assisted sintering [9]. Densities in excess of 90% can be obtained by this method. These results for grain growth were fit to a growth law of the form, d 3 - do3 = a t n exp{-Q/RT}

(1)

with the values of a, n and Q given in Table #1. Also shown in the figure is a single datum for the results of grain growth in a sample that had not been predensified. It can be seen that grain growth is greatly accelerated at high initial densities. TABLE 1 Parameters deduced from eqn. (1) for isothermal grain growth in n-TiO2. For the derivation of Q, n was taken to be 1.0. Specimen #2 values are in parentheses.

Temp [°CI

do [nm]

ro [% bulk]

n

ct [nm3-sl

Q [M/mole]

700

14 (14)

89 (94)

1.10(1.15)

7.29 x 1017

278

825

1.02

Control of Grain Growth: One of the strong motivations for developing nanophase ceramics is their potential for superplastic forming. This is described elsewhere in the proceedings [10], however, unless grain growth can be controlled, the prospects for superplastic deformation are seriously limited [11 ]. Various methods to control grain growth in fine-grained ceramic materials have been discussed by Brook [12]. One method for preventing grain growth during sintering employs pressure assisted sintering, and it was tried for n-TiO2. Pressure increases the driving force for densification as seen by the expression, Os = 2y/r + Oa

(2)

where Os is the total sintering stress, y is the surface energy and Oa is the applies stress. Since the increase in sintering rates is a consequence of the increase in the driving force for densification, rather than a change in atomic mobilities, no corresponding increase in the rate of grain growth is expected. The sintering treatment consisted of compacting the n-TiO2 powder in a powder press at 490 or 550 °C with an applied stress of I GPa. The results of these tests are included in Fig. 1. It is seen that the densities can be increased to ~, 95% by this treatment, without significant increase in grain size. Although the use of pressure to enhance densification rates can be effective, it is often not practical. Nor does it solve the problem of grain growth during superplastic deformation and therefore other means are necessary. We have begun attempts to control grain size using impurity

SINTERINGANDGRAINGROWTH

177

doping. As discussed by Brook, impurity doping can be very effective if the impurity segregates to grain boundaries. In this case, the grain boundary mobility, Msoh is reduced to, 1

Ms°l= M / l + MctCoa2)

(3)

where M is the grain boundary mobility in absence of impurities, a is the atomic spacing and Co is the bulk impurity concentration and a characterizes the impurity boundary interaction [12]. We have studied the influence of Y impurities in n-TiO2 [8,13]. The strong influence of Y is seen in SEM micrographs shown in Fig. 5." Whether the Y has indeed segregated to the grain boundaries and established a space charge that pins the boundaries is presently unknown. Bennison and Harmer, for example, have reported similar suppression of grain growth in A1203, doped with MgO, but they found no evidence for segregation. They suggested that the aliovalent MgO in alumina might prevent grain growth via alteration of the point defect chemistry rather than by MgO creating pinning forces on the boundaries [14]. Present work to study the chemistry of the grain boundaries in our n-TiO2 samples is presently being carried out. 11 1 l I I I

9

3

i 0

FIG. 5.

SEM micrographs o f a) pure TiO2 and b) Y doped TiO2 after annealing for 15 h at 1000°C. Note the different scales.

0.05

i

i

•

#2, P = 94%

•

#1,9=89% i

0.1 0.15 0.2 1Nd [1Nnm]

i 0.25

0.3

FIG. 6. Vickers Hardness as a function o f grain size [9].

