《复合材料 Composites》课程教学资源（学习资料）第五章 陶瓷基复合材料_Progress in Silicon-Based Non-Oxide Structural Ceramics

Int J. of Refractory Metals Hard Materials 15(1997)13-4 Printed in Great Britain. All rights reser ELSEVIER PII:S0263·4368(96)00046-7 Progress in Silicon-Based Non-Oxide Structural Ceramics Wolfgang Dressler*& ralf riedel Fachgebiet Disperse Feststoffe, Fachbereich Materialwissenschaft, Technische Hochschule Darmstadt, Petersenstr 23 A, D-64287 Darmstadt, Germany (Received 1 August 1996; accepted 16 August 1996) Abstract: The progress in monolithic Si,N, and SiC as well as in Si3N4/SiC composites for structural applications is reviewed. The conventional processing including the powder synthesis, densification and microstructural design is dis- cussed. The mechanical properties of the resulting silicon based non-oxide cera- mics and their industrial applications as structural components are summarized As an alternative route to fabricate Si,N /SiC composites the hybrid processing utilizing the thermal conversion or organosilicon precursors to amorphous and polycrystalline multicomponent materials is described. The hybrid processed cera- mics exhibit ultra-high temperature stability with respect to crystallization, oxida- tion and decomposition. o 1997 Elsevier Science Limited 1 INTRODUCTION oxidizing environments only up to temperatures in the range of 1200-13000C. At higher tem- high hardness and strength, excellent creep, oxi- dized. even in the bulk. This behavior is dation and corrosion resistance as well as their related to the presence of sintering promoting low density, silicon nitride(Si3N4)and silicon compounds like Mgo, Al2O3 or Y2O3 which nitride/carbide(Si, N,/SiC) based ceramics and react with SiOz formed during the oxidation ceramic composites are promising candidate reaction to give low viscous silicates, (iii)th materials for high temperature applications in use of conventionally processed secondary motor and turbine devices and are frequently phase free SiC in the high temperature field is ed as cutting tools, a typical applio limited to low strength applications, ,> du hard materials. However, these materials still the reduced flaw tolerance caused by the low mechanical performance at temperatures above uid phase sintered Sic reveals a higher sensitiv- 1200oC is required. These limitations are due to ity towards oxidation owing to the reaction of the intrinsic properties of ceramics or result the used sintering aids with the oxide product from the used processing process and basically formed on the SiC surface,(v)the conven- an be attributed to the following points: (i) tional fabrication of dense Si3N,/SiC-composites ceramics are difficult to apply with high relia- is difficult due to the distinct sintering behavior bility owing to their intrinsically brittle behavior. of the Si3 na and Sic powder particles used as In contrast to this, metals are limited by corro- the starting materials. .8 However, significant sion problems and by reduced performance at improvements of the room and high-tempera temperatures approaching their melting point, ture properties of Si3 Na- and Sic-based cera- (ii) commercial Si3N, parts can be applied in mics like strength, fracture toughness, creep and oxidation resistance have been achieved by *Present address: Robert Bosch GmbH, FVIFLW, Post- tailoring of microstructure, -I7 by generation of fach 106050, D-70049, Stuttgar Si,N,/SiC micro/micro-., 18-20 or micro/nano- 13

ELSEVIER Int. J. of Refractory Metals & Hard Materials 15 (1997) 13-47 © 1997 Elsevier Science Limited Printed in Great Britain. All rights reserved 0263-4368/97/$17.00 PII: S0263-4368(96)00046-7 Progress in Silicon-Based Non-Oxide Structural Ceramics Wolfgang Dressier* & Ralf Riedel Fachgebiet Disperse Feststoffe, Fachbereich Materialwissenschaft, Technische Hochschule Darmstadt, Petersenstr. 23 A, D-64287 Darmstadt, Germany (Received 1 August 1996; accepted 16 August 1996) Abstract: The progress in monolithic Si3N4 and SiC as well as in Si3Nn/SiC￾composites for structural applications is reviewed. The conventional processing including the powder synthesis, densification and microstructural design is dis￾cussed. The mechanical properties of the resulting silicon based non-oxide cera￾mics and their industrial applications as structural components are summarized. As an alternative route to fabricate Si3N4/SiC composites the hybrid processing utilizing the thermal conversion or organosilicon precursors to amorphous and polycrystalline multicomponent materials is described. The hybrid processed cera￾mics exhibit ultra-high temperature stability with respect to crystallization, oxida￾tion and decomposition. © 1997 Elsevier Science Limited 1 INTRODUCTION Owing to their advantageous properties, like high hardness and strength, excellent creep, oxi￾dation and corrosion resistance as well as their low density, silicon nitride (Si3N4) and silicon nitride/carbide (Si3N4/SiC) based ceramics and ceramic composites are promising candidate materials for high temperature applications in motor and turbine devices and are frequently used as cutting tools, a typical application for hard materials. However, these materials still have application limits where reliability or mechanical performance at temperatures above 1200°C is required. These limitations are due to the intrinsic properties of ceramics or result from the used processing process and basically can be attributed to the following points: (i) ceramics are difficult to apply with high relia￾bility owing to their intrinsically brittle behavior. In contrast to this, metals are limited by corro￾sion problems and by reduced performance at temperatures approaching their melting point, (ii) commercial Si3N4 parts can be applied in *Present address: Robert Bosch GmbH, FVIFLW, Post￾fach 106050, D-70049, Stuttgart. 13 oxidizing environments only up to temperatures in the range of 1200-1300°C. At higher tem￾peratures, the material creeps' and is oxi￾dized 2"3 even in the bulk. This behavior is related to the presence of sintering promoting compounds like MgO, A1203 or Y203 which react with SiO2 formed during the oxidation reaction to give low viscous silicates, (iii) the use of conventionally processed secondary phase free SiC in the high temperature field is limited to low strength applications, 4"5 due to the reduced flaw tolerance caused by the low fracture toughness of additive free SiC, (iv) liq￾uid phase sintered SiC reveals a higher sensitiv￾ity towards oxidation owing to the reaction of the used sintering aids with the oxide product formed on the SiC surface, 6 (v) the conven￾tional fabrication of dense Si3Nn/SiC-composites is difficult due to the distinct sintering behavior of the Si3N 4 and SiC powder particles used as the starting materials. 7"8 However, significant improvements of the room and high-tempera￾ture properties of Si3N4- and SiC-based cera￾mics like strength, fracture toughness, creep and oxidation resistance have been achieved by tailoring of microstructure, 9-'7 by generation of Si3N4/SiC micro/micro -7'8'18-26 or micro/nano-

D R. Riede y H Fig. 1. Crystal structure of trigonal a-Si, Na(P31c). View nearly along z-axis. Silicon atoms in red and nitrogen atoms in Fig. 2. Crystal structure of hexagonal B-Si, N,(P63). View nearly along z-axis. Silicon atoms in red and nitrogen atoms In green

14 W. Dressier, R. Riedel H ! ..J Z Fig. 1. Fig. 2. Crystal structure of trigonal ~-Si,N4 (P31c). View nearly along z-axis. Silicon atoms in red and nitrogen atoms in green. Crystal structure of hexagonal/~-Si3N4 (P6~). View nearly along z-axis. Silicon atoms in red and nitrogen atoms in green

