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Effects of Heating Rate on Densification and Grain Growth during Field-Assisted Sintering of ␣-Al2O3 and MoSi2 Powders L.A. STANCIU, V.Y. KODASH, and J.R. GROZA Two types of powders, electrically conductive MoSi2 and insulating ␣-Al2O3, were sintered by a fieldassisted sintering technique (FAST) using heating rates from 50 ⬚C to 700 ⬚C/min. The Al2O3 powders were sintered to 99 pct density at 1100 ⬚C for 2 minutes under 45 MPa pressure. For Al2O3, no exaggerated grain growth was observed and the final grain size inversely scaled with the heating rate. Such a grain growth behavior fits the literature models based on multiple transport mechanisms for constant-heating-rate sintering. The density reached by MoSi2 under similar sintering conditions was 91 pct. The grain size was independent of the heating-rate value. Specific electrical field and pressure effects are shown to contribute to enhanced densification and minimal coarsening in each material.

I. INTRODUCTION

IN ceramic powder processing, high heating rates are often used to enhance sintering and reduce grain growth. The essence of this concept is that fast heating rates favor densification instead of grain growth in coarsening sensitive systems (i.e., with higher activation energy for densification than for grain growth).[1–4] In these systems, a slow heating rate results in grain coarsening by surface diffusion, which is able to operate through low temperature ranges. Conversely, a high heating rate takes the system to high temperatures where a high densification rate is favored before surface diffusion causes grain growth. For instance, a high-heatingrate sintering was showed to be an effective means to reach high density and control grain size in ceramics such as ␣Al2O3 and ZnO.[3–6] The field-activated sintering technique (FAST) is a newly developed sintering method, in which an external electrical field is applied to enhance densification. The method applies pulsed direct current (DC) to powders subject to a modest pressure (⬍100 MPa). Electrical current application enables fast heating rates (up to 700 ⬚C/min). The sintering cycles are very short, typically less than 10 minutes for full densification of both conductive and nonconductive powders. A more detailed description of the process may be found in Reference 7. It is postulated that electrical current application creates favorable conditions for the removal of impurities and the activation of powder particle surfaces. This activation explains the high densities obtained in ceramics without additives and direct grain-to-grain contact at atom scale observed by high resolution electron microscopy (HREM) in ceramics and metals.[8,9,10] Usually, FAST-sintered materials are characterized by high densities and fine grain sizes. For instance, sintering to full density of AlN at 2000 K for 5 minutes maintained the initial submicron grain size (final 0.77 from initial 0.44 ␮m).[8] Fine grain sizes have been typically retained in FAST-sintered nanosize powders. In L.A. STANCIU, Graduate Research Assistant, V.Y. KODASH, formerly Postdoctoral Research Fellow, now Postdoctoral Fellow, and J.R. GROZA, Professor, are with the Chemical Engineering/Materials Science Department, University of California Davis, Davis, CA 95616. Manuscript submitted February 13, 2001. METALLURGICAL AND MATERIALS TRANSACTIONS A

TiN nanopowders, coarsening during FAST sintering resulted in a final grain size that was at least 1 order of magnitude less than in the conventional sintering of the same powders.[9] The limited coarsening observed in FAST sintering is attributed to a very short time at high temperatures and also reflects minimal coarsening during the initial and intermediate stages of sintering. Generally, significant densification with little coarsening was noticed in specimens during the heating phase. For instance, an enhancement of the sintering behavior was observed in WC-Co powders during the heating stage before melting for the typical liquidphase sintering. The use of FAST sintering in solid state (i.e., premelting) to 1525 K induced 98 pct density vs the 70 pct that resulted from the conventional sintering of the same powders.[11,12] This enhanced low-temperature sintering was collectively attributed to the surface activation and faster heating rate in FAST as compared to those in conventional sintering. The purpose of the present work is to study the heatingrate effects in the FAST sintering of two powders: Al2O3 and MoSi2. These materials exhibit different electrical conductivity behavior: Al2O3 ceramic is an insulator, while MoSi2 has more metallic bonding and, therefore, is electrically conductive.[13,14] Both materials are characterized by extensive grain growth during conventional sintering.[3,15] Alumina is often characterized by abnormal or exaggerated grain growth.[3] The choice of different electrical behaviors enables a distinction between physical processes of mass and charge transport both of which contribute to FAST densification and grain coarsening.

