CERAMICSINTERNATIONAL

Available online at www.sciencedirect.com

http://dx.doi.org/0272-8842/& 20

nCorrespondinTel.: þ91 512 2

E-mail addre

(2015) 12849–12860

Ceramics International 41 www.elsevier.com/locate/ceramint

Polycarbosilane based UD C/SiC composites: Effect of in-situ grown SiCnano-pins on mechanical properties

Suresh Kumara,n, M.K. Misraa, Samar Mondalb, R.K. Guptaa, Raghwesh Mishraa, Ashok Ranjana,A.K. Saxenaa

aDefence Material & Stores Research & Development Establishment, Kanpur 208013, IndiabAdvanced Systems Laboratory, Hyderabad 500058, India

Received 30 April 2015; received in revised form 22 June 2015; accepted 26 June 2015Available online 4 July 2015

Abstract

Uni-directional C/SiC composites were fabricated by thermal molding of carbon fiber tows with the in-house synthesized polycarbosilane. Themolding was carried out at 300 1C followed by pyrolysis at two different temperatures viz. 1400 and 1600 1C. The composite specimens weredensified by repeated vacuum infiltration of the polycarbosilane at 170 1C and pyrolysis at the pre-decided temperatures to achieve density up to1.8–1.85 g/cc. Flexural and tensile strength of the composites pyrolysed at 1600 1C were found to be 700 MPa and 275 MPa while the propertiesof the composites pyrolysed at 1400 1C were about 20% lower than these values. Micro-structure revealed that the in-situ growth of SiC nano-pins at 1600 1C, play an important role to enhance the mechanical properties. Growth mechanism of the SiC nano-pins has also been proposed.& 2015 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Keywords: C/SiC composite; Ceramic; Mechanical property; Polycarbosilane

1. Introduction

Continuous fiber reinforced Cf/SiC and SiCf/SiC compositehave been widely used for many high-temperature structuralapplications. These composites can be prepared by chemicalvapor infiltration (CVI), reaction bonding, hot pressing andpolymer impregnation and pyrolysis (PIP) process [1–4]. Anumber of studies have been carried out by many researcherson the use of the polymer derived ceramics precursors (PDC),for the SiC, SiNC and SiOC based ceramics [5]. The propertiesof the PDCs are related to their unique nano sized micro-structure which makes them of great use in many disciplines[5]. SiC has been considered for many high temperatureapplications due to its unique mechanical and thermal proper-ties [5,6]. Therefore, a variety of Polycarbosilanes (PCS) havebeen synthesized so far, to generate SiC fiber/matrix [6–9].

10.1016/j.ceramint.2015.06.12215 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

g author.451759-78x140; fax: þ91 512 2450404.ss: sureshtanwar@rediffmail.com (S. Kumar).

Molecular structure and type of the starting polymer controlsthe composition, the number of phases, phase distribution andthe microstructure of the ceramic obtained. Although the use ofthe PCS has been demonstrated way back in 1970 for SiC fibersynthesis, but, a uniform material is still not available in theopen market in all the countries. Therefore many types of PCShave been developed and being used for the different purposes[8,10–12]. Methyl-polycarbosilane ([MeHSiCH2]n) has beensuccessfully used for generation of SiC matrix for hightemperature composites as well as for the SiC fiber synthesis[1–3,10–12].Processing parameters which control the microstructures and

properties of the ceramic matrix composite have also beenstudied by few researchers [13–16]. Still, there is a scope forthe improvement in the PDC based ceramics composites, asthe translation of the reinforcement strength to the composite isonly about 15% [2,3,11]. The mechanical properties of theceramic matrix composite depend on many parameters whichare related to the knowhow of the polycarbosilane synthesis,process equipment features and the composite processing

S. Kumar et al. / Ceramics International 41 (2015) 12849–1286012850

parameters. Some studies have been reported on the modifica-tion of the matrix by adding micro/nano ceramic powders inorder to cut down the processing time and to enhance themechanical strength. The nanostructure growth during thepyrolysis of polycarbosilane (wires, belts, tubes, etc.) report-edly improves the properties of the final ceramic product [17].

In-situ SiC nano structure has been generated by addingferrosene in the PCS which has resulted in to SiC nanos-tructure growth in the matrix during the pyrolysis. It has showna significant effect on the properties of the C/SiC composites[18]. Although, the basic processing steps of the C/SiCcomposite remained similar but the reported properties aresignificantly different from one study to the other [2,3,11,17,18]. This might be due to the fact that, the starting PCS isdifferent in most of the reported cases w.r.t. one or more thanone parameters like molecular weight, softening point, func-tional groups and ceramic yield etc. In addition to this, thecomposite processing parameters like curing time, temperature,rate of heating, pyrolysis temperature, holding time weredifferent for different types of polycarbosilane which havenot been spelled out in detail. In this scenario it becomes verydifficult to adopt any reported process to realize PDC based C/SiC composite products. Some studies have been reported onthe fabrication of C/SiC composite structural parts by the PIPprocess [19,20]. Filament wound technique is proposed for theC/SiC composite nozzle fabrication [21,22]. In the presentwork it is planned to carry out the initial experiments toestablish the composite processing parameters and their effecton the unidirectional (UD) composite properties using the in-house synthesized polycarbosilane. The UD composite size ischosen based on the requirement for a typical filament woundprocess to realize a large size conical structure.

Therefore, it is aimed to fabricate a multiple numbers of UDC/SiC composite strips using the in-house synthesized methyl-polycarbosilane and to study the effect of pyrolysis tempera-ture on the bending and tensile strength. The compositeproperties would be discussed in the light of microstructureget evolved due to the different processing conditions.

2. Experimental and raw materials

2.1. Methyl-polycarbosilane

In-house synthesized methyl-polycarbosilanes of softeningpoint in the range 190–200 1C and 120–130 1C were used forC/SiC composite fabrication. PCS having softening point, inthe range 190–200 1C is termed as HSP–PCS and it was usedfor the thermal molding of the composite strips. The PCShaving softening point in the range 120–130 1C is termed asLSP–PCS and it is used for the densification of the composites.The PCS was characterized w.r.t. its functional groups,molecular weight, ceramic yield and the phases obtained atdifferent pyrolysis temperatures. FT-IR spectra of the assynthesized PCS and the PCSþDVB cured systems weremeasured using the KBr pellet method in a Perkin Elmerinstrument. The PCS was also analyzed for its elementalcomposition using the standard CHNS analysis method.

