polymer-derived ceramics and their space applications · 2020-07-01 · polymer-derived ceramics...

Polymer-Derived Ceramics and Their SpaceApplications

S. Packirisamy, K. J. Sreejith, Deepa Devapal, and B. Swaminathan

ContentsIntroduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4Synthesis, Characterization, and Ceramic Conversion of Preceramic Polymers . . . . . . . . . . . . . . . 6

Silicon Containing Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6Ceramics from Carbonaceous Polymers . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33Active Filler Controlled Pyrolysis (AFCOP) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 35

Polymer-Derived Ceramics for Space Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38Ceramic Matrix Composites . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38Protective Coatings . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51Ceramic Adhesives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 73Lightweight Ceramics . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

Concluding Remarks and Perspectives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 92

S. Packirisamy (*)Ceramic Matrix Products Division, Analytical Spectroscopy and Ceramics Group, PCM Entity,Vikram Sarabhai Space Centre, Indian Space Research Organization, Thiruvananthapuram, India

Advanced Polymeric Materials Research Laboratory, Department of Chemistry and Biochemistry,School of Basic Sciences and Research, Sharda University, Greater Noida, Indiae-mail: [email protected]; [email protected]

K. J. Sreejith · D. DevapalCeramic Matrix Products Division, Analytical Spectroscopy and Ceramics Group, PCM Entity,Vikram Sarabhai Space Centre, Indian Space Research Organization, Thiruvananthapuram, Indiae-mail: [email protected]; [email protected]; [email protected]

B. SwaminathanCeramic Matrix Products Division, Analytical Spectroscopy and Ceramics Group, PCM Entity,Vikram Sarabhai Space Centre, Indian Space Research Organization, Thiruvananthapuram, India

Morgan Advanced Materials, Murugappa Mogan Thermal Ceramics Ltd., Ranipet, Vellore District,Tamil Nadu, Indiae-mail: [email protected]

© Springer Nature Switzerland AG 2020Y. Mahajan, R. Johnson (eds.), Handbook of Advanced Ceramics and Composites,https://doi.org/10.1007/978-3-319-73255-8_31-2

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https://doi.org/10.1007/978-3-319-73255-8_31-2

Abstract

Inorganic and organometallic polymers capable of giving high ceramic residue(more than 50 wt%) on heat treatment in an inert atmosphere are called “pre-ceramic polymers.” As they are polymeric in nature, processing techniques usedfor conventional polymer processing can be easily adopted. They can be appliedas coating, cast into film and drawn into fiber and then converted intocorresponding ceramic material. Amorphous materials that are thermally stableto very high temperatures with compositions not obtainable with common syn-thetic methods can be obtained from preceramic polymers. Kinetic stabilizationof less stable phases, adaptability of various fabrication capabilities of polymerprocess engineering, formation of nanoceramics of desired composition, pres-sureless sintering, and machinability are the main advantages of obtainingceramics from polymeric precursors.

Polymer-derived ceramics find applications as oxidation resistant high tem-perature ceramic materials in the form of fiber, coatings and adhesives, and matrixof ceramic matrix composites for use by aerospace, nuclear, and defense estab-lishments. In addition, they are also being investigated for end-use in biomedicaldevices, drug delivery systems, water remediation, energy storage devices, micro-electronics, and nanosensors.

The present chapter deals with synthesis, characterization, and ceramic con-version of silicon-based preceramic polymers, and ceramics from carbonaceouspolymers, and their possible space applications. In view of the voluminousliterature, equal emphasis could not be given to many of the developments inthe area of preceramic polymers and the discussion is confined to relevantsystems which have the scope for space applications.

Keywords

Polymer-derived ceramics · Preceramic polymers · Space applications · Thermalprotection systems · Ceramic matrix composites · Oxidation resistant coatings ·Thermal barrier coatings · Atomic oxygen resistant coatings · Ceramicadhesives · Lightweight ceramics

Abbreviations

3-APTES 3-AminopropyltriethoxysilaneABSE Amino-bis(silylethane)AFCOP Active Filler Controlled Polymer PyrolysisAO Atomic OxygenAPMDEOS 3-AminopropylmethyldiethoxysilaneBCTS Boron-Modified CTSBLS Boundary Layer SplitterBMG Methyl and Glycidoxy Group Containing Poly(borosiloxane)BMV Methyl and Vinyl Group Containing Poly(borosiloxane)BN Boron Nitride

2 S. Packirisamy et al.

BP Poly(phenylborosiloxane)BPV Phenyl and Vinyl Group Containing Poly(borosiloxane)BSAS Barium-Strontium-AluminosilicateCMC Ceramics Matrix CompositeCNT Carbon Nano TubeCTE Coefficient of Thermal ExpansionCTS 1,3,5-Trimethyl-10,30,50-trivinylcyclotrisilazaneCVD Chemical Vapor DepositionCVI Chemical Vapor InfiltrationDMDCS DimethyldichlorosilaneDPDCS DiphenyldichlorosilaneEBC Environmental Barrier CoatingEC Eddy CurrentFESEM Field Emission Scanning Electron MicroscopyFM Fiber to MatrixGPTMOS GlycidoxypropyltrimethoxysilaneHDA High Density AblativeHFH Heat Flux HistoryHRTEM High Resolution Transition Electron MicroscopyILSS Interlaminar Shear StrengthIMI Internal Multiscreen InsulationKHS Kinetic Heat SimulationLAM Liquid Apogee MotorLEO Low Earth OrbitLSI Liquid Silicon InfiltrationMAS-NMR Magic Angle Spinning-Nuclear Magnetic ResonanceMAX Phase In MAX phase M is an early transition metal, A is an A-group

element, and X is either carbon and/or nitrogenMEMS Micro-Electromechanical SystemsMTEOS MethyltriethoxysilaneMTMOS MethyltrimethoxysilaneMVDCS MethylvinyldichlorosilaneMWCNT Multi Wall Carbon Nano TubeNMR Nuclear Magnetic ResonancePBDPS PolyborodiphenylsiloxanePBS PolyborosiloxanePCS PolycarbosilanePDC Polymer-Derived CeramicPHPS PerhydridopolysilazanePIP Polymer Infiltration and PyrolysisPSH PolysilahydrocarbonPTEOS PhenyltriethoxysilanePTMOS PhenyltrimethoxysilanePVD Physical Vapor DepositionRLV-TD Reusable Launch Vehicle-Technology Demonstrator

Polymer-Derived Ceramics and Their Space Applications 3

RMI Reactive Melt InfiltrationSEM Scanning Electron MicroscopySWLE Side Wall Leading EdgeTBC Thermal Barrier CoatingTEM Transmitting Electron MicroscopyTEOS TetraethoxysilaneTHF TetrahydrofuranTMOS TetramethoxysilaneTPhBS TitanophenylborosiloxaneTPS Thermal Protection SystemUHTC Ultra High Temperature CeramicVTEOS VinyltriethoxysilaneVTMEOS Vinyltris(2-methoxyethoxy)silaneWLE Wing Leading EdgeXRD X-Ray DiffractionZBS ZirconoborosiloxaneZPhBS ZirconophenylborosiloxaneZTBS Zirconotitanoborosiloxane

Introduction

The ever-increasing demand for more efficient and powerful aircraft and spacecraftengines which have to operate in extreme operating conditions such as high tem-peratures, fast speeds, high stresses and hostile environment has fueled the demandfor developing and qualifying newer engineered ceramics/advanced ceramic mate-rials [1, 2]. The feasibility of combining many metallic and nonmetallic atoms toform ceramics in several atomic structural arrangements provides enormous scope tomaterial scientists to invent new ceramic materials to meet the requirements ofaerospace industry [1].

Among the various routes available for obtaining advanced ceramics, polymer-derived ceramic (PDC) route offers the most convenient and easily adoptablemethod for obtaining advanced ceramics in desired composition and shape. Theadvancement of PDCs owes its origin to the pyrolysis of polymers to obtain solidmaterials such as graphite, glassy carbon, carbon fibers, and carbon-carbon compos-ites. Extension of this concept to inorganic and organometallic polymers led to thedevelopment of PDCs as fibers, composites, coatings, adhesives, and cellular mate-rials. The journey of PDCs began with Chantrell and Popper [3], when theydemonstrated for the first time that monolithic silicon nitride could be prepared bymolding polysilazane to the desired shape followed by pyrolysis. However, this fastemerging field got impetus only by mid-1970s when Verbeek et al. [4] in Germanyreported the preparation of non-oxide ceramic fiber comprising of silicon carbideand silicon nitride by pyrolyzing polysilazane and polycarbosilazane, and Yajimaand coworkers [5] in Japan prepared polycarbosilane (PCS) from polydimethylsilaneand evaluated it as precursor for silicon carbide fiber.


The production of ceramics by pyrolysis of polymeric precursors presents manyadvantages over the conventional methods of ceramic preparation [6]. The advan-tages include low temperature processability, high purity of the ceramics, the abilityto control the composition of the ceramics, easiness in the preparation of shapedobjects, and the ability to generate products in a variety of forms such as coatings,films, joints, monolithic bodies, foams, and fibers (Fig. 1). Over the years, PDCshave gained importance, as evidenced by the number of research publications,reviews, patents, international conferences and workshops conducted and bookspublished relating to this topic [7–13]. From the commercial point of view, theglobal market of PDCs in 2017 is only USD 437.16 million [14], very much lesscompared to that of a well-known polymer, polyethylene (USD 163 billion). Atpresent, the market of PDCs is primarily due to the demand for exploring high-endapplications in aerospace, defense, energy, and semiconductor sectors, and themarket survey predicts a compounded annual growth rate of 10.2% during the period2017–2022 [14]. Once the enabling technologies based on PDCs are inducted intoactual applications in the above sectors, the demand for PDCs is expected to grow ata much faster rate.

The present review highlights the developments on PDCs with a special focus onthe following space applications: (i) matrix resins for CMCs required for thermo-structural applications, (ii) protective coatings for nonmetallic and metallic sub-strates, and atomic oxygen resistant coatings, (iii) ceramic adhesives for joiningdifferent substrates, and (iv) lightweight ceramics for advanced thermal protectionsystems. This review is not an attempt to present a detailed account of the develop-ments in PDCs, and in view of the voluminous research work in this area, readersattention is drawn to excellent reviews, book chapters, and books on this topic.

Fig. 1 Polymer-derived ceramic (PDC) route to ceramic products. (Adopted from Ref. [6])


Synthesis, Characterization, and Ceramic Conversionof Preceramic Polymers

The composition, number of phases, phase distribution, and the microstructure of thefinal ceramic obtained are largely influenced by the type and the molecular structureof preceramic polymer. In view of this, research efforts toward synthesizing novelpreceramic polymers with the desired physicochemical properties, optimizing thethermolysis conditions to maximize the ceramic yield, and improving the hightemperature capability and environmental stability of ceramics obtained therefromcontinue unabated.

The transformation of preceramic polymer to ceramics involves mainly thefollowing three steps (Fig. 2): (i) shaping and cross-linking of these precursors toform two- or three-dimensional preceramic networks, (ii) transformation of thenetworks into amorphous covalent ceramics (organic/inorganic transition) by heattreatment, and (iii) crystallization of the amorphous solids into thermodynamicallystable phases via different metastable intermediates [6, 7].

Silicon Containing Polymers

This class of polymers is the most widely studied precursors for ceramics, and this isattributed to easy availability of variety of organosilicon monomers in large quan-tities at an affordable cost. Different Si-containing polymers having the generalformula, –(RR1Si-X)– can be synthesized by suitable choice of monomers andsynthetic routes as precursors for SiC, SiBC, SiOC, SiOBC, MOSiOC (where Mstands for metallic element), Si3N4, SiCN, and SiBCN ceramics (Fig. 3).

Polysilanes and PolycarbosilanesFollowing the pioneering work of Yajima et al. [5] on the conversion of intractablepolydimethylsilane to PCS (Scheme 1), the initial efforts were focused toward

Fig. 2 Conversion ofpreceramic polymer toceramic. (Adopted fromRef. [7])


Fig.3

Ceram

icconv

ersion

ofsilicon

containing

polymers.(A

dapted

from

[7])


obtaining SiC fibers and modifying PCS to incorporate metallic elements such Al,Ti, and Zr to improve the high temperature capability of SiC fibers [15, 16].

In 1980s, PCS gained its significance not only as precursors for SiC fibers ormetal modified SiC fibers but also for many other applications including that ofmatrix resins for CMCs. The research work on PCS has been carried out over theyears with the aim of achieving one or more of the following objectives: (i) obtainingPCS with near stoichiometric composition [17], (ii) cross-linking/modification/curing to improve the ceramic residue [18], (iii) minimizing oxygen incorporationduring curing [19], (iv) catalyst systems for pressureless synthesis of PCS [20], and(v) chemical modification of PCS to obtain new ceramic systems [21, 22]. Thoughthe synthesis of PCS from intractable polydimethylsilane was reported in 1975, theinterest on this precursor is not diminishing as indicated by the number of patentsand papers published in recent years.

Extensive studies on the ceramic conversion of PCS [23, 24] and metal modifiedPCS have been reported [25, 26]. PCS on heat treatment in inert atmosphere givesamorphous SiC at around 800 �C which on further heat treatment at temperaturesabove 1200 �C gets converted to β-SiC (Fig. 4) with crystallite size of 3–10 nm [22].If the starting PCS is rich in carbon, the PDC obtained is β-SiC with nanodomains ofcarbon dispersed in the ceramic matrix. In the case of metal (M) modified PCS, thePDC obtained is usually a mixed non-oxide ceramic comprising of metal carbide andβ-SiC or SiMCO [Fig. 5].

The research work on polysilastyrene, a soluble polysilane (Scheme 2), by Westet al. [27] and soluble vinyl-functionalized polysilanes and branched polysilanes/polycarbosilane (Scheme 3) by Schilling et al. [28] has motivated many researchersto synthesize soluble precursors directly without the need for converting polysilaneto polycarbosilane. Many such precursors gave poor ceramic yield, and cross-linkingby UV irradiation or by chemical means was required to improve the ceramic yield[29]. Schilling et al. [28] observed that branching at silicon improved the ceramicyield and reported the synthesis of branched or vinyl-functionalized silanes, poly-carbosilanes, and polysilahydrocarbons (PSH) through the reaction oforganochlorosilane monomers with vinylic monomers or through the reaction ofvinyl-functionalized chlorosilanes with organochlorosilane monomers under dechlo-rination conditions (Scheme 3).

Packirisamy et al. [30–33] synthesized PSHs through the dechlorination ofdiorganodichlorosilane using sodium in the presence of styrene (Scheme 4). Ingeneral, PSHs gave poor ceramic residue (5–25%), and some of them containing

Scheme 1 Yajima’s processfor the synthesis ofpolycarbosilane frompolydimethylsilane [5]


Fig. 4 XRD pattern of ceramics obtained from PCS at 1500 �C [22]

Fig. 5 XRD of PCS modified with Zr(acac)4. (Reproduced from Ref. [21] with permission fromSpringer Nature)


Scheme 2 Synthesis of poly(dimethylsilylene-co-methylphenylsilylene) [polysilastyrene] [27]

Scheme 3 Synthesis of (a) vinyl-functionalized polysilane, (b) polycarbosilane, and (c) branchedpolysilane [28]

Scheme 4 Synthesis of polydisilahydrocarbon from diorganodichlorosilane and styrene [30, 34]


polysilyl blocks and methylphenylsilyl units were converted to tractable poly-carbosilanes through heat treatment at 420 �C in a partially confined environmentand these polymers gave ceramic residue of 40–85% at 900 �C (Fig. 6) [34]. Thehigh ceramic residue (85%) of PCS from PSH containing methylphenylsilylene isattributed to phenylene insertion reaction (Scheme 5).

Polysiloxanes and PolysilsesquioxanesPolysiloxanes (Fig. 7) are the most studied ones among inorganic polymers[35]. Hydrolysis of diorganodichlorosilanes (R2SiCl2) is widely used for the syn-thesis of polysiloxanes. Synthesis of polysiloxanes by ring opening polymerizationof cyclic trimers and tetramers has largely replaced the hydrolysis approach[36]. Depending on the functionality of the side chains, polysiloxanes can be linearor cross-linked networks.

Cross-linked polysiloxanes [36] and polysilsesquioxanes (Fig. 8) have beenprepared by the sol-gel process through hydrolysis and condensation reactions oforganically modified silicon alkoxides [36, 37].

Polysiloxanes and polysilsesquioxanes on pyrolysis give siliconoxycarbide(SiOC) glasses [38]. Compared to fused SiO2, carbon-rich SiOC glasses havesuperior mechanical properties at elevated temperatures, in particular, improved

0 100 200 300 400 500 600 700 800 900 10000

20

40

60

80

100

4

32

1)

% (thgi eW

Temperature (oC)

Fig. 6 Thermogravimetriccurves of: 1. PSH containingdimethyldisilyl units (PSH-1),2. PSH containingdimethyldisilyl andphenylmethyldisilyl units(PSH-2), 3. PSH-1 heattreated at 420 �C, and4. PSH-2 heat treated at420 �C [34]

Scheme 5 Phenylene insertion taking place in PSH containing dimethyldisilyl andphenylmethyldisilyl units (PSH-2) during heat treatment [34]


creep resistance [39]. By the addition of ceramic or metallic particles, the mechanicaland physical properties of SiOC matrix can be modified [40].

The microstructural evolution of SiOC glasses from polysiloxanes and poly-silsesquioxanes has been studied in detail by many researchers [41–46]. Duringheat treatment at temperatures above 900 �C, redistribution reactions take placebetween Si-O and Si-C bonds as given in the general equation below [46]:

SiOxC4�x þ SiOyC4�y ¼ SiOx�1C5�x þ SiOyþ1C3�y

where x and y are integers, 0 < x � 4, and 0 � y < 4.

Fig. 7 Polysiloxane structure

Fig. 8 General types of poly(organosilsesquioxane)s. (Reproduced from Ref. [37] with permissionfrom American Chemical Society)


The free carbon phase formed at 600 �C remains unchanged up to 1350 �C asevidenced by 29Si-NMR spectra and chemical analysis. Investigation by Saha andRaj [44] suggests that the temperature regimes at which SiCO crystallizes into β-SiCoverlaps with the conditions where a carbothermal reaction between silica andcarbon as shown below takes place:

SiO2 að Þ þ 3C grð Þ ! SiC βð Þ þ 2CO gð Þ "Thus, SiOC may crystallize in two ways: (i) the amorphous structure may phase

separate into stoichiometric compositions of SiO2, SiC, and C and then crystallize;and (ii) a second pathway is the growth of β-SiC accompanying the carbothermalreduction of silica, as given above. Crystallization may occur by a combination ofdecomposition as well as phase separation. At temperatures above 1400 �C, the silicaphase can react with SiC to volatilize into SiO and CO as shown below:

2SiO2 að Þ þ SiC βð Þ ! 3SiO gð Þ " þCO gð Þ "Saha et al. [43] proposed a simple geometrical model based on surface-to-volume

ratio of the domains which was moderately successful in predicting the domain sizein SiOC (Fig. 9). The principal assumptions in this model are based on the frame-work of the amorphous structure justified on the basis of NMR data and on the creepand viscoelastic behavior of the SiOC ceramic at temperatures beyond 1000 �C.Subsequent to proposing this model, it was shown that amorphous SiOC ceramicscan be etched with hydrofluoric acid leaving behind a porous structure with a median

Fig. 9 Proposed model forthe nanodomain structures inpolymer-derived SiCO.(Reproduced from Ref. [43]with permission from JohnWiley and Sons)


pore diameter of approximately 3 nm. The change in the composition after etchingshows that only silica is removed.

PolyborosiloxanesPolyborosiloxanes have been an exploration target because of their thermodynam-ically stable backbone linkage. Though polyborosiloxanes could not be used as hightemperature polymers due to poor hydrolytic stability, they serve as cost-effectiveprecursors for siliconboronoxycarbide (SiBOC) ceramics. They are also investigatedin sol-gel chemistry as precursors for borosilicate glasses. Several methods havebeen described for the synthesis of polymers containing -Si-O-B-O-linkages usingdi- and tri-functional silicon compounds as the source of silicon, and boric acid,borontrioxide, borontrichloride, boron tribromide, borate esters, and phenylboronicacid as the source of boron [47–50].

Yajima et al. [48] synthesized poly(borodiphenylsiloxane) (PBDPS) by reactingboric acid with diphenyldichlorosilane (DPDCS) (Scheme 6). PBDPS was obtainedin 92% yield and found to be far more moisture resistant than its analog preparedfrom boric acid and dimethyldichlorosilane [48]. PBDPS was used as a catalyst forthe conversion of intractable polydimethylsilane to polycarbosilane [51]. It was alsoevaluated as a binder for making SiC and Si3N4 sintered bodies [52]. Synthesis ofPBDPS through the reaction of boric acid with diphenylsilanol in dibutylether orN-methyl-2-pyrrolidone has also been reported [53].

Sol-gel process was extensively used with the aim of preparing, viacocondensation reactions between TEOS and triethylborate or boric acid, a morehomogeneous SiO2-B2O3 gel as precursor for borosilicate glass and fibers [49, 50,54]. It is found that the boron atoms are present in the silica matrix as a separate B2O3

and/or B(OH)3 phase [55]. This is attributed to the well-known reactivity of theB-O-Si bonds toward hydrolysis. The borosiloxane linkages that might form in theearly stages of the sol-gel process are consumed at the end when increasing amountof water is produced via condensation [54]. Soraru et al. [56] demonstrated that if amodified silicon alkoxide, such as RSi(OEt)3, where R¼Me, Et, or Vi iscohydrolyzed with tetraethylborate, then a large amount of trigonal BO3 units areincorporated in the hybrid siloxane-network via B-O-Si bridges, leading toborosiloxane gels. The resulting gels appear homogeneous, clear, and no evidenceof the formation of a crystalline B(OH)3 can be detected. The retention of B-O-Si

ClCl Si +OH

B

OH

OHn-butylether

300°CSiA O B

O

O

A

An

Scheme 6 Synthesis of poly(borodiphenylsiloxane) [48]


bonds in the hybrid borosilicate gels is partly due to: (i) the low hydrolysis ratio usedin the sol-gel synthesis (H2O)/Si ¼ 1.5), (ii) the in situ protection that the hydro-phobic organic groups bonded to silicon provide against water attack, and (iii) thehigher electron density at the silicon atoms in the organically substituted gels whichstrengthen the Si-O bonds.

