Progress in Sialon Ceramics

Journal of the European Ceramic Society 20 (2000) 2275±2295

V.A. Izhevskiy a,1, L.A. Genova a, J.C. Bressiani a, F. Aldinger b,*

aInstituto de Pesquisas EnergeÂticas e Nucleares, IPEN/CNEN-SP, Trav. R, 400, Cidade UniversitaÂria, SaÄo Paulo, SP-05508-900, Brazil bMax-Planck-Institut fuÈr Metallforschung und UniversitaÈt Stuttgart, Institut fuÈr Nichtmetallische Anorganische Materialien, Heisenbergstrasse 5, D-70569, Stuttgart, Germany

Received 9 June 1999; received in revised form 21 December 1999; accepted 30 January 2000

Abstract In view of the considerable progress that has been made over the last several years on the fundamental understanding of phase relationships, microstructural design, and tailoring of properties for speci®c applications of rare-earth doped SiAlONs, a clear review of current understanding of the basic regularities lying behind the processes that take place during sintering of SiAlONs is timely. Alternative secondary phase development, mechanism and full reversibility of the a0 to b0 transformation in relation with the phase assemblage evolution are elucidated. Reaction sintering of multicomponent SiAlONs is considered with regard of wetting behavior of silicate liquid phases formed on heating. Regularities of SiAlON's behavior and stability are tentatively explained in terms of RE element ionic radii and acid/base reaction principle. # 2000 Elsevier Science Ltd. All rights reserved. Keywords: Phase equilibria; Phase transformations; SiAlONs; Sintering

Contents

1. Introduction...... 2276

2. Occurrence of nitrogen melilite and nitrogen woehlerite solid solutions containing aluminum ...... 2277

3. Subsolidus phase relationships in R2O3±Si3N4±AlN±Al2O3 systems ...... 2282

4. Thermal stability of RE-a-SiAlONs and the reversibility of a0 to b0 transformation ...... 2285

5. Reaction densi®cation of a and (a+b) SiAlONs ...... 2290

6. Conclusions...... 2293

Acknowledgements...... 2293

References ...... 2293

* Corresponding author. E-mail address: [email protected] (F. Aldinger). 1 Invited scientist, on leave from the Institute for Problems of Material Science, National Academy of Sciences of Ukraine, Kiev, Ukraine.

1. Introduction YxSi3(3x+n)-Al(3x+n)OnN4n, with 0.08silicon nitride. Initially, of 0.095 to 0.1 nm, and signi®cantly higher, x=0.25, for they were discovered in the early 1970s13 and have been small cations such as ytterbium with a radius of 0.085 developed actively since. In the following years, fully nm.17 More recent investigations,18,19 showed that the a dense polycrystalline bodies were prepared by pressureless phase area becomes larger with the decreasing size of the sintering techniques. Profound understanding of basic Me ion, i.e. in the order Nd (0.99A)oxygen-rich a- mulated was critically evaluated and summarized in SiAlON phase will be found in the Yb-a-SiAlON system. several reviews.46 Although the initial goal of creating By changing the overall composition in the Y-SiAlON a single-phase silicon nitride-based ceramics without system, (a+b)-SiAlON ceramics can be prepared with a any intergranular amorphous phases due to the tran- varying a-SiAlON:b-SiAlON phase ratio.20, 21 a-SiAlON sient liquid phase sintering was not achieved, SiAlON phase normally appears as equiaxed grains in the ceramics became one of the commercially produced microstructure of the material, while b-SiAlON phase high-tech ceramic materials. forms elongated grains with an aspect ratio of 4 to 10. There are two SiAlON phases that are of interest as However, some recent research of, in particular, Sm- engineering ceramics, a-SiAlON and b-SiAlON, which and Y-doped a-SiAlONs,22, 23 which provides even bet- are solid solutions based on a and b-Si3N4 structural ter possibilities of microstructure tailoring. It is evident modi®cations, respectively, and designated respectively from its chemical composition that a-SiAlON phase as a0 and b0. b-SiAlON is formed by simultaneous oers possibility of reducing the amount of residual equivalent substitution of Al±O for Si±N and has most glassy grain-boundary phase by incorporating con- commonly been described by the formula Si6zAlz- stituents of sintering aids into the crystal structure, OzN8z. In this formula z can be varied continuously which otherwise will be present as substantial amounts from zero to about 4.2.79 The homogeneity region of b- of residual glassy phase and deteriorate the high tem-

SiAlON extends along the Si3N4±(Al2O3AlN) tie line of perature performance of the material. Thus, in principle, the phase diagram. The unit cell of b0 contains two a-SiAlON enables the preparation of the practically

Si3N4 units. The a-SiAlON has a unit cell comprising four glass-free SiAlON composites, which may become Si3N4 units and forms a limited two-dimensional phase interesting candidates for high-temperature structural 4 region in the plane Si3N4 3 Al2O3 ÁAlNMeNÁ3AlN of and engineering applications. the Me±Si±Al±O±N phase diagram. In the latter diagram The possibility of varying the a-SiAlON:b-SiAlON the b-SiAlON falls on one border of the plane mentioned. phase ratio by slightly changing the overall composition Of special interest are the Me±Si±Al±O±N systems where was shown to open many possibilities to prepare the a-SiAlON phase is stabilized by Me ions, such as Y-SiAlON ceramics with desired properties. This is of lithium, magnesium, calcium, yttrium, and the rare-earth the utmost importance in connection with tailoring (RE) metals except lanthanum, cerium, praseodymium, the properties for speci®c applications. Hardness and europium.1012 The homogeneity range composition of increases markedly with increasing a-SiAlON phase a-SiAlON can be described by the general formula content, whereas the fracture toughness decreases. High p+ MemSi12(pm+n)Al(pm+n)OnN16n for a metal ion Me . a-SiAlON content also improves oxidation resistance of Two substitution mechanisms may act; the ®rst is simi- the material and the thermal shock resistance.4 Other lar to that of b-SiAlON with n (Si+N) being replaced ways of changing properties of SiAlON ceramics by by n (Al+O). The second mechanism is further repla- means of replacing Y2O3 by other RE-oxide additives cement of pmSi4+ by pmAl3+, the electron balance were also investigated.24 Precisely this direction of being retained by incorporation of mMep+ into the research lead to discovery of the full reversibility of the phase structure. a0 to b0 transformation for certain phase assemblies The a-SiAlON phase region was most extensively stu- under conditions of post-sintering heat treatment.2527 died in the Y±Si±Al±O±N system, where the two-dimen- It has been observed that the a-SiAlON phase is less sional phase ®eld extension can be expressed as stable at low temperatures and decomposes into rare- V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 2277 earth-rich intergranular phase(s) and b-SiAlON with a N in the bridging position and arranged lengthwise along remarkably elongated crystal morphology. The extent the a-axis. These units are linked by RE polyhedra. of the transformation is more pronounced with Both types of phases are very stable refractory com- increased heat-treatment temperature. The mechanism pounds and occur in almost all rare-earth SiAlON sys- of the fully reversible a0 to b0 transformation was tems.34 With high melting points (1900C for yttrium investigated by the same authors.28 This transformation N-melilite35) and a composition close to a nitrogen-rich provides an excellent possibility for optimizing phase limit of the liquid region, these phases ful®ll some of the content and microstructure without further additions ideal requirements for a grain boundary phase in SiA- of oxides and nitrides merely by heat treatment at lON ceramics. However, the most important distinction appropriately chosen temperatures. In this way, of melilite±woehlerite type of secondary phases is that premeditated values of hardness, strength and tough- they do not form a eutectic with the matrix Si3N4 while ness can be achieved from a single starting composition, the traditional oxide-base secondary phases form one which opens new perspectives for their successful unconditionally. This becomes absolutely obvious if one application. considers phase diagrams of R±Si±(Al/O)±N systems. A The objective of the present review, therefore, is to liquid phase region exists at high temperature on all provide a comprehensive overall view of the new devel- these diagrams, separating the Si3N4 corner and the opments in SiAlONs that will facilitate advances of this plane of additive oxides and their compounds. In con- subject, and allow areas for further study to be more trast, there is no liquid phase region between M(R)/J(R) clearly identi®ed. The review will cover such important and Si3N4 corner on any of the phase diagrams of R±Si± issues as thermal stability of dierent RE-doped (a+b) (Al/O)±N systems. Although some liquid does form SiAlON composites, phase relationships in this systems, during initial heating because of the ternary eutectic of alternative secondary phases for SiAlON ceramics, and SiO2,Al2O3,andR2O3, the liquid should be consumed 0 0 mechanisms of reversible a to b transformation. largely by Si3N4 to form M(R) or J(R). Thus, sintering of the latter compositions is relatively dicult at the intermediate temperature for lack of liquid, but it 2. Occurrence of nitrogen melilite and nitrogen woehlerite improves drastically once the melting point of M(R) solid solutions containing aluminum and J(R) is exceeded. Indeed, the lack of liquid at the intermediate temperature could be advantageous because Nitrogen-containing (or N-)melilites and nitrogen- it reduces the possibility of melt evaporation and massive containing (or N-) woehlerites (N-woehlerites are also powder relocation. Additionally, both melilite and known as J-phases) have been recently closely re-evaluated woehlerite type phases crystallize readily on cooling as the candidates for alternative secondary phases in directly from the liquid, which leads to quicker and

