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The microstructure of pressureless sintered silver-toughened alumina: an in situ TEM study
Materials Science and Engineering, A161 (1993) 119-126
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The microstructure of pressureless sintered silver-toughened alumina: an in situ TEM study J. Wang, C. B. Ponton and P. M. Marquis IRC in Materialsfor High PerformanceApplications & School of Metallurgy and Materials, The Universityof Birmingham, Birmingham B15 277"(UK) (Received July 27, 1992; in revised form September 30, 1992)
Abstract An in situ transmission electron microscopy (TEM) study has been made of the microstructure of a pressureless sintered alumina-silver composite, which consists of both intergranularly and intragranularly positioned silver inclusions in the alumina matrix. The silver grains at the grain boundaries and junctions of the alumina matrix are 1-4/~m in size and angular in morphology. Residual thermal strains result in the formation of either the well-established dislocation rings or deformation twins or both in these intergranularly positioned silver particles. The silver particles entrapped inside alumina grains are 0.1-2/~m in size and rounded in morphology. Deformation twins were not found in these intragranularly positioned silver particles. The weak interfacial bonding between silver inclusions and the alumina matrix is characterized by the occurrence of voids or pores at the interface. Thermal mismatch-induced microcracks were not observed at the interface.
I. Introduction
The proprietary studies on strengthening and/or toughening brittle matrices using a metallic second phase, mainly for cutting tool applications, dates back more than 30 years [1]. These earlier investigations did not, however, lead to a well-established understanding of the microstructure-property relationships in these ceramic-metal composites, partly due to the lack of certain appropriate experimental techniques such as transmission electron microscopy (TEM). There has been a new upsurge of interest in ceramic matrices toughened using a metallic second phase in the advanced engineering ceramic community since the early 1980s, and a class of metallic-phase toughened ceramic materials have been fabricated and microstructurally characterized, including lanxide aluminium-toughened alumina [2, 3], nickel-toughened alumina and mullite [4], and, more recently, silvertoughened alumina [5, 6]. A fracture toughness of 7-8 MPam °-5 is obtainable in these metallic-phase-dispersed ceramic matrices via proper compositional and microstructural control, compared with a fracture toughness of 2.5-3.0 MPam °5 for the monolithic ceramic matrices. The presence of a metallic inclusion in a ceramic matrix results in crack bridging by the metallic phase, and crack deflection at the interface between the metallic phase and the ceramic when the composite material is fractured [2, 7, 8]. 0921-5093/93/$6.00
The fabrication of silver-toughened ceramic matrices, which has been carried out using a conventional ceramic processing route, e.g. pressureless sintering of mixed alumina and silver oxide powder compacts, is both simple and economic [5, 6]. During the sintering of an alumina-silver oxide powder compact, a series of physical and chemical changes take place with increasing heating temperature. For example, the silver oxide particles undergo decomposition in the temperature range 400-450 °C, forming metallic silver inclusions in the alumina matrix. The metallic silver particles then melt at 962 °C [9]. It is apparent that the densification of silver-toughened alumina involves a molten silver phase at the grain boundaries and junctions of the alumina matrix. As has been reported [10], the interfacial energy between alumina and molten silver liquid is considerably high and little reaction occurs at the interface. Therefore, the densification mechanisms at the sintering temperature in silver-toughened alumina ceramics are different from those in conventional liquid phase-assisted ceramic systems [11]. Molten silver liquid exhibits high vapour pressure above its melting temperature. For example, a partial pressure of 6.6 mm Hg was measured for molten silver at 1500°C [12]. The high vapour pressure of molten silver liquid results in a large loss in silver content at the sintering temperature, and a silver-free aluminasintered surface, the thickness of which varies with the © 1993 - Elsevier Sequoia. All rights reserved
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composition and sintering temperature, was observed in each sintered alumina-silver composite [6]. On cooling from the sintering temperature, the molten silver inclusions undergo crystallization at temperatures around 962°C [9]. The thermal mismatchinduced residual strains may play an important role in the microstructure-mechanical property relationships in the alumina-silver composite. In comparison with other metallic-phase-toughened ceramics, little investigation has been made into the microstructural characterization of silver-toughened alumina ceramics, although the potentially high fracture toughness has been demonstrated in these materials. The aim of the present work is to carry out a microstructural study on a pressureless sintered alumina-silver composite using in situ TEM observation.
Fig. 1. A silver-free-aluminasintered surface forms in each sintered specimen pellet due to the silver loss resultingfrom the high vapour pressure of molten silver liquid at the sintering temperature.