II. EXPERIMENTAL Both ␣-Al2O3 powder (Sumitomo Chemical Co., Japan) with a particle size of 0.1 ␮m and MoSi2 powder (Aldrich Co., Milwaukee, WI) with a particle size of 2 ␮m were used. The actual chemistry (i.e., impurities given by the supplier) of pure ␣-Al2O3 powder is shown in Table I. The commercial MoSi2 powder was 99 pct pure, with only traces of impurities present. The FAST experiments were performed using a Spark VOLUME 32A, OCTOBER 2001—2633

Table I. Chemical Composition of ␣-Al2O3 Powder Element

Fe

Si

Mg

Cu

Na

Content (ppm)

9

150

10

⬍1

8

Table II. Densities and Grain Sizes of FAST Sintered Al2O3 and MoSi2 as a Function of Heating Rate Green Density Final Density (g/cm3) (g/cm3) Heating Rate (⬚C/min) Al2O3 MoSi2 Al2O3 MoSi2 50 250 700

0.52 0.51 0.52

0.6 0.59 0.59

3.95 3.95 3.93

5.73 5.73 5.76

Grain Size (␮m) Al2O3

MoSi2

6 to 9 2 to 4 3 to 4 2 to 4 0.5 to 0.6 2 to 4

Plasma Sintering machine (Sumitomo, Japan). Loose powders were filled in graphite die-and-punch units with no additives or lubricants. A vacuum of 3 Pa (22 mtorr) and uniaxial pressure of 45 MPa were used for all experiments. This way, the initial green compact densities were the same for all heating-rate runs (Table II). The powders were processed at heating rates of 50 ⬚C, 250 ⬚C, and 700 ⬚C/min up to 1100 ⬚C, and dwelled at this temperature for 2 minutes. The final compacts were approximately 19 mm in diameter and 4 to 6 mm in height. Sample densities were measured using Archimedes’ method. The error in density measurement is less than 1 pct. The X-ray diffraction (XRD) analysis was carried out using a Scintag X-ray Diffractometer XDS-2000 with monochromatic Cu K␣ radiation. The microstructures of freshly fractured FAST-sintered samples were studied using an ISI DS130 SEM microscope at a voltage of 10 kV. The densification maps for Al2O3 and MoSi2 powders were calculated using hot isostatic pressing 6.1 software developed by Ashby.[16] Data for MoSi2 were taken from References 17 and 18. III. RESULTS Figure 1 shows typical FAST densification and densification rate curves for Al2O3 and MoSi2. The densification of Al2O3 starts at 900 ⬚C and reaches a maximum rate at 950 ⬚C (Figure 1(a)). It is noteworthy that sample densification is already completed on reaching the dwell temperature of 1100 ⬚C (i.e., no more shrinkage is observed beyond the temperature ramp). The densification of MoSi2 starts at 660 ⬚C (Figure 1(b)). Its rate gradually increases with the increase in temperature and reaches a maximum at ⬃950 ⬚C. The final densities and grain sizes of Al2O3 and MoSi2 are shown in Table II. As expected from the shrinkage evolution, the density of Al2O3 was close to theoretical (99 pct). The highest density of molybdenum disilicide was 91 pct, calculated based on 100 pct MoSi2 content. However, XRD analysis showed traces of a lower silicide, Mo5Si3 (Figure 2(b)). The density calculation showed that Mo5Si3 traces do not significantly change the relative density. No Mo5Si3 was present in the MoSi2 starting powders (Figure 2(a)). Most probably, Mo5Si3 forms during FAST sintering. In any case, MoSi2 decomposes in vacuum at 1400 ⬚C, according to the reaction 2634—VOLUME 32A, OCTOBER 2001

(a)

(b) Fig. 1—Shrinkage (1) and shrinkage rate (2) vs temperature during FAST sintering of (a) Al2O3 and (b) MoSi2 (heating rate 50 ⬚C/min).