Si3N4-based ceramics after the Al-containing solid solu- much more complete crystallization as compared to any tions of this phases were discovered and investigated.2932 devitri®cation process. The latter occurs in solid state, Both phases were discovered relatively early in the and consequently all diusion processes are by several course of silicon nitride development and at the time orders of magnitude slower than in the liquid. were considered as undesirable secondary phases due to The catastrophic oxidation at intermediate tempera- their poor oxidation resistance at temperature between tures mentioned above can be avoided by formation in 900 and 1100C.33 Under such conditions they oxidize the material of melilite and woehlerite based solid solu- rapidly to crystobalite and y-R2Si2O7 (where R=yttrium tions containing aluminum instead of pure M(R) and or rare earth metal) with a 30% increase in speci®c J(R). The general formulas of these solid solutions are volume, and this results in catastrophic cracking and R2Si3xAlxO3xN4x and R4Si2xAlxO7xN2x, respec- total failure of material. N-melilites have the general tively. The solid solutions will be further referred to as 0 0 formula R2Si3O3N4 and occur as an intermediate phase M (R) and J (R). In these phases, Si±N bonds are pro- during sintering of a number of Si3N4-based ceramics. gressively being replaced by Al±O bonds, leading to the In particular, N-melilite is located in the phase compat- increase of the oxygen content, which in turn improves ibility regions of a(a+b)-SiAlONs, which in theory the oxidation resistance. can produce a good microstructure of SiAlONs N- Until now, the regularities of solid solubility of alu- melilite as the only secondary phase. N-melilite [M(R)], minum in M(R) and J(R) were studied for the majority has a tetragonal structure that is built up of an in®nite of RE elements (Nd, Sm, Gd, Dy, Yb, and Y) that are linkage of corner-sharing Si(O,N)4 units in tetrahedral used in SiAlON production. It was unequivocally sheets perpendicular to the [001] direction. Sandwiched shown that the solid solubility range of M0(R) phases in between these sheets are Y3+ or rare earth Ln3+ ions respect of aluminum are determined by ionic radius of the which hold the silicon oxynitride layers together.34 The RE element. With decreasing ionic radius the solubility structure of N-woehlerite is monoclinic and contains limits in melilite solid solution decrease. For elements either Si2O5N2 or Si2O6N ditetrahedral units containing with large ionic radii, i.e., R=Nd and Sm, up to one Si 2278 V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295

Table 1 Compilation of experimental data on the formation of N-melilite in dierent systems40

Cation Ionic radius r3:r2 Al-free N-melilite Formation Maximum aluminum solubility in M0(R) R O 3 a 0 rR A formation of M (R) La3+ 1.016 0.77 Does not form No study Experimental data currently unavailable Nd3+ 0.995 0.75 Observed Observed x = 1.0 Sm3+ 0.964 0.73 Observed Observed x = 1.0 Dy3+ 0.908 0.69 Observed Observed x = 0.7 Y3+ 0.893 0.68 Observed (not at 1500C) Observed x = 0.7 Yb3+ 0.858 0.65 Does not form No study Experimental data currently unavailable can be replaced by Al. Consequently, yttrium melilite the decrease of the RE ionic radius. The results of the has the lowest observed aluminum solubility (x=0.6) in M0(R) solubility range determination for dierent sys- rare earth elements melilite solid solutions.36 Both ytter- tems are summarized in Table 1.40 bium and lanthanum do not stabilize N-melilite.3739 It The data on the correlation between the x value in was, however, shown that ytterbium forms J0(R) solid M0(R) solid solution formula and the lattice parameters solution that can serve as an alternative secondary of M0(R) were determined for the majority of RE stabi- phase in this system. On the other hand, lanthanum lizing elements.29,36,41 Linear increase of the lattice forms only the M0 phase with the highest substitution parameters with increasing of the x values was observed level (x=1), which is a point phase in La±Si±Al±O±N for all systems investigated up to a certain x value for system, and the M0(La) homogeneity range does not respective M0(R) phases. However, these x values do not exist. Here it must be speci®cally emphasized that phase necessarily represent true solubility limits as all samples pure M(R) or M0(R) were never formed neither in the contained additional phases, as it was mentioned earlier, form of a secondary phase in any RE-doped SiAlON some of which were rare earth rich. It can, however, be system, nor if synthesized separately according to stoi- noted that the c/a ratios are almost constant with chiometry. Appreciable amounts of J0(R), REAG increasing x-value for each M0(R) phase which indicates

