Microstructural changes in alumina during hot isostatic pressing

Materials Science and Engineering, A 109 ( 1989 ) 37 - 43

37

Microstructural Changes in Alumina during Hot lsostatic Pressing* J. BESSON and M. A B O U A F Ecole Nationale SupOrieure des Mines de Paris Centre des Mat~riaux, BP 87 91003, Evry ('&lex (France)

(Received June 2, 1988)

Abstract

Grain growth enhancement during deformation is a well-known phenomenon which has already been observed during creep tests in superplastic alloys" and ceramics. In this study it is shown that similar effects are obtained during densification by hot isostatic pressing of a fine-grained alumina powder. Changes in grain shape during densification are also observed. Finally, a model is developed which allows the modelling of grain growth enhancement in finite element calculations. 1. Introduction Temperatures required to densify ceramic powders totally can be reduced when using hot isostatic pressing (HIP) rather than sintering. Therefore, HIP has the potential for providing ceramic parts with homogeneous, isotropic and fine microstructures. However, microstructural parameters such as grain size or grain shape may depend on the way densification is achieved. Earlier experiments on alumina [1, 2] or superplastic alloys [3, 4] have shown that grain growth can be enhanced during deformation, and can be related to strain, strain rate or stress. However, all these studies deal with dense materials tested under tensile or compressive uniaxial loading. The purpose of this work is to show that similar effects can be obtained during the densification of an alumina powder by comparing grain growth during sintering and HIP. On the other hand, an attempt is made to model grain growth during deformation. This allows existing models of the densification of powder compacts [5-7] to take into account grain growth and strain enhancement of grain growth during deformation. This could be very useful *Paper presented at the Symposium on Ceramic Materials Research at the E-MRS Spring Meeting, Strasbourg, May 31 -June 2, 1988. 0921-5093/89/$3.50

due to the fact that grain size changes can have drastic effects on the densification kinetics.

2. Experimental procedure A high-purity ex-alum alumina powder (HR8 Criceram; A1203 content greater than 99.98%) was used in this work; the chemical analysis, as stated by the supplier, is given in Table 1. HIP tests were carried out using as containers 304L stainless steel tubes a few centimetres long with an internal diameter of 16 mm and a thickness of 1 ram. The powder had to be precompacted because of its poor loose density; the relative green density of the powder compact was equal to 0.51 after precompaction. The container was outgassed at 600 °C and sealed under vacuum. The temperature used was 1300 °C, with varying pressures and holding times. In each case the heating rate was 20 °C min -l. At the end of the hold, the power supply was cut off, resulting in a rapid drop of both temperature and pressure. A typical HIP cycle is presented in Fig. 1. After HIP the container was dissolved in aqua regia. Sintering experiments were carried out in a vacuum furnace using cylindrical, precompacted ceramic samples. Densities were measured using helium pycnometry or by measuring the dimensions of the samples when open porosity was still present. Each sample was polished with diamond paste and thermally etched (1280°C, 20 min or 1350 °C, 20 min) to reveal the grain boundaries. Morphological parameters (grain size, shape

ChemicalanalysisofAI203

TABLE l

Impurity (ppm) Ca

Fe

Mg

Na

Si

Ti

Other

<5

10

<3

48

40

<3

<5

© Elsevier Sequoia/Printed in The Netherlands

38 I

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factor) were obtained by tracing grain boundaries and porosities on the SEM micrographs with black ink on an overlay suitable for image analysis (Nachet NS1500 apparatus). In every case, 600-1200 grains were randomly sampled. The edge effects bias in planar sampling was compensated by using the Miles procedure [8]. The diameter of a grain is calculated by converting each grain area into an equivalent circle. The mean grain size is the average of equivalent diameters calculated using planar sections. Mathematical procedures developed to calculate the three-dimensional distribution of particle size [9] cannot be used because the particles are nonspherical. However, we assume that the difference between two-dimensional-mean grain size and three-dimensional-mean grain size remains small. All structures are supposed to be isotropic because of the densification processes (HIP or sintering).

3.

Results

3.1. Densification Densification curves during sintering are shown in Fig. 2 for different temperatures and holding times (separate samples were used to obtain each data point). Final densities after HIP are shown in Fig. 3 as a function of the applied pressure. In most cases a relative density of 99% is obtained. It should be noted that a pressure of 200 MPa is less effective for the densification than lower pressures. Several holding times (0, 30 and 60 min) were used under a pressure of 100 MPa (Fig. 2); in each case a relative density of 99% was achieved.