Hydrolysis and condensation reactions in an anhydrous sol-gel system compris-ing TEOS, boric acid, and ethanol were studied by Zha et al. [57]. Boric acid is ableto hydrolyze TEOS directly, and subsequent condensation reactions formborosiloxane and siloxane linkages. Since water is absent from the initial mixture,the usual first step (as shown below) in a silica-based sol-gel reaction, hydrolysis ofthe silicon alkoxide, cannot occur.

� Si ORð Þ þ H2O !� Si OHð Þ þ ROH

Zha et al. [57] prepared borosiloxane sol by reacting boric acid and TEOS inethanol in multiple steps in order to achieve uniformity of the sol. The anhydrousmixture is stable indefinitely against gelation, but can be readily gelled by addition ofNaOH in ethanol, and this system is useful for preparing borosilicate glasses at lowertemperatures with good homogeneity. In a similar approach, Soraru et al. [58]prepared borosilicate gel by dissolving boric acid in organoalkoxysilane understirring at 78 �C without any external addition of water and then leaving the solutionleft open for gelation. 11B-MAS NMR spectroscopic studies of dried gels confirmedthe incorporation of boron atoms into the siloxane network via B-O-Si bridges. Onthe other hand, when a conventional tetrafunctional silicon alkoxide such as tetra-methoxysilane (TMOS) or TEOS is reacted with boric acid, the expected borosilicategels are not formed.

In a modification of the above approach, Ambadas et al. [34, 59] synthesizedborosiloxane oligomers through the condensation of boric acid with phenyltri-methoxysilane (PTMOS) and vinyltriethoxysilane (VTEOS) in diglyme (Scheme7). The extent of reaction was controlled by the removal of alkyl alcohol formed asthe by-product and the oligomers were soluble in diglyme. The advantage of thisapproach is that resinous borosiloxane oligomers are obtained in 3–5 h withoutwaiting for days for gelation. The oligomers, poly(vinylborosiloxane) (BSiO-I) and

R1

OR

Si

OR

OR

OH

OH B

OH

+Diglyme,HCl 150-160°C-3 ROH

A

O

Si OR1 BO

OO

Where R1 = Phenyl or Vinyl R = Methyl or Ethyl

Scheme 7 Synthesis of poly(borosiloxane)s from boric acid and organotrialkoxysilane [59]


poly(phenylborosiloxane) (BSiO-II), were soluble in common organic solvents suchas tetrahydrofuran, 1,4-dioxane, toluene, and diglyme. BOSi-I and BOSi-II gaveceramic residue of 85% and 65%, respectively, at 900 �C (Fig. 10).

Following the above work, synthesis of several poly(borosiloxane)s by reactingboric acid with different organotrialkoxysilanes and their mixtures was investigated[60–66] with a view to understanding the effect of organic substituent on silicon onthe processability, molecular weight, microstructure, thermal stability, ceramic res-idue, and the stability of ceramic phase at high temperatures. Devapal et al. [63]observed that T3 structures are formed more when boric acid is reacted withphenyltrimethoxysilane than with phenyltriethoxysilane (Fig. 11). The ratio ofpeak intensities of T3 to T2 (Fig. 11) is 1:0.34 for BSiPh-1 and 1:0.68 for BSiPh-2.This shows that T3 structures are formed more when PTMOS is used which isattributed to the higher reactivity of PTMOS compared to that of PTEOS.

When the synthesis of borosiloxane oligomers is carried out in the absence ofhydrochloric acid, hydrolysis of Si–OR to Si–OH and the subsequent reaction ofSi–OH with B–OH forming Si–O–B bond or its self-condensation forming Si–O–Sibond are ruled out, and the most predominant reaction would be the reaction ofSi–OR with B–OH forming Si–O–B bond with the elimination of ROH. Thisprocess has been successfully used for the synthesis of epoxy-, vinyl-, and epoxy-and vinyl- (in the same backbone) functionalized borosiloxane oligomers eliminat-ing premature gelation [60].

The purpose of incorporating epoxy and vinyl group in borosiloxane oligomers isto provide additional cross-linking sites which can be made use of during processingof these oligomers as matrix resins for composites and infiltrating resins. However,while attempting to remove the solvent, dioxane/diglyme, premature gelation wasobserved. In order to overcome this problem, the nonaqueous sol-gel process wasfurther modified by eliminating the use of solvent [66] and this process is termed as

Fig. 10 TG curves of poly(vinylborosiloxane) [BSiO-I] and poly(phenylborosiloxane) [BSiO-II].(Reproduced from Ref. [59] with permission from Springer Nature)


Fig.1

129SiN

MRspectraandstructuralun

itspresentinpo

ly(pheny

lborosilo

xane)ssynthesizedfrom

boricacid

andPTMOS/PTEOS.(Reprodu

cedfrom

Ref.

[63]

with

perm

ission

from

SpringerNature)


“solventless non-aqueous sol-gel process.”Molecular weight and ceramic residue ofsome of the poly(borosiloxane)s synthesized by reacting boric acid with organoalk-oxysilane or mixtures of organoalkoxysilanes are given in Table 1. It is obvious thatusing a mixture of VTEOS and MTEOS is advantageous as it gives borosiloxaneoligomer having both higher ceramic residue and molecular weight.

Schiavon et al. [67] synthesized novel poly(borosiloxanes), with a B/Si molarratio of 0.2 through a dehydrogen coupling reaction of 1,3,5,7-tetramethyl-1,3,5,7-tetracyclo-tetrasiloxane (D4H) and boric acid, and by hydrolysis/condensation reac-tions of VTEOS and boric acid, followed by a hydrosilylation reaction with D4H(Scheme 8). The ceramic yields for PBS-I and PBS-II at1000 �C were 90.8 wt% and89.8 wt%, respectively.

Rubinsztajn reported [68] a new synthetic procedure, viz., dehydrocarbon con-densation for the preparation of poly(borosiloxane). Trimethyl borate is reacted withdiphenylsilane in the presence of tris(pentafluorophenyl)borane in heptane leading tothe formation of Si–O–B linkages with the release of methane (Scheme 9). Thispolymer gave ceramic residue of 54% (at 800 �C in nitrogen atmosphere).

Polyborosiloxanes synthesized from boric acid or its derivatives have boroncontent of maximum 6–7 wt%. Attempts to incorporate higher boron content byincreasing the boric acid to alkoxysilane ratio beyond a certain limit have not beensuccessful as boric acid precipitates from the reaction mixture on standing. One wayof improving boron content is by incorporating boron-hydride cage structures in theborosiloxane network. This has been achieved by reacting decaborane (14) [B10H14]with 3-aminopropylmethyldiethoxysilane (APMDEOS) in diglyme to obtainethoxysilyl-functionalized adduct and then reacting it with boric acid [60] in tetra-hydrofuran (THF) at 65 �C. It is noticed that the adduct obtained by reactingdecaborane and APMDEOS is soluble in THF, whereas the product obtained afterthe reaction with boric acid is insoluble in organic solvents. This oligomer has boroncontent of 17.5% and gives a ceramic residue of 86% (boron content of 16.7%) at900 �C in inert atmosphere.

Table 1 Molecular weight data and ceramic residue of poly(borosiloxane)s synthesized from boricacid and organoalkoxysilanes (1:2 mole ratio)

Alkoxysilane(s)Mw Mpeak Mn

PDICeramic residue (wt%) at 900 �C Reference

MTEOS 2160 1460 1760 1.48 84.0 61

PTEOS 1950 1320 1440 1.35 69.0 60

VTEOS 3010 1000 1130 3.01 86.2 60

VTEMOS 2160 1360 1430 1.6 83.7 61,64

GPTMOS 4070 2020 1600 2.54 42.0 60

VTEOS + PTEOS(1:1 mole ratio)

1560 1260 1040 1.24 81.0 62

VTEOS + MTEOS(1:1 mole ratio)

9840 2360 2420 4.07 87.0 62

VTEOS + GPTEOS(1:1 mole ratio)

1810 2840 1760 1.61 56.0 62


Ceramic conversion studies of poly(borosiloxane)s: Schiavon et al. [69] inves-tigated the evolution of ceramics from SiBOC glasses and compared them withSiOC glasses using XRD, and 29Si-MAS-NMR and 11B-MAS-NMR spectroscopictechniques. SiBOC glass was obtained by pyrolysis of borosiloxane gel prepared bysol-gel process from MTEOS and boric acid, and SiOC glass was obtained bypyrolysis of siloxane gel prepared by sol-gel process of MTEOS. Based on theceramic yield and the chemical analysis it is concluded that SiOC and SiBOC glassesare thermally stable in the temperature range from 1000 �C to 1500 �C. The weightloss is very much limited up to 1500 �C, and the elemental composition is constant,suggesting that the carbothermal reduction between the SiO2-rich and carbon-richdomains is not active. Thus, the progressive formation of β-SiC nanocrystals, whichhas been revealed by the XRD study (Figs. 12 and 13) for the SiOC and SiBOC

Scheme 8 Synthesis of polyborosiloxanes from cyclic hydridosiloxane (D4H) and boric acid.(Reproduced from Ref. [67] with permission from Elsevier)


glasses, cannot be ascribed to a decomposition process. It has to be related to anucleation and crystallization phenomenon that occurs without compositionalchanges similar to the type of evolution well-known for ternary SiOC glasses.

The phase separation into SiC4 and SiO4 domains occurs through redistributionreactions of Si–C and Si–O bonds as shown below:

SiC3Oþ SiCO3 ! SiC4 þ SiO4

This mechanism is quite local with respect to the silicon atoms and can easilyaccount for the nucleation of very small β-SiC crystals (18 Å at 1500 �C), such asthose observed for the ternary SiOC glass. The growth rate of SiC in SiOC glasses isvery low because of the extremely high viscosity of the SiOC glass which can be upto 2 or 3 orders of magnitude higher than pure SiO2 glass. Such a high viscositysuggests an extremely low mobility of silicon and carbon atoms in the SiOCnetwork. The XRD study shows that the addition of boron atoms to the SiOCnetwork increases the crystallization kinetics of β-SiC (Figs. 12 and 13).

Based on 29Si- and 11B-MAS NMR spectral studies, Schiavon et al. [69]suggested that the SiBOC glasses produced at 1200 �C are formed by a homoge-neous amorphous network based on mixed siliconoxycarbide and boronoxycarbideunits. From 1200 �C to 1400 �C, that is, during the crystallization of β-SiC, themixed BC2O and BCO2 units are consumed and B(OSi)3 units are formed. Thisevolution is more pronounced for the SiBOC glass with the highest amount of boron(B/Si ¼ 0.3). The formation of SiC nuclei occurs through redistribution reactions ofSi–C, B–C, Si–O, and B–O bonds that involve, at the same time, mixed siliconoxycarbide (SiCxO4–x, 0 � x � 4) and boron oxycarbide (BCyO3–y, 0 � y � 2) unitsas shown below:

SiC3Oþ BCO2 ! SiC4 þ BO3

Scheme 9 Synthesis of polyborosiloxane by dehydrocarbon condensation. (Reproduced fromRef. [68] with permission from Springer Nature)


At the end of this process, boron atoms are present in triagonal borosilicate sites,B(OSi)3, and the amorphous matrix, in which the SiC nanocrystals are dispersed, isregarded as a borosilicate-based matrix [69].

Researchers at the Vikram Sarabhai Space Centre [61–65, 70, 71] investigated theceramic conversion at 900 �C, 1500 �C, and 1650 �C of borosiloxane oligomerssynthesized by nonaqueous sol-gel process and by solventless nonaqueous sol-gelprocess for different monomer feed ratios (described earlier). Varying the monomerfeed ratio is expected to influence B/Si ratio in the oligomer which in turn wouldinfluence the concentration and distribution of silicon oxycarbide units, boronoxycarbide units, and nanodomains of carbon. The formation of β-SiC nano-crystallite in these systems is explained in line with the exchange reactions proposedby Schiavon et al. [69]. The basic difference between the polyborosiloxane preparedby sol-gel process and the nonaqueous sol-gel process is that in the latter case B-O-Silinkages are formed in the initial (sol) stage itself which could influence the crystal-lization kinetics.

Angle (2Θ)

a b

c

1500 °C

β-SiC

1400 °C

1300 °C

1200 °C

1500 °C

β-SiC

1400 °C

1300 °C

1200 °C

1500 °C

β-SiC

1400 °C

1300 °C

1200 °C

Angle (2Θ)10 20 30 40 50 60 70 80 10 20 30 40 50 60 70 80

Angle (2Θ)10 20 30 40 50 60 70 80

Fig. 12 Evolution of XRD spectra with pyrolysis temperature for (a) pure SiOC glass, (b) SiBOC,B/Si ratio 0.1, and (c) SiBOC, B/Si ratio 0.3. (Reproduced from Ref. [69] with permission fromJohn Wiley and Sons)


XRD patterns of ceramics obtained at 900 �C, 1500 �C, and 1650 �C frompolyphenylborosiloxane and polyvinylborosiloxane are given in Fig. 14.

It is seen that up to 1500 �C, the ceramic residue remains amorphous and formsβ-SiC at 1650 �C. When the temperature is increased from 1500 �C to 1650 �C, thecarbothermal reduction between SiO2-rich and carbon-rich domains of theoxycarbide glass is more predominant resulting in the formation of crystalline SiCand volatile species (SiO, CO, and CO2) [72] as given below:

2CSiO3 þ 3C ! SiCþ SiO " þ3CO " þCO2 "In variance with the reported observation [63] that poly(phenylborosiloxane)

synthesized from PTEOS and boric acid by nonaqueous sol-gel process(in diglyme using HCl as catalyst) does not form β-SiC at 1500 �C, Sreejith [62]observed that SiBOC obtained from poly(phenylborosiloxane) synthesized by

Angle (2Θ)

10 hrs

2 hrs

1 h

c

a b

β-SiC

10 20 30 40 50 60 70 80Angle (2Θ)

10 hrs

2 hrs

1 h

10 20 30 40 50 60 70 80

β-SiC

Angle (2Θ)

10 hrs

2 hrs

1 h

10 20 30 40 50 60 70 80

Fig. 13 Evolution of XRD spectra with time during isothermal annealing at 1300 �C for (a) pureSiOC glass, (b) SiBOC, B/Si ratio 0.1, and (c) SiBOC, B/Si ratio 0.3. (Reproduced from Ref. [69]with permission from John Wiley and Sons)


solventless nonaqueous sol-gel process undergoes exchange reactions andcarbothermic reduction at 1500 �C resulting in the formation of β-SiC under similarheat treatment conditions. It is also observed that boric acid:PTEOS ratio in themonomer feed influences the crystallization characteristics at 1500 �C.

Sreejith [62] also studied the ceramic conversion of borosiloxane oligomersobtained from boric acid and mixtures of organoalkoxysilanes [(PTEOS+VTEOS)(BPV); MTEOS+VTEOS (BMV) and MTEOS+GPTMOS (BMG)] for varyingmonomer feed ratios synthesized by solventless nonaqueous sol-gel process. It isobserved that the crystallization characteristics of borosiloxane oligomers obtainedfrom mixtures of organoalkoxysilane appear to be different from that of borosiloxaneoligomers from the corresponding organoalkoxysilanes particularly at 1500 �C.Unlike polyphenylborosiloxane, poly-vinylborosiloxane, and BPV systems, forBMV and BMG systems, SiBOC ceramic obtained by pyrolysis at 900 �C wasquite stable at 1500 �C and even at 1650 �C considerable amount of SiBOC glassyphase was observed suggesting that reducing the carbon content in the SiBOCceramic has a profound effect on the crystallization characteristics. Zhang et al.[73] claimed that sol-gel derived SiBOC ceramic from methyltrimethoxysilane(MTMOS) and MTEOS mixture and boric acid on heat treatment at 1600 �C resultedin the formation of β-SiC nanocrystals and free carbon. Structural analysis by TEMand Raman spectral analysis suggests the presence of highly graphitized free carbonat 1600 �C. Turbostratic graphite with 10–20 layers of graphene is clearly observedin TEM, confirming the significant promoting effect of B on the graphitization offree carbon (Fig. 15).

Fig. 14 Evolution of XRD spectra with temperature for ceramic obtained from (a) poly(phenylbor-osiloxane) and (b) poly(vinylborosiloxane) [60]


Devapal [60] studied the ceramic conversion of B10H12-cage containingborosiloxane oligomer (BSiDB-1) having high boron content (17.5%). It givesmixed non-oxide ceramics containing nanocrystallites of β-SiC (53.9 nm) andh-BN (25.7 nm). It is interesting to note that BSiDB-1 on heat treatment at1650 �C does not give B4C. This is attributed to the absence of direct B-C bond inthe oligomer where boron atom of decaborane is directly attached to the nitrogenatom of the Lewis base.

Devapal [60] synthesized glass forming borosiloxane oligomers by nonaqueoussol-gel process. The advantage of this approach is that there is molecular leveldistribution of Si-O and B-O units and, hence, would give a glass with uniformcomposition compared to the one obtained by melting of SiO2-B2O3 mixture. Theoligomers give ceramic residue in the range of 73–80% at 900 �C in argon atmo-sphere. XRD studies suggest that the glass formed from this type of oligomer is quite

Fig. 15 HRTEM images of ceramic from sol-gel derived borosiloxane prepared fromMTMOS andMTEOS mixture and boric acid. (Reproduced from Ref. [73] with permission from Elsevier)


stable, and formation of crystalline SiO2 (crystoballite) is not noticed even afterannealing at 1650 �C for 2 h in argon atmosphere.

PolymetallosiloxanesInorganic polymers closely related to the polyorganosiloxanes are the poly-metallosiloxane polymers. Initial interest on polymetallosiloxane is to improve thethermal stability of siloxane backbone through the incorporation of metal. As themetallosiloxane bond in the main chain is easily converted to oxide ceramic mate-rials, polymetallosiloxanes are used as precursors for SiO2-MxOy ceramics in theform of fibers, thin films, bulk bodies, and fine powders [74]. Polymetallosiloxaneswere synthesized by the reaction of silicic acid or partially hydrolyzed tetra-ethoxysilane with a metal chelate compound, metal alkoxide, or metal chloride[54, 75]. The polymetallosiloxanes are, however, random copolymers of siloxanedomain and metallosiloxane domain, because it is difficult to control their micro-structure by conventional synthetic techniques. Over the years, with the advance-ments in the field of polymer-derived siliconoxycarbide ceramics,polymetallosiloxanes are considered more as precursors for metal modified SiOCceramics rather than as precursors for SiO2-MxOy oxide ceramics. The thermalalteration of hybrid polydimethylsiloxane–titania nanometer sized composites(nanocomposites) has been studied up to 1600 �C to determine the changes ofSiOC phase as a function of temperature and titania content. The samples wereprepared by the sol-gel method starting from molecular precursors,diethoxydimethylsilane, and titanium isopropoxide. Between 500 �C and 800 �C,the polymer-to-ceramic conversion takes place for all samples with the formation ofdifferent oxycarbide phases and the structure of SiOC and titanium oxycarbidephases has been influenced by the titania content in the gel [76].

Liu et al. [77] obtained SiOC and SiZrOC ceramics starting from the hydrosilanesystem without and with zirconium-n-propoxide [Zr(OPr)4] by sol-gel route. Theprecursor structure indicated that the presence of Si-H bonds led to a different cross-linking mechanism of SiZrOC precursor, where the Si-H bonds were involved in theformation of Si-O-Zr bonds as shown below:

� Si� Hþ � Zr � OPr !� Si� O� Zr � þC3H8

� Si� Hþ � Zr � OH !� Si� O� Zr � þH2

The presence of Zr caused a lower content of sp3 carbon in SiZrOC ceramics. Atabove 1400 �C, the sp3 carbon to silicon ratio in both ceramics increased due tocarbothermal reduction, especially rapidly for SiOC ceramics, suggesting thecarbothermal reduction in SiZrOC ceramics has been strongly suppressed by thepresence of ZrO2.

Ionescu et al. [78] prepared polymer-derived SiOC/ZrO2 ceramic nano-composites by modification of polysilsesquioxanes with zirconium alkoxide toincorporate Zr in the siloxane network structures. Umicevic et al. [79] studied theevolution of the environment of Hf sites in a hafnium-alkoxide-modified poly-silsesquioxane upon polymer-to-ceramic transformation via the perturbed angular


correlation (PAC) method. The results of the PAC measurements on samples ther-mally treated at temperatures from 400 �C to 1300 �C indicate that Hf is surroundedonly by oxygen at all studied temperatures. This finding is in agreement with theevolution pathway of polymer-derived SiHfOC ceramics, which were reported to begenerated as single-phase amorphous materials upon pyrolysis of alkoxide-modifiedpolysiloxanes and subsequently to phase separate and crystallize toward HfO2/SiOCnanocomposites.

SiZrOC fibers were prepared from polyzirconosiloxane gels by using polyvinyl-pyrrolidone as spinning agent [80]. Yan et al. [81] prepared polyhafnosiloxane(PHfSO) by a sol-gel process of a mixture of TEOS, HfCl4, anddimethyldiethoxysilane. Pyrolysis of PHfSO gel fibers yields dense SiOC/HfO2

fibers with high ceramic yield of 81 wt%. The SiOC/HfO2 fibers have tensilestrength of 930 MPa and exhibit good thermal stability up to 1500 �C.

