CHAPTER 10

SILICON CARBIDE AND OXYCARBIDEFIBERSR. Naslain

The major deficiency of carbon fibers is their sensitivity to oxidation even at relatively lowtemperatures. Although silicon carbide (SiC) fibers are also sensitive to oxidation, theiroxidation starts at higher temperatures and yields a protective silica coating. In an oxidativeenvironment, SiC and Si-C-O fibers are generally more useful than carbon fibers [1-3]. Largediameter silicon carbide fibers are obtained by chemical vapor deposition (Chapter 4). Smalldiameter silicon carbide and oxycarbide fibers are derived from solid polydimethylsilazaneprecursor fibers (this chapter) .10.1 General considerations

SiC isacovalent solid in which silicon istetrahedrally coordinated tocarbon (Sp3 bonds). TheSiC4 tetrahedra are interconnected only by corner sharing. SiC exhibits two simple crystalstructures, a cubic blende type modification or the p (or 3C) form, and a hexagonal wurtzitemodification orthe ex (or 2H) form. SiC also displays numerous polymorphs orpolytypes withhexagonal or rhombohedral unit cells which correspond to complex stacking sequences [4].SiCisrigid and brittle (Table I), and has high thermal and electrical conductitivities. SiC meltsat2500Cwith peritectic decomposition. Finally, the diffusion coefficients in SiC are low evenat relatively high temperatures, a feature which correlates with its resistance to sintering andcreep.

Table I. Silicon carbide fiber properties.

Materials g/~' E cr' IX. P.GPa GPa lO-6oC' n.anSingle crystal SiC whiskers 3.21 480-580 6-20 5.0SiC/C fibers made by CVD 360-390 3.6-4.6PC5-based Hi-Nicalon SiC fibers 3.10 420 2.6 3.1 0.1

Silicon carbide fibers are derived [1-3] [5-6] from a polymer precursor fiber bya process thatissimilar to the carbon fiber process. It starts offwith polydimethlysilane (PDMS) rather thanpolyacrylonitrile. First generation Si-C fibers are oxygen containing fibers [7-13]. Secondgeneration Si-C fibers are oxygen-free. They can be produced in a first generation processwherein the green fiber iseither cured by electron beam irradiation [27-29], orderived from ahigh molecular weight polycarbosilane precursor fiber [30]. These SiC fibers contain free

266 Chapter 10

carbon and possess a fine microstructure that is stable at elevated temperatures. Theirmodulus is low 270GPa) [27] [29-30] and they creep, but the creep rate is lower than that offirst generation Si-C(O) fibers [31-32].Third generation, oxygen-free Si-C fibers can also be produced bypolymer precursor routes.They are quasi-stoichiometric, Le., they have an atomic C/Si ratio close tounity. Free carbonthat is present in E-beam cured PCS fibers can be removed by performing the pyrolysis in adecarburizing atmosphere [33]. A precursor fiber can be used with an atomic C/Si ratio closetounity [34-36]. Boron oraluminum can be introduced into the green fiber to remove excesscarbon and oxygen as CO, wh ile the pyrolytic res idue undergoes sintering with boron oraluminum acting as sintering aids [37-40]. Finally, quasi-stoichiometric SiC fibers can beproduced by extrusion ofa viscous slurry ofa submicrometer SiC powder in suspension in aliquid [41-42]. Quasi-stoichiometric fibers have aYoung's modulus ofabout 420 GPa.The criteria for designing fibers for use in ceramic matrix composites (CMCs) are differentfrom those for designing fibers for use in polymer or metal matrix composites. The keyproperties are thermal stability and mechanical properties at high temperatures [43]. As aconsequence, relatively coarse microstructures are obtained at elevated temperatures,corresponding to somewhat lower failure strengths (-2 GPa), but high thermal stability andcreep resistance are preferable toultrafine microstructures.

Synthesis ofpolydimethylsilane

Autoclave(470C) or

Catalyst(PBDPSO)

Figure 1.Fabrication ofSi-C-O fibers from PCS precursors

10.2 Preparation of SiCO fibers

Si-C-Ofiber

The generic polymer based process that yields oxygen containing silicon carbide fibersconsists of five steps. (1) Polydimethylsilane, or PDMS, is synthesized. (2) PDMS isrearranged into polycarbosilane, or PCS. (3) PCS is melt spun and yields a solid, green, or

Chapter 10 267

(3)

precursor fiber. (4) The solid PCS precursor fiber is rendered infusible byoxygen crosslinkingand (5) the stabilized fiber is pyrolyzed and yields the final functional Si-C-O fiber.Polycarbosilane (PCS) is the most commonly used organosilicon polymer precursor forproducing SiC based fibers. Polycarbosilane is a generic name forpolymers containing Si-Cbonds in their backbones [5) [6) [44).10.2.1 The Yajima processIn the generic process, dimethyldichlorosilane is polymerized and yields polydimethylsilane(POMS), asshown inEquation 1, through alkali metal promoted dehalocoupling ofchlorosilanes. The polymerization of dimethyldichlorosilane is carried out in xylene in anitrogen atmosphere. POMS has a polymeric chain structure with a complex empiricalformula (SiCl940oo2Hs46). Itcannot beused directly as a SiC precursor. Its backbone has a Si-Si chain and itspyrolytic residue at800C isalmost nil [7].n(CH3)] SiCI] +2nM ~[(CH3hSi]" +2nMCI (1)with M=Li, Na

400C(CH3JJ Si-Si(CH3JJ ) (CH3)3 Si-CH]-Si-(CH3hH (2)POMS must therefore be converted into polycarbosiliane (PCS) by the Kumadarearrangement as shown in Equations 2-3 [6). The rearrangement inserts methylene (-CH2-)groups from pendent methyl groups into the Si-Si chain [6] [44). This produces a polymer witha -Si-C-Si- backbone, which prefigures that of silicon carbide, and lateral Si-H bonds, whichhelp to stabilize the green fibers. Insertion of methylene groups into the POMS backboneyields polysilapropylene (PSP), te., SiC2H6, with C/Si at. = 2, but there is no gas evolution.The PDMS/PCS conversion proceeds bya thermal, or catalytic thermal, rearrangement. Theactual mechanism is more complex than assumed in Equation 2; the PCSs are partlycrosslinked and their formulas are different from that ofPSP.

CH] CH] CH]-li--li--- +li--CHl

I) 1 1 ~CH] CH] H

The thermal rearrangement of POMS (Equation 4) was initially performed at 470C underatmospheric pressure in nitrogen (PC-N-470 in Table II). However, under such conditions,the reaction is slow. The rearrangement step can be shortened byone order of magnitude ifthe reaction is performed in an autoclave, the final internal pressure being about 10 MPaowing to the formation of gaseous species (PC-A-470 in Table II). The resulting product isdissolved in n-hexane and the solution is filtered. After removal of the solvent, the residue isvacuum distilled yielding a yellowish brown PCS. The chemical composition of PC-A-470,i.e.49.99 wt.%Si; 37.88 wt.%C; 0.99 wt.%0 and 6.60 wt.%H, corresponds toSiC17700 03H370[7). Such PCSs are solids at room temperature with high softening points (Ts ~300C) .

