hi-nicalontm-ssic fiber oxidation and scale crystallization kinetics

Hi-NicalonTM-S SiC Fiber Oxidation and Scale Crystallization Kinetics

R. S. Hay,‡,† G. E. Fair,‡ R. Bouffioux,§ E. Urban,¶ J. Morrow,k A. Hart,k and M. Wilson††

‡Air Force Research Laboratory, Materials and Manufacturing Directorate, WPAFB, Ohio

§New Mexico Tech. U., Socorro, New Mexico

¶Appalachian State University, Boone, North Carolina

kU. Cincinnati, Cincinnati, Ohio

††Ohio State University, Columbus, Ohio

The oxidation and scale crystallization kinetics of Hi-

NicalonTM

-S SiC fibers were measured after oxidation in dryair between 700° and 1400°C. Scale thickness, composition,

and crystallization were characterized by TEM with EDS,

supplemented by SEM and optical microscopy. TEM was usedto distinguish oxidation kinetics of amorphous and crystalline

scales. Oxidation initially produces an amorphous silica scale

that incorporates some carbon. Growth kinetics of the amor-

phous scale was analyzed using the flat-plate Deal-Grovemodel. The activation energy for parabolic oxidation was

248 kJ/mol. The scales crystallized to tridymite and cristoba-

lite, starting at 1000°C in under 100 h and 1300°C in under

1 h. Crystallization kinetics had activation energy of 514 kJ/molwith a time growth exponent of 1.5. Crystalline silica nucleated

at the scale surface, with more rapid growth parallel to the sur-

face. Crystalline scales cracked from thermal residual stressand phase transformations during cool-down, and during oxida-

tion from tensile hoop growth stress. High growth shear stress

was inferred to cause intense dislocation plasticity near the

crystalline SiO2–SiC interphase. Crystalline scales were thin-ner than amorphous scales, except where growth cracks

allowed much more rapid oxidation.

I. Introduction

THE oxidation of SiC to SiO2 is of concern for both struc-tural applications and for passivation of SiC semicon-

ductors. Literature on SiC oxidation is extensive, and thereare several recent reviews.1–4 The literature on silicon oxida-tion to SiO2 is even more extensive, and in many, but not inall ways, SiC and silicon oxidation are similar. Despite exten-sive research on the topic, much controversy remains aboutmechanisms, and many complicating factors exist. Factorsthat affect SiC oxidation rates include impurities in thefiber,5–7 particularly alkali and alkali earths. These changeoxidation rates, reduce scale viscosity,8–9 and lower tempera-tures for scale crystallization.3,10,11 Moisture has similareffects.3,12,13 However, rare-earth impurities cause rare-earthdisilicates to crystallize that reduce oxidation rates.14 Crystal-lization of SiO2 scales to cristobalite or tridymite reducesoxygen permeability,3,10,15,16 and is a source of confusion ininterpretation of SiC oxidation data that does not distinguishcrystalline and amorphous scales.6,17–19 Crystallization of

SiO2 scales on nominally pure SiC is reported to start attemperatures of 1400°C,6,10,11 at temperatures as low as 700°C,20 and at temperatures in between.2,3,21 Disparity in theseresults illustrates the extreme sensitivity of silica crystalliza-tion to SiC impurities and trace H2O and other gases in theenvironment.3,22,23

Passive oxidation of SiC in most cases is controlled byinterstitial diffusion of O2 molecules through the amorphoussilica scale.2,3 However, at temperatures above 1300°C con-trol by O� or other ionic species may occur.2,6,24,25 Highanisotropy for parabolic oxidation kinetics in SiC has beenextensively documented,7,26–31 and is related to differentamounts of carbon incorporated in the silica scale.3 This isdetermined by whether the SiC surface orientation is carbonor silicon rich. The slower oxidation rates of SiC comparedto silicon are attributed to carbon in the scale.2,3 Much con-troversy exists about the nature, extent, temperature depen-dence, and H2O, and other environmental dependence of thissilicon oxycarbide, or “carbon debris,” layer.2,3,7

For silicon, there are differences in oxidation between flat-plate and cylindrical geometries which are related to the dif-ferent growth stresses that develop from the large volumeexpansion during passive oxidation. Radial and axial com-pressive growth stress and tensile hoop growth stresses arepredicted for oxidation of silicon fibers.32–34 Externallyapplied tensile stress has been shown to increase oxidationrates of silicon,35–37 and recently the same has been demon-strated for SiC fibers.38

Data and analysis of the oxidation and scale crystal-lization kinetics Hi-NicalonTM-S SiC fiber are presented.Hi-NicalonTM-S fiber was chosen because it has near-stoichi-ometric SiC composition (~1 at% oxygen and ~2 at% carbon).39

It is a strong candidate for advanced ceramic matrix compos-ites (CMCs), and has the smoothest surface of currentlyavailable SiC fibers.40 Properties of this fiber are described inseveral publications.39,41–46 Most work suggests that passiveoxidation of SiC fibers reduces their strength.38,46–52 How-ever, recent work suggests that thin silica scales (<100 nm)increase SiC fiber strength,53,54 as might be expected fromthe residual compressive stress in the scales and surface flawhealing. Fiber strength defines CMC tensile strength,55 sofiber strength degradation is critical.

The goal is to thoroughly characterize both the oxidationand scale crystallization kinetics Hi-NicalonTM-S fiber in dryair, and analyze and compare those results with the bestmodels and data currently available. Transmission electronmicroscopy (TEM) is used to characterize oxidation ratesand to distinguish between amorphous and crystalline scales.Scale composition and microstructure evolution are charac-terized and discussed. Some complementary results arepresented elsewhere.56 The determined kinetics will be used

D. Butt—contributing editor

Manuscript No. 29310. Received December 09, 2010; approved April 23, 2011.†Author to whom correspondence should be addressed. e-mail: randall.hay@wpafb.

af.mil

3983

J. Am. Ceram. Soc., 94 [11] 3983–3991 (2011)

DOI: 10.1111/j.1551-2916.2011.04647.x

© 2011 The American Ceramic Society

Journal

as a baseline for further oxidation studies in CMC applica-tion environments that include steam, low pO2, and impuri-ties that are introduced into the CMC during fiber coatingand processing along with those already present in the fiber.The relationships between scale thickness, scale residualstress, scale crystallization, and fiber strength will be dis-cussed in a future article.