2. Experimental work The as-received alumina powder (average particle size: 0.75/zm, from BA Chemical Ltd., England) and the as-received silver oxide powder (average particle size: 0.2/tm, from BDH Ltd., Poole, England) were mixed together by ball milling in propanol, using zirconia balls as the milling medium, for 6 h to obtain a composition of AI203 + 40 wt.% mg20. The ball-milled mix was then dried using a combination of an infrared heating lamp and hot plate, followed by compaction in a steel die of 12.5 mm in diameter at a pressure of 120 MPa. The powder compacts were sintered in an electric furnace at 1600 °C for 2 h using heating and cooling rates of 4 °C min- l. As mentioned above, as a result of the high vapour pressure of molten silver liquid at the sintering temperature, the average silver content in the as-sintered sample is lower than that in the pre-sintered powder compact [6]. Furthermore, a silver-free aluminasintered surface is formed in each sintered pellet, see Fig. 1. It was shown, using XRD phase analysis, that the as-sintered samples consisted of alumina and metallic silver. The average silver content in the assintered sample was worked out to be 10.1 vol.% on the basis of total weight loss and the observed thickness of the silver-free alumina-sintered surface. The assintered alumina-silver composite exhibited a sintered density of 98.5% theoretical density, as measured using the Archimedes technique in distilled water, and an indentation toughness of 8.4 MPam °5 (indentation load: 196 N; the equation proposed by Anstis et al. [13] was used for calculation). The average grain size of the alumina matrix is 5.5/zm, as measured using the linear intercept method when more than 150 grains were counted on the polished and then etched surface. To prepare specimens for in situ TEM study, a slice
of 300 # m in thickness was cut from the as-sintered ceramic using a diamond slicing wheel, followed by grinding on a 400 grit SiC paper to reduce the thickness of the slice to approximately 150 gm. Discs of 3 mm in diameter were obtained from the slice using a diamond disc cutter. The specimen discs were further polished and then dimpled, followed by ion beam thinning. The in situ TEM microstructural study was carried out on the ion beam thinned foils using Joel4000CX at 400 kV.
3. Results and discussion Figure 2(a-c) has three bright field transmission electron micrographs showing the general view of the microstructure for the alumina-silver composite fabricated in the present work. It consists of an alumina matrix of 3-10/~m in grain size and silver inclusions of 1-4/~m in size, both at the grain boundaries and junctions of the alumina matrix. Relatively smaller silver particles, in the range of 0.1-2/zm in size, were observed to occur within certain alumina grains. The silver inclusions exhibit considerably different morphologies with respect to their sizes and locations in the alumina matrix: (a) the relatively larger silver grains at the grain junctions of the alumina matrix exhibit an elongated morphology along the grain boundaries; (b) the intermediate-sized silver phase at the grain boundaries of the alumina matrix exhibit a near-ellipsoidal morphology, Fig. 3(a, b); and (c) the relatively smaller silver particles entrapped inside alumina grains are near-spherical in morphology.
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Fig. 2. Three bright field transmission electron micrographs showing the microstructure of the alumina-silver composite fabricated in the present work. The silver grains at the grain boundaries and junctions of the alumina matrix tend to be angular in morphology, and those, or sections of those, entrapped by the aluminagrains are rounded in morphology.
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Voids ol pores often exist at the interface between the silver grains and the alumina matrix. At the sintering temperature, the microstructural development in the alumina-molten silver two-phase system involves two concurrent processes: the interaction between the alumina grain boundaries/junctions and molten silver inclusions, and the ripening of the molten silver inclusions in the alumina matrix. The occurrence of both the intergranularly and intragranulaxly positioned silver particles in the alumina matrix is a direct result of such an interaction between the moving alumina grain boundaries and the molten silver inclusions at the sintering temperature. The contact angle and interfacial energy between alumina and molten silver liquid are high (150 ° and 900 ergs cm-3 respectively) and little reaction takes place between the two phases [10]. Similar to a two-solid phase ceramic system [14, 15], the molten silver inclusions at the grain boundaries of the alumina matrix apply a dragging or pinning force to the mobile grain boundaries, although they are viscously deformable at the sintering temperature. The pinning effect of a silver inclusion to a mobile alumina grain boundary will be largely dependent on its size. Relatively smaller silver inclusions exert a smaller pinning force on the grain boundaries than the relatively larger ones do. Therefore, these small silver inclusions are more likely to be swallowed up by the mobile grain boundaries and therefore more likely to be entrapped within alumina grains than the relatively larger ones. The difference in morphology between the intergranularly and intragranularly positioned silver grains is due to the deformability and high surface tension of the molten silver liquid at the sintering temperature. Those silver inclusions at the grain junctions of the alumina matrix are subjected to a highly anisotropic stress and are likely to be squeezed by the growth of adjacent alumina grains, resulting in the formation of an elongate d morphology along the grain boundaries. Similarly, the near-ellipsoidal morphology of the silver grains at the grain boundaries of the alumina matrix is due to the balance between the pressure from the neighbouring alumina grains and their high surface tension (800-850 mN m -2) at the sintering temperature [9, 12]. Similar to the situation when pores are entrapped by grains of a monolithic ceramic, the relatively smaller silver particles entrapped within alumina grains exhibit a near-spherical morphology in order to minimize the total interfacial free energy [15]. These intragranulaxly positioned silver inclusions are subjected to a more isotropic stress from the alumina matrix than the silver inclusions entrapped at the grain boundaries and grain junctions of the alumina matrix. The formation of voids or pores at the interface between the silver grains and the alumina matrix is