Silicon-based non-oxide structural ce H 3 Fig. 3. Crystal structure of hexagonal a-SiC polytype 6H(P6, mc). View nearly along y-axis. Silicon atoms in red and Fig. 4. Crystal structure of cubic B-SiC polytype 6H 3C. View nearly along [101]-direction. Silicon atoms in red and carbon atoms in blue

Silicon-based non-Qxide,~ructural ceramics 15 + , , . YT" Z 3 4 Fig. 3. Fig. 4. Crystal structure of hexagonal a-SiC polytype 6H (P63 mc). View nearly along y-axis. Silicon atoms in red and carbon atoms in blue. Crystal structure of cubic fl-SiC polytype 6H 3C. View nearly along [101]-direction. Silicon atoms in red and carbon atoms in blue

16 W. Dressler, R. riedel composites, by using high melting sintering AB except that it is rotated by 180% on the c- additives, 4 by reduction of additive content* axis. s Consequently, the channels present in and by devitrification of grain boundary pha- the B-modification are closed off into two inter ses.as reviewed in this article. Additionally, stices and the c-dimension of a-si3N an alternative route for the fabrication of silicon (c=0. 5617 nm) is approximately twice that of nitride, silicon carbide ceramics and Si3N,/SIC- B-Si3N4(c=0. 29107 nm). The lattice parameters composites, the hybrid processing, 4-0, 43,44 in the a directions are similar: a(a-Si3N4) shows exceptional possibilities to meet the =0.7818 nm, a(B-Si3 NA)=0-7595 nm aforementioned requirements. Here, organosili The technical synthesis routes lead mainly to con polymers(hybrids)are converted to silicon- a-Si,N, which converts to B-Si3 Na during liquid based inorganic materials this processing phase sintering. Grun" calculated the free technique not only the phase composition and enthalpy of a-to B-transformation at 10 Pa and microstructure but also the thermo-mechanical 298 K to be -30 kJ mol. Hampshire and properties of the final ceramics can be tailored Jack*/ reported that the activation energy for by the choice of organosilicon system, the the a-to B-transformation during liquid phase design of the intermediate processing steps(to sintering is similar to the dissociation energy of convert the system to an inorganic amorphous the Si-n bond, 435+38 kJ mol Thus the intermediate) and the use of annealing treat- mechanism of the reconstructive transformation ments(to transform the inorganic intermediate seems to be the breaking of Si-N bonds, the into the desired final crystalline ceramics) solution of the less stable a-phase in the formed This article in general reviews the progress in liquid phase and finally the reprecipitation of Si3N4 and Sic based materials achieved in the less soluble, more stable b-modification past decade. In particular, the fabrication and The technical production of Si,n4 powder properties of Si,N,/SiC composites derived predominantly performed on four different from (i) conventional processing and (ii) from routes: (i) the direct nitridation of silicon pow advanced techniques are comparatively dis- der,(ii) the carbothermal reduction and subse quent nitriding of Sio2,(ii the diimide process and, (iv) the gas phase reaction of silanes with ammonia 2 SILICON NITRIDE (Si,N,)-INTRINSIC The direct nitridation of elemental silicon to STRUCTURAL PROPERTIES AND achieve stoichiometric Si, Na was developed SYNTHESIS from Weiss and Engelhardt in 1910 and is still the most common industrial processing Si3 N, is a highly covalent compound(70% cova-route lence)having a density of 3 19g cm-3and 3Si+2N2-100140c occurring in two different crystal structures the a-and B-modification. The a-structure possesses AH=-750 kJ mol the space group P3lc the B-modification is hex The resulting SiN, powder is subsequently agonal(P63/m)as depicted in Figs 1 and 2. The milled and consists mainly of the a-modifica- polymorphs consist of slightly distorted tetra tion. In this connection, the choice of starting hedral SiN,(sp hybridization of Si)and planar silicon powder quality (grain size and purity) NSi,(sp- hybridization of f N units The SiN determines on the one hand the price of the tetrahedrons are joined by sharing nitrogen cor- resulting Si,N, and on the other hand the purity ners so that each nitrogen is common to three of the product. Especially the use of semicon tetrahedra. The unit cells of a-Si3n4 and ductor silicon leads to extremely pure Si, NA- B-Si3NA are represented by Si12N16 and Sions, powders. 4 The carbothermal reduction method respectively, as shown in Figs 1 and 2. The starts from a mixture of fine Sio, and carbon B-structure is composed of puckered rings of powder. This mixture is converted into Si N4 at alternating Si andn atoms having a stacking 1500 C in flowing N2 quence of ABAB+ and forming channels 1500°C ( diameter about 0. 15 nm) along the c-direction 3SiO2+6C+2N2 Si3N4+6Co(2) The a-modification contains the same AB layer In order to convert the initial SiOz completel and an additional layer CD which is similar to into Si3n4 excess carbon is necessary. .>3 By

16 IV.. Dressier, R. Riedel composites, 27-39 by using high melting sintering additives, 4°'4~ by reduction of additive content 43 and by devitrification of grain boundary pha￾ses 3"42 as reviewed in this article. Additionally, an alternative route for the fabrication of silicon nitride, silicon carbide ceramics and Si3N4/SiC￾composites, the hybrid processing, 34-36"43"44 shows exceptional possibilities to meet the aforementioned requirements. Here, organosili￾con polymers (hybrids) are converted to silicon￾based inorganic materials. Using this processing technique not only the phase composition and microstructure but also the thermo-mechanical properties of the final ceramics can be tailored by the choice of organosilicon system, the design of the intermediate processing steps (to convert the system to an inorganic amorphous intermediate) and the use of annealing treat￾ments (to transform the inorganic intermediate into the desired final crystalline ceramics). This article in general reviews the progress in Si3N 4 and SiC based materials achieved in the past decade. In particular, the fabrication and properties of Si3N4/SiC composites derived from (i) conventional processing and (ii) from advanced techniques are comparatively dis￾cussed. 2 SILICON NITRIDE (Si3N4) -- INTRINSIC STRUCTURAL PROPERTIES AND SYNTHESIS Si3N4 is a highly covalent compound (70% cova￾lence) having a density of 3.19g cm -3 and occurring in two different crystal structures the a- and fl-modification. The a-structure possesses the space group P31c the fl-modification is hex￾agonal (P63/m) as depicted in Figs 1 and 2. The polymorphs consist of slightly distorted tetra￾hedral SiN4 (sp 3 hybridization of Si) and planar NSi3 (sp 2 hybridization of N) units. The SiN4 tetrahedrons are joined by sharing nitrogen cor￾ners so that each nitrogen is common to three tetrahedra. The unit cells of a-Si3N4 and fl-Si3N 4 are represented by Si12Ni6 and Si6Ns, respectively, as shown in Figs 1 and 2. The fl-structure is composed of puckered rings of alternating Si and N atoms having a stacking sequence of ABAB 45 and forming channels (diameter about 0.15 nm) along the c-direction. The a-modification contains the same AB layer and an additional layer CD which is similar to AB except that it is rotated by 180 ° on the c￾axis. 45 Consequently, the channels present in the fl-modification are closed off into two inter￾stices and the c-dimension of a-Si3N 4 (c=0.5617nm) is approximately twice that of fl-Si3N4 (c=0.29107 nm). The lattice parameters in the a directions are similar: a(a-Si3N4) =0.7818 nm, a(fl-Si3N4)=0.7595 nm. 46 The technical synthesis routes lead mainly to a-Si3N4 which converts to fl-Si3N4 during liquid phase sintering. Grfin 46 calculated the free enthalpy of a- to fl-transformation at 105 Pa and 298 K to be -30kJ mol-'. Hampshire and Jack 47 reported that the activation energy for the a- to fi-transformation during liquid phase sintering is similar to the dissociation energy of the Si-N bond, 435+38kJ mol -~. Thus the mechanism of the reconstructive transformation seems to be the breaking of Si-N bonds, the solution of the less stable a-phase in the formed liquid phase and finally the reprecipitation of less soluble, more stable fi-modification. The technical production of Si3N 4 powder is predominantly performed on four different routes: (i) the direct nitridation of silicon pow￾der, (ii) the carbothermal reduction and subse￾quent nitriding of SiO2, (iii) the diimide process and, (iv) the gas phase reaction of silanes with ammonia. The direct nitridation of elemental silicon to achieve stoichiometric Si3N4 was developed from Weiss and Engelhardt in 191048 and is still the most common industrial processing route 5,.52 3Si + 2N2 1 lOO-14oo°c ~" Si3N4 AH= -750 kJ mol-1. (1) The resulting Si3N4 powder is subsequently milled and consists mainly of the a-modifica￾tion. In this connection, the choice of starting silicon powder quality (grain size and purity) determines on the one hand the price of the resulting Si3N 4 and on the other hand the purity of the product. Especially the use of semicon￾ductor silicon leads to extremely pure Si3N n￾powders? 4 The carbothermal reduction method starts from a mixture of fine SiO2 and carbon powder. This mixture is converted into Si3N 4 at 1500°C in flowing N~. 3Si02+6C+2Nz 15oooc > Si3N4+6C0 (2) In order to convert the initial SiOz completely into Si3N4 excess carbon is necessary. 5~'53 By

