Silicon nitride and sialon ceramics

Silicon nitride and sialon ceramics

Current Opinion in Solid State and Materials Science 4 (1999) 453–459 Silicon nitride and sialon ceramics Anatoly Rosenflanz* 3 M Company, 3 M Center...

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Current Opinion in Solid State and Materials Science 4 (1999) 453–459

Silicon nitride and sialon ceramics Anatoly Rosenflanz* 3 M Company, 3 M Center, Building 201 -3 N-06, St. Paul, MN 55144 -1000, USA

Abstract The most intriguing recent development in the field of silicon nitride ceramics has undoubtedly been the discovery of a cubic form of silicon nitride. Major advances were made in a-SiAlON ceramics, including the development of thermally stable, in situ reinforced grades. Significant achievements were reported in tailoring the mechanical properties of silicon nitride ceramics through control of secondary phase chemistry and grain morphology.  2000 Elsevier Science Ltd. All rights reserved.

1. Introduction Silicon nitride and its derivative sialon ceramics continue to be subjected to the sustained attention of ceramics research groups worldwide. Although earlier hopes of utilizing the excellent thermomechanical properties of these materials in such demanding applications as gas turbines were not realized, the advancements of the last decade have led to implementation of silicon nitride materials in cutting tools, bearings, turbochargers, engine valves and other wear-resistant components. More applications are expected in the near future as exemplified by the title of the 1998 ACerS meeting Orton memorial lecture — ‘Commercialization of Advanced Structural Ceramics — Patience is a Necessity’ [1]. Basically, Si 3 N 4 ceramics possess several shortcomings which must be surmounted if the range of their applications is to expand. Among them are insufficient resistance to unpredictable brittle fracture, degradation of excellent mechanical properties at elevated temperatures and insufficient oxidation resistance. Typical silicon nitride ceramics developed to date to counteract one or the other deficiency contain either b-Si 3 N 4 , b9-SiAlON, a9-SiAlON or a mixture of a9 / b9. b-Si 3 N 4 ceramics provide the best fracture resistance through development of microstructure containing whisker-like grains capable of deflecting or bridging a crack. Historically, fracture resistance of these ceramics was believed to be governed mainly by the morphological *Tel.: 11-651-737-0451; fax: 11-651-575-1942. E-mail address: [email protected] (A. Rosenflanz)

features of the resultant microstructure, i.e. the grain diameters and aspect ratios of whisker-like grains [2]. Only a few publications attempted to clarify the effect of secondary phase chemistry on fracture properties of these ceramics [3]. Sialon ceramics that are isostructural with either bSi 3 N 4 (b9-SiAlON) or a-Si 3 N 4 (a9-SiAlON) offer the advantage of incorporating some of the sintering additives into the Si 3 N 4 lattice, thus reducing the overall amounts of secondary phase and potentially improving high temperature properties. Traditionally, despite being harder than b-Si 3 N 4 ceramics, single phase a9-SiAlON materials were not considered as serious candidates for structural applications because of their highly brittle nature, resulting from their inability to form acicular grains. In addition, most a9-SiAlONs stabilized by rare-earth cations were found to be unstable at temperatures typically expected in demanding applications, i.e. around 14008C, where they converted to softer b9-SiAlON and undesirable intergranular phase [4]. In the time frame of this review, several studies have conclusively documented that in addition to the effect of grain morphology on the mechanical properties of b-Si 3 N 4 ceramics, the role of secondary-phase chemistry must not be underestimated and in some cases could be overriding. At the same time, in situ reinforced a9-SiAlON ceramics, which combine both high hardness and high toughness, were discovered. Questions regarding the stability of a9SiAlON were clarified and shown to be related to changes of phase diagram with temperature. The large number of new developments in a9-SiAlON ceramics seems especially impressive when compared to the most recent review on

1359-0286 / 00 / $ – see front matter  2000 Elsevier Science Ltd. All rights reserved. PII: S1359-0286( 00 )00004-8


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these materials published in 1991 [5]. Finally, the very intriguing discovery of a superhard cubic form of silicon nitride was unveiled.