(RE3Al5O12) or other phases were present depending on that the structural changes introduced when Si±N pairs the RE oxide used, the amount of J0(R) increasing with are replaced by Al±O ones are isotropic in all M0(R) phases studied. Since the nominal x value, i.e., the one determined by the powder mixture composition, was not relevant due to the multiphase nature of the synthesized materials, SEM±EDS studies have been used to estimate solubility limits of Al in various M0(R) phases. The maximum aluminum content, as determined in Ref. 36, were x=0.98, 0.9, 0.78, 0.71, and 0.57 for R= Nd, Sm, Gd, Dy, and Y, respectively. Thus, in M0(Nd) and M0(Sm) systems one Si±N pair per formula unit can be replaced by one Al±O pair in accordance with previous ®nd- ings29,41,42 while the true solubility limits in the M0(Gd) and M0(Dy) systems seem to be somewhat less (0.75) and even smaller in the M0(Y) system (0.60). These ®ndings also con®rm that the limit of aluminum solubility decreases with the ionic radii of the RE element.43 It was also shown, that the substitution of Si±N by Al±O in the M(R) structure produces a lattice expansion.36 The fact that the maximum value of lattice expansion was found in the M0(Nd) system (1.66%) and the minimum one in the M0(Y) system (0.70%) further supports the conclusion that the investigated systems exhibit dierent solid solubility ranges. A tentative attempt to rationalize the interrelation between the ionic radii of RE element and the solubility Fig. 1. Schematic representation of the structure of yttrium N-melilite 0 showing separation of the negatively charged silicon oxynitride layers limits of dierent M (R) phases was made in Ref. 40 and by Y3+ cations.40 is based on the simple considerations regarding the V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 2279 structure of melilite-type phases. The projection of the composition, and M0(R) melts.29 This generates large yttrium melilite on (100) may be schematically repre- amounts of liquid which, in combination with silicon, sented as shown in Fig. 1. With non-bridging anions nitrogen and other components, promotes the pre- taking up positions on the apex of the tetrahedra each cipitation of a-SiAlON phase and accelerates the ®nal silicon oxynitride layer of the N-melilite structure pos- densi®cation. This eect has been widely observed in sesses an overall negative charge that has to be balanced several SiAlON systems.41,44 by the RE cations. These cations separate the silicon It is clear that the melting temperature of M0(R) var- oxynitride layers by a ®nite distance. In systems with ies in dierent systems according to the densifying small cation sizes such of as yttrium or dysprosium, the cation, but a low melting temperature for M0(R) is a key negatively charged layers are allowed to come relatively factor in the preparation of dense a and ab-SiAlON closely together and possibly create a repulsive force ceramics. According to Ref. 29 the M0(Sm) melts at which can destabilize the melilite structure. A structurally about 1700C, and on cooling at normal furnace cooling unstable M0(R) phase may result in a low aluminum rate M0(Sm) reprecipitates, suggesting that any liquid accommodation capability. However, the exact mechanism with M0(R) composition does not form a glass. involved in this process is currently unknown. Although it has not yet been con®rmed that dissolution One of the main issues of interest in regard of the of aluminum in M(R) actually lowers the melting tem- properties of the M0(R) phases is their behavior at high perature of these phases, but the closeness of the M0(R) temperatures primarily in relation with densi®cation of composition to the liquid-forming region certainly respective SiAlON compositions. Al-free melilites tend dominates the behavior of M0(R) phases at high tem- to become unstable at temperatures exceeding 1800C. peratures. As it will be discussed further, the high tem- Melting of melilites is thought to be impossible because perature behavior of M0(R) phases plays an important the edge of the liquid region in the four component R± role in the reversible a0 to b0 transformation in SiAlON Si±O±N plane does not reach as far as the melilite ceramics. composition (i.e., 67 eq.% nitrogen). Instead, the melilite Another important factor in designing and processing phases are believed to decompose with signi®cant loss of of various silicon nitride based materials, (a+b)-SiA- nitrogen, resulting in the formation of J phase. For lONs in particular, are the phase relations of M0(R) M0(R), stability at high temperature is determined by with the neighboring phases in dierent M±Si±Al±O±N two factors. With increasing aluminum content, the systems. The most thorough and systematic research of limiting composition of M0(R) gradually moves towards this subject was accomplished in Refs. 45 and 46. This the liquid region, and at the same time the liquid region research covered the majority of RE oxides (R=La, can accommodate a higher nitrogen concentration with Gd, Dy, Er, Y, Nd, Sm, and Yb) that are used as sin- increasing temperature, approaching the terminal com- tering aids in complex SiAlON systems. Firstly, it was position of M0(R). Current preparative evidence suggests shown that the nominal compositions of N-melilite do that the liquid region eventually reaches the M0(R) not form a single-phase material for neither of the RE

Table 2 46 a Reactions and phase formation in R2Si3O3N4 compositions

Rare-earth element, R Reaction temperature (C)/time (h) Appearance Phase analysis by XRD M(R) lattice parameter (A)

Lanthanum 1650/2 (Sintering) White K, 1:2 1550/1.5 (HP) Gray K, 1:2 1600/1.5 (HP) Gray K, 1:2 1650/1.5 (HP) Black 1:2, K Neodymium 1550/1.5 (HP) Blue M. K (vw) 7.721 5.036 1700/1.5 (HP) Blue M, K (w) Samarium 1700/2 (Sintering) White M, K (vw) 1700/1.5 (HP) Brown M, K (vw) 7.695 4.991 Gadolinium 1700/2 (Sintering) White M, K (vw) 7.650 4.961 Dysprosium 1700/2 (Sintering) White M, J (w) 7.618 4.925 Europium 1700/2 (Sintering) Pink M, J, 2:1 1700/1.5 (HP) Pink M, J (mw) 7.585 4.896 Ytterbium 1700/2 (Sintering) Black M, J, 2:1 1700/1.5 (HP) Black M, J (mw) 7.563 4.876

a mw, medium weak, w, weak, vw, very weak, HP, hot pressing. 2280 V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295

Table 3 46 a Reactions and phase formation in R2Si2AlO4N3 compositions

Rare-earth Reaction temperature Appearance Phase analysis M(R) lattice parameter (A) element, R (C)/time (h) by XRD ac

Lanthanum 1650/2 Partly melted 1:2, K 1600/2 White, not dense M0, K, 1:2 7.855 5.120 1550/1.5 (HP) Gray K, 1:2, M0 Neodymium 1650/2 Purple M0 7.766 5.055 Samarium 1650/2 Brown M0 7.730 5.014 Gadolinium 1650/2 Gray M0 7.689 4.993 Dysprosium 1650/2 Yellow M0,J0 (vw) 7.655 4.955 Yttrium 1650/2 White M0,J0 (w) 7.629 4.929 Europium 1650/2 Pink M0,J0 (mw) 7.608 4.920 Ytterbium 1650/2 Black J0

a mw, medium weak, w, weak, vw, very weak, HP, hot pressing.

elements investigated. Moreover, it was unequivocally J phase 2Y2O3 ÁSi2N2O weak base=weak acid shown that M(R) is not formed with lanthanum, which, depending on the synthesis conditions, forms a mixture For example, as the basicity of R2O3 decreases with . of K-phase (La2O3 Si2N2O) and 1:2 phase (La2O32- decreasing ionic radius, it prefers to react with a weak Si3N4) in this or that proportion. The results of the acid to form J phase rather than K phase or 1:2 phase. attempts of M(R) synthesis with dierent RE elements It is also interesting to note that the molar ratio of base are listed in Table 2. to acid in K(R) and J(R) increases from 1:1 to 2:1; this The tendency for forming dierent components in re¯ects the need for a larger amount of the weaker base high-temperature synthesis can be rationalized using the to react with the same acid. concept of acid-base reactions.47 Silicon nitride may be Interesting results and trends were discovered for 0 considered as more acidic than Si2N2O, which enters the M (R) phases when an attempt was made to synthesize 0 compositions of J (2R2O3Si2N2O) and K (R2O3Si2N2O) the nominal composition R2Si2AlO4N3.M(R) solid phases, formed by dierent RE elements as ``impurity'' solution was found in all cases except ytterbium (Table 3). or additional phases (see Table 2), and light RE oxides Overall, the systematic trend of phase distribution as more basic than heavy RE oxides. The various minor discovered was very similar with the one for M(R), with phases formed in the presence of melilite then can be the minor phases favoring K-phase, in the case of viewed as the product of the following reactions: lanthanum, and J-phase, in the case of heavy rare-earth elements. However, some dierences were also 1:2 phase La2O3 Á2Si3N4 strong base=strong acid discovered. First, in the case of lanthanum, melilite solid solution K phase La2O3 ÁSi2N2O strong base=weak acid was formed. This is in contrast with Al-free composition, which did not form La-melilite. The lattice parameters of M0(La) were larger than all of the other M(R) and M0(R) found so far (a=7.855A and c=5.120A). 0 The composition of M (La) was found to be La1.98Si1.82 Al0.96O4.21N3.08, i.e., essentially an x=1 compound in the R2Si3xAlxO3+xN4x series. Similar compositional investigations revealed that the x value decreases monotonically with the atomic number of the RE element. This can be interpreted as a size eect. From the interrelation between the ionic radii of RE elements and both x value and (cubic root of) unit cell volume, (a2c)1/3, from M(R) and M0(R) (Fig. 2) it is evident that the unit cell has expanded in M0(R) relative to M(R). This is accompanied by the substitution of Si±O bond (bond length 1.62A) and Si±N bond (1.74AÊ ) by Al-O bond (1.75AÊ ). Apparently, in the case of lanthanum, which Fig. 2. X(Al) atomic fraction in M0(R) and cell size as a function of has a very large radius, only after maximal substitution, ionic radii of rare-earth elements.46 x=1, is it possible to incorporate lanthanum into the V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 2281