II. Experiments

Hi-NicalonTM-S fiber (Lot #281201) sizing burn-off at800°–1000°C caused SiO2 scales to form that were non-uniform in thickness during subsequent oxidation. Contami-nation of the scale by trace inorganic decompositionproducts from the sizing was suspected to cause the nonuni-formity. Consequently, the sizing was removed from 50 in.of Hi-NicalonTM-S fiber tow (~0.025 m2 fiber surface) by dis-solution in 175 cc of boiling distilled deionized water in aPyrex glass beaker for 1 h. Fibers were then dried in an ovenfor 20 min at 120°C. The dissolution and drying steps werethen repeated. The water was collected for inductivelycoupled plasma (ICP) chemical analysis of sizing impurities,and compared to control experiments done without fibers.The desized SiC fibers were also analyzed for impurities byICP analysis. The main sizing impurities were Ca and Na.There was enough Ca in the sizing to yield ~1 monolayer ofCaO on the fiber surface after sizing burn-off. The main fiberimpurities were Cl (620 ppm), S (52 ppm), Ca (45 ppm), Na(35 ppm), and Fe (27 ppm).

The desized fiber tows (500 filaments/tow) were oxidizedin flowing dry air (<10 ppm H2O) in an alumina boat dedi-cated to these experiments, using an alumina muffle tube in aMoSi2 heating-element furnace. Sodium in commercial alu-mina tubes has been shown to contaminate SiO2 during SiCoxidation studies, causing higher oxidation rates and earlierscale crystallization.5 However, this effect gradually disap-peared as the tubes see exposure times >2000 h at 1200°–1400°C.5 Therefore, both the muffle tube and boat werebaked-out at 1540°C for 4 h in dry air prior to use at eachseparate temperature, and the highest temperature studieswere the last ones done. One or two fiber tows 2 in. in lengthwere placed in the boat using tweezers. The filaments in thetows were spread out so that filaments did not cementtogether during oxidation. This allowed easy filament separa-tion for filament mechanical testing,53,56 which will be dis-cussed more thoroughly in a future article. Fiber oxidationwas done at temperatures from 700° to 1400°C for exposuretimes up to 100 h. Fiber tows were heated and cooled at10°C/min. A total of 46 different heat treatments were done.A complete Table with all experiments including scale thick-nesses with standard deviations, number of observations,crystallization fractions, and crystal polymorphs is publishedelsewhere.56

The uniformity of the SiO2 scale was characterized usingreflected light interference fringes. Differences in oxide scalethickness and crystallization extent between filaments closeto the bottom of the alumina boat and those at the top, ifpresent, could be quickly discerned in this manner. Surfacemorphology and cracking of SiO2 were characterized bySEM and optical microscopy. Fiber cross-section TEM speci-mens were prepared from the oxidized fibers by publishedmethods.57,58 TEM sections were ion-milled at 5 kV andexamined using a 200 kV Phillips CM-200 LaB6-filamentTEM and a 300 kV FEI Titan TEM. SiO2 oxidation productthickness, cracking, and crystallization were characterized.When both crystallized and amorphous scales were present,separate thickness measurements were made for each. Thefraction of crystallized scale (f) was determined from thenumber of filaments observed with crystallized scalecompared to the total number of filament scales observed.The crystallized SiO2 phase (tridymite or cristobalite) wasidentified from selected area electron diffraction patterns.

Scale composition in both crystallized and amorphousregions was measured by EDS analysis in TEM using bothspot analysis (Phillips CM-200, 5–50 nm spot size) and X-raydot mapping (FEI Titan).

III. Results and Discussion

(1) General Characterization(A) Amorphous Scale: The SiO2 scales formed after

oxidation at low temperatures and at short times were amor-phous. Thinner scales (<50 nm) were more variable in thick-ness, in some cases, varying with the SiC grain surfaceorientation (Fig. 1). Scales formed at 700°–900°C had surfaceroughness inferred to have been inherited from the underlyingSiC grains. Samples oxidized at 1000°C and higher hadsmooth scale surfaces (Fig. 1); evidently the scale viscositywas low enough at those temperatures for surface tension tosmooth surface roughness. Systematic differences in oxidescale thickness for filaments at the bottom (close to the alu-mina boat) and top of a tow were usually not evident; how-ever, there were slight differences in the shade of the reflectedlight interference color (e.g., blue to blue-violet) in filamentswith very thick amorphous scales (1000°C, 100 h). For thickamorphous scales, similar differences in reflected light interfer-

Fig. 1. TEM images of SiC oxidation scale after 100 h at (a) 700°,(b) 800°, (c) 900°, (d) 1000°C.

3984 Journal of the American Ceramic Society—Hay et al. Vol. 94, No. 11

ence color diagnostic of a small thermal gradient were some-times evident along the 2 in. length of a tow. These interfer-ence color differences were estimated to correspondto thickness differences of 10–20 nm in thick scales, about2%–3% of average scale thickness. Occasionally, “rogue” fila-ments with significantly different scale thicknesses wereobserved by reflected light interference color. Presence of suchfilaments was not correlated with their position in the alumina

boat; they are more likely due to nonuniform impurity distri-bution. Roughly 1%–5% of the filaments fit this category.We did not attempt to quantify the population fraction orscale thicknesses differences of these filaments further.

High resolution TEM images of amorphous scales did nothave obvious graphitic structure in the scale or along theSiO2–SiC interface (Fig. 2). This observation does not ruleout the presence of network carbon in the scale, and theremay be subtle free carbon structures present that requireother characterization techniques.59 Small 5–20 nm intergran-ular pockets of graphitic carbon were present in the SiC justbeneath the SiC–SiO2 interface, even after oxidation at 1200°C (Fig. 3). No cracks were observed in amorphous scales,even in those a micrometer thick.

(B) Crystalline Scale: Crystallization always nucleatedat the scale surface. Growth was more rapid parallel to thesurface than through the scale thickness [Figs. 4(d) and (e)],and crystal spherulites that formed shortly after nucleationwere similar in cross-section. Spherulites were typically thesame size locally, but their size and abundance varied fromfilament to filament, and along long lengths of any particularfilament (Fig. 5). Systematic differences in crystallizationextent for filaments at the bottom (close to the alumina boat)and top of a tow were not evident. As mentioned for amor-phous scales, “rogue” filaments with much more extensivecrystallization were observed by reflected light optical micros-copy. These are not included in Fig. 5.