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PRESSURE (NPa)

Fig. 3. Density as a function of HIP pressure (holding time 60 min, temperature 1300 °C).

3.2. Grain size Grain-growth kinetics during annealing can be described using formula [10] k d----G= dt G - Gm

(1)

where k = k 0 exp( - Q/RT ), G being the grain size, k 0 a constant, R the gas constant, Q the activation energy and T the absolute temperature. We describe grain growth during sintering using eqn. (1). A least square fit adjustment gives: ( m + 1)k0= 1.7 x 10 -11 m 4 S - ] , Q = 570 kJ mol -l and m = 3. The initial grain size is 0.25 /~m. A similar grain size exponent was found during static annealing of dense alumina [1], of an alumina-zirconia mixture [11], and during sintering of an FeO-doped alumina [ 12]. Our results are plotted in Fig. 4, using eqn. (1) after integration G m+l - G 0 m + l = ( m + l ) k t

(la)

where G Ois the initial grain size and t the time. Figure 5 shows that most of the grain boundaries are free of porosity, although pore dragging

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100

150

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(MPo)

Fig. 7, Final grain size after HIP (holding time 60 min, temperature 1300 °C) as a function of pressure.

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Fig. 6. Pore-boundary separation in alumina sintered for 96 h at 1400 °C.

can be observed (Fig. 6). Therefore it seems that grain growth during sintering is not controlled by porosity. It should be noted that data obtained on already dense samples are also in good agreement

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Fig. 8. Final grain size after HIP as a function of holding time (pressure 100 MPa, temperature 1300 °C) and comparison with grain growth during static annealing.

with eqn. (1). It has been reported [1] that grain growth could be controlled by fl-A1203-type precipitates lying at grain boundaries. Such precipitates can be formed in alumina containing alkaline impurities such as sodium or potassium [13]; this is so in the case of the alumina used in this study (see Table 1 ). Grain size after an HIP cycle is plotted as a function of the applied pressure with a constant holding time (60 min) and temperature (1300 °C) in Fig. 7. The final grain size varies as a linear function of the pressure. Figure 8 shows the grain size as a function of the holding time (pressure 100 MPa, temperature 1300 °C); in the same figure we indicate grain size obtained by assuming pure static annealing (thus using eqn. ( 1 )). Although the relative density does not change in these samples (Fig. 3), the grains are growing faster than during pure annealing.

40

Figure 9 illustrates microstructures obtained during various HIP cycles. Under 200 MPa, abnormal grain growth was observed as shown in Fig. 10; this could possibly explain the lower density observed under this pressure (Fig. 3). To check the possible effect of pressure on grain growth, although pressure has been reported to inhibit grain growth when sufficiently high [14], we carried out post-HIP treatments on already dense materials obtained by hot pressing. The results are summarized in Table 2 and show that pressure has no observable effects on grain growth: calculated final grain sizes (assuming static annealing, eqn. (1)) are in good agreement with experimental observed grain sizes. It then becomes obvious that grain growth enhancement during HIP is present as a consequence of deformation rather than pressure.

3.3. Shape In order to estimate shape changes of grains during deformation we used the shape factor Fg

Fig. 10. Abnormal grain growth in alumina during HIP (holding time 60 min, temperature 1300 °C, pressure 200 MPa).

TABLE 2 Observed and calculated (eqn. (1)) grain growth in dense and porous materials during HIP (holding time 60 rain, temperature 1300 °C, heating rate 20 °C min- 1)

Pressure (MPa)

Fig. 9. Typical microstructures after HIP (holding time 60 min, temperature 1300°C) for different pressures: (a) 50 MPa, (b) 100 MPa, (c) 150 MPa, (d) 200 MPa.