MoSi2 → Mo5Si3 ⫹ Si In the initial stages of sintering, the electrical current flowing through the electroconductive MoSi2 generates Joule heating at contact points where the electrical resistance is high. Such local heating may generate temperatures in excess of 1400 ⬚C, thereby causing MoSi2 to decompose into Mo5Si3. The Si may partly evaporate from the sample and partly react with Mo5Si3, forming secondary MoSi2 when the temperature in contacting areas decreases and necks between particles are formed. The XRD analysis did not reveal any Si in the samples, but it may be present in a quantity smaller than the XRD detection limit (⬃1 to 2 pct). The microstructures of FAST-sintered samples are shown in Figures 3 and 4. Both Al2O3 and MoSi2 underwent intergranular fracture by bending revealing grain contours and pores randomly scattered on the faces of the grains and at grain boundaries. The Al2O3 contains small pores in all samples, regardless of the heating-rate value. No exaggerated or abnormal grains were observed. The grain size remained in the submicron regime for the fastest heating rate in Al2O3. At 700 ⬚C/min, the grain size was 5 to 6 times larger than in the initial powders. Grain size was 30 to 40 times larger than the initial grain size at 250 ⬚C/min and 60 to 90 times larger at 50 ⬚C/min. The final grain size in the slowest heating-rate experiment is 6 to 9 ␮m. For comparison, the rapid sintering of pure Al2O3 of a similar initial METALLURGICAL AND MATERIALS TRANSACTIONS A

(a) (a)

(b) (b) Fig. 2—XRD patterns of (a) MoSi2 starting powders and (b) FAST sintered specimen.

size using conventional sintering at 1850 ⬚C for 10 minutes resulted in both a grain size of ⬃30 ␮m and abnormal grain growth.[3] The grain size in MoSi2 was found to be independent of the heating rate (Table II). Pore sizes do not considerably change with the heating rate (Figure 5). A similar pore size gradient (i.e., coexisting large and small pores) is noticeable at each heating rate. IV. DISCUSSION Fast heating rates of ceramics are used to achieve sintering with controlled grain growth. To restrict grain growth, sintering is performed in conditions under which mass transport mechanisms leading to densification are favored over those leading to grain boundary migration. However, even for conventional sintering, it is difficult to identify clearly the mechanisms responsible for the two competing events: densification and coarsening. The problem is compounded for constant-heating-rate experiments. For instance, a careful analysis of the constant-heating-rate densification of alumina indicated that grain boundary diffusion cannot be neglected in the early sintering stages.[19] Its contribution must also be accounted for along with that of the more conventional METALLURGICAL AND MATERIALS TRANSACTIONS A

(c) Fig. 3—SEM micrographs of Al2O3 specimens FAST sintered at 1100 ⬚C for 2 min with heating rates of (a) 50 ⬚C/min, (b) 250 ⬚C/min, and (c) 700 ⬚C/min.

contribution of surface diffusion. Considering that multiple transport mechanisms are operating, Chu et al.[8] developed a model of constant-heating-rate sintering for the entire densification process. The model describes the interplay between densification and grain growth. When the interpore spacing is proportional to grain size, the coarsening function G(T, t) may be described by Gm ⫽ Gm0 ⫹ g(T )t [1] VOLUME 32A, OCTOBER 2001—2635

(a) Fig. 5—Final grain size of FAST-sintered Al2O3 as a function of heating rate.

i.e., for constant heating rate (␣), where the temperature is T ⫽ T0 ⫹ ␣t, Eq. [1] becomes[5] Gm(T ) ⫽ Gm0 ⫹

(b)

(c) Fig. 4—SEM micrographs of MoSi2 samples FAST sintered at 1100 ⬚C, with heating rates of (a) 50 ⬚C/min, (b) 250 ⬚C/min, and (c) 700 ⬚C/min.

where G0 is the starting grain size at time t ⫽ t0, and m is a constant. The value g(T ) is a coarsening function, which has a general Arrhenius behavior given by g(T ) ⫽ g0 exp (⫺Qg /kT )