268

CHJIci-si-ct

ICHJI

Dim ethyldichlorosilane1 NaCH JI

[-Si-] +NaCI)n CHJ

n~6Hs ,0 Polydimethysilane

[-Si-O-B / I ~I '0 N N2 gas

C6Hs

Chapter 10 269

are not true linear PSPs (Equation 2). A model linear PSP was synthesized [44) by a ring-opening polymerization of 1,1 ,3,3-tetramethyl-1 ,3-disilacyclobutane (Equation 5).

n

(5)

The resulting linear PSP is a liquid oligomer at room temperature, with Mn = 2300 Da, incontrast to the high softening point of PCSs with similar Mn- Furthermore, the pyrolyticresidue yield of linear PSP at1000C is less than 20%. However, this yield increases if PSPis further heat treated, i.e., thermally crosslinked. For example, at400-500C, PSP yields asoluble solid precursor (Mn =8600 Da) with ayield of65% and, at900C, aceramic materialwith a yield of 45%. Analysis of the preceramic polymer shows a loss of both carbon andhydrogen compared to SiC2H6, as did the Yajima type PCSs. A more severe heat treatment(e,g., at450C in an autoclave) results in a material which becomes progressively insolublewith astill higher, up to85%, ceramic yield (Equation 4) [44). All these features show that thePCSs produced in the Kumada rearrangement of PDMS by the Yajima route are not linearPSPs but branched polymers.

Viscosity measurements and spectroscopic data (NMR, IR, UV) have shown that the structureofPCS polymers is roughly planar with three different atomic bonding schemes for the siliconatoms [11 [6) [9). The respective fraction of each atomic bonding scheme depends upon thePDMS/PCS conversion conditions (Table II). The structure of PCSs resulting from thecatalytic thermal rearrangement ofPDMS (PC-B type polymers) contains more Si-Si bonds. Itcould thus be belter described as polysilane chains connected through highly branchedcarbosilane nodes [44).

10.2.2 Melt spinning ofPCS

PCSs are usually spun in the molten state at about 300C. When necessary, PCSs arepretreated in order to adjust their molecular weight distribution [29) [46). Spinning isperformed under nitrogen since PCSs are sensitive to oxygen and moisture. The moltenfilaments are drawn down by mechanical pulling in order toachieve adiameter of15-20 IJm.

10.2.3 Stabilization and curing

Since PCSs are melt spun, they have to be cured before being pyrolyzed. In the genericprocess, infusible fibers are achieved by oxidation of the green PCS fibers. Oxidation ofPCSin air begins atabout 150C. The reaction isexothermic and the weight gain, which dependson the nature of the precursor, ranges from 8 wt.% (PC-A-470) to 17 wt.% (PC-B-3.2). Theempirical formulas for the oxidized PCSs are SiC163H3340020 (PC-A-470) and SiC176HmOo6B(PC-B-3.2), respectively [8) [13).

270

Wave number v,crn-t

Chapter 10

Figure 2.Infrared spectra of PDMS, PCA470 and its pyrolytic residues at increasing pyrolysis temperaturesaccording to reference (8); reproduced with permission.

IR spectroscopy shows that curing proceeds mainly by oxidation of Si-H bonds and, to alesser extent, of Si-CH3 bonds to form OH , C = 0 , Si-O-Si or Si-O-C groups [6] [8] [13].Oxidation ofSi-H and Si-CH3 bonds increases the degree ofcrosslinking of the PCS units, viathe formation of Si-O-Si orland Si-o-C bridges, rendering the polymeric precursor fiberprogressively infusible. Radical mechanisms have been proposed to account for theformation ofsuch bridges via the preferred oxidation of the Si-H bonds [6].10.2.4 Pyrolysis ofPCS fibers

PCS precursor fibers are converted to SiC based ceramic fibers by pyrolysis in an inertatmosphere at1200-1300C [6] [8] [13] [18] [28] [46-47]. This conversion is accompanied byevolution of gas (Figure 3) and weight loss. A large fraction of the organic bonds break at800-900C, and the pyrolytic residue is an amorphous, still hydrogenated, Si-C (or Si-C-O)

Chapter 10 271

containing material. With increasing temperature, heteroelements (H and 0) are released asa result of the scission of the last C-H bonds and decomposition of SiC,Ct. At 1200C,oxygen cured PC-A-470 loses 20 wt.%. Above 1600-1800C, the pyrolytic residue reachesits thermodynamic equilibrium state. However, since the tensile strength of the fibersproduced from oxygen cured PCS undergoes a maximum at 1200C, the pyrolysis of thecommercial fibers ends atabout this temperature. They are below their equilibrium state andare metastable when heated above their processing temperature.

6

400 800 1200Temperature, K

1600 2000

Figure 3.Gas evolution during the vacuum pyrolysisunder ofPCS fibers cured inhelium by E-beam irradiation (28);reproduced with permission from the American Ceramic Society, PO Box 6136, Westerville,Ohio 43086-6136.

The pyrolysis of PCS can be discussed by considering four temperature domains overlappingone another to some extent: 200-600C (domain 1); 500-900C (domain 2); 700-1200C(domain 3) and above 1200C(domain 4). The majority of the weight loss occurs indomain 1as a result of volatilization of low molecular weight PCSs anddehydrogenation/dehydrocarbonation, condensation and radical polymerization reactions. Atthe end ofdomain 1, PCSs are almost fully crosslinked .

The organic/inorganic transition occurs within both the second and third temperature domains[28] [47] with a fresh but minor weight loss and a density increase. The gaseous specieswhich are formed are mainly CH4 and Hzfrom scission of the relatively weak Si-CH3 and Si-Hlateral bonds indomain 2, and mainly Hz from scission of the stronger C-H bonds from the Si-CH2-Si backbone in domain 3. At the end of domain 2, i.e., at 800-900C, the pyrolyticresidue isan amorphous Si-C (or Si-C-O) material with a composition close toSiC16Ho65 [6]. Itisstill hydrogenated and contains an excess ofcarbon with respect to the stoichiometric PSPhaving a C/Si at. = 1. Italso displays numerous structural defects and itsdensity (2.21 g/cm3)isstill low. A tentative structural model has been proposed [6].

272 Chapter 10

As the temperature is raised todomain 3, hydrogen corresponding to the residual C-H bondsfrom the Si-CH2-Si backbone is progressively released with a continuous but slow densityincrease. Above about 1000C, clusters of free carbon and nanocrystals of P-SiC areformed . At the end ofdomain 3, ceramic grade Nicalon Si-C-O fibers consist of a dispersionof free carbon clusters and SiC nanocrystals in an amorphous Si-C-O matrix [14). As thetemperature is increased to domain 4, growth of P-SiC nanocrystals and decomposition ofternary silicon oxycarbide occurs in the pyrolytic residue, and evolution of CO and SiO isaccompanied by further weight loss.

10.2.5 Related Si-C-O(Ti) fibersTitanium can be introduced in PCS precursors as titanium tetrabutoxide, Ti{OC4H9)4, to yieldSi-C-O{Ti) ceram ic fibers such as Tyranno (Equation 6) [48). Fibers produced from suchPCS{Ti) precursors have a slightly higher pyrolysis temperature (Tp) than that previouslymentioned for their pure PCS counterparts [49]. As a result, they also retain their amorphousstate and hence their tensile strength toaslightly higher temperature.

PCS{Ti) precursors are produced in a one (or two) step process by crosslinking PCS chainswith titanium tetrabutoxide [48) or isopropoxide [49). The PCS chains are formed with the useof PBDPSO catalyst and no autoclave. As a result, PCS{Ti) already contains a significantamount ofoxygen in the as-prepared state.

~ OC H ~~ I 4 9 ~HC-Si--O-- Ti--O--Si-CH

3 I I I 3 (6)CH OCH CH

~ 2 4 9 ~ 2

PCS{Ti) precursors with Mn '" 1600 Da are melt spun at 270C and the green continuousfibers are cured in air at170C. During the curing step, part of the Si-H bonds still present inthe precursor isused to further crosslink the polymer through the formation ofSi-O-Si bridges,resulting in an increase of the oxygen content. The cured fiber is pyrolyzed in an inertatmosphere. Ata pyrolysis temperature of 1300C, the chemical composition (34.14 at.%Si,43.67% C, 20.79% 0, 1.22% Ti, 0.13% Nand 0.04% B) corresponds to a formula ofSiC1260061Tio OJ. Thus, the Si-C-O-Ti fiber contains an excess of carbon and the pyrolyticresidue is amorphous. Crystallization occurs above 1300C with an evolution of gaseousoxides, formation ofSiC and TiC nanocrystals, and adecrease in tensile strength.