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Fig. 3. The silvergrains at the grain boundaries of the alumina matrix are near-ellipsoidalin morphology.There exist well-established dislocationtings in many silvergrains. closely related to the high vapour pressure of molten silver liquid at the sintering temperature [12]. Owing to this high vapour pressure, on the one hand, Ostwald ripening may be the dominant mechanism responsible for the growth of silver grains in the alumina matrix [16]. On the other hand, a back stress will be associated with the high vapour pressure of the molten silver phase, which will offset a degree of driving force for the densification of the alumina matrix [17]. As has been observed by the present authors [6], the sintering temperature for achieving a sintered density of 98.5% theoretical density in alumina matrices containing more than 5 vol.% silver is more than 100 °C, higher than that needed for achieving a similar sintered density in the monolithic alumina matrix, when a duration of 2 h at the sintering temperature is employed. The final sintered density of the alumina-silver composites was slightly lower than that for the alumina matrix. As mentioned earlier, the decrease in silver content with increased heating temperature is a direct result of the high vapour pressure of the molten silver phase. The disappearance of silver inclusion or a part of silver inclusion will leave a void in the alumina matrix, although further densification may eventually eliminate the void from the structure. The release of vapourized silver from the alumina matrix requires a well-interconnected channel network. It is likely that the grain boundaries will act as such diffusion channels for the vapourized silver phase entrapped at the grain boundaries and junctions of the alumina matrix. It is, however, almost impossible for vapourized silver entrapped inside an alumina grain to be released, as no interconnected channels are available. The entrapped silver vapour may generate an
internal pressure at the interface between the alumina and molten silver phase, limiting the densification attained in this region. This explains the fact that interfacial voids or pores are more likely to occur around silver grains or the sections of silver grains entrapped by the alumina grains. These grains are rounded in morphology. The existence of these interracial voids will assist the occurrence of crack deflection, which was observed to be the dominant toughening mechanism in the silver-toughened alumina ceramics. It is anticipated that the thermal mismatch between the metallic phase and the oxide phase is an important microstructural feature for the silver-dispersed alumina matrix [18, 19]. On cooling from 1600 °C (the sintering temperature) to 962 °C (the melting point of metallic silver), the thermal mismatch between alumina and molten silver liquid may not be important in forming residual strains in silver-toughened alumina, as any thermal stress will be constantly accommodated by the viscously deformable molten silver phase. The linear thermal expansion coefficient of metallic silver ((18-23)x 10 -6 C -1) is almost three times of that for alumina ceramics ((6-9)x 10 -6 C-1) o v e r the temperature range 25-962°C. On cooling at temperatures below 962 °C, a thermal stress will be generated in the composite structure. Owing to thermal mismatch, microcracks occur at the interface in many ceramic matrix composite systems [20]. As will be discussed later, the thermal stress did not lead to the formation of microcracks at the interface between silver inclusions and the alumina matrix in the as-sintered material. The distribution of the thermal mismatch-induced residual strain energy is dependent on the elastic properties of the component materials. The alumina grains exhibit a
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Fig. 4. Three bright field transmission electron micrographsand the associated SADP showingthe occurrence of (a) dislocation rings, (b) partial twinningand (c) complete twinning, of the silver particles at the grain junctions of the alumina matrix. Microcrackswere not observed at the interface between the silvergrains and the aluminamatrix. higher elastic modulus (380 GPa) than the metallic silver ( 105 GPa). Under a given stressing condition, the former will undergo a lower degree of deformation than the latter. As is shown in Figs. 2(a-c) and 3(a-c), few strain contours and dislocation networks were observed to occur in the alumina grains, indicating that these rigid alumina grains accommodate a very small percentage of the total residual strain caused as a result of the thermal mismatch. It has been well established that stress-induced dislocations and deformation twinning occur in metallic silver under an appropriate strain condition [21, 22]. Specifically, plane {111} and direction (112) are the twinning elements for the face-centred cubic silver [23]. It was observed that most silver grains entrapped at the grain boundaries and junctions of the alumina matrix exhibit either a well-defined dislocation network or deformation twins or both, see Figs. 3(b,c), 4(a-c), 5(a,b) and 6(a,b). Microcracks were not found at the interface between the silver inclusions and the alumina matrix, regardless of the difference in particle size and morphology of various silver grains observed. Figure 4(a-c) consists of three bright field TEM micrographs and the associated diffraction patterns showing the occurrence of dislocation rings, partial twinning and complete twinning, respectively, in silver grains entrapped at the grain junctions of the alumina matrix. Selected area electron diffraction confirmed that deformation twinning occurs on the {111} planes. Correlation was not found between the orientation of alumina grains and the twinning direction of silver inclusions, as the thermal stress at a grain boundary or junction of the alumina matrix is a highly anisotropic and complex
process. As an example, Fig. 6(a,b) comprises two bright field transmission electron micrographs and the associated selected area diffraction pattern (SADP) showing that the deformation twins in silver grains can be parallel and perpendicular to the c-axis of the alumina grains. As discussed earlier, certain small silver grains, which exhibit a rounded morphology, are entrapped inside alumina grains. The straining conditions in these intragranularly positioned silver grains may be different from those in the intergranularly positioned silver grains, in terms of their difference in size and in morphology [18, 19]. Well-established strain contours were observed to occur in almost all these rounded silver particles, regardless of their difference in size and in shape, see Fig. 7(a-d). Dislocation rings were found in the relatively large silver particles entrapped within alumina grains, Fig. 7(a,b). Deformation twins were, however, not found in these intragranularly positioned silver grains when more than 60 grains were observed, although some of these silver grains were highly strained. The level of thermal mismatch-induced residual strain energy in a two-phase material is dependent on parameters such as the differences in thermal and elastic properties of the two component phases, and in the size and shape of the secondary phase in the matrix [18, 19]. The residual strain energy increases with increasing particle size of the secondary phase. An angular inclusion is likely to result in a high strain concentration at the interface. As discussed above, the intragranularly positioned silver grains are relatively smaller in size and more rounded in morphology than
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(a)
Fig. 5. A single twin formed in a silver grain entrapped at the grain junction of the alumina matrix. Image (b) was taken when image (a) was tilted by 14".
