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《复合材料 Composites》课程教学资源（学习资料）第二章增强体_fiber51 Fine diameter ceramic fibres

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Fine diameter ceramic ®bres A.R. Bunsell *, M.-H. Berger Ecole des Mines de Paris, Centre des MateÂriaux, BP 87, 91003 Evry Cedex, France Received 28 January 2000; accepted 12 March 2000 Abstract Two families of small diameter ceramic ®bres exist. The oxide ®bres, based on alumina and silica, which were initially produced as refractory insulation have also found use as reinforcements for light metal alloys. The production of SiC based ®bres made possible the development of ceramic matrix composites. Improved understanding of the mechanisms which control the high temperature behaviour of these latter ®bres has led to their evolution towards a near stoichiometric composition which results in strength retention at higher temperatures and lower creep rates. The SiC ®bres will however be ultimately limited by oxidation so that there is an increasing interest in complex two phase oxide ®bres composed of a-alumina and mullite as candidates for the reinforcement of ceramic matrices for use at very high temperatures. These ®bres show low creep rates, comparable to the SiC based ®bres but are revealed to be sensitive to alkaline contamination. # 2000 Elsevier Science Ltd. All rights reserved. Keywords: Al2O3; Creep; Fibres; Microstructure; SiC 0955-2219/00/$ - see front matter # 2000 Elsevier Science Ltd. All rights reserved. PII: S0955-2219(00)00090-X Journal of the European Ceramic Society 20 (2000) 2249±2260 Contents 1. Introduction......................................................................................................................................................... 2250 2. Silicon carbide ®bres from organic precursors.....................................................................................................2250 2.1. SiC from an oxygen cured precursor route..................................................................................................2250 2.2. Electron cured precursor ®laments .............................................................................................................. 2251 2.3. Near stoichiometric ®bres ............................................................................................................................2252 3. Oxide ®bres..........................................................................................................................................................2254 3.1. Alumina silica ®bres..................................................................................................................................... 2254 3.1.1. The Sal ®bre...................................................................................................................................2254 3.1.2. The Altex ®bre .................................................................................................................................. 2255 3.1.3. The Nextel 312-440 ®bres..................................................................................................................2255 3.2. Alpha-alumina ®bres.................................................................................................................................... 2256 3.2.1. Fully dense a-alumina ®bres .............................................................................................................2256 3.2.2. Porous a-alumina ®bres ....................................................................................................................2257 3.3. Alpha-alumina ®bres containing a second phase.........................................................................................2257 3.3.1. Zirconia-reinforced alumina ®bres.................................................................................................... 2257 3.3.2. The Nextel 720 alumina mullite ®bre................................................................................................2257 4. Conclusions..........................................................................................................................................................2259 References ................................................................................................................................................................ 2259 * Corresponding author. E-mail address: anthony.bunsell@mat.ensmp.fr (A.R. Bunsell)