PolymetalloborosiloxanesA new class of precursors, viz., polymetalloborosiloxanes, was synthesized bySwaminathan [22] from boric acid, alkoxysilane, and metal alkoxides which wereevaluated as precursors for metal-modified SiBOC ceramics. Titanoborosiloxanesand zirconoborosiloxane oligomers were synthesized by reacting titanium tetra-ethoxide and zirconium tetrapropoxide, respectively, with boric acid andphenyltriethoxysilaneby solvent less process (Scheme 10).

Heat treatment of titanoborosiloxane oligomers in argon atmosphere at 900 �C,results in the formation of TiO2 (Anatase) and TiO2 (Rutile). When the heat-treatment temperature is increased to 1500 �C and 1650 �C, formation of titaniumsub-oxides is observed which is attributed to the reaction of TiO2 (Rutile) with freecarbon formed from the oligomer. Mixed non-oxide ceramics, viz., SiC-TiB2 andSiC-B4C-TiB2, are formed when titanoborosiloxane mixed with graphite powderand elemental boron, respectively, are heat treated at 1500 �C [22]. Zirconoboro-siloxane oligomer was completely converted into mixed non-oxide ultra-high tem-perature ceramic, viz., ZrB2-SiC-ZrC ceramics, by carbothermic reduction at1500 �C using phenolic resin as the source of carbon.

In an attempt to increase the carbon content of the ceramics, titanoborosiloxane(TPhBS) and zirconoborosiloxane (ZPhBS) oligomers were synthesized by reacting

Scheme 10 Synthesis of polymetalloborosiloxanes [22]


phenyl boronic acid (in place of boric acid), PTEOS and Ti(OEt)4/Zr(OPr)4 in 1:1:1mole ratio by solventless nonaqueous sol-gel process (Scheme 11) [22].

Unlike ZBS synthesized from boric acid, ZPhBS was soluble in common organicsolvents even after the removal of the by-products ethanol and propanol. Theimproved solubility of ZPhBS compared to that of ZBS is attributed to the presenceof phenyl group in the former. TPhBS and ZPhBS give ceramic residue of 72% and68%, respectively, at 900 �C. Ceramic conversion studies of TPhBS and ZPhBS at1500 �C suggest that the ceramic from TPhBS consists of Ti3O5 and TiC in thewt. ratio of 93:7, whereas the ceramic from ZPhBS consists of m-ZrO2, t-ZrO2, andZrSiO4 (Zircon) in the wt. ratio of 59:20:21. This suggests that under the chosenexperimental conditions, titanium oxide ceramics undergo carbothermic reductionmore readily than the zirconium oxide ceramics. Raman spectral studies suggest thatamorphous carbon present in TPhBS-derived ceramic is consumed during the heattreatment for carbothermic reduction whereas it remains unreacted in ZPhBS-derived ceramic.

In a recent study, Anurag et al. [82] reported the synthesis of zirconotitano-borosiloxane (ZTBS) which gave a ceramic residue of ~58% at 900 �C in argonatmosphere. On borothermal reduction at 1500 �C of ZTBS, oxide free mixed UHTCwas obtained (Scheme 12).

Zhang et al. [83] recently reported the synthesis of a precursor for SiZrBOCceramic. The precursor was synthesized by reacting boric acid withmethyltriethoxysilane and zirconylchloride octahydrate by sol-gel process followedby solvothermal process at 130 �C. SiZrBOC ceramic is obtained in 84% yield at1400 �C and is quite stable in air up to 1400 �C.

Scheme 11 Synthesis of metalloborosiloxanes from phenyl boronic acid [22]

Scheme 12 Synthesis of zirconotitanoborosiloxane oligomer. (Reproduced from Ref. [82] withpermission from Springer Nature)


Polysilazanes and PolycarbosilazanesPolysilazanes and polycarbosilazanes are characterized by having a –Si–N– back-bone or a –R–Si–N– backbone, respectively [84]. They are generally prepared eitherby ammonolysis or aminolysis of organochlorosilanes. Ammonolysis of MeHSiCl2resulted in a cyclic silazane (MeSiHNH)3–4 and a linear polymer (Scheme 13).

Synthesis of polysilazanes and polycarbosilazanes and their conversion to Si-Nand Si-C-N ceramics have been extensively reviewed [85–87], and hence, a detailedaccount of the chronological development has not been included in the presentchapter. However, some of the recent research work has been included in thischapter.

Wang et al. [88] reported the synthesis of ethynyl-terminated silazane precursorswhich gave ceramic residue of 84% at 900 �C. This precursor on heat treatment at1450 �C gave mixed non-oxide ceramics (SiC and Si3N4 and free C). Jun et al. [89]reported laser pyrolysis of polysilazane precursor using Nd:YAG laser to SiC, andthe complete transformation from organic to inorganic had taken place when thelaser scanning power was 600 W. This procedure was useful for obtaining SiCcoating on steel substrate. Sun et al. [90] studied the ceramic conversion of poly-vinylsilazane with ZrB2 filler. SiCN matrix formed from polyvinylsilazane remainedamorphous up to 1400 �C in argon, and XRD studies indicated the presence of ZrO2

at 1100 �C, ZrC/ZrN at 1200 �C, and Zr2CN at 1500 �C which account for theimproved high temperature stability of the ceramic.

Polymeric Precursors for SiBCN CeramicsIn order to enhance the properties of SiCN ceramics, boron atoms were added atmolecular level to obtain boron-modified silicon carbonitride ceramics. SiBCNceramics retain their amorphous nature even above 1440 �C due to the formationof B(C)N layers which prevented carbothermic reduction of SiCN ceramic phases bysurrounding them. A review by Viard et al. [91] highlights the preparation of PDCfibers with a special focus on amorphous SiCN and SiBCN networks using melt-spinning and electrospinning processes. Niebylski [92] synthesized organoboro-silazane polymers by reacting polysilazane with trialkoxy, triaryloxy, or tri(arylalkoxy)boroxine. Chemical modification of hydridopolysilazane was performedby reacting it with liquid borazine at low temperatures to obtain boron-modifiedpolysilazane [93]. Chemical modification of hydridopolysilazanes with mono-functional boranes such as pinacolborane (Me2CO)2BH, l,3-dimethyl-l,3-diaza-2-borocyclopentane (CH2NMe)2BH, and 2,4-diethylborazine (Et2B3N3H4) has alsobeen reported [94].

The synthesis and thermal behavior of boron-modified polysilazanes {B[C2H4Si(R)NH]3}n and polysilylcarbodiimides {B[C2H4Si(R)NCN]3}n were reported by

Scheme 13 Synthesis of polysilazane [84]


Aldinger and Weinmann (Scheme 14) [95]. These SiBCN ceramic precursors wereobtained by treating vinyl-substituted polysilylcarbodiimides [(H2C¼CH)(R)-SiNCN]n with borane-dimethyl-sulfide BH3.S(CH3)2. The polysilylcarbodiimidesthemselves were synthesized via treatment of vinyl-substituted chlorosilanes(H2C¼CH)(R)SiCl2 with cyanamide, H2N-C�N, in presence of pyridine, or viathe reaction of the vinyl-substituted chlorosilanes with bis(trimethylsilyl)carbodiimide. The synthesis by a novel reaction pathway and the polymer-to-ceramic conversion of boron-containing polysilylcarbodiimides were reported byWeinmann et al. [96]. Reaction of tris(hydridosilylethyl)boranes with differentamounts of cyanamide in THF followed by thermolysis of such compounds in anargon atmosphere affords SiBCN ceramics in 65–85% yield.

Sarkar et al. [97] synthesized SiBCN ceramic precursor through a simpledehydrocoupling and hydroboration reaction of an oligosilazane containing amineand vinyl groups with (BH3)Me2S. This modified precursor was pyrolyzed to obtainSiBCN ceramics. Solid-state NMR studies indicated that ceramics formed were amixture of hexagonal boron nitride along with turbostratic boron nitride and BN2Cceramics. Increasing pyrolysis temperature leads to the transformation of BN2Cgroups into BN3 and “free” carbon. Gao et al. [98] studied the significance of theprocessing route on the nanostructure evaluation of carbon-rich SiCN and SiBCN

Si

BCHR

R'CH3

ClX

CH2H

NH3

R=H,CH3,ClR'=C2H4Si(R)Cl2

X=NH,N=C=NR=H,CH3,X0.6 R=H,CH3

R'=C2H4Si(R)H2

Me3SiN=C=NSiMe3

X=NH,N=C=NR'=C2H4SI(R)X

Si-B-C-N

Cl

R

Si

Si

Si

C

R

R

R

BR'

R' CHCH3

H H

or BH3-SMe2 NH3,[nBuLi]

CH

CH3BR'

R'

R=H,CH 3,X0.6

Ceramization

ceramic

n

12

3

nX

Scheme 14 Synthesis of boron-modified silazanes. (Adopted from Ref. [95])


ceramics. Bulk carbon-rich SiCN and SiBCN ceramics were successfully producedby hot pressing of polyphenylvinyl and polyborophenyl-silylcarbodiimide.

Five different SiBCN polymeric precursors with B/N atomic ratios ranging from1:3 to 1:0.5 were synthesized by Muller et al. [99] and converted to ceramics bythermolysis at 1400 �C in 70–75% yield. Yu et al. [100] prepared polyborosilazaneprecursors for SiBCN ceramics by using 9-borabicyclo-[1,3,3] nonane andcopolysilazanes as starting materials, involving the hydroboration reaction betweenvinyl groups of polysilazane and B–H groups of 9-BBN under mild conditions andconverted to ceramics. The crystallization behavior and microstructures of PBSZ-derived SiBCN ceramics reveal that the crystallization begins at 1600 �C. Furtherheating at 1800 �C gave XRD patterns for mixed SiC, Si3N4, and BN(C). It isobserved that the introduction of boron improves the thermal stability and ceramicdensity while inhibiting the SiC crystallization. Singh et al. [101] reported boron-modified silazane liquid precursor by reacting polyureamethylvinylsilazane withethylborate (Scheme 15), and this precursor was used for preparing SiBCN-coatedMWCNT.

Using this precursor, nanocomposites with CNTs were also prepared which canbe applied as hard, lightweight coatings on turbine blades, or on clothing. Dispersionof CNTs, fullerenes, and other nanofillers within the ceramic matrix results innanocomposites with tunable semiconductor properties [102]. The introduction ofCNTs in nanocomposites resulted in increased heat resistance to 15,000 W/cm2-about 10 times greater than heat experienced by rocket nozzles.

A new kind of polyborosilazane was synthesized by Zeng et al. [103] fromhexamethylcyclotrisilazane by exposure to borontrichloride. Heat treatment of thepolymer at 1800 �C resulted in BN crystals. Schiavon et al. [104] synthesizedpolyborosilazanes with varying Si/B molar ratios, via hydroboration, of vinyl-substituted cyclotrisilazane, [CH2¼CH(CH3)SiNH]3, with BH3.N(CH3)3. The poly-borosilazanes were pyrolyzed up to 1000 �C under argon atmosphere, which resultedin amorphous SiBCN glasses. The SiBCN glass with B/Si ¼ 3 was amorphous up to1700 �C, and at this temperature, β-SiC and β-Si3N4 were formed. At 2000 �C, β-SiCand BN were formed.

Ganesh and Renjith [105] reported a simple and low cost route for preparing SiBCNprecursors by reacting boric acid with 1,3,5-trimethyl-10,30,50-trivinylcyclotrisilazane

Scheme 15 Synthesis of boron-modified polysilazane [101]


(CTS) in the molar ratio of 1:1, 1:3, and 1:5. The changes in CTS concentration haveaffected the ceramic phases, morphology, and elemental composition of ceramicsobtained from boron-modified CTS (BCTS). The results revealed the formation ofβ-SiC, β-Si3N4, and oxide ceramic phases with BCTS for the molar ratios of 1:1 and1:3, whereas turbostratic BN(C) ceramic phase was formed in addition to otherphases with BCTS in the molar ratio of 1:5.

Polyorganoborosilazane was synthesized via monomer route from a single-sourceprecursor and thermolyzed at 1300 �C in argon atmosphere by Kousaalya et al.[106]. The crystallite size was found to increase from 2 nm to 8 nm with increase indwelling time. A liquid polymeric precursor to SiBCN ceramics was synthesized viadehydrogenation of polymethylsilane and borazine by Wang et al. [107]. Si-H bondsin polymethylsilane react with N-H or B-H bonds in borazine to form cross-linkedpolymer with viscosity of 850 mPa.s which gave β-SiC and Si3N4 crystals on heattreatment at 1600 �C. Si/C/N/H polymers with Si:C atomic ratios �1 were synthe-sized by Gerstel et al. [108] through Wurtz-type coupling and subsequentammonolysis of ClSiH(Me)NHSiH(Me)Cl and MeSiHCl2 or H2SiCl2. These poly-mers were subsequently reacted with a borane to obtain three different precursorswhich were transformed into SiBCN ceramic materials by thermolysis at 1400 �C.The high-temperature behavior up to 2150 �C was investigated by thermo-gravimetric analysis and XRD.

Liu et al. [109] reported the self-healing mechanism of SiBCN ceramics derivedfrom hyperbranched polyborosilazanes. The self-healing ability of cracks wasenhanced at 1000 �C in air with the increase of oxidation time. The ceramics withhigh content of boron show much more obvious self-healing performance due to theformation of SiO2, B2O3, and B2O3.SiO2 which fill the crack and prevent the damageof substrates. Bechelany et al. [110] investigated spark plasma sintering in thetemperature range of 1500–1900 �C at 100 MPa to obtain SiBCN monoliths fromSiBCN powders prepared from boron-modified polycarbosilazane in nitrogen orammonia atmosphere. Lee et al. [111] reported an easy and cost-effective precursorfor SiBCN by reacting trichlorosilane, borotrichloride, and hexamethyldisilazane ata mole ratio of 1:1:4.

Zhao et al. [112] synthesized liquid polyborosilazane through a novel methodwhich used sodiumborohydride as boron source. The reaction sequence is given inScheme 16. Vinylmethyldichlorosilane was reacted with hexamethyldisilazane andthe resultant product was reacted with sodium borohydride to obtain poly-vinylsilazane with Si-H groups. This product was further reacted withdimethylvinylsilane and hexamethylenedisilazane. Vinyl groups present in the pre-cursor underwent hydroboration reaction with diborane formed from sodium boratein the earlier step. The SiBCN ceramic formed at 1400 �C contained B-N, Si-N, andSi-C bonds with smooth and dense surface and still retained principally amorphousstructure up to 1600 �C. In addition, the viscosity of the polyborosilazane was65 mPa.s, and hence, found application as precursor for PIP process for the fabrica-tion of CMCs.

In a simplified reaction process, SiBCN precursors were synthesized by one stepco-condensation of propylamine, BCl3, and SiCl4/SiHCl3 [113]. These polymers


gave ceramic residue of about 65% at 900 �C in nitrogen atmosphere. Type ofceramics formed from polyborosilazane precursor depends on the atmosphere usedfor pyrolysis. Ceramics formed under nitrogen undergoes carbothermal reductionwith Si-N bonds and free carbon at above 1650 �C to form SiC, whereas the sameprecursor sintered in ammonia atmosphere did not undergo carbothermal reductionand remained as Si3N4 and BN at 1650 �C [110]. Ceramic conversion of poly-borosilazane in air resulted in SiBCNO ceramics along with nano quartz andamorphous B2O3 [114]. Tamayo et al. [115] converted sol-gels prepared by reactingTEOS, 3-APTES, triethylborate, and hydroxyl-terminated polydimethylsiloxanewith SiBCNO ceramics by carrying out pyrolysis under nitrogen followed byammonolysis using NH3/N2 gas mixture. During pyrolysis nitrogen atom waseffectively incorporated in the ceramics without replacing carbon. Less graphitizedstructure was observed due to the reduced mobility of carbon atoms in the mixed(B)CN bonds [115]. SiBN ceramic fibers with less than 0.1 wt% carbon wereobtained from polyborosilazane precursor synthesized from hexamethyldisilazane,HSiCl3, BCl3, and CH3NH2 using a multistep polymerization process [116] and asingle step polymerization process [117].

Scheme 16 Synthetic route for poly(borosilazane) containing active Si-H group. (Reproducedfrom Ref. [112] with permission from Elsevier)


Metal Modified Precursors for SiCN CeramicsSimilar to the modification of polysilazane to boron-modified polysilazane, poly-silazanes can be easily modified with aluminum and transition metal alkoxides/metalamido complexes to form Si(M)CN ceramics where M is Al, Ti, Zr, Hf, etc. Differentsynthetic approaches for metal modified polysilazane or polysilacarbodiimides werereviewed by Mera et al. [12] and Ionescu et al. [118]. The ceramic formed from theseprecursors depends on the chemical composition of precursor and the atmosphere inwhich ceramic conversion was carried out. Aluminum-modified polysilazanesreferred to as polyaluminosilazanes were synthesized by reacting polysilazane withaluminum-containing monomers such as Al(CH3)3, AlH3, (CH3CH2)3Al,(CH3)2AlNH2, and aluminum trialkoxide, and these precursors on thermal conver-sion gave SiAlCN ceramics in high yield [119]. Polymer-derived SiAlCN exhibitedmuch better oxidation and hot-corrosion resistance than other silicon-based ceramicsat temperatures up to 1400 �C [120], which makes this material suitable for makingmicro-sensors and MEMS devices for harsh environments. Oxidation studiesrevealed that SiAlCN possesses a lower oxidation rate than SiCN. The betteroxidation resistance of SiAlCN is attributed to the oxide layer containing Al,which slowed down oxygen diffusion relative to that with pure SiO2.

Ceramics from Carbonaceous Polymers

Phenolic resins are extensively used in the production of composite materialsespecially for phenolic resin-carbon fabric composites. In line with the developmentsin the area of high temperature materials, phenolic resins were also evaluated asprecursors for carbide ceramics, especially SiC. The thermosetting property ofphenolic resins can be effectively made use of in the preparation of complex ceramiccomponents. For preparing SiC, phenolic resin is first mixed with silicon or siliconcompounds and cured, and then the cured material is heat treated at 1500 �C orabove.

SiC ceramic powders were synthesized by reacting/mixing phenolic resins withany of the silicon sources such as elemental silicon [121], silica formed by sol-gelprocess [122], fumed silica [123], and rice husk silica [124] followed by reactionsintering. Phenol-furfuraldehyde resin was also used as carbonaceous polymer inplace of phenol-formaldehyde resin for obtaining SiC [125]. SiC powders weresynthesized by the coat-mix process, with phenolic resin and silicon powders asstarting materials [121]. The results show that higher sintering temperature and closeto stoichiometric equivalence of carbon and silicon are required for producing highpurity SiC powders. As the sintering temperature is raised to 1500 �C, the formationof SiC occurs which is based on the liquid-solid reaction. High purity SiC(99.8 wt%) powders can be easily obtained by this method. Tetraethoxysilane andphenolic resin were used for preparing a homogeneous binary carbonaceous silicaxerogel, and nickel nitrate was employed as a pore-adjusting reagent.


From phenolic resin, SiC porous structures [122, 123], SiC coating [126], SiCceramic wires [124, 125], SiC hollow spheres [127], SiC fibers [128], and SiC/SiCcomposites [129] were prepared. Porous SiC ceramics with hybrid pore structure,which is a combination of tubular pores and network SiC struts in the tubular pores,were fabricated via sol-gel conversion of carbonized pine wood with phenolic resin/SiO2 sol followed by carbonization and carbothermal reduction reaction at elevatedtemperatures in argon atmosphere [122]. Phenol-furfuraldehyde resin-elemental Simix on sintering at 1500 �C gave nano β-SiC powder along with SiC wires/fibershaving diameter in nano to submicron range [125]. Prabhakaran et al. [125] observedthat in the XRD pattern of SiC wires/fibers the diffraction line corresponding to<220> plane at 2θ ¼ 59.95 has the highest relative intensity (Fig. 16), and it differsfrom the standard powder diffraction pattern of β-SiC where multi-oriented randomcrystals with <111> plane at 2θ ¼ 35.96 showed the highest intensity. This isattributed to formation of single crystalline SiC formed with orientation along<111> plane.

Hollow spheres of phenolic resin/silica composite were synthesized by macro-scopic phase separation during aerosol-assisted spraying [127]. The spheres wereconverted into carbon-silica hollow spheres by pyrolyzing at 900 �C, which onfurther heat treatment at 1450 �C were converted into SiC hollow spheres.

SiC fibers are particularly desirable reinforcement materials for a variety ofcomposites because of their thermal and oxidation resistance as well as high tensilestrength and Young’s modulus. However, SiC fibers are expensive both because ofthe costly precursors and the number of steps involved in this process. A cost-effective route is reported for SiC fibers, requiring only two steps: synthesis ofsilicon dioxide-organic hybrid fibers by sol-gel processing and the carbothermal

Fig. 16 XRD pattern of SiC ceramic wires from phenol-furfuraldehyde and elemental silicon.(Reproduced from Ref. [125] with permission from Springer Nature)


reduction of the hybrid fibers in an argon atmosphere [128]. Advantages of thisprocess include the relative low-cost of the starting materials, the two-step process,and the high purity of the SiC fibers produced.

Phenolic resins were also used for preparing multicomponent ceramics and ultra-high temperature ceramics [130–139]. Phenolic resin was used as an effective sourcethat could provide free carbon for fabrication of the SiC nanoparticles by in situreaction with Si3N4 and SiO2 [130]. Sizes of the formed SiC particles were in therange of 50–300 nm. Limited quantity addition of polymer resulted in high porositywhereas excessive additions caused agglomeration of residual carbon and SiC,creating fatal flaws in the final material.

