E≈S Journal of the European Ceramic Society 20(2000)2249-2260 fine diameter ceramic fibres A.R. Bunsell M.-H. Berger Ecole des Mines de Paris, Centre des Materiaux, BP87, 91003 Ewry Cedex, france Received 28 January 2000: accepted 12 March 2000 Two families of small diameter ceramic fibres exist. The oxide fibres, 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 fibres made possible the development of ceramic matrix composites. Improved understanding of the mechanisms which control the high tem- perature behaviour of these latter fibres has led to their evolution towards a near stoichiometric composition which results in strength retention at higher temperatures and lower creep rates. The Sic fibres will however be ultimately limited by oxidation so that there is an increasing interest in complex two phase oxide fibres composed of a-alumina and mullite as candidates for the reinforcement of ceramic matrices for use at very high temperatures. These fibres show low creep rates, comparable to the Sic based fibres but are revealed to be sensitive to alkaline contamination. C 2000 Elsevier Science Ltd. All rights reserved Keywords: Al2O3; Creep: Fibres; Microstructure: SiC Contents ntroduction Silicon caride fibres from urgd pr cursor route. 2.1. SiC 2250 2.2. Electron cured precursor filaments 2251 23. Near stoichiometric fibres 3. Oxide fibres 3. 1. Alumina silica fibres 3.1.1. The Saffil fibre 3. 1.2. The Altex fibre 2255 3.1. 3. The Nextel 312-440 fibres 2255 2. Alpha-alumina fibres 3. 2.1. Fully dense a-alumina fibre 3. 2.2. Porous a-alumina fibres 3.3. Alpha-alumina fibres containing a second phase 2257 2257 3.3. 1. Zirconia-reinforced alumina fibres 2257 3.3.2. The Nextel 720 alumina mullite fibre 4. Conclusions 2259 References 0955-2219/00/S. see front matter C 2000 Elsevier Science Ltd. All rights reserved PII:S0955-2219(00)00090-X
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
A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20(2000)2249-2260 Creep at 1300C began to be produced in the early 1970s for high tem- Hot Zone: 25 mm perature insulation. Silica was added to maintain the transitional forms of alumina and inhibit a-alumina Tyranno SA 0.32 GPa grain growth. Such fibres show poor creep resistance however. Later in that decade pure and stiffer a-alumina fibres were produced for the reinforcement of aluminium to obtain light weight composites with a high Youngs modulus. The brittleness of such fibres posed difficulties ns in Hi-Nicalon Type S0.67 GPa fibre diameter and by the additions of second phases. The latest fibres in this family of reinforcements consist of a- alumina together with mullite which show remarkable high temperature properties and good creep resistance if contamination from alkaline elements can be avoided Fig. 10. Typical creep curves obtained at 1300C for near stoichio- metric silicon carbide fibres 3. 1. Alumina silica fibres behaviours of the three fibres are shown in Fig. 10. It The difficulties in producing pure alumina fibres, can be seen that the Nicalon type S fibre shows lower which are the control of porosity and grain growth of the creep rates than the other two fibres and this fibre is also alpha phase, as well as the brittleness of these fibres can seen to maintain its room temperature strength up to be overcome by the inclusion of silica in the structure. 1400C. The use of the electron curing process for the PCs The microstructures of these fibres depend on the highest precursor is clearly of benefit although it imposes a cost temperature the fibres have seen during the ceramisation penalty. The sintering aids used in the other two fibres Very small grains of n-y- or 8-alumina in an amor- are seen to increase creep rates by increasing diffusion phous silica continuum are obtained with temperatures rates within the fibres at high temperatures. The char- below 1000-11000C. Above this range of temperatures a acteristics of the three near stoichiometric SiC fibres are rapid growth of a-alumina porous grains is observed in hown in Table 1 pure alumina fibres. 9 The introduction of silica allows The emerging generation of stoichiometric SiC fibres this transformation to be limited, as it reacts with alumina represents a solution to the instability of earlier fibres to form mullite(3Al2O3: 2SiO2). The presence of mullite however the accompanying increase in Youngs modulus at grain boundaries controls the growth of the a-alu and a slight loss in strength due to larger grain sizes mina which has not been consumed by the reaction. 20 leads to fibres which become more difficult to handle The Youngs moduli of these fibres are lower com and convert into structures. This difficulty may be pared to that of pure alumina fibres, and such fibres are overcome by transforming partially converted Si-C-o produced at a lower cost. This, added to easier handling fibres. into the woven or other form of fibre arrangement due to their lower stiffness makes them attractive for ollowed by pyrolysis and sintering to convert them into thermal insulation applications, in the absence of sig a stoichiometric dense form. The fibre structure could nificant load, in the form of consolidated felts or bricks, up then be infiltrated to form the matrix, giving an optimised to at least 1500 C. Such fibres are also used to reinforce ceramic matrix composite. However even stoichiometric aluminium alloys in the temperature range of 300- 1200C resulting in the formation of a silica surface o their lower Youngs moduli layer. CMCs rely for their tenacity on their ability to accumulate cracks which are however paths for oxida- 3. 