Oxidation behavior of SiC ceramics synthesized from processed cellulosic bio-precursor

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Oxidation behavior of SiC ceramics synthesized from processed cellulosic bio-precursor
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  Oxidationbehaviorof    SiCceramicssynthesizedfromprocessedcellulosicbio-precursor AnweshaMaity a ,DipulKalita b ,NijhumaKayal a ,TridipGoswami b ,OmprakashChakrabarti a, *,   ParuchuriGangadharRao b a Central   Glass   and    Ceramic    Research    Institute,   CSIR,   Kolkata   700032,   West     Bengal,    India b  North   East     Institute   of    Science   and    Technology,   CSIR,    Jorhat,    Assam,    India Received   12   August   2011;   received   in   revised   form   16   February   2012;   accepted   18   February   2012Available   online   25   February   2012 Abstract Oxidation   behavior   of    Si/SiC   ceramic   composite   synthesized   from   processed   cellulosic   bio-precursor   was   studied   in   dry   air   over   thetemperature   range   1200–1350   8 C.   The   material   was   synthesized   from   processed   bio-precursors   (bleached   bamboo   kraft   pulp   in   the   form   of    flatboard   of    bulk    density   0.58   g   cm À 3 )   and   had   a   bulk    density   of    2.66   g   cm À 3 ,   porosity   of    0.6   vol%   and   contents   of    Si   and   SiC   phases   of    39.1%   and60.3%   (v/v)   respectively.   The   process   of    oxidation   could   be   described   closely   by   a   parabolic   oxidation   equation.   An   activation   energy   of    141.4   kJ/ mol   was   obtained.   Both   the   SiC   and   Si   phases   oxidized   and   the   oxidation   was   mainly   controlled   by   the   transport   of    molecular   oxygen   through   thegrowing   oxide   layer.   Pre-oxidation   at   1300   8 C   for   24   h   in   ambient   air   increased   the   strength   of    Si/SiC   ceramics   by   around   46%   because   of    thehealing   of    the   surface   defects   created   during   surface   preparation   by   the   oxide   layer. #   2012   Elsevier   Ltd   and   Techna   Group   S.r.l.   All   rights   reserved. Keywords:   D.   Mechanical   property;   Processed   bio-precursor;   SiC   ceramic;   Oxidation   behavior 1.   Introduction Siliconized   silicon   carbide   (Si/SiC)   or   reaction   bondedsiliconcarbide   (RBSC)   is   a   well   known   industrial   ceramiccomposite   produced   by   infiltrating   a   porous   carbonaceouspreformwith   molten   silicon   atatemperature   in   the   range   of 1500–1600 8 C.Variouscarbonaceous   preforms   are   used   suchascarbon/SiC   powder   compacts,   carbon   felts,   carbon   fiber–carbonmatrix   composite   preforms,   and   carbon   powdercompacts[1–5].   Si/SiC   ceramics   are   also   produced   byreplicationof    naturally   occurring   plantbio-structures   (bio-preformsor   bio-precursors),   for   example,   woods.   The   obviousaimis   to   mimic   the   biological   cellular   morphology   in   a   ceramicmicrostructure.   Compared   to   the   technically   manufacturedceramicmaterials,   the   biomorphic   Si/SiC   ceramics   areexpectedto   provide   superior   combination   of    properties.   In   atwostep   processing,   bio-preforms   are   first   converted   to   carbontemplates   (C-templates)   which   on   subsequent   infiltration   withliquid   silicon   produce   Si/SiC   ceramics.   A   major   problem   of    thisprocessing   technique   arises   out   of    the   variation   of    micro-structureof    bio-preforms.   Depending   on   the   plant   age,   locality,climate,etc.   lot-to-lot   variation   of    the   microstructures   of    thestartingbio-preforms   can   occur   which   ultimately   affects   thepropertiesof    the   final   ceramics.   Processed   bio-preforms   areusedto   produce   industrial   goods   of    acceptable   uniformproperties(e.g.   paper).   