Silicon-based non-oride structural ceramics sing quartz sand and clay minerals low-priced pure and fine Si,N4 powder. The diimide pro oroducts are achievable 4, 55 Highly pure pow- cess is industrially used to process Sian, por ders can be produced by applying synthetic ders for advanced applications tarting materials derived from pyrolytic or Owing to the predominantly homogeneous ol- gel reactions nucleation, the reaction of SiCa or SiH4 with The liquid phase reaction of SiCl4 and NH3 NH3 in the gas phase provides fine and pure forming Si(NH)2 has been firstly performed by Si, powders with large specific surfaces Persoz in 1830 and has been investigated in (2-20 m" g)and high sintering activities detail by Blix and Wirbelauer, 7 Glemser and Similar to the liquid phase reaction of Sicl4 Naumann and Mazdiyasni and Cooke 9 later. with NH,(see eqn(3))amorphous Si3n4 pre- Billy showed that the ammonolysis of Sicla cursors are formed which have to be thermally under these conditions results in a polymerized converted to crystalline Si, N4. The use of the polysilicon diimide low priced Sicla requires the extraction of the n SiCl+6n nh C-RT [Si(NH)2I by-product NHCl(see eqn (3)). In contrast to this only H2 is evolved during the ammonolysis +anNeCI (3) of the expensive and spontaneous inflammable During the subsequent calcination the gener- (in air)SiHa. The synthesized amorphous pre- ated silicium diimide transforms into an amor- ceramic compounds are crystallized to mainly phous Si,N, accompanied by the evolution of -SiaNa at temperatures between 1200 and NH, or N,/Hz. At temperatures above 1200 C 1500C under nitrogen. Both methods are used the diffusion controlled crystallization to in industry to process commercially availab a-Si, n takes place possessing an activation Si,, qualities to produce engine and turbine energy of 306 kJ mol parts. In Table 1 the properties of some com mercially available Si3N4 powders are summa rized [SI(NH)]. 1200-1400°C a-Si3N +N2+3H 3 SILICON CARBIDE (SiC)-INTRINSIC The particle morphology, particle size and STRUCTURAL PROPERTIES AND phase composition is determined by the pro- SYNTHESIS cessing parameters such as temperature, reac tion time and impurities. 2. 0. The liquid phase The fundamental structural elements of the reaction of SiCl, with NH, gives extraordinary various polytypes of Sic are covalently(88% Table 1. Characteristics of commercially available Si, N,-powders derived from diffe measurements by the authors Powder type SN-E10 SN-ESP Grade GP LC 12-SX A200 Production process Liquid phas Liquid pha Gas phase Direct diimide nitridation Manufacturer Ube Industries, Ube Industries, H C Starck, H.C. Starck, Toshiba Ceramics, Tokyo Tokyo Berlin Tokyo Impurities(wt%) 18-21,(15) 0 <0001 <001 <0005 <0002 <0002 pecinc surtace are (98) Mean particle size 05,(055) 05,(079) 06,(048) 阝-SiN4(wt%) 5,(41) <10,(64)

Silicon-based non-oxide structural ceramics 17 using quartz sand and clay minerals low-priced products are achievable? 4"55 Highly pure pow￾ders can be produced by applying synthetic starting materials derived from pyrolytic 53 or sol-gel 52 reactions. The liquid phase reaction of SIC14 and NH3 forming Si(NH)2 has been firstly performed by Persoz 56 in 1830 and has been investigated in detail by Blix and Wirbelauer, 57 Glemser and Naumann 58 and Mazdiyasni and Cooke 59 later. Billy" showed that the ammonolysis of SiCI4 under these conditions results in a polymerized polysilicon diimide. nSiC14 + 6nNH30°C-- RT) [Si(NH)2], +4nNH4C1. (3) During the subsequent calcination the gener￾ated silicium diimide transforms into an amor￾phous Si3N 4 accompanied by the evolution of NH3 or NJH2. At temperatures above 1200°C the diffusion controlled crystallization to ~-Si3N461 takes place possessing an activation energy of 306 kJ moli 1200-- 14OO°C -- [Si(Nn)2], , ) ~-Si3N4 n pure and fine Si3N4 powder. The diimide pro￾cess is industrially used to process Si3N4 pow￾ders for advanced applications. Owing to the predominantly homogeneous nucleation, the reaction of SiCl4 or SiH4 with NH3 in the gas phase provides fine and pure Si3N4 powders with large specific surfaces 64"65 (2-20m 2 g-') and high sintering activities. Similar to the liquid phase reaction of SiCl4 with NH~ (see eqn (3)) amorphous Si3N4 pre￾cursors are formed which have to be thermally converted to crystalline Si3N 4. The use of the low priced SiCl4 requires the extraction of the by-product NH4C1 (see eqn (3)). In contrast to this only H2 is evolved during the ammonolysis of the expensive and spontaneous inflammable (in air) Sill4. The synthesized amorphous pre￾ceramic compounds are crystallized to mainly ~-Si3N 4 at temperatures between 1200 and 1500°C under nitrogen. Both methods are used in industry to process commercially available Si3N 4 qualities to produce engine and turbine parts. In Table 1 the properties of some com￾mercially available Si3N4 powders are summa￾rized. +N2+3H2 (4) The particle morphology, particle size and phase composition is determined by the pro￾cessing parameters such as temperature, reac￾tion time and impurities. 62"63 The liquid phase reaction of SiCl4 with NH3 gives extraordinary 3 SILICON CARBIDE (SIC) -- INTRINSIC STRUCTURAL PROPERTIES AND SYNTHESIS The fundamental structural elements of the various polytypes of SiC are covalently (88% Table 1. Characteristics of commercially available Si3N4-powders derived from different production processes. () Means measurements by the authors Powder type SN-E I O SN-ESP Grade G P LC 12-SX A 200 Production process Liquid phase Liquid phase Gas phase Direct Carbothermal diimide diimide nitridation reduction Manufacturer Ube Industries, Ube Industries, H.C. Starck, H.C. Starck, Toshiba Ceramics, Tokyo Tokyo Berlin Berlin Tokyo Impurities (wt%) O <2 (1.1) (1.0) 1.1-1.6 1.8-2.1, (1.5) 2.0 C <0-2 <0.2 <0.05 <0.2 0.9 CI < 0-01 0.01 < 0.1 < 0.001 -- Fe < 0.01 0.01 < 0.01 < 0.008 0.007 A1 < 0-005 -- < 0.004 < 0.005 0.2 Ca < 0-005 < 0.002 < 0.002 < 0.002 0.01 Specific surface area (9.8) (7.5) (12.2) (21.4) -- (m 2 g ') Mean particle size ds,, (#m) fl-Si3N4 (wt%) 0.5, (0.55) (0.64) 0.5, (0.79) 0.6, (0.48) 0.9 <5, (4.1) (<3) <10, (6.4) s8, (5-6) 2