[2]

where g0 is a material constant, Qg is the activation energy for grain growth, k is the Boltzmann constant, and T is the absolute temperature. When temperature changes with time, 2636—VOLUME 32A, OCTOBER 2001

g0 ␣

兰

冢 冣

Qg exp ⫺ dT T0 kT T

[3]

The experimental grain size values for Al2O3 obey the law described by Eq. [3] and the final grain size is inversely proportional to the heating rate ␣ (Table II). A plot of grain size vs the inverse of the heating rate shows a best fit for an exponent value of m ⫽ 3 (Figure 5). A similar cube law was reported for some other ceramics, which display the same coarsening-to-densification relationship (e.g., ZnO).[5] The external electrical field does not seem to alter this behavior. The main assumption for deriving Eq. [3], interpore spacing equal to grain size, may hold in the FAST-sintered alumina. The microstructure of the present alumina samples shows pores located predominantly at grain boundaries (Figure 3). Therefore, grain-boundary mobility may be influenced mostly by porosity. Moreover, no additives to alter grain-boundary mobility were used in the present experiments. This alumina is either reasonably pure (Table I) or the electrical field application may have a cleaning effect on the powder, thus reducing the impurity influence on the mobility of grain boundaries. Since the pores are small and only a few are left after FAST sintering (Figure 3), they cause less drag on grain boundaries. In other words, the rapid reduction of porosity by activated grain-boundary and volume diffusion in FAST sintering gives rise to high grainboundary mobility. This high mobility makes the sintering of pure powders sensitive to heating rate; therefore, a rapid heating is essential to prevent grain growth. Indeed, Wang et al.[20] showed that 50 ⬚C/min is the minimum heating rate needed to avoid coarsening predominance in the field sintering of Al2O3. The present results at heating rates equal to or exceeding this minimum value compare well with those obtained by Wang et al. Field-charge transport along the Al2O3 particle surface, as well as at the grain boundary, may accelerate surface- and grain-boundary diffusion. At high temperatures, when the electrical resistivity of alumina decreases, volume diffusion may also be enhanced, due to field activation. As a collective result, the FAST sintering METALLURGICAL AND MATERIALS TRANSACTIONS A

of Al2O3 reaches densities close to theoretical before reaching the dwelling temperature of 1100 ⬚C. For comparison, conventional sintering with rapid heating of undoped alumina of similar submicron size resulted in 94.7 pct density at 1850 ⬚C and 10 min of dwelling.[3] Although no quantitative description of the field effects on the specific diffusivities is possible at this time, it seems that the favorable ratio of activation energies for volume-vs-surface diffusion remains the same as in conventional heating. Therefore, rapid heating in FAST sintering results in a slower kinetics of the grain growth in pure alumina. The grain size of FAST-sintered MoSi2 does not show any dependence on heating rate. The observed coarsening was minimal and the final grain size in FAST-sintered MoSi2 is 3 to 4 ␮m. By comparison, a grain size of 20 ␮m at a density of 94 pct was obtained by the conventional sintering at 1700 ⬚C of a powder with a similar initial grain size (1 ␮m).[15] Despite the unfavorable particle size (i.e., 1 order of magnitude larger than that of Al2O3), the densification of electrically conductive MoSi2 powders starts at a lower temperature (660 ⬚C), as compared to nonconductive Al2O3 (900 ⬚C) (Figure 1). This may be accounted for by two main phenomena in electroconductive powders: arcing and constriction resistance.[21] First, arcing or electrical discharges at contact areas are responsible for surface “cleaning” but, at the same time, may enhance surface transport by surface diffusion and/or vapor transport. The evaporation of cobalt was usually noticed in the FAST sintering of WC-Co alloys.[11] Second, a high constriction resistance is created at the small particle contact areas. The higher resistance of these contacts generates local Joule heating. This way, the temperature of the ductile brittle transition in MoSi2 may be reached and local yielding or creep may be initiated while the temperature of the sample is still low. Possible local high temperatures are indicated by the presence of small amounts of Mo5Si3 (Figure 2(b)). Therefore, since pressure-assisted consolidation is facilitated, no grain growth is expected. All of the preceding events enhance neck formation and growth in the early sintering stages. In the intermediate stages of sintering, a charge gradient is developed in the vicinity of pores of different sizes. Similar to initial particle contacts, the electrical field increases as the concentration of equipotential lines increases (Figure 6). The electrical current density is higher next to large pores than next to small ones. This creates a temperature gradient, i.e., the temperature is higher next to large pores than next to small ones. Raichenko[22] calculated the temperature gradient, ⵜT, which develops in the vicinity of these pores under pulsed-field application: ⵜT ⬇