Fig. 6. Two bright field TEM micrographs and the associated SADP showing that the deformation twins in silver grains can be parallel and perpendicular to the C-axis of neighbouring alumina grains.
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Fig. 7. Two small silver particles entrapped within alumina grains. The residual thermal stress results in the formation of strain contours and dislocation rings in these relatively small and rounded silver particles. The highly strained structure in (d) is indicated by the slightly distorted SADP. Deformation twins were not found when more than 60 grains were observed.
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those entrapped at the grain boundaries and junctions of the alumina matrix. The thermal mismatch-induced residual strain energy level in these intragranularly positioned silver particles will be lower than in those intergranularly positioned silver grains. A minimum level of residual strain energy is required in order to generate either a dislocation ring or a deformation twin in the face-centred cubic silver. Specifically, the strain energy needed to induce a deformation twin is generally higher than that needed to induce a dislocation ring in metallic silver, if it is assumed that the formation of the deformation twin is via a dislocation mechanism [23]. The thermal mismatch-induced residual strain energy in the intragranularly positioned silver particles may not be high enough to result in the formation of deformation twinning on cooling from the sintering temperature. This explains the absence of deformation twins in these small and rounded silver grains.
4. Summary The pressureless sintered alumina-silver composite consists of both intergranularly positioned silver grains of 1-4 p m in size and intragranularly positioned silver particles of 0.1-2 p m in size in the alumina matrix. The latter are more rounded in morphology than the former. The weak interface between the silver grains and the alumina matrix is characterized by the occurrence of voids or pores around the silver grains or part of the silver grains entrapped by the alumina grains. In situ TEM observation showed that the thermal mismatch-induced residual strains in the alumina-silver composite were largely accommodated by the silver particles, which are much less rigid than the alumina grains. Microcracks were not observed at the interface between silver inclusions and the alumina matrix. The residual thermal stress results in the formation of either the well-established dislocation rings or deformation twins or both in silver particles entrapped at the grain boundaries and junctions of the alumina matrix. Deformation twins were not observed in the intragranularly positioned silver particles, as they are relatively smaller in size and more rounded in morphology than those at the grain boundaries and junctions of the alumina matrix and therefore a limited level of thermal strain is induced at the interface on cooling from the sintering temperature.
Acknowledgments The authors thank Dr. T. J. Doel of the IRC in Materials for High Performance Applications, University of Birmingham, for reading the manuscript of this paper. The financial support given for this project by the SERC is acknowledged.
References 1 A. E. S. White and T. H. Blakeley, Chem. Ind. (London), (1962) 1740. 2 M. K. Aghajanian, N. H. MacMillan, C. R. Kennedy, S. J. Luszcz and R. Roy, J. Mater. Sci., 24 (1989) 658. 3 M. Bannister, H. Shercliff, G. Bao, F. Zok and M. F. Ashby, Acta Metall., 40 (1992) 1531. 4 W. H. Tuan and R. J. Brook, Reactive sintering of alumina/Ni composites, in R. Morrell (ed.), Advanced Engineering with Ceramics, The Institute of Ceramics, Stoke-on-Trent, 1991, p. 335. 5 W.H. Tuan and W. B. Chou, Toughening alumina with silver inclusions, presented at the 2nd European Ceramic Society Conference, Augsburg, Germany, September 11-14, 1991. 6 J. Wang, C. B. Ponton and P. M. Marquis, Silver toughened alumina, Trans. J. Brit. Ceram. Soc. (1992) in press. 7 M. E Ashby, E J. Bunt and M. Bannister, Acta Metall., 37 (1989) 1847. 8 L. S. Sigl, P. A. Mataga, B. J. Dalgleigh, R. M. McMeeking and A. G. Evans, Acta Metall., 36 (1988) 945. 9 D. R. Lide, in CRC Handbook of Chemistry and Physics, CRC Press, Boca Raton, 1992, pp. 4-123. 10 S.P. Mehrotra and A. C. D. Chaklader, Metall. Trans., 16(B) (1985)567. 11 W. J. Huppmann and G. Petzow, The elementary mechanisms of liquid phase sintering, in G. C. Kuczynski (ed.), Sintering Processes, Plenum Press, New York, 1980, p. 189. 12 C. B. Alcock, in CRC Handbook of Chemistry and Physics, CRCF Press, Boca Raton, 1992, pp. 5-70. 13 G. R. Anstis, P. Chantikul, B. R. Lawn and D. B. Marshall, J. Am. Ceram. Soc., 64 (1981) 533. 14 H. Hillert, Acta Metall., 36 (1988) 3177. 15 E F. Lange, J. Am. Ceram. Soc., 67(1984) 164. 16 R.D. Wengrenovitch, Acta Metall., 30 (1982) 1079. 17 L. C. De Jonghe and M. N. Rahaman, Acta Metall., 36
(1988) 223. 18 F. F. Lange, in R. C. Bradt, D. P. H. Hasselman and F. F. Lange (eds.), Fracture Mechanics of Ceramics, Vol. 1, Plenum Press, New York, 1978. 19 R.W. Davidge and D. J. Green, J. Mater. Sci., 3 (1968) 69. 20 P.F. Beclier, J. Am. Ceram. Soc., 74 (1991) 255. 21 R.E. Reed-Hill, J. P. Hirth and H. C. Rogers (eds.), Deformation Twinning. Metall. Soc. Conf., Vol. 25, Gordon and Breach, New York, 1965. 22 H. Hu, S. Cline and S. tL Goodman, J. AppL Phys., 32 ( 1961 ) 1392. 23 R. E. Smallman, Modern Physical Metallurgy, 4th edn., Butterworths, London, 1985.