A.R. Bunsell, M.-H. Berger/ of the European Ceramic Society 20(2000)2249-2260 1. Introduction is around I um as at this diameter the fibres become a health hazard, if inhaled, as they block the alveolar Small diameter ceramic fibres have undergone structure of the lungs. Also related to the ease of con- hanges since their early development due to the verting the fibres into preforms is the desire for a strain to for reinforcements in structural ceramic matrix failure of around 1% and as a Youngs modulus of 200 te(CMC) materials to be used in air at temperatures GPa or more is required, this imposes a room tempera above 1000C. There now exists a range of oxide and non- ture strength of more than 2 GPa. Competing material oxide fibres with diameters in the range of 10 to 20 um are usually dense so that a specific gravity of less than 5 which are candidates as reinforcements. Applications would be desirable. The fibres are destined to be used at envisaged are in gas turbines, both aeronautical and high temperature and in air so that long term chemical, ground based, heat exchangers, first containment walls microstructural and mechanical stability up to and pre for fusion reactors as well as uses for which no matrix is ferably above, 1500C is required. This means that the necessary such as candle filters for high temperature gas structure of the ceramic fibre should not evolve and it filtration should exhibit low creep rates no greater than those of Initially ceramic fibres were produced in the early nickel based alloys. Lastly, low reactivity with the 1970s for use as refractory insulation which required the matrix is required if the crack stopping process which is material to withstand high temperatures, typically up to the basis of CMC tenacity is to be achieved 1600 C, in air but under no applied load. The fibres which were developed at this time were the discontinuous Saffil fibres, with diameters of 3 um, introduced by ICI 2. Silicon carbide fibres from organic precursors in 1972. 2 These fibres were made by the blow extrusion of a sol and consisted of y-alumina with 3% SiO2 to 2. 1. SiC from an oxygen cured precursor route inhibit grain growth and control porosity. The Nextel 12 fibre introduced by 3M around 1974 was also made The work of Yajima and his colleagues in Japan was by a sol route and was an amorphous fibre with a mullite first published in the mid-1970s. The Nicalon and Tyr composition. Both of these fibres are still in production. anno fibres produced, respectively, by Nippon Carbon and Later in the same decade Du Pont produced Fibre FP Ube Industries are the commercial results of this work. which was the first continuous a-alumina fibre and was These fibres are produced by the conversion of, respec made specifically for the reinforcement of aluminium. 3 tively polycarbosilane(PCS)and polytitanocarbosilane However it was the commercial production of the Sic(PTC) precursor fibres which contain cycles of six atoms based Nicalon fibres" in 1982 by Nippon Carbon which arranged in a similar manner to the diamond structure allowed ceramic matrix composites to be developed. The of B-SiC. The molecular weight of this polycarbosilane fibres were initially used by SEP to replace carbon fibres is low, around 1500, which makes drawing of the fibre in carbon-carbon composites used in rocket nozzles. difficult. The addition of around 2% wt. of titanium This improved the oxidative resistance of the material. achieved by the grafting of titanium alkoxide between The carbon matrix was then replaced by Sic to make the PCs chains in the Ube precursor, increases molecular the first Sic-Sic composites which could be considered weight slightly, which may help with drawing and also for applications at higher temperatures than those at may slightly increase thermal resistance by the creation which nickel based super alloys could be used. Efforts of Ti-C bonds at high temperatures In these polymers have been made since this time to improve the high tem- methyl groups(CH3) in the polymer are included as perature behaviour of small diameter SiC fibres by mak side groups to the-(Si-C)- main chain so that during ing them with compositions increasingly approaching pyrolysis hydrogen is produced, leaving a residue of free stoichiometry. 