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Analytical investigation of grain boundaries of compressive deformed Al-doped sintered β-SiC
Materials Science and Engineering A285 (2000) 136 – 143 www.elsevier.com/locate/msea
Analytical investigation of grain boundaries of compressive deformed Al-doped sintered b-SiC Kenji Kaneko a,*, Sawao Honda a,1, Takayuki Nagano a, Tomohiro Saitoh b a
Ceramics Superplasticity, ICORP, Japan Science and Technology Corporation (JST), c/o Japan Fine Ceramics Center 2F, 2 -4 -1, Mutsuno, Atsuta-ku, Nagoya 456 8587, Japan b Japan Fine Ceramics Centre, 2 -4 -1, Mutsuno, Atsuta-ku, Nagoya 456 -8587, Japan
Abstract Series of analytical investigations were carried out to study the effect of deformation process at grain boundaries of Al-doped b-SiC (b-SiC[Al]). Three specimens, as-sintered, annealed and compressed b-SiC[Al] specimens were provided. Grain boundaries were observed by a high-resolution transmission electron microscope (HRTEM) and analytically investigated by energy dispersive X-ray spectroscopy (EDS) also electron energy loss spectroscopy (EELS) attached on a scanning transmission electron microscope (STEM). Glass phases Al–Si–O were commonly observed from the grain boundaries of the as-sintered specimen. It was also analyzed by EELS measurements that the amorphous phases were occasionally modified to form the fourfold coordinated aluminum-oxide by the compressive deformation. On the other hand, there were little traces of Al and O atoms observed from annealed specimen which suggests the vaporization of amorphous glass phase from the grain boundaries, due to the elevated temperature. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Deformation; Grain boundary; SiC; High resolution transmission electron microscopy; Scanning transmission electron microscope; Energy dispersive X-ray spectroscopy; Electron energy loss spectroscopy; Sintering additives
1. Introduction In general, the non-oxide ceramics, such as SiC and Si3N4, are difficult to densify by normal sintering without additives due to their extremely low self-diffusiveness, usually B, C and Al composites are chosen as sintering additives for SiC [1]. The additives have effects to lower the grain boundary energy then to promote sintering process. In the case of Al composites, they are in the form of metal Al, Al2O3 or AlN. Al2O3-doped specimen shows stronger toughness than the B- or C- doped specimen, and these liquid phases act as the mass transport media during the densification via a solution-reprecipitation mechanism. Al increases the lattice diffusion coefficient of SiC by incorporation into SiC grains as a solid solution which contributes to the densification in the sintering process. Incorporation of Al does not effect the grain growth rate which is proportional to the rate of grain boundary diffusion. * Corresponding author. 1 Department of Materials Science and Technology, Nagoya Institute of Technology, Gokiso, Showa-ku, Nagoya 466-85555, Japan.
The transport properties of the liquid phase will be dependent on its volume fraction and chemical environments. These are determined by the characteristics of additives, the relevant eutectic temperatures, and the densification parameters such as the temperature and the atmosphere [2,3]. Additives alter the chemical stability of the final material and degrade its mechanical properties at high temperatures, due to their segregation at the grain boundaries and the formation of glass phase. The glass phase softens at elevated temperatures then becomes as the negative effects on the mechanical and the chemical properties. It is commonly known that Al2O3 usually participates in liquid phase formation so that SiC degrades its strength at elevated temperatures, at around 2200 K. The relationship between the mechanical properties and additives has been studied carefully by many researchers, with respect to the micro- and nano-structures for both a- and b-SiC [4–9]. The grain boundary has been known as a critical factor to determine the strength of structural ceramics. For example, Si3N4 is known to be strongly dependent
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on the secondary phases, commonly present at the multiple grain junctions and the grain boundaries [10,11]. Also in the case of SiC, it is known that the high temperature strength and the deformability strongly depend on the types of sintering additives and their amount [8,12].
(a)
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In this paper, addition to the grain boundary structural investigations by HRTEM, both compositions and chemical environments are examined for three types of Al-doped b-SiC, b-SiC[Al], specimens. Addition to the EDS analysis showing that there were traces of both O and Al at the grain boundaries of the as-sintered and the compressed specimens, it was shown (b)
(c)
Fig. 1. (a) HRTEM image presenting grain boundary of the as-sintered specimen. (b) HRTEM image presenting grain boundary of the compressed specimen. (c) HRTEM image presenting grain boundary of the annealed specimen.
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Fig. 2. (a) EDS Al – K edge Line profile results over grain boundary. (b) EDS O – K edge Line profile results over grain boundary.
by EELS measurements that the chemical environments of the grain boundaries were different from as-sintered to compressed b-SiC[Al] specimen. The clean grain boundaries were observed from the annealed specimen, with little traces of A1 and O at the grain boundaries, by EELS.
