JOURNAL OF MATERIALS SC1ENCE LETTERS 12 (1993) 70-72

Transmission electron microscopy characterization of a ceria-fluxed silicon nitride H.-J. KLEEBE, M. K. C I N I B U L K

Max-Planck-lnstitut für Metallforschung, Institut für Werkstoffwissenschaft, Seestraße 92, D-7000 Stuttgart 1, FRG

The main interest in silicon nitride ceramics is in their application as structural materials at elevated temperatures. Si3N4 ceramics consist of highly refractory grains surrounded by a less refractory vitreous phase due to liquid-phase sintering that occurs during densification. Therefore, the hightemperature behaviour depends primarily on the volume fraction, composition, distribution and crystalline state of the intergranular phase. Orte method of improving the high-temperature performance of this material is to form a more refractory intergranular phase. The substitution of lanthanide oxides for MgO, and Y203 or Y203 + A1203 as sintering additives has been investigated as a method of improving the refractoriness of this intergranular phase [1-11]. In addition to the refractoriness of the intergranular phase present at multiple-grain junctions, the thickness of the amorphous film located at grain and phase boundaries is important with respect to high-temperature properties [12-14]. It has been shown both theoretically [13] and experimentally [15] that the thickness of the intergranular film depends on the composition of the adjacent grains as well as on the composition of the residual amorphous phase itself. In this letter we report on the microstructure of a CeO2-doped Si3N4 ceramic, placing emphasis on the nature of the amorphous phase present at multiple-grain junctions and along two-grain boundaries. The material investigated was fabricated by twostep gas-pressure sintering. The 0:-Si3N4 powder, with 5 wt % CeO2 and an additional i wt % A1203 for improved sinterability, was first sintered at 1840°C for l h under 2 M P a Nz followed by a second-stage sintering at 2050 °C for 1 h under 6 MPa N 2. The processed material had a density of 3.22gcm -3 [16], which is equivalent to 98.1% theoretical density using the rule-of-mixtures and assuming that CeO2 reduces to Ce203 during sintering. This material was reported to exhibit high toughness and good creep behaviour [16]. Foils for electron microscopy were prepared by the standard techniques of grinding, dimpling and ion-beam thinning, followed by coating with a thin layer of carbon to minimize charging under the electron beam. Conventional transmission electron microscopy and analytical electron microscopy were performed using a JEOL 2000FX operated at 200 kV, equipped for energy-dispersive X-ray spectroscopy (EDS) with al~ ultrathin window germanium detector. To determine the nature and thickness of the 70

intergranular films at two-grain boundaries accurately, high-resolution electron microscopy (HREM) was carried out at 400 kV on a JEOL 4000EX (0.18 nm point-to-point resolution). The general microstructure of this material consisted of /3-Si3N4 grains surrounded by an amorphous secondary phase, shown in Fig. 1. The grains were about 0.2-1/xm in diameter across the basal plane, with a small fraction of elongated grains exhibiting aspect ratlos of up to 15. The triple-grain junctions containing the secondary phase were of the order of 20-200 nm in diameter. After a careful search for crystalline secondary phases by tilting the specimen and looking for contrast variations, as well as using microdiffraction and diffuse dark-field imaging, it was concluded that no devitrification of the secondary phase had occurred. During densification CeO2 has been proposed to reduce to Ce203 by the reaction [4] 12CEO2 + ySiO2 + xSi3N4 ~ 6Ce203 + (y + 3)SIO2 + (x - 1)Si3N4 + 2N2 leading to an increase in the volume fraction and Si:Ce ratio of the liquid during sintering. The final solid-state reaction products usually consist of Si3N4 grains and a C e - S i - O - N amorphous phase. The presence of an amorphous secondary phase in the material under investigation is consistent with that reported in the literature [1-3, 5-8] for Si3N4 with small ceria additions when no additional heat treatment is performed to induce crystallization.

Figure 1 Low-magnification bright-field image showing the general microstructure of the CeO2-sintered Si3N4, including high aspect ratio /%Si3N4 grains and Ce-rich amorphous secondary phase.

0261-8028 © 1993 Chapman & Hall

However, Lange [4] and Guha et al. [3] detected Ce-apatite (CesSi3012N), and Ce-apatite plus Ce-wollastonite (CeSiO2N), respectively, by X-ray diffraction, but only in compositions containing at least 10 wt % CeO2. Babini et al. [5] reported the presence of Ce-orthosilicate (Ce4.67Si3013), but also in compositions fluxed with > 10wt% CeO2. Huseby and Petzow [1] reported that a small quantity of CeO2 was present in a hot-pressed material and was attributed to a small fraction of unreacted additive; however, no other crystalline secondary phases were detected. The amorphous nature of the secondary phase in the present material can be seen in Fig. 2. The diffuse dark-field technique was used to detect directly the amorphous phase in bright contrast. The intergranular phase, present at the triple-grain junctions, can also be seen at the grain boundaries. EDS microanalysis of the amorphous phase at multiplegrain pockets indicates that it is a C e - S i - A 1 - O - N glass as shown in Fig. 3a. The quantification of spectra obtained from several multiple-grain junctions indicated the glass phase to be homogeneous throughout, having an approximate composition of 1 5 + 3 a t % N, 5 8 + 3 a t % O, 2 . 5 + 0 . 5 a t % A1, 17.5 + 1.0 at % Si and 7.0 + 0.5 at % Ce. The Ce:A1 ratio of 2.5, based on the relative amounts of CeO2 and A1203 (5 and 1 wt %, respectively) added to the system, compared favourably with the value of 2.8 obtained from the relative amount of Ce and A1 detected by EDS in the amorphous phase. Since a small amount of A1 (about i at %) could be detected to be in solid solution with Si3N4, the concentration of A1 in the intergranular phase was reduced, thereby increasing the Ce:A1 ratio as shown above. The concentration of Ca (as a possible impurity) in the amorphous phase was below the detection limit. Detection of Mg and Fe, also common impurities in the amorphous intergranular phases of Si3N4 ceramics, was prevented since their peaks oveflapped those of Ce. EDS analysis of two-grain boundary fiims revealed the presence of Ce (Fig. 3b).