2. b-Si 3 N 4 / b9-SiAlON The common thread of the latest works in these ceramics was the establishment of separate links between mechanical properties and grain morphology, on the one hand, and the characteristics of glass / grain interface as dictated by the secondary phase chemistry, on the other hand. Even though these questions had been addressed in literature in the past [3,6], they were never clarified to the extent that they have been very recently. In a first part of a two-part article addressing the first link, i.e. propertymicrostructure, Becher et al. [*7] reported superior K IC , steeply rising R-curve and fracture strength in materials with a distinct bimodal grain diameter distribution. Microstructures with a large number of elongated grains but with broad (monomodal) grain diameter distribution are less favorable, resulting in lower short crack fracture resistance, flatter R-curve and lower strength. Seeding with b-grains with an average size exceeding that of the starting powder was shown to be one of the possible methods of generating the distinctly bimodal grain diameter distribution [*7,8]. Alternatively, as reported in Ref. [9] decreasing the initial a-Si 3 N 4 powder size below a threshold value of 0.3 mm also resulted in distinctly bimodal grain size distribution. Both of these methods essentially rely on generation of grains of different sizes at the beginning of grain growth stage, which provides a driving force for abnormal grain growth. In resolving the question of the combined role of grain morphology and secondary phase chemistry, Kleebe et al. [10] reported an increase in fracture toughness with the increasing size of b-grain diameters, with no evidence of any correlation to aspect ratios of b-grains. Additive composition was shown to be the determining factor in microstructure development as grain coarsening occurred via diffusion-controlled Ostwald ripening. In addition, the authors established that crystallization of secondary phase in triple junctions can contribute to higher K IC by setting up residual tensile stresses along grain boundaries. Lastly, the choice of sintering aids was shown to affect glass network structure along grain boundaries, altering debonding behavior and fracture resistance. The second part of the above-mentioned article [**11] further expanded on the role of the additive composition in the debonding behavior of a glass / grain interface during crack propagation. The authors found that samples processed with higher Y 2 O 3 :Al 2 O 3 ratios exhibited better fracture resistance, which was attributed to a greater tendency for formation of epitaxial b9-SiAlON on b-Si 3 N 4 grains in compositions with lower ratios of Y 2 O 3 :Al 2 O 3 .

(Formation of b9-SiAlON in compositions with lower Y 2 O 3 :Al 2 O 3 ratios is in agreement with known phase relationships in this system since such compositions are closer to the base of the Janecke prism and, therefore, to b9-SiAlON.) The detrimental nature of the epitaxial layer of b9-SiAlON on b-Si 3 N 4 was documented when incorporation of Al–O in b-Si 3 N 4 lattice was linked to establishment of matched bonding with oxynitride glass and, therefore, strengthened interface [12]. Interestingly, the strength of this interface as inferred from debonding studies is very sensitive to the z value (Al–O incorporation) in b9-Si 62z Al z O z N 82z , increasing rapidly with the increase in z from 0 to 0.2 and then leveling off [**11,13]. By taking precautions to exclude the b9-SiAlON / glass interface effect, the debonding behavior between b-whiskers and Ln-bearing oxynitride glasses (Ln5La, Yb) was further analyzed to clarify the effect of rare-earth cation size, and it was found that the interfacial debond length increases as smaller Yb is replaced by La [13]. Thus, interfacial bonding between b-Si 3 N 4 whisker and Yb-glass is stronger than that of b-whisker and La-glass. As a result of the above studies, summarized pictorially in Fig. 1, we now understand that there is much more behind a fracture-resistant silicon nitride than microstructure containing whisker-like grains. While the positive effect of bimodal grain diameter distribution was previously known [14] we now understand that the characteristics of the glass / grain interface as determined by bonding matching (in turn a function of phase relationships) are equally important. An additional role that liquid phase can play in microstructure development was demonstrated by Wang et al. [*15], who reported a new growth mechanism for Si 3 N 4 crystals in a liquid environment. This mechanism involves atomic diffusion through the liquid to the side surface of a growing crystal and further migration of these atoms to the end caps, feeding axial grain growth. In this scenario, the increased amount of intergranular phase is directly linked to enhanced grain growth in the axial direction; as the side surfaces of the crystal become more exposed to the liquid, they act as an additional pathway for matter transport to the end caps. The techniques described above which improve the fracture resistance of silicon nitride ceramics unfortunately tend to compromise their high-temperature properties. Comparison of room and high temperature properties of Si 3 N 4 prepared with different amounts of Y 2 O 3 :SiO 2 ratio [16] showed that the best creep resistance can be achieved in Si 3 N 4 sinter-hiped with 2 wt.% of Y 2 O 3 . However, the room temperature properties of this material were found to be inferior to those of materials prepared with a higher amount of sintering aids. In this study, the authors also found that additive compositions leading to crystallization of secondary phase greatly inhibited densification at the sintering stage. Analogous observations of the detrimental effects of secondary phase crystallization on densification