Fig. 3. Phase relations of M0(R) with neighboring phases in R±Si±Al± O±N system.46 Fig. 6. Phase relations of M0(R) with neighboring phases in Yb±Si± Al±O±N system.46

a A6:795 0:892r 0:045x

c A4:121 0:874r 0:033x

where r is the ionic radius (in angstroms) of the RE ion using Ahrens scale. These formulas were obtained by a regression analysis of the experimental data on XRD- determined lattice parameters of M(R) and M'(R) formed by dierent RE elements and can be used for

identifying the composition of R2Si3xAlxO3+xN4x Fig. 4. Phase relations of M0(R) with neighboring phases in R±Si±Al± once the identity of the rare-earth element is known. O±N (R = Gd, Dy, Er, Y) system.46 As to the two other phases, K and J, that were present in the synthesized samples the mechanism of formation and the trends of stability were also formulated in Ref. 46. Formation of K phase was attributed to the reaction

of R2O3 with the surface-oxidized Si3N4, forming Si2N2O. In the case of the Al-containing compositions, the oxidized nitride was supposed to be incorporated into M0(R) solid solution thus preventing the formation of the K phase. For J phase, formation of J0(R) solid solutions was con®rmed for a number of RE elements. Moreover, it was shown that the stability of both J(R) and J0(R) increases with the decrease of ionic radii of RE elements. The x value for J0(R) was shown to increase with the ionic radius of the RE element, which is similar to the trend found for M0(R). Recognizing the systematic trends in M0(R), J0(R), a0, Fig. 5. Phase relations of M0(R) with neighboring phases in R±Si±Al± and b0 compositions, which vary with dierent RE ele- O±N (R = Nd, Sm) system.46 ments, a series of tentative diagrams were suggested in Ref. 46 to delineate the phase relationships between melilite and melilite structure. In case of ytterbium, which has a very neighboring phases. The RE aluminoxynitride, R2AlO3N, small radius in contrast, it is apparently stable only in denoted as 1:1, was included into consideration, based the limit of smallest volume realizable at x=0. on the ®ndings of Ref. 48. Four prototypical diagrams, The lattice parameters of melilite and its solid solu- presented in Figs. 3±6, illustrate these relationships: tion obtained for dierent RE elements at dierent x For lanthanum, there is no a0 solid solution; the x=1 0 values can be correlated to each other, according to Ref. compound of M (R) forms. In addition, La2Si6O3N8 46, by the following formula: (1:2 phase) exists as a neighboring phase (see Fig. 3); 2282 V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295

For neodymium and samarium, there is an extensive reducing the amount of residual grain boundary glass. solid solution of both a0 and M0(R), but only J phase, Therefore, phase relationships in Ln±Si±Al±O±N sys- not J0(R) solid solution, is compatible with M0(R). In tems are of particular interest. 0 addition, R2AlO3N forms and is compatible with M (R) Y±Si±Al±O±N system was investigated most thor- and J (see Fig. 4). The compatibility of RAlO3 (R=Nd, oughly. Initially, research in this system has been 0 Sm) with M (R) were not resolved. According to Ref. restricted to the region bounded by Si3N4, b-SiAlON, 0 46, M (R) is compatible with RAlO3, while according to Al2O3, SiO2 and Y2O3, which does not include the solid Ref. 49, SmAlO3 may be just a product that forms on solution a-SiAlON. Huang et al. reported the a-SiAlON 50 cooling at some speci®c conditions and, in fact, is not formation in the system Si3N4±AlN±Y2O3 and Si3N4± 0 17 compatible with M (R). In Ref. 42 only coexistence of AlN±RE. The systems Si3N4±AlN±Y2O3 and Si3N4± 0 M (R), R2AlO3N and J phase at 1550 C was found. (AlNAl2O3)±(YN3AlN) have also been studied by the For gadolinium, dysprosium, europium, and yttrium, same authors.51 Later, with the emergence of a-SiAlON, there was found an extensive solid solution of a0, but the information on the nitrogen-rich part of the Y±Si±Al±N± solid solution range of M0(R) was shown to be compa- O system became necessary. Thirty-nine four-phase 0 tible with M (R), while the compound R2AlO3N failed equilibria, also named compatibility tetrahedra, had to form in this case (see Fig. 5). Ytterbium was shown to enter a0 solid solution and J0(R) solid solution. However, only M(R), not M0(R), was proven to form in this case (see Fig. 6).

3. Subsolidus phase relationships in R2O3±Si3N4±AlN± Al2O3 systems

As already mentioned, the importance of using RE oxides for the densi®cation of silicon nitride ceramics was recognized in recent years. Not only are they very eective along with alumina for densi®cation, either singly or in combination with yttria, but they can also Fig. 7. Representation of Y±Si±Al±O±N system showing phases be accommodated in the aSi3N4 lattice forming a-SiA- occurring in the region bound by Si3N4,Y2O3,Al2O3, and AlN, and lON, thus providing the opportunity for decreasing the Si±Al±O±N behavior diagram at 1700C.52 transient liquid phase content after sintering and hence

Table 4 52 a Subsolidus compatibility tetrahedra in Si3N4±AlN±Al2O3±Y2O3

0 Al2O3±b60±15R-YAG Al2O3±15R±15R ±YAG 0 0 0 0 Al2O315R ±12H ±YAG Al2O3±12H ±21R ±YAG 0 0 0 Al2O3±21R ±AlN±YAG 15R±15R ±12H±12H ±YAG 12H±12H0±21R±21R0±YAG 21R±21R0±27R±27R0±YAG 27R±27R0±2Hd±2Hd0±YAG 2H±2Hd±AlN±YAG 21R0±27R0±AlN±YAG 27R0±2Hd0±AlN±YAG 21R±27R±YAG±J0(R) 27R±2Hd±YAG±J0(R) 2Hd±AlN±YAG±J0(R) AlN±YAG±J0(R)±YAM

AlN±YAM±J±Y2O3 b60±b25±15R±YAG b25±15R±12H±YAG b25±b10±12H±YAG 0 0 b10±a ±12H±YAG a ±12H±21R±b10 0 0 a ±21R±b10±b8 a ±21R±b8±27R 0 0 d a ±b8±27R±b2 a ±27R±b5±2H 0 d 0 d a ±b5±2H ±b2 a ±2H ±b2±AlN 0 0 a ±b2±AlN±Si3N4 a ±12H±21R±YAG a0±21R±YAG± M a0±21R±27R±M a0±27R±2Hd±M a0±2H±AlN±M M±21R±YAG±J0(R)M±21R±27R±J0(R) M±21R±27R±J0(R) M±27R±2Hd±J0(R) M±2Hd±AlN±J0(R) M±AlN±J0(R)±J

a . . 0 . YAM, 2Y2O3 Al2O3;J,2Y2O3 Si2N2O; J (R) = 2Y2O3 Al2O3± . . d Y2O3 Si2N2O; M=Si3N4 Y2O3; 15R, 12H, 21R, 27R, 2H are Si±rich terminals of AlN polytypoids; 15R0, 12H0, 21R0, 27R0,2Hd0 are Al± Fig. 8. Representation of compatibility of the YAG with polytypoid 52 rich terminals of AlN polytypoids. phases, AlN, and Al2O3, 12 compatibility tetrahedra are formed. V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 2283 . been established in the region Si3N4±AlN±Al2O3±Y2O3 melilite (Y2Si3O3N4), J phase (2Y2O3 Si2N2O), and other (Table 4).52 The subsolidus phase relationships in the oxygen-containing phases. The solubility limits of the a- region Si3N4±AlN±YN±Y2O3 were also determined. SiAlON on the Si3N4±YN3AlN join were determined to Only one compound, 2YN:Si3N4, was con®rmed in the range from m=1.3 to 2.4 in the formula Ym/3Si12m binary system Si3N4±YN contrary to the early results of AlmN16. No quinary compounds were found. Seven Thompson53 who reported the existence of three com- compatibility tetrahedra were established in the region . . . pounds: 6YN Si3N4, 2YN Si3N4,andYNSi3N4. This Si3N4±AlN±YN±Y2O3 (Table 4). contradiction could be explained by the purity of the Sixty-eight compatibility tetrahedra were established YN used as a starting material, and in that case the X- in the system Y±Si±Al±N±O: 39 in the region bounded . . ray diraction lines of 2YN Si3N4 and YN Si3N4 by Si3N4, SiO2, AlN, Al2O3,andY2O3, seven in the reported by Thompson may be regarded as a mixture of region bounded by Si3N4, AlN, YN, and Y2O3, and 22 in the region Si3N4, b60,Al2O3, SiO2, and Y2O3. (here and further bn with n=0±60 is the denomination of b- SiAlON solid solution proposed in Ref. 7; n value indicates the degree of Si±N for Al±O substitution). All this data are graphically presented in Figs. 7±11.