1.1. The Saffil fibre tion of the fibres. Silica layers created on the fibres in The Saffil fibre20,21 is a discontinuous fibre with a the vicinity of cracks would fuse the fibres to the matrix diameter of 3 um and was introduced by ICI in 1972. It seriously reducing fibre pull-out and the absorption of consists of 8-alumina and 4% of silica and is produced failure energy of the composite. For this reason this by the blow extrusion of partially hydrolysed solutions family of fibres are likely to be limited to a maximum of some aluminium salts with a small amount of silica. emperature of I400°C The fibre contains mainly small 8-alumina grains of around 50 nm but also some a-alumina grains of 100 The widest use of the Saffil type fibre in composit 3. Oxide fibres is in the form of a mat which can be shaped to the form desired and then infiltrated with molten metal, usually Oxide fibres are clearly of interest for high tempera- aluminium alloy. It is the most successful fibre reinfor ture use in air and alumina combined with silica fibres cement for metal matrix composites
behaviours of the three ®bres are shown in Fig. 10. It can be seen that the Nicalon type S ®bre shows lower creep rates than the other two ®bres and this ®bre is also seen to maintain its room temperature strength up to 1400C. The use of the electron curing process for the PCS precursor is clearly of bene®t although it imposes a cost penalty. The sintering aids used in the other two ®bres are seen to increase creep rates by increasing diusion rates within the ®bres at high temperatures. The characteristics of the three near stoichiometric SiC ®bres are shown in Table 1. The emerging generation of stoichiometric SiC ®bres represents a solution to the instability of earlier ®bres however the accompanying increase in Young's modulus and a slight loss in strength due to larger grain sizes leads to ®bres which become more dicult to handle and convert into structures. This diculty may be overcome by transforming partially converted Si±C±O ®bres, into the woven or other form of ®bre arrangement followed by pyrolysis and sintering to convert them into a stoichiometric dense form. The ®bre structure could then be in®ltrated to form the matrix, giving an optimised ceramic matrix composite. However even stoichiometric silicon carbide ®bres will suer from oxidation from 1200C resulting in the formation of a silica surface layer. CMCs rely for their tenacity on their ability to accumulate cracks which are however paths for oxidation of the ®bres. Silica layers created on the ®bres in the vicinity of cracks would fuse the ®bres to the matrix seriously reducing ®bre pull-out and the absorption of failure energy of the composite. For this reason this family of ®bres are likely to be limited to a maximum temperature of 1400C. 3. Oxide ®bres Oxide ®bres are clearly of interest for high temperature use in air and alumina combined with silica ®bres began to be produced in the early 1970s for high temperature insulation. Silica was added to maintain the transitional forms of alumina and inhibit a-alumina grain growth. Such ®bres show poor creep resistance however. Later in that decade pure and stier a-alumina ®bres were produced for the reinforcement of aluminium to obtain light weight composites with a high Young's modulus. The brittleness of such ®bres posed diculties of handling. This has been ameliorated by reductions in ®bre diameter and by the additions of second phases. The latest ®bres in this family of reinforcements consist of aalumina together with mullite which show remarkable high temperature properties and good creep resistance if contamination from alkaline elements can be avoided. 3.1. Alumina silica ®bres The diculties in producing pure alumina ®bres, which are the control of porosity and grain growth of the alpha phase, as well as the brittleness of these ®bres can be overcome by the inclusion of silica in the structure. The microstructures of these ®bres depend on the highest temperature the ®bres have seen during the ceramisation. Very small grains of Z±g± or d±alumina in an amorphous silica continuum are obtained with temperatures below 1000±1100C. Above this range of temperatures a rapid growth of a-alumina porous grains is observed in pure alumina ®bres.19 The introduction of silica allows this transformation to be limited, as it reacts with alumina to form mullite (3Al2O3:2SiO2). The presence of mullite at grain boundaries controls the growth of the a-alumina which has not been consumed by the reaction.20 The Young's moduli of these ®bres are lower compared to that of pure alumina ®bres, and such ®bres are produced at a lower cost. This, added to easier handling due to their lower stiness, makes them attractive for thermal insulation applications, in the absence of signi®cant load, in the form of consolidated felts or bricks, up to at least 1500C. Such ®bres are also used to reinforce aluminium alloys in the temperature range of 300± 350C. Continuous ®bres of this type can be woven due to their lower Young's moduli. 