Molecular weight of the HSP–PCS and LSP–PCS wasdetermined using the gel permeation equipment. To determinethe ceramic yield of both types of the PCS, TGA was carriedout up to 900 1C, as most of the decomposition gets completedby this temperature [5,6]. It was carried out using hightemperature thermo-gravimetric equipment (TGA Q500V20.10 Build 36, Universal V4.7 A TA instrument). Approxi-mately 8–10 mg sample was placed in an open alumina panand the sample was heated at 10 1C/min in following N2,30 mL/min. Both types of the PCS samples were pyrolysed at1200, 1400 and 1600 1C in vacuum for 2 h in order to finalizethe ultimate pyrolysis temperature. Phase analysis of theceramized powder was carried out using a Philips XRDinstrument with Cu Kα radiation, at 40 kV, 2θ range of 20–901, step size of 0.051 and the scan rate of 11/min.

2.2. UD C/SiC composite fabrication

T300, 6 K fiber tows having an average strength up to3.2 GPa and modulus up to 3.0 GPa were used as reinforce-ment. Five tows (30000 filaments), were wound on a slottedgraphite fixture to obtain multiple numbers of uni-directionallyaligned fibrous preforms having fiber volume fraction (Vf) inthe range of 50–55%.The graphite fixture along with the fibers,was kept in a chemical vapor deposition reactor to apply a thininterface of pyro-carbon (PyC) on the fibers. The furnace washeated up to 1100 1C in flowing nitrogen at the heating rate of10 1C/min. Methane and hydrogen gases were passed throughthe furnace in the ratio of 1:1.2 (vol./vol.). The reactor pressurewas maintained at 10 mbar with the help of two stage vacuumpumping and throttle valve system.The deposition conditionswere maintained for 2 h.The furnace was allowed to cool undervacuum/inert and the fixture was taken out from the furnace atroom temp.The resin solution was prepared using the HSP–PCS and Di-

venyl Benzene (DVB) in a ratio of 5:1 (wt./wt. basis). Thesolution was further diluted with petroleum ether (50:50, vol./vol.). Few drops of Platinum spere catalyst were added to thesolution and it was impregnated in to the PyC coated UDfibrous preforms using the standard vacuum infiltration tech-nique. The fixture was taken out from the resin solution andwas allowed to cure at 300 1C under the pressure (1–2 t) in ahydraulic press to obtain UD carbon fiber reinforced compositestrips. The cured composite strips were pyrolysed undervacuum/inert at 1400 and 1600 1C temperature for 2 h. Theheating was carried out at a rate of 0.5–2 1C/min up to 1000 1Cand at 10 1C/min beyond 1000 1C during the pyrolysis. Thepyrolysed C/SiC composite strips were further densified byinfiltrating the LSP–PCS without DVB addition at 170 1Cunder vacuum followed by the pressurizing the system up to6 bar with inert gas and heated further up to 300 1C. In order toimprove C/Si ratio in the matrix DVB was not mixed in theLSP–PCS as DVB generates additional carbon and disturbs theC/Si ratio of the PCS. However, during molding it was feltnecessary to add DVB in HSP–PCS in order to generate abetter cross-linked network of the matrix. Infiltration andpyrolysis cycles were repeated several times to obtain C/SiC

Fig. 1. Process flow sheet of fabrication of UD C/SiC composite.

Fig. 2. FTIR spectra of PCS before and after thermal molding.

Fig. 3. TGA analysis of the as synthesized PCS and PCS/DVB mixture.

S. Kumar et al. / Ceramics International 41 (2015) 12849–12860 12851

composite strips of density in the range of 1.8–1.85 g/cc. Thecomposite strips pyrolysed at 1400 1C are termed as set-1,while, the ones pyrolysed at 1600 1C are termed as set-2. Theeach composite strip has an average Vf of about 0.50–0.55 andan average cross-sectional area of about 2.8 mm2.The sche-matic and sequence of the UD C/SiC composite processingsteps are shown in Fig. 1.

2.3. Composite characterization

Room temperature flexural strength (three point bendstrength) was determined for the strips of the both sets overa span of 20 mm with a crosshead speed of 0.5 mm/minute.The flexural strength was calculated as per the ASTMC-1341.Room temperature tensile strength was also measured in agauge length of 50 mm. Crosshead speed was maintained at0.5 mm/minute. From the each set, a minimum number of 10strips were tested for the flexural and tensile tests. The testingwas carried out at Instron, model 5967 machine. The micro-structure of the tested composite strips was studied using SEM,ZEISS and FEI Quanta 400. EDAX analysis was also carriedout using the same instruments.

3. Results and discussions

3.1. Polycarbosilane characterization

3.1.1. Physical propertiesUsing the procedure described by Saxena et al. [8] PCS

having controlled molecular weight and softening temperaturecould be synthesized in adequate quantity. The numberaverage molecular weight of the HSP–PCS was found to be1800 where Mw/Mn was 5.0 while for the LSP–PCS, it wasfound to 1250 with Mw/Mn¼5.36. Weight fraction of C, H, N,

O and S was found to be 38.23%; 7.81%; 0.0%; 0.01% and0.001% respectively. Whereas, silicon percentage is deter-mined by subtracting the % weights of C, H, N, O and S from100; it is calculated to be 53.949%. Density of the HSP–PCScrystals was found to be about 1.0 g/cc.

3.1.2. FTIR analysisIt is evident from the FTIR spectra (Fig. 2) that the PCS,

primarily contains, C–H, Si–CH3, Si–CH2–Si and Si–H bonds.The peaks 2956 cm�1 and 2896 cm-1corresponds to the C–Hbond stretching of Si–CH3; 2092 cm�1 of Si–H stretching; thepeak located at 1411 cm�1 is attributed to CH2 deformation ofSi–CH2–Si; 1248 cm�1corresponds to Si–CH3 stretching and1043 cm-1corresponds to Si–CH2–Si. However Si–O–Si bondpeak is also evident at1100 cm-1. Si–O–Si might be resultantof the partial surface oxidation of the PCS during thecomposite molding. Although, the primary curing of the PCSis performed by its crosslinking with the DVB, but atmo-spheric oxygen would also react with Si–H reactive bond. It

Fig. 4. XRD analysis of the matrix obtained at different pyrolysis temperatures.