Hybrid gels were synthesized from liquid mixtures composed of TEOS, titaniumtetraisopropoxide/titanium tetrakis(2,4-pentadionate), and phenolic resin [131, 132],and these gels on pyrolysis at 1500 �C yielded SiC-TiC ceramics. Si-B-C ceramiccomposites were prepared using SiB6, B4C, and phenolic resin as a carbon source bypressure-less sintering in argon atmosphere to determine their potential for applica-tions as high hardness and high temperature composites [133]. XRD patterns of thesintered bodies show the presence of B4C and SiC. Aluminum-containing carbide,Al4SiC4, was successfully synthesized using a mixture of Al(OH)3, SiO2, andphenolic resin by a carbothermal reduction process [134]. Phenolic resin–TaCl5precursor was recently used for preparing nanocrystalline TaC powder [135]. PorousSi3N4 ceramics with gradient distributions of SiC and pores were fabricated bydirectional nitridation sintering of diatomite preforms with phenolic resin as a carbonsource and pore forming agent [136]. Ganesh Babu and Devasia [137, 138] modifiedphenolic resin with boric acid and the modified resin was reacted with Si to obtainself-healing ceramic matrix comprising of SiC, B4C, and SiB6. On oxidation,borosilicate glass is formed which heals the cracks thereby preventing the inwarddiffusion of oxygen.

In a recent study, ZrB2-SiC ceramic powder was synthesized by mixing zircon(ZrSiO4) and B4C powders with phenolic resin in 6.6:1:3.6 wt. ratio followed by heattreatment at 1500 �C and 1650 �C [140]. XRD patterns of the ceramic indicate thatthe ceramic conversion was incomplete at 1500 �C and contains unreacted zirconalong with SiO2 and ZrB2. When the temperature was increased to 1650 �C, theoxide impurities were converted into ZrB2-SiC. It is worth noting that zirconobor-osiloxane oligomers mixed with phenolic resin was converted completely into ZrB2-SiC ceramics at 1500 �C [22].

Active Filler Controlled Pyrolysis (AFCOP)

Production of bulk ceramics from preceramic polymers suffers from the disadvan-tages that extraordinary volume shrinkage and pronounced increase in density occurduring polymer to ceramic conversion due to which extensive cracking and poreformation take place. Shrinkage and porosity formation could be overcome partly byfilling the polymer matrix with inert fillers such as SiC, B4C, Si3N4, andBN. However, when reactive filler particles are used in place of inert fillers, near


net shape bulk ceramic can be obtained due to the compensation of shrinkage of thepolymer by appropriate filler expansion resulting from the reaction of the fillerparticles with decomposition products of the polymer phase or reaction atmosphere.This approach is known as Active Filler Controlled Polymer Pyrolysis (AFCOP)[141–143]. The concept of AFCOP is depicted in Fig. 17 [143].

Dispersion of Ti, Cr, V, Mo, Si, B, CrSi2, MoSi2, etc. was used as active fillers forpoly(silsesquioxane) [142]. The concept introduced by Greil [141, 143] has beenextensively used to obtain dense bulk ceramic bodies, crack free coatings, CMCswith reduced porosity, and for obtaining a variety of novel composite ceramicsystems including UHTCs. A detailed discussion on density calculations and per-centage of active filler to be used in polymer pyrolysis is beyond the scope of thischapter, and readers’ attention is drawn to related research publications [141–144].

Vipin et al. [145] applied AFCOP concept for obtaining “shrinkage-free ceramic”from vinyl-group containing poly(borosiloxane) (BMV) using titanium silicide[TiSi2] as active filler. The optimum concentration of TiSi2 active filler was foundto be 15.6 wt% for heat treatment of BMVat 1500�C (Fig. 18a). Addition of excessTiSi2 beyond the optimum concentration resulted in a volume expansion. The BMVsystem exhibited the evolution of SiC phase without any stacking faults but theBMV+TiSi2 systems exhibited the evolution of SiC phase with stacking faults(Fig. 18b). It is evident from Fig. 18b that the TiSi2 filler-incorporated polymersystem at 1500 �C provides a mixture of non-oxide ceramic phases (TiB2, TiB, andTiC), as a major constituent, and oxide phases (SiO2, TiOC, and B2O3-SiO2) asminor constituents.

Fig. 17 Microstructuralchanges during polymer-ceramic conversion: (a) filler-free polymer pyrolysisinvolving extended porosityor shrinkage and (b) active-filler-controlled polymerpyrolysis with near-net-shapeprecursor-ceramic conversion.(Reproduced from Ref. [143]with permission from JohnWiley and Sons)


At still higher heat treatment temperatures of 1800 �C and 2000 �C, the evolutionof α-SiC, TiB2, β-SiC, TiB, TiC ceramic phases, and Ti3SiC2 MAX phase isobserved (Fig. 19) [146]. Ti3SiC2 MAX phase is normally produced by expensiveand complicated hot isostatic pressing technique. Hence, the evolution of MAXphase from preceramic polymer route offers a very promising approach for theproduction of this advanced engineering material.

Fig. 18 Effect of TiSi2 filler addition on the ceramic conversion of poly(methylvinylborosiloxane):(a) Volume expansion vs. wt% of TiSi2 and (b) XRD patterns: (A) 0 wt%TiSi2, (B) 5 wt% TiSi2,(C) 10 wt% TiSi2, (D) 15 wt% TiSi2, (E) 20 wt% TiSi2, (F) 25 wt% TiSi2, (G) 30 wt% TiSi2.(Reproduced from Ref. [145] with permission from authors)

Fig. 19 (a) XRD patterns and (b) SEM images of the ceramic obtained from TiSi2 filler added poly(methylvinylborosiloxane), showing the formation of MAX phase. (Reproduced from Ref. [146]with permission from Elsevier)


Polymer-Derived Ceramics for Space Applications

Space applications require materials, which are capable of withstanding extremeservice conditions, for example, nose-cone, leading edges, and rocket nozzles madeof ceramic composites and oxidation-protected C/C composites which should with-stand high temperatures in an oxidizing environment. Likewise, materials used in theconstruction of spacecraft and space station placed in Low Earth orbit (LEO) shouldbe capable of withstanding attack by atomic oxygen. Lightweight ceramic materials,as advanced thermal protection systems and ceramic adhesives for joining ofceramic components are also required for various space applications. This sectionfocuses on the use of preceramic polymers and carbonaceous polymers in thefabrication of CMCs and lightweight ceramics, preparation of ceramic coatingsand ceramic adhesives.

Ceramic Matrix Composites

CMCs are a subgroup of composite materials as well as a subgroup of technicalceramics. They consist of ceramic fibers embedded in a ceramic matrix, thus forminga ceramic fiber reinforced ceramic material. Among CMCs, carbon fiber and siliconcarbide fiber reinforced composites with SiC matrix have gained importance as hightemperature CMCs. SiC, a non-oxide covalent ceramic, exhibits high creep resis-tance and, hence, retains high strength up to 1650 �C. In an oxidizing atmosphere atthese temperatures, SiC undergoes passive oxidation to silica (SiO2) which preventsfurther oxidation of the remaining SiC. This property of SiC makes it superior to Cand eventually, C/SiC composite superior to C/C composite. Carbon, in any form –graphite matrix or fiber – starts oxidizing above 450 �C giving gaseous CO2, whicherodes the entire material and, thus, needs to be protected from oxidizing environ-ment above 450 �C. A suitable oxidation resistant coating (generally SiC coating) isused to protect the C/C composite, particularly, when used in space reentry vehicles.However, due to difference in coefficient of thermal expansion (CTE) of C and SiC,the SiC coating spalled-off leading to the failure of the mission. Thus, fiberreinforced CMCs are considered more reliable for reentry missions. Besides theexcellent oxidation resistance, higher specific mechanical properties of CMCs athigher temperatures, compared to that of metal alloys, qualify them as high temper-ature structural materials for aerospace applications [147, 148].

Among the well-known CMCs, viz., C/SiC, SiC/SiC, and C/C-SiC composites,C/SiC and C/C-SiC composites, are extensively used as hot structures. SiC/SiCcomposites, though has the best oxidation resistance, has limitation in strengthretention beyond 1400 �C. A wide variety of thermal protection systems and hotstructures has emerged (Fig. 20) as the sustainable approach, as the space industryhas moved from rocket launchers to shuttle orbiters to scramjet-based hypersonicvehicles [149].

The application areas of CMCs include aero-engine combustor liners, ducts,nozzle flaps, acoustic liners, turbine vanes, turbine blades, and turbine disks [150,


151], re-entry vehicles [152–154], as substrate for SiC mirrors [155], brakes formilitary and commercial aircrafts [156], as high temperature ceramic fasteners [157],and in radiation-cooled nozzle extensions and combustion chambers for smallthrusters [158]. The major end-use of C/SiC composites is as thermal protectionsystems (TPS) [159–162]. CMCs also showed good performance in rocket motorparts reaching temperatures of 2700 �C, withstanding the chemical and physicalerosion [163].

The potential use of SiC-matrix composites as hot structures of spacecraft wasdemonstrated at the prototype level some years ago with the “Hermes” (Europeanspace shuttle project). Here, the temperature envisaged was from 800 �C to 1600 �C,during the ascent and reentry phases of the flight, and the structures withstoodthermal shocks and cyclic mechanical loading under ablative or passive oxidizingatmospheres [164].

C/SiC was proposed as a suitable candidate for the nose-cap of X-38 reentryvehicle by NASAwhich was designed to be a technology demonstrator for the CrewReturn Vehicle (CRV) from the International Space Station (ISS) (Fig. 21)[165]. During the reentry phase the nose cap of X-38 experiences surface tempera-tures of up to 1750 �C. C/C-SiC nose-cap was envisaged in the EXPERT program ofESA. ESA has developed the intermediate experimental vehicle (IXV) project foraddressing the basic European needs for reentry from LEO, which uses C/SiC TPScomponents [166] (Fig. 22).

Existing satellite thrusters and liquid apogee motor (LAM) engines use hightemperature alloys where the permissible operating temperature is 1300 �C. Theuse of C/SiC composite as combustion chamber and nozzle material can increase themaximum combustion temperature above 1650 �C without any oxidation protection

Fig. 20 Thermal structures for aerospace applications. (Reproduced from Ref. [149])


coating [167]. With suitable functionally gradient coating the operating temperaturecan be further increased which will result in enhancing the specific impulse of thethruster. In view of the advantages of C/SiC composites, they are being evaluated assatellite thrusters (Fig. 23), nozzles, nozzle extensions (Fig. 24), and combustionchambers [167, 168].

Fig. 21 Nose-cap of X-38(NASA). (Reproduced fromRef. [165] with permissionfrom John Wiley and Sons)

Nose

Windward Assembly

Lateral Ablative

Flap Assembly

Hinge TPS

BaseLeeward Ablative

Lateral Ablative

Fig. 22 C/SiC Components for IXV (ESA): (a) Nose assembly, (b) shingles-based windwardassembly, and (c) flap assembly. (Reproduced from Ref. [166] with permission from Elsevier)


C/SiC composite is also used for the fabrication of lightweight mirrors and as facesheets in the construction of internal multiscreen insulation (IMI) [149] and ceramicsandwich structures (Fig. 25) [169, 170].

Fiber reinforced CMCs are fabricated by producing a dense ceramic matrixaround carbon or silicon carbide fibers. The build-up of the dense matrix is mainlydone by three techniques using: (i) gaseous reactants by chemical vapor infiltration(CVI), (ii) polymeric reactants by polymer infiltration and pyrolysis (PIP) process,and (iii) molten elements reacting with the preforms (by reactive melt infiltration(RMI)/liquid silicon infiltration (LSI)), each having their own advantages andlimitations.

In general, the matrix should be homogeneously distributed in the preform withlimited residual porosity. Another important aspect is the fiber to matrix(FM) bonding. In conventional composites (like Kevlar/epoxy or carbon/epoxy),the matrix (like epoxy) is ductile enough to pass on the stress to the fibers. However,in CMCs, the matrix is as stiff as the fibers, which makes it impossible for the matrix

Fig. 23 (a) 150 N bipropellant C/SiC rocket engine, (b) hot fire of 150 N bipropellant C/SiC rocketengine. (Reproduced from Ref. [167] with permission from Elsevier)

Fig. 24 C/SiC nozzleextensions developed forVulcan cryogenic engine(ULA). (Reproduced fromRef. [168] with permissionfrom John Wiley and Sons)


to pass on the load to the fibers. The fiber reinforcements can enhance the toughnessof the CMC, only if the FM bonding is weak. The concept of weak FM bondinghelping in crack deflection is explained in Fig. 26. The weak FM bonding is mainlyestablished by providing a thin layer (less than 1 μm) of another material on thecarbon fibers by the CVI process [171].

In the case of C/SiC composite without interphase coating fiber-matrix adhesionis strong and hence, the fracture propagates through the fiber, and the compositebehaves like a monolithic material. When an interphase coating such as BN or pyrocarbon is given onto the fiber reinforcement matrix, cracks get diverted around thefibers as shown in Fig. 26 and the composite exhibits good fiber-pullout contributingto improved fracture toughness of the composite.

The design of CMCs calls for selection of a suitable reinforcement and acompatible matrix, appropriate fiber-matrix (FM) interphase, the combination capa-ble of exhibiting high mechanical and thermal properties, and high chemical com-patibility with the high service temperature. Carbon fiber or SiC fiber woven asfabric or in multiple dimensions is the most preferred reinforcement. Carbon and SiCare the mostly preferred material for matrix. As FM interphase, hexagonal boronnitride (BN) and pyrolytic carbon are mainly used [172–175]. Multilayer interphasecoatings have also been attempted to improve the thermomechanical properties ofCMCs [171, 176].

The matrix phase in CMC can be single or multiple phases with compositionalvariations. C/SiC composites are the simplest in terms of the composition of thematrix. C/SiC composites are usually fabricated by the well-established CVI processand by this process both the interphase coating of the reinforcement fiber/fabric andSiC matrix are obtained. Though CVI process is useful in fabricating C/SiC com-posites with required mechanical properties, the fabrication of composites by thisprocess suffers from the following disadvantages: (i) quite high initial investment inestablishing CVI furnace and its high maintenance cost, (ii) use of corrosivechemicals such as boron trichloride and methylchlorosilanes, (iii) requirement ofspecially designed graphite molds for uniform deposition of interphase coating anduniform densification of matrix, (iv) difficulty in fabricating components having

Fig. 25 Schematic of a ceramic sandwich structure


complicated shape and larger size, (v) blockage of the surface pores due to matrixdeposition preventing the penetration of gas into the interior of the componentswhich limits the thickness of the components that could be fabricated by this process,(vi) low efficiency of the process, for example, the conversion of gaseous precursorsto SiC is less than 15%, and (vii) environmental/pollution-related issues relating toneutralizing the unreacted gaseous monomers. Though the basic CVI process hasundergone several modifications/improvements to overcome the above problems,fabrication of CMCs by CVI process remains to be costly and time consuming.

Polymer-derived ceramics route offers solution to most of the problems associ-ated with CVI process. As discussed earlier, preceramic polymer can be used as theprecursor matrix for the desired ceramics and any one of the polymer matrixcomposite preparation methods such as hand-lay-up, autoclave molding, resin trans-fer molding, and filament winding can be used to form the green composite. Thismethod takes the full advantage of the low temperature processability of preceramicpolymers. After pyrolysis, the porous composite is subjected to infiltration and

Fig. 26 Schematicrepresentation of the role ofinterphase in crack deflection


pyrolysis cycles to obtain a dense CMC, and this procedure is known as polymer-infiltration-pyrolysis (PIP) process (Fig. 27) [177].

Even for the fabrication of CMCs by PIP process, it may be required to useCVI/CVD furnace for interphase coating onto fiber reinforcement. Though PIPprocess is relatively simple and cost effective, care must be taken in each processingstep, particularly the pyrolysis step. During pyrolysis, evolution of gases due to thedecomposition of precursor matrix takes place leaving behind the amorphousceramic matrix which results in shrinkage of the matrix and formation of pores,voids, and micro-cracks. Sudden evolution of gases may lead to delamination of thegreen composite, and to overcome this problem, the first pyrolysis step is to becarried out at a slow heating rate. Repeated PIP process may also deteriorate thestrength of fiber reinforcement due to the possible reaction of matrix with theinterphase coating/fiber. Pyrolysis temperature should be carefully chosen suchthat the reinforcing fiber of choice does not lose its strength [177].

Efforts to visualize and to assess the distribution of voids created in the matrixduring polymer to ceramic conversion, by using X-ray computed tomography ispromising [178]. These approaches can help in formulating physics-based models incomposite fabrication processes (Fig. 28).

Silicon carbide matrix: Various preceramic polymers such as polysilane [179],polymethylsilane [180], polyvinylsilane [181], polycarbosilane [182–184], poly-carbosilane with AlN-Y2O3 as sintering additive [185], co-polymer of

Fig. 27 Flow chart for fabrication of composite by PIP. (Reproduced from Ref. [177] withpermission from Elsevier)


polycarbosilane and polymethylsilane [186], polytitanocarbosilane [187], poly-zirconocarbosilane [188], and allylhydrido-polycarbosilane [189, 190] were usedto prepare SiC matrix. Starfire SMP-10 was one of the initial commercially availablepolycarbosilane precursors to stoichiometric SiC. High cost of these resins hasimpeded the commercial applications of polymer-derived CMCs as in the case ofautomobile brake discs [191].

SiC matrix has also been obtained by carbothermal reaction of Si or SiO2.Prabhakaran [61] and Ajith [192] have investigated the fabrication of C/SiC com-posites from carbonaceous polymer such as phenol-formaldehyde and phenol-furfural resins [61], as a cost-effective approach. As no interphase coating wasused to protect the carbon fiber, the reaction of Si with carbon fiber could not beeliminated. However, the composite fabricated by this process retained the mechan-ical property at 1500 �C (Table 2).

Infiltrating C/SiC composite with polycarbosilane and borosiloxane oligomerfollowed by pyrolysis at 1500 �C [192] considerably improved the oxidation resis-tance. 2D C/SiC composites with and without pyrocarbon interphase were fabricatedusing boron modified phenol-formaldehyde resin as matrix precursor and thesecomposites exhibited flexural strength of 102 MPa and 38 MPa, respectively [138].

SiOC matrix: The search for a low-cost resin ended up in a range of SiOCprecursors. Among the commercially available resins for SiOC, Starfire system’sPolyramic range of resins were the widely used ones. Comparing the oxidationresistance of these resins, Polyramic RD-212a and RD-036a resulted in 100% massretention after oxidation while RD-688a, RD-642, and RD-684 polymers, which

Fig. 28 In situ 3D visualization of composite microstructure during polymer-to-ceramic conver-sion. (Reproduced from Ref. [178] with permission from Elsevier)

Table 2 Flexural strengthof C/SiC composite fromphenol-formaldehyde resins(Reinforcement: Toray 3 Kcarbon fabric withoutinterphase coating) [192]

Test temperature (�C)Flexural strength (MPa)

Mean SD

RT 72.79 10.92

1000 76.14 1.47

1300 79.20 10.35

1500 79.40 5.15


yield progressively higher carbon content ceramics, resulted in progressively lowermass retention and therefore lower oxidation resistance [191]. Considerable amountof work is carried out by researchers on SiOC matrix composites from siloxaneresins [193–196]. Polysiloxane with titanium [197] and silicon [198] and poly-silsesquioxane with boron [199] as filler were reported to address the problem ofvoids and shrinkage to certain extent. Srinivasan and Tiwari [199] noted that despiteits high temperature property, polysilsesquioxane has several limitations whichprevent its use as matrix resins for CMCs and advocated the use of alternate resinssuch as borosiloxanes, carborane-siloxanes, and other boron-nitrogen containingpolymers. For the fabrication of C/SiOC CMCs, “Fast Sol-Gel Process” is alsoreported [200]. In this method, green composites were fabricated by coating fast-sol-gel derived resins onto fabric preforms followed by hot press and pyrolysis.Carbon content in the matrix influenced the flexural strength. In C/SiOC composites,the matrix contains nano-crystallites of SiC embedded in an amorphous SiOCmatrix. The SiOC matrix has good oxidation resistance but has poor creep resistanceat higher temperatures due to the presence of silica nano domains. Use of SiBOCmatrix in place of SiOC improves the oxidation resistance further, as SiBOC matrixon oxidation forms stable borosilicate glass and also influences the crystallizationkinetics as discussed earlier. The ability of borosiloxane oligomers to give SiBOCceramic in good yield on pyrolysis is an advantage to use them as matrix resin forCMCs [201, 202].

CMC was fabricated by impregnating pitch-based carbon fiber strand (hightensile strength type) with poly(diphenylborosiloxane) solution containing β-SiCpowder, followed by molding in a metallic mold, heat treating in air at 500 �C andthen at 1000 �C in nitrogen atmosphere [53]. The mechanical property increases withthe increase in Si/B ratio in the precursor. The composite prepared from the precursorhaving Si/B ratio of 3:1 exhibited maximum bending strength of 320 MPa andYoung’s modulus of 55 GPa. CMCs were also fabricated from 2.5-dimensionallywoven carbon fabric using polyborosiloxane as matrix precursor [203]. This com-posite having density of 1.62 g/cm3 gave a bending strength of 83MPa and exhibitedmore toughness and higher fracture energy compared to CMC fabricated from satinweave carbon fabric.

Schiavon et al. [204] evaluated poly(methylborosiloxane) prepared from boricacid and methyltriethoxysilane as matrix resin for preparing C/SiBOC unidirectionalcomposite. SEM studies indicate good adhesion between carbon fiber and ceramicphase. This composite has improved oxidation resistance compared to C/SiOCcomposite.

Borosiloxane oligomers synthesized by nonaqueous sol-gel process have beenevaluated by researchers at the Vikram Sarabhai Space Centre as matrix resin forcarbon fiber and SiC fiber reinforced composites and arrived at a propriety compo-sition which gives maximum flexural strength at room temperature and at elevatedtemperatures. To begin with poly(phenylborosiloxane) [BP], BP with pyrolyzed BPand BP with elemental silicon were evaluated as matrix precursors to obtain CMCsdesignated as CMC-1, CMC-2, and CMC-3, respectively [201]. A number of PIPand heat treatment cycles are required to get a densified composite (Fig. 28). CMC-1


requires maximum number of processing cycles to achieve the saturation density. Onthe other hand, CMC-2 prepared using pyrolyzed PBS as filler requires lesserprocessing cycles compared to CMC-1. CMC-3 prepared using elemental siliconas filler requires the least number of densification cycles (Fig. 29a). SEM images ofcross section of CMC-1 are shown in Fig. 29b, c. Figure 29b shows the presence ofboth densified matrix and carbon fiber reinforcement. In spite of many PIP cyclesvoids are seen in the composite (Fig. 29c) which is attributed to the flaws in the greencomposite.