A. Rosenflanz / Current Opinion in Solid State and Materials Science 4 (1999) 453 – 459


Fig. 1. A pictorial summary of the combined effect of: (a) microstructure, and (b) secondary phase chemistry on fracture resistance of b-ceramics. Information for this figure is drawn mainly from Refs. [*7,10] and [13].

were made in sintering studies of a9-SiAlON ceramics [17]. A trade-off between room and high temperature properties is unfortunate, and one would hope to find a Si 3 N 4 ceramic combining both excellent room and high temperature performance. In this light, the development of in situ reinforced a9-SiAlONs appears to be quite timely and promising.

3. a9-SiAlON

3.1. In situ reinforced a9 -SiAlONs The first reports documenting the occurrence of elongated grains in a9-SiAlONs [18,19] proved that anisotropic grain growth can occur in these materials under appropriate conditions. Indeed, formation of a-Si 3 N 4 grains of different morphologies, including whisker, needle, prism and platelets was observed as early as 20 years ago [20].

However, no significant improvements in the fracture toughness of this ceramic were reported until very recently. Shen et al. [**21] demonstrated that formation of elongated grains is promoted by a high fraction of transient liquid phase. They found that Nd and Sm systems were more likely to lead to the formation of elongated a9SiAlONs than Y and Yb systems. Differences in the amount of transient phase were proposed as the cause of this phenomenon (more transient liquid exists at sintering temperatures in systems with larger stabilizing cation). Improvements in fracture toughness were reported for some compositions. Chen and Rosenflanz [**22] reported that a9-SiAlON material with elongated microstructure that has improved fracture resistance can be prepared in a variety of compositions by practicing nucleation control. The development of elongated a9-grains was traced to the nucleation rates of a9-SiAlON phase which depended on: (a) choice of starting powder; (b) type of stabilizing cation; and (c) temperature control of nucleation rates. Essentially, a9-SiAlONs with elongated grains can be prepared by


A. Rosenflanz / Current Opinion in Solid State and Materials Science 4 (1999) 453 – 459

practicing nucleation control in order to achieve a small population of relatively large a9-grains which can be coarsened without concurrent transformation and production of new grains, i.e. minimization of grain impingement. Alternatively, such techniques as seeding with pre-existent a9-grains to confine nucleation events to a small fraction of externally introduced nucleation sites were shown to lead to the formation of elongated grains as well [*23,*24] (analogous to the seeding that commonly occurs in bSi 3 N 4 ceramics). Supporting this concept of nucleationcontrolled microstructure evolution, Wang et al. [25] recently showed that b-Si 3 N 4 ceramics in situ reinforced with elongated Si 2 N 2 O grains were a result of low nucleation rate and fast grain growth of Si 2 N 2 O phase. Furthermore, kinetics of all phase transformation occuring in SiAlON systems, discussed by authors in terms of traditional theory of solidification from melt, was shown to be governed by the nucleation rates of product phases, which were determined in turn to be a function of relative stability and structural correspondence of parent and product phases [26,27].

3.2. Ca-a9 -SiAlONs In the grain growth stage of microstructural evolution, grain impingement would likely depend on the amount of intergranular glassy phase. Zhao et al. [*28] and Wood et al. [29] have recently shown that a9-grains with elongated morphologies can readily be obtained in Ca-stabilized a9-ceramics with compositions rich in AlN and Al 2 O 3 . These materials were found to have fracture toughness values exceeding 6 Mpa m 1 / 2 , but very low hardness values, which are due to large quantities of glassy phase. Wang et al. [30] further investigated a range of compositions lying on the Si 3 N 4 -CaO:3AlN line with x in Ca x Si 122x Al 3x O x N 162x ranging from 0.3 to 2.0 and also found that as composition shifted toward the CaO:3AlN point, fracture toughness increased and hardness decreased. This is to be expected, since more elongated grains grow at higher x values because there is more liquid phase present. In light of these new improvements in the mechanical performance a9-SiAlONs, clarification of the thermal stability of these materials becomes ever more important. In Ca-system most works established a large compositional range of stability of Ca-a9-SiAlON, which appears to extend further toward the Al-rich side of the a9-plane than most of the other metal oxide SiAlON systems [30,31]. This stable character of Ca-a9-phase further manifests itself in the ability of the material to resist reverse a9→b9 transformation. For instance, Hewett et al. [32] demonstrated that even compositions containing b9 and large proportions of glassy grain boundary phase did not exhibit reverse transformations up to 24 h at temperatures ranging between 11008C and 15008C. Furthermore, Mandal et al. [33,34] found that a9-SiAlONs stabilized with both Nd and Ca exhibited much greater resistance to reverse trans-