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The microstructure of pressureless sintered silver-toughened alumina: an in situ TEM study J. Wang, C. B. Ponton and P. M. Marquis IRC in Materialsfor High PerformanceApplications & School of Metallurgy and Materials, The Universityof Birmingham, Birmingham B15 277"(UK) (Received July 27, 1992; in revised form September 30, 1992)
Abstract An in situ transmission electron microscopy (TEM) study has been made of the microstructure of a pressureless sintered alumina-silver composite, which consists of both intergranularly and intragranularly positioned silver inclusions in the alumina matrix. The silver grains at the grain boundaries and junctions of the alumina matrix are 1-4/~m in size and angular in morphology. Residual thermal strains result in the formation of either the well-established dislocation rings or deformation twins or both in these intergranularly positioned silver particles. The silver particles entrapped inside alumina grains are 0.1-2/~m in size and rounded in morphology. Deformation twins were not found in these intragranularly positioned silver particles. The weak interfacial bonding between silver inclusions and the alumina matrix is characterized by the occurrence of voids or pores at the interface. Thermal mismatch-induced microcracks were not observed at the interface.
I. Introduction
The proprietary studies on strengthening and/or toughening brittle matrices using a metallic second phase, mainly for cutting tool applications, dates back more than 30 years [1]. These earlier investigations did not, however, lead to a well-established understanding of the microstructure-property relationships in these ceramic-metal composites, partly due to the lack of certain appropriate experimental techniques such as transmission electron microscopy (TEM). There has been a new upsurge of interest in ceramic matrices toughened using a metallic second phase in the advanced engineering ceramic community since the early 1980s, and a class of metallic-phase toughened ceramic materials have been fabricated and microstructurally characterized, including lanxide aluminium-toughened alumina [2, 3], nickel-toughened alumina and mullite [4], and, more recently, silvertoughened alumina [5, 6]. A fracture toughness of 7-8 MPam °-5 is obtainable in these metallic-phase-dispersed ceramic matrices via proper compositional and microstructural control, compared with a fracture toughness of 2.5-3.0 MPam °5 for the monolithic ceramic matrices. The presence of a metallic inclusion in a ceramic matrix results in crack bridging by the metallic phase, and crack deflection at the interface between the metallic phase and the ceramic when the composite material is fractured [2, 7, 8]. 0921-5093/93/$6.00
The fabrication of silver-toughened ceramic matrices, which has been carried out using a conventional ceramic processing route, e.g. pressureless sintering of mixed alumina and silver oxide powder compacts, is both simple and economic [5, 6]. During the sintering of an alumina-silver oxide powder compact, a series of physical and chemical changes take place with increasing heating temperature. For example, the silver oxide particles undergo decomposition in the temperature range 400-450 °C, forming metallic silver inclusions in the alumina matrix. The metallic silver particles then melt at 962 °C [9]. It is apparent that the densification of silver-toughened alumina involves a molten silver phase at the grain boundaries and junctions of the alumina matrix. As has been reported [10], the interfacial energy between alumina and molten silver liquid is considerably high and little reaction occurs at the interface. Therefore, the densification mechanisms at the sintering temperature in silver-toughened alumina ceramics are different from those in conventional liquid phase-assisted ceramic systems [11]. Molten silver liquid exhibits high vapour pressure above its melting temperature. For example, a partial pressure of 6.6 mm Hg was measured for molten silver at 1500°C [12]. The high vapour pressure of molten silver liquid results in a large loss in silver content at the sintering temperature, and a silver-free aluminasintered surface, the thickness of which varies with the © 1993 - Elsevier Sequoia. All rights reserved
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composition and sintering temperature, was observed in each sintered alumina-silver composite [6]. On cooling from the sintering temperature, the molten silver inclusions undergo crystallization at temperatures around 962°C [9]. The thermal mismatchinduced residual strains may play an important role in the microstructure-mechanical property relationships in the alumina-silver composite. In comparison with other metallic-phase-toughened ceramics, little investigation has been made into the microstructural characterization of silver-toughened alumina ceramics, although the potentially high fracture toughness has been demonstrated in these materials. The aim of the present work is to carry out a microstructural study on a pressureless sintered alumina-silver composite using in situ TEM observation.