10.3 Preparation ofoxygen-free SiC fibers

Nearly-oxygen-free Si-C fibers can be made by two routes. In one process, the precursorfibers are dry spun from a Yajima type PCS precursor, but the green fibers are cured byy-rays or E-beam irradiation [25-29). In the other process, curing is not needed. The greenprecursor fibers are dry spun from asolution ofsoluble but infusible PCSs [30).10.3.1 From radiation cured PCS precursor fibers

The first approach to nearly-oxygen-free SiC fibers is to mell spin green fibers from Yajimatype PCSs ( Mn = 1600-2000 Da) and to rendered them infusible by anaerobic curing with y-

Chapter 10 273

ray (e.g. , GOCo y-rays) orelectron beam (e.g., 2MeV) irradiation atroom temperature. E-beamcuring isusually preferred since the electron beam can be easily deflected toscan large fibersamples and provide ahigher irradiation dose in amuch shorter time.

Radiation curing of PCS requires aminimum dose of radiation, which isexpressed in gray orin rad (1 Gy =100 rad =1 J/kg). The degree ofcrosslinking can be estimated by measuringthe gel fraction, i.e., the percentage ofPCS insoluble in tetrahydrofuran (THF) after irradiation(Figure 4) [26]. In order to keep the fiber sample from melting, the sample containers aremaintained atroom temperature by water cooling.

75~~c0

'31 50Qi

274 Chapter 10

composition close to SiC141 which consist of P-SiC nanocrystals and free carbon and arevirtually free ofoxygen [25J [50J.

6.------------------.,

5

- Electron beam crossllnking (1.1 wt"k 0) Oxldatlve curing (13wt% 0) .c--;

..../ / \ .\.

/ \

....

400 800 1200 1600 2000

Heat treatment temperature,K

Figure 5.Variations of the entrapped radical concentration at room temperature as a function of the pyrolysistemperature in E-beam cured PCS fibers (1 .1 wt.% 0), from (47); reproduced with permission fnom the AmericanCeramic Society, Westerville, Ohio 43086.

Evolution ofH2 and CH4 (Figure 3) and weight loss can be observed during the pyrolysis ofE-beam cured PCS [28]. Evolution ofCH4 occurs as one single peak whereas that ofhydrogenoccurs as two overlapping peaks, suggesting a two step pyrolysis mechanism. The secondpeak at 1025C is almost the same for E-beam and oxygen cured fibers. Conversely, theintensity of the first peak (at 700C) strongly depends upon the percentage of Si-H bondsalready involved in the curing step.

The change in free radical concentration as a function of the pyrolysis temperature (Figure 5)shows that free radicals are formed and remain entrapped in the fibers [47]. The curve for theE-beam cured fibers can be decomposed into two peaks at 625C and 1025C, the formerbeing more intense than the latter (whereas it is the reverse for the oxygen cured fibers).There is a correlation between the gas evolution (Figure 3) and the entrapped radicalconcentration (Figure 5), suggesting a two step pyrolysis process.A radical reaction mechanism has been proposed. The evolution of H2 and CH4at lowtemperatures is related tothe formation of=Sibased free radicals corresponding tothe scission of theweaker Si-H and Si-CH3 bonds; evolution of H2 at higher temperatures is related to theformation of=Cbased free radicals, which result from the scission of the stronger C-H bondsin the Si-CH2-Si groups.

Chapter 10

10.3,2 From infusible PCS precursor fibers

275

The second approach to the production of Si-C fibers with a low oxygen content is based onthe use of PCS precursors which are infusible (Mw = 5000-10000) but completely soluble inorganic solvents such as toluene [30] [51-52]. Such PCSs can be prepared via the thermalrearrangement of POMS in an autoclave at 435-480C, as in the Yajima route, but theduration of the heat treatment is longer.

Green fibers are dry spun from the PCS solution with spinning aids such as high molecularweight polyvinylsilazane (PVSZ) and a small amount of a hydrosilylation catalyst, such asdicumylperoxide [52], to generate free radicals. The vinyl groups contribute to thecrosslinking of PCS chains through reaction with Si-H bonds, The solution of PCS and PVSZis filtered and is concentrated todevelop appropriate rheological properties.

The solution is dry spun at room temperature and the green fibers are wound onto a drum.The solvent is partly evaporated between the spinnerets and the drum, producing infusiblegreen fibers. These fibers are slowly heated to 200C in an inert atmosphere to release allresidual solvent and to fully crossl ink the polymer. They are finally pyrolyzed in argon below1000C to yield a black Si-C ceramic fiber. This fiber still contains a small amount ofhydrogen and oxygen , and a significant amount of excess carbon. It is amorphous and hasan empirical formula ofSiC178HolSOoosNoo4, but tends to crystallize above 1200C into a mixtureofP-SiC and free carbon .

10.4 Preparation of quasi-stoichiometric SiC fibers

When the pyrolysis of PCS precursor fibers is carried out in the presence of hydrogen, orwhen boron or aluminum doped PCS precursor fibers are pyrolyzed, it is possible to obtainquasi-stoichiometric silicon carbide fibers, Alternatively, quasi-stoichiometric fibers are alsoobtained from precursor fibers consisting ofSiC powder reinforced polymers.

10.4.1 Pyrolysis ofpes precursor fibers under hydrogenThe excess ofcarbon in Si-C fibers can be lowered bypyrolyzing Yajima type PCS precursorfibers in an atmosphere of hydrogen [33]. Hydrogen is assumed to favor release of thependent methyl groups and formation of CH4 during the organic/inorganic transition .Complete pyrolytic release of pendent methyl groups from the ideal polysilapropylene chainpolymer (Equation 1)would yield stoichiometric SiC, assuming that the carbon from the Si-CH2-Si backbone remains in the solid state. The residue after pyrolysis at 1000C is thoughttobe more hydrogenated than the residue prepared inargon. Quasi-stoichiometricSiC fibershave been recently produced from fusible [53] and infusible [34-35] PCSs. The processingdetails have not been disclosed, but it is likely that a hydrogen or hydrogen containingpyrolysis atmosphere was used.

10.4.2 Pyrolysis ofboron doped PCS precursor fibers

Yajima type fibers, produced from oxygen cured PCS, decompose above 11 00-1200C inaninert atmosphere with evolution of CO, This process can be used to simultaneously removeoxygen and carbon from the fiber. If the temperature is sufficiently high (e .g., 1800C),removal of oxygen is almost complete leaving behind a pyrolytic residue that is either amixture ofSiC +Corstoichiometric SiC having less than 0,5wt.% residual oxygen .

276 Chapter 10

Unfortunately, the Yajima type fibers, when treated at 1800C,become porous and thereforeextremely weak unless they are doped with SiC sintering aids. For example, doping Si-C-OYajima type fibers with 0.2 - 0.6wl.% Byields dense fibers after pyrolysis at1800C [37-39J.Ifenough oxygen is introduced when the green fibers are cured, stoichiometric SiC fibers mayresult having 97% of the theoretical density and less than 0.5wt. %O. The boron dopant canbe introduced in the green fiber by curing it in an atmosphere containing a gaseous boronbearing species [38-54J.Boron is a well-known sintering aid for silicon carbide, but temperatures of at least 2000Care required for sintering polycrystalline SiC powders with a grain size of the order of 1 IJm.Its effectiveness in sintering SiC based fibers at lower temperatures (1600-1800C) isprobably related to the extremely small SiC grain size ofPCS based materials.