2. Experimental procedure
2.1. Specimens b-SiC[Al] was produced by using the precursorderived ceramic method from polycarbosilane at Hitachi Research Laboratory [6,13]. Before pyrolyzing this powder, it was heated at 463 K for 1.8 ks in the air to promote cross-linking of the polymer structure. Pyrolization of the oxidized polycarbosilane was performed at 1573 K for 3.6 ks under vacuum, then fine b-SiC powders were prepared by grinding polycarbosilane powder. The grain size of the specimen at 2023 K was estimated from the sharpness of the XRD lines, as around 0.1 mm [13]. b-SiC powder was then mixed with 2 mol% of metal Al powder as a sintering additive. The mixed powder was molded and hot-pressed at 2173 K under 50 MPa
for 5.4 ks to fabricate the ceramic body. The fabrication process of the polycrystalline SiC was starting from polycarbosilane, then powder synthesis (oxidation pyrolysis), adding 2 mol% of sintering additives as Al, followed by a sintering process (mixing molding hot-pressing). Sintering additives of the Al compounds have been known to densify SiC well with suppressing the grain growth, in comparison with B or Be systems [14]. As a result, the final average grain sizes of the ceramics were measured as about 0.7 mm and maximum fracture toughness was found 5.1 Mpa·m1/2 [6]. The impurity content of the specimen was measured by inductively coupled plasma spectrometry, showed 1.67 wt.% of Al, and by Fourier transform infrared spectroscopy, showed 2.35 wt.% of O. Although Al and O may have been dissolved into SiC during the sintering process, the segregations of Al and/or Al2O3 at grain boundaries were expected for this case. Since, the solubility limit of Al in SiC was measured by powder X-ray diffraction method (Debye–Scherrer method) as almost 0.4 wt.% Al2O3, thus about 0.2 wt.% of Al, at around 2200 K [8,15]. Those three specimens available for this investigations were; as-sintered specimen, compressed specimen at 2173 K with strain rate 5·10 − 5 s − 1 for 4.2 ks with final strain 0.1, and annealed specimen with the same condition as at 2173 K for 4.2 ks, both in Ar.
2.2. Electron microscopy TEM specimens were prepared by the standard mechanical thinning method of a 2.1 ×2.1 mm square cut from each bulk specimens and mechanically thinned up to 50 mm. Then twenty to thirty hours of Argon ion milling at 5 kV, 0.5 mA and a beam angle of 14° was carried out until there was a minute hole at the center of the specimen disk. The applied voltage and the beam angle were gradually decreased to 2 kV and 12° respectively, to produce large area of appropriate specimen thickness. The accelerating voltage for ion-milling was reduced to prevent the radiation damage of the foils. HRTEM imaging was carried out on a Topcon 002B equipped with a LaB6 gun having a point resolution of 0.18 nm operated at 200 kV. AEM was performed with a dedicated STEM (Thermo, Vacuum-Generators HB 601 UX) operated at 100 kV, equipped with a high-resolution pole piece and a cold field emission gun. An energy resolution of better than 0.4 eV was obtained from the full-width half-maximum (FWHM) of zeroloss peak with the spectrometer entrance aperture limited to about 12 mrad. The STEM is equipped with an Oxford Instruments EDS system with a Super ATW [Si(Li)] window and a Parallel-EELS (PEELS) spectrometer (Gatan 766 with photodiode array).
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For measuring the EELS edges, the beam current was kept around approximately 0.2 nA with a mean beam diameter of almost 0.3 nm (FWHM) [16,17]. Almost 50 raw Si – L, Al – L and O – K EELS edges were obtained from each specimens, then processed using standard calculation tool provided with PEELS equipment, EL/P 3.0 (Gatan, USA). The thickness does not degrade the spatial resolution in EELS significantly, though the multiple scattering occurs in the thick specimen which causes difficulties to analyze the signal [18,19].
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the secondary phase completely wets the grain boundaries. The formation of amorphous grain boundary films is assumed to be the release of residual strain at the grain boundaries. Some significant changes compared with as-sintered specimen were observed from both compressed and annealed specimens, such as no amorphous phases observed, as can be seen from Fig. 1(b,c) respectively. The secondary phases may have been completely removed from the grain boundaries at elevated temperatures.
3.1. Energy dispersi6e X-ray spectroscopy 3. Results and discussion HRTEM imaging of grain boundaries showed that there were amorphous secondary phase from the as-sintered specimen, as shown in Fig. 1(a). The occurrence of the amorphous films in the specimen indicates that
Line profile method near grain boundaries of specimens was achieved along a direction vertical to the grain boundaries, then the elemental compositions of grain boundaries were analyzed. The impurity contents were detected in terms of the intensity of X-ray radia-
Fig. 3. ELNES obtained from the as-sintered specimen. Around Si-L edges and O – K edge, Fig. 3(a) and 4(a) respectively. ELNES obtained from the compressed specimen. Around Si-L edges and O–K edge, Fig. 3(b) and 4(b) respectively. ELNES obtained from the annealed specimen. Around Si – L edges and O–K edge, Fig. 3(c) and 4(c) respectively.