Compositions of amorphous scales had small and varyingcarbon concentrations, but no carbon was detected in adjacentcrystalline scale (Fig. 6). The nature of the carbon (networkor free) and amount were not determined. Unfortunately,measured carbon concentrations were highly variable, depen-dent on sample orientation with respect to the EDS detector,and subject to beam contamination. Problems with quantifica-tion of the “soft” carbon X-rays by EDS in analytical TEMare well known,60–62 and difficult to ameliorate. Qualitatively,there is less carbon in crystallized silica scales than in adjacentamorphous scale, but the data are not of sufficient quality tobe quantified.

The SiO2 crystallized to various polymorphs of tridymitewith minor a-cristobalite at 1000° and 1400°C, but at 1200°Cthe phase abundance was roughly evenly split between a-cris-tobalite and tridymite. Both tridymite and cristobalite wereextensively twinned; a-cristobalite twinning on (101) was par-ticularly common. Uncracked crystalline silica scales werealways thinner than amorphous scales formed at the sametimes and temperatures, as observed in other studies.10

Thicknesses of crystalline scales varied more than thicknessesof amorphous scales. Presumably variation is caused by thevarying amounts of time spent in the amorphous and crystal-line states, but anisotropy of O2 diffusion in non-cubic SiO2

polymorphs may also be responsible.Crystallized silica was often cracked [Figs. 4(f) and (j)].

Through thickness cracks were often along twin boundaries[Fig. 4(j)]. Debond cracks between SiO2 scale and SiC werealso common [Fig. 4(a)]. Cracked crystalline scales were pre-viously observed for Hi-NicalonTM and Hi-NicalonTM-S, andwere suggested to cause lower fiber strengths.47,63 Unlikeamorphous SiO2, cristobalite and tridymite have much largercoefficients of thermal expansion than SiC. Cracks wereassumed to form from thermal stress during cooling, andfrom volume contraction during the b ? a cristobalite phasetransformation.47,63

Thicker crystalline scales often had growth cracks thatwere parallel with the fiber axis. These cracks had wide aper-tures; they clearly formed during oxidation (Figs. 4 and 7).They are assumed to be caused by hoop tensile growth stres-ses that develop during oxidation in cylindrical geometry,which in turn is caused by the large volume expansion forSiC oxidation.32–34,64–66 There are dimples in SiC under thetips of growth cracks, and the crystalline scale was thicker inareas adjacent to such cracks. When these cracks form, the

Fig. 2. HRTEM images of the SiC–SiO2 interface after oxidation at100 h at 800°C.

Fig. 3. HRTEM images of SiC oxidation scale formed after 3 h at1200°C, showing small pockets of graphite beneath the oxide scale.

November 2011 Oxidation and Scale Crystallization Kinetics of SiC Fiber 3985

scale is no longer passivating, and the growth crack is ashort-circuit path for O2 ingress.64 SEM and optical reflectedlight images (Fig. 7) show growth cracks with wide aperturesin the axial direction, and thinner aperture cracks formedduring cooling that are often, but not always, perpendicularto the growth cracks.

High densities of dislocation debris are evidence of intenseplastic deformation in the 100–200 nm of crystalline scalenext to the SiC–SiO2 interface [Figs. 4(a) and (k)]. Thesedeformed regions tended to be in crystalline scales that werethinner than average, indicating that they had grown a longtime in the crystalline state. The outer scale without disloca-tions presumably formed in an amorphous state, deformedunder growth stresses by viscous flow, and then crystallized.Inner scale that grew in a crystalline state must deform bydislocation plasticity. Deformation twinning may also relaxgrowth stress.

Crystalline scales formed at 1300°C and higher often con-tained porosity (Fig. 8). Such porosity was not uniformlydistributed. In some cases a pore layer was in the middle ofthe scale and there was intense dislocation plasticity beneaththe pore layer, so this layer may mark the SiO2–SiC interfaceat the time the scale crystallized.

(2) Crystallization KineticsKolmogorov–Johnson–Mehl–Avrami (KJMA) analysis wasused to analyze silica crystallization kinetics (Fig. 9)67,68:

f ¼ 1� exp½� Kt n� ð1ÞK ¼ K0exp½� Q=RT� ð2Þ

where f is the fraction crystallized, K is a rate constant, t istime, n is the time growth exponent, K0 is a pre-exponentialfactor, and Q an activation energy for growth, and RT hasthe usual meaning in an Arrhenius expression. The analysisgives a parameter best fit for t, T, and f of:

Q ¼ 514 kJ/mol ð3aÞ

K0 ¼ 2� 1012 ð3bÞ

n ¼ 1:5 ð3cÞ

To our knowledge, the only KJMA analysis for nominallypure SiO2 finds Q = 408 kJ/mol, K0 = 1 9 106, and n = 1.63in flowing dry air, with unspecified furnace materials.69 Thisis a much lower K0 than our results [3]. Many factors mayinfluence crystallization kinetics. Crystallization kinetics isextremely sensitive to impurity levels in vitreous SiO2.

70,71

Impurities inherited from the Hi-NicalonTM-S SiC fiber areexpected to be in SiO2 scales, and there is also some carbonincorporated in the scale that is rejected during crystalliza-tion. Furthermore, the SiC cylindrical fiber substrate causestensile growth stresses not present for planar surfaces.64,65

Fig. 4. TEM images of SiC oxidation scale formed at 1100° and 1200°C. (a) 1100°C, 100 h. Debond crack, with intense dislocation plasticity ina-cristobalite adjacent to the crack. Twinning in the scale. (b) 1100°C, 100 h. Growth crack in tridymite scale. (c) 1100°C, 100 h. Tridymite scale,with incipient debond crack on right hand side, and some porosity along a grain boundary. (d) 1200°C, 1 h. Partially crystallized scale. (e) 1200°C, 3 h. Partially crystallized scale. (f) 1200°C, 3 h. Crystallized scale with two through thickness cracks. (g) 1200°C, 3 h. a-Cristobalite scale withtwinning. (h) 1200°C, 10 h. Crystallized scale with two grain boundaries. Grain on left hand side has intense dislocation plasticity in SiO2

adjacent to the crack. (i) 1200°C, 10 h. Amorphous scale. (j) 1200°C, 10 h. a-cristobalite scale with through thickness crack along twinboundary. (k) 1200°C, 10 h. Tridymite scale with twinning and some dislocation plasticity. (l) 1200°C, 30 h. Growth crack in crystalline scale.


These tensile stresses may influence cristobalite or tridymitenucleation.