18 W. Dressler: R. riedel covalence)bonded SiCA(sp" hybridization of Si) leads to the 2 H type whereas the doping with and CSi,(sp'hybridization of C)tetrahedra N and P forces the crystallization of the cubic (Figs 3 and 4). These tetrahedra are arranged in B-SiC. In contrast to the above mentioned irre slanes having common edges and one apex in versibility of the b/ la-transformation under the next plane of tetrahedra connecting the standard conditions Kieffer, Jepps and Page stacks. If the stacking sequence of the tetra- describe the reversible transformation of a-siC hedra is ABC a cubic zinc blend structure, des-(6 H)to B-Sic(3 C) by increasing the N2-pres gnated as B-SiC(Fig. 4), results, whilst the sure. sequence ABAB provides a hexagonal wurtzite Besides the synthetic SiC produced by carbo- structure,denoted as a-SiC (Fig 3). The hex- thermal reduction of SiO2, vapour phase reac gonal or rhombohedral a-SiC exists in many tions or thermal decomposition of silanes or polytypes(most frequent polymorphs: 2 H, 4 H, carbosilane some natural deposits are known 6 H and 15 R)composed of intermixed more a-SiC can be found in association with diamond complex arrangements of the tetrahedra planes in iron meteorites of the Canyon Diablo type resulting in large periods of stacking. The most and is denoted as Moissanite. Further natural common hexagonal a-SiC polytype 6 H can occurences are located in Bohemian volcanic derived from the cubic form by insertion of a breccias and Siberian Kimberlites. The cubic rotation(111> twin boundary after every three B-SiC polymorph was struck in the Green-River layers so that after six sequences the initial layer district, USA. position is obtained again. Despite this struc The industrial production of a-silicon carbide tural difference the density of all Sic-poly- is performed by the Acheson-process, -75a morphs is constant at 3. 17 gcm carbothermal reduction of Sio2. Using a graph The thermodynamic stability, the conditions ite electrode surrounded by a Sic rim for the of formation and the phase transformations of electrical coupling, a mixture of quartz sand or the Sic polymorphs have been intensively inves- crushed quartzite (58-65%0), graphite, petrol tigated by Knippenberg, Kieffer, Page, eum coke or ash-free anthracite(35-42%), Jepps and Page%, o and Heine. Controversly sodium chloride (1-2%)and wood chips to previous investigations, postulating B-SiC as (0. 5-1%)as additives is fused at temperatures a stable low temperature modification Knippen- between 2200 and 2400C, whereby the follow berg reported on B-SiC formation not only at ing reaction takes place low temperatures of about 1400oC but also at higher temperatures. Above 2000C surface dif- SiO, +3C-2200-240> SiC+2CO fusion leads to the irreversible transformation of B-Sic to the 6 H polytype of a-SiC. This 528 kJ mol- SiC behavior indicates that B-sic is a metastable SiC-modification which is only formed at lower The resulting a-SiC is coarse grained (Fig. 5), temperatures owing to the very small self diffu- has to be milled to the desired grit size and is sion coefficients of Si and C in SiC. Addition- divided into several qualities depending on the ally, the free enthalpy of B/oz-transformation amount of impurities. The inner part having a acting as the driving force of tra formation green color contains the purest material. The amounts to only -2 kJ mol- at T=2000 K amount of carbon, aluminium and other impur Heine?also calculated the hexagonal a-poly- ities increases continuously with the distance to types to be thermodynamically stable (lower the core and is accompanied by a change of the free energy) in comparison to the cubic B-SiC. SiC color from green to black. In order to Here, the formation of B-SiC is explained by the reduce the amount of metallic residues the pro- energetically preferred building up of parallel duced SiC powder is washed and leached. Sub ayers instead of the incorporation of twin sequently, the excess carbon is oxidized at lay boundaries by the insertion of antiparallel ori- 400C and the resulting oxide layer is removed ented layers typical for the a-SiC-structure by hydrofluoric acid Moreover, Jepps and P d p B-SiC can be produced by a modified Ache- showed that the addition of B stabilizes the 6H son- process at temperatures in the range of polytype whereas Al enhances the formation of 1500-1800C where a fine grained B-Sic is the 4 H polymorph. The presence of Al andn formed via a solid phase reaction. Gas phase