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Fig. 3. (Continued)
tion from the Al– K edge and O – K edge relative to the intensity of the Si – K edge signal. The elements observed during this experiments of among as-sintered, compressed and annealed specimens were, O and Al atoms (seen in the EDS spectrum, Fig. 2(a,b) respectively) with respect to Si atoms at grain boundaries were observed from the as-sintered and also from the compressed specimen. In all cases, EDS spectra were acquired for 50 s from each spots. Both as-sintered and compressed specimen showed the presence of Al and O clearly at grain boundaries. It was assumed from these results that the composition of grain boundaries were either AlxSiyOz glass or aluminum-oxides. During the compression experiment, the creation of cavities may have been restrained by the cause of stress concentration and interlocking at triple points. However, the presence of Al or O were not observed from the annealed specimen, which indicate the occurrence of vaporization of secondary phase during annealing process.
4. EELS/ELNES The results of EDS only indicate the chemical composition of grain boundaries but not their chemistries. Therefore, addition to the EDS results, EELS measurement had to be carried out to determine the chemistry of grain boundaries. EELS allows to extract information about the chemical composition of light elements, the bonding, the valence states, the atomic coordination, the symmetry, and other useful information of the specimen [19,20]. By extrapolation of the pre-edge background region (up to 20–30 eV above the onset of characteristic ionization edges fitted by a power law in the form of A·e − r, where A and r are constants), energy loss near edge structure (ELNES) can be achieved to optimize the specific information of materials from the original data [19]. The spectra were taken with the beam scanning a 2.0× 2.0 nm2 [17] and also with the stationary spot beam method having a probe size of about 3 A, . In comparison to a stationary spot-mode beam, the scan-
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ning method not only reduces the beam damage on the specimen but also allows a manual correction of the specimen drift by minor adjustments to the beam deflection coils, since the grain boundary can simultaneously be imaged on the dark field detectors. Three sets of spectra were recorded essentially with 0.1 eV/channel energy dispersion from various energy ranges to observe Al – L2,3 edges, Si – L2,3 edges and O –K edge. The regions of 20 – 30 nm thickness were carefully chosen as an optimum area to obtain EELS by Zero-loss/Plasmon-loss measurement. The original spectrum is usually dominated by the bulk matrix component due to the finite size of the beam.
4.1. As-sintered specimen It was clearly shown by EDS results that there were segregation of Al and O atoms at grain boundaries, where oxygen is believed to be located on the exterior
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of SiC grains originally as a contaminant during the material processing. ELNES of energy range between 50 and 150 eV showed that there are edges starting from almost 75 eV. The pre-edge of Si-L2,3 edges, an indicative of Al–L2,3 edges, see Fig. 3.(a), were still too weak to justify the chemical composition at grain boundaries, and only verified the presence of Al atoms. The O–K edge was also analyzed to determine the oxygen environment, and compared with O–K ELNES obtained from other form of alumino-oxides. The production of alumino-silicate-glass at grain boundaries are favored, due to the occurrence of oxygen molecule vaporization due to the irradiation effect (known as electron-hole drilling, see Fig. 4(a), marked X). The HRTEM image obtained, see Fig l.(a), showed the presence of glass phase at the grain boundaries and the O–K edge, Fig. 4(a), strongly indicated the fabrication of alumino-silicate-glass at the grain boundaries.
Fig. 4. ELNES obtained from the a –Al2O3 specimen. Around Al – L edges and O – K edge, Fig. 3(d) and 4(d) respectively. ELNES obtained from the gamma-Al2O3 specimen. Around Al–L edges and O–K edge, Fig. 3(e) and 4(e) respectively. Si-L ELNES edge obtained from beta-SiC, Fig. 3(f) and O – K ELNES edge obtained from SiO2-glass, Fig. 4(f).
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Fig. 4. (Continued)
4.2. Compressed specimen The same procedure was carried out for the compressed specimen as the as-sintered specimen. As can be seen from Al-L ELNES spectrum obtained from the grain boundaries, Fig. 3(b), the creation of alumino-oxides at grain boundaries. The O – K ELNES spectra, Fig. 4(b), was then compared with various types of Al2O3 spectra as in the case of as-sintered specimen, to confirm that it was very similar to the spectrum obtained from the four-fold coordinated alumino-oxides, such as gAl2O3.
at the grain boundaries, no sign of O or Al atoms were detected from the annealed specimen. According to the annealing experiment with respect to weight loss, Fig. 5, the decrease of weight started from 1400 K. The vaporization process [21] may be taking place above 1400 K; SiC (s)+ 2SiO2(l)3SiO(g)+ CO(g)
4.3. Simply annealed specimen Although the presence of Al or O at grain boundaries was discarded by EDS, see Fig. 2(a,b), that ELNES was also obtained from the annealed specimen. It was again denied by EELS, as can be seen from Fig. 3(c) and 4(c), Al–L and O – K edges respectively. Even though spatial difference – EELS/ELNES [16] was applied to characterize the ELNES contribution of atoms
Fig. 5. The annealing experiment with respect to the weight loss.
K. Kaneko et al. / Materials Science and Engineering A285 (2000) 136–143
SiC (s)+ Al2O3(s) Al2O (g) + SiO(g) +CO(g) so that the Al–Si – O glass becomes vaporized. This process may be assisted by the cavitation damage and promoted by the vaporization of the amorphous phase. On the other hand, in the case of compressed specimen, the compressive stress would suppress the vaporization process. Meantime, the Al and O atoms remaining at the grain boundaries may alter its forms from the Al–Si–O glass to four-fold coordinated material. Figs. 3(d–f) and 4(d – f) are the reference spectra obtained from standard materials to apply fingerprint technique for core-loss fine structures [19].
