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 tetrahedra is ABC a cubic zinc blend structure, designated as //-SIC (Fig. 4), results, whilst the sequence ABAB provides a hexagonal wurtzite structure, denoted as a-SiC (Fig. 3). The hexagonal 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 structural difference the density of all SiC-polymorphs 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 investigated 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 Knippenberg"" reported on//-SIC formation not only at low temperatures of about 1400°C but also at higher temperatures. Above 2000°C surface diffusion 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 diffusion coefficients of Si and C in SiC. Additionally, 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 ~-polytypes 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 oriented 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 irreversibility 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-pressure. Besides the synthetic SiC produced by carbothermal reduction of SiO2, vapour phase reactions 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 graphite electrode surrounded by a SiC rim for the electrical coupling, a mixture of quartz sand or crushed quartzite (58-65%), graphite, petroleum 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 following 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 impurities 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 produced SiC powder is washed and leached. Subsequently, 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 Acheson-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 decomposition of organosilanes CH~SiC13, (CH3)2 SiCI2 (CH3)3SiCI and (CH3)4Si or polycarbosilanes [-R2Si-CH2-],, (R=CH3, H) leads to nanocrystalline {/-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-stoichiometry. 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 sintering of covalent substances in general is difficult. 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 temperatures (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 elements. 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 subsequent 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 fracture 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 additives 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 particles 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 densification by using high volume fractions of additives having low melting temperatures and low viscosities. On the other hand high temperature resistance, which is deteriorated by the softening of a secondary phase. Additionally, the microstructural development has to be controlled 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 applications. 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 combined with high reliability and strength if controlled 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 characteristics, the used additive system and the sintering conditions. Additionally, the morphology of the //-Si3Na-crystals in the final microstructure is determined by their growth anisotropy. The preferred growth direction is perpendicular to the basal plane ({001}) of the formed hexagonal prisms. "4 The influence of the intrinsic powder properties on the final microstructure was shown by several authors in different additive 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 apowders (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 (determined 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 growing 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 microstructure was performed by measuring the twodimensional 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 distributions (weighted by volume) reveal that the //-doping leads to a decrease of the volume fraction 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) starting powders an equiaxed microstructure is produced 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 concluded that, if the fl-Si3N 4 nuclei density reaches a certain value, depending on the grain size distribution of the starting powder, a dissolution 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 exaggerated grown grains, which would deteriorate the materials strength. Therefore, these ceramics are expected to combine both high strength and high fracture toughness. The relation between microstructure and mechanical properties is discussed in the following paragraph. 4.1.3 Microstructure and mechanical properties of Si3N4-ceramics The interconnection between the Si,N4-ceramics 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. Accordingly, 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 modulus 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 conclusion 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 toughness on the volume weighted diameter of elongated 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 influence 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 propagation experiments "5 in fl-Si3Nn-whisker doped aluminium-yttrium oxynitride glasses with a fixed nitrogen content revealed that the debonding 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 microstructure 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 distributions of gas pressure (10 MPa N2 pressure) sintered AI203//Y203 (10 vol%) containing Si3Na-ceramics with fine and coarse microstructures 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 maximum 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