Crystallization rates for 100 h at 1000°C and 10 h at1100°C are comparable to some rates found in the litera-ture, albeit with more tridymite formation than cristoba-lite.3,72 Presence of impurities, particularly sodium, is oftensuggested to promote tridymite crystallization over that ofcristobalite,3,5,11,73 but the literature is inconsistent on this;some studies are ambiguous or suggest that sodium promotescristobalite crystallization.74–76 Both higher21 and muchlower20 crystallization temperatures were reported for oxidescales formed on a single crystal and CVD (b) SiC. Neitherstudy reported the type of furnace tube that was used, orother details helpful in assessment of the possible role of

Fig. 5. Optical images, reflected light, of fibers oxidized for 100 hat 1000°C. The variation in crystal nucleation and growth rates inthe oxide scale are evident.

Fig. 6. TEM image of a partially crystallized SiO2 scale formedafter 10 h at 1100°C. Crystallized scale is lenticular in the lower partof the image. A summation of four EDX spectra from (top) the SiCfiber, (middle) the amorphous scale, and (bottom) the crystallizedscale is shown to the right. Carbon spectra were consistently strongerin the amorphous scale than in the crystallized scale.

Fig. 7. Optical micrograph (reflected light) and the correspondingarea imaged by SEM. Cracks with wide aperture, with tracesrunning in the axial direction are growth cracks. Cracks with smalleraperture formed during cooling. Bright areas in the opticalmicrographs are interference fringes from debond cracks between thescale and the SiC.

Fig. 8. TEM images showing oxide scale formed after 1 h at 1300°C. Images (a) and (b) show porosity in the middle of the scale,extensive twinning in the outer cristobalite scale, and evidence ofintense dislocation plasticity in the inner scale. Images in (c) and (d)show healed growth cracks, with relict porosity.


impurities. Usually, 1200° or 1250°C is cited as the tempera-ture at which scales crystallize,3,64 without specifying time orfraction crystallized. The only comparable study on Hi-NicalonTM-S reports no crystallization, as measured byX-ray diffraction, in scales formed after 100 h at 1000°C indry air, but complete crystallization to cristobalite in 10 h at1300°C.47 This study also does not report the type of furnacetube or sample holder that was used.

Best fits of the crystallization fraction data to the KJMAmodel [1,2] converged strongly on a growth exponent n of1.5. This is diagnostic of three-dimensional growth from site-saturated nucleation.68 Microstructural evidence (Fig. 5) forsite-saturated nucleation was discussed. Locally, spherulitesizes were similar, suggesting synchronous nucleation consis-tent with site-saturated nucleation. But spherulite size andabundance varied from filament to filament, perhaps as aconsequence of filament to filament inhomogeneity in impu-rity levels. Change in the number of nucleation sites wouldnot affect the growth exponent n, but would affect K0. Three-dimensional, rather than two-dimensional, growth kineticsrequires that spherulites impinge each other before they growthrough the full thickness of the scale. This requires nucleispacings to be smaller than the film thickness. Maximumamorphous scale thickness was ~1 lm, and nuclei spacingsare clearly much larger than that (Fig. 5). However, it is notclear how the high crystal growth rate anisotropy wouldaffect the time growth exponent.

(3) Oxidation Kinetics(A) Models: Fiber oxidation kinetics does not deviate

significantly from those for flat plate geometry until the oxi-dation products for 12 lm diameter fibers are severalmicrons thick,56 as found previously for NicalonTM fiber77

and silicon cylinders.33,78 Since the SiO2 scales generally crys-tallized when they were greater than 1 lm thick, the flat-plate geometry is adequate for analysis. The thickness of theSiO2 scale (x) described by Deal–Grove kinetics for the flatplate geometry is:

dx=dt ¼ B=ðAþ 2xÞ ð4Þ

A ¼ A0 exp½�Qa=RT� ð5Þ

B ¼ B0 exp½�Qb=RT� ð6Þwhere A0 and B0 are pre-exponential constants, Qa and Qb

are activation energies. B is the parabolic rate constant andB/A is the linear rate constant. The quantity A is the SiO2

thickness at which crossover from linear to parabolic kineticsoccurs. For an initial SiO2 thickness of xi, the solution to[4] is:

x ¼ 1=2A f½1þ ðtþ sÞ=ðA2=4BÞ�1=2 � 1g ð7Þ

s ¼ ðx2i þ AxiÞ=B ð8Þ

where τ is a time shift that corrects for the presence of aninitial oxide layer. For long times, [7] becomes the simpleexpression for parabolic oxidation kinetics:

x2 ¼ Bt ð9Þ

The initial SiO2 thickness xi is often formed during heat-up. Its calculation is described elsewhere.56

(B) AnalysisThe scale thickness—time—temperature data were firstanalyzed assuming parabolic kinetics [6, 9] for amorphousscales only. Two fitting methods were used. The usualmethod of finding B for each temperature from x and t using[7] with linear regression to fit B0 and Qb to a ln(B) versus1/T plot was employed. The B0 and Qb parameters were alsodirectly fit to [6] and [9] for all data without calculation ofseparate B values for each temperature. This method hassome data weighting advantages that are discussed next forDeal–Grove analysis.

The scale thickness—time—temperature data were ana-lyzed using Deal–Grove kinetics [4–8] for amorphous scalesonly. Instead of calculating A and B values for each tempera-ture and using those values to fit parameters in the Arrheniusexpressions [5,6], analysis was done by fitting the four vari-ables A0, B0, Qa, and Qb to log(x), t, and T data using [6], assuggested recently.6 This method weights all data equally forparameter fitting, and avoids problems associated with exces-sive influence of data at temperatures for which there are fewmeasurements. Thickness measurements spanned almostthree orders of magnitude. Fitting to log(x), rather than x,weights the fit by proportional rather than absolute magni-tude, and avoids excessive influence of large x values on theparameter fit. Using numerical methods discussed else-where,56 oxidation times (t) were corrected by τ [8] for theextra time at temperature required to form the oxide thick-ness (xi) that actually formed during heat-up and cool-down.