5. Conclusions Effect of deformation process on grain boundaries, particularly on sintering additives were investigated by using both HRTEM and STEM. HRTEM images presented that the presence of amorphous secondary phase at the grain boundaries of as-sintered specimen, disappearance of amorphous phase at the grain boundaries of the compressed specimen and of the annealed specimen. Although EDS results clearly showed the presence of Al and O at grain boundaries from both as-sintered and compressed specimens, EELS analysis confirmed that the chemical environments of them are different. The disappearance of the glass phases may be caused by the vaporization by the application of high temperature in both the compressed and the annealed specimen, though Al and O atoms still remained in the case of the compressed specimen. Sintering conditions influences the additives at the grain boundaries by altering the phases of sintering additives. This leads to the change of mechanical properties in macroscopic sense. Although it is still too early to conclude that the effects of sintering additives at the grain boundaries are directly related to the deformation mechanisms or not, specimens with other sintering additives suitable for SiC (either boron or carbon, or mixture of Y2O3 and Al2O3) must be studied to compare with the results obtained from both macroscopic and microscopic investigations in the future. Addition to using different sintering additives, it is suggested to carry out other series of analysis by electron microscopes on specimens with deformation at low temperatures (so that there are no grain growth during the deformation process), and specimens with static growth (so that there are no stress but with grain growth). The grain nucleation might have occurred with the grain boundary sliding as dynamic recrystallization during deformation. .
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The work presented here distinctively demonstrated the combination of HRTEM imaging and AEM analysis as a powerful tool for material investigations, specially to investigate structures and chemistry of the grain boundaries. Furthermore, it is advantageous and recommended to compare experimental results with theoretical modeling of the grain boundary structures, like ab initio calculation [20], and theoretical calculations of electronic structures at the grain boundaries, molecular orbital method [22] and multiple scattering method [23]. Acknowledgements We would like to acknowledge Dr H. Kodama from Hitachi research Lab. for providing specimen, Dr Tsurekawa (Tohoku University) and Dr K. Kakimoto (JST) for valuable discussions. References [1] R.A. Alliegro, L.B. Coffin, J.R. Tinklepaugh, J. Am. Ceram. Soc. 39 (1956) 386. [2] F.F. Lange, J. Mater. Sci. 10 (1975) 314. [3] M.H. Lewis, R.J. Lumby, Powder Metall. 26 (1983) 73. [4] D.H. Stutz, S. Prochazka, J. Lorenz, J. Am. Ceram. Soc. 68 (1985) 479. [5] T.B. Jackson, A.C. Hurford, S.L. Bruner, R.A. Cutler, in: J.D. Cawley, C.E. Semler (Eds.), Silicon Carbide ‘87, American Ceramics Society, Westerville, (1989) 309. [6] H. Kodama, T. Miyoshi, J. Am. Ceram. Soc. 73 (1990) 3081. [7] K. Suzuki. In: S. Somiya, Y. Inomata (Eds), Silicon Carbide Ceramics-2, Elsevier, Amsterdam, p. 162. [8] T. Kinoshita, S. Munekawa, S.-I. Tanaka, Acta. Metall. Mater. 45 (1997) 801. [9] L.K.L. Falk, J. Euro. Ceram. Soc. 17 (1997) 983. [10] A. Tsuge, K. Nishida, M. Komatsu, J. Am. Ceram. Soc. 58 (1975) 323. [11] T. Watanabe, S. Kimura, S. Karashima, Phil. Mag. A. 49 (1984) 845. [12] M.A. Mulla, V.D. Krstic, Acta. Metall. Mater. 42 (1994) 303. [13] H. Kodama, T. Miyoshi, Adv. Ceram. Mater. 3 (1988) 177. [14] Y. Takeda, K. Nakamura, K. Maeda, M. Ura, Adv. Ceram. Mater. 1 (1986) 162. [15] Y. Tajima, W.D. Kingery, J. Am. Ceram. Soc. C 65 (1982) 27. [16] K. Kaneko, I. Tanaka, M. Yoshiya, Appl. Phys. Lett. 72 (1998) 191. [17] K. Kaneko, T. Gemming, I. Tanaka and H. Mullejans, Phil. Mag. A. (in press). [18] D.C. Joy, In: D.C. Joy, A.D.J. Roming, J.I. Goldstein (Eds.), Principles of Analytical Electron Microscopy, Plenum Press, New York, (1989). [19] R.F. Egerton, Electron Energy Loss spectroscopy in the Electron Microscope, 2nd edition, Plenum Press, New York, 1996. [20] P.R. Kenway, J. Am. Ceram. Soc. 77 (1994) 349. [21] T. Grande, H. Sommerset, E. Hagen, K. Wiik, M-A. Einarstrud, J. Am. Ceram. Soc. 80 (1997) 1047. [22] I. Tanaka, H. Adachi, J. Phys. D: Appl. Phys. 29 (1996) 1725. [23] R. Brydson, J. Micros. 180 (1995) 238.