5,6 However these fibres are inherently carbon. The production of the first generations of SiC limited by oxidation at very high temperatures. As a based fibres involved subjecting the precursor fibres to result of this limitation a renewal of interest has occurred heating in air at around 200 C to produce cross-linking in oxide systems as a means of making reinforcements of the structures. This oxidation makes the fibres infu capable of operating in air at even higher temperatures sible but has the draw back of introducing oxygen into than the sic fibres the structure which after pyrolysis. The ceramic An important characteristic needed in a ceramic fibre fibres are obtained by a controlled increase in tempera reinforcement is flexibility so that preforms can be made ture in an inert atmosphere up to 1200C. The properties by weaving or other related technologies. This is ensured, and composition of these fibres are shown in Table 1 with materials having even the highest Youngs modul The fibres obtained by this route have a by a small diameter, as flexibility is related to the reci- appearance when observed in SEM, as can be seen procal of the fourth power of the diameter. A diameter of Fig. 1, however a closer examination reveals that they the order of 10 um is therefore usually required for contain a majority of B-siC, of around 2 nm but also ceramic reinforcements. A lower limit in fibre diameter significant amounts of free carbon of less than Inm and

1. Introduction Small diameter ceramic ®bres have undergone great changes since their early development due to the need for reinforcements in structural ceramic matrix composite (CMC) materials to be used in air at temperatures above 1000C. There now exists a range of oxide and nonoxide ®bres with diameters in the range of 10 to 20 mm which are candidates as reinforcements.1 Applications envisaged are in gas turbines, both aeronautical and ground based, heat exchangers, ®rst containment walls for fusion reactors as well as uses for which no matrix is necessary such as candle ®lters for high temperature gas ®ltration. Initially ceramic ®bres were produced in the early 1970s for use as refractory insulation which required the material to withstand high temperatures, typically up to 1600C, in air but under no applied load. The ®bres which were developed at this time were the discontinuous Sal ®bres, with diameters of 3 mm, introduced by ICI in 1972.2 These ®bres were made by the blow extrusion of a sol and consisted of g-alumina with 3% SiO2 to inhibit grain growth and control porosity. The Nextel 312 ®bre introduced by 3M around 1974 was also made by a sol route and was an amorphous ®bre with a mullite composition. Both of these ®bres are still in production. Later in the same decade Du Pont produced Fibre FP which was the ®rst continuous a-alumina ®bre and was made speci®cally for the reinforcement of aluminium.3 However it was the commercial production of the SiC based Nicalon ®bres4 in 1982 by Nippon Carbon which allowed ceramic matrix composites to be developed. The ®bres were initially used by SEP to replace carbon ®bres in carbon±carbon composites used in rocket nozzles. This improved the oxidative resistance of the material. The carbon matrix was then replaced by SiC to make the ®rst SiC±SiC composites which could be considered for applications at higher temperatures than those at which nickel based super alloys could be used. Eorts have been made since this time to improve the high temperature behaviour of small diameter SiC ®bres by making them with compositions increasingly approaching stoichiometry.5,6 However these ®bres are inherently limited by oxidation at very high temperatures. As a result of this limitation a renewal of interest has occurred in oxide systems as a means of making reinforcements capable of operating in air at even higher temperatures than the SiC ®bres. An important characteristic needed in a ceramic ®bre reinforcement is ¯exibility so that preforms can be made by weaving or other related technologies. This is ensured, with materials having even the highest Young's moduli, by a small diameter, as ¯exibility is related to the reciprocal of the fourth power of the diameter. A diameter of the order of 10 mm is therefore usually required for ceramic reinforcements. A lower limit in ®bre diameter is around 1 mm as at this diameter the ®bres become a health hazard, if inhaled, as they block the alveolar structure of the lungs. Also related to the ease of converting the ®bres into preforms