3.1.1. The Sal ®bre The Sal ®bre20,21 is a discontinuous ®bre with a diameter of 3 mm and was introduced by ICI in 1972. It consists of d-alumina and 4% of silica and is produced by the blow extrusion of partially hydrolysed solutions of some aluminium salts with a small amount of silica. The ®bre contains mainly small d-alumina grains of around 50 nm but also some a-alumina grains of 100 nm. The widest use of the Sal type ®bre in composites is in the form of a mat which can be shaped to the form desired and then in®ltrated with molten metal, usually aluminium alloy. It is the most successful ®bre reinforcement for metal matrix composites. Fig. 10. Typical creep curves obtained at 1300C for near stoichiometric silicon carbide ®bres. 2254 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 2255 Heat treatments of the fibre above 1000c induce the heated in air to 760oc. a treatment which carbonises the 8-alumina to progressively change into a-alumina. After organic groups to give a ceramic fibre composed of 85% 100 h at 1200C or I h at 1400C acicular a-alumina alumina and 15% amorphous silica. The fibre is then grains can be seen on the surface of the fibre and mullite heated to 970 C and its microstructure, as shown in Fig is detected. After 2 h at 1400 C the transformation is l1, consists of small y-alumina grains of a few tens of complete and the equilibrium mullite concentration of nanometres intimately dispersed in an amorphous silica 13% is established. Shrinkage of the fibre and hence the phase. The fibre has a room temperature tensile strength dimensional stability of bricks is controlled up to at of 1.8 GPa and an elastic modulus of 210 GPa as shown astl500°C in Table 2. Subsequent heat treatment produces mullite above 1 100 c. at 1400oc the conversion to mullite is 3. 1.2. The Altex fibre completed and the fibre is composed of 55% mullite and The Altex fibre is produced by Sumitomo Chemicals.22 45% a-alumina by weight. The presence of silica in the It is produced in two forms with diameters of 9 and 17 Altex fibres does not reduce their strength at lower tem- um. The fibre is obtained by the chemical conversion of a peratures compared with pure alumina fibres, however a polymeric precursor fibre, produced from a sol and then lower activation energy is required for the creep of the fibre. 2 At 1200C the continuum of silica allows new tonian creep and the creep rates are higher than those observed with fibres composed solely of a-alumina and which are described below. The Altex fibre is produced as a reinforcement for aluminium alloys The 3M corporation produces a range of ceramic fibres under the general name of Nextel. The Nextel 312 and 440 fibres are produced by a sol-gel process. 24 They are composed of 3 mols of alumina for 2 mols of silica with various amounts of boria to restrict crystal growth. Solvent loss during the rapid drying of the fila ment produces oval cross sections with the major diameter up to twice the minor diameter. They are available with average calculated equivalent diameters of 8-9 and 10- 12 um. Their mechanical properties are reported in Table 2 The Nextel 312 fibre, first appeared in 1974, is com- posed of 62% wt. Al2O3, 24% SiO2 and 14% B2O3 and appears mainly amorphous from TEM observation 100nm although small crystals of aluminium borate have been reported. It has the lowest production cost of the three Fig. Il. TEM dark field image of the Altex fibre showing y-alumina fibres and is widely used but has a mediocre thermal stability as boria compounds volatilise from 1000oC Table 2 Properties and compositions of alumina based fibres Fibre type Manufacturer Trade Composition(wt % ensity Strength Strain to Youngs a-Al2O3 fibres Du pont de Nemours FP 9.9 Al,O3 Mitsui minin Almax 99.9 AlO3 61099%Al2O;0.2-0.3;SiO2 0.40.7:Fe,O Alumina-silica fibres ICI Saffil 0.67 Sumitomo Altex 85 Al2O3: 15 SiO2 62Al2O324SiO2;14B2O310-12or892.7 440 A12O3;28SiO2;2B2O3 1.l1 720 85A1-O3: 15 SiO2 0.81
Heat treatments of the ®bre above 1000C induce the d±alumina to progressively change into a-alumina. After 100 h at 1200C or 1 h at 1400C acicular a-alumina grains can be seen on the surface of the ®bre and mullite is detected. After 2 h at 1400C the transformation is complete and the equilibrium mullite concentration of 13% is established. Shrinkage of the ®bre and hence the dimensional stability of bricks is controlled up to at least 1500C. 3.1.2. The Altex ®bre The Altex ®bre is produced by Sumitomo Chemicals.22 It is produced in two forms with diameters of 9 and 17 mm. The ®bre is obtained by the chemical conversion of a polymeric precursor ®bre, produced from a sol and then heated in air to 760C, a treatment which carbonises the organic groups to give a ceramic ®bre composed of 85% alumina and 15% amorphous silica. The ®bre is then heated to 970C and its microstructure, as shown in Fig. 11, consists of small g-alumina grains of a few tens of nanometres intimately dispersed in an amorphous silica phase. The ®bre has a room temperature tensile strength of 1.8 GPa and an elastic modulus of 210 GPa as shown in Table 2. Subsequent heat treatment produces mullite above 1100C. At 1400C the conversion to mullite is completed and the ®bre is composed of 55% mullite and 45% a-alumina by weight. The presence of silica in the Altex ®bres does not reduce their strength at lower temperatures compared with pure alumina ®bres, however a lower activation energy is required for the creep of the ®bre.23 At 1200C the continuum of silica allows Newtonian creep and the creep rates are higher than those observed with ®bres composed solely of a-alumina and which are described below. The Altex ®bre is produced as a reinforcement for aluminium alloys. 