40TH ANNIVERSARY The existence of porosity in the first generation ceramic If improved fibres were to be made, which could go fibres was proposed so as to account for the outgassing some way to meeting the promise of a pure SiC, it was of the hundreds of volumes of gas per volume of fibre, clear that a more stoichiometric composition had to be which occurs during the transformation from an organic achieved. As mentioned above. it had been seen that the to a mineral structure, which occurs during the manu- presence of oxygen in the fibre resulted in a poorly or- facture of the fibre [25]. It was reasoned that nanometric ganised phase in which the Sic grains were embedded channels must exist in the fibres during the transformation It was clear that the oxygen content had to be consider stage as diffusion through a solid phase would be too slow. ably reduced if improvements were to be made. As the These nanochannels collat further heating above the primary reason that the oxygen was in the fibres was that temperature at which the gas is evolved. A porosity of it was introduced to render the precursor fibres infusible 6.2% was shown to exist experimentally by X-ray scatter- it was clear that other means of cross-linking the polymer ing measurements on Nicalon NLM-202 fibres pyrolized precursors had to be investigated at1400°C[26,27 The result of the studies on the first generation fibres as the conclusion that it was the non-stoichiometric com- 4. Second generation small diameter SiC fibres position of the fibres which were limiting their physical The route adopted to reduce the oxygen content in the characteristics. The presence of the amorphous intergran- second generation of small diameter SiC fibres was to in the fibres after pyrolysis making a Si-o-C phase. The render the PCS precursor fibre infusible. This was possible low fraction of a granular Sic phase accounted for the by using different types of irradiation which could interact Youngs modulus of the fibres being only half that of with the precursor polymers to produce free radicals and bulk SiC. The amorphous phase also explained why the gaseous products by the scission of the chemical bonds fibres began to lose strength and creep at temperatures of Si-CH3, Si-H and C-H. This allowed Si-Si and Si-C around 1000 or 1100oc whereas bulk sic would be bonds to be formed. a number of different types of radi pected to resist to higher temperatures [28]. Heating to ation were investigated by both fibre producers working above 1500.C induced rapid grain growth and outgassing in collaboration with the Japanese Atomic Energy Re- of Sio and Co which came from the oxygen in the amor- search Institute. Gamma irradiation was investigated but phous Si-o-C phase and also the free carbon in the fibre, finally electron radiation in a helium atmosphere was used as described by the following relationship to make the second generation fibres. The cross-linking step was followed by heat treatment at 327 C for a short time to eliminate the remaining free radicals which SiC,O,- SiC()+ Sio(g)+Co(g) apped in the irradiated precursor fibre [35]. An example of how the hydrogen atom, bonded to the silicon atom, as It had been noticed in early studies that the rate of shown in Fig. 2, is removed by electron bombardment, so strength loss of first generation fibres was lower in as to allow direct bonding between the two silicon atoms idizing atmospheres than in an inert argon atmosphere in neighbouring molecules, is shown in Fig. 6[36] [17, 29, 30]. It was concluded that heating the fibres in air Below 550 C cross-linking between main chains dom- produced a silica coating which hindered outgassing of inates and is induced by the dehydrogenation condensa the products of the decomposition of the Si-o-C phase; tion of the Si-H groups. From 550oC to 800C, the side this was further studied by heating fibres made from both chains on the cross-linked polymer begin to decompose the pCS and PtC precursors over a wide range of oxygen and CH4 and H2 are given off producing an inorganic fi partial pressures and in carbon dioxide [31, 32]. These bre. Above 800C and up to 1000C, hydrogen is given authors also showed that the formation of a barrier coat- off, most probably associated with the decomposition ing of Sic laid down by chemical vapour deposition on the fibre surface suppressed gas evolution and slowed the degradation process. They also showed that, what had been presumed in earlier studies was actually the case and that a solid layer of Sio2 formed on the surface of the Nicalon fibres and a Sioz layer with a small amount of TiO2 was formed on the surfaces of the Tyranno fibres It should however be noted that increasing the pressure an argon environment also hindered the onset of decomposition. When the first generation fibres wer heated in Argon at a pressure of 138 MPa the onset of Figure 6 Direct cross-linking of the PCS precursor polymer by irradiation fibre weight loss increased from 1200 to 1500C [34]40TH ANNIVERSARY The existence of porosity in the first generation ceramic fibres was