Recently   we   used   processed   bio-preformsmade   from   bamboo   pulp   to   synthesize   Si/SiCceramics[6].The   aim   was   to   develop   a   process   of    synthesisofbio-morphic   Si/SiC   ceramics   with   the   avoidance   of    structuralinhomogeneity   and   related   problems.   This   paper   reports   theoxidationbehavior   of    Si/SiC   ceramics   synthesized   from   theprocessedbio-precursor;   the   effect   of    oxidation   on   mechanicalpropertiesof    biomorphic   Si/SiC   ceramics   will   also   bediscussed. 2.Experimental Pulp   stock    was   made   from   well-matured   Bamboo   (  Bambusatulda Roxb.   collected   from   Jorhat   district   of    Assam,   India)followingkraft   pulping   method.   Pulp   slurry   was   prepared   by www.elsevier.com/locate/ceramint  Available   online   at   www.sciencedirect.com Ceramics   International   38   (2012)   4701–4706*   Corresponding   author.   Tel.:   +91   033   2473   3496;   fax:   +91   033   2473   0957. E-mail   address:   omprakash@cgcri.res.in   (O.   Chakrabarti).0272-8842/$36.00   #   2012   Elsevier   Ltd   and   Techna   Group   S.r.l.   All   rights   reserved.doi:10.1016/j.ceramint.2012.02.054  beatingthe   bleached   pulp   stock    followed   by   sizing   usingchemicalagents   and   used   for   making   rectangular   boards   bycastingand   drying   in   pressedcondition   (for   processing   detailssee[6]).The   pulp   boards   had   a   very   high   content   of    cellulose   of  $ 78.1%(w/w)   and   were   reasonably   strong.   The   properties   of thepulp   fiber   boards   are   presented   in   Table   1.The   cast   pulpboardswere   subsequently   converted   to   C-templates   bypyrolysisat   800   8 Cunder   flowing   nitrogen.   The   dimensionalshrinkages,pyrolytic   weight   changes   and   other   properties   of theC-templates   are   listed   in   Table   1.C-templates   were   furtherinfiltratedand   reacted   with   moltenSi   at   1600   8 Cunder   vacuumforconversion   to   SiC   ceramics.   The   properties   of    the   SiinfiltratedC-templates   are   also   summarized   in   Table   1.Thedensityand   porosity   of    the   Si-infiltrated   specimens   weremeasuredby   standard   water   immersion   technique.   Rectangularsliceswere   cut   from   the   infiltrated   specimen   for   the   preparationofsamples   for   oxidation   behavior   study   and   mechanicalpropertyevaluation.   These   samples   were   properly   ground   andpolishedusing   diamond   paste   up   to   1   m m   finish.   Oxidationbehaviorwas   studied   by   thermogravimetric   method   (STA409C,Netzsch-Geratebau,   GmbH,   Germany)   in   flowing   dry   air( P =   100   kPa,   flowrate   =   1   L/h)   at1200,   1250,   1300   and1350 8 C.In   each   case   the   sample   (rectangular   polished   chips   of 10mm   Â   10   mm   Â   3mm)   was   heated   at   a   constantrate   (20   K/ min)and   an   isothermal   hold   was   given   for   7   h   at   the   peak temperature.   The   mass   variation   (%   mass   retained)   wasautomaticallyrecorded   as   afunction   of    time.   Rectangularpolishedbars   (45   mm   Â   3.5   mm   Â   2.5mm,edges   of    thetensilesurface   were   rounded   off)   were   tested   at   roomtemperature   for   bending   strength   measurement   (3-point   mode,span40   mm,   crosshead   speed   0.5   mm/min)   and   Young’smoduluswas   obtained   from   the   load–deflection   data   using   astandardsoftware   (Instron   Bluhill-2,   UK).   Average   values   of fivereadings   weretaken   for   strength   and   modulus   determina-tion.In   aseparate   experiment   bar   samples   were   heated   in   air   at1300 8 Cfor   24   h   and   subsequently   tested   for   measurement   of    3-pointflexural   strength.   The   fracture   surfaces   offailed   sampleswereexamined   under   microscope. 