Flexural strength values of the composites are given in Table 3. Among the threecomposites, CMC-3 has the lowest flexural strength and this is attributed to thereaction of elemental silicon with carbon fiber. Use of pyrolyzed BP as filler has

Fig. 29 (a) PIP Densification of phenyl borosiloxane derived CMC, (b) SEM of cross section ofCMC-1, and (c) SEM of cross section of CMC-1 showing voids [201]

Table 3 Flexural strengthvalues of CMCs [201]

Sample Density (g/cc) Flexural strength (MPa)

CMC-1 1.78 145

CMC-2 1.81 173

CMC-3 1.83 110


contributed toward increasing the flexural strength and reducing the number ofinfiltration cycles. The higher flexural strength of CMC-2 is attributed to propersintering of the matrix which helps in improving the inherent load bearing capacityof the matrix, thereby increasing the bending strength [205]. Yet another reason forimproved flexural strength of CMC-2 is less number of infiltration and heat treat-ment cycles, which helps in minimizing the microcracks and differential shrinkageof the matrix [206]. CMC-1, CMC-2, and CMC-3 on oxidation in air at 1000 �C inair undergo weight loss of 23.2–26.2% as against 82.9% for C/C composite.

Devasia et al. [207, 208] explored the possibility of obtaining pyrocarbon inter-phase coating onto carbon fiber through polymer pyrolysis route. Carbon fiber wascoated with poly(acrylonitrile-co-itaconic acid) by dip coating process and thenpyrolyzed in the temperature range of 1000–1500 �C to obtain the PyC coating.The advantage of this approach is that it eliminates the need to use CVI furnace forobtaining interphase coating. Maximum improvement in flexural strength isobserved for C/SiBOC composite prepared from PyC coated carbon fabric andpoly(phenylborosiloxane) for the interphase coating thickness of 0.5 μm. SEMstudies indicate that PyC coating of 0.5 μm improves the fiber-pullout which isresponsible for the improvement in flexural strength (Fig. 30).

A detailed evaluation of different borosiloxane oligomers as precursor matrixsuggests that mechanical properties of C/SiBOC composite are influenced by theorganoalkoxysilane used for the synthesis of borosiloxane oligomers, nature ofadditives, and interphase coatings. The flexural strength of different C/SiBOCcomposite systems from borosiloxane oligomers is compared in Table 4.

It is observed that C/SiBOC composite from BMV resin gives the maximumflexural strength. SiBOC ceramics obtained from BMVoligomers are less prone toSiC crystallization as discussed earlier. The increased stability of SiBOC phase fromBMV is responsible for improving the load bearing capacity of the composite [209].

As C/SiBOC composite prepared from BMV resin gives the maximum flexuralstrength of 265 MPa, attempts have been made to make use of this resin forfabricating C/SiBOC composite for space applications. Initial trials have been

Fig. 30 SEM of fractured surface of C/SiBOC composite from polyphenylborosiloxane: (a)without and (b) with PyC coating (0.5 μm). (Reproduced from Ref. [208] with permission fromElsevier)]


focused on fabricating shaped composites, avoiding delamination and warping ofcomposites by using suitable molds, and adjusting the heat treatment process. Afterstandardizing the process, BMV has been successfully used for fabricating pylonleading edge and prototype air-breathing engine combustion chamber and wingleading edge trial segment for reusable launch vehicle (Fig. 31) by vacuum assistedresin transfer molding.

Table 4 Comparison of flexural strength of C/SiBOC composite prepared from differentborosiloxane oligomers

SlNo. Alkoxysilane(s)a System code

Flexural strength(MPa) Reference

1 PTEOS PBS 145 [62, 208]

2 PTEOS PBS (PyCb coated C fiber) 163 [207, 208]

3 PTEOS PBS+Si 110 [62, 201]

4 PTEOS PBS+PyBPS 173 [62, 201]

5 VTEOS VBS 124 [210]

6 VTEOS VBS+MWCNT (3%) 171 [210]

7 VTEOS + MTEOS BMV 265 [62, 209]

8 VTEOS + MTEOS BMV (BN coated carbon fiber) 285 [62, 209]aBoric acid: Alkoxysilane ratio ¼ 1:2 except in Sl. No. 7bPyrocarbon interphase coating obtained from polyacrylonitrile

Fig. 31 Shaped CMC components from borosiloxane [209]


In order to qualify BMV-based composites for inducting into the space program,it is necessary to study the high temperature characteristics and oxidation resistance.The composite retains mechanical properties up to 1000 �C (Table 5).

Long duration oxidation resistance test of BMV-based CMC was carried out at1000 �C. The CMC samples were exposed to 1000 �C in a muffle furnace. Thecomposite shows a cumulative mass loss of 39% after 60 min (Fig. 32a).

Mass loss of the C/C composite is observed to be four times as compared to thatof C/SiBOC composite. The higher oxidation resistance of C/SiBOC is attributed tothe oxidation stable matrix and protection of carbon fiber bundles by SiBOC. Themass loss of 39% for C/SiBOC is mainly due to the burnout of carbon fibers getting

Table 5 Comparison ofMechanical properties ofC/SiBOC composite atroom temperature and at1000 �C [209]

TemperatureFlexural strength(MPa)

Tensile strength(MPa) ILSS

RT 236 153 10

1000 �C 264 142 13

Fig. 32 Evaluation of oxidation resistance of BMV-based C/SiBOC at 1000 �C in air: (a)comparison of C/SiBOC with C/C composite, (b) matrix after oxidation, and (c) burnout of carbonfiber bundles [62]


exposed on machining as evident from SEM. If proper oxidation protection isprovided, this mass loss can be considerably minimized.

Protective Coatings

Oxidation Resistant and Environmental Barrier Coatings for Carbon/Carbon and Ceramic Matrix CompositesAdvanced C/C composites have attractive properties such as high strength up to2200 �C in inert atmosphere, low density, low coefficient of thermal expansion andthermal shock resistance [211]. For these reasons, they are considered as the mostsuitable candidate for hot structures such as nosecone and leading edges of reentryvehicles and reusable space shuttles [212]. However, in an oxidizing environment,they undergo degradation above 500 �C, which drastically brings down their prop-erties. Hence, they need protection against oxidation if they are to be used for theabove applications.

Multilayered oxidation resistant coating is the popular choice for the protection ofC/C composites against oxidation [211]. A typical multilayer coating consists ofthree parts. An initial sealant layer is applied to the substrate by either slurrypainting, spraying, or CVD. A primary oxidation barrier is then applied by CVD,consisting of near stoichiometric SiC or siliconized SiC. Because of the thermalexpansion mismatch between the fiber-reinforced substrate and the primary oxida-tion barrier, a network of cracks develop within the coating upon cooling down fromthe deposition temperature of CVD coating. Higher the deposition temperature,wider the cracks that are developed and these cracks provide paths for the diffusionof oxidizing species to the substrate. In order to seal the cracks on SiC coating, in thelast step of the coating process borosilicate/silicate glass coating is applied bypainting, spraying, or impregnation using B2O3-SiO2 mixture, tetraethoxysilane, orby sol-gel process. Cracks also develop in the SiC coating when SiC-coated C/Ccomposites are subjected to high temperatures during reentry, due to thermal expan-sion mismatch between SiC coating and C/C composite. The glass coating that isapplied over SiC coating should be capable of sealing the cracks which are formedduring the processing of the coating and also the cracks which are formed during theflight. Thus, it is necessary that the glass coating should melt at a slightly lowertemperature than the service temperature, in order to ensure its melting and flowinginto the cracks. In addition to sealing the cracks, the glass coating offers yet anotheradvantage, viz., protection against moisture absorption by SiC coating and by thecomponent [213]. If glass coating is not applied over SiC coating, it would absorbmoisture and the moisture absorbed at the interface between the SiC coating and thesubstrate would form water vapor/steam when heat load is applied onto the substrate[214]. This would result in the development of internal pressure at the interfacecausing rupture of the coating from the substrate when the SiC coated compositeexperiences higher temperatures.

Several variants of the oxidation protection methodology have been reported inpatent literature [215–217] and in research publications [218–222] to improve the


performance of the coating with respect to the following aspects: (i) to eliminate/reduce the cracks on the oxidation protection coating by way of minimizing thethermal expansion mismatch between the composite substrate and the coating, (ii) toimprove the adhesion of the coating with the substrate, (iii) to increase the hightemperature capability of the coating, (iv) to improve the service life and reusability,and (v) to modify the coating to incorporate environmental barrier capability,particularly, resistance to degradation of the ceramic coating to water vapor.

Though ceramic matrix composites such as C/SiC and C/C-SiC and UHTCcomposites are more resistant to oxidation compared to C/C composites, they alsorequire oxidation protection in order to extend their service life. The oxidationprotection methods of C/C composites are adopted as such or suitably modified tomeet the requirements of other composite systems. It is likely that the protectivecoatings may be damaged during installation or operation process. The destructionof these coatings may cause oxidation of C/C composites leading to serious accident[223], and hence, suitable repair systems are to be developed. Unlike the hightemperature coating process, the in situ repair process has to be a low temperatureprocess and should be easily adoptable. Polymeric precursors would be the bestchoice for the repair work as they can be cured into a cross-linked network by UVlight [224] or by slight heating [225].

The shrinkage of the films during pyrolysis at high temperatures and the resultantcracks could be avoided by adding inert or active fillers [226]. It is reported thatpolymer-derived ceramics show excellent oxidation resistance [120, 227] and watervapor corrosion resistance [228, 229]. In view of the obvious advantages of polymer-derived ceramics, preceramic polymers such as PCS [230], polysilazanes [231],polysiloxanes [227], boron containing polymers [232, 233], and boron and siliconcontaining polymers [92, 234, 235] have emerged as promising alternatives for theprotection of C/C composites, CMCs, carbon fibers, ceramic fibers, and metalsagainst oxidation.

Oxidation resistant coating for C/C-SiC composite was reported by Bill andHeimann [231] using liquid polysilazane with silicon powder as filler. The compos-ite was coated by dip coating process and then subjected to pyrolysis at 1100 �C at aslow heating rate (0.7–1.2 �C/min). Ceramic coating thickness of ~5 μm wasachieved in each dip coating and pyrolysis cycle. The adherence of the ceramiclayer on the C/C-SiC material depends on the pretreatment of the substrate as well ason the pyrolysis conditions. Adhesion of the ceramic layer on the C/C-SiC substratewas determined by the direct pull method and the adhesive strength was in the rangeof 6–14 MPa. The adhesion of the ceramic layers could be increased by 30% byadditional annealing of the samples in nitrogen atmosphere at 1350 �C. The observedimprovement was probably due to nitridation of the silicon filler.

The oxidation behavior of the coating and the ability to protect the C/C-SiCsubstrate against oxidation were investigated using thermogravimetric analysis[231]. Both samples were heated in air up to 1000 �C at a rate of 1 �C/min andthen cooled down at a rate of 8 �C/min. The reference sample exhibited a weight lossof more than 50% whereas the sample coated with 10 μm thick ceramic layer lostabout 30% of its original weight. Oxidation resistance test conducted at 1160 �C and


1260 �C suggested that all the carbon fiber reinforcement was burned off both in thecoated and uncoated samples though it took more time for the coated sample. Thereason for the weight loss in the case of coated sample is attributed to the porosityand micro-cracks in the coating.

Wang et al. [236] formulated a single layer and multilayer repair coating based onpolysilazane. Single layer coating slurry was prepared by mixing of polysilazanesolution in ethanol and ZrC powder followed by ball milling. The slurry was brushedon the damaged areas of C/C composites followed by curing using an infrared lamp.For the multilayer coating structure, silicon powder, mullite powder, and borosilicateglass powder were used as the raw materials for different layers. The fabricationprocess for each layer is similar to the single-layer process. Oxidation resistance testindicated that single layer coating was not able to effectively protect the coating-damaged C/C composites. On the other hand, the multilayer coatings could protectthe C/C composites from static oxidation at 1300 �C and 1500 �C. The plasma windtunnel test was carried out with heat flux 4.0 MW/m2, stagnation pressure 3.5 kPa,and air velocity of 2 Mach. The multilayer coating withstood this test condition for1000 s.

Manocha and Manocha [237] investigated the oxidation resistance of methyl-silsesquioxane-based SiOC coating on C/C composites. Methylsilsesquioxane resinwas dissolved in a mixture of toluene and isopropylalcohol, and the cross-linkingagent 3-aminopropyltriethoxysilane (2%) was mixed with the solution. C/C com-posite coupons were coated with the resin solution and cured at room temperaturefollowed by post curing at 150 �C. The samples were heat treated at 1000 �C innitrogen atmosphere. The composites were further coated with the resin to fill cracks,if any, developed during the first heat treatment. The coated samples were heattreated to 1250 �C in nitrogen. Oxidation behavior of the coated and uncoatedsamples was studied thermo-gravimetrically, by heating the samples to 1200 �C ata heating rate of 10 �C/min under the flow (100 mL/min) of compressed air andrecording the weight loss at different temperatures. Uncoated C/C composite hadexperienced complete weight loss at 950 �C whereas the SiOC coated composite hadexhibited less than 1% weight loss in the temperature region 800–1200 �C. Theresults suggest that SiOC coating offers very good protection to C/C compositeagainst oxidation.

Niu et al. [238] fabricated MoSi2-SiOC-Si3N4 coating over SiC coated C/Ccomposite by slurry method. The slurry was prepared by dispersing MoSi2 andSi3N4 powders in a high hydrogen silicone oil (H-PSO) and using tetramethyltetravinylcyclotetrasiloxane (Vi-D4) as cross-linker and platinum chloride as cata-lyst. The coating was heat treated at 1000 �C for 2 h in nitrogen atmosphere to obtainMoSi2-SiOC-Si3N4. The average coating thickness was 140 μm. For comparison,MoSi2-SiOC outer coating was prepared by the same slurry method. The isothermaloxidation curves of the C/C composites with different outer coatings at 1500 �C inair are shown in Fig. 33. It is observed that MoSi2-SiOC-Si3N4 exhibits superioroxidative protective abilities compared to MoSi2-SiOC. The antioxidation propertiesof MoSi2-SiOC-Si3N4 multiphase coatings are influenced by Si3N4 content. FromFig. 33a it is evident that the MoSi2-SiOC/SiC coated C/C specimen exhibits an


excellent oxidation resistance in the first 20 h with a negligible weight loss of 0.5%.However, as seen from Fig. 35b the weight loss rate increases gradually until theweight loss reaches up to 5.52% for 80 h. The above data infers that MoSi2-SiOCcoating offers limited protection. The reason for this observation is that the CTE ofMoSi2 is 8.1 � 10�6 K�1 is much higher than that of the inner SiC(4.45 � 10�6 K�1), and hence, stress-induced micro-cracks are formed with theincrease in temperature. These cracks provide the diffusion pathways for oxygen toget access to the matrix, leading to the increase in rate of oxidation at the later stage.The mismatch of CTE between the outer and inner coating is minimized by intro-duction of Si3N4with a lower CTE (2.7 � 10�6 K�1). It is seen that MoSi2-SiOC-Si3N4/SiC coating with Si3N4 content of 30% or 45% could protect C/C matrix fromoxidation at 1500 �C for 100 h. However, the coating with 60% Si3N4 could give theoxidation protection to C/C composites just for a short period of time.

XRD studies of the coating after exposure to 1500 �C suggest that MoSi2-SiOCand MoSi2-SiOC-(45%)Si3N4 coatings consist of SiO2, Mo5Si3, and MoSi2 phases.Based on this inference, the following reaction scheme was proposed [238]:

5MoSi2 sð Þ þ 7O2 gð Þ ! Mo5Si3 sð Þ þ 7SiO2 sð Þ2Mo5Si3 sð Þ þ 21O2 gð Þ ! 10MoO3 gð Þ þ 6SiO2 sð ÞSi3N4 sð Þ þ 3O2 gð Þ ! 3SiO2 sð Þ þ 2N2 gð Þ

The above reactions generate SiO2. Especially, the oxidation of MoSi2 is charac-terized by the selective oxidation of Si. SiO2 formed in sufficient quantity forms adense film, which effectively serves as a barrier against oxygen attacking the matrix.This is evident from SEM studies of the samples subjected to oxidation at 1500 �C.At medium temperature of 800 �C, MoSi2-SiOC/SiC coated sample has a significantweight loss of 6.76% after oxidation for 10 h whereas MoSi2-SiOC-(45%)Si3N4/SiCcoated C/C sample has a better oxidation resistance with weight loss of 2.11% afteroxidation for the same time. The oxidation resistance of both the coatings at 1500 �C

Fig. 33 Weight loss curves of different coated samples at 1773 K in air: (a) 20 h and (b) 80 h (1:MoSi2-SiOC-(15%)Si3N4; 2:MoSi2-SiOC-(30%)Si3N4; 3:MoSi2-SiOC-(45%)Si3N4; 4:MoSi2-SiOC-(60%) Si3N4). (Reproduced from Ref. [238] with permission from Elsevier)


is better than that observed at 800 �C mainly because SiO2 formed at this temper-ature has a poor self-healing ability.

C/SiC composites exhibit excellent oxidation resistance in dry air, due to theformation of a dense silica layer on their surface [239, 240]. However, the watervapor containing in combustion environment removes the formed silica quickly,resulting in the mass loss and failure of the C/SiC composites [241, 242]. Hence,environmental barrier coatings (EBCs) are applied in order to protect the C/SiCcomposites from corrosion in combustion gas [243]. Liu et al. [225] investigatedtwo-component polysiloxane-derived SiOC-barium-strontium-aluminosilicate(SiOC-BSAS) coating as environmental barrier coating for C/SiC composites.Polysiloxane-BSAS (40 wt%) slurry was uniformly brushed on C/SiC composite.After curing at 100 �C, the coating was subjected to pyrolysis at 900 �C followed byheat treatment at 1350 �C under the flow of argon to obtain SiOC-BSAS coatedC/SiC composite. The oxidation behavior of SiOC-BSAS coated C/SiC compositewas studied at 1250 �C in dry air for 100 h. The weight loss and the residual flexuralstrength of coated and uncoated C/SiC composites are compared in Fig. 34.

The weight loss for C/SiC composites without SiOC–BSAS coatings increaseswith the increase in exposure time. The weight loss is 8.5% after the composites areoxidized for 100 h whereas the weight loss for the composite coated with SiOC–BSAS is only about 3 wt% for the same duration. In the case of residual flexuralstrength, the composites with SiOC–BSAS coatings retain 72% of their originalstrength after oxidized for 100 h, while the composites without coating retain only25% of their original strength.

The water vapor corrosion behavior of uncoated and SiOC–BSAS-coated C/SiCcomposites was carried out at the temperature of 1250 �C in an atmosphere of 50%O2–50% H2O with a flowing rate of 8.5 � 10�4 m s�1 (Fig. 35).

C/SiC composites without SiOC–BSAS coating undergo a considerable weightloss, while the ones with SiOC–BSAS coatings show little weight loss duringcorrosion in water vapor [225]. The results indicate that the polymer-derived

Fig. 34 The oxidation behavior of C/SiC composites with and without SiOC–BSAS coatings indry air at 1250 �C as a function of oxidation time: (a) weight change, (b) residual flexural strength.(Reproduced from Ref. [225] with permission from John Wiley and Sons)


SiOC–BSAS coatings can effectively block oxygen and the water vapor attack. Inthe case of uncoated composite, SiC matrix reacts with oxygen and water vapor toform SiO2 scale on the surface of the composite at the very beginning of water vaporcorrosion. Though this process results in weight gain, the silica formed reacts withwater vapor resulting in the formation of gaseous Si(OH)4 which causes weight loss.This recession process leads to the formation of cracks in the composite through whichoxygen and water vapor diffuse into the composite. As far as the residual flexuralstrength is concerned, C/SiC composite with SiOC-BSAS coating retains 80% of theoriginal flexural strength even after exposure to water vapor at 1250 �C for 200 h.

Liu et al. [244] have studied SiCN–Sc2Si2O7 as environmental barrier coatingsfor C/SiC composite using polysilazane as binder and Li2CO3 as sintering aid. Thedense SiCN–Sc2Si2O7 coatings were obtained after heat treatment at 1250 �C for 2 hunder argon. Addition of Li2CO3 as sintering aid lowered the sintering temperatureby about 200 �C. Water vapor corrosion of the C/SiC composites with and withoutSiCN–Sc2Si2O7 coatings were carried out at 1250 �C for 200 h. The results showedthat the SiCN–Sc2Si2O7 coatings could effectively block the water vapor attack athigh temperatures, and the sintering additive did not deteriorate the corrosionresistance of the SiCN–Sc2Si2O7 coatings.