is the desire for a strain to failure of around 1% and as a Young's modulus of 200 GPa or more is required, this imposes a room temperature strength of more than 2 GPa. Competing materials are usually dense so that a speci®c gravity of less than 5 would be desirable. The ®bres are destined to be used at high temperature and in air so that long term chemical, microstructural and mechanical stability up to and preferably above, 1500C is required. This means that the structure of the ceramic ®bre should not evolve and it should exhibit low creep rates no greater than those of nickel based alloys. Lastly, low reactivity with the matrix is required if the crack stopping process which is the basis of CMC tenacity is to be achieved. 2. Silicon carbide ®bres from organic precursors 2.1. SiC from an oxygen cured precursor route The work of Yajima and his colleagues in Japan was ®rst published in the mid-1970s.4 The Nicalon and Tyranno ®bres produced, respectively, by Nippon Carbon and Ube Industries are the commercial results of this work. These ®bres are produced by the conversion of, respectively polycarbosilane (PCS) and polytitanocarbosilane (PTC) precursor ®bres which contain cycles of six atoms arranged in a similar manner to the diamond structure of b-SiC. The molecular weight of this polycarbosilane is low, around 1500, which makes drawing of the ®bre dicult. The addition of around 2% wt. of titanium, achieved by the grafting of titanium alkoxide between the PCS chains in the Ube precursor, increases molecular weight slightly, which may help with drawing and also may slightly increase thermal resistance by the creation of Ti±C bonds at high temperatures. In these polymers methyl groups (±CH3) in the polymer are included as side groups to the±(Si±C)n± main chain so that during pyrolysis hydrogen is produced, leaving a residue of free carbon. The production of the ®rst generations of SiC based ®bres involved subjecting the precursor ®bres to heating in air at around 200C to produce cross-linking of the structures. This oxidation makes the ®bres infusible but has the drawback of introducing oxygen into the structure which remains after pyrolysis. The ceramic ®bres are obtained by a controlled increase in temperature in an inert atmosphere up to 1200C. The properties and composition of these ®bres are shown in Table 1. The ®bres obtained by this route have a glassy appearance when observed in SEM, as can be seen from Fig. 1, however a closer examination reveals that they contain a majority of b-SiC, of around 2 nm but also signi®cant amounts of free carbon of less than 1nm and 2250 A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20 (2000) 2249±2260

A.R. Bunsell, M.-H. Berger Journal of the European Ceramic Society 20(2000)2249-2260 Table I Properties and compositions of silicon based fibres Fibre type Manufacturer Trade mark Composition(wt % Diameter Density Strength Strain to Youn (um) (g/cm3)(GPa) failure(%)modulus(GPa) Si-C based Nippon Carbide Nicalon NLM 202 56.6 Si: 31.7 C: 11.70 2.55201.05 Nippon Carbide Hi-Nicalon 624Si:37.1C:0.50 Ube Industries Tyranno Lox-M 54.0 Si: 31.6 C: 12.40: 2.0 Ti 8.5 2.3 Ube Industries Tyranno Lox-E 54.8 Si: 37.5 C: 5.80: 1.9 Ti I1 Near stoichiometric Nippon Carbon Hi-Nicalon S SiC+o+C 13 Ube Industries Tyranno SA SiC+C+o+Al 10 Dow Corning Sylramic SiC+TiB2+C+O 0.75 390 excess silicon combined with oxygen and carbon as an matrix material and with a high fibre volume fraction intergranular phase. 7. Their strengths and Young's mod- under the name of Tyranno Hex. 12 Bundles of fibres, uli show little change up to 1000C. Above this tempera which have been pre-oxidised to give them a thin surface ture, in air, both these properties show a slight decrease up layer of silica, are hot-pressed leading to a dense hex to 1400C. Titanium carbide grains are seen in the Tyr- agonal packing of the fibres, the cavities being filled by anno fibres from 1200 C 0 Between 1400 and 1500 C the silica and TiC particles. The strength of Tyranno Hex intergranular phases in both Nicalon and Tyranno fibres measured in bending tests has been reported to be stable begin to decompose, carbon and silicon monoxides are up 1400C in air evacuated and a rapid grain growth of the silicon carbide grains is observed. The densities of the fibres decrease 2. 