3.Results   and   discussion 3.1.Material    preparation Despite   avast   loss   in   weight   during   pyrolysis,   C-templatesshowedstructural   integrity   without   any   sign   of    visible   cracks.Nearly   uniform   pyrolytic   shrinkages   were   obtained   in   all   majordimensions.Liquid   silicon   infiltration   produced   almost   densesamples.XRD-analysis   showed   the   presence   of    only   b -SiC   andSias   the   major   ceramic   phases   (Fig.   1).In   our   earlier   work    wehavecarriedout   detailed   XRD-   and   microstructure   analysis   of C-templatesand   Si-infiltrated   C-templates   [6].Microstructureexamination   of    the   C-templates   showed   that   the   structuralmorphology   of    cellulosic   fibers   was   retained   in   the   pyrolyzedsamples.The   contents   of    the   SiC   and   Si-phases   were   estimatedfromthe   densityand   porosity   data   and   found   to   be   39.09   and60.31vol%   respectively.   Microstructure   examination   revealedthatthe   fibrous   carbons   were   converted   to   the   fibrous   SiCstructureswhich   wereseen   to   be   embedded   in   the   solidifiedresidualSi   phase.   All   the   details   of    results   were   described   in   ourpublished   work    [6]. 3.2.   Oxidation   behavior  TheSi/SiC   ceramic   samples   exhibited   similar   trends   duringoxidativeheating   at   four   different   temperatures   (Fig.   2).Theyshowed   an   initial   weight   loss   followed   by   a   progressive   weightgain.The   maximum   loss   in   weight   occurred   at   temperaturesbetween   967   and   1017   8 C;the   weight   loss   was   found   to   vary   intherange   of    0.01–0.04%.   Also   achange   was   noticed   in   theweightgain   trend.   Following   a   rapid   gain   in   weight,   the   gain Table   1Post-ceramization   properties   of    C-templates   and   its   bio-preform   formulation.Characteristics   of    processed   bio-preform   Characteristics   of    C-templates   Characteristics   of    siliconized   C-templatesBulk    density a (g   cm À 3 )Flexuralstrength b (MPa)Young’smodulus c (MPa)Dimensional   shrinkages   (%)   Pyrolyticweightloss   (%)Bulk density(g   cm À 3 )Bulk density(g   cm À 3 )Porosity(vol%)Flexuralstrength(MPa)Young’smodulus(GPa)Length   Width   Thickness0.58   4.4   Æ   0.8   165.8   Æ   26.9   20.26   21.33   21.43   71.21   0.60   2.66   0.60   207.7   Æ   7.8   199.5   Æ   5.1 a Determined   by   measuring   weight   and   dimension. b Determined   in   an   Universal   Testing   Machine   (Instron   Bluehill-2,   UK)   in   3   point   mode. c Determined   from   load–deflection   curve   obtained   during   strength   testing. 9080706050403020100 CC   a  r   b .  u  n   i   t 2 Theta ( o )SiCSiCSiCSiCSiSiSi ab Fig.   1.   XRD-profiles   of    (a)   the   silicon   infiltrated   C-template   and   (b)   the   C-template   prior   to   infiltration.  A.    Maity   et    al.    /    Ceramics    International   38   (2012)   4701 – 4706  4702  became   slow   and   finally   the   weight   gain   curves   flattened   out.Thischange   became   more   pronounced   at   higher   temperatures.Isothermal   heating   was   continued   for   a   period   of    7   h   ateachtemperature(1200,   1250,   1300   and   1350   8 C).The   weight   gainrecordedwas   converted   to   weight   gain   per   unitarea   and   wasplottedagainst   time   (Fig.   3).Theinitial   weightloss   indicatedthat   residual   carbonwaslikely   to   be   presentin   theinfiltrated   specimens.Duringinvestigation   on   themicrostructure   of    densebiomorphicSi/ SiCceramicssynthesized   from   woodprecursors   byorientingimaging   microscopy,our   group   detected   thepresence   of unreactedcarbon   [7].   Theoxidationprocess   is   assumed   tobeginwithoxidationofresidual   carbon,evolving   gaseousCOorCO 2 ;   themass   loss   observed   can   be   assignedto   theoxidation   ofcarbon.   Atlow   temperatureoxidationof    carboniscontrolled   by   therate   ofaccompanyingchemicalreactionwhereasat   hightemperatures(forexample,   T    >   700 8 C)masstransferby   diffusioncontrols   theprocess   ofoxidationofcarbon   [8].   