18 W. Dressier, R. Riedel covalence) bonded SiC4 (sp 3 hybridization of Si) and CSi 4 (sp 3 hybridization of C) tetrahedra (Figs 3 and 4). These tetrahedra are arranged in planes having common edges and one apex in the next plane of tetrahedra connecting the stacks. If the stacking sequence of the tetra￾hedra is ABC a cubic zinc blend structure, des￾ignated as //-SIC (Fig. 4), results, whilst the sequence ABAB provides a hexagonal wurtzite structure, denoted as a-SiC (Fig. 3). The hex￾agonal or rhombohedral a-SiC exists in many polytypes (most frequent polymorphs: 2 H, 4 H, 6 H and 15 R) composed of intermixed more complex arrangements of the tetrahedra planes resulting in large periods of stacking. The most common hexagonal a-SiC polytype 6 H can be derived from the cubic form by insertion of a rotation (111) twin boundary after every three layers so that after six sequences the initial layer position is obtained again. Despite this struc￾tural difference the density of all SiC-poly￾morphs is constant at 3.17 g cm -3. The thermodynamic stability, the conditions of formation and the phase transformations of the SiC polymorphs have been intensively inves￾tigated by Knippenberg, "6 Kieffer, 67 Page, 68 Jepps and Page 6''~'' and Heine. 73 Controversly to previous investigations, postulating //-SIC as a stable low temperature modification Knippen￾berg"" reported on//-SIC formation not only at low temperatures of about 1400°C but also at higher temperatures. Above 2000°C surface dif￾fusion leads to the irreversible transformation of //-SIC to the 6 H polytype of a-SiC. This behavior indicates that //-SIC is a metastable SiC-modification which is only formed at lower temperatures owing to the very small self diffu￾sion coefficients of Si and C in SiC. Addition￾ally, the free enthalpy of ///a-transformation acting as the driving force of transformation amounts to only -2kJ mol ' at T--2000K. 72 Heine ~' also calculated the hexagonal ~-poly￾types to be thermodynamically stable (lower free energy) in comparison to the cubic//-SIC. Here, the formation of//-SIC is explained by the energetically preferred building up of parallel layers instead of the incorporation of twin boundaries by the insertion of antiparallel ori￾ented layers typical for the a-SiC-structure. Moreover, Jepps and Page 6''~'' and Page °8 showed that the addition of B stabilizes the 6 H polytype whereas A1 enhances the formation of the 4 H polymorph. The presence of A1 and N leads to the 2 H type whereas the doping with N and P forces the crystallization of the cubic //-SIC. In contrast to the above mentioned irre￾versibility of the ///a-transformation under standard conditions Kieffer, 67 Jepps and Page 69 describe the reversible transformation of a-SiC (6 H) to//-SIC (3 C) by increasing the N2-pres￾sure. Besides the synthetic SiC produced by carbo￾thermal reduction of SiO2, vapour phase reac￾tions or thermal decomposition of silanes or carbosilanes some natural deposits are known. a-SiC can be found in association with diamond in iron meteorites of the Canyon Diablo type and is denoted as Moissanite. Further natural occurences are located in Boehemian volcanic breccias and Siberian Kimberlites. The cubic /~-SiC polymorph was struck in the Green-River district, USA. The industrial production of a-silicon carbide is performed by the Acheson-process, 73-~5 a carbothermal reduction of SiO2. Using a graph￾ite electrode surrounded by a SiC rim for the electrical coupling, a mixture of quartz sand or crushed quartzite (58-65%), graphite, petrol￾eum coke or ash-free anthracite (35-42%), sodium chloride (1-2%) and wood chips (0.5-1%) as additives is fused at temperatures between 2200 and 2400°C, whereby the follow￾ing reaction takes place: SiO2 + 3C 22,,o 2400°( ~ ~ SiC + 2CO -528 kJ mol-' SiC. (5) The resulting a-SiC is coarse grained (Fig. 5), has to be milled to the desired grit size and is divided into several qualities depending on the amount of impurities. The inner part having a green color contains the purest material. The amount of carbon, aluminium and other impur￾ities increases continuously with the distance to the core and is accompanied by a change of the SiC color from green to black. In order to reduce the amount of metallic residues the pro￾duced SiC powder is washed and leached. Sub￾sequently, the excess carbon is oxidized at 400°C and the resulting oxide layer is removed by hydrofluoric acid. //-SIC can be produced by a modified Ache￾son-process at temperatures in the range of 1500-1800°C where a fine grained //-SIC is formed via a solid phase reaction. 5"~" Gas phase

Silicon-based non-oxide structural ceramics such as CH, result in ultra fine and pure B-Sic the production proce powder properties dependent on reactions of SiH, or SiCl4 with hydrocarbons Table 2. Typical SiC powders. Additionally, the gas phase decom Production process Acheson Modified Gas phase position of organosilanes CH3SiCl3,( CH3)2 process Acheson reaction SiCl2(CH3)3 SiCI and(CH3)4Si or polycarbosi lanes [-R2Si-CH2-I,(R=CH3, H) leads to nano- Composition(wt%) crystalline B-SiC. The characteristic data of , sio 97 a-SiC >98 B-SiC 97.2 B-SiC SiC-powders produced by the different methods are listed in Table 2 00 0006 00009 4 CONVENTIONAL PROCESSING Diameter(um) 045 0·27 175 MICROSTRUCTURING AND MECHANICAL Specific surface area 14 (m2g) PROPERTIES OF MONOLITHIC Si,, AND SiC CERAMICS covalent and strongly directional chemical Classical ceramic materials consist mostly of bondings in Si, N, and Sic cause very low self oxides which are predominantly ionic materials. diffusion coefficients. 77.76 Hence the conditions Additionally, the bondings are nondirectior for bulk diffusion are unfavourable and sinter and the densification of these ceramics takes ing of covalent substances in general is diffi place by volume or grain boundary diffusion cult.79 chanced by vacancy formation due to non-stoi- Several methods have been used to overcome chiometry. In contrast to that, the highly this low sinterability of covalent ceramics. One of them is to enhance the sintering by applying an external isostatic pressure at high tempera- tures(HIP)which is unfortunately limited to small parts and high cost applications. Another part of the mater is produced in-situ by reaction from its ele- ments. The third way described in the following paragraphs is to add sintering aids. This process includes deagglomeration and mixing of starting powders, drying and sieving of the resulting mixtures, moulding into green bodies and sub- sequent sintering 4.1 Sili 4.1.1 Si of si In order to achieve elongated grain structures in the final microstructure increasing the frac ture toughness as described later Si,- ceramics are mainly produced from a-Si3N4-powders Oxides like MgO, Al2O3, Y2O3, rare earth oxides and mixtures of them are used as addi tives to produce dense ceramics by liquid phase sintering, which stages:.(i) the particle rearrangement by the development of capillary forces among the par ticles 4 due to the formation of an eutectic melt consisting of the used additives and the Sio2 on he si3N4 surface, (ii)the solution of a-Si Fig. 5. Crystalline SiC as received from the Acheson the diffusion of Si and n through the liquid phase and the reprecipitation on B-Si3N

Silicon-based non-oxide structural ceramics 19 reactions of Sill4 or SIC14 with hydrocarbons such as CH 4 result in ultra fine and pure fl-SiC powders? Additionally, the gas phase decom￾position of organosilanes CH~SiC13, (CH3)2 SiCI2 (CH3)3SiCI and (CH3)4Si or polycarbosi￾lanes [-R2Si-CH2-],, (R=CH3, H) leads to nano￾crystalline {/-SIC. The characteristic data of SiC-powders produced by the different methods are listed in Table 2. 4 CONVENTIONAL PROCESSING, MICROSTRUCTURING AND MECHANICAL PROPERTIES OF MONOLITHIC Si3N4 AND SiC CERAMICS Classical ceramic materials consist mostly of oxides which are predominantly ionic materials. Additionally, the bondings are nondirectional and the densification of these ceramics takes place by volume or grain boundary diffusion enhanced by vacancy formation due to non-stoi￾chiometry. In contrast to that, the highly Fig. 5. Crystalline SiC as received from the Acheson process. Table 2. Typical SiC-powder properties dependent on the production process Production process Acheson Modified Gas phase process Acheson reaction process Composition (wt%) SiC 97 a-SiC > 98 fl-SiC 97.2/~-SiC Free C l'4 0.4 1.0 Free SiO2 0.7 0.3 1.3 Fe 0.06 0.04 0.006 Ai 0.01 0-03 0.0017 Ca -- -- 0.0009 Diameter (/~m) 0.45 0-27 0'3 Specific surface area 14 17.5 -- (m 2 g ') covalent and strongly directional chemical bondings in Si3N, and SiC cause very low self diffusion coefficients. 77"76 Hence, the conditions for bulk diffusion are unfavourable and sinter￾ing of covalent substances in general is diffi￾cult. TM Several methods have been used to overcome this low sinterability of covalent ceramics. One of them is to enhance the sintering by applying an external isostatic pressure at high tempera￾tures (HIP) which is unfortunately limited to small parts and high cost applications. Another is reaction sintering, where part of the material is produced in-situ by reaction from its ele￾ments. The third way described in the following paragraphs is to add sintering aids. This process includes deagglomeration and mixing of starting powders, drying and sieving of the resulting mixtures, moulding into green bodies and sub￾sequent sintering. 4.1 Silicon nitride ceramics 4. l. 1 Sintering of Si,N~-ceramics In order to achieve elongated grain structures in the final microstructure increasing the frac￾ture toughness as described later Si~N4-ceramics are mainly produced from ct-Si~N4-powders. Oxides like MgO, A120~, Y203, rare earth oxides and mixtures of them are used as addi￾tives to produce dense ceramics by liquid phase sintering, which can be subdivided into three stages: 8'''~' (i) the particle rearrangement by the development of capillary forces among the par￾ticles 84 due to the formation of an eutectic melt consisting of the used additives and the Si02 on the Si3N4 surface, (ii) the solution of ~-Si~N~, the diffusion of Si and N through the liquid phase and the reprecipitation on /~-Si~N4