The best fit for Deal–Grove parameters for oxidationkinetics for Hi-NicalonTM-S fiber are shown in a scale thick-ness—time plot (Fig. 10). A best fit to the data was:

A0 ¼ 4:7� 10�4m ð10aÞ

Qa ¼ 110 kJ/mol ð10bÞ

B0 ¼ 1:2� 10�8m2=s ð10cÞ

QbðparabolicÞ ¼ 248 kJ/mol ð10dÞ

Qb=aðlinearÞ ¼ 138 kJ/mol ð10eÞ

The fit had strong convergence to the Qb value (±5 kJ/mol), but was less sensitive to the other parameters, particu-larly Qa and A0. Since most data points are described by par-abolic kinetics, this is to be expected. The best fits forparabolic oxidation kinetics are also plotted; the valuesfound for B0 and Qb were 2.5 9 10�8 m2/s and 258 kJ/mol,respectively. The difference in activation energy with [10d] is

Fig. 9. KJMA analysis of crystallization of the SiO2 scale.


not significant, and the 29 higher pre-exponential factor[10c] is presumed to be an artifact of including linear regimedata in the parabolic analysis. Superimposed on this plot arelines of constant crystallization fraction (f) calculated fromthe KJMA analysis using [1–3].

(C) Discussion: Above 1100°C the calculated para-bolic oxidation rate constants (B) of Hi-NicalonTM-S SiCfiber are higher than most rate constants reported for SiC,including the data for other SiC fibers. There is not as muchoxidation data reported below 1100°C. Here, as expected, Bfor Hi-NicalonTM-S oxidation is lower than B for less-stoichiometric SiC fibers. Many factors may contribute tothe relatively high oxidation rates of Hi-NicalonTM-S:

(1) Data for high temperatures are from crystallized orpartially crystallized scales. Crystallization reduces oxidationrates.3,10,18 Reduced oxidation rates are prominent for SiCfiber oxidation data above 1200°C, suggesting that crystalli-zation may be common at those temperatures. Many of thesestudies report crystallization of the scale,56 but with insuffi-cient t–T–f data to characterize crystallization kinetics. (2)The impurities in Hi-NicalonTM-S may cause higher oxida-tion rates than those for single crystal oxidation.3 Greaterthan 20 ppm Cl, S, Ca, Na, and Fe were present in Hi-NicalonTM-S. The presence of “rogue” filaments with signifi-cantly different oxidation and crystallization rates suggeststhat these impurities may not be uniformly distributed. Con-versely, the higher purity of Hi-NicalonTM-S may causeslower oxidation than that for older, less stoichiometricfibers. (3) Texture and/or SiC polytype dependence may alsobe a factor, but the latter is controversial.1–2 (4) Tensile stressenhances O2 permeability in amorphous SiO2.

35–38 Geometri-cal effects on growth stress that cause tensile hoop stress dur-ing oxidation of cylindrical substrates may contribute tohigher oxidation rates. (5) Furnace bake-out times may nothave been long enough to eliminate contamination, so ouroxidation rates may be higher than those done in fused silicaor quartz tube furnaces.5 Oxidation at higher temperaturescould cause correspondingly higher sodium contaminationfrom the alumina tube and boat, leading to systematicallyhigher oxidation rates at higher temperatures with an appar-ent higher Qb.

6,79,80 However, as discussed earlier, the bake-out temperature of 1540°C should be sufficient to removeNa2O(g) from alumina that would otherwise contaminateSiC oxidation experiments done at 700°–1300°C.

Literature values for activation energies for parabolic oxi-dation kinetics (Qb) vary widely.56 Reviews of oxidation datafind that Qb can vary from 80 to 612 kJ/mol,3 from 134 to498 kJ/mol,81 or from 69 to 498 kJ/mol,82 depending on thematerial studied and the methods used. Qb can vary from~100 to ~270 kJ/mol°K for C-face (fast oxidation) andSi-face (slow oxidation) surfaces, respectively, in dry air.3

For fibers, our Qb of 248 kJ/mol is the highest yet reported,although it is very similar to one reported for Hi-Nicalon,63

and 30% larger than one reported for TyrannoTM SA.54 Athorough comparison of our Qb and Qb/a values is presentedelsewhere56; our values are close to the median valuesreported in a variety of SiC oxidation studies. For almost allother fiber data, crystallization was reported at high tempera-tures, and is probably a significant source of discrepancy.56

Most of the data are also for carbon-rich fibers measured byweight gain, which is another source of discrepancy.

IV. Summary and Conclusions

The oxidation kinetics for formation of amorphous scaleobeys a Deal–Grove relationship, with activation energy forparabolic oxidation (Qb) of 248 kJ/mol and a pre-exponentialfactor (B0) of 1.2 9 10-8 m2/s. The activation energy for lin-ear oxidation (Qb/a) was 138 kJ/mol and the pre-exponentialfactor (A0) was 4.7 9 10�4 m. The measured parabolic oxi-dation kinetics and activation energy (Qb) are higher thanthose observed for many SiC fiber oxidation studies, particu-larly for higher temperatures. Many of these studies do notclearly distinguish oxidation rates for crystalline and amor-phous scales, so lower O2 permeability after crystallization athigh temperatures is at least a partial explanation for thehigh SiC fiber Qb.

The KJMA crystallization kinetics of the amorphous SiO2

scale had an activation energy Q of 514 kJ/mol, a timegrowth exponent n of 1.5, and a pre-exponential factor K0 of2 9 1012. These kinetic parameters are clearly extremely sen-sitive to SiC impurities and the oxidation environment. Thescales crystallized to cristobalite and tridymite, with tridymitepredominating at most temperatures. Significant crystalliza-tion began within 100 h at 1000°C and 1 h at 1300°C, andwas complete within 80 h at 1100°C, 30 h at 1200°C and 1 hat 1300°C. Crystallization always nucleated at the SiO2 sur-face, and growth through the scale thickness was muchslower than growth parallel to the surface. Nucleation wasinterpreted to be site-saturated, but rates varied significantlyfrom filament to filament.

Amorphous scales were smooth and uncracked. Crystallinescales were often cracked, with both through thickness anddebond cracks from tensile thermal stress driven by the largerthermal expansion of cristobalite and tridymite than SiC, andfrom polymorphic phase transformations in cristobalite andtridymite. Thicker crystalline scales had growth cracks thatare interpreted to form under tensile hoop growth stressespresent for the cylindrical fiber geometry, driven by the largevolume expansion during SiC oxidation. When these cracksform, the scale is no longer passivating; enhanced oxidationtakes place underneath these short-circuits for O2 diffusion.High dislocation densities in crystalline SiO2 near the SiO2–SiC interface are evidence of high SiO2 growth shear stresses.