Analytical investigation of grain boundaries of compressive deformed Al-doped sintered b-SiC Kenji Kaneko a,*, Sawao Honda a,1, Takayuki Nagano a, Tomohiro Saitoh b a
Ceramics Superplasticity, ICORP, Japan Science and Technology Corporation (JST), c/o Japan Fine Ceramics Center 2F, 2 -4 -1, Mutsuno, Atsuta-ku, Nagoya 456 8587, Japan b Japan Fine Ceramics Centre, 2 -4 -1, Mutsuno, Atsuta-ku, Nagoya 456 -8587, Japan
Abstract Series of analytical investigations were carried out to study the effect of deformation process at grain boundaries of Al-doped b-SiC (b-SiC[Al]). Three specimens, as-sintered, annealed and compressed b-SiC[Al] specimens were provided. Grain boundaries were observed by a high-resolution transmission electron microscope (HRTEM) and analytically investigated by energy dispersive X-ray spectroscopy (EDS) also electron energy loss spectroscopy (EELS) attached on a scanning transmission electron microscope (STEM). Glass phases Al–Si–O were commonly observed from the grain boundaries of the as-sintered specimen. It was also analyzed by EELS measurements that the amorphous phases were occasionally modified to form the fourfold coordinated aluminum-oxide by the compressive deformation. On the other hand, there were little traces of Al and O atoms observed from annealed specimen which suggests the vaporization of amorphous glass phase from the grain boundaries, due to the elevated temperature. © 2000 Elsevier Science S.A. All rights reserved. Keywords: Deformation; Grain boundary; SiC; High resolution transmission electron microscopy; Scanning transmission electron microscope; Energy dispersive X-ray spectroscopy; Electron energy loss spectroscopy; Sintering additives
1. Introduction In general, the non-oxide ceramics, such as SiC and Si3N4, are difficult to densify by normal sintering without additives due to their extremely low self-diffusiveness, usually B, C and Al composites are chosen as sintering additives for SiC [1]. The additives have effects to lower the grain boundary energy then to promote sintering process. In the case of Al composites, they are in the form of metal Al, Al2O3 or AlN. Al2O3-doped specimen shows stronger toughness than the B- or C- doped specimen, and these liquid phases act as the mass transport media during the densification via a solution-reprecipitation mechanism. Al increases the lattice diffusion coefficient of SiC by incorporation into SiC grains as a solid solution which contributes to the densification in the sintering process. Incorporation of Al does not effect the grain growth rate which is proportional to the rate of grain boundary diffusion. * Corresponding author. 1 Department of Materials Science and Technology, Nagoya Institute of Technology, Gokiso, Showa-ku, Nagoya 466-85555, Japan.
The transport properties of the liquid phase will be dependent on its volume fraction and chemical environments. These are determined by the characteristics of additives, the relevant eutectic temperatures, and the densification parameters such as the temperature and the atmosphere [2,3]. Additives alter the chemical stability of the final material and degrade its mechanical properties at high temperatures, due to their segregation at the grain boundaries and the formation of glass phase. The glass phase softens at elevated temperatures then becomes as the negative effects on the mechanical and the chemical properties. It is commonly known that Al2O3 usually participates in liquid phase formation so that SiC degrades its strength at elevated temperatures, at around 2200 K. The relationship between the mechanical properties and additives has been studied carefully by many researchers, with respect to the micro- and nano-structures for both a- and b-SiC [4–9]. The grain boundary has been known as a critical factor to determine the strength of structural ceramics. For example, Si3N4 is known to be strongly dependent
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on the secondary phases, commonly present at the multiple grain junctions and the grain boundaries [10,11]. Also in the case of SiC, it is known that the high temperature strength and the deformability strongly depend on the types of sintering additives and their amount [8,12].
(a)
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In this paper, addition to the grain boundary structural investigations by HRTEM, both compositions and chemical environments are examined for three types of Al-doped b-SiC, b-SiC[Al], specimens. Addition to the EDS analysis showing that there were traces of both O and Al at the grain boundaries of the as-sintered and the compressed specimens, it was shown (b)
(c)
Fig. 1. (a) HRTEM image presenting grain boundary of the as-sintered specimen. (b) HRTEM image presenting grain boundary of the compressed specimen. (c) HRTEM image presenting grain boundary of the annealed specimen.
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Fig. 2. (a) EDS Al – K edge Line profile results over grain boundary. (b) EDS O – K edge Line profile results over grain boundary.
by EELS measurements that the chemical environments of the grain boundaries were different from as-sintered to compressed b-SiC[Al] specimen. The clean grain boundaries were observed from the annealed specimen, with little traces of A1 and O at the grain boundaries, by EELS.