Bytheevolution   ofgaseous   products   poresareformed   and   thereacting   specieshaveto   diffuse   along   theporepriorto   reactingwiththecarbonlyingat   theporebottom.   Atthe   low   temperatures,thechemical   reactionrateislikely   to   below   and   theoxygen   consumption   along   theporeis   alsolow;   at   hightemperatureschemical   reactionratecorrespondingto   oxidation   ofcarbonis   highand   theoxygenconsumptionis   alsohighmaking   diffusion   mass   transfertocontrol   theprocessof    oxidation   of    carbon.Aslongas   carbonispresent   andreacts   with   oxygen,   thelength   ofthediffusionpathforthegaseousreactantwillincrease,   and,   as   aresult,theconsumptionrate   ofcarbonwill   decrease   as   thefunctionoftime.   This   is   evident   in   theweightlosstrend   as   shown   inFig.2.   Theincreasein   weightis   related   to   theoxidation   of boththe   Siand   SiCphasesgiving   riseto   theformation   of silica.   Thegrowth   of    thesilicaonthepore   wall   likely   resultsinaprogressivedecrease   in   thepore   size[9].   A   timeeventually   comes   when   theporesare   sealedwith   thesilicanearthesample   surface.   Thereactant   has   to   diffuse   throughacoherent   layerof    thesilica   depositedon   thesample   surfaceforreaction   to   occur.   We   observed   a   change   in   theweightgaintrend   in   theTGcurves   likely   because   of    theonset   of    theprotectivescale   formation.Finally   theweightgain   curveflattensout.   This   featureis   very   much   similar   to   theoxidationweightgaintrend   ofreaction   bonded   siliconcarbide,   a   dense   duplex   ceramic   compositecontaining   SiandSiC   phases   [10].For   theentire   periodof    isothermalheatingateach   of    thefourpeak    temperaturestheweightgain   datawere   found   to   be   best   fitby   aparabolic   oxidation   rate   curveofthetype:   w 2 =   kt  ,where   w   is   the   weightgainper   unit   area, k  theoxidation   rateconstant   and t  thetime   (Fig.3).   Theoxidationrate   constant k    at   each   temperaturewasdeterminedfromtheslope   ofthelinear   plotof    w   against  ffiffi  t  p  (Fig.   4).   Theactivation   energy   was   calculatedfromtheArrhenius   equation k  =    A e À E   /   RT  ,   where    A   is   thepre-exponentialfactor   indepen-dentof    temperature,   E    theactivation   energy,  R   thegasconstantand   T    theabsolute   temperature.   TheArrheniusplotoflog k    vs   1/  T  is   shown   in   Fig.5.   All   thepointscorrespondingto   thetest   temperaturesfitto   a   straightline 5004003002001000 100.0100.1100.2100.3100.4100.5100.6100.7 1350 o C1300 o C1250 o C1200 o C    %    M  a  s  s  r  e   t  a   i  n  e   d Time, min Fig.   2.   Oxidation   mass   change   vs   time. 20151050 0.00.20.40.60.81.0 1300 o C1350 o C1250 o C1200 o C    M  a  s  s  c   h  a  n  g  e  p  e  r  a  r  e  a ,  m  g  c  m   -   2 Square root of time, min 1/2 Fig.   4.   Oxidation   mass   change   vs   time 1/2 . 4003002001000 0.00.20.40.60.81.0 1350 o C1300 o C1250 o C1200 o C    M  a  s  s  c   h  a  n  g  e  p  e  r  a  r  e  a ,  m  g  c  m   -   2 Time, min Fig.   3.   Oxidation   mass   gain   vs   time.  A.    Maity   et    al.    /    Ceramics    International   38   (2012)   4701 – 4706    4703  wherethegradientofthelinegivesthevalue   of    theactivationenergy.   Thevalue   of    E    obtainedin   thepresent   workwas141.4kJ/mol.Incase   of    oxidation   ofSiit   is   widely   acceptedthattheprocessof    oxidation   is   controlled   bydiffusionofdiatomicoxygen   throughthegrowingsilicafilm[11].Theactivationenergy   foroxidation   of    Siagrees   wellwithvalue   of activation   energy112.5kJ/molfor   molecularoxygen   diffu-sionthroughfusedsilica   [12].Atlow   temperatures   theactivationenergiesfor   oxidationof    sinteredand   hot   pressedSiCare   reported   to   be   133.5   and   154.4   kJ/mol,   whereasathightemperatureshighvalues   are   obtained.Thissuggeststhat   oxidationofSiCat   low   temperaturesmaybecontrolledby   inward   diffusion   ofoxygenmolecules,similar   in   processoccurringin   thecase   ofoxidation   of    Si.Highactivationenergiesobtainedforoxidation   ofSiC   at   hightemperaturesare   not   in   favor   ofoxygentransportmechanism.   