Boron and silicon containing polymers are ideally suited for protection of C/Ccomposites as they would form mixed non-oxide ceramics on heat treatment, havingbetter thermo oxidative stability than simple non-oxide ceramics [245]. On oxida-tion, the mixed non-oxide ceramics would form borosilicate glass which can meltand flow into the cracks thereby sealing the pathways for oxygen entry[246]. Devapal [60] evaluated poly(phenylborosiloxane)-based SiC coating for theprotection of C/C composites. Initial attempts to prepare SiC coating by heattreatment of poly(phenylborosiloxane) with SiC inert filler has not been successfuldue to the poor adhesion of the coating onto C/C composite. In order to overcomethis problem, poly(phenylborosiloxane) was blended with addition-curable phenolicresin. Addition-curable phenolic resin was chosen in place of conventional phenolic

Fig. 35 Corrosion behavior of C/SiC composites (a) weight change and (b) residual flexuralstrength as a function of corrosion time with and without SiOC–BSAS coating. (Reproduced fromRef. [225] with permission from John Wiley and Sons)


resin as the former gives higher carbon residue. Elemental silicon and SiC were usedas fillers. The amount of silicon was chosen in such a way that there is no fee carbonleft in the sintered coating. The precursor coating was applied onto C/C compositecoupons by spray coating, cured at 150 �C, pyrolyzed at 900 �C, and then heattreated at 1450 �C. The coating process was repeated to get a crack free coating ofrequired thickness (150–200 μm). SiC coated and uncoated C/C composite couponswere subjected to oxidation test in a muffle furnace maintained at 1000 �C. Thesamples were removed at 5 min interval and allowed to cool to room temperature andweighed. The samples were exposed to the oxidizing environment for a total periodof 50 min. This exposure is a severe test as the sample is taken to high temperaturerapidly and brought back to room temperature suddenly which would cause thermalshock. Scanning electron micrographs of uncoated C-C composite and SiC coatedC-C composite before and after exposure to oxidizing environment are shown inFig. 36. The mass loss of uncoated C/C composite is about four times higher thanthat of SiC coated C-C composite. After 50 min of exposure to the oxidizingenvironment, SiC coated C/C composite loses 45.77 mg/cm2 whereas uncoatedC/C composite loses 197.05 mg/cm2. It is noticed that while the dimension of thesample remained the same in the case of SiC coated C/C composite, for uncoatedC/C composite the dimension of the sample decrease continuously on exposure tooxidizing environment. The mass loss of 45.77 mg/cm2 after 50 min of exposure for

Fig. 36 SEM of (a) C/C composite, (b) SiC coated C/C before, (c) C/C composite, and (d) SiCcoated C/C composite after oxidation [60]


SiC coated C/C composite suggests that there are certain defects in the SiC coatingthrough which oxygen gets entry to C/C composite causing erosion.

From the micrographs it is quite evident that the uncoated C/C compositeundergoes extensive surface erosion leading to degradation of both the matrix andthe reinforcement. No such surface erosion is noticed on the SiC coated C/Ccomposite. As discussed earlier, about 9.71% mass loss is noticed for the coatedcomposite on exposure to oxidizing environment. A closer look at the SEM of SiCcoated C/C composite before exposure to oxidizing environment reveals the pres-ence of micro-cracks (Fig. 37).

These cracks were probably formed when the sample was allowed to cool aftersintering. Though the cooling rate was slow (3 �C/min), the formation of crackscould not be eliminated probably due to the thermal expansion mismatch of C-Ccomposite and SiC coating and also due to the shrinkage caused by the evolution ofvolatiles from borosiloxane oligomer and addition-curable phenolic resin. Thecracks on SiC coating formed during sintering and in actual service can be sealedby providing an overlay glass coating. For this purpose, a borosiloxane oligomerdevoid of carbon was prepared by reacting boric acid with tetraethoxysilane bynonaqueous sol-gel process. At temperatures above 1400 �C, it melts and forms asmooth glass coating (Fig. 38).

Fig. 37 SEM images of SiC coated C/C composite exposed to oxidizing environment [60]

Fig. 38 SEM of glass coating from boric acid and tetraethoxysilane [60]


Possibility of using phenolic resin-based SiC coating for protecting C/C compos-ites as thermal protection system/thermostructural material for reusable launchvehicle, pylon leading edges, and splitter for air-breathing engines and reentrymissions have been explored by researchers at the Vikram Sarabhai Space Centre[247, 248]. Phenolic resins such as phenol-formaldehyde and phenol-furfuraldehydeor addition-curable phenolic resins with elemental silicon and glass forming addi-tives were used as the precursor for oxidation resistant coatings. In a multilayercoating process, the composition of each layer is suitably adjusted to achieve a finalcoating of desired thickness with good adhesion to the substrate. Glass formingadditives in the top most layer ensure that the coating is free from cracks. Eddycurrent (EC) methodology was used for nondestructive assessment of thickness ofSiC coating on C/C composite [249]. Unlike in other techniques, where approximatecoating thickness is determined by weight gain method, this technique can be usedfor evaluation of the coating thickness with error less than�5 μm. The methodologycould identify undercoated (thickness < 20 μm) specimens. Thus, it is possible toassess the efficacy of coating process and to readily identify thin coatings using Eddycurrent imaging. Based on this input, rework on the coating can be carried out toachieve uniform/desired thickness.

For qualifying the multilayer SiC coated C/C composite for use in reusablelaunch vehicle demonstrator, C/C composite coupons with multilayer SiC coating(thickness: coating of ~200–250 μm) were subjected to kinetic heat simulation(KHS) and arc-jet tests. The actual temperature encountered by the coating was~1600 �C. SEM images of the SiC coated samples before and after kinetic heatingsimulation test are shown in Fig. 39. Practically no mass loss was observed afterKHS test.

In order to improve the capability of SiC coating, after two layers of SiC coatingsilicide-based glass coating is given as the outermost layer. C/C composites with thismultilayer coating were subjected to the required heat flux in the plasma wind tunnelfacility. The coating withstood the test without undergoing erosion.

Fig. 39 SiC coated C/C sample (a) before, (b) during, and (c) after KHS test


C/C composite nosecap was coated with multilayer SiC coating of thick-ness ~ 250 μm [247] and was successfully flown in the India’s first Reusable LaunchVehicle Demonstrator (RLV-TD) (Fig. 40) [248, 250].

Yet another area of application for SiC coated C/C composites is in air breathingengines (scramjet engines). They are being developed as an advanced alternative forconventional jet engines, for improving the long-term performance and reusability.Conventional jet engines use a compressor to squeeze air into the engine, then sprayfuel into the compressed air and ignite it to produce thrust by funneling it through theback. On the other hand, a scramjet engine uses the speed of the aircraft to compressthe air into the engine and, hence, requires very few moving parts for its operation. Ascramjet engine consists of a constricted tube through which inlet air is compressedby the high speed of the vehicle, a combustion chamber where fuel is combusted,and a nozzle through which the exhaust jet leaves at higher speed than the inlet air.

Air breathing engines should have significantly higher specific impulse within theatmosphere than rocket engines. To achieve this, the incoming air should be directedinto the combustion chamber to attain supersonic compression. This is accomplishedby using leading edge structures as the tip portion of the inlet body. Its main functionis to direct the incoming air into the combustion chamber. One of the key issues indeveloping an air breathing engine is the oxidation protection of these leading edgesas the engine is operated at high temperatures. Predictions for the top speed of ascramjet engine vary between Mach 12 and Mach 24 (orbital velocity) [251]. Studiesoften plan on “active cooling,” where coolant circulating throughout the vehicle skinprevents it from disintegration. Often the coolant is the fuel itself, in much the sameway that modern rockets use their own fuel and oxidizer as coolant for their engines.All cooling systems add weight and complexity to a launch system and reduce itsefficiency. The increased cooling requirements of scramjet engines result in lowerefficiency. In place of such a complex system, a single material can be used whichcan withstand the following: (i) very high temperatures, (ii) steady-state and

Fig. 40 (a) RLV-TD with SiC coated C/C composite nosecap (inset), (b) RLV-TD in launch pad[248, 250]


transient localized heating from shock waves, (iii) high aerodynamic loads, (iv) highfluctuating pressure loads, (v) severe flutter, vibration, fluctuating, and thermallyinduced pressures, and (vi) erosion from airflow over the vehicle and through theengine.

SiC coated C/C composite is one of the primary candidate materials for use inair-breathing engine as combustion chamber, pylon leading edges, and sidewallleading edges. As the leading edge would face a severe thermooxidative environ-ment compared to the rest of the components, the focus has been to develop asuitable process for SiC coating of C/C composite leading edges. SiC coatingprocess similar to the one used for reusable launch vehicle using phenolic resin ascarbonaceous binder was used for obtaining SiC coated C/C composites. Both brushcoating and spray coating techniques were used to achieve a coating thickness (forthree layers) of ~200 μm. The SiC coated C/C leading edge (Fig. 41) obtained afterbuffing was subjected to arc-jet test, simulating the operational environment of an airbreathing engine. During the test, the leading edge was exposed to a heat flux of250 W/cm2 using a jet of argon plasma containing 25% O2 for duration of 14.5 s.The leading edge was allowed to cool in the test stand. The coating was free fromcracks. The test was repeated for a duration of 20 s and then for 40 s. The surfacetemperature profile of the leading edge during the third test (40 s duration) is shownin Fig. 42.

Even after three exposures, the coating was intact without any erosion. Thus, thesample was able to withstand the arc-jet test for the required flight conditions. As theSiC coated leading edge withstood the above test, another leading edge with SiCcoating was subjected to heat flux corresponding to air breathing propulsion and thecoating withstood the test without any erosion. During the actual use, the full lengthof the leading edge would experience the severity of the environment. In an attempt,

Fig. 41 SiC coated C/C composite pylon leading edge: (a) uncoated C/C composite, (b) SiCcoated C/C composite, (c) SiC coating after buffing


to simulate this condition, the width of the leading edge was reduced from 67 mm to15 mm so that the full length of the leading edge gets exposed to arc-jet. No erosionis observed, and the color change observed at tip is attributed to the loss of excesscarbon during arc-jet test (Fig. 43).

After successful testing of SiC coated C/C pylon leading edges, they were furthersubjected to ground level test under pressure. SiC coated pylon leading edges (width67 mm and radius of curvature 1 mm) were subjected to hot test under the followingtest conditions for a duration of 16 s as against the actual requirement of 7 s:temperature 1527 �C, air pressure 12 bar, oxygen pressure 2.5 bar (20.8%) byvolume, and Mach No. 2.5. The coating was intact without any erosion. Boundarylayer splitter (BLS) and Side Wall Leading Edge (SWLE) with two layers of SiCcoating and one top layer silicide coating were tested in arc-jet facility for qualifying

Fig. 42 Arc-jet testing (third test) of SiC coated pylon leading edge

Fig. 43 SiC coated leading edge (width 15 mm; radius of curvature 1 mm): (a) before and (b) afterexposure to arc-jet


these components in scramjet engine test. SiC coated pylon leading edges, BLS andSWLE were successfully tested in the scramjet engine technology demonstrator atMach 6 using Advanced Technology Vehicle in August 2016 [252].

Depending on the mission, the oxidation protection coating thickness and com-position required are to be fine-tuned to meet the mission objectives. The SiC coatingprocedure described for reusable launch vehicle and air-breathing engine dual moderamjet engine module could not meet the requirements of ballistic reentry. In aballistic reentry, TPS experiences stagnation temperature well above 2000 �C. Atypical temperature history of SiC coated C/C composite during ballistic reentrymission is given in Fig. 44 [253].

SiC coating is suited for use at temperatures below 1620 �C and above thistemperature under low partial pressures of oxygen, considerable erosion takesplace. When the temperature increases above 1620 �C the oxidation completelyshifts from passive to active oxidation, causing sudden increase in temperaturewhich is known as “Temperature Jump Effect” [254]. The reactions involving SiCwhich take place at different temperature regimes are given below:

SiCþ 2SiO2 ! 3SiOþ CO 1420�C

� �

SiCþ SiO2 ! 2SiOþ C 1550�C

� �

2SiCþ SiO2 ! 3Siþ 2CO 1620�C

� �

Below 1420 �C passive oxidation of SiC takes place forming SiO2 which pre-vents further oxidation of SiC. Transition from active oxidation to passive oxidationtakes place forming gaseous SiO in the temperature range of 1420–1620 �C, andabove 1620 �C, active oxidation only occurs resulting in considerable erosion of thecoating. Increasing the coating thickness from 250 μm to 500 μm could not protect

Fig. 44 Temperature historyof SiC coated C/C compositeduring ballistic reentry [253]


C/C composite thermal protection system. Sreeja et al. [255] investigated thepossibility of giving an unsintered coating over SiC coating which will take awaythe initial heat by way of pyrolysis of carbonaceous polymer used as a binder andbring down the temperature below 1620 �C. This concept was tested using graphitesubstrate (Fig. 45).

It is observed that the unsintered sealant coating is able to bring down thestagnation temperature by about 600 �C [255]. Hence, this sealant coating wasgiven over SiC coating to shift the active oxidation to passive oxidation. Yet anotheradvantage of giving a sealant coating is that it would seal the cracks in SiC coatingformed during processing. This approach has been successfully adopted for oxida-tion protection of C/C composite noseshell (Fig. 46).

Carbonaceous polymer-based coating process which was used successfully forprotecting C/C composites was not adhering well to C/SiBOC composites and amultilayer coating process was developed in which pyrolyzed borosiloxane oligomer

0 10 20 30 40 50 60 70 80 90 1000

500

1000

1500

2000

2500

(b)

)C°(

erutarepmeT

Time (sec)

(a)

Fig. 45 Temperature profileof (a) uncoated and (b) sealantcoated sample [255]

Fig. 46 SiC coated C/Cnoseshell (dia. 960 mm) withsealant coating [255]


was used as a filler in addition to glass forming additives. The coating withstood theheat flux history of reentry vehicle [256].

Oxidation Resistant/Environmental Barrier Coatings on MetalsHigh temperature metals and alloys are extensively used as hot structural compo-nents in aerospace industry as turbine blades, exhaust nozzles, actively cooledcombustion chambers, etc. [257]. These metal surfaces are to be protected at hightemperatures in order to prevent oxidation by corrosive gases and the catalyticactivity of the metal surface which enhances the surface temperature. Hence, it isrequired to design a suitable protection system comprising of inhibitors, sealants,and coatings for these metallic hot structures at high temperature [258]. Differentprecursors such as polymethylsilsesquioxane [259], poly(hydridomethylsiloxane)[260], polysiloxane [261, 262], polymethoxymethylsiloxane [262], and hydroxy-terminated linear polysiloxane [262] were investigated on substrates such as steel,stainless steel, titanium, aluminum, magnesium, and zinc using various coatingmethods such as spray coating, dip coating, spin coating, and PE-CVD. Polymer-derived ceramic coatings for oxidation protection of metal substrates have beendescribed in recent reviews [13, 263], and hence, the discussion is confined to a fewof the important developments in this area.

Günthner et al. [264] prepared PHPS by the ammonolysis of dichlorosilane andABSE polycarbosilazane by ammonolysis of bis(dichloromethyl)silylethane, andevaluated them as protective coatings for SS. PHPS and ABSE, on heat treatmentat 800 �C, form ceramic-like gradient coatings, which successfully protected stain-less steel sheets from oxidation up to 1000 �C. The gradient Si(N)CrO layers act as adiffusion barrier against oxygen and reduce the weight gain and the oxidation kineticconstant up to two orders in magnitude (Fig. 47).

Pyrolysis of precursor to ceramics results in porosity and residual stress caused bythe shrinkage. These problems can be overcome by the use of passive fillers like BN,ZrO2, or Al2O3 and active fillers like ZrSi2 and TiSi2 [264]. Günthner et al. [265]formulated a high performance environmental barrier coating system consisting of aperhydridopolysilazane (PHPS) bond coat and a polysilazane-based glass/ceramiccomposite top coat for mild steel substrates. The thermal treatment of the coatingwas performed in air at 800 �C. The critical coating thickness of the coating systemswas increased up to 100 μm by optimizing the volume fraction of the fillers and theprocessing parameters. Cyclic oxidation tests showed that the coating system pro-tects mild steel against oxidation up to 700 �C. Wang et al. [266] studied PDC matrixcomposite coatings for oxidation protection of Inconel 617, a superalloy, using asimple dip coating method. Three coating systems, viz., SiON coating, particle filledSiOC coating, and their combined coating system, were evaluated by conductinglong time static oxidation test at 800 �C for 100–200 h. The double layer coatingsystem comprising of SiON bond coat and particle filled SiOC top coat was found tobe the most effect one in maintaining the integrity of the substrate subjectedoxidation test for 200 h.

Parchoviansky et al. [267] also investigated a double layer coating system with aPDC bond coat and a glass/ceramic filled PDC top coat as an effective solution for


overcoming the disadvantages of PDC-based coatings. The bond coat has multipleroles in the coating system: it protects the metallic surface, increases the adhesionbetween the surface and the top coat, and reduces the CTE mismatch between themetal and the top coat. A double layer coating consisting of polysilazane basedbond-coat was prepared from both PHPS and Durazane1800® (prepared by theco-ammonolysis of dichloromethylvinylsilane and dichloromethylsilane), and testedfor its high temperature oxidation resistance. The top coat contained high meltingSiO2-Al2O3-ZrO2 glass microspheres. At 900 �C, the TBC has reduced the massgain due to oxidation to about one-half of that in uncoated steel oxidized under thesame conditions.

Riffard et al. [268] evaluated commercially available polysilazane resin (withoutany filler) for protecting AISI 304 Stainless Steel from oxidation. The coating wasapplied either by dip or spray coating process. During pyrolysis of the coating at800 �C for 1 h in air, the initial PHPS-based coating reacted with oxygen leading tothe formation of an amorphous silica layer at the steel surface which acts as adiffusion barrier for the chemical diffusing species. This amorphous silica layeracts as a protective glass which provides complete separation of the AISI 304 steelsurface from the oxidizing environment.

Polysiloxane-based EBC system for steel in oxidizing environments was devel-oped by Nguyen et al. [269]. Different fillers such as alumina, Al flakes, Cr, Si, W,Zn, and glass frits were used to get a coating thickness of ~36.7 μm. The coatingshowed good adhesion with the substrate, owing to the lower shrinkage of theceramic due to the presence of fillers. The presence of glass frits improved the

Fig. 47 (a) PHPS coated steel substrate oxidized at 1000 �C for 1 h in air, (b) SEM of the coatedsubstrate oxidized at 500 �C for 1 h in air, and (c) Oxide layer thickness after oxidation at 1000 �Cfor 10 h in air. (Reproduced from Ref. [264] with permission from Elsevier)


adhesion because of the dissolution of the oxide layer already present on the metal atthe interface into glass.

Molybdenum alloys containing boron and silicon are replacing nickel-basedsuper alloys, due to their high physical-mechanical characteristics at high tempera-tures [270]. On exposure to air, Mo alloys form MoO3, which evaporates in thetemperature range between 400 �C to 800 �C resulting in poor oxidation resistance ofthese alloys. Mo5SiB2 filled PHPS was used as coating for protecting Mo-Si-B alloy[270]. This coating was effective in reducing the oxidation rate by 50% at 1100 �C.The improvement in oxidation resistance is attributed to the presence of amorphousSiOxNy phase and silica phase formed from free Si present in the PHPS-derivedceramic.

Thermal Barrier Coatings for MetalsOxidant resistant coatings and environmental barrier coatings for metals described inthe preceding section may not bring down the surface temperature of metallicsubstrate. In order to induce a temperature drop on the metal surface, it is essentialto develop thermal barrier coatings (TBCs) [271, 272] the concept of which is givenin Fig. 48.

The earliest TBCs were frit enamels based on alumina/zirconia which wereapplied to aircraft engine components in the 1950s [273]. The effectiveness ofthese coatings was limited by the relatively high thermal conductivity of aluminaand low phase stability of zirconia-based materials. The use of TBCs has resulted ina significant improvement in the efficiency of aircraft gas turbines [274] and thesame is achievable in land-based gas turbines. By using TBCs temperature differ-entials across the coating of as much as ~300 �C can be achieved. This temperaturedrop reduces the thermally activated oxidation rate of the bond coat applied to metalcomponents and hence, delays failure by oxidation. It also retards the onset ofthermally induced failure mechanisms that contribute to component durabilityand life.

Fig. 48 Schematic of a typical TBC showing distribution of temperature (BC: bond coat; TGO:thermally grown oxide). (Reproduced from Ref. [272] with permission from Elsevier)


https://www.sciencedirect.com/topics/materials-science/boron

One of the important developments was the introduction of NiCrAlY bond coatsand yttria-stabilized zirconia (YSZ) topcoats in the mid-1970s. YSZ has become thepreferred TBC layer material because of its low thermal conductivity, relatively highcoefficient of thermal expansion (CTE), and good erosion resistance [275]. Whilethe low thermal conductivity facilitates in bringing down the surface temperature,the high CTE helps to increase the thermal stress capability. In practice, yttriaconcentration is kept in the range of 6–8 wt%, since this composition results in theformation of a metastable tetragonal phase, which resists the monoclinic transfor-mation of the composition upon cooling. This maximizes spallation life of coatingsby inhibiting crack propagation and volume change upon cooling.

By attaching an adherent layer of a low thermal conductivity material to the metalsurface, a temperature drop can be induced across the thickness of the layer. Thisresults in a reduction in the metal temperature of the component to which it isapplied. Using this approach, temperature drops up to 170 �C at the metal surfacehave been estimated for 300–400 μm thick yttria-stabilized zirconia coatings. Thistemperature drop reduces the (thermally activated) oxidation rate of the bond coatapplied to metal components, and so delays failure by oxidation. It also retards theonset of thermally induced failure mechanisms that contribute to component dura-bility and life.

In launch vehicles for regions where the flight-induced heating is substantial,thermal protection system (TPS) is used on structural members so that the surfacetemperature is brought below the temperature where the structural strength degra-dation is substantial. On the other hand, a hot structural material is that whichfunctionally sustains both the thermal loads as well as the structural loads. The hotstructural members in reentry/reusable vehicles are its nose-cap, wing leading edges(WE), elevons, and rudders of which the nose-cap is made of C/C and WLE andvertical tail edges are made of 15CDV6/Inconel metallic alloy. These alloys are alsocandidate materials in the construction of air-breathing engines. The protective TBCcoating required for enabling the application of these alloys must have high emis-sivity and be thermally and structurally stable during heating and cooling phase.

Different techniques are used for deposition of TBCs [275–278]: (i) physicalvapor deposition (PVD), (ii) electron beam physical vapor deposition (EB-PVD),(iii) chemical vapor deposition (CVD), (iv) plasma spraying, and (v) polymer-derived ceramic (PDC) route. Of the methods used for TBCs, PDC route is themost recent one which takes advantage of high ceramic yield of the precursors, highpurity of starting materials, lower capital investments, lower processing tempera-tures, variation of micro porosity to tailor the microstructure, and optimization oferosion resistance.