formation as compared to Nd-a9-SiAlON, once again proving the stabilization ability of Ca cation. The positive effect of stabilizing a9-SiAlON with both large and small rare-earth cations (Dy and Sm) was also demonstrated by Zhang et al. [35]. Interestingly, even cations that are typically considered too large (Sr and La) can enter a9SiAlON structure when co-doped with Ca [18,36].

3.3. Reverse transformations While Ca-a9-SiAlONs do appear to resist reverse transformation, most other a9-SiAlONs suffer from it. A number of studies in the past five years have amassed evidence demonstrating that reverse transformation (a9→b9) in a9-SiAlON materials occurs more readily in systems stabilized with larger rare-earth cation, higher amounts of liquid phase with lower viscosity and in compositions containing some b9-SiAlON (i.e. outside of the range of stability of single phase a9-SiAlON) [37–39]. All of these observations have recently been explained and rationalized in terms of changes with temperature in the phase relationships in materials with compositions near the Si 3 N 4 corner [**40]. The authors, using Yb, Y and Nd as modifiers, demonstrated that the size of single phase of a9-SiAlON decreases not only with an increase in size of the stabilizing cation, also documented by other researchers [41,42], but also with a decrease in temperature. In addition, all of the phase lines on the a9-plane were determined to be a function of size of stabilizing cation and temperature, shifting toward the Si 3 N 4 corner with the increase in cation size and decrease in temperature (see Fig. 2). This decrease in stability of a9-phase with temperature decrease supplies direct thermodynamic driving force for reverse transformation. The magnitude of this driving force is greater for larger cations, which explains the greater tendency of a9-sialons with larger cations to transform. In addition, systems with larger stabilizing cations very likely have a kinetic advantage for reverse transformation due to the presence of larger amounts of intergranular phase. Decreased solubility (x) in M x Si 122x Al 3x O x N 162x with lower temperature was also ruled to be a determining factor in the occurrence of reverse transformation in a recent study of Mitomo and Ishida [43]. Summarizing the current understanding of the issue of reverse transformation in a9-SiAlON ceramics one can say that while the sheer number of various experimental observations can leave a novice in the field overwhelmed with its complexity, the underlying reason for it is now clear: (1) stability of a9-phase is temperature and composition dependent; (2) manifestations of these variations of stability of a9-SiAlON are governed by transformation kinetics which are a function of nucleation (pre-existent b9 grains) and growth (type and amount of intergranular phase) rates.

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Therefore, the discovery of an even harder form of Si 3 N 4 was truly intriguing news, even though it is currently of purely academic interest.

4. Cubic Si 3 N 4 In 1999, in Nature, Zerr et al. [**46] reported the exciting discovery of cubic Si 3 N 4 synthesized at high temperature (2200 K) and high pressure (15 GPa). This new form of Si 3 N 4 has a spinel structure in which nitrogen atoms adopt a cubic, closed-packed configuration and silicon atoms occupy interstitial sites in both tetrahedral and octahedral coordination in a 1:2 ratio. Preliminary results revealed metastability of c-Si 3 N 4 to at least 700 K at ambient pressure. The calculated bulk (300 GPa) and shear c44 (350GPa) moduli for c-Si 3 N 4 exceeded those of a and b-structures, being quite similar to those of stishovite, a high-pressure phase of SiO 2 . It is therefore expected that the hardness of c-Si 3 N 4 will be close to that of stishovite, i.e. on the order of 33 GPa. Such high hardness values obviously suggest that this new form of Si 3 N 4 should offer significant advantages in industrial applications where hardness and wear resistance are crucial. However, the economic feasibility of producing c-Si 3 N 4 and its thermal and pressure metastability must be addressed first.

5. Conclusion Fig. 2. Variation of compositional range of stability of single phase a9-SiAlON with temperature. (a) Large stability range is evident at high temperature. (b) Much decreased stability is apparent at lower temperature. Note also the shift in a9-b-tie lines toward the Si 3 N 4 -RN:3AlN line with temperature decrease. Data is taken from Ref. [**40] for Nd system.