3.1.3. The Nextel 312-440 ®bres The 3M corporation produces a range of ceramic ®bres under the general name of Nextel. The Nextel 312 and 440 ®bres are produced by a sol±gel process.24,25 They are composed of 3 mols of alumina for 2 mols of silica with various amounts of boria to restrict crystal growth. Solvent loss during the rapid drying of the ®lament produces oval cross sections with the major diameter up to twice the minor diameter. They are available with average calculated equivalent diameters of 8±9 and 10± 12 mm. Their mechanical properties are reported in Table 2. The Nextel 312 ®bre, ®rst appeared in 1974, is composed of 62% wt. Al2O3, 24% SiO2 and 14% B2O3 and appears mainly amorphous from TEM observation although small crystals of aluminium borate have been reported. It has the lowest production cost of the three ®bres and is widely used but has a mediocre thermal stability as boria compounds volatilise from 1000C Table 2 Properties and compositions of alumina based ®bres Fibre type Manufacturer Trade mark Composition (wt.%) Diameter (mm) Density (g/cm3 ) Strength (GPa) Strain to failure (%) Young's modulus (GPa) a-Al2O3 ®bres Du Pont de Nemours FP 99.9 Al2O3 20 3.92 1.2 0.29 414 Mitsui Mining Almax 99.9 Al2O3 10 3.6 1.02 0.3 344 3M 610 99% Al2O3;0.2±0.3;SiO2 0.4±0.7;Fe2O3 10±12 3.75 2.6 0.7 370 Alumina±silica ®bres ICI Sal 95 Al2O3;5 SiO2 1±5 3.2 2 0.67 300 Sumitomo Altex 85 Al2O3;15 SiO2 15 3.2 1.8 0.8 210 3M 312 62 Al2O3;24 SiO2;14 B2O3 10±12 or 8±9 2.7 1.7 1.12 152 3M 440 70 Al2O3;28 SiO2;2 B2O3 10±12 3.05 2.1 1.11 190 3M 720 85 Al2O3;15 SiO2 12 3.4 2.1 0.81 260 Fig. 11. TEM dark ®eld image of the Altex ®bre showing g-alumina grains. A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20 (2000) 2249±2260 2255
A.R. Bunsell, M.-H. Berger Journal of the European Ceramic Sociery 20(2000)2249-2260 inducing some severe shrinkage above 1200C. To improve the high temperature stability in the Nextel 440 fibre, the amount of boria has been reduced This latter fibre is composed of 70%Al2O3, 28% SiO2 and 2% B2O3 by weight and is formed in the main of small y-alumina in morphous silica. The fibre is of interest for the rein- forcement of aluminium 3.2. Alpha-alumina fibres Alpha-alumina is the most stable and crystalline form of alumina to which all other phases are converted upon heating above around 1000 C. As we have seen above m fibres based on alumina can contain silica as its presence allows the rapid growth of large and porous a-alumina Fig. 12. Fracture morphology of fiber FP revealing its granular grains to be controlled. However the presence of silica structure. reduces the Youngs modulus of the fibre and reduces their creep strength. High creep resistance implies the applied stress equal to 17% of the failure stress. There production of almost pure a-alumina fibres however to was no overall preferential direction for the grain growth obtain a fine and dense microstructure is difficult. The but the development of cavities at some triple points was control of grain growth and porosity in the production of noticed due to the pile up of intergranular dislocations at a-alumina fibres is obtained by using a slurry consisting triple points caused by insufficient accommodation of the of a-alumina particles, of strictly controlled granulo- deformation. The external surfaces of the FP fibres bro- metry. in an aqueous solution of aluminium salts these ken in creep at 1300C after large deformations showed alumina particles act as seeds to lower the formation numerous transverse microcracks, which were not temperature and rate of growth of the a-alumina grains. observed for smaller strains. No modification of the The rheology of the slurry is controlled through its water granulometry was observed after heat treatment without content. The precursor filament which is then produced load at 1300C for 24 h but large deformations resulting by dry spinning is pyrolysed to give an a-alumina fibre from tensile and creep tests conducted at 1300 C were observed to induce grain growth. The strain rates from 3. 