Fig. 1. A silver-free-aluminasintered surface forms in each sintered specimen pellet due to the silver loss resultingfrom the high vapour pressure of molten silver liquid at the sintering temperature.
2. Experimental work The as-received alumina powder (average particle size: 0.75/zm, from BA Chemical Ltd., England) and the as-received silver oxide powder (average particle size: 0.2/tm, from BDH Ltd., Poole, England) were mixed together by ball milling in propanol, using zirconia balls as the milling medium, for 6 h to obtain a composition of AI203 + 40 wt.% mg20. The ball-milled mix was then dried using a combination of an infrared heating lamp and hot plate, followed by compaction in a steel die of 12.5 mm in diameter at a pressure of 120 MPa. The powder compacts were sintered in an electric furnace at 1600 °C for 2 h using heating and cooling rates of 4 °C min- l. As mentioned above, as a result of the high vapour pressure of molten silver liquid at the sintering temperature, the average silver content in the as-sintered sample is lower than that in the pre-sintered powder compact [6]. Furthermore, a silver-free aluminasintered surface is formed in each sintered pellet, see Fig. 1. It was shown, using XRD phase analysis, that the as-sintered samples consisted of alumina and metallic silver. The average silver content in the assintered sample was worked out to be 10.1 vol.% on the basis of total weight loss and the observed thickness of the silver-free alumina-sintered surface. The assintered alumina-silver composite exhibited a sintered density of 98.5% theoretical density, as measured using the Archimedes technique in distilled water, and an indentation toughness of 8.4 MPam °5 (indentation load: 196 N; the equation proposed by Anstis et al. [13] was used for calculation). The average grain size of the alumina matrix is 5.5/zm, as measured using the linear intercept method when more than 150 grains were counted on the polished and then etched surface. To prepare specimens for in situ TEM study, a slice
of 300 # m in thickness was cut from the as-sintered ceramic using a diamond slicing wheel, followed by grinding on a 400 grit SiC paper to reduce the thickness of the slice to approximately 150 gm. Discs of 3 mm in diameter were obtained from the slice using a diamond disc cutter. The specimen discs were further polished and then dimpled, followed by ion beam thinning. The in situ TEM microstructural study was carried out on the ion beam thinned foils using Joel4000CX at 400 kV.
3. Results and discussion Figure 2(a-c) has three bright field transmission electron micrographs showing the general view of the microstructure for the alumina-silver composite fabricated in the present work. It consists of an alumina matrix of 3-10/~m in grain size and silver inclusions of 1-4/~m in size, both at the grain boundaries and junctions of the alumina matrix. Relatively smaller silver particles, in the range of 0.1-2/zm in size, were observed to occur within certain alumina grains. The silver inclusions exhibit considerably different morphologies with respect to their sizes and locations in the alumina matrix: (a) the relatively larger silver grains at the grain junctions of the alumina matrix exhibit an elongated morphology along the grain boundaries; (b) the intermediate-sized silver phase at the grain boundaries of the alumina matrix exhibit a near-ellipsoidal morphology, Fig. 3(a, b); and (c) the relatively smaller silver particles entrapped inside alumina grains are near-spherical in morphology.
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Fig. 2. Three bright field transmission electron micrographs showing the microstructure of the alumina-silver composite fabricated in the present work. The silver grains at the grain boundaries and junctions of the alumina matrix tend to be angular in morphology, and those, or sections of those, entrapped by the aluminagrains are rounded in morphology.