Interestingly, thermal instability of a9-SiAlON modified by large cations (Sm, Nd) was recently linked to decreased oxidation resistance of these materials. It was shown that release of rare-earth cations during the reverse transformations led to the formation of oxidation-prone mellilite phase [44]. Study of high temperature properties of a9 / b9SiAlON ceramics [45] demonstrated improved oxidation resistance but worse creep properties of materials with less intergranular glassy phase and decreased a9-SiAlON content. This seemingly contradictory observation (usually, less glass means better creep properties) could also be caused by the reverse transformation effect during oxidation studies. The surge of activity in a9-research, described above, is partially fueled by the possibility of utilizing the increased hardness of a9-SiAlON phase compared to b-Si 3 N 4 .

The collection of the reviewed works outlined the active involvement of ceramics research in developing ever better Si 3 N 4 -based materials. Major advancements in both bSi 3 N 4 / b9-SiAlON and a9-SiAlON materials were presented in this review. It was shown that mechanical properties of b-Si 3 N 4 based ceramics are governed by the combined effect of microstructure assemblage and secondary phase chemistry, each having a distinct contribution. These findings may prove instrumental in future development of other materials in the silicon nitride family that contain intergranular phase. In particular, a9-SiAlON seems to be a good candidate material to consider. Recent discoveries, described in this review, proved the possibility of forming in situ reinforced microstructures in a9SiAlONs and verified the importance of additive composition. Future developments in a9-SiAlONs thus can focus on clarification of the role of secondary phase chemistry and composition of a9-SiAlON near grain / glass interface on debonding, for instance. Good toughness and ease of fabrication of some Ca-a9-SiAlONs certainly merits further studies in improving hardness of these materials. Finally, the possibilities of utilizing a superhard cubic Si 3 N 4 are intriguing indeed.


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Acknowledgements I would like to thank Ross Plovnick for reviewing the manuscript.

References Papers of particular interest, published within the annual period of review, have been highlighted as: * of special interest; ** of outstanding interest. [1] Savitz M. Commercialization of advanced structural ceramics — Patience is a necessity. Am Ceram Soc Bull 1999;78(1):53–6. [2] Kawashima T, Okamoto H, Yamamoto H, Kitamura A. Grain size dependence of the fracture toughness of silicon nitride ceramics. J Ceram Soc Jpn 1991;99(4):320–3. [3] Peterson IM, Tien T-Y. Effect of the grain boundary thermal expansion coefficient on the fracture toughness in silicon nitride. J Am Ceram Soc 1995;78(9):2345–52. [4] Ukyo Y, Sugiyama N, Wada S. Thermal stability of Y-a9-SiAlON coexisting with b9-SiAlON. In: Kimura S, Niihara K, editors. Proc. 1st Intl Symp Sci Eng Ceram. Ceram Soc Jpn, 1991:141–6. [5] Cao GZ, Metselaar R. a9-SiAlON ceramics, a review. Chem Mater 1991;3:242–52. [6] Hoffmann MJ. Analysis of microstructural development and mechanical properties of Si 3 N 4 ceramics. In: Hoffmann MJ, Petzow G, editors. Tailoring of mechanical properties of Si 3 N 4 ceramics, Dordrecht: Kluwer Academic Publishers, 1994:59–72. [*7] Becher PF, Sun EY, Plucknett KP, Alexander KB, Hsueh C-H, Lin H-T, Waters SB, Westmoreland CG, Kang E-S, Hirao K, Brito ME. Microstructural design of silicon nitride with improved fracture toughness: I. Effects of grain shape and size. J Am Ceram Soc 1998;81(11):2821–30. This study delineates the effect of microstructure on mechanical properties of b-ceramics without complications resulting from a glass / grain interface effect. [8] Dressler W, Kleebe HJ, Hoffmann MJ, Ruhle M, Petzow G. Model experiments concerning abnormal grain growth in silicon nitride. J Eur Ceram Soc 1996;16:3–14. [9] Lee CJ, Kim DJ. Effect of a-Si 3 N 4 particle size on the microstructural evolution of Si 3 N 4 ceramics. J Am Ceram Soc 1999;82(3):753–6. [10] Kleebe HJ, Pezzotti G, Ziegler G. Microstructure and fracture toughness of Si 3 N 4 ceramics: combined roles of grain morphology and secondary phase chemistry. J Am Ceram Soc 1999;82(7):1857– 67. [**11] Sun EY, Becher PF, Plucknett KP, Hsueh C-H, Alexander KB, Waters SB, Hirao K, Brito ME. Microstructural design of silicon nitride with improved fracture toughness: II. Effects of yttria and alumina additives. J Am Ceram Soc 1998;81(11):2831–40. Very important and interesting findings regarding the role of additive composition in determining the overall mechanical properties of b-ceramics are presented in this work. The results may prove far-reaching in terms of applicability to other silicon nitride materials. [12] Sun EY, Alexander PF, Becher PF, Hwang S-L. b-Si 3 N 4 whiskers embedded in oxynitride glasses: interfacial microstructure. J Am Ceram Soc 1996;79(10):2626–32. [13] Sun EY, Becher PF, Hsueh C-H, Painter GS, Waters SB, Hwang S-L, Hoffmann MJ. Debonding behavior between b-Si 3 N 4 whiskers and oxynitride glasses with or without an epitaxial b-SiAlON interfacial layer. Acta mater 1999;47(9):2777–85.