Similarly, an experimental, quasi-stoichiometric, oxygen-free SiC fiber has been producedfrom a Si-AI-C-O precursor. In this process, aluminum is introduced into the polymer asaluminum (III) acetylacetonate, and acts like boron as a SiC sintering aid during heattreatment at1800C [40J.10.4.3 From extruded SiC pOWder/polymer mixtures

Continuous quasi-stoichiometric SiC fibers can also be produced from SiC powder/polymermixtures [41J [57-58J. These mixtures consist of a fine SiC powder (usually a-SiC) with agrain size less than about 1 IJm and a boron bearing species, i.e., boron carbide, B,C. Thesource ofcarbon is typically a phenolic resin and various polymers to render the mixture meltspinnable. After melt spinning, the green fibers can be directly fired or they can be woven andthen fired . During firing in nitrogen, residual oxygen is released as CO, and B,C acts as asintering aid. Since the grain size of the SiC powder is relatively large, firing has to beperformed at 200Q-2300C to achieve a high degree of densification. The draw ratio duringmelt spinning is low and the ultimate fiber diameter remains high 100 IJm).10.5 Structure of siliconcarbide fibers

Fibers resulting from the pyrolysis of PCS and related precursor fibers consist of pure SiC,SiC + Cor SiC + SiO,Cv + C, depending on their processing conditions. Silica is not presentexcept eventually as adiscontinuous thin film at the fiber surface.

10.5.1 Silicon oxycarbide fibers

Fibers resulting from the pyrolysis of oxygen cured PCS precursor fibers at 850-1000Cremain amorphous and still contain significant amounts of hydrogen and C-C bonds. With afurther increase inpyrolysis temperature to1200-1300C, hydrogen is progressively releasedand SiC nanocrystals and free carbon clusters are formed. Specifically, the resulting fibersconsistofa mixture of P-SiC nanocrystals and partly hydrogenated carbon clusters which aredispersed inan amorphous silicon oxycarbide phase, usually formulated as SiO,Cv

The formation of both P-SiC and carbon is evidenced by transmission electron microscopy(TEM) [1J [16J. The P-SiC phase appears as tiny crystals with a mean size of 2 nm. Thecarbon phase appears as randomly oriented 0.7 to 0.8 nm BSUs, which consist ofpolyaromatic species associated face to face and which could be related tocoronene, C2,H,2[16). PCS pyrolytic residues contain 20 at. %residual hydrogen after pyrolysis at850C and 4

Chapter 10 277

at.% after pyrolysis at 1000C (46). Since the C-H bond is a strong bond, it is likely thatresidual hydrogen is present as C-H bonds. Finally, free carbon can also be observed byRaman spectroscopy. The Raman spectrum of Si-C-O fibers exhibits an unresolved bandafter pyrolysis at 850C. This Raman band splits into two lines at 1580 and 1350 crrr' afterpyrolysis at1000C. (1).The formation ofa ternary phase in Si-C-O fibers isevidenced in x-ray photoelectron spectraorXPS (14) [59-63]. The C1s and Si2p photopeaks recorded from the surface ofa typical Si-CoO fiber are asymmetrical (Figure 6) and can be decomposed into several components.Component I of the C1s and Si2p photopeaks (binding energy, BE, of 283 .3 and 100.5 eV,respectively) corresponds to carbon and silicon from SiC (BE = 282.9 and 100.5 eV,respectively). Similarly, component III of the Si2p photopeak (BE = 103.1 eV) is assigned tosilicon in silica (BE = 103.4 eV).The fact that its intensity decreases dramatically after argon ion etching shows that silica isonly present at the fiber surface [16] [63]. Component II of the Si2p photopeak whichcorresponds toa binding energy, i.e., 101 .5eV, intermediate between those for silicon in SiC(100.5 eV) and silicon in silica (103.4 eV) , is assigned toa ternary SiO,C, phase, where Si istetrahedrally bonded to both C and 0 atoms. Component II of the C1s photopeak (BE =284.6 eV) corresponds either to C atoms from the hydrogenated free carbon phase or to Catoms from the ternary SiO.C, species.

(b) C1s

290

SiC

106

Binding energy, eV

SiC

98

Figure 6.C1s and Si2p core level spectra recorded from: (a) unetched and (b) Ar-etched Si-C-O Nicalon fibersderived from PCS, --, experimental; - - - background substraction; - - calculated (63);reproduced with permission.

Determin ing the molarcomposition of the Si-C-O fibers is complicated [15-16] [59-60) (63).There is already some dispersion in the elemental overall analysis data even for the majorelements (Table III). Neglecting residual hydrogen, the mean composition of the Si-C-O fiber

278 Chapter 10

(Nicalon NL 200) is close to 38 a1.% Si, 48 a1.% C and 14 a1.% and corresponds to theoverall formula SiCl260037.

Table ill. Elemental composition of Si-C-C fibers derived from PCS precursor fibers

Silicon oxy- Silicon Carbon Oxygencarbide fibers at% wt% at% wt% at% wt%

Nicalon NL-200'I/ 37.70 57.90 43.10 28.40 15.00 13.20PeN-based 470 34.40 51.30 53.40 34.10 11.20 9.55Nicalon NL-200(21 39.00 58.00 47.00 30.00 14.00 12.00Nicalon NL-22~' 36.49 55.64 49.95 32.57 13.56 11.78Nicalon NL-20i" 37.00 54.20 47.40 29.80 15.60 13.00Nicalon NL-202(21 39.50 58.90 48.50 30.90 12.00 10.20Nicalon NL-20i41 37.80 57.10 49.80 32.20 12.40 10.70

Hydrogenat% wt%3.7 0.200.9 0.05

Nitrogenat% wt%0.5 0.3

,II [16J; (21 EPMA-data, [59J[15J[50J; W chemical analysis, [61J; (41 XPS-data, [60J

The calculation of the molar composition supposes astructural model for the SiO,Ct phase. Itis usually assumed that the oxycarbide consists of tetrahedral units, SiCaO"" in which thesilicon atom is bonded to carbon and oxygen . The compositional domain of the tetrahedralunit ranges continuously from Si04 (a =0) to SiC4 (a =4). Additionally, it isassumed that thecarbon atom remains tetrahedrally bonded to 4Si (as in SiC), and the oxygen atom is bondedto2Si (as in silica), with no other crossbonding. Under such assumptions, the formula for thesilicon oxycarbide derived from that ofthe tetrahedral unit is:

SiCa / .PI / 2( -I-a)If Y=a/4 and x =1/2{4-a), the general overall formula of the oxycarbide SiO,Ct can bewritten as [59]:

a=4y and x =2(1- y) or xy=J--2

(7)

The variations of x and y as a function of a, and reciprocally those of a and 4 - a as afunction of y, which give the correlation between the nature of the tetrahedral unit, SiCaO",.and the composition ofthe silicon oxycarbide, SiO,Ct, are shown in Figure 7.

Si02 Si03l2C1 /4 SiOC,12 Si0112C3I4 SiC

400.25 0.50 0.75 14

~3 3

~2 2~tf~.~

--- 1i ......... ......_........-...-00 -

~:a:~--t----- ~ ~2 a 3

Si04 Si03C Si02C2 SiOC3 SiC4

Figure 7.Correlation between the nature ofthe tetrahedral SiC.04-a un~ and the chemical overall formula inthe siliconoxycarbide.