2. Electron cured precursor filaments rapidly and the tensile properties exhibit a dramatic fall. When a load is applied to the fibres, it is found that a A later generation of Nicalon and Tyranno fibres has creep threshold stress exists above which creep occurs. been produced by cross-linking the precursors by elec he fibres are seen to creep above 1000oC and no stress tron irradiation so avoiding the introduction, at th enhanced grain th is observed after deformation stage, of oxygen. These fibres are known as Hi-Nicalon, 5 Creep is due to the presence of the oxygen rich inter- which contains 0.5% wt oxygen and Tyranno LOX-El3 granular phase. The fibres made by the above process are which contains approximately 5% wt. oxygen. 0, 14The the Nicalon NL-200 and Tyranno LOX-M fibres. The higher value of oxygen in the LOX-E fibre is due to the Lox M fibres have been successfully used for the forma introduction of titanium alkoxides for the fabrication of tion of composite material, without the infiltration of a the PtC. The decrease in oxygen content in the hi- Nicalon compared to the NL-200 fibres has resulted in an increase in the size of the Sic grains and a better organisation of the free carbon. This can be see in Fig. 2. The size of the Sic grains is 5 to10 nm Carbon aggre- gates appear by the stacking of four distorted layers over a length of 2 nm on average. A significant part of the Fig. 2. Lattice fringe image showing regions of lower order between Fig 1. Fracture morphology of a first generation Nicalon fibre, hav- ng a diameter of 15 um and revealing a glassy appearance crystallised e sed lay.s of around 10 nm and free carbon in the form of several distorted layers with lengths of around between 2 and 5 nm

excess silicon combined with oxygen and carbon as an intergranular phase.7,8 Their strengths and Young's moduli show little change up to 1000C.9 Above this temperature, in air, both these properties show a slight decrease up to 1400C. Titanium carbide grains are seen in the Tyranno ®bres from 1200C.10 Between 1400 and 1500C the intergranular phases in both Nicalon and Tyranno ®bres begin to decompose, carbon and silicon monoxides are evacuated and a rapid grain growth of the silicon carbide grains is observed. The densities of the ®bres decrease rapidly and the tensile properties exhibit a dramatic fall. When a load is applied to the ®bres, it is found that a creep threshold stress exists above which creep occurs.11 The ®bres are seen to creep above 1000C and no stress enhanced grain growth is observed after deformation. Creep is due to the presence of the oxygen rich intergranular phase. The ®bres made by the above process are the Nicalon NL-200 and Tyranno LOX-M ®bres. The Lox-M ®bres have been successfully used for the formation of composite material, without the in®ltration of a matrix material and with a high ®bre volume fraction, under the name of Tyranno Hex.12 Bundles of ®bres, which have been pre-oxidised to give them a thin surface layer of silica, are hot-pressed leading to a dense hexagonal packing of the ®bres, the cavities being ®lled by silica and TiC particles. The strength of Tyranno Hex measured in bending tests has been reported to be stable up 1400C in air. 2.2. Electron cured precursor ®laments A later generation of Nicalon and Tyranno ®bres has been produced by cross-linking the precursors by electron irradiation so avoiding the introduction, at this stage, of oxygen. These ®bres are known as Hi-Nicalon,5 which contains 0.5% wt. oxygen and Tyranno LOX-E13 which contains approximately 5% wt. oxygen.10,14 The higher value of oxygen in the LOX-E ®bre is due to the introduction of titanium alkoxides for the fabrication of the PTC. The decrease in oxygen content in the HiNicalon compared to the NL-200 ®bres has resulted in an increase in the size of the SiC grains and a better organisation of the free carbon. This can be see in Fig. 2. The size of the SiC grains is 5 to10 nm. Carbon aggregates appear by the stacking of four distorted layers over a length of 2 nm on average. A signi®cant part of the Table 1 Properties and compositions of silicon based ®bres Fibre type Manufacturer Trade mark Composition (wt.%) Diameter (mm) Density (g/cm3) Strength (GPa) Strain to failure (%) Young's modulus (GPa) Si±C based Nippon Carbide Nicalon NLM 202 56.6 Si;31.7 C;11.7 O 14 2.55 2.0 1.05 190 Nippon Carbide Hi-Nicalon 62.4 Si;37.1 C;0.5 O 14 2.74 2.6 1.0 263 Ube Industries Tyranno Lox-M 54.0 Si;31.6 C;12.4 O;2.0 Ti 8.5 2.37 2.5 1.4 180 Ube Industries Tyranno Lox-E 54.8 Si;37.5 C;5.8 O;1.9 Ti 11 2.39 2.9 1.45 199 Near stoichiometric Nippon Carbon Hi-Nicalon S SiC+O+C 13 3.0 2.5 0.65 375 SiC Ube Industries Tyranno SA SiC+C+O+Al 10 3.0 2.5 0.75 330 Dow Corning Sylramic SiC+TiB2+C+O 10 3.1 3.0 0.75 390 Fig. 1. Fracture morphology of a ®rst generation Nicalon ®bre, having a diameter of 15 mm and revealing a glassy appearance. Fig. 2. Lattice fringe image showing regions of lower order between crystallised b-SiC grains of around 10 nm and free carbon in the form of several distorted layers with lengths of around between 2 and 5 nm. A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20 (2000) 2249±2260 2251