[14] Ohji T, Hirao K, Kanzaki S. Fracture resistance behavior of highly anisotropic silicon nitride. J Am Ceram Soc 1995;78(11):3125–8. [*15] Wang LL, Tien T-Y, Chen I-W. Morphology of silicon nitride grown from a liquid phase. J Am Ceram Soc 1998;81(10):2677–86. An interesting work describing a new mechanism for grain growth in silicon nitride materials. [16] Hoffmann MJ, Geyer A, Oberacker R. Potential of the sinter-HIPtechnique for the development of high-temperature resistant Si 3 N 4 ceramics. J Eur Ceram Soc 1999;19(13–14):2359–66. [17] Bandyopadhyay S, Hoffmann MJ, Petzow G. Effect of different rare-earth cations on the densification behaviour of oxygen rich a-SiAlON composition. Ceramics International 1999;25:207–13. [18] Hwang CJ, Susintzky DW, Beaman DR. Preparation of multication a9-SiAlON containing strontium. J Am Ceram Soc 1995;78(3):588– 92. [19] Shen Z, Ekstrom T, Nygren M. Temperature stability of samariumdoped a9-SiAlON ceramics. J Eur Ceram Soc 1996;16:43–53. [20] Niihara K, Hirai T. Growth, morphology and slip system of a-Si 3 N 4 single crystal. J Mater Sci 1979;14:1952–60. [**21] Shen ZJ, Nordberg L-O, Nygren M, Ekstrom T. a9-SiAlON grains with high aspect ratio — utopia or reality? In: Babini GN, editor. Proc. Nato AST Engineering Ceramics ’96 — Higher Reliability through Processing. Dordrecht: Kluwer Academic Publishers; 1997:169–178. A remarkable report on the formation of in situ reinforced a9SiAlON with improved fracture resistance. Roles of large amounts of liquid phase and large cations are outlined. [**22] Chen I-W, Rosenflanz A. A tough SiAlON ceramic based on a-Si 3 N 4 with a whisker-like microstructure. Nature 1997;389:701– 4. First work to outline main principles of formation of elongated microstructure in a9-SiAlON ceramics. Much improved mechanical properties are reported for these ceramics. [*23] Kim J, Rosenflanz A, Chen I-W. Microstructure control of in situ toughened a-SiAlON ceramics. J Am Ceram Soc, in press. First observation of applicability of seeding concept to a9-SiAlON ceramics. Improved techniques can lead to substantial improvements in properties. [*24] Rosenflanz A. a9-SiAlON: Phase stability, phase transformations and microstructural evolutions. PhD Thesis, Univ Michigan 1997. An important thesis delineating both the microstructural development of a9-SiAlON ceramics and their thermal stability. [25] Wang C, Emoto H, Mitomo M. Nucleation and growth of silicon oxynitride grains in a fine-grained silicon nitride matrix. J Am Ceram Soc 1998;81(5):1125–32. [26] Rosenflanz A, Chen I-W. Kinetics of phase transformations in SiAlON ceramics: I. Effects of cation size, composition and temperature. J Eur Ceram Soc 1999;19(13-14):2325–35. [27] Rosenflanz A, Chen I-W. Kinetics of phase transformations in SiAlON ceramics: II. Reaction paths. J Eur Ceram Soc 1999;19(13– 14):2337–48. [*28] Zhao H, Swenser SP, Cheng Y-B. Elongated a-SiAlON grains in pressureless sintered SiAlON ceramics. J Eur Ceram Soc 1998;18(8):1053–7. A noteworthy work on elongated Ca-a9-SiAlONs. Fracture toughness exceeding 6 Mpa m 1 / 2 is reported. [29] Wood CA, Zhao H, Cheng Y-B. Microstructural development of calcium a-SiAlON ceramics with elongated grains. J Am Ceram Soc 1999;82(2):421–8. [30] Wang PL, Zhang C, Sun WY, Yan DS. Characteristics of Ca-aSiAlON-phase formation, microstructure and mechanical properties. J Eur Ceram Soc 1999;5(5):553–60. [31] Hewett CL, Cheng Y-B, Muddle BC. Phase relationships and related microstructural observations in the Ca–Si–Al–O–N system. J Am Ceram Soc 1998;81(7):1781–8. [32] Hewett CL, Cheng Y-B, Muddle BC, Trigg MB. Thermal stability of calcium a-SiAlON ceramics. J Eur Ceram Soc 1998;18(4):417–27.