2.1. Fully dense a-alumina fibres 1000 to 1300oC were seen to be a function of the square of The FP-fibre, manufactured by Du Pont in 1979, was the applied stress and the activation energy was found to the first wholly a-alumina fibre to be produced. It is no be in the range of 550 to 590 kJ mol. The creep longer produced but is a useful reference for a pure a- mechanism of Fibre FP has been described as being based alumina fibre. It was continuous with a diameter of 18 on grain boundary sliding achieved by an intergranular um. This fibre was composed of 99.9% a-alumina and movement of dislocations and accommodated by several had a density of 3.92 g/cm and a polycrystalline micro- interfacial controlled diffusion mechanisms, involving structure with a grain size of 0.5 um, a high Youngs boundary migration and grain growth. The failure of the modulus 410 GPa, a tensile strength of 1.55 GPa at 25 fibre, at high temperature, occurred after a short period mm but a strain to failure of only 0.4%.26 This brittle- of damage by the growth of transverse intergranular ness together with its diameter made it unsuitable for microcracks from the cavities, the coalescence of which weaving and although showing initial success as a rein- led to a non- flat failure surface forcement for light alloys, production did not progress This fibre was seen to be chemically stable at high beyond the pilot plant stage. Up to 1000C the Fibre FP temperature in air, however its isotropic fine grained showed linear macroscopic elastic behaviour in tension. microstructure led to easy grain sliding and creep Its granular fracture morphology can be seen in Fig. 12. excluding any application as a reinforcement for ceramic Above 1000oC, the fibre was seen to deform plasticity in structures. tension and the mechanical characteristics decreased Other manufacturers have modified the production apidly At 1300oC, strains in tension increased and could technique to reduce the diameter of the alpha-alumina sometimes reach 15%. Creep was observed from 1000C. fibres that they have produced. This reduction of diameter Very little primary creep was reported, but steady state has an immediate advantage of increasing the flexibility creep was seen followed by tertiary creep. This small and and hence the weaveability of the fibres continuous increase of the strain just before failure A continuous a-alumina fibre, with a diameter of 10 um, indicated an accumulation of damage in the fibre that was introduced by 3M in the early 1990s with the trade- preceded failure Grain growth of 40% was revealed for name of Nextel 610 fibre. 7It is composed of around 99% a fibre which failed with a strain of 30%, with an Imina although a more detailed chemical analysis
inducing some severe shrinkage above 1200C. To improve the high temperature stability in the Nextel 440 ®bre, the amount of boria has been reduced. This latter ®bre is composed of 70% Al2O3, 28% SiO2 and 2% B2O3 by weight and is formed in the main of small g-alumina in amorphous silica. The ®bre is of interest for the reinforcement of aluminium. 3.2. Alpha-alumina ®bres Alpha-alumina is the most stable and crystalline form of alumina to which all other phases are converted upon heating above around 1000C. As we have seen above, ®bres based on alumina can contain silica as its presence allows the rapid growth of large and porous a-alumina grains to be controlled. However the presence of silica reduces the Young's modulus of the ®bre and reduces their creep strength. High creep resistance implies the production of almost pure a-alumina ®bres however to obtain a ®ne and dense microstructure is dicult. The control of grain growth and porosity in the production of a-alumina ®bres is obtained by using a slurry consisting of a-alumina particles, of strictly controlled granulometry, in an aqueous solution of aluminium salts. These alumina particles act as seeds to lower the formation temperature and rate of growth of the a-alumina grains. The rheology of the slurry is controlled through its water content. The precursor ®lament which is then produced by dry spinning is pyrolysed to give an a-alumina ®bre. 3.2.1. Fully dense -alumina ®bres The FP-®bre, manufactured by Du Pont in 1979, was the ®rst wholly a-alumina ®bre to be produced.3 It is no longer produced but is a useful reference for a pure aalumina ®bre. It was continuous with a diameter of 18 mm. This ®bre was composed of 99.9% a-alumina and had a density of 3.92 g/cm3 and a polycrystalline microstructure with a grain size of 0.5 mm, a high Young's modulus 410 GPa, a tensile strength of 1.55 GPa at 25 mm but a strain to failure of only 0.4%.26 This brittleness together with its diameter made it unsuitable for weaving and although showing initial success as a reinforcement for light alloys, production did not progress beyond the pilot plant stage. Up to 1000C the Fibre FP showed linear macroscopic elastic behaviour in tension. Its granular fracture morphology can be seen in Fig. 12. Above 1000C, the ®bre was seen to deform plasticity in tension and the mechanical characteristics decreased rapidly. At 1300C, strains in tension increased and could sometimes reach 15%. Creep was observed from 1000C. Very little primary creep was reported, but steady state creep was seen followed by tertiary creep. This small and continuous increase of the strain just before failure indicated an accumulation of damage in the ®bre that preceded failure. Grain growth of 40% was revealed for a ®bre which failed with a strain of 30%, with an applied stress equal to 17% of the failure stress. There was no overall preferential direction for the grain growth but the development of cavities at some triple points was noticed due to the pile up of intergranular dislocations at triple points caused by insucient accommodation of the deformation. The external surfaces of the FP ®bres broken in creep at 1300C after large deformations showed numerous transverse microcracks, which were not observed for smaller strains. No modi®cation of the granulometry was