Evidence of carbon in amorphous SiO2 scales was foundby analytical TEM. Unfortunately the soft X-rays used forcarbon EDS detection, carbon contamination, and the prox-imity of carbon in SiC do not allow quantification of carbonconcentration and determination of whether it is network orfree carbon. High resolution TEM did not show obvious gra-phitic structures in the oxide scale. These observations are,however, qualitatively consistent with a large number ofobservations of carbon-rich layers that vary with SiC surfaceorientation. Future work will explore relationships betweenthe oxide scale thickness, microstructure, and the strength ofthe oxidized fibers.

Fig. 10. Deal-Grove kinetics (thick lines) for Hi-NicalonTM-S fiberoxidation for amorphous SiO2 scale formation. Individual datapoints are shown as dots. The thin solid and dotted lines show thebest fits for parabolic oxidation kinetics only, calculated by twodifferent methods. The KJMA analysis for scale crystallizationkinetics (grey) is superimposed.


References

1C. Raynaud, “Silica Films on Silicon Carbide: A Review of ElectricalProperties and Device Applications,” J. Non. Cryst. Sol., 280, 1–31 (2001).

2I. Vickridge, J. Ganem, Y. Hoshino, and I. Trimaille, “Growth of SiO2 onSiC by Dry Thermal Oxidation: Mechanisms,” J. Phys. D: Appl. Phys., 40,6254 (2007).

3V. Presser and K. G. Nickel, “Silica on Silicon Carbide,” Crit. Rev. SolidState Mater. Sci., 33, 1–99 (2008).

4T. Narushima, M. Kato, S. Murase, C. Ouchi, and Y. Iguchi, “Oxidationof Silicon and Silicon Carbide in Ozone-Containing Atmospheres at 973K,”J. Am. Ceram. Soc., 85, 2049–55 (2002).

5E. Opila, “Influence of Alumina Reaction Tube Impurities on the Oxida-tion of Chemically-Vapor-Deposited Silicon Carbide,” J. Am. Ceram. Soc., 78,1107–10 (1995).

6C. E. Ramberg and W. L. Worrell, “Oxygen Transport in Silica at HighTemperatures: Implications of Oxidation Kinetics,” J. Am. Ceram. Soc., 84,2607–16 (2001).

7C. E. Ramberg, G. Cruciani, K. E. Spear, R. E. Tressler, and C. F. Ram-berg, “Passive-Oxidation Kinetics of High-Purity Silicon Carbide from 800 to1100 C,” J. Am. Ceram. Soc., 79, 2897–911 (1996).

8R. H. Doremus, “Viscosity of Silica,” J. Appl. Phys., 92, 7619–29 (2002).9G. Pezzotti and G. S. Painter, “Mechanisms of Dopant-Induced Changes

in Intergranular SiO2 Viscosity in Polycrystalline Silicon Nitride,” J. Am.Ceram. Soc., 85, 91–6 (2002).

10V. Presser, A. Loges, Y. Hemberger, and K. G. Nickel, “MicrostructuralEvolution of Silica on Single-Crystal Silicon Carbide. Part I: Devitrificationand Oxidation Rates,” J. Am. Ceram. Soc., 92, 724–31 (2009).

11V. Presser, A. Loges, Y. Hemberger, and K. G. Nickel, “MicrostructuralEvolution of Silica on Single-Crystal Silicon Carbide. Part II: Influence ofImpurities and Defects,” J. Am. Ceram. Soc., 92, 1796–805 (2009).

12E. J. Opila and R. E. Hann, “Paralinear Oxidation of CVD SiC in WaterVapor,” J. Am. Ceram. Soc., 80, 197–205 (1997).

13E. J. Opila, “Oxidation Kinetics of Chemically Vapor-Deposited SiliconCarbide in Wet Oxygen,” J. Am. Ceram. Soc., 77, 730–6 (1994).

14F. Rodrı́guez-Rojas, A. L. Ortiz, F. Guiberteau, and M. Nygren, “Oxida-tion behaviour of pressureless liquid-phase-sintered a-SiC with additions of5Al2O3 + 3RE2O3 (RE = La, Nd, Y, Er, Tm, or Yb),” J. Eur. Ceram. Soc.,30, 3209–17 (2010).

15J. Rodriguez-Viejo, F. Sibieude, M. T. Clavaguera-Mora, and C. Monty,“18O diffusion through amorphous SiO2 and cristobalite,” Appl. Phys. Lett.,63, 1906–8 (1993).

16L. J. T. Ogbuji and M. Singh, “High-Temperature Oxidation Behavior ofReaction-Formed Silicon Carbide Ceramics,” J. Mater. Res., 10, 3232–40 (1995).

17L. U. J. T. Ogbuji, “Effect of Oxide Devitrification on Oxidation Kineticsof SiC,” J. Am. Ceram. Soc., 80, 1544–50 (1997).

18L. U. J. T. Ogbuji, “Subparabolic Oxidation Behavior of Silicon Carbideat 1300C,” J. Electrochem. Soc., 145, 2876–82 (1998).

19L. Filipuzzi and R. Naslain, “Oxidation Kinetics of SiC‐ Based CeramicFibers;” pp. 35–46. in Advanced Structural Inorganic Composites, Edited byP. Vincenzini. Elsevier Science Publishers, Amsterdam, 1991.

20M. J. F. Guinel and M. G. Norton, “Oxidation of Silicon Carbide and theFormation of Silica Polymorphs,” J. Mater. Res., 21, 2550–63 (2006).

21F. Sibieude, J. Rodrı́guez, and M. T. Clavaguera-Mora, “Kinetics andCrystallization Studies by In Situ X-Ray Diffraction of the Oxidation ofChemically Vapour Deposited SiC,” Thin Solid Films, 204, 217–7 (1991).

22D. J. Choi and W. D. Scott, “Devitrification and Delayed Crazing of SiO2

on Single Crystal Silicon and Chemically Vapor Deposited Silicon Nitride,”J. Am. Ceram. Soc., 70, C269–72 (1987).

23W. Zhou, H. Fu, L. Zhang, X. Sun, S. She, and J. Ma, “Effects of Impuri-ties and Manufacturing Methods on the Devitrification of Silica Fibers,”J. Am. Ceram. Soc., 74, 1125–8 (1991).

24Z. Zheng, R. E. Tressler, and K. E. Spear, “Oxidation of Single-CrystalSilicon Carbide, Part I. Experimental Studies,” J. Electrochem. Soc., 137, 854–8 (1990).

25J. Li, T. J. Lenosky, C. J. Forst, and S. Yip, “Thermochemical andMechanical Stabilities of the Oxide Scale of ZrB2 + SiC and Oxygen TransportMechanisms,” J. Am. Ceram. Soc., 91, 1475–80 (2008).