S. Kumar et al. / Ceramics International 41 (2015) 12849–1286012852

can also be noted that the peak intensity corresponds to Si–Hbond decreased after curing with DVB. This shows that themost of the Si–H bonds available in the PCS got consumeddue to curing reaction with the DVB. However, all Si–H couldnot be consumed during curing as some of the Si–H bondsmight not be approachable for the cross-linking. Oxygencuring of the PCS is well reported for the PCS fiberstabilization which might have helped the surface crosslinkingduring the thermal molding of the composite strips. Similarfindings have been reported elsewhere also [10].

3.1.3. TGA studiesThermal decomposition of LSP–PCS, HSP–PCS and HSP–

PCSþ20% DVB solution mixture is shown in Fig. 3. Thedecomposition nature of all the samples is similar, but theinitiation temperature of the DVB mixed PCS is higher thanthe HSP–PCS and LSP–PCS. It is evident that, the HSP–PCSonly loses about 8% of its mass between 100 and 440 1C andits decomposition rate is relatively low beyond 440 1C. Weightloss in this temperature range is due to the evaporation of thelower molecular weight fraction. This evaporation creates a lotof pores in the matrix. The porosity need to be filled byrepeated infiltration/pyrolysis to impart a reasonable strength inthe composite. The weight decreased rapidly until about800 1C. This large weight loss is interpreted as the decom-position of the side chains which leads to the shrinkage of thematrix. Although, the shrinkage increases the density of thematrix but due to fiber network, the matrix does not shrinkuniformly. Multiple cracks get develop in matrix, perpendi-cular to the fiber axis due to CTE mismatch. Beyond 800 1C,the rate of decomposition gets lowered further, which can beseen from the flattened TG curves.

Between 100- and 440 1C, the LSP–PCS loses about 20%weight against 8% of HSP–PCS. The decomposition rate is thehighest between 400 and 500 1C. Similar to HSP–PCS thedecomposition rate is low beyond 800 1C. This indicates that,most of the decomposition gets completed by 900 1C in boththe cases. The ultimate ceramic yield at 9000 C is about 75% ofHSP–PCS against 63% in case of LSP–PCS.

Decomposition behavior of the HSP–PCSþ20% DVBmixture is found to be slightly different from HSP–PCS andLSP–PCS. The weight loss of the PCS/DVB mixture between100 to 440 1C is only about 2–3%. Similar to the HPS–PCS,the decomposition of the mixture also starts at about 450 1C.Si–H reactive group gets consume in three dimensionalnetworking during the curing with DVB. Therefore, highercrosslinking network hinders in the removal of the smaller sizegroups as they might get entrapped in the network. The finalchar yield is found to be about 82% by weight against 75% ofHSP–PCS. However, the higher ceramic yield of mixture maybe in the form of additional carbon in the matrix.

From the TGA curves, it is evident that, most of the weightloss takes place between 440 and 800 1C. In other words, therate generation of the pyrolysis gases is the highest in between440 and 800 1C. These gases have to diffuse back from thecomposite into the bulk. During the diffusion the gases maylead to delamination as the interfilament binding at the first

pyrolysis time is the least. Therefore, it is preferred to havelower diffusion rate of the gases which should not haveadverse effect on the composite integrity. This has to befinalized by experimentation. Therefore, to cater sufficient timefor the back diffusion of the pyrolysis gases and to avoid thede-lamination of the composite, the rate of the heating duringthe pyrolysis, was kept between 0.5 and 1 1C/min up to 800 1Cand 5 1C between 800and 200 1C. Beyond 1200 1C the ratewas kept at 10 1C up to the ultimate heating temperature tohave a relatively poor interfacial bonding.

3.1.4. XRD analysisIn order to study the ceramic phase formation, the HSP–PCS

and LSP–PCS were pyrolysed at three different temperaturesviz. 1200, 1400 and 1600 1C. The XRD profiles of thepyrolysed PCS ceramic powder of both the cases were similar;for the HSP–PCS it is shown in Fig. 4.The appearance of the diffraction peaks at 2θ¼35.61, 601,

and 71.61, correspond to the lattice planes of (1 1 1), (2 2 0),and (3 1 1) for the β-SiC polymorph (JCPDC:75-0254). Theappearance of the diffraction peaks at 2θ¼34.11, 41.51 and75.61, correspond to the lattice planes of (1 0 1), (1 0 4) and(0 0 1 2) for the 6 H–SiC polymorph (JCPDC:75-1541). Thebroad peaks indicate that the crystal grain size of the SiCmatrix might be in nano size. The peak position corresponds toβ-SiC polymorph are present at all temperatures but their shapeand intensity are different. The peaks become sharper astemperature increased. In other words it indicates that thecrystallite size increases with temperature. Peaks correspond to6H–SiC polymorph phase is seen only at 1600 1C; howeverthey are comparatively of low intensity. In addition to this,some of the peaks which were present at 1400 1C got mergedin the spectra at 1600 1C.Origin software was used for the qualitative analysis of the

2θ values of the main peaks. The mean nano-crystallites size(L) of the ceramic powder was estimated using the Scherrer

S. Kumar et al. / Ceramics International 41 (2015) 12849–12860 12853

equation (1) and the XRD data (Fig. 4).

L¼ Kλ

β Cos θð1Þ

where λ is the X-ray wavelength in nanometer (nm), β is the peakwidth of the diffraction peak profile at half maximum heightresulting from the crystallite size in radians, K is a constant relatedto the crystallite shape, normally taken as 0.9 [23]. Averagecrystallites size based on the first peak of β-SiC at 1200, 1400and 1600 1C were found to be 1.12 nm; 1.91 nm and 4.36 nmrespectively.

In addition to the β-SiC peaks, there are some faint peaks of silicaat 1400 1C, while there is no indications of formation of silica at1200 1C. The silica peaks are identified at the diffraction peaks at2θ¼41.71, 49.41, and 66.61, correspond to the lattice planes of(2 0 0), (1 1 2) and (2 1 2) for the hexagonal SiO2 polymorph(JCPDC:89-8951). The formation of silica at 1400 1C might bedue to the presence of some inherent oxygen in the PCS and additionof some oxygen during the curing of the PCS sample in air.