Chapter10 279

The molar composition of Nicalon NL 200 can be computed as 46.06% SiC, 23.03% SiOxC,.xI2[X= 1.1053] and 30.64% free carbon, assuming: (1) the Nicalon NL 200 Si-O-C fiber consistsof SiC, SiOxC,.x12 and free carbon; (2) its atomic composition is 38% Si, 48% C and 14% 0;and (3) the intensity ratio of the two photopeaks assigned toSiC and SiOxCv (and reported tobe 2 for Nicalon NL 200) gives the molar ratio between the two species.Thus, the fibers resulting from the pyrolysis at 1200-1300C of oxygen cured Yajima typePCS precursor fibers in an inert atmosphere are far from consisting of pure SiC since themolar fraction of SiC is about 50% . They also contain significant amounts of partlyhydrogenated free carbon and silicon oxycarbide. It is noteworthy that the value, or perhapsmean value [14] [64], of the term 1-x/2 = 0.447 is close to 0.5 in Nicalon NL-200 (Figure 7)and that the main tetrahedral units are therefore close toSi02C2 [15].The atomic distribution of silicon, carbon and oxygen in the three phases (SiC, SiOxC,.x12 andfree carbon) can be calculated as a function of x for a given overall atomic composition [59].Other approaches have been used to calculate the molar composition of PCS based Si-C-Ofibers from the overall atomicpercentages [16] [60].10.5.2 Silicon carbide fibers

Silicon carbide fibers with low oxygen content fall in two groups. One group of fibers displaysa C/Si ratio higher than one and contains free carbon , e.g., 30 mol.% for Hi-Nicalon. Theother group of fibers displays a C/Si ratio close toone, e.g., 1.05-1.07, for quasi-stoichiometricSiC fibers having a free carbon content of

280 Chapter 10

For apyrolysis temperature of1600-1800C and C/Si ratio of1.05, the fibers made from PCSprecursor fibers have an improved crystallization state [34-36] [53-54] [65]. Their SiC grainsize is 5-20 nm. Although the dominant phase is still P-SiC (or 3C), the fibers also containsmall amounts of hexagonal a-polytypes. Numerous crystal defects such as twins andstacking faults are observed, features which are common in SiC ceramics. The amount offree carbon is lower and "clean" SiC/SiC boundaries are locally present. The carbon phase isbetter organized, and the carbon layer stacks are thicker and more extended. Finally, thefiber density (3.1 g/cm3) isclose tothat ofdense SiC (d lheor. =3.21 g/cm3) .

Figure 8.TEM-micrograph of a nearly-stoichiometric SiC UF-HM fiber showing the SiC grain size. The electrondiffraction pattern isshown inthe inset; reproduced with permission from the American Ceramic Society, Westerville,Ohio.

For a pyrolysis temperature of 20002300C and a C/Si ratio assumed to be slightly higherthan one, fibers produced from sintered powders display a state ofcrystallization resemblingthat of sintered bulk SiC ceramics [41] [57-58]. The microstructure of such fibers is muchcoarser, with SiC grains of about 1 urn in size and impurities (e.g., B4C) at the SiC grainboundaries. Finally, the SiC fibers prepared at high temperatures usually exhibit a surfaceconsisting ofalmost pure carbon over adepth of50-200 nm [54] [65].10.6 Thermal stability of silicon fibers

The Si-C-O fibers produced from oxygen cured PCS precursor fibers at 1200-1300C areunstable at high temperatures. Si-C fibers, i.e., almost pure SiC or SiC + C mixtures, arethermodynamically stable to 2500C. The only cause of microstructural instability is theextremely small grain size of the phases. When such fibers are maintained at high

Chapter 10 281

temperatures, grain size tends to increase. The driving force is the reduction in surfaceenergy. Chemical reactions occurring in Si-C-O fibers and SiC grain growth in Si-C-O, Si-Cand SiC fibers cause adramatic reduction in tensile strength.

10.6.1 Silicon oxycarbide fibers

When a Si-C-O fiber is heated in vacuum or an inert atmosphere above its processingtemperature, an added weight loss occurs that is accompanied by evolution of gaseousspecies and by a profound change in its microstructure [16-17] [21] [23-24] [29] [66-71]. Theweight 1055 rate remains slow up toabout 1250C (Figure 9). The overall asymptotic weight1055(~W

282 Chapter 10

other by C cages and amorphous SiO,C,.The kinetics ofSi-C-O fiber degradation, expressed as the variation in decomposition ratio, X,as a function of time have been reported toobey an Avrami-Erofeev type equation [69]:- Ln ( I - X) =kt"

Ln {- Ln ( I - X)1= Ln k +m Ln tor

(8a)

(8b)where k isa rate constant and maconstant whose value is in the range 0.5 - 4 and dependsupon the mechanism of the transformation. It appears from Figure 10 that the fiberdecomposition kinetics obey Equation 8b with a value ofmof the order of 1.5 for sufficientlyhigh temperatures. Thus, decomposition might be kinetically governed by the rate ofgrowthof the SiC nanocrystals bydiffusion, the nanocrystals growing 3-dimensionally. Finally, thenoted decomposition is thermally activated and the apparent activation energy, 633-791kJ/mol, isvery similar to that reported for the diffusion ofCinSiC (560 - 840 kJ/mol).

4 '--"'-__'--_--l.__---'-__...J6.0 6.2 6.4 6.6 6.8

Reciprocal temperature T-1 , 10-4K-1

Figure 10. Decomposition kinetics of Si-C-O (Nicalon NL 200) fibers: Avrami-Erofeev plot of the data[69):reproduced w~h permission from the CeramicSociety ofJapan,Tokyo.

From the experimental data reported above, the decomposition might occur by a mechanisminvolving the two following equations:4 SiOxCI-x I2~ (4 -3x) SiC+x CO+3 x SiO (9)y SiO +2yC~ y SiC +Y CO (10)4 SiOxCI-x l2+2yC~ (4-3x+ y ) SiC+(x+ y) CO+(3x- y ) SiO (11)

Chapter 10 283

This mechanism accounts for the evolution ofa gaseous SiO + CO mixture, the growth of theSiC crystals and the decrease in the amounts of free carbon and silicon oxycarbide. Therelative amounts of CO and SiOwhich are formed and the composition of the final residue(i.e., SiC orSiC + C) depend upon the relative amounts of free carbon and silicon oxycarbidein the fiber. For the Nicalon fiber (NL 200), the main species in the gas phase is CO and thesolid residue is SiC. There is therefore enough SiO formed bydecomposition of the siliconoxycarbide phase (Equation 9) to consume all the free carbon by Equation 10 [731. Thismechanism also accounts for one of the processes used to produce oxygen-free nearly-stoichiometric SiC fibers [38] [54]. Since CO and SiO diffuse and escape from the fiber, thedecomposition starts near itssurface and the decomposition front moves radially towards thefiber axis, yielding askin/core microstructure [16].The decomposition of Si-C-O can be impeded, or at least shifted to higher temperatures, bysubjecting the fibers to a pressurized gas or a gas tight coating. The decomposition ofNicalon fibers is shifted to about 1800C when an isostatic argon pressure of 138 MPa isapplied to the fibers [68] [74-76]. Similarly, the stability of the fibers is enhanced at 1300Cwhen they are treated inAr-CO gas mixtures with Peo =40 kPa. In addition , a carbon buildupon the fiber surface results from adecomposition ofCO:

(12)Slightly pre-oxidized Si-C-O fibers have a uniform glassy silica coating . Their decompositionoccurs at higher temperatures. This suggests that the silica coating remains gas tight andprevents the evolution ofCO and SiO up to at least 1400C [69].10.6.2 Silicon carbide fibers

Oxygen-free Si-C fibers, consisting of SiC and free carbon, are expected to be stable up to2500C, but experimental data show that the fibers experience surface decomposition andgrain growth when aged athigh temperatures. Surface decomposition, which is related to asurface vaporization of Si, is driven by the formation of a surface layer of carbon. Graingrowth of the nanometer size SiCcrystals isdriven bya reduction ofsurface energy.