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Voids ol pores often exist at the interface between the silver grains and the alumina matrix. At the sintering temperature, the microstructural development in the alumina-molten silver two-phase system involves two concurrent processes: the interaction between the alumina grain boundaries/junctions and molten silver inclusions, and the ripening of the molten silver inclusions in the alumina matrix. The occurrence of both the intergranularly and intragranulaxly positioned silver particles in the alumina matrix is a direct result of such an interaction between the moving alumina grain boundaries and the molten silver inclusions at the sintering temperature. The contact angle and interfacial energy between alumina and molten silver liquid are high (150 ° and 900 ergs cm-3 respectively) and little reaction takes place between the two phases [10]. Similar to a two-solid phase ceramic system [14, 15], the molten silver inclusions at the grain boundaries of the alumina matrix apply a dragging or pinning force to the mobile grain boundaries, although they are viscously deformable at the sintering temperature. The pinning effect of a silver inclusion to a mobile alumina grain boundary will be largely dependent on its size. Relatively smaller silver inclusions exert a smaller pinning force on the grain boundaries than the relatively larger ones do. Therefore, these small silver inclusions are more likely to be swallowed up by the mobile grain boundaries and therefore more likely to be entrapped within alumina grains than the relatively larger ones. The difference in morphology between the intergranularly and intragranularly positioned silver grains is due to the deformability and high surface tension of the molten silver liquid at the sintering temperature. Those silver inclusions at the grain junctions of the alumina matrix are subjected to a highly anisotropic stress and are likely to be squeezed by the growth of adjacent alumina grains, resulting in the formation of an elongate d morphology along the grain boundaries. Similarly, the near-ellipsoidal morphology of the silver grains at the grain boundaries of the alumina matrix is due to the balance between the pressure from the neighbouring alumina grains and their high surface tension (800-850 mN m -2) at the sintering temperature [9, 12]. Similar to the situation when pores are entrapped by grains of a monolithic ceramic, the relatively smaller silver particles entrapped within alumina grains exhibit a near-spherical morphology in order to minimize the total interfacial free energy [15]. These intragranulaxly positioned silver inclusions are subjected to a more isotropic stress from the alumina matrix than the silver inclusions entrapped at the grain boundaries and grain junctions of the alumina matrix. The formation of voids or pores at the interface between the silver grains and the alumina matrix is
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Fig. 3. The silvergrains at the grain boundaries of the alumina matrix are near-ellipsoidalin morphology.There exist well-established dislocationtings in many silvergrains. closely related to the high vapour pressure of molten silver liquid at the sintering temperature [12]. Owing to this high vapour pressure, on the one hand, Ostwald ripening may be the dominant mechanism responsible for the growth of silver grains in the alumina matrix [16]. On the other hand, a back stress will be associated with the high vapour pressure of the molten silver phase, which will offset a degree of driving force for the densification of the alumina matrix [17]. As has been observed by the present authors [6], the sintering temperature for achieving a sintered density of 98.5% theoretical density in alumina matrices containing more than 5 vol.% silver is more than 100 °C, higher than that needed for achieving a similar sintered density in the monolithic alumina matrix, when a duration of 2 h at the sintering temperature is employed. The final sintered density of the alumina-silver composites was slightly lower than that for the alumina matrix. As mentioned earlier, the decrease in silver content with increased heating temperature is a direct result of the high vapour pressure of the molten silver phase. The disappearance of silver inclusion or a part of silver inclusion will leave a void in the alumina matrix, although further densification may eventually eliminate the void from the structure. The release of vapourized silver from the alumina matrix requires a well-interconnected channel network. It is likely that the grain boundaries will act as such diffusion channels for the vapourized silver phase entrapped at the grain boundaries and junctions of the alumina matrix. It is, however, almost impossible for vapourized silver entrapped inside an alumina grain to be released, as no interconnected channels are available. The entrapped silver vapour may generate an
internal pressure at the interface between the alumina and molten silver phase, limiting the densification attained in this region. This explains the fact that interfacial voids or pores are more likely to occur around silver grains or the sections of silver grains entrapped by the alumina grains. These grains are rounded in morphology. The existence of these interracial voids will assist the occurrence of crack deflection, which was observed to be the dominant toughening mechanism in the silver-toughened alumina ceramics. It is anticipated that the thermal mismatch between the metallic phase and the oxide phase is an important microstructural feature for the silver-dispersed alumina matrix [18, 19]. On cooling from 1600 °C (the sintering temperature) to 962 °C (the melting point of metallic silver), the thermal mismatch between alumina and molten silver liquid may not be important in forming residual strains in silver-toughened alumina, as any thermal stress will be constantly accommodated by the viscously deformable molten silver phase. The linear thermal expansion coefficient of metallic silver ((18-23)x 10 -6 C -1) is almost three times of that for alumina ceramics ((6-9)x 10 -6 C-1) o v e r the temperature range 25-962°C. On cooling at temperatures below 962 °C, a thermal stress will be generated in the composite structure. Owing to thermal mismatch, microcracks occur at the interface in many ceramic matrix composite systems [20]. As will be discussed later, the thermal stress did not lead to the formation of microcracks at the interface between silver inclusions and the alumina matrix in the as-sintered material. The distribution of the thermal mismatch-induced residual strain energy is dependent on the elastic properties of the component materials. The alumina grains exhibit a