Specifically, for space application, a four-layered coating system (Fig. 49)consisting of bond coat, thermal insulative layer, alumina layer, and final highemissive layer using siloxane as binder has been developed [279] for 15CDV6steel substrate and the coating performance was evaluated using IR lamp heatingsimulations at a heat flux of 8.5 W/cm2 for 1035 s. The tests showed good results forthe ceramic coating system with a back wall temperature of 299 �C. But the coatingpeeled off during cooling process. This problem has been overcome by using


polysilsesquioxane as binder and by modifying the coating process [280]. Themodified TBC was tested successfully at 8.5 W/cm2 for 1035 s (Fig. 50).

Atomic Oxygen Resistant CoatingsThe LEO space is a complex, energetic environment, which ranges from an altitudeof 200–700 km and from an inclination of 0–60� above the equatorial plane[281]. Effects of different constituents of the space environment on spacecraftmaterials play a crucial role in determining the system function, reliability, andlifetime. Direct entry to the low earth space environment causes profound alterationsof thermal, mechanical, and optical properties of polymeric materials [282]. Theseeffects typically depend on interactions between the spacecraft and the diverseelements of the space environment itself such as: (i) atomic oxygen (AO) onspacecraft materials, (ii) micrometeroid and debris impact, (iii) ultra-high vacuum,(iv) far ultraviolet hard radiation and particles generally, (v) charged particlesbombardment, and (vi) space craft charging. In particular, one of the most importantfactors that gives serious damage to many polymeric materials is AO which is a

Fig. 49 Schematic of TBCon metallic substrate usingsiloxane as binder [279]

Fig. 50 TBC coating developed via spray coating technique (a) before and (b) after IR lampheating test [280]


dominant neutral species in LEO, formed by the photo-dissociation of molecularoxygen as shown below:

O2 þ hυ ! O þ O

O2 þ hυ ! Oþ þ O� 1%ð ÞAO formed by photodissociation has a very high probability of long-term sur-

vival in the low earth altitudes because there is an appropriate O2 density here tofacilitate reasonable AO production and low probability of interaction with neigh-boring atoms or molecules. The number density of AO at about 300 km altitude is~2 � 109 (minimum) to 8 � 109 atoms/cm3 (maximum) during solar activity[281]. At LEO altitudes, the neutral atmosphere consists primarily of AO (80%)and nitrogen molecules (20%).

Though the average thermal velocity of the gas molecules at low earth altitudes isvery low, the spacecraft orbiting in the LEO rams into the AO environment at avelocity of 8 km/s producing collision energy in the range of 4.5–5 eV which isadequate enough to cause bond scission in spacecraft materials. Thus, the interactionof AO with materials used in the construction of spacecrafts and space stations mayresult in surface erosion, changes in chemical composition and surface morphology,changes in optical properties and formation of particulate and molecular contamina-tion of the spacecraft/space station surfaces [282–284]. To over come this problem, itis necessary to develop materials that are resistant to AO attack or protect thematerials susceptible to AO attack with AO resistant coatings. Metal oxide coatingsthat have negligible rates can serve as AO resistant coatings. Lack of flexibility,requirement of elaborate sputtering techniques, and pinhole defects of metal oxidecoatings are of major concern. Inorganic and inorganic-organic hybrid polymerssuch as polysiloxanes [285, 286], polysilsesquioxanes [287], polysilahydrocarbons[288], polyphosphazenes [289], phosphorus containing polyimides [290–292],siloxane-imides [293], polyhedral oligomeric silsesquioxane (POSS)-polyimides[294], siloxane-epoxy [295], siloxane-imide-epoxy [296], carborane-siloxanes[297], decaborane-based polymers [298], and polysilazanes [299–301] are capableof overcoming these problems. These polymers on exposure to AO form a thin oxidelayer, preventing further reaction of the coating with AO. Even if, AO penetrates intothe next layer due to surface defects, the unoxidized virgin polymer underneath theoxide layer would provide self-healing. The thermal expansion coefficients of thevirgin polymer and the polymeric substrate are usually close to each other due towhich the crack formation on the surface of the coating would be prevented. In thissection, the discussion is limited to a few representative precursor systems toexemplify the use of virgin preceramic polymers for yet another space applicationand a detailed account on AO resistant materials/coating is beyond the scope of thissection.

Devapal et al. [288] evaluated the AO resistance of a polydisilahydrocarbon, viz.,poly(tetramethyldisilylene-co-styrene) [PTMDSS], whose structure is shown inFig. 51. PTMDSS offered good protection to polyimide film and carbon compositeagainst AO attack which is clearly seen from the SEM images (Fig. 51). However,


this coating was somewhat tacky in nature and this problem was overcome by partlyreplacing dimethylsilyl units with phenylmethyl silyl units [60].

Chang et al. [301] synthesized a polysilazane by aminolysis ofdichloromethylvinylsilane, and the polymer was evaluated for its AO erosion onKapton film in a ground-based simulator. The mass loss, optical properties, and SEMsurface morphology of the samples before and after AO exposure indicate that thepolysilazane coating protects Kapton from AO erosion effectively (Fig. 52).

AO resistant coatings have also been obtained by sol-gel process. Arjun et al.[285] synthesized a polysiloxane by acid catalyzed hydrolysis reaction ofmethyltriethoxysilane and 1,1,3,3-diethoxytetramethyldisiloxane and used thisresin for protecting aluminized Kapton, carbon polyimide (C-PI) composite, andglass polyimide composite from AO attack. The mass loss data (Table 6) andFESEM of uncoated and polysiloxane coated samples exposed to AO suggest thatpolysiloxane obtained by sol-gel process offers excellent protection against AOattack.

Fig. 51 Evaluation of PTMDSS as atomic oxygen resistant material: SEM of (a) uncoatedaluminized Kapton

®

film, (b) PTMDSS coated aluminized Kapton®

film, (c) uncoatedC-polyimide composite, and (d) PTMDSS coated C-polyimide composite exposed to atomicoxygen. (Reproduced from Ref. [288] with permission from John Wiley and Sons)


Fig. 52 AO erosion kinetics for Kapton and Silazane coated Kapton and SEM of (a) Kapton film,(b) exposed Kapton, (c) coated Kapton, and (d) exposed coated Kapton. (Reproduced from Ref.[301] with permission from John Wiley and Sons)


Ceramic Adhesives

The engineering designs of ceramics and composite components in aeronautics,space transportation, energy, and nuclear industries require fabrication andmanufacturing of large and complex shaped parts of various thicknesses. In manyinstances, it is more economical to build up complex shapes by joining simplegeometrical shapes. Thus, joining and attachment have been recognized as enablingtechnologies for successful utilization of ceramic components in various demandingapplications [302–304]. In space applications, C/C-, CMC-, and UHTC-basedmaterials are being evaluated in structural parts such as leading edges, nosecone,and nozzles for hypersonic vehicles, which require proper joining technologies.Requirement also exists for non-space applications such as joining of C/C to Cualloys and SiC/SiC to SiC/SiC for nuclear components and to join radiator subelements for heat rejection systems. Different types of joining techniques, viz.,brazing, diffusion bonding, riveting, bolting, and adhesive bonding, are currentlyunder usage/progress for ceramic components. Advanced technologies such asceramic joining using affordable robust ceramic joining (ARCJ) technology, bond-ing of ceramics to metal by brazing, microwave assisted joining, transient liquidphase bonding, and spark plasma joining have also been developed. The majordisadvantage of these techniques lies in huge energy wastage to attain high temper-atures, high facility cost of the process, and severe damage to the substrates.

The main technical and performance requirements of a ceramic adhesive joint are:(i) use temperature> 1200 �C, (ii) good strength and oxidation resistance properties,(iii) low CTE mismatch to minimize residual stresses and good thermal shockresistance, (iv) leak tight joints, (v) reliable and affordable technique adaptable toin-field installation, and (vi) service and repair [304]. The other technical challengesinvolved are in the design and selection of joints, materials related issues, thermo-mechanical properties of the joints, and environmental effects on joint properties. Ingeneral, on comparing with other joining techniques, adhesives contribute highly tostructural integrity, ease of manufacturing, enhanced performance, improved safety,and cost and time savings. The continuity of the adhesive bond plays a vital role asthe loading is borne by the whole bonding plane, not by one or several points.However, the main problem with the application of adhesive for spacecraft is hightemperatures. The high temperature adhesive should possess good bonding proper-ties with the various substrates and should withstand the high temperature and

Table 6 Mass loss data of aluminized Kapton, C-polyimide composite, and Glass-polyimidecomposite without and with polysiloxane coating exposed to AO fluence of 1.5 � 1020 atoms/cm2.(Adopted from Ref. [285])

SubstratesMass loss (mg/cm2) ofuncoated samples

Mass loss (mg/cm2) ofpolysiloxane coated

Aluminized Kapton 0.48 0.08

C-polyimide composite 4.02 0.09

Glass-polyimide composite 1.77 0.16


various stresses that the system may encounter during reentry. A class of materialwith such high temperature resistance is ceramic adhesives. Like any other adhesivebonding, the uniform stress area is maintained with a low strain rate during mechan-ical tests. Further, due to the thin layer of the adhesive, it essentially behaves like athin film and this would limit the generation of triaxial state of stress within thebonded area leading to overcoming the brittle behavior of ceramics, resulting in itsusefulness for structural applications in aerospace industry. The polymeric adhesives(modified epoxies, polyimides, etc.) would undergo cross-linking or chain scissionwithin the polymeric matrix on exposure to high energy radiations, and there is aconsiderable influence on the mechanical strength of the polymer, while the ceramicadhesives after processing are capable of withstanding high energy radiation andthere is no deterioration observed in joint strength. The normal high temperatureceramic adhesives comprise of inorganic materials based on fluorides, sulfides, andoxides of transition and non-transition metals of nearly eutectic composition dis-persed within the ceramic matrix. However, if the strain rates are high, there wouldbe limitations in applying this adhesive although the other performances are good.Thus, a series of heat resistant and high structural adhesive for space applications forfrequent usage in heat resistant environments has to be developed.

As moved by the tragedy of failure of Columbia in February 2003, worldwideextensive research has been initiated for developing ultra-high temperature ceramicadhesives to repair cracks and holes on the shuttle wings and nose during voyage ofmanned mission. NASA has developed an adhesive called Non-Oxide AdhesiveeXperimental (NOAX), a high-tech formulation of preceramic sealant (hydridopo-lycarbosilane) mixed with silicon carbide powder and other fillers designed to meetNASA’s stringent specifications for space flight. The adhesive was used to fill cracksor losses in protective coatings up to 0.5 mm wide and 100 mm long, and it wasdesigned to get semirigid after exposure to several hundred degrees of heat, butwould not harden completely until it experienced the extreme temperature of reentryinto earth’s atmosphere [305–307]. Another ceramic adhesive, viz., Glenn Refrac-tory Adhesive for Bonding and Exterior Repair (GRABER), was developed at theNASA Glenn Research Centre for the repair of reinforced carbon/carbon (RCC)composite thermal protection systems structures [308]. The GRABER material hasshown multiuse capability for the in-space repair of small cracks and as an adhesiveand sealant for RCC-leading edge material. It is a refractory adhesive paste withdesired ceramic fillers in polymer/phenolic resin matrix with appropriate additivessuch as surfactants. The paste was applied to the damaged or cracked area of theRCC composite components with a caulking gun, and cured at 100–120 �C whichtransformed into a high temperature ceramic during reentry conditions. The repairedsamples were successfully tested (Fig. 53) under the arc-jet facility at the NASAJohnson Space Centre.

Sreeja et al. [255] formulated a ceramic adhesive useful for joining differentsubstrates for space applications. Phenol-formaldehyde/phenol-furfuraldehyderesins were mixed with elemental silicon and glass forming additives to formulateadhesives which cured at temperatures below 200 �C. The adhesive underwent insitu polymer to ceramic conversion at high temperatures/re-entry conditions with the


simultaneous formation of SiC and silica-based glass which sealed the cracks formedduring the conversion process. This was proved by carrying out arc jet testing ofgraphite specimens joined together using phenolic resin at reentry conditions. Thehigh temperature ceramic adhesive was also evaluated for crack repair of C/Ccomposites and CMCs. A cavity of dimension 20 mm � 2 mm � 1 mm was madein SiC coated C/C coupon and sealed with the high temperature adhesive, cured andtested under arc-jet (Heat flux of 300 W/cm2 for 50 s; 20% O2) (Fig. 54). Therepaired part remained intact (Fig. 55). There was negligible mass loss observed andthe joint was intact.

Besides, carbonaceous polymers and hydridopolycarbosilane, other preceramicpolymers have also been evaluated for applications as high temperature adhesives[309–317]. The development of ceramic adhesives from preceramic polymer route isnoted for its easy application and significantly lower processing temperature. Also,there is no problem of residual reactant products, which may adversely affect thejoining as in the case of reaction based approach. However, large volume shrinkageassociated with the conversion of preceramic precursors to covalently bonded

Fig. 53 RCC specimen with crack (width: 0.89 mm) repaired: (a) pre-arc jet test and (b) post-arcjet test. (Adopted from Ref. [308])

Fig. 54 Graphite test specimen assembled together using high temperature adhesive: (a) arc jettesting and (b) tested specimen [255]


network structures leads to formation of cracks during processing. The use of inertfillers such as SiC, B4C, Al2O3, and silica and active fillers such as Al, Si, Si-Al, Ti,MoSi2, CrSi2, and TiSi2 improves the strength and quality of joints [309]. The finalproperties of joints such as maximum temperature of use, chemical compatibilitywith substrates or fillers, thermal expansion, gas evolution, yield, wetting behavior,processing procedure, joint density, presence of cracks, joint thickness, shrinkagestresses, joint microstructure, and composition are found to be affected by several

Fig. 55 (a) Cavity in SiC coated C/C composite, (b) cavity filled with high-temperature adhesive,and (c) SEM of repaired part after test [255]


parameters. These parameters are the following: polymer composition, polymerceramic yield, rheology of joining mixture, heat treatment (maximum temperature,dwell time, heating rate, atmosphere), type and amount of fillers, and type ofsubstrate (composition, surface morphology) [309].

A new class of material, inorganic geopolymer, has been intensively studied as aviable economical alternative to organic polymers and inorganic cements in devel-oping refractory adhesives for military and aircraft applications [318]. Geopolymer-based adhesives have exceptionally high thermal and chemical stability, compressivestrength, and adhesive behavior. Geopolymers are inorganic aluminosilicates formedby polycondensing tetrahedral silica (SiO4) and alumina at ambient or low temper-atures (<150 �C) into polymeric structures. Lower ratios of SiO2:K2O ¼ 1–3 resultin three-dimensional cross-linked rigid networks with brittle properties, whereashigher ratios of SiO2:K2O ¼ 3–15 give raise to two-dimensional/linearly linkedstructures with adhesive and rubbery properties. Compared to silicate-based adhe-sive compositions, aluminosilicate-based adhesive compositions (geopolymers)have excellent physicochemical and mechanical properties, which include lowdensity, micro-porosity or nano-porosity, negligible shrinkage, high strength, ther-mal stability, high surface hardness, and fire and chemical resistance. They have highmelting points but can undergo structural transformations at lower temperatureswithout causing the materials to change dimension and porosity. These desirableproperties render these materials as ceramic adhesives. The alkali silicate solid orliquid solution is used for dissolution of raw materials to form reactive precursorsrequired for geopolymerization.

Sreeja et al. [319] developed an alkali activated aluminosilicate geopolymer forbonding metals to ceramics and metal to metal, for high temperature applications.The bonding was achieved at 175 �C by solid state reaction of alkaline solution ofalkalisilicate precursor with the refractory filler, contributing to the bulk aluminosil-icate matrix. This adhesive was used for bonding silica tile to Inconel (Fig. 56). Thebonding remains intact even after heating to 500 �C in a muffle furnace.

Fig. 56 Inconel/silica tilebonded specimen after heatedto 500 �C in a muffle furnace[319]


Lightweight Ceramics

Lightweight ceramics include ceramic foam, ceramic honeycomb, ceramic sandwichstructures, and shingles. These materials have been evaluated for space applicationsas lightweight thermal protection systems, micrometeoroid and debris protectionsystems, and space mirrors.

Cellular Ceramics/Ceramic FoamCellular ceramics are ceramic materials with good level of porosity, where theporosity is present as repeating or non-repeating hollow cells. The cells can beregular in size and shape as in a honeycomb structure or irregular as in a foamstructure. Other architectures also exist such as interconnected structures [320]. Cel-lular materials are characterized by low density, high specific area, high tortuosity offlow paths, and low thermal conductivity [321].

The characteristics of cellular ceramic depend on the morphology of the cellswhich is strongly influenced by the processing methodology. Each fabricationmethod is best suited for producing a specific range of cell sizes, cell size distribu-tion, and overall amount of porosity [322]. Wide range of processing routes such asreplication technique or polymeric sponge method, direct foaming method, andburn-out technique were developed to produce ceramic foams. The basic principlesof these methods are schematically presented in Fig. 57 [323].

CVD/CVI method is also used for the preparation of ceramic foam [324]. Thepreparation of ceramic foams adopting these techniques has been extensivelyreviewed by several authors [320–326], and hence, a detailed account of thesetechniques has been avoided in this review.

The use of preceramic polymers for the fabrication of cellular ceramics offersvarious advantages over the conventional route using ceramic powders in a slurrycontaining stabilizers, binders, and sintering aids. Some of the problems faced in theconventional route are: difficulty in maintaining good suspension of the solidparticles, achieving good infiltration into porous templates, controlling the viscosity,structure collapse due to binder burnout, and high temperatures required forsintering. Use of preceramic polymers successfully addresses these issues and pro-vides additional benefits such as control over the nano-level ceramic composition,ability to adapt to novel processing methods, and the option to easily machine theprecursor foam after curing.

Fabrication of ceramic foams using preceramic polymers is done via threedifferent routes: (i) using replica, (ii) using porogens, and (iii) direct foaming [327,328]. The replica method and the use of porogens result in open cell foams, anddirect foamingmethod can result in both closed and open celled foams [329]. Thoughreplica method is widely used industrially for the fabrication of open cell ceramicfoams, mechanical properties are lower due to the presence of hollow strut leftbehind by the fugitive replica. In this context, the other two methods offer greatpotential for improving the properties.

Among the preceramic polymers, polycarbosilane [330–332] and polysiloxane[333–335] are highly utilized for the fabrication of cellular ceramic materials.


Polysiloxanes have the advantage over PCS in terms of lower cost, availability ofmultitude of variants in bulk, and the ability to self-foam. The self-foaming nature ofpolysiloxane results from the presence of –OH and –R groups in the polymerstructure, which on heating cross-links and generates alcohols and water whichescape during the curing, forming the porosity [336].

In general, it is observed that the compressive strength of SiOC foams decreasewith increase in pyrolysis temperature from 1000 �C to 1400 �C [334]. This may beattributed to the carbothermal reaction, resulting in SiO evolution, which in turnweakens the struts. Various attempts have been made to improve compressivestrength by nano-casting approach [337], by compressive molding followed bysintering at 1750 �C [338] or by using sintering aid [339] or SiC filler [340]. Aninteresting point to note here is that whereas SiOC foams lose strength on heattreatment at higher temperature due to carbothermal reaction, causing damage to theload bearing struts, SiC foam can retain the strength due to the absence ofcarbothermal reaction. In fact, it is reported that SiC foams from polycarbosilaneshow a decrease in pore size [341], which can enhance its load bearing

Fig. 57 Schematic representation of fabrication process of ceramic foam. (Reproduced from Ref.[323] with permission from John Wiley and Sons)


characteristics, with increased strut density. PIP process also provides an option todensify the struts as the polymers can infiltrate the polyurethane template resulting indense struts [342].

Silicon-boron-oxycarbide (SiBOC) foam was prepared from poly-methylborosiloxane resin [343]. The foam showed good mechanical integrity andshape retention with shrinkage. The real density of the struts was measured to be1.93 g/cc. Geometrical density showed a change from 0.15 g/cc for the polymericfoam to 0.18 g/cc for the ceramic foam. SEM images of the cross-section of strut areshown in Fig. 58. The images clearly depict the evolution of SiBOC hollow strutfrom the triangular polyurethane strut.

Kim et al. [344] converted polysiloxane-phenolic resin mixture into microcellularSiC ceramics with cell density greater than 109 cells/cm3 and cells smaller than20 μm. This involved the following steps: (i) fabricating a formed body from a

Fig. 58 SEM images of the strut cross-section of the foam obtained at different stages of itspreparation: (a) of the template PU foam, (b) of the template after first cycle of coating withmethylborosiloxane and cross-linking, and (c) of the final SiBOC foam. (Reproduced from Ref.[343] with permission from John Wiley and Sons)


mixture of polysiloxane, phenolic resin (used as a carbon source), polymer micro-beads (used as sacrificial templates), and Al2O3–Y2O3 (an optional sintering addi-tive), (ii) cross-linking the polysiloxane in the formed body, (iii) transforming thepolysiloxane and phenolic resin by pyrolysis into silicon oxycarbide and C, respec-tively; and (iv) synthesizing SiC by carbothermal reduction. By controlling themicrobead and additive contents, it was possible to adjust the porosity so that itranged from 60% to 95%.

Novel routes for the fabrication of cellular materials are also emerging. A methodfor sizing and aligning pores via freeze casting is reported [345]. The pores align inthe freezing direction of the polymer and the pores become narrower with freezingtemperature [346]. Macro/mesoporous SiOC ceramic monoliths of anisotropic struc-ture were prepared by freeze casting, using methylphenyl polysiloxane or methylpo-lysiloxane and (3-aminopropyl)triethoxysilane as precursors. The very low thermalconductivities and relatively high compressive strength are very promising for use incryogenic engineering [347]. Porous SiC ceramics with low oxygen content wereprepared by freeze casting using divinylbenzene (DVB) as a cross-linker for PCS. Aclear improvement in compressive strength with increasing PCS content is seen,with the highest compressive strength of 18.7 MPa found for samples with 20 wt%[348].