Initial relative density

Initial grain size (/am)

Observed final grain size (/am)

Calculated final grain size (/am)

25

0.51 0.98

0.25 0.59

0.35 0.62

0.32 0.60

100

0.51 0.98

0.25 0.59

0.75 0.62

0.32 0.60

41

defined by the geodesic diameter Og [15] D 2

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(3)

where d*(A, B) denotes the distance between the two points A and B along a path included in the particle, represented by the set E An illustration of the concept of geodesic diameter is given in Fig. 11. Variations of the mean shape factor during sintering are shown as a function of time in Fig. 12. The initial shape factor was calculated using the projected area of initial individual particles (Fig. 13). The data of Fig. 12 can be explained as follows: initial particles have worm-like shapes; they become rounded during densification. For sufficient annealing times, grain growth along specific crystallographic directions can occur, resulting in more elongated grains.

Similar measurements were carried out on samples subjected to HIE The mean shape factor as a function of the applied pressure for a constant holding time (60 min) is shown in Fig. 14. Increasing pressure promotes the rounding of the particles but at sufficiently high pressure, abnormal grain growth can occur, resulting in a higher shape factor. Under a pressure of 100 MPa, increasing the holding time (0, 30 and 60 min) does not change the final mean shape factor; in each case, we have Fg ~ 2. 4. Grain growth enhancement: interpretation and discussion Authors working on superplastic alloys have proposed several explanations for the strain enhancement of grain growth. The model pro-

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Fig. 14. Mean shape factor during HIP (holding time 60 rain, temperature 1300 °C) as a function of pressure.

42 posed by Clark and Alden [3] assumes that deformation produces an excess of vacancies at grain boundaries, resulting in an enhancement of mobility. Another explanation was given by Wilkinson and Caceres [4]: superplastic deformation resulting from grain sliding causes damage at grain boundary triple junctions. Grain boundary migration can recover this damage; since there is a geometrical bias (grain boundaries are moving from big to small grains), grain growth enhancement can occur. However, since these models are dealing with metallic superplastic alloys, it is difficult to use them in order to check our experiments. Concerning ceramic materials, experiments carried out on MgO-doped alumina have shown that grain growth can be enhanced during compressive creep tests [1, 2]. Experimental results show that the applied stress does not influence the grain size exponent in eqn. (1) but it does increase constant k. The authors have proposed the following formula, taking into account stress enhancement of grain growth G = Ko exp( Gm \

V*o-Qg] ~ ]

(4)

where Qgis the activation energy, o the applied compressive stress, K 0 a constant and V* an activation volume. This model deals only with compressive uniaxial creep tests carried out on dense material. Therefore it would seem difficult to extend eqn. (4) to the case of multiaxial stress states: instead of o we could use the pressure (but we have seen that pressure itself has no effect on grain growth) or the Mises equivalent stress. However, the Mises equivalent stress represents a shear stress, and it would not be consistent to use an activation volume since shear deformation does not produce volume changes. We propose a model of grain growth enhancement during deformation which uses the macroscopic concept of power density produced during visco-plastic deformation. This model takes into account every kind of deformation, and can be applied to porous materials. The power per unit volume w is given by

w=0:

vp

(5)

where evp is the visco-plastic strain rate tensor and 0the stress tensor. We assume that part of the power will not be dissipated, and will create defects or damage capable of enhancing grain

growth. Note that the observed abnormal grain growth (Fig. 10) could be explained by a higher defect concentration around big grains. The evolution of the defect concentration D is governed by (6)

1) = k I w - k2D

where kl and k 2 are temperature-dependent variables. Finally, we assume that grain growth rate is given by G = Gsa+/~D

(7)

where ,u is a temperature-dependent variable, and Gsa is the growth rate during pure static annealing. One can rewrite eqn. (6) using eqn. (7)

1)= k, w -

(G_ G a)

(s)

Assuming a given geometrical mechanism for the production and healing of damage, we may consider that k~ and k2/l~ do not depend on temperature. This means that the activation energies of k 2 and/~ are the same. In the case of HIP we have

(9)

w = P '6

P where P is the applied pressure and p the relative density. Using eqns. (6), (7) and (9)we obtain

=

Gsa -

k2

k2

(10)

p

Neglecting Gsa, we obtain, after integration of eqn.(10) f

f

AG=- f dO+P kl f d(lnf,)=Z,+I2P i

(11) where i and f in the integral represent the initial and final states. For a given final density and a sufficiently long time to reduce defect concentration, both integrals in eqn. (11) (I1 and /2) are constant. The final grain size is therefore a linear function of the applied pressure as shown in Fig. 7. Note, however, that delayed effects during damage recovery can be observed, as is clearly shown in Fig. 8. It may also be interesting to use the model in the case of uniaxial creep tests (o=constant,