Manyinvestigationsare   reported   which   revealed   thatseveralfactorsare   responsible   for   higher   oxidationactivation   energyathigher   temperaturesfor   SiC[13–19].   The   activationenergy   is   dependenton   many   factors,such   as   crystal-lographic   plane,   dopantconcentration,   crystallization   of oxide   films,   and   grain   orientation   onthesurface   of    oxidation.Harris   examinedtheinfluenceof    crystallographic   plane({000   1}planes)   on   theoxidationofsingle   crystalSiC(6Hpolytype)   [20].According   to   theauthorthe(0   0   0 ¯  1)cfaceexhibitedparabolic   oxidationwiththeactivation   energyof 197   kJ/mol.   Theoppositeface   exhibitedreducedoxidationrates   and   theauthor   observed   linearkinetics   andmeasuredthe   activationenergyto   be356kJ/mol(itcorrespondsto   achemicalreaction   at   theinterface).Difference   in   oxidationbehavior   forthetwo   {00   0   1}faces   was   also   reported   byVonMunch   and   Pfaffender   [21],similar   to   theobservations   madebyHarris   [20].Ramberg   et   al.examined   theoxidationbehavior   of    thecubic   CVDSiCand   observed   that   theoxidationratesof    the(1   1   1)   face   of    cubic   CVD-SiC   were   thesame   as   those   ofthe(0   0   0   1)   face   of    thesingle   crystal   SiC,the   oppositefaces   of    thetwo   materials,( ¯  1 ¯  1 ¯  1)   and(0   0   0 ¯  1),alsooxidized   at   thesamerates,   but   much   faster   than   theiropposite(0   0   1)/(00   0   1)   faces   [22].Many   investigationsreportedextremelyhighactivation   energies( > 400   kJ/mol)fortheoxidation   process   of    polycrystalline   SiCmaterialsabove   1400   8 C.Several   studies   haveshown   that   appreciablelattice   diffusionofoxygenoccurs   during   oxidationof polycrystalline   SiCmaterials   at   hightemperatures(1300   8 Cand   above)[23].Theactivation   energyforlattice   diffusioninfusedsilica   was   reported   by   Sucovto   be   $ 300   kJ/mol[24];otherstudiesindicated   thisvalue   could   be   as   high   as   450   kJ/ mol[25].Lattice   diffusionof    oxidant   may   be   oneof    thereasonsforhighactivation   energiesfor   high   temperatureoxidation   of    SiC.Thedecrease   in   oxidation   rate   at   and   above1400   8 C   exhibitedbythesintered   polycrystalline   SiCmaterialswasfound   to   berelated   to   thecrystallization   of theoxidefilm.   It   wasindicated   that   thetransport   of    oxidantthroughthecrystalline   phasewas   considerably   slowerthanthrough   theamorphous   phase.Decreases   in   oxidationrate   of thesinteredmaterialwerecorrelatedwiththeformationof    acontinuouslayerof    thecrystals   overtheSiC   surface   [14].Singhalsuggested   outwarddesorptionof    COmolecules   fromthegrowth   interfaceas   thereason   behindthehighactivationenergyfortheoxidationprocess   of    SiCat   hightemperatures[13].In   our   material   both   theSiandSiCphases   oxidizeandtheactivation   energy   obtained   in   thepresent   study   indicatesthattheoxidation   process   occurs   by   inward   oxygentransportmechanism.   Theoxidationrate   constantsobtainedin   thepresentworkoverthetemperaturerange   of1200–1300   8 Care   listedin   Table2.   