26Y. Hijikata, T. Yamamoto, H. Yaguchi, and S. Yoshida, “Model Calcula-tion of SiC Oxidation Rates in the Thin Oxide Regime,” Mater. Sci. Forum,600-603, 663– (2009).

27Y. Song, S. Dhar, L. C. Feldman, G. Chung, and J. R. Williams, “Modi-fied Deal Grove Model for the Thermal Oxidation of Silicon Carbide,”J. Appl. Phys., 95, 4953–7 (2004).

28J. A. Costello and R. A. Tressler, “Oxidation Kinetics of Silicon CarbideCrystals and Ceramics: I. In Dry Oxygen,” J. Am. Ceram. Soc., 69, 674–81 (1986).

29K. Christiansen and R. Helbig, “Anisotropic Oxidation of 6H-SiC,”J. Appl. Phys., 79, 3276–81 (1996).

30J. N. Shenoy, M. K. Das, J. A. Cooper, and M. R. Melloch, “Effect ofSubstrate Orientation and Crystal Anisotropy on the Thermally OxidizedSiO2/SiC Interface,” J. Appl. Phys., 79, 3042–5 (1996).

31T. Goto and H. Homma, “High-Temperature Active/Passive Oxidationand Bubble Formation of CVD SiC in O2 and CO2 Atmospheres,” J. Eur.Ceram. Soc., 22, 2749–56 (2002).

32T. J. Delph, “Intrinsic Strain in SiO2 Thin Films,” J. Appl. Phys., 83,786–92 (1998).

33D.-B. Kao, J. P. McVittie, W. D. Nix, and K. C. Saraswat, “Two-Dimen-sional Thermal Oxidation of Silicon - II. Modeling Stress Effects in WetOxides,” IEEE Trans. Electron. Dev., 35, 25–37 (1988).

34C. S. Rafferty, L. Borucki, and R. W. Dutton, “Plastic Flow During theThermal Oxidation of Silicon,” Appl. Phys. Lett., 54, 1516–8 (1989).

35A. Mihalyi, R. J. Jaccodine, and T. J. Delph, “Stress Effects in the Oxida-tion of Planar Silicon Substrates,” Appl. Phys. Lett., 74, 1981–3 (1999).

36J.-Y. Yen and J.-G. Hwu, “Enhancement of Silicon Oxidation Rate Dueto Tensile Mechanical Stress,” Appl. Phys. Lett., 76, 1834–5 (2000).

37J.-Y. Yen and J.-G. Hwu, “Stress Effect on the Kinetics of Silicon Ther-mal Oxidation,” J. Appl. Phys., 89, 3027–2 (2001).

38W. Gauthier, F. Pailler, J. Lamon, and R. Pailler, “Oxidation of SiliconCarbide Fibers During Static Fatigue in Air at Intermediate Temperatures,”J. Am. Ceram. Soc., 92, 2067–73 (2009).

39S. M. Dong, et al., “Characterization of Nearly Stoichiometric SiCFibres,” J. Mater. Sci., 36, 2371–81 (2001).

40T. Hinoki, L. L. Snead, E. Lara-Curzio, J. Park, and A. Kohyama, “Effectof Fiber/Matrix Interfacial Properties on Mechanical Properties of Unidirec-tional Crystalline Silicon Carbide Composites,” Ceram. Eng. Sci. Proc., 23,511–8 (2002).

41C. Sauder and J. Lamon, “Tensile Creep Behavior of SiC-Based FibersWith a Low Oxygen Content,” J. Am. Ceram. Soc., 90, 1146–56, doi:10.1111/j.1551-2916.2007.01535.x (2007).

42A. R. Bunsell and A. A Piant, “Review of the Development of ThreeGenerations of Small Diameter Silicon Carbide Fibres,” J. Mater. Sci., 41,823–39 (2006).

43T. Ishikawa, “Advances in Inorganic Fibers,” Adv. Polym. Sci., 178,109–44 (2005).

44J. J. Sha, T. Nozawa, J. S. Park, Y. Katoh, and A. Kohyaman, “Effect ofHeat-Treatment on the Tensile Strength and Creep Resistance of AdvancedSiC Fibers,” J. Nucl. Mater., 329–333 592–6 (2004).

45T. Tanaka, S. Shibayama, M. Takeda, and A. Yokoyama, “Recent Pro-gress of Hi-Nicalon Type S Development,” Ceram. Eng. Sci. Proc., 24, 217–3(2003).

46T. Shimoo, H. Takeuchi, and K. Okamura, “Oxidation Kinetics andMechanical Property of Stoichiometric SiC Fibers (Hi-Nicalon-S),” J. Ceram.Soc. Jpn., 108, 1096–2 (2000).

47M. Takeda, A. Urano, J. Sakamoto, and Y. Imai, “Microstructure andOxidation Behavior of Silicon Carbide Fibers Derived from Polycarbosilane,”J. Am. Ceram. Soc., 83, 1171–6 (2000).

48H.-E. Kim and A. J. Moorhead, “Strength of Nicalon Silicon CarbideFibers Exposed to High-Temperature Gaseous Environments,” J. Am. Ceram.Soc., 74 666–9 (1991).

49J. J. Brennan, “Interfacial Characterization of a Slurry-Cast Melt-Infil-trated SiC/SiC Ceramic-Matrix Composite,” Acta Mater., 48, 4619–28(2000).

50Y. Gogotsi and M. Yoshimura, “Oxidation and Properties Degradation ofSiC Fibres Below 850 C,” J. Mater. Sci. Lett., 13, 680–3 (1994).

51E. Lara-Curzio, “Stress-Rupture of Nicalon/SiC Continuous Fiber Cera-mic Composites in Air at 950 C,” J. Am. Ceram. Soc., 80, 3268–72 (1997).

52E. Lara-Curzio, “Oxidation Induced Stress-Rupture of Fiber Bundles,”J. Eng. Mater. Tech., 120, 105–9 (1998).

53R. S. Hay, G. E. Fair, E. Urban, J. Morrow, J. Somerson, and M. Wilson,“Oxidation Kinetics and Strength Versus Scale Thickness for Hi-Nical-onTM-S Fiber,” Ceramic Transactions, 215, Singh, Zhu, Zhou, and SIngh(eds) (2010).

54P. Mogilevsky, E. E. Boakye, R. S. Hay, and R. J. Kerans, “MonaziteCoatings on SiC Fibers II: Oxidation Protection,” J. Am. Ceram. Soc., 89,3481–90 (2006).