1 R

冪

␴0 T0 E0 2 ⌬␶ ⭈ 2CM n

[4]

where R is the pore radius, ␴0 the electrical conductivity, CM the specific heat, T0 the initial temperature, Eo the intensity of electric field, ⌬␶ the time of electric field effect, and n the number of electrical impulses. In turn, either this temperature gradient generates a vacancy gradient, ⵜCv , or more vacancies are created in the vicinity of the large pores. The vacancy flow J is given by[22] J ⫽ Dv (kT /T ⵜT ⫺ ⵜCv) METALLURGICAL AND MATERIALS TRANSACTIONS A

[5]

Fig. 6—Schematic of electrical field density and temperature gradients in the vicinity of small and large pores.[22]

where Dv is the diffusion coefficient of vacancies and kT is the thermal diffusivity. Therefore, vacancy diffusion occurs from the large toward the small pores; the final result is that the large pores shrink. This is the opposite of conventional sintering, in which a large pore grows at the expense of small pores. The microstructure of fast-sintered MoSi2 indicates a nonequilibrium pore structure (i.e., large and small pores (Figure 4)) with large pores characterized by convex surfaces. Such surfaces point to large-pore shrinkage or dissolution (i.e., moving toward the center of curvature). The minimal coarsening observed in MoSi2 is indicative of the negligible effect of surface transport or nondensifying mechanisms in the early sintering stages at the heating rates used in this work. The possible pressure contribution to densification in the early stages also results in little or no coarsening. This contribution is independent of heating rate. The relatively large concentration of pores (with a size up to 1 ␮m (Figure 4)) on the grain boundaries considerably decreases the mobility of grain boundaries in the late sintering stages, as well. The final fine grain size may be, at least partially, the result of the pressure effects during FAST sintering. To account for the pressure contribution to densification, Figure 7 shows the calculated densification maps for Al2O3 and MoSi2. These maps were built using the same combination of temperature and pressure as is used in FAST experiments, and the same initial particle size as is found in the experimental powders (i.e., 0.1 ␮m for Al2O3 and 2 ␮m for MoSi2). As expected, pressure-driven densification by power-law creep is predominant and depends on the applied pressure in MoSi2. This pressure contribution may also explain the densification onset at temperatures as low as 600 ⬚C. The VOLUME 32A, OCTOBER 2001—2637

gives rise to high grain-boundary mobility that makes the sintering of fine pure Al2O3 powders sensitive to the heating rate during FAST. In Al2O3, the grain size is inversely related to the heating rate. Grain growth behavior fits the literature models based on multiple transport mechanisms for constant-heating-rate sintering. In MoSi2, no dependence of final grain size on heating rate was observed. FAST sintering of conductive MoSi2 enables rapid resistance heating, vacancy generation to shrink large pores, and pressure application to restrict grain growth. ACKNOWLEDGMENTS The authors are grateful to the NSF support of this work (DMI 9978699) and partial ARO support. (a) REFERENCES

(b) Fig. 7—Densification maps of (a) Al2O3 powder with a particle size of 0.1 ␮m and (b) MoSi2 powder with a particle size of 2 ␮m.