A Si-C fiber experiences no noticeable weight loss when aged up to about 1800C in anatmosphere of argon and itschemical composition does not change much. However, someevolution ofhydrogen continues tooccur [50] and acarbon layer builds up on the fiber surfacewhose thickness increases with time ata given aging temperature, e.g., to30 nm for one hourofpyrolysis at1600Cand to70 nm for ten hours ofpyrolysis.

Some growth of the SiC grains occur when the pyrolysis temperature is raised, even in Hi-Nicalon fibers containing a large amount of free carbon. However, the maximum grain sizeremains relatively small, e.g., 20 nm at1400Cand 50 nm at1600C, compared with those ofheat treated Si-C-O fibers. Free carbon is in equilibrium with SiC. However, it undergoessome reorganization with an increase of both the carbon layer size (La) and the stackthickness (N). This might occur by de-wrinkling of the layer stacks and their edge-to-edgeassociation into larger layers tending to lieflat on the faces of the growing SiC crystals.

Near stoichiometric, dense SiC fibers are produced at higher temperatures (1800 to 2300C)with a boron bearing sintering aid; they do not contain significant amounts of free carbon [34-35] [41] [53] [72]. As a result, the SiC grain size is relatively large inthese fibers, e.g., 50-200

284 Chapter 10

nm, and a few crystals are as large as 1.0 IJm (35). Some of the secondary intergranularphase, e.g., carbon, may impede the growth of the SiC crystals and the SiC fibermicrostructure may have ahigher thermal stability between 1200 and 1600C when they havebeen processed at higher temperatures. Therefore, the microstructure is to some extentstabilized during processing.

10.7 Mechanical properties of SiC fibers

SiC based fibers display a linear elastic behavior at room temperature. Their Young'smodulus, E, and tensile strength, OR, depend on the processing conditions, particularly thepyrolysis temperature, Tp.

10.7.1 Atroom temperature

At first, both the modulus and tensile strength of Si-C-O fibers increase as the pyrolysistemperature is raised to >850C, i.e., as hydrogen is released and densification occurs.Then, they decrease sharply above 1200C when the fibers undergo decomposition. There isthus an optimum Tp value. The modulus and tensile strength of the oxygen-free SiC + Cfibers produced from E-beam cured PCS precursor fibers also show maxima as the pyrolysistemperature is raised but the maxima are shifted tohigher temperatures and the decrease inboth E and OR is gradual. The fibers still have high stiffness and strength at 2000C (Figure11 ).

4r-----------------,

.., SiC+Cfiber

~ ---)(--)(--.,.; --oK ---....Si-C-O fiber

1200 1600Temperature. C

2000

Figure 11 . Variations of the room temperature tensile properties of SiC-based fibers as a function of thetemperature. The pyrolysis temperature forSi-C-O fibers derived from PCS and Si-G fibers derived from radiation-cured PCS with0.4 wi.%ofoxygen (77]: reproduced with permission from the Woodhead Publishing Ltd.

The modulus of SiC based fibers strongly depends upon the occurrence of intergranularphases. It is of the order of 580 GPa for SiC whiskers and 400-450 GPa for the nearlystoichiometric, oxygen-free polycrystalline SiC fibers [54) (65). Conversely, it is much lower

Chapter 10 285

when the fibers contain free carbon and/or amorphous silicon oxycarbide. The modulus ofoxygen-free SiC +C fibers decreases when the C/Si ratio increases from 1 to1.6. Finally, forthree reasons, the Si-C-O fibers consisting ofa SiC +C+ SiO,C 1.J

286 Chapter 10

(16b)

or the occurrence ofspecific flaw populations, e.g., a population of internal flaws (ml/crol) andapopulation ofsurface flaws (mJcr02) [79-82] :

PR =l- exP[-~(~J",J _~(~)",1] (16c)Vo O"oJ So 0"02

The data for recent Nicalon NL-200 fibers ofagiven gauge length (L=25 mm) are spread overawide range from 1500 to 4000 MPa. The average isabout 3000 MPa (Figure 12). In a firstapproximation, these experimental data can be fitted to a unimodal two parameter Weibulldistribution. There are often a few data points, usually at low stress levels, which do notactually fall on the straight line predicted by equations of type 15a or 15b, suggesting thatmore than one flaw ispresent in the fiber.

0 .90

0.60l!!::l

S(; 0.30.~I 0 .10D..

Nicalon NL200

1500 2500Tensile strength. MPa

3500

Figure 12. Weibull plot of tensile strength data for Si-C-O (Nicalon NL 202) fibers. with a un imodal Weibulldistribution (80); reproduced with permission from VSP, Zeist, NL.

About 96% of the fracture origins in Nicalon NL200 fibers are surface flaws and only 4% areinternal flaws. Since the internal flaws are minimal, the corresponding data can be treated ascensored data and the tensile strength of the fiber can be statistically analyzed (Figure 12)with a unimodal Weibull distribution, where m = 4.5 and cro = 2670 MPa [80]. Data for anolder Nicalon fiber. however, gave a better fit with a bimodal Weibull distributioncorresponding toa family ofsurface flaws (m, = 3.64; crOl =4.64 GPa) and a familyof internalflaws (rn, =9.41 and cr02 =5.08) [79].More recently, the observation of two partly concurrent populations of flaws in Nicalon NL200fibers has been attributed [82] to extrinsic and intrinsic flaws. Accordingly, a family ofextrinsicflaws (severe flaws at the fiber surface) is responsible for fa ilure at the lowest stress levels

Chapter 10 287

(17)

(with m2 =1,92), and a family of intrinsic flaws (located both atthe surface and in the volume)control failure athigher stress levels (with m,=4,5),The identification of families of flaws, which are responsible for the failure of a batch of fiberspecimens, is not straightforward. Before testing, the specimens must receive a solubledamping coating ofawax ora polymer such as 1-3 polypropanediol, or they must be tested ina liquid medium such as glycerol to absorb the shock wave energy due tospecimen burstingatfailure. Under such precautionary conditions, the primary failure surfaces can be recoveredand the origin ofthe fracture determined [79-80].The surface flaws are mainly microvoids and microcracks induced by spinning andmechanical abrasion whereas the internal flaws are microvoids or inclusions originating fromthe PCS precursor fiber or formed during pyrolysis. The failure surface often shows thefracture origin-mirror-mist-hackle-crack branching morphology characteristic of brittlematerials. The depth a and the width c of surface flaws in Nicalon NL200 fell in the ranges0.09 < a < 0.6 IJm and 0.25 < c < 1.7 IJm, respectively. The surface flaws were modeled assemi-elliptical cracks, allowing calculation ofthe fiber toughness, KIC, by the Griffith equation:

I K/cO"R =y' (J(Qc/ 12where Yisageometric parameter depending on the shape of the defect (here, Y=0.8331 , foralc = 0.3). K,C, calculated from failure stress (crR) and flaw depth (a) data, is of the order of1.55 -1 .9MPa m".

The tensile strength of Si-C-O fibers decreases after exposure to elevated temperatures.When Nicalon NL 200 fibers are exposed for 1 hour to 1300C in argon (P = 100 kPa), theirmean tensile strength and scale parameter, cr., decrease by 45% while their Weibull modulusremains unchanged [80-83]. Fibers exposed to more severe conditions (e.g ., for 5hours inavacuum at1500C)are so weak that they cannot be tested. Finally, the fact that oxygen-freefibers maintain their tensile strength under similar conditions relates to the absence of siliconoxycarbide and its decomposition process.