A. Rosenflanz / Current Opinion in Solid State and Materials Science 4 (1999) 453 – 459 [33] Mandal H, Thompson DP. a(b-Sialon transformation in calciumcontaining a-SiAlON ceramics. J Eur Ceram Soc 1998;5(5):543–52. [34] Mandal H. New developments in a-SiAlON ceramics. J Eur Ceram Soc 1999;19(13–14):2349–57. [35] Zhang C, Sun WY, Yan DS. Optimizing mechanical properties and thermal stability of Ln-a-b-SiAlON by using duplex Ln elements (Dy and Sm). J Eur Ceram Soc 1999;19(1):33–9. [36] Mandal H, Hoffmann MJ. Preparation of multiple-cation a-SiAlON ceramics containing lanthanum. J Am Ceram Soc 1998;82(1):229– 32. [37] Mandal H, Thompson DP, Ekstrom T. Reversible a(b-SiAlON transformation in heat-treated SiAlON ceramics. J Eur Ceram Soc 1993;12:421–9. [38] Shen Z, Ekstrom T, Nygren M. Temperature stability of samariumdoped a9-SiAlON ceramics. J Eur Ceram Soc 1996;16:43–53. [*39] Camuscu N, Thompson DP, Mandal H. Effect of starting composition, type of rare-earth sintering additive and amount of liquid phase on a(b-SiAlON transformation. J Eur Ceram Soc 1997;17:599–613. An important study summarizing most of the experimental observations regarding reverse transformation in a9-SiAlON ceramics. [**40] Rosenflanz A, Chen I-W. Phase relationships and stability of a9SiAlON. J Am Ceram Soc 1999;82(4):1025–36. First study to explain the phenomenon of reverse transformation by



[43] [44]




variation of a9-stability range with temperature. It is shown that decreased stability of a9-SiAlON with temperature decrease is the direct cause for this transformation. All phase relationships are found to depend on temperature. Shen Z, Nygren M. On the extension of the a-SiAlON phase area in yttrium and rare-earth doped systems. J Eur Ceram Soc 1997;17(13):1639–45. Huang Z-K, Tien T-Y, Yen T-S. Subsolidus phase relationships in Si 3 N 4 –AlN–rare-earth oxide systems. J Am Ceram Soc 1986;69(10):C241–2. Mitomo M, Ishida A. Stability of a-SiAlONs in low temperature annealing. J Eur Ceram Soc 1999;19(1):7–15. Nordberg L-O, Nygren M. Stability and oxidation properties of RE-a-SiAlON ceramics (RE5Y, Nd, Sm, Yb). J Am Ceram Soc 1998;81(6):1461–70. Klemm H, Herrmann M, Reich T, Schubert C, Frassek L, Wotting G, Gugel E, Nietfeld G. High-temperature properties of mixed a- / b9-SiAlON materials. J Am Ceram Soc 1998;81(5):1141–8. Zerr A, Miehe G, Serghiou G, Schwarz M, Kroke E, Riedel R, Fuess H, Kroll P, Boehler R. Synthesis of cubic silicon nitride. Nature 1999;400:340–2. The truly remarkable discovery of a cubic form of Si 3 N 4 is discussed in this paper. It remains to be seen whether this form of Si 3 N 4 will be of practical use in industrial applications.