2. Experimental procedure
2.1. Specimens b-SiC[Al] was produced by using the precursorderived ceramic method from polycarbosilane at Hitachi Research Laboratory [6,13]. Before pyrolyzing this powder, it was heated at 463 K for 1.8 ks in the air to promote cross-linking of the polymer structure. Pyrolization of the oxidized polycarbosilane was performed at 1573 K for 3.6 ks under vacuum, then fine b-SiC powders were prepared by grinding polycarbosilane powder. The grain size of the specimen at 2023 K was estimated from the sharpness of the XRD lines, as around 0.1 mm [13]. b-SiC powder was then mixed with 2 mol% of metal Al powder as a sintering additive. The mixed powder was molded and hot-pressed at 2173 K under 50 MPa
for 5.4 ks to fabricate the ceramic body. The fabrication process of the polycrystalline SiC was starting from polycarbosilane, then powder synthesis (oxidation pyrolysis), adding 2 mol% of sintering additives as Al, followed by a sintering process (mixing molding hot-pressing). Sintering additives of the Al compounds have been known to densify SiC well with suppressing the grain growth, in comparison with B or Be systems [14]. As a result, the final average grain sizes of the ceramics were measured as about 0.7 mm and maximum fracture toughness was found 5.1 Mpa·m1/2 [6]. The impurity content of the specimen was measured by inductively coupled plasma spectrometry, showed 1.67 wt.% of Al, and by Fourier transform infrared spectroscopy, showed 2.35 wt.% of O. Although Al and O may have been dissolved into SiC during the sintering process, the segregations of Al and/or Al2O3 at grain boundaries were expected for this case. Since, the solubility limit of Al in SiC was measured by powder X-ray diffraction method (Debye–Scherrer method) as almost 0.4 wt.% Al2O3, thus about 0.2 wt.% of Al, at around 2200 K [8,15]. Those three specimens available for this investigations were; as-sintered specimen, compressed specimen at 2173 K with strain rate 5·10 − 5 s − 1 for 4.2 ks with final strain 0.1, and annealed specimen with the same condition as at 2173 K for 4.2 ks, both in Ar.
2.2. Electron microscopy TEM specimens were prepared by the standard mechanical thinning method of a 2.1 ×2.1 mm square cut from each bulk specimens and mechanically thinned up to 50 mm. Then twenty to thirty hours of Argon ion milling at 5 kV, 0.5 mA and a beam angle of 14° was carried out until there was a minute hole at the center of the specimen disk. The applied voltage and the beam angle were gradually decreased to 2 kV and 12° respectively, to produce large area of appropriate specimen thickness. The accelerating voltage for ion-milling was reduced to prevent the radiation damage of the foils. HRTEM imaging was carried out on a Topcon 002B equipped with a LaB6 gun having a point resolution of 0.18 nm operated at 200 kV. AEM was performed with a dedicated STEM (Thermo, Vacuum-Generators HB 601 UX) operated at 100 kV, equipped with a high-resolution pole piece and a cold field emission gun. An energy resolution of better than 0.4 eV was obtained from the full-width half-maximum (FWHM) of zeroloss peak with the spectrometer entrance aperture limited to about 12 mrad. The STEM is equipped with an Oxford Instruments EDS system with a Super ATW [Si(Li)] window and a Parallel-EELS (PEELS) spectrometer (Gatan 766 with photodiode array).
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For measuring the EELS edges, the beam current was kept around approximately 0.2 nA with a mean beam diameter of almost 0.3 nm (FWHM) [16,17]. Almost 50 raw Si – L, Al – L and O – K EELS edges were obtained from each specimens, then processed using standard calculation tool provided with PEELS equipment, EL/P 3.0 (Gatan, USA). The thickness does not degrade the spatial resolution in EELS significantly, though the multiple scattering occurs in the thick specimen which causes difficulties to analyze the signal [18,19].
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the secondary phase completely wets the grain boundaries. The formation of amorphous grain boundary films is assumed to be the release of residual strain at the grain boundaries. Some significant changes compared with as-sintered specimen were observed from both compressed and annealed specimens, such as no amorphous phases observed, as can be seen from Fig. 1(b,c) respectively. The secondary phases may have been completely removed from the grain boundaries at elevated temperatures.
3.1. Energy dispersi6e X-ray spectroscopy 3. Results and discussion HRTEM imaging of grain boundaries showed that there were amorphous secondary phase from the as-sintered specimen, as shown in Fig. 1(a). The occurrence of the amorphous films in the specimen indicates that
Line profile method near grain boundaries of specimens was achieved along a direction vertical to the grain boundaries, then the elemental compositions of grain boundaries were analyzed. The impurity contents were detected in terms of the intensity of X-ray radia-
Fig. 3. ELNES obtained from the as-sintered specimen. Around Si-L edges and O – K edge, Fig. 3(a) and 4(a) respectively. ELNES obtained from the compressed specimen. Around Si-L edges and O–K edge, Fig. 3(b) and 4(b) respectively. ELNES obtained from the annealed specimen. Around Si – L edges and O–K edge, Fig. 3(c) and 4(c) respectively.