Coppola   et   al.studied   theoxidationbehavior   ofpressure-less   sintered a -SiC   ceramics(98–99%dense,sinteredwithboronand   carbon-based   sinteringadditives)   at   temperaturesbetween1200   8 C   and1500   8 Cin   static   airandobtainedrate   constantvalues   of 7.08 Â   10 À 13 ,   5.09   Â   10 À 12 ,7.64 Â   10 À 11 kg 2  /m 4 s   at1200,   1350and1500   8 C   respectively   [26].Oxidationbehavior   of    reaction   sinteredsilicon   carbide   havingabulk density   of    3.00g   cm À 3 and   unreactedSi-content   of  $ 23%(v/ v)wasstudiedin   airovera   temperaturerangeof    1200–1350   8 C   and   parabolicoxidationrate   constantsof 3.4 Â   10 À 11 ,   4.3 Â   10 À 11 ,1.03 Â   10 À 10 and1.45 Â   10 À 10 kg 2  /m 4 s   were   reported   at   temperatures1200,1250,   1300and   1350 8 C   respectively   [27].ReactionsinteredSiChas   higher   oxidation   rate   constants   than   pressure-lesssinteredSiCbecause   of    thepresence   ofresidual   silicon   in   thematerial.Thebiomorphic   SiC   has   higher   residual   siliconcontent   thanreaction   sintered   SiC   which   is   likely   related   to 6.86.76.66.56.46.36.26.1 -1.9-1.8-1.7-1.6-1.5-1.4-1.3 10 4 1/T    l  o  g   k 1350130012501200 Temperature ( o C) Fig.   5.   Plot   of    log   k    vs   1/  T  .Table   2Values   of    oxidation   rate   constant   at   different   temperatures.Temperature   ( 8 C)   Oxidation   rate   constant   (kg 2  /m 4 s)1200   3.85   Â   10 À 10 1250   1.19   Â   10 À 9 1300   1.51   Â   10 À 9 1350   3.31   Â   10 À 9  A.    Maity   et    al.    /    Ceramics    International   38   (2012)   4701 – 4706  4704  thehigher   oxidationrate   constants   obtainedin   thepresentstudy. 3.3.Flexural   strength The   effects   of    pre-oxidation   at1300   8 Cfor   24   h   in   air   on   thestrengthof    the   Si/SiC   ceramics   were   examined.   The   failure   of thespecimen   in   bending   occurred   without   any   deformation   andamean   fracture   strength   value   of    303.9   Æ   38.3   MPa   was   found.Pre-oxidationat   1300   8 Csubstantially   increased   the   strength   bynearly46%.   Limited   information   is   available   on   the   effect   of pre-oxidationon   the   strength   of    biomorphic   SiC   ceramics.   Incaseof    conventional   Si/SiC   ceramic   composites   synthesizedfollowingthe   powder   route   conflicting   information   is   available.Trantina   observed   that   the   strength   of    the   siliconized   SiC   couldbeincreased   by   about   25%   by   aheat-treatment   at   1200   8 Cfor1h   [28].   Tomlinson   et   al.   showed   that   the   oxidation   of siliconizedsilicon   carbide   always   decreased   the   strength,   theextentof    which   could   go   up   to   50%   depending   on   period   of oxidationat   1350   8 C[29].   For   strengthening   due   to   oxidation   itisgenerally   accepted   that,   after   oxidation   the   silica   scale   tendstofill   in   or   blunt   surface   flaws,   resulting   in   increasing   thestrength.Tomlinson   et   al.   observed   that   oxidation   of    siliconizedsiliconcarbide   specimens   at   1350   8 Cresulted   in   formation   of anoxide   layer   following   aparabolic   oxidation   law   andpreferential   inter-grain   oxidation   roughened   the   surface   andformeddeep   pits.   Extensive   cracking   was   observed   on   thesurfaceof    the   oxide   layer   that   could   result   in   decrease   instrength.Our   sampleis   a   duplex   ceramic   compositesynthesizedfrom   processed   cellulosic   precursor   that   containedboththe   SiC   and   Si   phases.   During   microstructure   examinationnocracks   were   evident   on   the   surface   of    the   specimen   oxidized.Anumber   of    fractured   surfaces   were   examined.   The   typicalfracturesurface   of    pre-oxidized   specimen   tested   at   roomtemperature   for   flexure   strength   measurement   is   shown   inFig.6.Hardly   any   initiation   site   could   be   located   or   crack    pathcouldbe   identified.   It   indicated   that   cracking   could   not   occurnearthe   surface   after   oxidation.   It,   therefore,   appeared   that,   inthepresent   case   pre-oxidation   at1300   8 Ccould   improve   theroomtemperature   strength   of    the   Si/SiC   ceramics   by   healing   orrounding   off    the   crack    or   flawsby   the   flow   of    the   silica   scale.Retentionof    strength   of    Si/SiC   ceramics   in   service   conditionsinvolving   high   temperature   ambient   atmosphere   is   a   veryimportantissue.   