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Fig. 4. Three bright field transmission electron micrographsand the associated SADP showingthe occurrence of (a) dislocation rings, (b) partial twinningand (c) complete twinning, of the silver particles at the grain junctions of the alumina matrix. Microcrackswere not observed at the interface between the silvergrains and the aluminamatrix. higher elastic modulus (380 GPa) than the metallic silver ( 105 GPa). Under a given stressing condition, the former will undergo a lower degree of deformation than the latter. As is shown in Figs. 2(a-c) and 3(a-c), few strain contours and dislocation networks were observed to occur in the alumina grains, indicating that these rigid alumina grains accommodate a very small percentage of the total residual strain caused as a result of the thermal mismatch. It has been well established that stress-induced dislocations and deformation twinning occur in metallic silver under an appropriate strain condition [21, 22]. Specifically, plane {111} and direction (112) are the twinning elements for the face-centred cubic silver [23]. It was observed that most silver grains entrapped at the grain boundaries and junctions of the alumina matrix exhibit either a well-defined dislocation network or deformation twins or both, see Figs. 3(b,c), 4(a-c), 5(a,b) and 6(a,b). Microcracks were not found at the interface between the silver inclusions and the alumina matrix, regardless of the difference in particle size and morphology of various silver grains observed. Figure 4(a-c) consists of three bright field TEM micrographs and the associated diffraction patterns showing the occurrence of dislocation rings, partial twinning and complete twinning, respectively, in silver grains entrapped at the grain junctions of the alumina matrix. Selected area electron diffraction confirmed that deformation twinning occurs on the {111} planes. Correlation was not found between the orientation of alumina grains and the twinning direction of silver inclusions, as the thermal stress at a grain boundary or junction of the alumina matrix is a highly anisotropic and complex
process. As an example, Fig. 6(a,b) comprises two bright field transmission electron micrographs and the associated selected area diffraction pattern (SADP) showing that the deformation twins in silver grains can be parallel and perpendicular to the c-axis of the alumina grains. As discussed earlier, certain small silver grains, which exhibit a rounded morphology, are entrapped inside alumina grains. The straining conditions in these intragranularly positioned silver grains may be different from those in the intergranularly positioned silver grains, in terms of their difference in size and in morphology [18, 19]. Well-established strain contours were observed to occur in almost all these rounded silver particles, regardless of their difference in size and in shape, see Fig. 7(a-d). Dislocation rings were found in the relatively large silver particles entrapped within alumina grains, Fig. 7(a,b). Deformation twins were, however, not found in these intragranularly positioned silver grains when more than 60 grains were observed, although some of these silver grains were highly strained. The level of thermal mismatch-induced residual strain energy in a two-phase material is dependent on parameters such as the differences in thermal and elastic properties of the two component phases, and in the size and shape of the secondary phase in the matrix [18, 19]. The residual strain energy increases with increasing particle size of the secondary phase. An angular inclusion is likely to result in a high strain concentration at the interface. As discussed above, the intragranularly positioned silver grains are relatively smaller in size and more rounded in morphology than
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(a)
Fig. 5. A single twin formed in a silver grain entrapped at the grain junction of the alumina matrix. Image (b) was taken when image (a) was tilted by 14".
Fig. 6. Two bright field TEM micrographs and the associated SADP showing that the deformation twins in silver grains can be parallel and perpendicular to the C-axis of neighbouring alumina grains.
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Fig. 7. Two small silver particles entrapped within alumina grains. The residual thermal stress results in the formation of strain contours and dislocation rings in these relatively small and rounded silver particles. The highly strained structure in (d) is indicated by the slightly distorted SADP. Deformation twins were not found when more than 60 grains were observed.
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those entrapped at the grain boundaries and junctions of the alumina matrix. The thermal mismatch-induced residual strain energy level in these intragranularly positioned silver particles will be lower than in those intergranularly positioned silver grains. A minimum level of residual strain energy is required in order to generate either a dislocation ring or a deformation twin in the face-centred cubic silver. Specifically, the strain energy needed to induce a deformation twin is generally higher than that needed to induce a dislocation ring in metallic silver, if it is assumed that the formation of the deformation twin is via a dislocation mechanism [23]. The thermal mismatch-induced residual strain energy in the intragranularly positioned silver particles may not be high enough to result in the formation of deformation twinning on cooling from the sintering temperature. This explains the absence of deformation twins in these small and rounded silver grains.
4. Summary The pressureless sintered alumina-silver composite consists of both intergranularly positioned silver grains of 1-4 p m in size and intragranularly positioned silver particles of 0.1-2 p m in size in the alumina matrix. The latter are more rounded in morphology than the former. The weak interface between the silver grains and the alumina matrix is characterized by the occurrence of voids or pores around the silver grains or part of the silver grains entrapped by the alumina grains. In situ TEM observation showed that the thermal mismatch-induced residual strains in the alumina-silver composite were largely accommodated by the silver particles, which are much less rigid than the alumina grains. Microcracks were not observed at the interface between silver inclusions and the alumina matrix. The residual thermal stress results in the formation of either the well-established dislocation rings or deformation twins or both in silver particles entrapped at the grain boundaries and junctions of the alumina matrix. Deformation twins were not observed in the intragranularly positioned silver particles, as they are relatively smaller in size and more rounded in morphology than those at the grain boundaries and junctions of the alumina matrix and therefore a limited level of thermal strain is induced at the interface on cooling from the sintering temperature.
Acknowledgments The authors thank Dr. T. J. Doel of the IRC in Materials for High Performance Applications, University of Birmingham, for reading the manuscript of this paper. The financial support given for this project by the SERC is acknowledged.
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