observed after heat treatment without load at 1300C for 24 h but large deformations resulting from tensile and creep tests conducted at 1300C were observed to induce grain growth. The strain rates from 1000 to 1300C were seen to be a function of the square of the applied stress and the activation energy was found to be in the range of 550 to 590 kJ molÿ1 . The creep mechanism of Fibre FP has been described as being based on grain boundary sliding achieved by an intergranular movement of dislocations and accommodated by several interfacial controlled diusion mechanisms, involving boundary migration and grain growth. The failure of the ®bre, at high temperature, occurred after a short period of damage, by the growth of transverse intergranular microcracks from the cavities, the coalescence of which led to a non-¯at failure surface. This ®bre was seen to be chemically stable at high temperature in air, however its isotropic ®ne grained microstructure led to easy grain sliding and creep, excluding any application as a reinforcement for ceramic structures. Other manufacturers have modi®ed the production technique to reduce the diameter of the alpha-alumina ®bres that they have produced. This reduction of diameter has an immediate advantage of increasing the ¯exibility and hence the weaveability of the ®bres. A continuous a-alumina ®bre, with a diameter of 10 mm, was introduced by 3M in the early 1990s with the tradename of Nextel 610 ®bre.27 It is composed of around 99% a-alumina although a more detailed chemical analysis Fig. 12. Fracture morphology of ®ber FP revealing its granular structure. 2256 A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20 (2000) 2249±2260
A.R. Munsell M-H. Berger / Journal of the European Ceramic Society 20(2000)2249--2260 gives 1. 15% total impurities including 0.67% Fe2O3 used without elimination of porosity and internal stresses. As a as a nucleating agent and 0.35% Sioz as grain growth consequence, grain growth at 1300C is activated without inhibitor. It is believed that the silica which is introduced an applied load and reaches 40% after 24 h, unlike that does not form a second phase at grain boundaries with the other pure alpha alumina fibres, for which grain although the suggestion of a very thin second phase growth is related to the accommodation of the slip by separating most of the grains has been observed by diffusion transmission electron microscopy. As can be seen from The fibre exhibits linear elastic behaviour at room Fig. 13, the fibre is polycrystalline with a grain size of temperature in tension and brittle failure. The mechan 0. 1 um, 5 times smaller than in Fibre FP. As shown in ical properties of the Almax fibre are shown in Table 2. Table 2. this smaller grain size together with the smaller Its Youngs modulus is lower than that of the Fibre FP, diameter leads to a fibre strength which is twice the because of the greater amount of porosity. The reduction tensile strength measured for Fibre FP of the measured failure stress of the Almax compared to Creep occurs from 900C and strain rates are 2 to 6 those of the FP and the more pronounced intragranular times larger than those of Fibre FP, due to the finer gran- failure mode for this fibre compared to the FP fibre ulometry and possibly to the chemistry of its grain show a weakening of the grains by the intragranular boundaries. A stress exponent of approximately 3 is found porosity. The fibres exhibit linear macroscopic elastic between 1000 and 1200C with an apparent activation behaviour up to 1000C. Above 1000C the mechanical energy of 660 kJ/mol. Failure frequently occurs via the characteristics decrease rapidly, with a more severe drop coalescence of cavities into large cracks over the whole than for Fibre FP. Tensile failure of the Almax fibre at gauge length and the failure surfaces are rougher compared 1250 C has revealed isotropic grain growth up to 55% to those obtained at room temperature, in the same No extended regions of damage could be observed, as manner as was seen with fibre fp local cavitations seen on the fibre fp surface nd necking around heterogeneities sometimes induce Creep occurs from 1000C for the Almax fibres and the fibres show a lower resistance to creep than Fibre FP. Diffusion and grain boundary sliding are facilitated by 3.2. 2. Porous a-alumina fibres the growth of intergranular porosity and much higher An a-alumina fibre which is still commercially avail- strain rates are obtained. This intergranular porosity able was produced first in the early 1990s by Mitsui which appears in the fibre during creep may have been Mining. 28 It is composed of almost pure a-alumina and created by the interception of intragranular pores by the has a diameter of 10 um. The fibre has a lower density boundaries of the growing grains. This intergranular pore of 3.60 g/cm compared to Fibre FP. Like Fibre FP, the th, rapidly induces fail Imax fibre consists of one population of grains of around 0.5 Hm; however, the fibre exhibits a large 3.3. Alpha-alumina fibres containing a second phase amount of intragranular porosity, and associated with numerous intragranular dislocations without any peri- 3. 3.1. Zirconia-reinforced alumina fibres odic arrangement. 