43

T = constant). The flow stress equation of the dense material can be expressed as geqVP= A exp( - Q/RT )Ocq"G -P

(12)

where Ocq is the Mises equivalent stress, and geqvp the equivalent visco-plastic strain rate. Assuming quasi-steady state, the grain growth rate is given by (~

O ~ a + / t kl

pressure will trigger abnormal grain growth: in both cases, the final grains will be more elongated. Finally, we propose a model of grain growth enhancement during deformation which uses macroscopic variables and which can take into account multiaxial loadings and densification. This model could therefore be used in finite element computations.

k~ Acknowledgments

k, exp [[-R~- Q I ] O~q"+' G -p : 0~, + Akt ~2

(13)

In the case of fine-grained alumina, we have n = 1 and p = 3 (CoNe creep [16] or diffusion-accommodated superplasticity [17]), and grain growth rate during static annealing is given by eqn. (1). With these data eqn. (13) can be rewritten .

=--

6 3

(14)

where C is a constant used to simplify eqn. (13). This result is in good agreement with previously reported results [1, 2] (see above): the stress increases the constant k in eqn. (1) and does not change the grain size exponent m. It has also been recently reported that grain size exponents in eqn. (1) can be characterized by m = 1 during creep and by m = 3 during static annealing [18]; eqn. (12) can also take into account these results (with p=l). 5.

Conclusions

The present results show that grain growth can be enhanced during densification by HIP of a fine-grained ceramic powder; this result can be compared with those obtained during creep tests on alumina [1]. On the other hand, our results concerning grain growth during static annealing are in good agreement with those proposed in the literature. The study of grain shape changes during densification shows that, primarily, particles are rounded. However, sintering for a sufficiently long time will produce anisotropic grain growth, and hot isostatic pressing with a sufficiently high

The authors are indebted to J. L. Koutny for carrying out the HIP experiments and to the Soci&6 Crramiques Techniques Desmarquest for providing the powder. References

1 J.D. Fridez, C. Carry and A. Mocellin, in W. D. Kingery (ed.), Advances in Ceramics, Vol. 10, American Ceramic Society, 1984, p. 720. 2 C. Carry and A. Mocellin, in B. Baudelet and M. Surry (eds.), Superplasticity, Editions du CNRS, 1985, p. 16.1. 3 M. A. Clark and T. H. Alden, Acta Metall., 21 (1973) 1195.

4 D. S. Wilkinson and C. H. Caceres, Acta Metall., 32 (1984) 1335. 5 A. S. Helle, K. E. Easterling and M. E Ashby, Acta Metall., 33 (1985) 2163. 6 M. Abouaf, J. L. Chenot, G. Raisson and P. Bauduin, Int. J. Numer. Methods Eng., 25 (1988) 213. 7 J. Besson and M. Abouaf, Proc. Int. Conf. on HIP of Materials, Antwerp, 25-27April, 1988, p. 1.17. 8 R. E. Miles, Math. Biosci., 6 (1970) 85. 9 E. E. Underwood, in R. T. DeHoff and E N. Rhines (eds.), Quantitative Microscopy, McGraw-Hill, New York, 1968, p. 149. 10 R. J. Brook, in F. F. Y. Wang (ed.), Ceramic Fabrication Processes, Vol. 9, Academic Press, New York, 1976, p. 331. 11 S. Hori, R. Kurita, M. Yoshimura and S. Somiya, J. Mater. Sci. Lett., 4 (1985) 1067. 12 J. Zhao and M. P. Harmer, J. Am. Ceram. Soc., 70 (1987) 860. 13 M. Blanc, A. Mocellin and J. L. Strudel, J. Am. Ceram. Soc., 60 (1977) 403. 14 H. Hahn and H. Gleiter, Scr. Metall., 13 (1979) 3. 15 M. Coster and J. L. Chermant, PrOcis d'Analyse d'lmage, Editions du CNRS, pp. 194,298. 16 R.L. Coble, J. Appl. Phys., 34 (1963) 1679. 17 M. E Ashby and R. A. Verall, Acta Metall., 21 (1973) 149. 18 R.S. Mishra and G. S. Murty, J. Mater. Sci. Lett., 7(1988) 185.