Subsolidus phase relationship in R2O3±Si3N4±AlN- Al2O3 systems until now were thoroughly investigated for most of RE elements, although not in such a detailed and profound way as the Y-doped one. The determination of the whole system, represents a very large amount of work and most initial studies have restricted to the sub-systems R±Si±O±N, R±Al±O±N (R=Ce, Pr, Nd and Sm) and some planes involving a and b-SiAlON.17,41 It is known that elements in the lanthanide series are similar to yttrium in compound formation and phase relationships. Previous research on R±Si±O±N subsystems,54 indicates that the phase relationships are nearly the same as in the Y±Si±O±N system, but in R±Al±O±N systems55 the phase relationships were shown to vary with atomic number of the rare earth Fig. 9. Representation of compatibility of the a-SiAlON with poly- element. The phase relationships in the systems with typoid phases (from 2Hd to 12H), AlN, and b0, eight compatibility high Z-value of the rare earth element are similar to tetrahedra are formed.52 those in Y±Al±O±N.56 In the systems R±Al±O±N (where R=Ce to Sm) to nitrogen containing compounds exist Ð

R2AlO3N and R12Al12O18N2 (magneto-plumbite, MP compound) Ð and no garnet phase occurs; instead the 55 1:1 aluminate phase becomes stable. The R2AlO3N and MP compounds do not occur in those systems containing rare earth oxides with cations smaller than that of Gd.17,55 Therefore, in the ®ve-component systems with the low Z-value rare earth elements, the phase

Fig. 11. Subsolidus phase relationships in the region bound by Si3N4, 0 52 52 Fig. 10. Compatibility terahedron a -b10-12H-YAG. AlN, YN, and Y2O3. 2284 V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295

Table 5

Subsolidus compatibility tetrahedra in the systems Si3N4±AlN±Al2O3± 59 a R2O3 (R = Nd and Sm)

0 Al2O3±b60±15R±MP Al2O3±15R±15R ±MP 0 0 0 0 Al2O315R ±12H ±MP Al2O3±12H ±21R ±MP 0 0 0 Al2O3±21R ±AlN±MP 15R±15R ±12H±12H ±MP 12H±12H0±21R±21R0±MP 21R±21R0±27R±27R0±MP 27R±27R0±2Hd±2Hd0±MP 2Hd±2Hd0±AlN±MP 21R0±27R0±AlN±MP 27R0±2Hd0±AlN±MP d d AlN±2H ±MP±LnAlO3 2H ±27R±MP±LnAlO3 27R±21R±MP±LnAlO3 21R±12H±MP±LnAlO3 12H±15R±MP±LnAlO3 b60±15R±MP±LnAlO3 b60±Al2O3±MP±LnAlO3 b60±b25±15R±LnAlO3 Fig. 12. Representation of Sm(Nd)±Si±Al±O±N system showing pha- b25±15R±12H±LnAlO3 b25±b10±12H±LnAlO3 0 0 ses occurring in the region bound by Si N , Sm(Nd) O ,AlO , and b10±12H±LnAlO3±M (R) b10±12H±21R±M 3 4 2 3 2 3 0 0 0 0 AlN, and Si±Al±O±N behavior diagram at 1700C.59 a ±b0±b10±M a ±b10±21R±M 0 0 a ±b10±b8±21R a ±b8±21R±27R 0 0 d a ±b8±b5±27R a ±b5±27R±2H 0 d 0 d a ±b5±b2±2H a ±b2±2H ±AlN 0 0 d 0 a ±b2±b0 AlN a ±AlN±2H ±M a0±2Hd±27R±M0 a0±27R±21R±M0 0 0 M ±12H±21R±LnAlO3 M ±21R±27R±LnAlO3 0 d 0 d M ±27R±2H ±LnAlO3 M ±2H ±AlN±LnAlO3 0 0 M ±AlN±LnAlO3±J(R) M ±AlN±J(R)±M AlN±LnAlO3±J(R)±Ln2AlO3N LnAlO3±J(R)±Ln2AlO3N±Ln2O3

a 0 MP; LnAl12O18N; M ,Ln2Si3xAlxO3+xN4x; J(R), Ln4Si2O d 7N2; 15R, 12H, 21R, 27R, 2H are Si±rich terminals of AlN polytypoids; 15R0, 12H0, 21R0, 27R0,2Hd0 are Al±rich terminals of AlN polytypoids. relationships become slightly dierent from those in Y± Fig. 13. Sm(Nd)2O3 is compatible with all polytypoid phases (Si-rich Si±Al±O±N. In the oxygen-rich region where two ®ve- terminal), AlN, and Sm(Nd)Al12O18N (MP compounds) forming ®ve 57,58 component phases Ð U phase (R3Si3xAl3+x O12+xN2x) and W phase (R4Si9 Al5O10N) Ð exist, the former is stable for rare earth cations between La and dysprosium the latter stable only in lanthanum, cerium and neodymium systems. None of the phases above occurs in the yttrium system. The stability of N-melilite

(R2Si3xAlxO3+xN4x) in dierent RE systems was discussed in the previous section of this paper. Most closely and most recently the subsolidus phase relationships were studied in neodymium and samar- 59 60 ium, and dysprosium containing R2O3±Si3N4±AlN± Al2O3 systems. The attention paid to the former two elements is caused by the eorts to substitute yttrium by Fig. 14. M0 (melilite solid solution) is compatible with b-SiAlON b -b some cheaper RE elements in Si3N4based ceramics 0 10 without hindering the properties. Dysprosium contain- and a±SIAlON forming a0-b0-M0 compatibility tetrahedron.59 ing system was chosen because dysprosium is one of the central elements in the RE series: in this region, phase samarium-containing system and consequently, at a relationships in R±Si±Al±O±N systems are changing given temperature, the size of the liquid-forming region from those observed for the low-Z elements (La, Nd, is larger. For both samarium and neodymium containing Sm) to those observed for high-Z elements (Er, Yb... systems, 44 compatibility tetrahedra were established and Y). It was therefore interesting to ®nd out whether (Table 5). Some speci®c features of the systems investi- these changes occur simultaneously or whether there is gated as compared to the Y±Si±Al±O±N system were overlap in this region between the dierent phase rela- discovered. Unlike the yttrium containing system where tionships observed for the high-Z and low-Z rare earths. J phase, or woehlerite (Y4Si2O7N2), forms a continuous According to Ref. 59, samarium has the same beha- solid solution with YAM (Y4Al2O9) in neither the Sm- vior as neodymium in regard of phase formation, but it nor the Nd- containing system was the formation of was noted that melting temperatures were lower in the such solid solutions observed, although the J phase itself V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 2285