Fig. 5. Mechanical properties of the com

Fig. 6. Mechanical properties of the com

At 1200 1C, the reaction might not have been completed with theoxygen and would have resulted in some broad peaks. These peaksmight be resultant of the mixture of SiC and SiO2 amorphous phases.At 1400 1C the peaks of SiO2 are evident but their proportion is veryless compared to the β-SiC. At 1600 1C, the peaks of SiO2 seem tobe absent. The SiO2 formed at 1400 1C, might have been eitherevaporate in the form of SiO after reacting with SiC (reaction(2)) orwould have converted into SiC following the reaction(3). Formationof a smaller amount of crystallized silica in the pyrolyzed ceramics at1400 1C is being reported elsewhere [24].

SiCþ2SiO2-3SiOþCO (2)

SiOþC-SiCþCO (3)

3.2. Composite processing and characterization

As mentioned, at two pyrolysis temperatures, UD compositestrips of density up to 1.85 g/cc were obtained after 6 pyrolysis

posites pyrolysed at 1400 1C (set-1).

posites pyrolysed at 1600 1C (set-2).

S. Kumar et al. / Ceramics International 41 (2015) 12849–1286012854

cycles. A minimum of 10 samples are tested for flexural andtensile strength from each set. The samples are designated asS-1–S-10 within the sets. The representative stress–straincurves of minimum three samples of tensile are shown inFig. 5a while for flexural, it is shown in Fig. 5b. Similar curvesfor the set-2 are shown in Fig. 6a and b.

The arithmetic mean average tensile and flexural strengthvalues found to be the higher for the set-2 which waspyrolysed at 1600 1C compared to the ones pyrolysed at1400 1C (set-1). The average tensile strength of the set-1 andset-2 was found to be 235 MPa (s.d.¼17.6) and 265 MPa(s.d.¼15.0) respectively. Similarly, the mean bending strengthof the set-1 and set-2 was found to be 595 MPa (s.d.¼25.3)and 690 MPa (s.d.¼23.4) respectively.

The stress–strain curves of the tensile tested specimens ofboth the sets are similar (Figs. 5a and 6a). It shows almost alinear behavior up to the ultimate failure and sudden drop inthe stress. The linear behavior indicates that damage insensi-tive nature of the composite similar to the carbon–carboncomposite [25]. A damage insensitive stress strain behavior isobserved when the deformation of the composite is mainlydominated by the fiber in the loading direction. The elasticmodulus of the composite would be almost independent of thematrix. SEM images of the as synthesized and tensile testedspecimens of set-1 and set-2 are shown in Fig. 7(a)–(d).

Big pores are not observed in the matrix of the set-1 alsonano-pins are not present. The crack propagation is observedalong the fiber periphery (Fig. 7a). The fibers of differentlengths get pulled out from the matrix during tensile tests ofset-2 (Fig. 7b). This shows that the fiber matrix interface ismoderate and the stress generated cracks have deflected alongthe interface before the bulk failure of the composite. Highmagnification image of the set-2, revealed that the matrix havemultiple, equally spaced and uniform transverse cracks per-pendicular to the fiber axis. Similar to the set-1the matrix ofthe set-2 seems to be free from the big pores in the vicinity ofthe filament (Fig. 7c). The dense matrix build up around thefilament is the resultant of the high char yield HSP–PCS andDVB resin system and it gets develop during the firstpyrolysis. Subsequent infiltration helps in development of thematrix in the interfilament pores. The fiber pull out is ofseveral tens of microns (Fig. 7b and c) is observed in case ofset-2. The fiber pull out seems to be at the PyC/SiC matrixinterface as the PyC coating remains are observed on thesurface of the fractured fibers (Fig. 7d). The damage of theunidirectional composite involves multiple micro-cracks or thecracks that form in the matrix perpendicular to the loadingdirection (Fig. 7c). The cracks are obvious due to theconversion of the lower density (1.0 g/cc) polycarbosilane tothe higher density (3.2 g/cc) SiC; there is net volume shrink-age. CTE of the carbon fiber along the axis is 0–1� 10�6/Kwhile in the radial direction it is about 9� 10�6/K. Therefore,during pyrolysis and cooling, the composite have non-uniformshrinkage due to the CTE mismatch between Carbon fiber andSiC [4,26,27]. Although, these micro-cracks get filled to anextent during the subsequent infiltration/pyrolysis cycles, but,still some pores remain unfilled (Fig. 7a). The failure might

have started at the crack tip in the matrix due to stressconcentration.Progress of the matrix cracks which are already present in

the matrix and the fiber failure are random phenomena duringthe tensile loading. The failure of the specimens is almostinstantaneous (at same strain) as the strength gets reduced fromabout 600 MPa to about 20 MPa (Figs. 5 and 6). The fibersseem to be broken almost at similar extension level. This alsoindicates that the load carried by the matrix was much lowercompared to the fibers. Due to this reason the load transfermechanism is different than that of the composites like SiC/SiC or CVI/CVD based Cf/SiC [25] where matrix does nothave multiple fragments as in PIP based composites. In CVIbased composites the SiC matrix get deposited in solid form onthe fibers and remains as such while in case of PIP the matrixgets build up in multiples numbers of infiltration cycles and ismade up of small fragments. In CVI based C/SiC or SiC/SiCcomposites, the ratio of the elastic modulus of the matrix to thefiber is about 2.0 while for C/C category composite (like PIPbased C/SiC) this ratio is about 1/20. From the failurebehavior, it can be interpreted that PIP based C/SiC compositedo have matrix insensitive behavior similar to the C/Ccomposites.The flexural stress–strain curves are slightly different from