A.R. Bunsell, M.-H. Berger /Journal of the European Ceramic Sociery 20(2000)2249-2260 Sic is not perfectly crystallised and surrounds the ovoid the oxygen content to be reduced. The resulting fibres, B-Sic grains. Further heat treatment of the fibres at known as Tyranno ZE and which contains 2% wt. of 1450C induces the Sic grain to grow up to a mean size of oxygen, show increased high temperature creep and che- 30 nm, to develop facets and be in contact with adjacent mical stability and resistance to corrosive environments Sic grains, as is shown in Fig 3. Turbostratic carbon has compared to the LOX-E fibre. 5, 16 been seen to grow preferentially parallel to some of these facets and could in some cases form cages around Sic 2.3. Near stoichiometric fibres grains limiting their growth. Significant improvements in the creep resistance are found for the Hi-Nicalon Efforts to reduce the oxygen content by processing in fibre compared to the NL-200 fibre which can further be inert atmospheres and cross linking by radiation have enhanced by a heat treatment so as to increase its crys- produced fibres with very low oxygen contents. These talinity. The LOX-E fibre has a microstructure and fibres are not, however, stoichiometric as they contain sig- creep properties which are comparable to those of the nificant amount of excess free carbon affecting oxidative LOX-M and Nicalon NL200 fibre A comparison of the stability and creep resistance. Near stoichiometric SiC creep behaviour of the LOX-E and Hi-Nicalon fibres is fibres from polymer precursors are produced by the two shown in Fig 4. Despite the electron curing process, the Japanese fibre producers and in the Usa by Dow use of a PtC does not allow the reduction of the oxygen in Corning by the use of higher pyrolysis temperatures the intergranular phase of the ceramic fibre to the extent This leads to larger grain sizes and the development of a seen in the Hi-Nicalon so that, as in the lox-M, grain sintered material growth is impeded below 1400C and creep is enhanced. A Nippon Carbon has obtained a near-stoichiometric more recent polymer, polyzirconocarbosilane(PZt) has fibre, the Hi-Nicalon S, from a PCS cured by electron allowed the titanium to be replaced by zirconium and irradiation and pyrolysed by a modified Hi-Nicalon pro- cess in a closely controlled atmosphere above 1500C. As a result it is claimed by the manufacturer that excess car bon is reduced from C/Si=1.39 for the Hi-Nicalon to 1.05 for the Hi-Nicalon S. The fibre has a diameter of 12 um and Sic grain sizes of between 50 and 100 nm. The microstructure of the type S fibre is shown in Fig. 5. Con- siderable free carbon, which could help pin the structure at high temperature, can be seen between the Sic grains Ube Industries has developed a near stoichiometric fibre made from a polyaluminocarbosilane. The precursor fibre is cured by oxidation, pyrolysed in two stages, first to 1300C, to form an oxygen rich SiC fibre, then up to 1800 C to allow first the outgassing of CO, between 1500 and 1700 C, and sintering. The addition of aluminium as a Fig 3. Growth of faceted B-Sic grains and carbon aggregates parallel to sintering aid allows the degradation of the oxicarbide phase the faces of a Sic grain in the Hi-Nicalon fibre heated at 1400 C for 24h. at high temperature to be controlled and catastrophic grain 1350° LOX-E Hi-Nic 020000400006000080000100000120000140000 200nm Time(s) of the hi-nicalon ig. 4. A comparison of creep curves obtained at 1350C with the metric SiC fibre, revealing Sic grains of 50-100 nm and free turbos. Tyranno LOX-E and Hi-Nicalon fibres tratic carbon at triple points