W Dressler, R. Riedel nuclei followed by p-Si, N4 grain coarsening" mined by X-ray analysis 2)from 6 5N um and (iii) the coalescence of ystals, (undoped)to 9 3 and 12. 1 N um. Under the which is of limited importance due to the low assumption that B-Si, N4 nucleation is negligible volume diffusion in si this was explained by a higher number of grow In this context beneath others particularly ing grains during aB-transformation if B-Si, N two contradictory requirements are to be taken doping is applied. The influence of seeding of into account. On the one hand optimal densifi- fine grained a-rich UBE SN-E10(E10, mean cation by using high volume fractions of addi- crystallite size r=0-06 um) with B-Si, NA Denka tives having low melting temperatures and low (mean crystallite size r=0. 14 um) is shown in iscosities. On the other hand high temperature Fig. 6. The quantitative analysis of the micro- resistance, which is deteriorated by the soften- structure was performed by measuring the two ing of a secondary phase. Additionally, the dimensional size shape distribution of more microstructural development has to be control- than 2000 grains on polished and plasma etched led with respect to high toughness and specimen and by the subsequent stereological strength. In general, it is impossible to meet computation of the three-dimensional grain size all requirements and therefore silicon nitride shape distributions. The size-shape distribu- ceramics have to be designed for specific appli- tions (weighted by volume)reveal that the cations. Some basic principles used to control B-doping leads to a decrease of the volume frac- the microstructural development and to tion of grains having a length smaller than improve the mechanical properties of silicon 0.5 um from 11. 5 vol%(E10)to 0. 5 vol% nitride ceramics are presented below (Denka). Simultaneously, the mean grain length and grain diameter increases from 0-36to 4.1.2 Microstructural design 0-80 um and from 0- 12 to 0- 45 um, respectively The tailoring of Si,NA microstructure can be Additionally, low B-doping(Denka/E10 4/96) used to produce high fracture toughness com- leads to an enhanced grain growth in the length bined with high reliability and strength if con- direction, but to a decrease in aspect ratio of trolled grain growth can be achieved. Otherwise the coarser grains owing to the large grain abnormally grown B-Si3N4-crystals act as crack width of the added B-particles. Higher amounts initiation sites and thus reduce the strength of of B-nuclei(Denka/E10 20/80) results in a the material. Herewith it has to be taken into reduction of the maximum grain length and account that the microstructural development aspect ratio USing pure B-Si3N4(Denka)start- of Sia na depends on the starting powder charac- ing powders an equiaxed microstructure is pro- teristics, the used additive system and the sin- duced possessing a low mean aspect ratio of 1. 8 tering conditions. Additionally, the morphology in comparison to specimens sintered from a-rich of the B-Si, N4-crystals in the final microstruc- powder(E10) having a mean aspect ratio of 2.7 ture is determined by their growth anisotropy. Moreover, an overall grain coarsening was The preferred growth direction is perpendicular observed. From this investigation it was con- to the basal plane())of the formed hex- cluded that, if the B-Si3N4 nuclei density agonal prisms. The influence of the intrinsic reaches a certain value, depending on the grain powder properties on the final microstructure size distribution of the starting powder, a dis- was shown by several authors in different addi- solution of the smaller B-Si3 N, particles in the tive systems. 3 B3, 85- Particularly, the phase liquid secondary phase occurs at an early stage ratio and crystallite size of a-and B-phase in the of a/B-transformation resulting in a coarsening starting powder have been found to be the key of the final microstructure. 3 Additionally, it factors in the microstructural development. was deduced that the often observed abnormal On the one hand the doping of coarse a-Si3N4- grain growth in Siana is due to a kinetic and powders (UBE SN-ESP, mean crystallize size energetic growth advantage of B-Si3N4-crystals r=0-10 um) with coarse B-Si3N2-nuclei(Denka, having a large basal plane(1001)). .These mean crystallite size r=0 14 um) led to a grain results show that tailoring of the final Si3 na refinement in the sintered ceramics. The grains microstructure becomes possible by controlling er unit area increased from 0 56N um for the B-Si3Na-nuclei density, morphology and size the undoped material to 0 72 and 0-86N um distribution in the starting powder. In order to by increasing the B-Si3Na-nuclei density(deter- optimize mechanical properties of the final