Table 6

Subsolidus compatibility tetrahedra in the system Si3N4±AlN±Al2O3± 60 a Dy2O3

0 Al2O3±b60±15R±DyAG Al2O3±15R±15R ±DyAG 0 0 0 0 Al2O315R ±12H ±DyAG Al2O3±12H ±21R ±DyAG 0 0 0 Al2O3±21R ±AlN±DyAG 15R±15R ±12H±12H ±DyAG 12H±12H0±21R±21R0±DyAG 21R±21R0±27R±27R0±DyAG 27R±27R0±2Hd±2Hd0±DyAG 2H±2Hd±AlN±DyAG 21R0±27R0±AlN±DyAG 27R0±2Hd0±AlN±DyAG

b60±b25±15R±DyAG b25±15R±12H±DyAG 0 b25±b10±12H±DyAG b10±12H±21R±a 0 0 21R±b10±b8±a b8±21R±27R±a 0 d 0 27R±b8±b5±a b5±27R±2H ±a 0 0 59 d 0 d 0 Fig. 15. Compatibility tetrahedron a -b21-RM . 2H ±b5±b2±a b2±2H ±AlN±a 0 0 0 AlN±b2±b0±a b0±b10±a ±M 0 0 b10±12H±a ±DyAG 12H±21R±a ±DyAG 21R±27R±a0±M0 27R±2Hd±a0±M0 2Hd0±AlN±a0±M0 AlN±2Hd±M0±DyAP 2Hd0±27R±M0±DyAP 27R±21R±M0±DyAP AlN± 2Hd0±DyAP±DyAM 2Hd±27R±DyAP±DyAG 27R±21R±DyAP±DyAG 21R±M0±DyAG±DyAP 0 0 0 0 21R±a ±M ±DyAG a ±b10±DyAG±M 0 0 J (middle)±DyAM±DyAP±AlN Dy2O3±J (whole)±AlN J±J0(middle)±M±M0±AlN J0(middle)±M0±DyAP± AlN

a DyAM = Dy4Al2O9; DyAP = DyAlO3; DyAG = Dy3Al5O12; 15R, 12H, 21R, 27R, 2Hd are Si±rich terminals of AlN polytypoids; 15R0, 12H0, 21R0, 27R0,2Hd0 are Al±rich terminals of AlN polytypoids.

(from b10 to b60) and a-SiAlON (oxygen-rich) forming 12 Al-polytypoids-containing compatibility tetrahedra 0 59 0 0 Fig. 16. Phase relationships of b -NdAlO3 plane at 1750 C. (Fig. 17) and one a ±b -containing compatibility tetra- 0 hedron, a ±b10±DyAG12H (Fig. 18) unlike the Sm(Nd)- containing systems where the YAG does not occur. The was readily formed. The N-melilite solid solution M0(R) melilite solid solution M0 is compatible with both b and was found to be compatible with b-SiAlON phase only a-SiAlON phases, and with DyAG forming two com- 0 0 0 0 in the range b0±b10. Moreover, according to the XRD patibility tetrahedra: a b0-b10-M and a b10-M -DyAG data Nd- and Sm- melilite (Ln2O3±Si3N4) seems to have (Fig. 19). Moreover, in the Dy-containing system the priority over bSi3N4 in accommodating aluminum and J-phase (Dy4Si2O7N2) forms a continuous solid solution 0 oxygen [thereby forming M (R) phase] and therefore the with Dy4Al2O9 unlike the Sm(Nd)-containing ones and tie lines between melilite and b-SiAlON do not run similar with Y and Yb.32,56,61. It was also shown that 0 parallel. The tie line between bSi3N4 and SmAlO3 was the compound DyAlO3 is compatible only with M proven not to exist on account of the results of heat (melilite solid solution) and Dy-rich terminals of AlN treatment of a number of relevant compositions at 1350 polytypoids (from AlN to 21R), and therefor the tie line 0 0 and 1550 C that resulted in formation of M (R) and between LnAlO3 and b that occurs in the Nd(Sm)±Si± apatite [Sm10(SiO4)6N2] only. Graphically these results Al±O±N systems does not exist in the Dy±Si±Al±O±N are presented in Figs. 12±16. system. All the results are summarized in Fig. 20. In the Dy±Si±Al±O±N system the region de®ned by Subsolidus phase relationships for other RE elements, the four end-members Si3N4, AlN, Al2O3 and Dy2O3, although not studied in such detail as the ones presented that includes the phases a-SiAlON, b-SiAlON and AlN above, could be easily derived by analogy if the RE ele- polytypoids, and covers the range of compositions used ment ionic radius is taken into account appropriately. for the design and processing of commercial multi-phase SiAlON ceramics was investigated.60 The binary tie- lines and 42 compatibility tetrahedra were established 4. Thermal stability of RE- -SiAlONs and the based on experimental results and on the more extensive reversibility of 0 to 0 transformation knowledge of phase relationships in the Y±Si±Al±O±N and Sm(Nd)±Si±Al±O±N systems (Table 6). The RE densi®cation additives used for sintering of

The DyAG (Dy3Al5O12) phase is stable and compa- SiAlON ceramics are generally insoluble in the b-SiA- tible with all the AlN polytypoid phases, b-SiAlON lON structure and even though most of them are a- 2286 V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295

0 Fig. 19. M (melilite solid solution) is compatible with b-SiAlON b0- 0 0 0 60 b10 and a±SiAlON forming a -b -M compatibility tetrahedron.

Fig. 17. Graphic representation of compatibility of the DyAG with polytypoid phases, AlN, and Al2O3, 12 compatibility tetrahedra are formed.60 Fig. 20. Representation of Dy±Si±Al±O±N system showing phases

occurring in the region bound by Si3N4,Dy2O3,Al2O3, and AlN, and Si±Al±O±N behavior diagram at 1700C(a-SiAlON plane = the plane 60 with a-SiAlON composition DyxSi12(m+n)Alm+nOnN16-n).

been paid to the eects of the post-sintering heat treatment on the stability of SiAlON phases, while the characterization of the ®nished product has always focused on the account of the residual glass and the type, amount and microstructure of the resulting crystalline oxynitride phase. However, in the course of the study of densi®cation and post sintering heat treatment of (a+b)-SiAlON compositions containing selected RE Fig. 18. DyAG is compatible with b-SiAlON b10±b60 and a±SiAlON 0 60 additives (Yb O Dy O ,SmO either singly or in (oxygen-rich) forming a -b10-12H-DyAG compatibility tetrahedron. 2 3, 2 3 2 3 25 combination with Y2O3). Mandal et al. reported that the transformation from a0 to b0 readily occurred in SiAlON formers, some amount of additives remains as some RE-(a0+b0) composite materials when heat trea- a residual M±Si±Al±O±N intergranular glassy phase, ted between 1000 and 1600C and that extent of the which degrades the mechanical and chemical properties transformation was more pronounced with increasing of the material above the glass-softening temperature, temperatures. which normally is about 900±1100C. Although the structure of both a0 and b0 phases are

Post-sintering heat treatment is one of the accepted basically built up of corner sharing (Si, Al)(O, N)4 tetra- methods to eliminate or minimize the glassy grain hedra, there is a distinct dierence in the atomic boundary phase in nitrogen ceramics by converting it arrangement between the two phases.63,64 The transfor- into refractory crystalline phase(s), thus improving the mation between a0 and b0 phases has a reconstructive materials performance at elevated temperatures. This is nature, which involves the breaking of chemical bounds usually achieved via a conventional glass-ceramic pro- and substantial atomic diusion, and hence requires cess during which the glass devitri®es at a temperature signi®cant amounts of thermal energy. As a result of the 58,62 above its Tg point but without being melted. There strong covalent nature of the bounding associated with has been much work in understanding of the grain both the a0 and b0 structures, the atomic diusivity of boundary crystallization, but very little attention has the species making up the lattices is inherently low. It is, V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 2287

Fig. 21. Heat treatment of 1775 C pressureless sintered (a) b-SiAlON and (b) a+b-SiAlON materials. B,B-phase {(Y, R)2SiAlO5N}; W, wollasto- nite {(Y, R)SiO2N}; U, U-phase {R3Si3Al3O12N2}; Y, J(R)-phase {(Y