55W. A. Curtin, B. K. Ahn, and N. Takeda, “Modeling Brittle and ToughStress-Strain Behavior in Unidirectional Ceramic Matrix Composites,” ActaMater., 46, 3409–20 (1998).

56R. S. Hay, G. E. Fair, R. Bouffioux, E. Urban, J. Morrow, A. Hart, andM. Wilson, “Relationships between Fiber Strength, Passive oxidation andScale Crystallization Kinetics of Hi-NicalonTM-S SiC Fibers,” Ceram. Eng.Sci. Proc., 32 (2011) in press.

57R. S. Hay, J. R. Welch, and M. K. Cinibulk, “TEM Specimen Preparationand Characterization of Ceramic Coatings on Fiber Tows,” Thin Solid Films,308-309, 389–92 (1997).

58M. K. Cinibulk, J. R. Welch, and R. S. Hay, “Preparation of ThinSections of Coated Fibers for Characterization by Transmission ElectronMicroscopy,” J. Am. Ceram. Soc., 79, 2481–4 (1996).

59H.-J. Kleebe, C. Turquat, and G. D. Soraru, “Phase Separation in anSiCO Glass Studied by Transmission Electron Microscopy and ElectronEnergy-Loss Spectroscopy,” J. Am. Ceram. Soc., 84, 1073–80 (2001).

60A. Armigliato and R. Rosa, “X-Ray Microanalysis Combined with MonteCarlo Simulation for the Analysis of Layered Thin Films: The Case of CarbonContamination,” Micros. Microanal, 15, 99–105 (2009).

61T. C. Isabell, P. E. Fischione, C. O’Keefe, M. U. Guruz, and V. P.Dravid, “Plasma Cleaning and Its Applications for Electron Microscopy,”Micros. Microanal., 5, 126–35 (1999).

62P. Rolland, V. L. Carlino, and R. Vane, “Improved Carbon Analysis withEvactron Plasma Cleaning,” Micros. Microanal., 7, 964–5 (2004).

63T. Shimoo, T. Hayatsu, M. Takeda, H. Ichikawa, T. Seguchi, andK. Okamura, “Mechanism of Oxidation of Low-Oxygen SiC Fiber Preparedby Electron Radiation Curing Method,” J. Ceram. Soc. Jpn., 102, 617–22(1994).

64G. Chollon, R. Pallier, R. Naslain, F. Laanani, M. Monthioux, andP. Olry, “Thermal Stability of a PCS-Derived SiC Fibre with a Low OxygenContent (Hi-Nicalon),” J. Mater. Sci., 32, 327–47 (1997).

65C. H. Hsueh and A. G. Evans, “Oxidation Induced Stresses and SomeEffects on the Behavior of Oxide Films,” J. Appl. Phys., 54, 6672–86 (1983).


66D. K. Brown, S. M. Hu, and J. M. Morrissey, “Flaws in Sidewall OxidesGrown on Polysilicon Gate,” J. Electrochem. Soc., 129, 1084–9 (1982).

67F. Liu, F. Sommer, and E. J. Mittemeijer, “Analysis of the Kinetics ofPhase Transformations; Roles of Nucleation Index and Temperature Depen-dent Site Saturation, and Recipes for the Extraction of Kinetic Parameters,”J. Mater. Sci., 42, 573–87 (2007).

68F. Liu, F. Sommer, C. Bos, and E. J. Mittemeijer, “Analysis of Solid StatePhase Transformation Kinetics: Models and Recipes,” Int. Mater. Rev., 52,193–212 (2007).

69X. Li, X. Yin, L. Zhang, and S. He, “The Devitrification Kinetics of SilicaPowder Heat-Treated in Different Conditions,” J. Non. Cryst. Sol., 354,3254–9 (2008).

70M. L. F. Nascimento and E. D. Zanotto, “Mechanisms and Dynamics ofCrystal Growth, Viscous Flow, and Self-Diffusion in Silica Glass,” Phys. Rev.B, 73, 024209 (2006).

71F. Nascimento and E. D. Zanotto, “Diffusion Processes in Vitreous SilicaRevisited,” Phys. Chem. Glasses, 48, 201–17 (2007).

72E. Gugel, H. W. Hennicke, and P. Schuster, “Zur Bildung der SiO2-Schicht auf SiC,” Berichte der Deutschen Keramischen Gesellschaft, 46, 481(1969).

73R. B. Sosman, The Properties of Silica. 97–106 The Chemical CatalogCo., NY, 1927.

74M. Muroya and S. Kondo, “The Effect of Impurities Upon the PhysicalProperties of SIlica Gel by the Heat-Treatment,” Bull. Chem. Soc. Japan, 43,3453–6 (1970).

75J. Phalippou, M. Prassas, and J. Zarzycki, “Crystallization of Gelsand Glasses Made from Hot-Pressed Gels,” J. Non. Cryst. Sol., 48, 17–30 (1982).

76Z. Congshan, J. Phalippou, and J. Zarzycki, “Influence of Trace AlkaliIons on the Crystallization Behaviour of SIlica Gels,” J. Non. Cryst. Sol., 82,321–8 (1986).

77Y. T. Zhu, S. T. Taylor, M. G. Stout, D. P. Butt, and T. C. Lowe,“Kinetics of Thermal, Passive Oxidation of Nicalon Fibers,” J. Am. Ceram.Soc., 81, 655–60 (1998).

78L. O. Wilson and R. B. Marcus, “Oxidation of Curved Silicon Surfaces,”J. Electrochem. Soc., 134, 481–90 (1987).

79D. S. Fox, “Oxidation Behavior of Chemically-Vapor-Deposited SiliconCarbide and Silicon Nitride from 1200 to 1600 C,” J. Am. Ceram. Soc., 81,945–50 (1998).

80U. J. T. Ogbuji and E. J. A Opila, “Comparison of the Oxidation Kineticsof SiC and Si3N4,” J. Electrochem. Soc., 142, 925–30 (1995).

81K. L. Luthra, “Some New Perspectives on Oxidation of Silicon Carbideand Silicon Nitride,” J. Am. Ceram. Soc., 74, 1095–103 (1991).

82T. Shimoo, F. Toyoda, and K. Okamura, “Oxidation Kinetics ofLow-Oxygen Silicon Carbide Fiber,” J. Mater. Sci., 35, 3301–5 (2000). h


hi-nicalontm-ssic fiber oxidation and scale crystallization kinetics

Documents