20 W. Dressier, R. Riedel nuclei ~° followed by//-Si3N4 grain coarsening "3 and (iii) the coalescence of //-Si3N4 crystals, which is of limited importance due to the low volume diffusion in Si3N 4. In this context beneath others particularly two contradictory requirements are to be taken into account. On the one hand optimal densifi￾cation by using high volume fractions of addi￾tives having low melting temperatures and low viscosities. On the other hand high temperature resistance, which is deteriorated by the soften￾ing of a secondary phase. Additionally, the microstructural development has to be control￾led with respect to high toughness and strength."-" In general, it is impossible to meet all requirements and therefore silicon nitride ceramics have to be designed for specific appli￾cations. Some basic principles used to control the microstructural development and to improve the mechanical properties of silicon nitride ceramics are presented below. 4.1.2 Microstructural design The tailoring of SigN4 microstructure can be used to produce high fracture toughness com￾bined with high reliability and strength if con￾trolled grain growth can be achieved. Otherwise abnormally grown //-Si3Na-crystals act as crack initiation sites and thus reduce the strength of the material." Herewith it has to be taken into account that the microstructural development of Si3N 4 depends on the starting powder charac￾teristics, the used additive system and the sin￾tering conditions. Additionally, the morphology of the //-Si3Na-crystals in the final microstruc￾ture is determined by their growth anisotropy. The preferred growth direction is perpendicular to the basal plane ({001}) of the formed hex￾agonal prisms. "4 The influence of the intrinsic powder properties on the final microstructure was shown by several authors in different addi￾tive systems.' "' 3. ~3. ~5-,,, Particularly, the phase ratio and crystallite size of a- and//-phase in the starting powder have been found to be the key factors in the microstructural development.' "'3 On the one hand the doping of coarse a-Si3N a￾powders (UBE SN-ESP, mean crystallize size r=0"10 pm) with coarse //-Si3N4-nuclei (Denka, mean crystallite size r=0"14 #m) led to a grain refinement in the sintered ceramics. The grains per unit area increased from 0.56 N pm 2 for the undoped material to 0.72 and 0.86 N #m 2 by increasing the fi-Si3N4-nuclei density (deter￾mined by X-ray analysis '2) from 6-5 N #m -3 (undoped) to 9.3 and 12.1 N pm 3. Under the assumption that fl-Si3N4 nucleation is negligible this was explained by a higher number of grow￾ing grains during a/fl-transformation if fl-Si3N4 doping is applied. The influence of seeding of fine grained a-rich UBE SN-E10 (El0, mean crystallite size r=0.06 #m) with fl-Si3N4 Denka (mean crystaUite size r=0.14/~m) is shown in Fig. 6.'3 The quantitative analysis of the micro￾structure was performed by measuring the two￾dimensional size shape distribution of more than 2000 grains on polished and plasma etched specimen and by the subsequent stereological computation of the three-dimensional grain size shape distributions. 92 The size-shape distribu￾tions (weighted by volume) reveal that the //-doping leads to a decrease of the volume frac￾tion of grains having a length smaller than 0.5pm from 11.5vo1% (El0) to 0.5vo1% (Denka). Simultaneously, the mean grain length and grain diameter increases from 0.36 to 0.80 pm and from 0"12 to 0-45 #m, respectively. Additionally, low //-doping (Denka/E10 4/96) leads to an enhanced grain growth in the length direction, but to a decrease in aspect ratio of the coarser grains owing to the large grain width of the added//-particles. Higher amounts of //-nuclei (Denka/E10 20/80) results in a reduction of the maximum grain length and aspect ratio. Using pure //-Si3N 4 (Denka) start￾ing powders an equiaxed microstructure is pro￾duced possessing a low mean aspect ratio of 1.8 in comparison to specimens sintered from a-rich powder (El0) having a mean aspect ratio of 2.7. Moreover, an overall grain coarsening was observed. From this investigation it was con￾cluded that, if the fl-Si3N 4 nuclei density reaches a certain value, depending on the grain size distribution of the starting powder, a dis￾solution of the smaller fl-Si3N 4 particles in the liquid secondary phase occurs at an early stage of a///-transformation resulting in a coarsening of the final microstructure.'" ,3 Additionally, it was deduced that the often observed abnormal grain growth in Si3N 4 is due to a kinetic and energetic growth advantage of //-Si3Na-crystals having a large basal plane ({001}). ''''3 These results show that tailoring of the final Si3N4 microstructure becomes possible by controlling the//-Si3Nn-nuclei density, morphology and size distribution in the starting powder. In order to optimize mechanical properties of the final

Silicon -based non-oxide structural ceramics 21 product the starting powders should possess can be analyzed in terms of the pullout mode narrow B-Si3N4 grain size distribution and have developed by becher et al.This model faceted, elongated B-Si3n4 crystals explains the toughening behavior of whisker Ceramics prepared from such Si3N4-powders reinforced ceramic matrix composites. Accord should exhibit microstructures containing a ingly, the fracture toughness( K,d)of materials large amount of elongated grains increasing the which reveal mainly debonding and pullout fracture toughness of the material without exag- (crack deflection is neglected) depends on the gerated grown grains, which would deteriorate matrix toughness(KTe), a constant(A)as wel the materials strength. Therefore, these cera the volume fraction (V) and the diameter mics are expected to combine both high (Dmi)of the reinforcing particles (egn(6)) strength and high fracture toughness. The rela Herewith the constant (A)depends on the tion between microstructure and mechanical strength and elastic moduli of the reinforcing properties is discussed in the following para graph phase as well as the Poisson ratio, elastic mod ulus and fracture energy of the matrix and the 4.1.3 Microstructure and mechanical properties of fracture energy of the interface between the Si,N ceramics reinforcing phase and the matrix The interconnection between the Si, Na-cera mics microstructure and its fracture toughness KIe=[(KI)2+A V Dmin] 10 D/E104/96 0102030405,060 010203,04,0506.0 Leng th lum Length fum D/E1020/80 20 886=0E 01,02,03040506,0 °102080405060 gth fum] Length lum ig. 6. Microstructural development of Si, N.-ceramics. E10: a-Si, N,(UBE SN-E10, UBE Industries, Japan) containi 4. 1 vol% B-Si3N4; D/E10 4/ 96: a-Si3N4(E10) seeded with 4 vol %o B-Si3N )DE102080 vol% B-Si,N, (D); D: B-Si, N,(SN-BS, Denka, Japan) containing 2. 5 vol% x-Si, N

Silicon-based non-oxide structural ceramics 21 product the starting powders should possess a narrow fl-Si3N4 grain size distribution and have faceted, elongated fl-Si3N4 crystals. Ceramics prepared from such Si3N4-powders should exhibit microstructures containing a large amount of elongated grains increasing the fracture toughness of the material without exag￾gerated grown grains, which would deteriorate the materials strength. Therefore, these cera￾mics are expected to combine both high strength and high fracture toughness. The rela￾tion between microstructure and mechanical properties is discussed in the following para￾graph. 4.1.3 Microstructure and mechanical properties of Si3N4-ceramics The interconnection between the Si,N4-cera￾mics microstructure and its fracture toughness can be analyzed in terms of the pullout model developed by Becher et al. 93 This model explains the toughening behavior of whisker reinforced ceramic matrix composites. Accord￾ingly, the fracture toughness (K,c) of materials which reveal mainly debonding and pullout (crack deflection is neglected) depends on the matrix toughness (K','c), a constant (A) as well as the volume fraction (V 0 and the diameter (D,,~n) of the reinforcing particles (eqn (6)). Herewith the constant (A) depends on the strength and elastic moduli of the reinforcing phase as well as the Poisson ratio, elastic mod￾ulus and fracture energy of the matrix and the fracture energy of the interface between the reinforcing phase and the matrix. KI~=[(K¢]~c)2-kA. Vf. Dm,~]'/2 (6) / ..~'° ....""" ." .. ............................................................................................. ~' 4o ..." .. ........................................................................................................................ ....."" SO ' , " ' ........................................................................................................................................... T SO : .."" ... • ........................................................................................................................... ..~Pl~ -- ..... - .-. --" - """""""'"I 0 , ... ...... ." . ..... P=~+:4 ° .... 0 ~ i I.'" v' "(::" i:::~""i:::"~i""i::~V~1 2 ~,~X 0 1,0 2,0 3,0 4,0 5,0 8,0 -- 0 1,0 2,0 3,0 4,0 5,0 6,0" -- Length rum] Length rum] Z o,o,o,l l .................................................................................... o ....................... ~)~o~ ~ .............. I ~ ...... ..... • = " • , :: if:! > o~~::f_2L~ ,, ..... :::: :::i:,ii I 0 1,0 2,0 3,0 4,0 5,0 6,01 Y~ 00¢" 1,0 Length rum] Length rum] Fig. 6. Microstructural development of Si.~N4-ceramics. El0:0~-Si3N4 (UBE SN-E10, UBE Industries, Japan) containing 4.1 vol% fl-Si~N4; D/E10 4/96: 0~-8i3N 4 (El0) seeded with 4 vol% fl-Si3N4 (D); D/E10 20/80:~-Si~N4 (El0) seeded with 20 vol% fi-Si3N4 (D); D: fl-Si~N4 (SN-BS, Denka, Japan) containing 2.5 vol% ct-Si3N4. '~