the tensile ones. Although the curves are similar for both thesets and for all the samples but bending characteristics such asinitial slope and failure stress are different. The curves can bedivided into two regions: an initial linear step followed by anonlinear step. The C/SiC composites pyrolysed at 1600 1C,exhibit better flexural properties than the ones pyrolyzed at1400 1C. In contrast to the tensile curves the flexural stress–strain curves are zig-zag which shows the phenomena of crackarresting, branching and crack deflection. This shows clearly,that the failure is non-brittle. In the flexural stress strain curvesit can be seen that it is not only the matrix alone which fails insteps, there may be some outer fibers which might got failedalong with the matrix. Therefore at many instances stress camedown sharply due to the failure of outermost fibers where crackgets initiated by a pore/crack in the matrix. On transfer of theload to the remaining unbroken fibers the composite strip bearthe load till further bulk failure of the fiber tows.The variation in the strength between the two processing

conditions might be due to the different morphological featuressuch as micro-cracks and fiber/matrix local debonds. Thesedifferences are unavoidable due to the different CTEs of thecomposite constituent's and non-uniformity in the shrinkage ofthe composite during the pyrolysis (Fig. 7c). The variation mayalso be due to the change of the interfacial conditions,microstructure at two different temperatures and heating rates.A wide variation in the flexural strength values has been

reported for the C/SiC composites [25–27]. Flexural strengthof the C/SiC and SiC/SiC composite fabricated by differenttypes of polycarbosilane precursor and at different pyrolysistemperatures is reported in between 130 and 662 MPa[18,28,29]. The density, pyrolysis temperature and precursortype play a vital role for the ultimate mechanical properties.The flexural strength depends on the ultimate density. It was

Fig. 7. SEM images of set-1 and set-2 showing different aspects.

S. Kumar et al. / Ceramics International 41 (2015) 12849–12860 12855

up to 662 MPa for 3D C/SiC composite of density 1.95 g/ccwhile at 1.79 g/cc it was only 557 MPa [30]. In another study,PIP based, 2D C/SiC composite have demonstrated flexuralstrength up to 319.2 MPa [31]. With the increasing heatingrate, the density of Cf/SiC composites increases and theinterfacial bonding weakens. During the pyrolysis of the set-1 and set-2 heating rate was kept sufficiently high beyond1200 1C. The rate of cooling was as high as 20 1C/min from1600 1C to 1000 1C which would have resulted into theweakened interface. The heating/cooling effect would behigher on the set-2. The flexural strength of Cf/SiC compositereportedly enhanced from 145 to 480 MPa when the heatingrate was increased from 0.5 to 15 1C/min. The strength couldbe further enhanced to (557 MPa) when the pyrolysis tem-perature of the last densification cycle was elevated from 1200to 1600 1C, The improvement was attributed to the desirableinterfacial structure and the increased density [32]. Initialimperfections can be minimized by selecting suitable processparameters and process route but cannot be prevented in the

PIP based composites, due to the complex pyrolysis mechan-ism and matrix/reinforcement interactions.Microstructure revealed that, the matrix in set-2 is bound by

the additional in-situ generated nano-structured material(Fig. 8a) which was absent in case of set-1 (Fig. 8b). Thenano-material might have bridge the cracks which are devel-oped during the pyrolysis of the polycarbosilane matrix. Theadditional strengthened SiC matrix due to in-situ grown nano-pins might have resulted into the finer reinforcement densityand smaller radii crack tips. These nano-pins might haveplayed a decisive role in controlling the crack initiation duringthe mechanical tests. It is visible that the matrix did not allowthe straight cut as the SiC matrix showed additional strengthwhere nano-pins are present (Fig. 8c). The nano-structuredreinforced matrix diverted the crack (Fig. 8c). Due to thisreason, the crack had to follow a long path along the matrixcrack then along the interface to reach the point where it couldfail the fiber. The shattered fiber image, obtained from thecompletely broken sample tested at higher strain rate (Fig. 8d)

Fig. 8. Uniformly grown sic-nano-pins in set-2 and set-1 matrix microstructure.

S. Kumar et al. / Ceramics International 41 (2015) 12849–1286012856

further proved that the composite have absorbed a lot ofenergy along the fiber/matrix interface.

The nano-pins form a complex bulk network where these arenaturally oriented in different directions (Fig. 9a). Theorientation of the nano-pins in different direction would helpin dispersion of the mechanical stress in multiple directionsbefore the composite fails due to the bulk failure of the fiberswhich any how carries maximum applied load. Also, the nano-pins would act like a bridge between the two edges of thematrix cracks (Fig. 9b) and might have been instrumental forthe enhanced mechanical properties of the set-2.

The pyrolysis conditions have a strong effect: for instance,vacuum promotes carbothermal reduction reactions and cantherefore promote crystallization. While the added pressure inany form like mechanical-hot pressing or gaseous-hot iso-staticpressing reportedly hinder the carbothermal reduction reactionsand leading to a suppression of crystallization [33].

The heating rate affects the ceramic yield and composition

[34]. Heating rate, dwelling temperature and time influence the

composition and microstructure of the matrix, as they affect

the extent of crystallization, carbothermal reduction reactions

and filler reactions occurring in the material. Therefore, the

properties of set-1 and set-2 are found to be different. Higherpyrolysis temperature and higher heating rates reportedlyimprove the mechanical properties [32] similar to the findingof this manuscript.

3.3. Nano-pin microstructure and growth mechanism

As mentioned, the process parameters remained same forboth the sets except the pyrolysis temperature. The highertemperature has resulted in situ nano-structure generationwhich strengthened the matrix. The higher temperature andheating rates weakened the interface which has resulted inhigher strength due to large energy dissipation along theinterface. Before proposing the nano-structure formationmechanism it is desired to discuss the theories that have beenproposed for the formation of nano-SiC fiber and particles inSi–C–O system [35–40].

a)

Growth of β-SiC fiber and whisker has been explained byVLS method [35]. Where, V stands for vapor feed gases, Lfor liquid catalyst, and S for solid crystalline material. Theliquid catalyst forms an interface with the crystalline

Fig. 9. SiC nano-pins complex bulk network present in the matrix crack of the set-2.