SiC is not perfectly crystallised and surrounds the ovoid b-SiC grains. Further heat treatment of the ®bres at 1450C induces the SiC grain to grow up to a mean size of 30 nm, to develop facets and be in contact with adjacent SiC grains, as is shown in Fig. 3. Turbostratic carbon has been seen to grow preferentially parallel to some of these facets and could in some cases form cages around SiC grains limiting their growth. Signi®cant improvements in the creep resistance are found for the Hi-Nicalon ®bre compared to the NL-200 ®bre which can further be enhanced by a heat treatment so as to increase its crystalinity. The LOX-E ®bre has a microstructure and creep properties which are comparable to those of the LOX-M and Nicalon NL200 ®bre. A comparison of the creep behaviour of the LOX-E and Hi-Nicalon ®bres is shown in Fig. 4. Despite the electron curing process, the use of a PTC does not allow the reduction of the oxygen in the intergranular phase of the ceramic ®bre to the extent seen in the Hi-Nicalon so that, as in the LOX-M, grain growth is impeded below 1400C and creep is enhanced. A more recent polymer, polyzirconocarbosilane (PZT) has allowed the titanium to be replaced by zirconium and the oxygen content to be reduced. The resulting ®bres, known as Tyranno ZE and which contains 2% wt. of oxygen, show increased high temperature creep and chemical stability and resistance to corrosive environments compared to the LOX-E ®bre.15,16 2.3. Near stoichiometric ®bres Eorts to reduce the oxygen content by processing in inert atmospheres and cross linking by radiation have produced ®bres with very low oxygen contents. These ®bres are not, however, stoichiometric as they contain signi®cant amount of excess free carbon aecting oxidative stability and creep resistance. Near stoichiometric SiC ®bres from polymer precursors are produced by the two Japanese ®bre producers and in the USA by Dow Corning by the use of higher pyrolysis temperatures. This leads to larger grain sizes and the development of a sintered material. Nippon Carbon has obtained a near-stoichiometric ®bre, the Hi-Nicalon S, from a PCS cured by electron irradiation and pyrolysed by a modi®ed Hi-Nicalon process in a closely controlled atmosphere above 1500C.5 As a result it is claimed by the manufacturer that excess carbon is reduced from C/Si=1.39 for the Hi-Nicalon to 1.05 for the Hi-Nicalon S. The ®bre has a diameter of 12 mm and SiC grain sizes of between 50 and 100 nm. The microstructure of the type S ®bre is shown in Fig. 5. Considerable free carbon, which could help pin the structure at high temperature, can be seen between the SiC grains. Ube Industries has developed a near stoichiometric ®bre made from a polyaluminocarbosilane.17 The precursor ®bre is cured by oxidation, pyrolysed in two stages, ®rst to 1300C, to form an oxygen rich SiC ®bre, then up to 1800C to allow ®rst the outgassing of CO, between 1500 and 1700C, and sintering. The addition of aluminium as a sintering aid allows the degradation of the oxicarbide phase at high temperature to be controlled and catastrophic grain Fig. 3. Growth of faceted b-SiC grains and carbon aggregates parallel to the faces of a SiC grain in the Hi-Nicalon ®bre heated at 1400C for 24 h. Fig. 5. The microstructure of the Hi-Nicalon type S, near stoichiometric SiC ®bre, revealing SiC grains of 50±100 nm and free turbostratic carbon at triple points. Fig. 4. A comparison of creep curves obtained at 1350C with the Tyranno LOX-E and Hi-Nicalon ®bres. 2252 A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20 (2000) 2249±2260

A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20(2000)2249-2260 2253 rowth and associated porosity, which occurred with the 1600oC to form a near stoichiometric fibre called SYL- previous oxygen rich fibres, avoided. The precursor fibre RAMIC fibre. 8 Such a fibre has a diameter of 10 um and can then be sintered at high temperature so that the Sic grain sizes ranging from 0. 1 to 0. 2 um with smaller excess carbon and oxygen are lost as volatile species to grains of TiB, and B4C. Fig 8 shows the microstructure yield a polycrystalline, near-stoichiometric, SiC fibre. of the Sylramic fibre which has also be seen to contai This Tyranno SA fibre has a diameter of 10 um and Sic excess carbon grain sizes of about 200 nm. The microstructure of the A comparison of the Youngs moduli of the three fibre is shown in Fig. 6 which also reveals considerable near stoichiometric Sic fibres gives different results excess carbon. Less than 1% wt of Al has been added suggesting that they are not fully dense sic materials as a sintering aid and the manufacturer claims that it The Hi-Nicalon-s fibre has an elastic modulus 375 GPa gives better corrosion resistance compared to other and a density of 3.0 g/cm The Tyranno SA fibre has a metals. A fracture surface of the Tyranno SA fibre is Youngs modulus of 330 GPa and a density of 3.0 g/ shown in Fig. 7 and can be seen to be noticeably more cm The elastic modulus of the Sylramic fibres is 390 granular than the earlier generations of fibres. iPa and its density is 3. 