Vickers Hardness: Mechanical properties, including hardness, bulk modulus, etc. are very sensitive to the porosity and microstructure of ceramic materials [15]. This was recently observed using Vickers hardness and elastic moduli measurements in n-TiO2 [16,17]. We have measured Vickers hardness on n-TiO2 samples that were prepared by pressure-assisted sintering and therefore were of high initial density and small grain size [9]. These mechanical properties measurements are the first on high density nanophase ceramics. Fig. 6 shows the hardness as a function of grain size. The grain size was varied by isothermal annealing at 700 °C as described above. Two features are particularly noteworthy in view of the hardness measurements performed during isochronal annealing treatments [13,16,17]: (i) The Vickers hardness in dense nanophase n-TiO2 samples is about equal to that in single crystal TiO2. (ii) Annealing at 700 °C leads to softening rather than to hardening. During isochronal sintering treatments, the hardness increased monotonically with temperature, and the highest value attained was lower than that found in our pre-densified specimens. The reasons for the Hall-Petch-like behavior, i.e., decreasing hardness with d1/2 are presently unclear. At the smaller grain sizes, where the decrease with increasing grain size is small,

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dislocations are unlikely to be involved, but they may play a role in the rapid decrease in hardness at d > 0.4/am. Conclusions This work provides some of the first systematic results on the sintering characteristics and properties of nanophase ceramics. The main results are: (i) both n-TiO2 and ZrO2 undergo densification at temperature much lower than they do in more conventional sized powders; (ii) they densify without significant grain growth until the density reaches ,- 90% bulk density and the porosity becomes closed; (iii) at densities above 90%, grain growth can be rapid; abnormal grain growth, however, was never observed; (iv) grain growth can be controlled by pressure-assisted sintering or by Y doping in n-TiO2, although complete densification appears to require some grain growth. (v) Vickers hardness in dense nanophase ceramics are as high as in single crystal TiO2, but it decreases with increasing grain size on annealing. Although this research provides a wide range of information about nanophase ceramics and begins to characterize their properties and kinetic behavior, it should be realized this research is still in its rudimentary stages and a fundamental understanding of the reasons for these properties remains to be elucidated. Acknowledgments

This work was supported in part by the U.S. Army Research Office under contract DAAL03-88K-0094 and the U.S. Department of Energy, Basic Energy Sciences under contract DE-AC0276ER0119. References

1. 2. 3. 4. 5. 6. 7. 8. 9. 10. 11. 12. 13. 14. 15. 16. 17.

R. Birringer, H. Gleiter, H.P. Klein and P. Marquart, Phys. Lett. 102A, 365 (1984). see e.g., W.D. Kingery, H.K. Bowen and D.R. Uhlmann, Introduction to C e ~ i c s , 2nd ed., J. Wiley and Sons, New York, (1975). H. Hahn and R.S. Averback, J. Appl. Phys. 67. 1113 (1990). C.G. Granqvist and R.A. Buhrman, J. Appl. Phys. 47, 2200 (1976). H. Hahn, J. Logas and R.S. Averback, J. Mater. Res. 5. 609 (1990). A. Zangvil, Y.-R. Xu, H.J. H6fler and R.S. Averback, unpublished. W. Wagner, R.S. Averback, H. Hahn, W. Petty and A. Wiedenmarm, unpublished. J.C. Logas and R.S. Averback, unpublished; J.C. Logas, Ph.D. dissertation, University of Ilinois at Urbana-Champaign. H.J. H6fler and R.S. Averback, Scripta Metall. et Mater. in press. H. Hahn and R.S. Averback, this volume. see e.g., Superplastic Forming of Structural Alloys, ed. N.E. Paton and C.H. Hamilton, Metallurgical Society of AIME Press, Warrendale PA. (1982). R. Brook, Proc. Brit. Cer. Soc.,ed. R.W. Davidge, No. 32, 7 (1982). R.S. Averback, H. Hahn, H.J. H/Sfler,J.C. Logas and T.C. Shen, Mats. Res. Soc. Symp. Proc. Vol. 153, p.3. S.J. Bennison and M.P. Harmer, J. Am. Ceram. Soc. 68, C-22 (1985). R.W. Rice, Mater. Sci. and Eng. 73, 215 (1985). R.W. Siegel, S. Ramaswamy, H. Hahn, Li. Zongquan, Lu Ting and R. Gronsky, J. Mater. Res. 3, 1367 (1988). M.J. Mayo, R.W. Siegel, A. Narayanasamy and W.D. Nix, J. Mater. Res. 5, 1073 (1990).