10.7.2 Athigh temperatures

The high temperature (HT) properties of SiC based fibers depend upon their pyrolysistemperature, their thermal stability, the time and the test atmosphere. When mechanical testsare performed at a test temperature that is higher than the pyrolysis temperature, somechange occurs in the composition orland microstructure ofthe fibers which may strongly affecttheir mechanical behavior.

(a) Tensile testsYoung's modulus and tensile strength of Si-C-O and Si-C fibers decrease when the testtemperature is increased (Figure 13). The room temperature tensile strength of the Si-Cfibers (Hi-Nicalon) is higher than that of Si-C-O fibers (Nicalon NLM 200), but the two fibershave almost the same strength, (1200 MPa), when tested at 1400C. At any testtemperature, the modulus of the Si-C fiber ishigher than that ofthe Si-C-O fiber [31] [84]

288

4

3caQ.e

~15>

Chapter 10 289

1.2 t l 1350~

o~ 0.8 ff

.= o 0.4 200

-

o I I I I Pure Argon

O 400 1200 2000 Time, min

Figure 14. Creep of Si-C-O (Nicalon NL 200) fibers under an applied stress of 0.45 GPa as measured in argon [59]; reproduced with permission from the American Ceramic Society, PO Box 6136, W esterville, Ohio 43086-6136.

Si-C-O fibers are prone to creep above 1100~ (Figure 14). The creep rate does not reach steady state prior to failure during tensile tests performed in argon. It decreases continuously with time, indicating that primary (or logarithmic) creep dominates the entire creep life of the fibers. The strain-time curves obey the following classical law:

1 c = - - L n ( 1 + f l ~o t ) (18)

P where p is a constant and /; o is the strain rate at zero strain. Creep is rate controlled by the viscous flow of the glassy silicon oxycarbide containing free carbon and SiC nanocrystals in suspension [21] [31]. Above 1100~ silicon oxycarbide decomposes with formation of more SiC, resulting in a continuous increase of the viscosity and thus a decrease of the creep rate.

Conversely, the creep curves for fibers tested in CO-rich atmospheres also display a large steady state domain (secondary creep) [59]. This linear domain is related to the fact that, under such conditions, the silicon oxycarbide and the fibers are stable. For higher test temperatures (Tt ___1400"C), the fibers are no longer stable for Pco _

290 Chapter 10

SiC + C fibers (e.g., Hi-Nicalon) creep at temperatures as low as 1000C although they arealmost totally free ofsilicon oxycarbide. This suggests a different creep mechanism (32). Asteady state domain isalways observed ata test temperature of ::;;;1400C, i.e., when the testtemperatures is lower than the assumed processing temperature (Figure 15). Conversely, ata test temperature Tt ~1500C, the strain rate decreases continuously with time up to failureand primary creep dominates the entire creep life of the fiber. The apparent activation energyis 0, = 220 kJ/mol for 1000 < Tt < 1250C and ~ = 700 kJ/mol for 1250 < Tt < 1400C,suggesting that an important change occurs in the creep mechanism.

The creep rate of Hi-Nicalon type fibers is significantly lowered when the fibers are heattreated at 14001600C in argon prior to the creep tests (also performed in argon). Further,fibers which are heat treated at 1600Cexhibit a steady state creep domain atTt =1500C,whereas the untreated fibers do not creep when tested at Tt = 1500C (32). The creepresistance of heat treated fibers at 1400C in air is higher than that of the untreated fibers(87).In the family ofTyranno Si-M-C-O fibers (M =Ti, ZrorAI) byUbe, aquasi-stoichiometric fiberwith acomposition close to68 wt.% Si, 31% C, 0.3% 0, and 0.6% AI (C/Si = t07) offers thehighest creep resistance at1300Cunder a tensile load of1GPa. The creep mechanism hasbeen reported tobe transgranular diffusion (89). Finally, a sintered a-SiC fiber produced byCarborundum has a measurable creep rate ranging from of5.1 x 10-8 to 1.4 X 107S1 between1300 and1500Catstress levels ranging from 70 to300 MPa (88). From preliminary data, n=1.86 0.50 and 0 =273 12 kJ/mol.

2.0 ...------------------,

1 GPa Ar

2.0

3.0

1.0

105

l- ....L... ....I..... ---l 0.0

15

EE. 1.0

Chapter 10

(c) Bend stress relaxation test291

A simple bend stress relaxation test can be used to compare the creep resistance ofindividual filaments [43]. A bending stress is applied with a graphite jig to the fiber whichacquires a radius of curvature, Ro. The assembly is submitted to a heat treatment. Aftercooling , the graphite jigpieces are separated revealing a fiber with acreep induced curvatureofradius R (with R2:Ro). The stress relaxation isquantified by aparameter mdefined as m=1 - (RJR). For given test conditions, creep resistant fibers are characterized by m-valuesclose to1(i.e.,R Ro) whereas creeping fibers display low mvalues (R "" Ro).Si-C-O fibers (Nicalon NLM-200) creep at low temperatures (1000-1200C), whereas thesintered a-SiC fibers (Carborundum) creep at much higher temperatures (1300-1600C).The other fibers, which are derived from PCS precursor fibers, i.e., oxygen-free Si-C fibers(Hi-Nicalon) and the quasi-stoichiometric fibers (Hi-Nicalon S or Dow Corning fiber) fallbetween these two limits. The fibers which have been heat treated at 1400-1600C anddisplay a larger SiC grain size (e.g., 30-50 nm) exhibit acreep resistance similar to that of thea-SiC sintered fiber.

In summary, there isaclose relationship between the creep resistance and the SiC grain size.The larger the grain size, the higher the creep resistance. However, since tensile strengthusually decreases as the grain size increases, it appears that optimizing the fibermicrostructure in order to achieve both a high tensile strength and a good creep resistancemay be contradictory [90].10.8 Oxidation ofsilicon carbide fibers

Depending on the oxidation conditions, the oxidation of silicon-rich non-oxide ceramicsproceeds by passive or active oxidation . At low temperatures and high oxygen partialpressures, aprotective layer ofsilica is formed by passive oxidation (Equations 20a and 20b).SiC(s)+202(g)~ Si02(s) + CO2( g)SiC(s)+3 /2 02 (g)~Si02(s,l)+CO(g)

(20a)(20b)

At high temperatures and low oxygen partial pressures, silica is no longer formed; it isconverted to gaseous SiO by active oxidation (Equation 21).SiC(S)+02 (g) ~SiO(g)+CO(g) (21)In the passive oxidation regime, the ability of Si based ceramics to form a continuousprotective layer of silica depends upon their silicon content and the volume change occurringduring oxidation. It isoften characterized with aparameter, Ll, defined as:

~=A._d_dSi0 1

M siowith A =CSi .__1

M.w

(22a)

(22b)

where CSi is the weight fraction of silicon in the material, MSi02 and MSi, the molar weights ofsilica and silicon, and d and dSi02, the densities of the ceramic material and silica. Theoxidation yields a silica scale covering the material when ~ 2: 1. For dense SiC (CSi = 0.70;

292 Chapter 10

dsiC =3.2 g/cm3) , t!. '" 2.2, assuming that glassy silica is actually formed (with dSio2 =2.20g/cm3) .