26 This indicates rapid grain growth Du Pont further developed the Fibre FP to produce a of a-alumina grains during the fibre fabrication process fibre ca RD-166 which consisted of 80% wt. a alumina with 20% wt. of partially stabilised zirconia. 29 The introduction of tetragonal zirconia resulted in a toughening of the fibre and a decrease of its Youngs modulus allowing some increase in strain to failure. The fibre began to creep at 1100oC, 100 C above the tem- perature creep threshold for Fibre FP and showed lower strain rates, however this advantage, compared to Fibre FP, was lost at 1300C. 0 PRD-166 was not produced commercially as its large diameter prohibited weaving despite the small increase in flexibility 3M has announced the development of a zirconia reinforced u mInd lextel 650. with a small diameter than that of the PRD-166 fibre which may allow it to be woven 100n 3.3.2. The Nextel 720 alumina mullite fibre The Nextel 720, produced by 3M, contains the same Fig 13. TEM image of Nextel 610 fibre revealing a-alumina grains of alumina to silica ratio as in the Altex fibre, that is around 0. I um around 85% wt. A12O3 and 15% wt SiO, 31 The fibre
gives 1.15% total impurities including 0.67% Fe2O3 used as a nucleating agent and 0.35% SiO2 as grain growth inhibitor. It is believed that the silica which is introduced does not form a second phase at grain boundaries although the suggestion of a very thin second phase separating most of the grains has been observed by transmission electron microscopy. As can be seen from Fig. 13, the ®bre is polycrystalline with a grain size of 0.1 mm, 5 times smaller than in Fibre FP. As shown in Table 2. this smaller grain size together with the smaller diameter leads to a ®bre strength which is twice the tensile strength measured for Fibre FP. Creep occurs from 900C and strain rates are 2 to 6 times larger than those of Fibre FP, due to the ®ner granulometry and possibly to the chemistry of its grain boundaries. A stress exponent of approximately 3 is found between 1000 and 1200C with an apparent activation energy of 660 kJ/mol. Failure frequently occurs via the coalescence of cavities into large cracks over the whole gauge length and the failure surfaces are rougher compared to those obtained at room temperature, in the same manner as was seen with Fibre FP. Local cavitations and necking around heterogeneities sometimes induce failure. 3.2.2. Porous -alumina ®bres An a-alumina ®bre which is still commercially available was produced ®rst in the early 1990s by Mitsui Mining.28 It is composed of almost pure a-alumina and has a diameter of 10 mm. The ®bre has a lower density of 3.60 g/cm3 compared to Fibre FP. Like Fibre FP, the Almax ®bre consists of one population of grains of around 0.5 mm; however, the ®bre exhibits a large amount of intragranular porosity, and associated with numerous intragranular dislocations without any periodic arrangement.26 This indicates rapid grain growth of a-alumina grains during the ®bre fabrication process without elimination of porosity and internal stresses. As a consequence, grain growth at 1300C is activated without an applied load and reaches 40% after 24 h, unlike that with the other pure alpha alumina ®bres, for which grain growth is related to the accommodation of the slip by diusion. The ®bre exhibits linear elastic behaviour at room temperature in tension and brittle failure. The mechanical properties of the Almax ®bre are shown in Table 2. Its Young's modulus is lower than that of the Fibre FP, because of the greater amount of porosity. The reduction of the measured failure stress of the Almax compared to those of the FP and the more pronounced intragranular failure mode for this ®bre compared to the FP ®bre show a weakening of the grains by the intragranular porosity. The ®bres exhibit linear macroscopic elastic behaviour up to 1000C. Above 1000C the mechanical characteristics decrease rapidly, with a more severe drop than for Fibre FP. Tensile failure of the Almax ®bre at 1250C has revealed isotropic grain growth up to 55%. No extended regions of damage could be observed, as seen on the Fibre FP surface. Creep occurs from 1000C for the Almax ®bres and the ®bres show a lower resistance to creep than Fibre FP. Diusion and grain boundary sliding are facilitated by the growth of intergranular porosity and much higher strain rates are obtained. This intergranular porosity which appears in the ®bre during creep may have been created by the interception of intragranular pores by the boundaries of the growing grains. This intergranular pore growth, rapidly induces failure. 3.3. Alpha-alumina ®bres containing a second phase 3.3.1. Zirconia-reinforced alumina ®bres Du Pont further developed the Fibre FP to produce a ®bre called PRD-166 which consisted of 80% wt. aalumina with 20% wt. of partially stabilised zirconia.29 The introduction of tetragonal zirconia resulted in a toughening of the ®bre and a decrease of its Young's modulus allowing some increase in strain to failure. The ®bre began to creep at 1100C, 100C above the temperature creep threshold for Fibre FP and showed lower strain rates, however this advantage, compared to Fibre FP, was lost at 1300C.26 PRD-166 was not produced commercially as its large diameter prohibited weaving despite the small increase in ¯exibility. 3M has announced the development of a zirconia reinforced alumina ®bre, Nextel 650, with a smaller diameter than that of the PRD-166 ®bre which may allow it to be woven.30 3.3.2. The Nextel 720 alumina mullite ®bre The Nextel 720, produced by 3M, contains the same alumina to silica ratio as in the Altex ®bre, that is around 85% wt. Al2O3 and 15% wt. SiO2. 31 The ®bre Fig. 13. TEM image of Nextel 610 ®bre revealing a-alumina grains of around 0.1 mm. A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20 (2000) 2249±2260 2257