low-sintering onset temperature may be also a macroscopic result of the local heating and deformation effects at the interparticle contacts due to field application. In contrast, Al2O3 densification occurs primarily because of surfacetension-driven diffusion along interpaticle boundaries, with little influence from the applied pressure up to the level used in FAST sintering (45 MPa) (Figure 7a). V. CONCLUSIONS Field-activated mass transport during FAST sintering enhances the densification rate of both insulating Al2O3 and conductive MoSi2. Submicron Al2O3 was densified to neartheoretical density. A final grain size of 0.5 to 0.6 ␮m was obtained when the heating rate was 700 ⬚C/min. The final density in MoSi2 was 91 pct; grain size was 2 to 4 ␮m. No exaggerated or abnormal grains were observed in either material. Rapid reduction of porosity by activated sintering

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1. D.L. Johnson: in Advanced Ceramics II, S. Somiya, ed., Elsevier Applied Science, London, 1987, pp. 1-6. 2. R.L. Coble and T.K. Gupta: in Sintering and Related Phenomena, G.C. Kuczynski, N.A. Hooton, and C.F. Gibbon, eds., Gordon Breach Science Publisher Inc., New York, NY, 1967, pp. 423-44. 3. M. Harmer, W.W. Roberts, and R.J. Brook: Trans J. Br. Ceram. Soc., 1979, vol. 78, pp. 22-25. 4. R.J. Brook: Proc. Br. Ceram. Soc, 1982, vol. 32, pp. 7-24. 5. M.Y. Chu, M.N. Rahaman, and L.C. De Jonghe: J. Am. Ceram. Soc., 1991, vol. 74, pp. 1217-25. 6. D. Dadon, L.P. Martin, M. Rosen, A. Birman, D. Gershon, J.P. Calame, B. Levush, and Y. Carmel: J. Mater. Synthesis Processing, vol. 4, 1996, pp. 95-103. 7. J.R. Groza: ASM Handbook, [1998, vol. 7,] Powder Metal Technology and Applications, Metals Park, OH, pp. 583-87. 8. J.R. Groza, S.H. Risbud, and K. Yamazaki: J. Mater. Res., 1992, vol. 7, pp. 2643-45. 9. J.R. Groza, J.D. Curtis, and M. Kra¨mer: J. Am. Ceram. Soc., 2000, vol. 83 (5), pp. 1281-83. 10. K.R. Anderson, J.R. Groza, M. Fendorf, and C.J. Echer: Mater. Sci. Eng., 1999, vol. A270, pp. 278-82. 11. J.R. Groza and J. Oakes: Adv. Powder Metall. Particulate Mater., 1997, vol. 12, pp. 12-51. 12. J.M. Doh, J.R. Groza, and J. Oakes: Adv. Powder Metall. Particulate Mater., 1998, vol. 1, pp. 105-14. 13. B. Bhattacharyya, D.M. Bylander, and L. Kleineman: Phys. Rev. B, 1985, vol. 32, pp. 7973-78. 14. V.S. Neshpor and G.V. Samsonov: Poroshkovaya Metallurgiya, 1962, vol. 6, pp. 14-19. 15. P.S. Kisly, M.A. Kyzenkova, V.G. Kayuk, and O.V. Pshenichnaya: in High Temperature Borides and Silicides, T.Y. Kosolapova, ed., Naukova Dumka, Kiev, 1978, pp. 104-08. 16. M.F. Ashby: HIP 6.1 Software for Constructing Maps for Sintering and Hot Isostatic Pressing, Cambridge University, Cambridge, United Kingdom, 1990. 17. R. Suryanarayanana and S.M.L. Sastry: Acta Mater., 1999, vol. 47, pp. 3079-98. 18. R. Suryanarayanana, S.M.L. Sastry, and K.L. Jerina: Acta Mater., 1994, vol. 42, pp. 3741-50. 19. S.H. Hillman and R.M. German: J. Mater. Sci., 1992, vol. 27, pp. 2641-48. 20. S.W. Wang, L.D. Chen, and T. Hirai: J. Mater. Res., 2000, vol. 15, pp. 982-87. 21. J.R. Groza and A. Zavaliangos: Mater. Sci. Eng., 2000, vol. A287, pp. 171-77. 22. A.I. Raichenko: “Fundamental Processes in Powder Sintering,” Metalurgyia, Moscow, 1987, pp. 11-14.

METALLURGICAL AND MATERIALS TRANSACTIONS A