1 g/cm. A comparison of the Dow Corning has produced stoichiometric SiC fibres strengths as a function of temperature of the three near using PTC precursors containing a small amount of tita- stoichiometric SiC fibres is shown in Fig 9. The three nium, similar to the precursors described above for the fibres show much improved creep properties with creep earlier generation Ube fibres. These fibres are cured by rates are of the order of 10-s at 1400C when com oxidation and doped with boron which acts as a sinter pared to the earlier generations of fibres which have ing aid. The precursor fibre is pyrolysed at around rates of 10-7 s-I at the same temperature. The creep Fig. 6. The microstructure of the Tyranno SA, near stoichiometric Fig. 8. The microstructure of the Sylramic, near stoichiometric Sic Sic fibre, revealing Sic grains of 200 nm and free turbostratic carbon fibre, revealing Sic grains of around 200 nm. TiB, grains of 50 nm and free turbostratic carbon at triple points. 2,5 g苏g aHSylramic 一H| Nicalon-● Tyrann 0,0 1000 1500 Fig. 7. Fracture morphology of a Tyranno SA fibre revealing its Fig 9. Comparison of mean failure stress for near stoichiometric sili- granular structure con carbide fibres

growth and associated porosity, which occurred with the previous oxygen rich ®bres, avoided. The precursor ®bre can then be sintered at high temperature so that the excess carbon and oxygen are lost as volatile species to yield a polycrystalline, near-stoichiometric, SiC ®bre. This Tyranno SA ®bre has a diameter of 10 mm and SiC grain sizes of about 200 nm. The microstructure of the ®bre is shown in Fig. 6 which also reveals considerable excess carbon. Less than 1% wt. of Al has been added as a sintering aid and the manufacturer claims that it gives better corrosion resistance compared to other metals. A fracture surface of the Tyranno SA ®bre is shown in Fig. 7 and can be seen to be noticeably more granular than the earlier generations of ®bres. Dow Corning has produced stoichiometric SiC ®bres using PTC precursors containing a small amount of titanium, similar to the precursors described above for the earlier generation Ube ®bres. These ®bres are cured by oxidation and doped with boron which acts as a sintering aid. The precursor ®bre is pyrolysed at around 1600C to form a near stoichiometric ®bre called SYLRAMIC ®bre.18 Such a ®bre has a diameter of 10 mm and SiC grain sizes ranging from 0.1 to 0.2 mm with smaller grains of TiB2 and B4C. Fig. 8 shows the microstructure of the Sylramic ®bre which has also be seen to contain excess carbon. A comparison of the Young's moduli of the three near stoichiometric SiC ®bres gives dierent results suggesting that they are not fully dense SiC materials. The Hi-Nicalon-S ®bre has an elastic modulus 375 GPa and a density of 3.0 g/cm3 . The Tyranno SA ®bre has a Young's modulus of 330 GPa and a density of 3.0 g/ cm3 . The elastic modulus of the Sylramic ®bres is 390 GPa and its density is 3.1 g/cm3 . A comparison of the strengths as a function of temperature of the three near stoichiometric SiC ®bres is shown in Fig. 9. The three ®bres show much improved creep properties with creep rates are of the order of 10ÿ8 sÿ1 at 1400C when compared to the earlier generations of ®bres which have rates of 10ÿ7 sÿ1 at the same temperature. The creep Fig. 7. Fracture morphology of a Tyranno SA ®bre revealing its granular structure. Fig. 8. The microstructure of the Sylramic, near stoichiometric SiC ®bre, revealing SiC grains of around 200 nm, TiB2 grains of 50 nm and free turbostratic carbon at triple points. Fig. 9. Comparison of mean failure stress for near stoichiometric silicon carbide ®bres. Fig. 6. The microstructure of the Tyranno SA, near stoichiometric SiC ®bre, revealing SiC grains of 200 nm and free turbostratic carbon at triple points. A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20 (2000) 2249±2260 2253

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