Recently, ceramic aerogels were prepared by PDC route [349, 350]. A combina-tion of ultra-low density and low thermal conductivity makes aerogel a sought afterTPS material for space applications. Radiative conduction in foam core TPS mate-rials can be reduced by infiltrating aerogels into the pores of the foam, which in turnadds to the strength of the foam structure also.

In addition to conventional preceramic polymers, phenolic resin has also beeninvestigated as precursor for SiC foam [192]. Phenolic resin was blended with SiCand elemental silicon, and acetone was used as solvent to adjust the viscosity of theslurry. Polyurethane foam was infiltrated with this slurry, cured, pyrolyzed and thenheat treated at 1500 �C to obtain SiC foam having density of 0.15 g/cc and averagepore diameter of 350 μm. SiC foam specimens having density varying from 0.15 to0.75 g/cm3 were prepared by controlling the number of densification cycles using thesame slurry.

Porous SiC bodies were prepared by powder compaction method [192]. Thepolyurethane foam infiltrated with precursor slurry and cured at 75 �C was crushedto a fine powder. It was mixed with phenolic resin-based precursor slurry and moldedunder pressure to form the desired shape and cured at a temperature of 175 �C. Thisgreen body was then removed from the mold and pyrolyzed at 900 �C followed bysintering at 1500 �C under inert atmosphere for allowing the formation and densi-fication of SiC. Further densification was done by repeated infiltration with pre-ceramic slurry followed by pyrolysis and sintering. By this process, porous SiCbodies having density in the range 0.9 g/cc to 1.75 g/cc were prepared.

Porous SiC bodies prepared by densification of SiC foam and powder compactionmethod described above were machined to get leading edges and subjected to arc-jettest at 250 W/cm2 in air for 15 s (Figs. 59 and 60). Leading edge prepared bydensification of SiC foam withstood the test without any erosion or mass loss


whereas slight erosion was observed for the leading edge prepared by powdercompaction method.

Ajith [192] explored the possibility of using SiC foam tile as a hot structure in anoxidizing environment. SiC foam tile of density 0.75 g/cm3 (size: 150 mm �150 mm � 20 mm), prepared by densifying SiC foam of density 0.2 g/cm3 bypolymer-infiltration-pyrolysis (PIP) process using phenolic precursor slurry, wassubjected to kinetic heat simulation (KHS) test.

KHS tests were conducted at two different heat flux history (HFH) for thewindward region of a typical reentry vehicle with a maximum heat flux of4.1 W/cm2 and 17.5 W/cm2, corresponding to two different trajectories. Maximumtemperature of 426 K and 742 K was experienced at the backwall of SiC tile uponexposure after ~500 s and ~ 600 s, respectively, when the tile was subjected to twodifferent heat flux history. Inspection of the tile after these two tests indicates that thetile retains its integrity without any weight loss (Fig. 61).

Fig. 59 Leading edge after arc-jet test: (a) densification method and (b) powder compactionmethod [192]

1800

1600

1400

1200

1000

800

600

400–4 0 4

Arc on

Arc off

8

Time (s)

Tem

pera

ture

(°C

)

12 16 20

1800

1600

1400

1200

1000

800

600

400

–4 0 4

Arc on

Arc off

8

Time (s)

Tem

pera

ture

(°C

)

12 16 20

a b

Fig. 60 Temperature profile of the leading edge during arc-jet test: (a) for leading edge prepared bydensification method and (b) for leading edge prepared by compaction method [192]


Refractory foams also find application in transpiration cooling in hypersonicvehicles [351]. Ultramet’s transpiration-cooled combustion chambers with ceramicfoam (or metallic foam) increase performance and reduce cost. By allowing a smallportion of the fuel to flow directly through the wall, wall temperatures are kept low,and oxidation of the chamber wall is virtually eliminated. A foam core transpiration-cooled engine offers all the advantages of a foam core regeneratively cooled enginewith certain additional benefits. Increased coolant flow can be easily directed at areassubjected to the greatest heat flux. The inner wall of the entire chamber is maintainedunder nonoxidizing conditions, thus eliminating concern about oxidation. The needfor fuel-film cooling is eliminated, which simplifies the injector design. Becauseonly 2–5% of the total fuel flow is needed for transpiration and since no fuel-filmcooling is needed, most of the fuel can be used for combustion and burned effi-ciently. The lower coolant volume reduces cost and weight associated with turbo-machinery. In transpiration cooling, open-cell structural foam is combined with afine porosity liner at the inner surface with tailorable permeability. Only 2% of thetotal fuel flow is required for cooling.

SiC foam has also been used for preparing lightweight space mirrors [352]. In theseapplications closed cell SiC foam is coated with SiC and then coated with silicon. Theceramic closed-cell foam possesses as little as 5% of the density of its solid counter-part, which gives it attractive thermal and inertial properties. The foam-core optic’slow mass can dramatically improve closed loop performance for active beam steering.The advantage of the foam core is that it is physically incapable of producing print-through which is a concern in geometrical light weighting of SiC mirrors.

Ceramic HoneycombCeramic honeycomb structures are conventionally produced by the process ofextrusion. The material to be extruded is usually a mixture of ceramic powderswhich has been plasticized to give favorable extrusion performance [353]. Ceramic

Fig. 61 (a) Backwall temperature of SiC foam tile subjected to KHS, (b) SiC foam tile after twoKHS tests [192]


honeycomb was also prepared by embedding ceramic fiber tape in a ceramic matrix[354, 355]. A fabric tape or sheet formed of fibers of high-temperature material isimpregnated with the liquid ceramic precursor, which contains either a preceramicpolymer dissolved in a liquid vehicle, a colloidal suspension of a preceramicmaterial, or any other suitable liquid form of a precursor. The impregnated tape orsheet is then heated to evaporate the vehicle, the tapes or sheets are formed into thecorrugated shape in accordance with the desired dimensions of a honeycombstructure and cured to retain the corrugated shape. A high temperature ceramic-based adhesive or preceramic adhesive/binder is applied to the nodes of the corru-gations where the strips will be joined to form the honeycomb, and the sheets arestacked in the honeycomb arrangement. The adhesive is then cured at 100–300 �C,followed by pyrolysis at 1100–1200 �C. The resulting honeycomb is then dippedagain in the same slurry used initially, and pyrolyzed is once again.

In another approach, aluminum and paper honeycomb were dipped in a precursorslurry comprising of phenolic resin, SiC powder, and elemental silicon powder andcured at 150–200 �C, pyrolyzed at 900 �C, and then heat treated at 1500 �C to obtainSiC honeycomb [62].

Ceramic Sandwich StructuresThe concept of the construction of sandwich structures using polymeric foam andhoneycomb can be extended to the construction of ceramic sandwich structures forhigh temperature applications. Light weight ceramics such as ceramic foam andceramic honeycomb can be made use of as core materials since they have significantweight advantage (low density) and high thermal insulation characteristics [356,357]. In addition to the core materials, the fabrication of ceramic sandwich structuresrequires high temperature resistant face sheets, ceramic adhesives, and/or ceramicfasteners. A ceramic sandwich structure constructed using C/SiC composite as facesheets and SiC foam as the core material was evaluated as an internal multiscreenthermal insulation in place of silica tile in reentry vehicles [358].

Sreejith [62] evaluated three ceramic sandwich structures with C/SiC compositeas face sheets and SiC foam (FCS-01), SiC honeycomb prepared from paperhoneycomb (HCS-01) and SiC honeycomb prepared from aluminum honeycomb(HCS-02) as core material as thermal protection system for the windward region ofreusable launch vehicle. For the construction of the sandwich panels, phenolic resin-based ceramic adhesive was used for bonding C/SiC face sheets with the core. Thesesandwich structures having thickness of 20–30 mm and density of 0.48–0.52 g/ccwere subjected to KHS test, and the results are summarized in Table 7. All the threesandwich structures withstood the test maintaining the structural integrity. From theresults presented in Table 7, it is evident that FCS-01 with the lowest thickness of20 mm shows the least backwall temperature of 412 �C, though it has been subjectedto higher heat flux compared to the other two structures. Compared to honeycombsandwiches (HCS-01 and HCS-02), in the foam sandwich (FCS-01), the openchannel across the core is minimized to a greater extent or totally avoided.

Ajith [192] prepared two sandwich structures of dimension 150 mm� 150 mm�25 mm with C/SiC face sheets using SiC foam prepared from phenolic resin-based


precursor slurry and SiC Durofoam® (Fig. 62). These sandwich structures weresubjected to KHS test with the maximum heat flux of 4.1 W/cm2. Backwalltemperature of 139 �C and 173 �C was observed for IMI-1 and IMI-2, respectively.As the backwall temperature was lower for IMI-2, it was subjected to second testwith maximum heat flux of 24 W/cm2 and the backwall temperature observed was412 �C. It is worth noting that IMI-1 experienced a higher backwall temperature of601 �C even when tested at a maximum heat flux of 17.5 W/cm2. Less temperatureexperienced at the back-wall of IMI-2 may be due to the less contact area as the poresize of SiC foam in IMI-2 is more than that of Durofoam®, and hence, heat transferby conduction from the facesheet to the foam will be less than that of Durofoam®. Inaddition, the discontinuity of the damaged struts of SiC foam of IMI-2, as evidencedby SEM, also probably contributes in reducing the heat conduction through the core.

Ajith [192] investigated ceramic sandwich-ablative hybrid system for high heatflux conditions. A hybrid TPS (Fig. 63) (Density: 0.54 g/cm3) was fabricated bybonding a high density ablative (HDA) composite onto ceramic sandwich structurewith C/SiC face sheets and Durofoam® core which was subjected to sandwich of size150 mm � 150 mm � 25 mm using the high temperature ceramic adhesive.

The HDA composite of 5 mm thickness (Density: 1.38 g/cm3) was made using90 wt% chopped silica fiber and 10 wt% phenolic resin. The total thickness of thehybrid system was 30 mm with an overall density of 0.54 g/cm3.

Fig. 62 Sandwich structureswith SiC foam core and C/SiCface sheets: (a) IMI-1 withDurofoam

®

and (b) IMI-IIwith SiC foam prepared bytemplate method usingphenolic slurry

Table 7 KHS test results of ceramic sandwich structures [62]

SampleHeight(mm) Emissivity

Density(g/cc)

Maximumfront wall heatflux, Hmax

(W/cm2)

Maximumbackwalltemperature,Tmax (�C)

Time lagbetweenHmax andTmax (s)

HCS-01 25 0.87 0.52 16 495 100

HCS-02 30 0.82 0.48 20 420 125

FCS-01 20 0.89 0.50 24 412 180


Maximum backwall temperature of 320 �C was observed for the hybrid systemwhen subjected to KHS test with maximum heat flux of 90 W/cm2. It is worth notingthat the sandwich structure without ablative has exhibited a back wall temperature of601 �C at the maximum heat flux of 17.5 W/cm2 as discussed earlier. Thus, it isevident that the ablative layer takes away part of the heat load resulting in thereduction of the backwall temperature. The hybrid system withstood the test withoutdelamination of the sandwich structure. However, the ablative composite underwentcharring resulting in the development of surface cracks and loosening of silica fibers(Fig. 64).

Ortona et al. [170] described an integrated assembly method for preparingceramic sandwich structures with CMC face sheets and structural grade ceramicfoam, having density of 0.36–0.38 g/cm3 and porosity of 86–87%. SiCf preformswere obtained by laying up fabric plies, pre-impregnated with a small amount ofphenolic resin, cured in an autoclave, and pyrolyzed in argon atmosphere. Ceramicfoam was bonded to pyrolyzed preform face sheets using SiC-filled polysilazaneadhesive and cured in an autoclave. After pyrolysis, assembled sandwiches were

Fig. 63 Ceramic-Ablativehybrid TPS

Fig. 64 Ceramic-ablative hybrid TPS after KHS test (two views)


then impregnated on both sides of the sandwich skins with polysilazane using abrush. PIP was repeated five times to get the required strength for the CMC facesheets. Joining skins to ceramic core before their actual densification allows a stronggrip of the bonding layer both to the CMC laminates and to the foam. Both CMC andfoam are embedded into the polymer-derived ceramic layer and, thus, mechanicallyfastened to it.

Wang et al. [359] described a procedure for the fabrication of ceramic sandwichstructures in which the SiC core foam was modified with SiO2 to reduce the thermalconductivity. The advantage of modifying SiC foam core with SiO2 is explained in theheat transfer model (Fig. 65). From Fig. 65a it is obvious that the radiation can passthrough the large pores between SiC skeletons easily and directly. On the other hand, itis noticed from Fig. 68b that the radiation is blocked by SiO2 powder. SiO2 particleshave high extinction coefficient and can absorb the radiation energy significantly afterradiation through multiple reflection. Thus, most of the heat transferred through solidSiO2 powders is by heat conduction. As the SiO2 particles are not completely fused,solid heat conduction efficiency is much lower than heat radiation, and therefore thecompound core exhibits a better performance in the sandwich structure.

Ceramic foam core sandwich structures are also considered for application asspace mirrors [360]. Foam-core sandwich panels offer exceptional stiffness-to-weight ratios, and Ultramet has used this architecture for advanced mirror systems.Foam-core mirrors employ an optical faceplate that is attached to the surface of afoam billet, and, if stiffness requirements demand, a second face plate can be addedto the rear surface.

CMC ShinglesPichon et al. [162] proposed a new concept for the construction of lightweight CMCstructures, viz., CMC shingles as advanced thermal protection system for thewindward region of reusable launch vehicles. This concept is explained in Fig. 66.The shingle concept is divided into two sets of elements: (i) the ones with mechan-ical functions (mechanical shell, fasteners, and stand-offs) and (ii) the ones withthermal functions (inner insulation layers, seals, and insulating washers). The mate-rial needed for the mechanical shell has to be mechanically very efficient and

Fig. 65 The heat transfer models for SiC foam core and compound core. (Reproduced from Ref.[359] with permission from Elsevier)


resistant to highly constrained thermal environments, but its thermal conductivitycharacteristics are not important. The internal insulation and seals that do not requirehigh mechanical properties can be composed of low weight, flexible, and higherperformance insulating materials. The attachment system of the panel to the airframestructure must resist relatively high temperatures to enable the thermal expansion ofthe panel and transmit out-of-plane mechanical loads between the panel and the coldstructure.

Pichon et al. [162] constructed the shingles using carbon fiber reinforced preform.This preform was based on one ply multilayer woven fabric which was assembled byweaving to result in a self-stiffened panel (Fig. 67). The molding of these panels wasperformed on a carbon/epoxy composite material mold which was capable ofwithstanding at least the temperature needed for hardening the process. Then thepreform was hardened by a liquid route. The SiC matrix was added by CVIdensification on a simple graphite frame holding the part. After grinding, deburring,and drilling of the holes, the SiC densification process by CVI was performed once

Fig. 66 The shingle concept.(Reproduced from Ref. [162]with permission fromElsevier)

Fig. 67 Shinglesmanufactured from C/SiCcomposite. (Reproduced fromRef. [162] with permissionfrom Elsevier)


again and oxidation protection coating was added, if necessary. The shingle wassubjected to mechanical testing, dynamic vibrational testing, acoustic testing, thermal,and thermomechanical tests. The singles withstood the tests without noticeable damage.

In a variation of the above singles approach, Ajith [192] prepared C/SiC paneledlightweight ceramic structures by using prefabricated C/SiC composite panels,ceramic adhesive, and zirconia felt (Fig. 68). This shingle specimen was subjectedto two KHS tests with maximum heat flux of 4 W/cm2 and 17.5 W/cm2 and backwalltemperatures of 150 �C and 445 �C. It is noticed that the backwall temperature isslightly higher than that of IMI-2 prepared using SiC foam from phenolic resinslurry. This may be due to following reasons: (i) higher insulating characteristics ofSiC foam than that of zirconia felt and (ii) C/SiC partition walls may conduct heatfrom the front wall to the back-wall.

Concluding Remarks and Perspectives

The field of polymer-derived ceramics has witnessed a tremendous growth over thelast 50 years resulting in the creation of vast knowledge base with voluminous datawhich provide valuable inputs for taking forward this fascinating area of interdisci-plinary research to the next level. In the initial years, the focus was to synthesizesilicon-based processable precursors for non-oxide ceramic fibers for high

Fig. 68 Construction of C/SiC shingles (150 mm � 150 mm � 25 mm) [192]


temperature applications. Gradually, the emphasis got shifted toward synthesizingspeciality monomers or synthesizing newer precursors, understanding the transfor-mation of polymeric phase to ceramic phase and the effect of heat treatmentconditions on the microstructure and phase stability of the ceramic formed. In anattempt to increase the phase stability and high temperature capability of ceramics,precursors were designed to give multicomponent ceramics/mixed non-oxideceramic systems. Over the years, considerable progress is made in this direction,and a large number of UHTC precursors have been synthesized with the mainpurpose of obtaining ceramic systems which can go well beyond the capability ofSiC in terms of thermooxidative stability and retention of properties at temperaturesabove 1600 �C.

Though the initial efforts in the area of PDCs were focused on synthesizingpreceramic polymers devoid of oxygen, in the last three decades, polysiloxaneshave gained importance for obtaining siliconoxycarbide (SiCO) ceramics. Therelative ease of synthesis of such polymers coupled with the fact that SiCO ceramicscan be obtained only through PDC route has motivated researchers to explore therealm of polysiloxanes and related materials. The surge in research activities in thisarea led to exploring other polymeric systems such as polysilsesquioxanes, poly-borosiloxanes, polymetallosiloxanes, and polymetalloborosiloxanes as potential pre-cursors for SiOC, SiBOC, and Si(M)BOC. A great deal of work on polysiloxanesand related materials is concerned with the phase stability of SiOC glasses and therole played by boron, metallic constituent, and nanodomains of carbon present inthese systems.

With the advancements in the field of nanotechnology, PDCs have attained muchmore significance in terms of both scientific understanding and their utilization invarious fields. Though the initial interest on PDCs was to obtain high temperaturematerials such as ceramic fibers, coatings, and matrix for CMCs, in recent years,PDCs have been explored for various applications which include biomedical appli-cations, drug delivery systems, water remediation, energy storage devices, micro-electronics, and nanosensors.

The main objective of this chapter is to focus on preceramic polymers which findapplications in space research. While discussing this aspect, ceramics obtained fromcarbonaceous polymers are also included in the discussion, as they offer cost-effective approach for obtaining high temperature materials. In spite of concertedefforts for developing high temperature materials such as CMCs and UHTCs, C/Ccomposites with oxidation protective coatings, obtained through conventional pro-cess, are preferred for most of the high temperature materials requirement for spaceapplications. This is mainly due to the pedigree of C/C composites and the avail-ability of carbon fibers of different grades to meet the specific requirements. Despitethe fact that ceramic matrix composites, ceramic coatings, and lightweight ceramicsfrom PDCs have caught the attention of researchers in the area of materials scienceand technology, the acceptance of PDC-based components for space applications asreplacement for oxidation protected C/C composites and silica tiles, a lightweightceramic material, is still evasive though they have been flown as candidate materialsin experimental space vehicles. The scenario is expected to change in the coming


years with the advancement in manufacturing processes such as additivemanufacturing/3-D printing, hybrid CVI+PIP process, and joining techniques.

In order to penetrate into the domain of high temperature materials, the cost ofprecursors for PDCs should be brought down drastically. Yet another major concernis the porosity of PDC-based components particularly for applications as exit conesand satellite thrusters. Though the porosity aspect is partly addressed throughAFCOP approach, this method could not be effectively utilized as the active fillerstend to react with reinforcement/interphase coating adversely affecting the mechan-ical properties of CMCs. To take advantage of AFCOP approach, it is imperative todevelop suitable protective coatings onto carbon fiber so that the number of infiltra-tion steps can be reduced. Another concern that requires attention is eliminating/reducing the deterioration of reinforcement due to reaction with pyrolysis products.This problem can be minimized by using flash pyrolysis technique. However, thistechnique is to be adopted from making miniature components to large-sizedcomponents required for space applications.

In the domain of PDC coatings, there is considerable scope particularly forproviding interphase coating onto fiber reinforcement and this approach can reducethe cost of production of CMCs by eliminating the need for CVI/CVD furnace forinterphase coating. With the advancements on UHTC precursors, new functionallygraded coatings for protection of composite structures and metal substrates can beenvisaged. As far as the thermal barrier coatings are concerned, there is a need todevelop precursors for top coat capable of withstanding higher temperatures. Thereis considerable scope for fine tuning the microstructure of the coating to increase thetemperature gradient between the front wall and back wall. Yet another area ofinterest is use of preceramic polymers-based atomic oxygen resistant coatings andcomposites for the construction of spacecrafts and low earth orbit space structures/space stations.

In the area of ceramic adhesives, researchers should take advantage of theavailability of new precursors for UHTCs. Suitable adhesive formulations basedon UHTC precursors are to be developed for joining of CMCs and ultra-hightemperature CMCs with super alloys and other substrates. As it becomes increas-ingly difficult to prepare large components, joining of ceramic components becomesinevitable. Reliability of such joints is to be addressed through health monitoringprocedures involving nondestructive testing techniques.

The advancements in the field of additive manufacturing are expected to boost thedevelopment and production of lightweight ceramics as advanced thermal protectionsystems. Production of near net shaped lightweight ceramic materials with improvedhigh temperature capability are expected to replace silica tile in the near future.Scope exists on realizing hybrid advanced thermal protection systems combiningablative systems with CMCs resulting in mass saving with improved reliability forfuture manned missions and interplanetary missions.

In conclusion, interdisciplinary approach with active participation of researchersin the field of chemistry, materials science, and engineering is the way forward torealize the vision of using PDC-based materials for space applications and outerspace exploration.


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