W. Dressler, R. Riedel observed the chemistry of the grain boundary Experimental results. 2 showed that Si,N phase to be an important parameter in the field ceramics having exclusively grains with aspect of toughening of SiaNa-ceramics. Crack propa ratios below 4 reveal fracture toughness values gation experiments in p-Si3NA-whisker doped of about 5.5 MPa m /. This led to the concl aluminium-yttrium oxynitride glasses with sion that the grain fraction with an aspect ratio fixed nitrogen content revealed that the deb- smaller than 4 can be regarded as the matrix onding length and, hence, the toughening and grains having an aspect ratio higher than 4 response of the whiskers decreases with risin can be considered as the reinforcing particles. Al2 O3 content due to an increase in interface By means of the above described quantitative energy. This experiment points out that fracture microstructural analysis method the volume toughness tailoring requires both the micro- weighted diameter of grains(2, Vr Dmin )has structure design and control of grain boundary been calculated from the measured grain size chemistry distributions According to Irwin the strength of ceramic In Fig. 7 the dependence of fracture tough- materials(oB) showing no plastic deformation ness on the volume weighted diameter of elon- can be correlated with the materials fracture gated grains(aspect ratio >4)is shown for toughness(Kle), a geometric factor(Y) and the Si,n4 ceramics densified by 10.7 wt% defec Y, O, and 3 6 wt% Al,O3 as sintering aids This plot shows that the fracture toughness grain diameter rises as described by eqn(6)and that fracture toughness values of 10-3 MPa'm can be achieved. It was concluded that the This reveals that the strength of brittle ceramics pull-out model fits in principle the experimental is not constant but depends on the defect size data. The deviations are explained by the influ- distribution incorporated in the ceramics due to ence of the neglected crack deflection which has processing defects and microstructure features been calculated from Bengisu et al. 4 to be It has been shown that the failure probability about 2 MPa m/2 even in composite materials (P)of brittle ceramics can be described by the having a low difference in the elastic properties two parametric Weibull-distribution. 00 Io1 between the matrix and the reinforcing phase Thus, very high strengthened Si,Na ceramics can be produced by grain coarsening due to p long heat treatments at high temperatures additionally, Becher et al. 95 and others b,, 96-9K Here(m)is the Weibull modulus and (oo) is a parameter. By plotting InIn [1/(1-P)] vs(oB) 200 straight line having the slope (m) results Therefore, the Weibull parameter (m)describes the strength variation of the particular material In Fig. 8 the four point bending strength di tributions of gas pressure(10 MPa N2 pressure) sintered Al2O/Y2O3(10 vol%) containing Si, -ceramics with fine and coarse microstruc tures derived from different starting powders are shown. i In the case of the material derived from UBE SN-E10 starting powder(UBE E-10 0,6 0.8 UBE Industries, Japan) the coarsening of the Dmin(a>4) um microstructure leads to a decrease of the maxi- mum and mean strength o an ig.7. Relation between fracture toughness and volume increase of the Weibull modulus from 13. 5 to weighted diameter of elongated(aspect ratio >4) grains for Si, N. ceramics densified by liquid phase sinteri 46, which means that the reliability of the (107wt%Y2O3+36wt%Al2O3)112 material can be improved by high temperature

22 W. Dressier, R. Riedel Experimental results TM showed that Si3N4 - ceramics having exclusively grains with aspect ratios below 4 reveal fracture toughness values of about 5.5 MPa'm ~/2. This led to the conclu￾sion that the grain fraction with an aspect ratio smaller than 4 can be regarded as the matrix and grains having an aspect ratio higher than 4 can be considered as the reinforcing particles. By means of the above described quantitative microstructural analysis method the volume weighted diameter of grains (Z,~'DL~,) has been calculated from the measured grain size distributions. In Fig. 7 the dependence of fracture tough￾ness on the volume weighted diameter of elon￾gated grains (aspect ratio >4) is shown for Si3N4 ceramics densified by using 10.7wt% Y20~ and 3.6wt% A1203 as sintering aids. ''''2 This plot shows that the fracture toughness increases by square if the volume weighted grain diameter rises as described by eqn (6) and that fracture toughness values of 10.3 MPa. m '/2 can be achieved. It was concluded" that the pull-out model fits in principle the experimental data. The deviations are explained by the influ￾ence of the neglected crack deflection which has been calculated from Bengisu et al. 94 to be about 2 MPa.m '/2 even in composite materials having a low difference in the elastic properties between the matrix and the reinforcing phase. Thus, very high strengthened Si3N4 ceramics can be produced by grain coarsening due to long heat treatments at high temperatures. Additionally, Becher et al. 95 and others 8"'~6-'~" 200 E 150 -L, n 100 50 Fig. 7. 00 I , i ' 0,2 0,4 0 6 0,8 Dmin(a>4) [pro] Relation between fracture toughness and volume weighted diameter of elongated (aspect ratio >4) grains for Si,N4-ceramics densified by liquid phase sintering (10"7 wt% Y203 +3"6 wt% A1203).' i, J2 observed the chemistry of the grain boundary phase to be an important parameter in the field of toughening of Si3N4-ceramics. Crack propa￾gation experiments "5 in fl-Si3Nn-whisker doped aluminium-yttrium oxynitride glasses with a fixed nitrogen content revealed that the deb￾onding length and, hence, the toughening response of the whiskers decreases with rising A1203 content due to an increase in interface energy. This experiment points out that fracture toughness tailoring requires both the micro￾structure design and control of grain boundary chemistry. According to Irwin ~° the strength of ceramic materials (o-,3) showing no plastic deformation can be correlated with the materials fracture toughness (K,c), a geometric factor (Y) and the defect size (a): glc a,= (7) Y" £d:a This reveals that the strength of brittle ceramics is not constant but depends on the defect size distribution incorporated in the ceramics due to processing defects and microstructure features. It has been shown that the failure probability (P) of brittle ceramics can be described by the two parametric Weibull-distribution.'"" '"' E P=l-exp - -- . (8) \(70/ J Here (m) is the Weibull modulus and (a,,) is a parameter. By plotting lnln [1/(I-P)] vs (aB) a straight line having the slope (m) results. Therefore, the Weibull parameter (m) describes the strength variation of the particular material. In Fig. 8 the four point bending strength dis￾tributions of gas pressure (10 MPa N2 pressure) sintered AI203//Y203 (10 vol%) containing Si3Na-ceramics with fine and coarse microstruc￾tures derived from different starting powders are shown." In the case of the material derived from UBE SN-E10 starting powder (UBE E-10: UBE Industries, Japan) the coarsening of the microstructure leads to a decrease of the maxi￾mum and mean strength as well as to an increase of the Weibull modulus from 13.5 to 46, which means that the reliability of the material can be improved by high temperature