S. Kumar et al. / Ceramics International 41 (2015) 12849–12860 12857

material to be grown at one side and liquid-vapor interfacein another side. The catalyst is a preferred site for thedeposition of the vapor reactants which causes the liquid tobecome supersaturated. Crystal growth occurs by theprecipitation from the supersaturated liquid at the solid-liquid interface. At sufficiently high temperature, the solidcatalyst particle melts and forms the liquid catalyst ball.Carbon/silicon source gas being absorbed into the liquidcatalyst, which soon get supersaturated and solid SiCprecipitates from the liquid catalyst onto the growthsubstrate. As precipitation continues, the structure grows,lifting the catalyst ball from the substrate and forming awhisker/nano structure [35].

b)

Growth of β-SiC nanowires is also carried out by using theactivated carbon powder and silicon substrate in vacuum at1200–1350 1C for 1–4 h. The following mechanism wasproposed: (1) diffusion of C/CO into silicon substrate, (2)weakening of silicon bond and atomic kick-out, (3) forma-tion of Si–C in vapor phase, (4) formation of saturated SiClayer, (5) formation of pyramid-like SiC nanostructure, and(6) formation of SiC nanowires [39].

c)

Another growth mechanism has also been proposed whereoxygen containing PCS/pyrolyzed PCS decomposes as perthe following reactions (4)–(6) at high temperature and lowpressure [40].

SiCxOy-SiC(s)þSiO(g)þCO(g) (4)

Furthermore, the generation of SiO(g) can react with CO(g)and/or free carbon to form crystals of SiC in the outer regions

SiOþ3CO-SiCþ2CO2 (5)

SiOþC-SiCþCO (6)

Once formed, these crystals of SiC may undergo grain growth.The large SiC grain formation on the outer edges has beenobserved by other also [40,41]. The formation of SiC at outer

surfaces was also observed from SiC/SiO2 reaction [42]. It ismost likely due to the interaction of SiO(g) with carbon.

In the present study the pyrolysis has been carried out undervacuum at 1400 and 1600 1C. At 1400 1C formation of SiO2

was observed (Fig. 4), although the peak intensities are lowwhich show that the SiO2 content might be in traces. In thecomposite, some oxygen might have got inserted into thematrix lattice during thermal molding of the composite in airwhich is evident from the FTIR also (Fig. 2) in the form of Si–O–Si bond peak. However, there is no sign of SiO2 at 1600 1Cas shown in XRD (Fig. 4). This indicates that the SiO2 mightgot consumed between 1400 and 1600 1C. From the micro-structure images (Figs. 7 and 8), it is evident that the nano-pinsare visible only at 1600 1C (set-2).The following sequence of the events/reactions has been

proposed for the growth of the SiC nano-pins at 1600 1C.Some of the events might have occurred in quick successionsor simultaneously.

(1)

Induction of oxygen in limited quantity in the PCS latticeduring the molding of the composite or due to its inherentnature;

(2)

Bulk evolution of CO (g) between 1200 and 1400 1C andsubsequently its reaction with the amorphous SiC to formSixOyCz flux layer locally [43,44];

(3)

Localized evolution of secondary CO (g) between 1400and 1600 1C [43] and lift the softened SixOyCz flux layerlocally and it would act as nucleation for SiC nano-pingrowth (step 1 in Fig. 10a);

(4)

Decomposition of the SixOyCz as per reaction (4) to enrichthe nucleation site with SiC and evaluation of SiO and CO(g) globally;

(5)

Lifting the nucleation site till the residual CO(g) evolved(step 2 in Fig. 10a) and simultaneously precipitation of SiCat the nucleation site and further growth (step 3 inFig. 10a). Growth of the nano-pins continues till evolution

Fig. 10. Growth of nano-pins and their EDAX analysis.

S. Kumar et al. / Ceramics International 41 (2015) 12849–1286012858

of the Co(g) and growth ceases after saturation (step 4 inFig. 10a). In some places branching of nano-pins is alsobeing observed, it is probably due to the merging of twodifferent sources of CO(g).

It is further summarized as: all fully grown nano-SiC pins

are almost of similar size and shape. All the pins having lengthbetween 800–2500 nm and head ball diameter of about 180nm(Fig. 10); Growth of the nano-SiC pins is random and mostlyoriginated from the matrix surface; The body of the nano-SiCpin is carbon rich SiC. The EDAX analysis shows that thenano-pin head comprises about 76% silicon against 18% in thebody (Fig. 10c); The nano-pins are present at the surface aswell as inside the matrix (Fig. 8c).

Since the quantity of CO(g) get evolved between 1400 and1600 1C is low and its evolution is local therefore the size andshape of all the nano-pins are similar (Fig. 10b). Synthesis of

SiC nanoparticles and nano fibers via carbothermal reductionof mesoporous carbon-silica nano composites is reportedelsewhere [45].The composite properties are more improvedwhen SiC particles are in nano-scale range or the morphologyis fiber-like [46]. This also corroborates with the EDAX resultsthat the head of the nano-pins is SiC rich where decompositionof SixOyCz have enriched the head with SiC. It may also beinferred that at 1400 1C, secondary CO(g) does not evolvetherefore nano-pins could not be grown.

4. Conclusions

UD C/SiC composite strips have been fabricated success-fully using the in house developed polycarbosilane following athermal molding method. High temperature vacuum/pressureinfiltration of LSP–PCS could densify the strips in 6 cycles to

S. Kumar et al. / Ceramics International 41 (2015) 12849–12860 12859

obtain composites of the density up to 1.85 g/cc. The hightemperature LPS–PCS infiltration has eliminated the use ofirritating DVB and any solvent which finally results into theporosity and lowers the densification efficiency. The compositepyrolysed at 1600 1C is composed of a mixture of β-SiC andα-SiC phase matrix while it was β-SiC phase at 1400 1C. Thematrix was relatively amorphous at 1200 1C and the tempera-ture was not considered for composite pyrolysis. The high charyield HSP–PCS and 20%DVB resin has resulted into the denseSiC matrix surrounding the fibers. Flexural and tensile strengthwere found to be 700 MPa and 275 MPa respectively for thecomposites pyrolysed at 1600 1C while at 1400 1C the proper-ties were about 20% lower than these values. The nano-pinswere found to be uniformly grown in the set-2. The higherstrength and lower scatter in the set-2 g are attributed due to in-situ growth of SiC nano-pins in the matrix which has cateredan additional crack length/surface. The growth of nano-SiCpins has been explained and a growth mechanism sequenceproposed. The size and shape of the nano-pins was almostsame and originating from the matrix surface. The tip of theSiC nano-pin was found to be silicon rich while the body wasfound to be carbon rich. The parameters of the set-2 would beuseful for the development of UD configuration based C/SiCcomposite structures. In the future studies, TEM/HRTEM ofthe nano-pins would be attempted to investigate the internalfeatures of the nano-pins. Some efforts should also be made infuture studies to increase the length of the nano-pins in order tomake them more effective for ceramic composites.