Figure2 Diffuse dark-field image revealing the amorphous nature of the intergranular phase (in bright contrast). Note inclined grain boundaries (indicated by arrows). The inset is the microdiffraction pattern of the intergranular phase.

Energy (keV)

10.24

(b)

si

Ce

t::i f, co o

Energy (keV)

lO.24

Figure3 EDS microanalyses of (a) the intergranular phase, obtained from the glass located at a triple-grain junction and (b) a two-grain boundary with vertical scale expanded to emphasize the presence of Ce in the intergranular film.

However, since the size of the probe (a minimum of about 20 nm for practical EDS) resulted in sampling primarily the adjacent /3-Si3N4 grains, the relative amount of Ce detected in this area was extremely low (< 0.5 at%), precluding any estimate of its concentration. Moreover, no conclusive results could be obtained regarding the N, O or A1 content of the film at the interface since these elements all are present to some extent in the adjacent Si3N4 grains. The thickness of the grain-boundary film separating two/3-Si3N 4 grains was determined from HREM images. Fig. 4a contains a lattice image of the intergranular film emanating from a three-grain junction. The amorphous intergranular phase seems to acquire a characteristic thickness for a given composition once it enters the grain boundary. Measurements of grain-boundary films in this material, such as that presented in Fig. 4b, indicate an equilibrium film thickness of 0.9-1.0 nm. The two grains shown in Fig. 4 have a misorientation of 67° with respect to the (10 i 0) plane and (12 3 0) plane (lower and upper grains of Fig. 4b, respectively). From the HREM observations the intergranular film thickness appears to be independent of grain misorientation, with the exception of low-energy boundaries or special orientations where no intergranular film may exist [17]. These results are consistent with previous observations indicating that the major parameters influencing the intergranular film thickness are the chemistry of the adjacent grains and the chemical composition of the amorphous film itself. The grain-boundary film thickness 71

Figure 4 (a) H R E M image of the amorphous phase penetrating a two-grain boundary and (b) higher-magnification image of the fl-Si3N4-/3-Si3N4 interface. The inset is the corresponding electron diffraction pattern.

in this material is less than that found in either Y203or Sc~O3-doped Si3N4, which contained film thicknesses of about 1.5 nm [15]. It can be concluded that the good high-temperature behaviour of this CeO2sintered Si3N 4 can be attributed to both the small thickness of the grain-boundary films as well as to their refractoriness due to the presence of Ce.

5. 6. 7. 8. 9.

Acknowledgements

10.

We thank D. R. Clarke for providing the silicon nitride material and M. Rühle for helpful discussions.

11.

References I. 2. 3. 4.

72

I . C . H U S E B Y and G. P E T Z O W , Powder Metall. Int. 6 (1974) 17. K. S. M A Z D I Y A S N I and C. M. C O O K E , J. Amer. Ceram. Soc. 57 (1974) 536. J. P. G U H A , P. G O U R S A T and M. BILLY, ibid. 63 (1980) 119. F . F . L A N G E , Amer. Ceram. Soc. Bull. 59 (1980) 239.

12. 13. 14. 15. 16. 17.

G. N. BABINI, A. BELLOSI and P. V I N C E N Z I N I , Ceram. Int. 6 (1980) 91. W . A . S A N D E R S and T. P. H E R B E L L , J. Amer. Ceram. Soc. 66 (1983) 835. W. A. SANDERS and D. M. M I E S K O W S K I , ibid. 64 (1985) 304. N. H I R O S A K I , A. O K A D A and K. M A T O B A , ibid. 71 (1988) C-144. J, S. V E T R A N O , H.-J. K L E E B E , E. H A M P P , M. J. H O F F M A N N and R. M. CANNON, J. Mater. Sei. Lett. 11 (1992) 1249. M . K . C I N I B U L K , G. THOMAS and S. M. JOHNSON, J. Amer. Ceram. Soc. 75 (1992) 2037. J. S. V E T R A N O , H.-J. K L E E B E , E. H A M P P , M. J. H O F F M A N N , R. M. CANNON and M. R Ü H L E , J. Mater. Sei. submitted. R. RA J, J. Geophys. Res. B87 (1982) 4731. D . R . C L A R K E , J. Amer. Ceram. Soc. 70 (1987) 15. 1dem, ibid. 72 (1989) 1604. H.-J. K L E E B E , J. Eur. Ceram. Soc. in press. D . R . C L A R K E , Private communication (1991). H. SCHMID and M. R Ü H L E , J. Mater. Sei. 19 (1984) 615.

Received 2 January and accepted 6 July 1992