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Fig. 3. (Continued)
tion from the Al– K edge and O – K edge relative to the intensity of the Si – K edge signal. The elements observed during this experiments of among as-sintered, compressed and annealed specimens were, O and Al atoms (seen in the EDS spectrum, Fig. 2(a,b) respectively) with respect to Si atoms at grain boundaries were observed from the as-sintered and also from the compressed specimen. In all cases, EDS spectra were acquired for 50 s from each spots. Both as-sintered and compressed specimen showed the presence of Al and O clearly at grain boundaries. It was assumed from these results that the composition of grain boundaries were either AlxSiyOz glass or aluminum-oxides. During the compression experiment, the creation of cavities may have been restrained by the cause of stress concentration and interlocking at triple points. However, the presence of Al or O were not observed from the annealed specimen, which indicate the occurrence of vaporization of secondary phase during annealing process.
4. EELS/ELNES The results of EDS only indicate the chemical composition of grain boundaries but not their chemistries. Therefore, addition to the EDS results, EELS measurement had to be carried out to determine the chemistry of grain boundaries. EELS allows to extract information about the chemical composition of light elements, the bonding, the valence states, the atomic coordination, the symmetry, and other useful information of the specimen [19,20]. By extrapolation of the pre-edge background region (up to 20–30 eV above the onset of characteristic ionization edges fitted by a power law in the form of A·e − r, where A and r are constants), energy loss near edge structure (ELNES) can be achieved to optimize the specific information of materials from the original data [19]. The spectra were taken with the beam scanning a 2.0× 2.0 nm2 [17] and also with the stationary spot beam method having a probe size of about 3 A, . In comparison to a stationary spot-mode beam, the scan-
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ning method not only reduces the beam damage on the specimen but also allows a manual correction of the specimen drift by minor adjustments to the beam deflection coils, since the grain boundary can simultaneously be imaged on the dark field detectors. Three sets of spectra were recorded essentially with 0.1 eV/channel energy dispersion from various energy ranges to observe Al – L2,3 edges, Si – L2,3 edges and O –K edge. The regions of 20 – 30 nm thickness were carefully chosen as an optimum area to obtain EELS by Zero-loss/Plasmon-loss measurement. The original spectrum is usually dominated by the bulk matrix component due to the finite size of the beam.
4.1. As-sintered specimen It was clearly shown by EDS results that there were segregation of Al and O atoms at grain boundaries, where oxygen is believed to be located on the exterior
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of SiC grains originally as a contaminant during the material processing. ELNES of energy range between 50 and 150 eV showed that there are edges starting from almost 75 eV. The pre-edge of Si-L2,3 edges, an indicative of Al–L2,3 edges, see Fig. 3.(a), were still too weak to justify the chemical composition at grain boundaries, and only verified the presence of Al atoms. The O–K edge was also analyzed to determine the oxygen environment, and compared with O–K ELNES obtained from other form of alumino-oxides. The production of alumino-silicate-glass at grain boundaries are favored, due to the occurrence of oxygen molecule vaporization due to the irradiation effect (known as electron-hole drilling, see Fig. 4(a), marked X). The HRTEM image obtained, see Fig l.(a), showed the presence of glass phase at the grain boundaries and the O–K edge, Fig. 4(a), strongly indicated the fabrication of alumino-silicate-glass at the grain boundaries.
Fig. 4. ELNES obtained from the a –Al2O3 specimen. Around Al – L edges and O – K edge, Fig. 3(d) and 4(d) respectively. ELNES obtained from the gamma-Al2O3 specimen. Around Al–L edges and O–K edge, Fig. 3(e) and 4(e) respectively. Si-L ELNES edge obtained from beta-SiC, Fig. 3(f) and O – K ELNES edge obtained from SiO2-glass, Fig. 4(f).
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Fig. 4. (Continued)
4.2. Compressed specimen The same procedure was carried out for the compressed specimen as the as-sintered specimen. As can be seen from Al-L ELNES spectrum obtained from the grain boundaries, Fig. 3(b), the creation of alumino-oxides at grain boundaries. The O – K ELNES spectra, Fig. 4(b), was then compared with various types of Al2O3 spectra as in the case of as-sintered specimen, to confirm that it was very similar to the spectrum obtained from the four-fold coordinated alumino-oxides, such as gAl2O3.
at the grain boundaries, no sign of O or Al atoms were detected from the annealed specimen. According to the annealing experiment with respect to weight loss, Fig. 5, the decrease of weight started from 1400 K. The vaporization process [21] may be taking place above 1400 K; SiC (s)+ 2SiO2(l)3SiO(g)+ CO(g)
4.3. Simply annealed specimen Although the presence of Al or O at grain boundaries was discarded by EDS, see Fig. 2(a,b), that ELNES was also obtained from the annealed specimen. It was again denied by EELS, as can be seen from Fig. 3(c) and 4(c), Al–L and O – K edges respectively. Even though spatial difference – EELS/ELNES [16] was applied to characterize the ELNES contribution of atoms
Fig. 5. The annealing experiment with respect to the weight loss.
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SiC (s)+ Al2O3(s) Al2O (g) + SiO(g) +CO(g) so that the Al–Si – O glass becomes vaporized. This process may be assisted by the cavitation damage and promoted by the vaporization of the amorphous phase. On the other hand, in the case of compressed specimen, the compressive stress would suppress the vaporization process. Meantime, the Al and O atoms remaining at the grain boundaries may alter its forms from the Al–Si–O glass to four-fold coordinated material. Figs. 3(d–f) and 4(d – f) are the reference spectra obtained from standard materials to apply fingerprint technique for core-loss fine structures [19].