SiC based fibers have t!. values ranging from 1.35 for the Si-C-O fibers (Nicalon NLM 200)and 1.66 for oxygen-free SiC + C fibers (Hi-Nicalon) to almost 2.2 for quasi-stoichiometricfibers (Hi-Nicalon S). Thus, oxidation yields aprotective silica layer, which remains glassy uptoabout 1200C, but becomes cristobalite at1400C. The glassy sheath contains nanoporesand probably some hydrogen as hydroxyl ions arising from the residual hydrogen in the fibers.Conversely, the cristobalite sheath may crack after cooling. When the silica sheath iscontinuous, it impedes evolution of CO/SiO and fiber decomposition as long as the internalCO/SiO pressure is lower than the external pressure [22-23) [29) [32) [91-92].The oxidation of SiC based fibers becomes noticeable above about BOOC. It occurs with anoverall fractional weight gain t!.rrlmo. The oxidation kinetics can be determined by TGA underflowing oxygen (or air) or/and by SEM of the thickness of the silica sheath (Figure 16). Whenthe weight gain remains

Chapter 10 293

suggests the oxidation process is rate controlled by diffusion, the rate limiting step beingpresumably the diffusion ofoxygen across the silica layer.

In cylindrical geometry, the surface through which diffusion occurs is not constant. Theexternal surface of the silica layer is larger than the inner surface and both change with time.As long as &1m. (thus, e) remains small 10%), this difference and the gas evolution can beneglected toa first approximation, and the kinetic constant can be calculated byapplying theclassical parabolic law for diffusion across aplanar interface:

(24a)(24b)

kl and k2 are kinetic constants, &Jmo and ei are the weight gain and silica thickness att =O.kl is expressed in reciprocal time unit whereas k2 has the dimension of a diffusion constant(L2t1) . k1 ork2 iscalculated byfitting the experimental data toEquations 24a or24b. Thermalvariations of the kinetic constants obey Arrhenius law since it is a diffusion process (Figure16):Kr = Kooexp - (E; / RT)or Ln Kr =Ln Koo - Ea / RT

(25a)(25b)

The corresponding apparent activation energy, Ea, is 107 kJ/mol for Si-C (Hi-Nicalon) fibers[50] and 69-77 kJ/mol for Si-C-O (Nicalon) fibers [22]. These values are lower than thosereported for pure dense SiC (e.g.,128 kJ/mol for SiC) deposited by CVD/CVI at1000C.A different treatment ofTGA based &1m. data accounts for the decrease of the surface areabetween the growing silica layer and the unreacted fiber, assuming the oxidation process isalso rate controlled by diffusion (90). The kinetic constant, I

294 Chapter 10

202) fibers [50) at room temperature. For a given precursor, the electrical properties dependstrongly on the pyrolysis temperature. Selected Si-C-O [93) and Si-C-o-Ti [49) fibers areavailable which exh ibit either a low (10-6 to 107n-1cm')ora high (0.5 to 101'11 crrr') electricalconductivity atroom temperature.

The electrical conductivity of a PCS residue increases dramatically when the pyrolysistemperature is increased from 700-800C to 1400-1600C. This increase is as large as 10orders ofmagnitude for uncured bulk PCS [62) or6orders ofmagnitude for oxygen cured andelectron beam cured PCS fibers [18) [29). Further, the activation energy, Ea, decreases withincreasing pyrolysis temperature. It is0.4eV for Si-C-Ofibers produced at1000-1200Candapproaches zero for 1400-1600C.

The dramatic change observed in the electrical behavior of SiC based fibers with increasingpyrolysis temperature isrelated tothe formation orland the organization of free carbon aroundthe SiC crystals. In Si-C-O fibers it occurs above 1200C. Below 1200C, the microstructureofthe fibers iseither amorphous ornanocrystalline, and carbon exists as isolated small BSUs.The material has semiconducting properties (Ea = 0.4 eV). Above 1200C, decompositionoccurs with formation of ~-SiC crystals and free carbon. The increase in electricalconductivity might be related to the removal of the glassy silicon oxycarbide and the formationofacontinuous network ofcarbon around the SiC crystals [18).

o

Eo -22:t:iOl -4.3

-6

Si-Cfiber(Hi-Nicalon)

Si-C-O fiber(NicalonNL202)

2 4 6 81oo0rr, K_l

10 12 14

Figure 17. Electrical conductivity ofSiC-based fibers asa function of test temperature: Si-C-O (Nicalon, NL 202)and Si-C (H i-Nicalon) fibers [50): reproduced with permission.

The electrical conductivity of Si-C fibers with low oxygen content increases sharply at firstwhen the pyrolysis temperature is increased from 1200 to 1400C, and then more slowly.Initially, most of the residual hydrogen is released and the carbon is formed. SUbsequentlythe SiC crystals grow and an intervening carbon network is formed . The thin carbon layer (orsheath) on the surface of Si-C fibers is not alone responsible for the observed gain inelectrical conductivity since its removal by a brief oxidation treatment at 600C does notmarkedly affect its electrical behavior [291.Very little is known about the thermal properties of SiC fibers. A thermal conductivity of12W/m.K and acoefficient of thermal expansion of3.1 x 10-6 IKhave been reported for Si-C-Ofibers (Nicalon NL200) [93).

Chapter 10

10.10 Applications

295

Applications for silicon carbide fibers are discussed in Chapter 12, along with applications forcarbon fibersand ceramic oxide fibers.

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[1) K. Okamura, Ceramic fibers from polymer precursors, Composites, 18[2], 107-120 (1987).[2] J. Lipowitz, J..A.Rabe and R. M. Salinger, Ceramic fibers derived from organosilicon polymers, in Handbook of

Fiber Science and Technology: Vol. //I. High Technology Fibers, Part C, M. Lewinand J. Preston, eds., MarcelDekker, New York, 207-273 (1993).

[3] W.H. Atwell, Polymeric routes tosilicon carbide and silicon nitride fibers, inAdvances inChemisfry Series 224:Silicon-Based Polymer Science, AComprehensive Resource, J.M. Zeigler, F.W. Gordon and M. Fearon, eds.,The Amer. Chern. Soc., Wash.,593-606 (1990).

[4] A.A. Verma and P. Krishna,Polymorphism and polytypism incrystals, Wiley, NY (1966).[5] D. Seyferth, Synthesis ofsome organosilicon polymers and their pyrolytic conversion toceramics, inAdvances

in Chemistry Series 224: Silicon-Based Polymer Science, A Comprehensive Resource, J. M. Zeigler, F. W.Gordon and M. Fearon, eds.,The American Chemical Soc., Washington, 565-59 (1990).

[6] R. M. Laine and F. Babonneau, Preceramic polymer routes tosilicon carbide, Chern. Mater., 5, 260-279 (1993).[7] S. Yajima, Y. Hasegawa, J. Hayashi and M. limura, Synthesis of continuous silicon carbide fibre with high

tensile strength and high Young's modulus. Part 1: Synthesisof polycarbosilane as precursor, J. Mater. Sci.,13,2569-76 (1978) .

[8) Y. Hasegawa, M. Iimura and S. Yajima, Synthesis of continuous silicon carbide fibre. Pt 2: conversion ofpolycarbosilane fibre into silicon carbide fibres, J.Mater. Sci., 15,720(1980).

[9) Y. Hasegawa and K. Okamura, Synthesis of continuous silicon carbide fibre. Part 3: Pyrolysis process ofpolycarbosilane and structure ofthe products, J. Mater. Sci., 18,3633-48 (1983).

[10) Y. Hasegawa and K. Okamura, Synthesis of continuous silicon carbide fibre. Part 4: The structure ofpolycarbosilane asthe precursor,J. Mater. Sci., 21 , 321 -328 (1986).

[11] H. Ichikawa, F. Machino, S. Mitsuno, T. Ishikawa, K. Okamura and Y. Hasegawa, Synthesis of continuoussilicon carbide fibre. Part 5:Factors affecting stability ofpolycarbosilane tooxidation, J. Mater. Sci., 21, 4352-58(1986).

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