The   experimental   results   obtained   in   thepresentstudy   indicated   the   suitability   of    the   Si/SiC   materialsynthesizedfrom   processed   bio-precursors   for   such   uses. 4.Conclusions Oxidative   heating   of    Si/SiC   ceramics   synthesized   fromprocessedcellulosic   bio-precursors   in   dry   air   indicated   initiallossof    weight   of    the   samples   likely   because   of    burning   of    anyresidualcarbon   presentin   them.   Maximumweight   loss   variedinthe   range   of    0.1–0.4%   and   it   was   found   to   occur   attemperatures   in   the   range   of    978–1016   8 C.Afterwards,   thesampleswere   found   to   gain   in   weight   and   following   rapidweightgain,   the   weight   change   curves   flattened   out   indicatingformationof    aprotective   oxide   scaleon   the   surface.   The   weightgaindata   obtained   during   the   isothermal   heating   at   tempera-turesin   the   range   of    1200–1250   8 Cwere   found   to   be   best   fittedtoaparabolic   rate   equation.   Activation   energy   of    141.4kJ/molobtained   in   the   present   study   indicated   that   oxidation   mightoccurvia   transport   of    molecular   oxygen   through   the   growingoxidelayer.   Si/SiC   samples   pre-oxidized   at   1300   8 Cfor   24   h   inambientair   showed   an   increase   of    flexure   strength   by   46%,probablybecause   of    the   healing   of    the   surface   cracks   createdduringsample   surface   preparation   by   the   flowof    oxide   scaleandit   further   indicated   suitability   of    the   material   for   hightemperatureapplication   in   challenging   atmospheres. Acknowledgements The   support   of    the   Coir   Board,   Ministry   of    Micro,   Small   andMediumEnterprises,   Govt.   of    India,   for   carrying   out   the   work andin   granting   one   of    the   authors   (AM)   a   fellowship   (projectassistant)is   acknowledged   with   thanks. References [1]   P.   Popper,   The   preparation   of    dense   self    bonded   silicon   carbide,   in:   P.Popper   (Ed.),   Special   Ceramics,   The   British   Ceramic   Research   Associa-tion/Heywood   and   Company   Ltd.,   London,   1960,   pp.   209–219.[2]   C.W.   Forrest,   F.Kennedy,   J.V.   Shennan,   The   fabrication   and   properties   of self-bonded   silicon   carbide   bodies,   in:   P.   Popper   (Ed.),   Special   Ceramics,vol.   5,   British   Ceram.   Res.   Assoc.,   UK,   1972,   pp.   99–123.[3]   T.   Hase,   H.   Suzuki,   T.   Iseki,   Formation   process   of    b -SiC   during   reactionsintering,   J.   Nucl.   Mater.   59   (1976)   42–48.[4]   J.   Schulte-Fischedick,   A.   Zern,   J.   Mayer,   M.   Ruhle,   M.   Frieß,   W.Krenkel,R.   Kochendorfer,   The   morphology   of    silicon   carbide   in   C/C–SiC   compo-sites,   Mater.   Sci.   Eng.   A   332   (2002)   146–152.[5]   W.B.   Hillig,   Tailoring   of    Si/SiC   composite   for   turbine   applications,   in:   J.Burke,   E.N.   Lenoe,   R.   Nathan   Katz   (Eds.),   Ceramics   for   High   Tempera-ture   Applications.   II,   1st   edition,   Brook    Hill   Publishing   Company,   Chest-nut   Hill,   MA,   USA,   1978,   pp.   989–1000.[6]   A.   Maity,   D.   Kalita,   T.K.   Kayal,   T.   Goswami,   O.P.   Chakrabarti,   H.S.   Maiti,P.G.   Rao,   Synthesis   of    SiC   ceramics   from   processed   cellulosic   bio-precursor,   Ceram.   Int.   36   (2010)   323–331.[7]   V.   Pancholi,   D.   Mallick,   Ch.   AppaRao,   I.   Samajdar,   O.P.   Chakrabarti,   H.S.Maiti,   R.   Majamdar,   Microstructural   characterization   using   orientationimaging   microscopy   of    cellular   Si/SiC   ceramics   synthesized   by   replica-Fig.   6.   Fracture   surface   of    pre-oxidized   Si/SiC   specimen.  A.    Maity   et    al.    /    Ceramics    International   38   (2012)   4701 – 4706    4705
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