Table 7 Mechanisms in¯uencing a0 to b0 transformation69

Driving force

Sample O:N ratio Sintering a0 to b0 R cation Amount and viscosity b-SiAlON nucleation (%) additive transformation size of liquid phase sites

R1 13.7 Nd2O3 or Sm2O3 + +++ ++ + Dy2O3 or Yb2O3 + No +++ +

R2 10.6 Nd2O3 or Sm2O3 + +++ ++ + Dy2O3 or Yb2O3 + No +++ +

R3 9.12 Nd2O3 or Sm2O3 + +++ ++ + Dy2O3 or Yb2O3 + No +++ +

R4 11.88 Nd2O3 or Sm2O3 + +++ ++ + Dy2O3 or Yb2O3 + No ++ +++

R5 10.34 Nd2O3 or Sm2O3 + +++ ++ No Dy2O3 or Yb2O3 No No No +++ 2290 V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 region, the regularities of the ease of phase transforma- polytypoid. These observations are fully consistent with tion are of a more complex nature. A combination of the results of Ref. 69. Overall, the above described three factors in¯uences the ease of transformation: the experimental results on the stability of a0 and a0 to b0 cation size of the sintering additive, the presence of b- transformations strongly indicate that for duplex SiAlON grains and also the amount and viscosity of (a+b)-SiAlON composites designed for high-temperature liquid phase present during heat treatment. The role of applications RE elements with smaller ionic radii present the liquid phase in the phase transformation for all an optimal choice. compositions was elucidated in the same way: an extra It must be also noted that in spite of extensive inves- amount (10 and 20%) of glass powders of overall com- tigation of the transformations in question it is not position R: Si: Al=1:1:1 was added to the samples. completely clear what is the driving force behind the Such additions led to unequivocal intensi®cation of the transformations. Therefore, it is quite possible that the a0 to b0 transformation in all samples with the exception resulting phase assemblies are formed due to kinetic of the ones from the ®fth group (see above), that were rather than thermodynamic reasons. doped with small ionic size of sintering additives (Dy2O3 and Yb2O3). For these materials the presence of b-SiAlON grains was necessary for triggering the transformation. 5. Reaction densi®cation of and ( + ) SiAlONs Moreover, the increase of the number of such grains (10% b-SiAlON with z=0.8 was added to one of the Densi®cation of SiAlONs occurs by reaction liquid- mixtures) enhanced the transformation substantially, phase sintering or ``transient liquid sintering''. Intensive while such grains might possibly act as a nucleation site investigations have been done on the reaction liquid- 0 0 for a to b transformation. Only then the amount and phase sintering of SiAlON in the metal oxide±Si3N4± 7182 viscosity of liquid became important factors for the overall AlN±SiO2±Al2O3 systems. Some of the earliest reaction a-SiAlON+b-SiAlON+(21RAlN)+liquid b- work was carried out in the MgO±Si3N4±AlN±SiO2± 7174 SiAlON+(21RAlN)+grain-boundary phase(s). b-SiAlON Al2O3 system. It was found that a transient phase, nucleation sites provided more favorable kinetic condi- which formed during hot pressing of a-Si3N4 with MgO, tions for transformation in case of large ionic size sintering reacted with a-Si3N4 above 1400 C to form b-Si3N4 and additives, but the process itself proceeded without any a vitreous phase.71 This reaction was found to expedite additional ones. The overall results are summarized in the sintering kinetics and the reaction pathway of a

Table 7, where increasing number of ``+'' marks indi- mixture of MgO±Si3N4±AlN±SiO2±Al2O3 depended on 0 0 7173 cate the more dominant mechanism in¯uencing a to b the amount of MgO added. The addition of Y2O3, transformation. instead MgO to a Si3N4±AlN±SiO2 mixture resulted in a Somewhat similar research was carried out for duplex material with diphasic microstructure consisting of b0-

(a+b)-SiAlON composites stabilized with neodymium, SiAlON and YAG (Y3Al5O12) with improved mechan- ytterbium, and yttrium. The former two RE elements ical properties.74 Later the attempts were made to were chosen because they represent the size end members obtain dense b-SiAlON ceramics in the Si3N4±AlN± 0 7577 of rare earth cations known to stabilize a , while yttrium SiO2±Al2O3 system. In these studies, it was shown served as a reference. The long-term high-temperature that the sintering kinetics depended on the composition stability of the dierent duplex (a+b)-SiAlON ceramics of starting powder mixtures, even though the ®nal was of special interest. Phase and microstructural composition of the sintered material was the same. In development and the volume fraction and chemistry of these and other reaction sintering and hot-pressing stu- residual intergranular phases were characterized and dies, the densi®cation mechanisms were identi®ed to be related to mechanical properties. It was shown that on solution-reprecipitation,7174 fast particle rearrange- sintering the duplex (a+b)-SiAlON starting powder ment,80 Cobble creep,80 or grain boundary sliding.79,81 compositions result in equilibrium phase compositions Recently, it was shown by Hwang and Chen that consisting of a0, b0 and a residual glassy phase. The reaction hot pressing of a0- and (a+ b)-Y-SiAlON took analytical XRD investigations showed that a0 and b0 place in three stages.82 Further, it was found out that substitution levels and the equilibrium composition and the wetting properties of the Y2O3±Al2O3±SiO2 eutectic volume fraction of the glass are dependent upon the melt controlled the densi®cation behavior of powder starting composition as well as the a0 stabilizing cation. compacts. The ®rst stage was identi®ed with the forma- However, the b0/(a0+b0) weight ratio was shown to tion of the ternary oxide eutectic and YAG. The depend mostly on the RE element used. The long-term shrinkage in this stage was due to the redistribution of heat treatment experiments at 1450C clearly indicate liquid in the powder compact and slight improvement in that yttrium and ytterbium give a stable a-SiAlON the particle packing densities. The second stage was phase, while neodymium gives a less stable one that identi®ed with the wetting of AlN particle and forma- during prolonged heat treatment at 1450 C decomposes tion of b60. The preferential wetting of AlN by the into crystalline phases: melilite, b-SiAlON, and 21R ternary oxide melt caused localization of the liquid, V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 2291 leading to a delaying eect on the second shrinkage heating was studied by means of heating up the mix- step. The third stage occurred with the dissolution of tures to some intermediate temperatures, cooling down, aSi3N4 and formation of the ®nal phase. Massive particle and subsequent XRD characterization. The wetting rearrangement was found to be the dominant densi®ca- experiments were to study the trends in the wetting tion mechanism. It also caused aluminum enrichment in behavior of the ternary oxide melts on the nitrides. For the liquid, leading to formation of b60 as a transient this purpose ternary eutectic compositions were pre- phase. Thus the physical/chemical characteristics of the pared and pressed into pellets, which were placed on liquid, in particular its wetting behavior, and the kinetic substrates prepared from Si3N4 and AlN by hot press- pathway of the intermediate reactions are clearly ing without additives. These ``sandwiches'' were then important factors in reaction densi®cation of multi- heated to dierent temperatures, in a nitrogen atmo- component systems. sphere, where the pellets fused. On cooling, the contact A survey of the literature indicates that until recent angles were measured from photomicrographs. Since time there was very limited knowledge of the wetting the wetting characteristics are sensitive to the composi- characteristics of various ternary oxide eutectic melts on tion (including the oxygen and nitrogen content in the nitrides. Most studies on reaction sintering have been melt) and the temperature, the true equilibrium was not restricted to the documenting of the evolution of phases achieved during the experiments.91 Therefore, these and densities. Only limited information may be gleaned experiments can not be considered as yielding the actual from the reports of transient/intermediate phases of values of the wetting angle. these studies. The reaction sequence were studied in Ca-, It was shown that the wetting behavior of the ternary