58 A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20(2000)2249-2260 has a circular cross section and a diameter of 12 um. fibres, at 1200C and oxide fibre so far pro- The sol-gel route and higher processing temperatures duced commercially. Results from creep tests conducted have induced the growth of alumina rich mullite, com- at 1400C are shown in Fig. 17 and give a creep rate of posed of 2 mols of alumina to one of silica(2: 1 mullite) the order of 10-6 s-. This improved creep resistance is and alpha alumina. Unlike other alumina-silica fibres attributed to the particular microstructure of the fibre the Nextel 720 fibre is composed of mosaic grains of which is composed of aggregates of mullite, which is about 0.5 u with wavy contours, consisting of several known as possessing very good creep properties, rather slightly mutually misoriented mullite grains in which are than the two-phase nature of the structure embedded a-alumina grains some of which are elon gated, as shown in Fig. 14. At room temperature its Youngs modulus and tensile strength at 25 mm are 260 and 2.1 GPa, respectively, as shown in Table 2. Post heat treatment leads to an enrichment of a-alumina in the fibre as mullite rejects alumina to evolve towards a 3: 2 equilibrium composition. Grain growth occurs from 1300C and at 1400 C the wavy interfaces are replaced by straight boundaries, as shown in Fig. 15. The fibre has been shown to be sensitive to contamination by alkalines which are thought to create a silicate phase at the fibre surface. Such phases can have melting points lower than 1000C, and if formed, allow rapid diffusion of elements and grain growth so that strength has been shown to depend on the test conditions. In the presend of such contaminants strength falls at temperatures above 1000oC due to crack initiation at large a-alumina platelets which develop at the surface, as Fig. 16 reveals If contamination is avoided and creep experiments are carried out Dorn plots indicate that 2 orders of magni tude exist between the strain rates of the Nextel 720 500nn Fig. 15. Microstructure of Nextel 720 fibre after heat treatment at 1400 C for 4 days 500nm 1200C. Failure is initiated by the growth of large a-alumina graln e %e Fig. 14. Microstructure of Nextel 720 fibre exhibiting mullite aggre- Fig. 16. Fracture morphology of Nextel 720 fibre failed in tension a es enclosing elongated a-alumina grains
has a circular cross section and a diameter of 12 mm. The sol±gel route and higher processing temperatures have induced the growth of alumina rich mullite, composed of 2 mols of alumina to one of silica (2:1 mullite) and alpha alumina. Unlike other alumina±silica ®bres the Nextel 720 ®bre is composed of mosaic grains of about 0.5 m with wavy contours, consisting of several slightly mutually misoriented mullite grains in which are embedded a-alumina grains some of which are elongated, as shown in Fig. 14.32 At room temperature its Young's modulus and tensile strength at 25 mm are 260 and 2.1 GPa, respectively, as shown in Table 2. Post heat treatment leads to an enrichment of a-alumina in the ®bre as mullite rejects alumina to evolve towards a 3:2 equilibrium composition. Grain growth occurs from 1300C and at 1400C the wavy interfaces are replaced by straight boundaries, as shown in Fig. 15. The ®bre has been shown to be sensitive to contamination by alkalines which are thought to create a silicate phase at the ®bre surface. Such phases can have melting points lower than 1000C, and if formed, allow rapid diusion of elements and grain growth so that strength has been shown to depend on the test conditions. In the presence of such contaminants strength falls at temperatures above 1000C due to crack initiation at large a-alumina platelets which develop at the surface, as Fig. 16 reveals. If contamination is avoided and creep experiments are carried out Dorn plots indicate that 2 orders of magnitude exist between the strain rates of the Nextel 720 ®bres, at 1200C and any other oxide ®bre so far produced commercially. Results from creep tests conducted at 1400C are shown in Fig. 17 and give a creep rate of the order of 10ÿ6 sÿ1 . This improved creep resistance is attributed to the particular microstructure of the ®bre which is composed of aggregates of mullite, which is known as possessing very good creep properties, rather than the two-phase nature of the structure. Fig. 15. Microstructure of Nextel 720 ®bre after heat treatment at 1400C for 4 days. Fig. 16. Fracture morphology of Nextel 720 ®bre failed in tension at 1200C. Failure is initiated by the growth of large a-alumina grains. Fig. 14. Microstructure of Nextel 720 ®bre exhibiting mullite aggregates enclosing elongated a-alumina grains. 2258 A.R. Bunsell, M.-H. Berger / Journal of the European Ceramic Society 20 (2000) 2249±2260