5. Conclusions Effect of deformation process on grain boundaries, particularly on sintering additives were investigated by using both HRTEM and STEM. HRTEM images presented that the presence of amorphous secondary phase at the grain boundaries of as-sintered specimen, disappearance of amorphous phase at the grain boundaries of the compressed specimen and of the annealed specimen. Although EDS results clearly showed the presence of Al and O at grain boundaries from both as-sintered and compressed specimens, EELS analysis confirmed that the chemical environments of them are different. The disappearance of the glass phases may be caused by the vaporization by the application of high temperature in both the compressed and the annealed specimen, though Al and O atoms still remained in the case of the compressed specimen. Sintering conditions influences the additives at the grain boundaries by altering the phases of sintering additives. This leads to the change of mechanical properties in macroscopic sense. Although it is still too early to conclude that the effects of sintering additives at the grain boundaries are directly related to the deformation mechanisms or not, specimens with other sintering additives suitable for SiC (either boron or carbon, or mixture of Y2O3 and Al2O3) must be studied to compare with the results obtained from both macroscopic and microscopic investigations in the future. Addition to using different sintering additives, it is suggested to carry out other series of analysis by electron microscopes on specimens with deformation at low temperatures (so that there are no grain growth during the deformation process), and specimens with static growth (so that there are no stress but with grain growth). The grain nucleation might have occurred with the grain boundary sliding as dynamic recrystallization during deformation. .
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The work presented here distinctively demonstrated the combination of HRTEM imaging and AEM analysis as a powerful tool for material investigations, specially to investigate structures and chemistry of the grain boundaries. Furthermore, it is advantageous and recommended to compare experimental results with theoretical modeling of the grain boundary structures, like ab initio calculation [20], and theoretical calculations of electronic structures at the grain boundaries, molecular orbital method [22] and multiple scattering method [23]. Acknowledgements We would like to acknowledge Dr H. Kodama from Hitachi research Lab. for providing specimen, Dr Tsurekawa (Tohoku University) and Dr K. Kakimoto (JST) for valuable discussions. References [1] R.A. Alliegro, L.B. Coffin, J.R. Tinklepaugh, J. Am. Ceram. Soc. 39 (1956) 386. [2] F.F. Lange, J. Mater. Sci. 10 (1975) 314. [3] M.H. Lewis, R.J. Lumby, Powder Metall. 26 (1983) 73. [4] D.H. Stutz, S. Prochazka, J. Lorenz, J. Am. Ceram. Soc. 68 (1985) 479. [5] T.B. Jackson, A.C. Hurford, S.L. Bruner, R.A. Cutler, in: J.D. Cawley, C.E. Semler (Eds.), Silicon Carbide ‘87, American Ceramics Society, Westerville, (1989) 309. [6] H. Kodama, T. Miyoshi, J. Am. Ceram. Soc. 73 (1990) 3081. [7] K. Suzuki. In: S. Somiya, Y. Inomata (Eds), Silicon Carbide Ceramics-2, Elsevier, Amsterdam, p. 162. [8] T. Kinoshita, S. Munekawa, S.-I. Tanaka, Acta. Metall. Mater. 45 (1997) 801. [9] L.K.L. Falk, J. Euro. Ceram. Soc. 17 (1997) 983. [10] A. Tsuge, K. Nishida, M. Komatsu, J. Am. Ceram. Soc. 58 (1975) 323. [11] T. Watanabe, S. Kimura, S. Karashima, Phil. Mag. A. 49 (1984) 845. [12] M.A. Mulla, V.D. Krstic, Acta. Metall. Mater. 42 (1994) 303. [13] H. Kodama, T. Miyoshi, Adv. Ceram. Mater. 3 (1988) 177. [14] Y. Takeda, K. Nakamura, K. Maeda, M. Ura, Adv. Ceram. Mater. 1 (1986) 162. [15] Y. Tajima, W.D. Kingery, J. Am. Ceram. Soc. C 65 (1982) 27. [16] K. Kaneko, I. Tanaka, M. Yoshiya, Appl. Phys. Lett. 72 (1998) 191. [17] K. Kaneko, T. Gemming, I. Tanaka and H. Mullejans, Phil. Mag. A. (in press). [18] D.C. Joy, In: D.C. Joy, A.D.J. Roming, J.I. Goldstein (Eds.), Principles of Analytical Electron Microscopy, Plenum Press, New York, (1989). [19] R.F. Egerton, Electron Energy Loss spectroscopy in the Electron Microscope, 2nd edition, Plenum Press, New York, 1996. [20] P.R. Kenway, J. Am. Ceram. Soc. 77 (1994) 349. [21] T. Grande, H. Sommerset, E. Hagen, K. Wiik, M-A. Einarstrud, J. Am. Ceram. Soc. 80 (1997) 1047. [22] I. Tanaka, H. Adachi, J. Phys. D: Appl. Phys. 29 (1996) 1725. [23] R. Brydson, J. Micros. 180 (1995) 238.