Acknowledgments

The authors wish to acknowledge the support of Mr. AnilKumar, Head, CMCD/ASL who allowed us to use their CVD/CVI furnace and Head CAF, DMSRDE for TGA and SEMcharacterization of the few composite specimens. The authorswould like to acknowledge the support of Dr. Shrikanth Sc ‘D’of ASL, Hyderabad and Dr. Bhanumathi Sc ‘D’ of DMRLHyderabad for their support in FESEM and XRD characteriza-tion respectively.

References

[1] M. Yuan, Z.R. Huang, S.M. Dong, et al., Microstructure of multilayeredinterphases processed by temperature-pulsing chemical vapor infiltration,Phys. Status Solidi 203 (2006) R58–R60.

[2] K. Jian, Z.H. Chen, Q.S. Ma, et al., Effects of pyrolysis processes on themicrostructures and mechanical properties of Cf/SiC composites usingpolycarbosilane, Mater. Sci. Eng. A 390 (2005) 154–158.

[3] Z.S. Rak, A process for Cf/SiC composites using liquid polymerinfiltration, J. Am. Ceram. Soc. 84 (2001) 2235–2239.

[4] Suresh Kumar, Anil Kumar, Ramesh Babu Mala, Raghvendra RaoMokhasunavisu, Fabrication and ablation studies of 4D C/SiC compositenozzle under liquid propulsion, DOI: 10.1111/ijac.12392.

[5] Paolo Colombo, Polymer-derived ceramics: 40 years of research andinnovation in advanced ceramics, J. Am. Ceram. Soc. 93 (2010)1805–1837.

[6] L. Shilling Jr., J.P. Wesson, T.C. Williams, Polycarbosilane precursorsfor silicon carbide, Am. Ceram. Soc. Bull. 62 (1983) 912–915.

[7] Marc Birot, Jean-Paul Pilot, Jacques Dunogues, Comprehensive chem-istry of polycarbosilanes, polysilazanes, and polycarbosilazanes as pre-cursors of ceramics, Chem. Rev. 95 (1995) 1443–1477.

[8] A.K. Saxena, Ashok Ranjan, Rajesh Kumar Tiwari, G.N. Mathur, Aprocess for the preparation of polycarbosilanes for use as precursor ofceramic materials, Indian Patent No 249015, 2011.

[9] Seyferth, M. Tasi, H.G. Woo, Poly(Vinylsilane), [CH2CH(SiH3)]n:preparation, characterization, and utilization as a preceramic polymer,Chem. Mater. 7 (1995) 236–243.

[10] B.e.i. Yang, Xingui Zhou, Jinshan Yu, The properties of Cf/SiCcomposites prepared from different precursors, Ceram. Int. 41 (2015)4207–4217.

[11] M. Berbon, M. Calabrese, Effect of 1600 1C heat treatment on C/SiCcomposites fabricated by polymer infiltration and pyrolysis with allylhy-dridopolycarbosilane, J. Am. Ceram. Soc. 85 (2002) 1891–1893.

[12] Chunlei Yan, Rongjun Liu, Yingbin Cao, Changrui Zhang, Fabricationand properties of PIP 3DCf/ZrC–SiC composites, Mater. Sci. Eng. A 591(2014) 105–110.

[13] Zheng Luo, Xingui Zhou, Jinshan Yu, F.e.i. Wang, High-performance 3DSiC/PyC/SiC composites fabricated by an optimized PIP process with anew precursor and a thermal molding method, Ceram. Int. 40 (2014)6525–6532.

[14] A. Ortona, A. Donato, G. Filacchioni, U. De Angelis, A. La BarberaC.A. Nannetti, B. Riccardi, J. Yeatman, SiC–SiCf CMC manufacturingby hybrid CVI–PIP techniques: process optimization, Fusion Eng. Des.51–52 (2000) 159–163.

[15] Shuang Zhao, Xingui Zhou, Jinshan Yu, Paul Mummery, Effect of heattreatment on microstructure and mechanical properties of PIP-SiC/SiCcomposites, Mater. Sci. Eng. A 559 (2013) 808–811.

[16] Jin-Chul Bae, Kwang-Youn Cho, Dea-Ho Yoon, Seung-Soo Baek, Jong-Kyoo Park, Jung-Il Kim, Dong-Won Im, Doh-Hyung Riu, Highlyefficient densification of carbon fiber-reinforced SiC-matrix compositesby melting infiltration and pyrolysis using polycarbosilane, Ceram. Int. 39(2013) 5623–5629.

[17] Y.Tange Otoishi, Growth Rate and morphology of silicon carbidewhiskers from polycarbosilane, J. Cryst. Growth 200 (1999) 467–471.

[18] Bingbing Pei, Yunzhou Zhu, Ming Yuan, Zhengren Huang, Yinsheng Li,Effect of in situ grown SiC nanowires on microstructure and mechanicalproperties of C/SiC composites, Ceram. Int. 40 (2014) 5191–5195.

[19] R. Kochendorfer, W. Krenkel, High-temperature ceramic-matrix compo-sites, Ceram. Trans. 57 (1995) 13–22.

[20] Kiyoshi Sato, Atsushi Tezuka, Osamu Funayama, Takeshi Isoda,Yoshiharu Terada, Shinji Kato, Misao Iwata, Fabrication and pressuretesting of a gas-turbine component manufactured by a preceramic-polymer-impregnation method, Compos. Sci. Technol. 59 (1999)853–859.

[21] S. Schmidt, S. Beyer, H. Knabe, H. Immich, R. Meistring, A. Gessler,Advanced ceramic matrix composite materials for current and futurepropulsion technology applications, Acta Astronaut. 55 (2004) 409–420.