Y- and Nd-a-SiAlON, and, according to Refs. 83 and oxide melt towards Si3N4 and AlN was strikingly dif- 84, gehlenite (Ca2Al2SiO7) was found to be transient ferent, depending on the metal oxide used. Such a pre- phase in Ca-a-SiAlON, whereas N-melilite (R2Si2O3N4) ference in reaction sequence was explained in terms of formed during the densi®cation of Y- and Nd-a-SiA- acid-base chemistry.91 According to the Pearson's prin- lON. N-melilite was also reported to be a transient ciple, hard acids preferentially react with soft bases.47 phase in Nd- and Sm-doped a-SiAlON.85,86 The same Empirically, polarizability, electronegativity, and charge phase was identi®ed as a transient one during the pres- to size ratio47 can be used as a measure of hardness of sureless sintering of lanthanide (Y and Nd to Lu, except 87 Pm)- modi®ed a-SiAlON. K phase (RSiO2N) and b- Si3N4 phase along with melilite were determined in the case of Nd-a-SiAlON in Ref. 38. Furthermore, in the same work J phase (R4Si2O2N) and b-Si3N4 were reported to occur along with melilite in the case of Sm- and Gd-a-SiAlON, and J phase and b-Si3N4 in case of Dy-, Y-, Er- and Yb-a-SiAlON. The reaction sequence in hot-isostatic-pressed and pressureless-sintered, RE- doped (a+b) SiAlONs have been studied.9,8790 For- mation of 12H and La±Si±oxynitride as the transient phase in the case of Y- or Y/La-doped samples were reported. Generally it has not been known what controls the formation and selection of various phases. The ®rst indication that a general understanding of reaction sequence during sintering of SiAlON materials may, after all, be based on simple physical chemical considerations governing surface activities of nitrides with reactive liquids was given in Ref. 82. There it was shown that in

Y-SiAlON system, Y2O3±Al2O3±SiO2 melt preferentially wets AlN, leading to the formation of b60- SiAlON as a transient phase, which delays densi®cation in the early stages due to liquid trapping. A more systematic research of the wetting behavior of ternary eutectic melts Me2O3±Al2O3±SiO2 (Me=Li, Mg, Ca, Y, Nd, Sm, Gd, Dy, Er, and Yb) and their in¯uence on the phase development during reaction hot pressing of aSi3N4, AlN, Al2O3, and Me2O3 powder 91 mixtures was accomplished. Later these results were Fig. 22. Schematic shrinkage curves when the eutectic melt wets (a) 92 92 tied in with densi®cation. The evolution of phases on Si3N4, and (b) AlN ®rst. 2292 V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295

93 an acid or a base. Recently, a model of pH0 (the point The schematic representation of expected shrinkage of zero charge that corresponds to the pH at which the curves when the eutectics melt preferentially wets Si3N4 immersed solid has a zero net surface charge) values and AlN are shown in Fig. 22 (a) and (b), respectively. calculation for ionic metal oxides based on the ioniza- In these curves, the beginning of various shrinkage steps tional potential of the metal ion was proposed.94 If the are identi®ed with some of the above characteristic oxides used in Ref. 91 are put in order of decreasing temperatures. Wetting of AlN is shown to lead to small basicity according to this model, and if the acidity of the or no immediate shrinkage in the powder compact nitrides is assumed to be determined by the acidity of because of the low AlN content of the typical SiAlON 82 their native oxide layers (in the case Si3N4 will be more compact. As it is shown in Fig. 22 (b), though, a sub- acidic than AlN), than both the experimental wetting sequent shrinkage step occurs when the majority Si3N4 data and the sequence of the intermediate reaction pro- is wetted. Variation of these shrinkage curves may also ducts in relation to their Si/Al ratio will be in total be possible. For example, it is known that the initial accordance with the above mentioned acid-base princi- precipitation of the intermediate phase usually starts ple. Therefore, the harder acid, Si3N4, reacts with hard with partial wetting of the ®rst nitride, i.e., T2 could be bases of oxide melts of lithium, calcium, magnesium, signi®cantly lower than the temperature when complete and lighter rare earth metals up to gadolinium, whereas wetting is achieved. Also, the dissolution of inter- the soft acid AlN reacts with the soft bases of the heavier mediate phase may or may not be higher than T5. RE metals down to dysprosium. In the order of Indeed, T4 may even be lower than T3, as in case of decreasing Si/Al ratio and additive basicity, typical lithium and calcium, where wetting and dissolution of 0 intermediate phases are O (SiN2O2), M-cordierite the second nitride (AlN) occurs after the intermediate 0 0 (Mg2Al4Si5O18), M (R) (R2Si2AlO4N3), and b60-SiAlON phase (O ) is dissolved and partial formation of a-SiA- 92 (Si0.8Al4.2O4.2N3.8). lON is already achieved. Experimental data on the The temperatures for the various reactions, which can characteristic temperatures of shrinkage determined for be used to understand the shrinkage behavior, were also a number of systems are presented in Table 8. Experi- identi®ed based on the observed wetting and reaction mental shrinkage curves obtained by means of hot behavior.87 The reactions identi®ed during the reaction pressing with a constant treating rate91 were in good densi®cation of a-SiAlON are: agreement with the ones, predicted by the sequence of wetting reactions.

. the eutectic formation (at temperature T1), Reaction densi®cation of Me±SiAlON occurs in three . wetting of a nitride powder and intermediate phase stages. The ®rst stage is associated with the formation of the

precipitation (at temperature T2), SiO2±Al2O3±Me2O3 ternary eutectic (at temperature T1). . secondary wetting of the other nitride powder (at The second stage is associated with the preferential

temperature T3), wetting of majority nitride powder, Si3N4. This occurs . dissolution of the intermediate phase (at tempera- at T2 when Si3N4 is wetted ®rst. If AlN is wetted ®rst, ture T4), and then no shrinkage step is observed at T2 and the . precipitation of the ®nal phase, a-SiAlON (at densi®cation is delayed until T3, when Si3N4 is wetted temperature T5). ®nally. The third stage involves the dissolution of the inter-

mediate phase if T4 is low; otherwise a distinct third stage is not seen while a gradual densi®cation continues following

Table 8 the second stage. In the extreme case of very high T4, Summary of features of the shrinkage curves for samples of Me±Si± this may result in poor density as in case of magnesium, Al±O±N (Me = Li, Ca, Mg, Nd, Sm, Gd, Dy, Er, Yb) hot pressed at a where cordierite precipitates as an intermediate phase. 92 heating rate of 15 C/min The formation temperature of the ®nal a-SiAlON

Volume shrinkage phase (T5) is not crucial in most cases, unless a sudden secondary precipitation of a-SiAlON at this later stage System T1 T2 T3 T4 T5 First Second 0 (C) (C) (C) (C) (C) stage stage (T 5) drains the liquid and retards densi®cation. Such 0 material must be densi®ed at temperatures below T 5. Li 1010 1225 1600 1500 1420 8.0 15.0 To achieve full densi®cation at relatively low tempera- Ca 1360 1450 1660 1610 1550 5.0 28.0 tures, a low wetting temperature for the second nitride Mg 1350 1500 1630 >1850 1575 5.0 9.0 Nd 1400 1450 1580 1800 1420 4.0 15.0 (T3) and low dissolution temperature of the inter- Sm 1380 1460 1600 1800 1420 7.0 17.0 mediate phase (T4) are required. The dominant process Gd 1370 1500 1620 1700 1575 7.0 10.0 for densi®cation in SiAlON ceramics was shown to be Dy 1360 1500 1560 ± 1575/1675 5.0 20 massive particle rearrangement although some dissolu- Er 1400 1500 1580 1640 1580 4.0 13.0 tion facilities full densi®cation by providing the neces- Yb 1380 1500 1560 1640 1560 9.0 10.0 sary liquid.92 V.A. Izhevskiy et al. / Journal of the European Ceramic Society 20 (2000) 2275±2295 2293

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