proposed so as to account for the outgassing of the hundreds of volumes of gas per volume of fibre, which occurs during the transformation from an organic to a mineral structure, which occurs during the manu￾facture of the fibre [25]. It was reasoned that nanometric channels must exist in the fibres during the transformation stage as diffusion through a solid phase would be too slow. These nanochannels collapse on further heating above the temperature at which the gas is evolved. A porosity of 6.2% was shown to exist experimentally by X-ray scatter￾ing measurements on Nicalon NLM-202 fibres pyrolized at 1400◦C [26, 27]. The result of the studies on the first generation fibres was the conclusion that it was the non-stoichiometric com￾position of the fibres which were limiting their physical characteristics. The presence of the amorphous intergran￾ular phase was clearly due to the oxygen which remained in the fibres after pyrolysis making a Si–O–C phase. The low fraction of a granular SiC phase accounted for the Young’s modulus of the fibres being only half that of bulk SiC. The amorphous phase also explained why the fibres began to lose strength and creep at temperatures around 1000 or 1100◦C whereas bulk SiC would be ex￾pected to resist to higher temperatures [28]. Heating to above 1500◦C induced rapid grain growth and outgassing of SiO and CO which came from the oxygen in the amor￾phous Si–O–C phase and also the free carbon in the fibre, as described by the following relationship. SiCxOy → SiC (s) + SiO(g) + CO(g) It had been noticed in early studies that the rate of strength loss of first generation fibres was lower in ox￾idizing atmospheres than in an inert argon atmosphere [17, 29, 30]. It was concluded that heating the fibres in air produced a silica coating which hindered outgassing of the products of the decomposition of the Si–O–C phase; this was further studied by heating fibres made from both the PCS and PTC precursors over a wide range of oxygen partial pressures and in carbon dioxide [31, 32]. These authors also showed that the formation of a barrier coat￾ing of SiC laid down by chemical vapour deposition on the fibre surface suppressed gas evolution and slowed the degradation process. They also showed that, what had been presumed in earlier studies was actually the case and that a solid layer of SiO2 formed on the surface of the Nicalon fibres and a SiO2 layer with a small amount of TiO2 was formed on the surfaces of the Tyranno fibres [33]. It should however be noted that increasing the pressure of an argon environment also hindered the onset of decomposition. When the first generation fibres were heated in Argon at a pressure of 138 MPa the onset of fibre weight loss increased from 1200 to 1500◦C [34]. If improved fibres were to be made, which could go some way to meeting the promise of a pure SiC, it was clear that a more stoichiometric composition had to be achieved. As mentioned above, it had been seen that the presence of oxygen in the fibre resulted in a poorly or￾ganised phase in which the SiC grains were embedded. It was clear that the oxygen content had to be consider￾ably reduced if improvements were to be made. As the primary reason that the oxygen was in the fibres was that it was introduced to render the precursor fibres infusible, it was clear that other means of cross-linking the polymer precursors had to be investigated. 4. Second generation small diameter SiC fibres The route adopted to reduce the oxygen content in the second generation of small diameter SiC fibres was to eliminate the oxygen induced cross-linking step used to render the PCS precursor fibre infusible. This was possible by using different types of irradiation which could interact with the precursor polymers to produce free radicals and gaseous products by the scission of the chemical bonds of Si–CH3, Si–H and C–H. This allowed Si–Si and Si–C bonds to be formed. A number of different types of radi￾ation were investigated by both fibre producers working in collaboration with the Japanese Atomic Energy Re￾search Institute. Gamma irradiation was investigated but finally electron radiation in a helium atmosphere was used to make the second generation fibres. The cross-linking step was followed by heat treatment at 327◦C for a short time to eliminate the remaining free radicals which were trapped in the irradiated precursor fibre [35]. An example of how the hydrogen atom, bonded to the silicon atom, as shown in Fig. 2, is removed by electron bombardment, so as to allow direct bonding between the two silicon atoms in neighbouring molecules, is shown in Fig. 6 [36]. Below 550◦C cross-linking between main chains dom￾inates and is induced by the dehydrogenation condensa￾tion of the Si–H groups. From 550◦C to 800◦C, the side chains on the cross-linked polymer begin to decompose and CH4 and H2 are given off producing an inorganic fi- bre. Above 800◦C and up to 1000◦C, hydrogen is given off, most probably associated with the decomposition of Figure 6 Direct cross-linking of the PCS precursor polymer by irradiation curing. 829