[22] S. Ochiai, S. Kimura, H. Tanaka, M. Tanaka, M. Hojo, K. Morishita,H. Okuda, H. Nakayama, M. Tamura, K. Shibata, M. Sato, Residualstrength of PIP-processed SiC/SiC single-tow mini composite exposed athigh temperatures in air as a function of exposure temperature and time,Compos. Part A: Appl. Sci. Manuf. 35 (2004) 41–50.

[23] Ahmad Monshi, Mohammad Reza Foroughi, Mohammad Reza Monshi,Modified Scherrer equation to estmate more accurately nano-crustallinesize using XRD, World J. Nano Sci. Eng. 2 (2012) 154–160.

[24] K. Okamura, T. Shimoo, K. Suzuya, K. Suzuki, SiC Based ceramic fibersprepared via organic-to-inorganic conversion process—a Review, Jpn.J. Ceram. Soc. 114 (2006) 445–454.

[25] J. Lamon, A micromechanics-based approach to the mechanical behaviorof brittle matrix composites, Comp. Sci. Technol. 61 (2001) 2259–2272.

[26] N.P. Bansal, Hand Book of Ceramic Composites, Kluwer AcadamicPublishers, 2005.

[27] K. Suresh, K. Sweety, K. Anil, S. Anupam, A.K. Gupta, G. Rohini Devi,Scr. Mater. 58 (2008) 826–829.

[28] Z.h.u. Yun-zhou, Huang Zheng-ren, Dong Shao-ming, Yuan ming,Jiang Dong-liang, Correlation of PyC/SiC interphase to the mechanical

S. Kumar et al. / Ceramics International 41 (2015) 12849–1286012860

properties of 3D HTA C/SiC composites fabricated by polymer infiltra-tion and pyrolysis, New Carbon Mater. 22 (2007) 327–331.

[29] J.i.e. Yin, Sea-Hoon Lee, L.u.n. Feng, Yunzhou Zhu, Xuejian Liu,Zhengren Huang, Se-Young Kim, In-Sub Han, The effects of SiCprecursors on the microstructures and mechanical properties of SiCf/SiC composites prepared via polymer impregnation and pyrolysisprocess, Ceram. Int. 41 (2015) 4145–4153.

[30] K.e. Jian, Zhao-Hui Chen, Qing-Song Ma, Hai-feng Hu, Wen-Wei Zheng, Effects of polycarbosilane infiltration processes on themicrostructure and mechanical properties of 3D-Cf/SiC composites,Ceram. Int. 33 (2007) 905–909.

[31] K.e. Jian, Zhao-Hui Chen, Qing-Song Ma, Hai-feng Hu, Wen-Wei Zheng, Effects of pyrolysis temperatures on the microstructure andmechanical properties of 2D-Cf/SiC composites using polycarbosilane,Ceram. Int. 33 (2007) 73–76.

[32] K.e. Jian, Zhao-Hui Chen, Qing-Song Ma, Wen-Wei Zheng, Effects ofpyrolysis processes on the microstructures and mechanical properties ofCf/SiC composites using polycarbosilane, Mater. Sci. Eng. A 390 (2005)154–158.

[33] E. Breval, M. Hammond, C.G. Pantano, Nanostructural characterizationof silicon oxycarbide glasses and glass–ceramics, J. Am. Ceram. Soc. 77(1994) 3012–3018.

[34] P. Colombo, A. Martucci, O. Fogato, P. Villoresi, SiC Films by LaserPyrolysis of Polycarbosilane, J. Am. Ceram. Soc. 84 (2001) 224–226.

[35] R.S. Wagner, W.C. Ellis, Vapor liquid solid mechanism of single crystalgrowth, Appl. Phys. Lett. 4 (1964) 89–90.

[36] H.J. Choi, H.K. Seong, J.C. Lee, Y.M. Sung, Growth and modulation ofsilicon carbide nanowires, J. Cryst. Growth 269 (2004) 472–478.

[37] H.O..Pierson, Handbook of refractory carbides and nitrides. Sandia Park,New Mexico, 1996.

[38] Xiumin Yao, Shouhong Tan, Zhengren Huang, Shaoming Dong,Dongliang Jiang, Growth mechanism of β-SiC nanowires in SiCreticulated porous ceramics, Ceram. Int. 33 (2007) 901–904.

[39] K.Y. Cheong and Z. Lockman, Growth Mechanism of Cubic-SiliconCarbide Nanowires, Journal of Nanomaterials, http://dx.doi.org/10.1155/2009/572865.

[40] T. Shimoo, I. Tsukada, T. Seguchi, K. Okamura, Effect of firingtemperature on the thermal stability of low-oxygen silicon carbide fibers,J. Am. Ceram. Soc. 81 (1988) 2109–2115.

[41] P. Le Coustumer, M. Monthioux, A. Oberlin, further-studies of thestability of Pcs-based ceramic fibers at high-temperatures: II effect of all-carbon environments, Br. Ceram. Trans. 94 (1995) 185–190.

[42] N. Jacobson, K. Lee, D. Fox, Reactions of silicon carbide and silicon (IV)oxide at elevated temperatures, J. Am. Ceram. Soc. 75 (1992)1603–1611.

[43] Yoshio Hasegawa, Synthesis of continuous silicon carbide fibre, J. Mater.Sci. 24 (1989) 1177–1190.

[44] G.X. Wang, G.Q. Max Lu, Benyan Pei, A.B. Yu, Oxidation mechanismof Si3N4 bonded SiC ceramics by CO, CO2 and steam, J. Mater. Sci. 33(1998) 1309–1317.

[45] J.F. Yao, H.T. Wang, X.Y. Zhang, Role of pores in the carbothermalreduction of carbon-silica nanocomposites into silicon carbide nanos-tructures, J. Phys. Chem. C 111 (2007) 636–641.

[46] Y.F. Chen, Factors affecting the growth of SiC nano-whiskers, J. Mater.Sci. Technol. 26 (2010) 1041–1046.

本文献由“学霸图书馆-文献云下载”收集自网络，仅供学习交流使用。

学霸图书馆（www.xuebalib.com）是一个“整合众多图书馆数据库资源，

提供一站式文献检索和下载服务”的24 小时在线不限IP

图书馆。

图书馆致力于便利、促进学习与科研，提供最强文献下载服务。

图书馆导航：

图书馆首页 文献云下载 图书馆入口 外文数据库大全 疑难文献辅助工具