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Surface & Coatings Technology
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Structural changes of CVD Si–B–C coatings under thermal annealing
Camille Pallier a, Georges Chollon a,⁎, Patrick Weisbecker a, Francis Teyssandier a,Christel Gervais b, Fausto Sirotti c
a Laboratoire des Composites Thermostructuraux (CNRS, Safran, CEA, UB1), Université de Bordeaux I, 3, allée de La Boétie, 33600 PESSAC, Franceb Laboratoire de Chimie de la Matière Condensée de Paris, UPMC Univ Paris 06, UMR 7574, Collège de France, 11, place Marcelin Berthelot, 75231 PARIS Cedex 05, Francec Synchrotron SOLEIL, L'Orme des Merisiers, Saint-Aubin, BP 48, 91192 GIF-sur-YVETTE CEDEX, France
⁎ Corresponding author. Tel.: +33 5 56 84 47 27; faxE-mail address: [email protected] (G. Cho
Si–B–C coatings were prepared by chemical vapor deposition (CVD) from CH3SiCl3/BCl3/H2 gas mixtures at900–1000 °C and under low total pressure (5 kPa). A selection of coatings was synthesized at different depo-sition temperatures and initial compositions of the gas phase. They were characterized in terms of morphol-ogy, elemental composition and structure, both at the atomic scale (by nuclear magnetic resonance and X-rayabsorption) and long range (by X-ray diffraction and transmission electron microscopy). The as-depositedmaterials consist of an amorphous boron carbide phase (a-BxC) including icosahedron-like units andenclosing SiC nanocrystals. The SiC grain size increases with the silicon concentration in the solid, i.e. withthe initial CH3SiCl3 concentration in the gas phase. The amorphous structure and the excess carbon are partlyaccommodated by unusual BC3 environments in boron carbides. Thermal annealing in inert atmosphere at in-creasing temperature and time gradually modifies their structure at short and long range. A disordered poly-cyclic carbon phase appears first, while the SiC grain size increases progressively, and rhombohedral boroncarbide (B4C) finally crystallizes around 1300 °C. The organization of the SiC4 sites is improved and the BC3sites are partially transformed into more usual intra and intericosahedral environments. Whereas the tem-perature has a strong influence on the structural changes, the annealing time has only a limited effect, thestructure being rapidly “frozen” in a metastable state.
Multilayered (Si)–(B)–Cmatrices (with alternating layers of SiC, B–Cand Si–B–C ceramics) have been developed to improve the durability ofceramic matrix composites (CMCs) for future aeronautic engines. Boroncontaining Si–B–C (respectively B–C) ceramics naturally generate aSiO2–B2O3 (resp. pure B2O3) protectivefilm in air, which heals thematrixcracks and protects the fibers against oxidation [1]. Si–B–C ceramics arecurrently deposited by chemical vapor infiltration (CVI) from a BCl3–MTS–H2 gas mixture (MTS stands for methyltrichlorosilane: CH3SiCl3).The fiber cloths are placed in a hot-wall reactor operating at high tem-perature (T>850 °C) and reduced total pressure (pb10 kPa). The CVDparameters (T, p and the various gas flow rates) are adjusted to obtaina homogeneous Si–B–C matrix infiltration of desired composition.
Only a few studies deal with the chemical vapor deposition (CVD) orCVI of Si–B–C ceramics [2–4]. Goujard et al. [4] originally developed Si–B–C outer coatings, using the BCl3–MTS–H2 system, in order to protectCMCs against oxidation. Uniform layers of various B and Si-contents
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have been deposited at relatively high growth rates. The infiltration ofthick fibrous structure specifically requires an efficient mass transfer ofthe reactants through the porous substrate and accordingly sufficientlylow chemical reaction rates [5]. This is why CVD conditions such asthose investigated by Goujard et al. [4] or Golda and Gallois [3] are un-suitable for CVI. Later on, a few adjustments of the process lead to the de-velopment of Si–B–C multilayer matrices, by CVI, using the same gassystem [6].
The so called “self-healing” composites were submitted to high tem-peraturemechanical testing in inert and oxidizing atmosphere, and theirmicrostructures were examined before and after testing [7–12]. Little in-formationwas revealed about the initial composition and structure of Si–B–C and B–C layers. Except at the interfaces, the Si–B–C layers werefound only partly crystallized [8] and composed of SiC nanocrystals anda disordered boron carbide phase [11–13]. The creep deformation ofthe composites was attributed to a complex combination of fiber creep,matrix microcracking, interface debonding and also, in air, to the viscousoxide flow [8,9,11]. A crystallization of the Si–B–C layers into SiC and B4Cwas reported after creep under argon or air at 1200–1300 °C, but thisphenomenonwas not supposed to have amajor influence on the overallcreep deformation [8,11]. However, an atypical primary creep stage wasobserved during the first hours at 1100 °C/200 MPa and 1200 °C/
150 MPa [12]. This transient behavior was related to the crystallizationand densification of the Si–B–C and B–C layers, possibly promoted bystress.
Very few studies aimed at the analysis of individual Si–B–C layers.Michaux et al. prepared 1D model composites in order to investigatethe high temperature tensile properties of SiC, B–C and Si–B–C coatings[14]. A time-dependant non-linear behavior was observed at 1200 °Cfor both B–C and Si–B–C systems, which was related to the amorphousnature of the boron-containing phase.
More than 10 years after Goujard et al., Berjonneau et al. dedicatedtheir work to experimental and fundamental aspects (homogeneous re-actions, growth kinetics and mechanism) of the deposition of Si–B–Ccoatings from BCl3–MTS–H2. This time, the CVD experiments were runat reduced pressure (2≤p≤12 kPa) and temperature (850≤T≤1050 °C), i.e. in conditions compatible with CVI [15]. The materialsobtained are disordered except at high T, p and MTS concentration, forwhich SiC nanocrystals appear embedded in an amorphous B andC-rich phase. Although the compositions were all found within theSiC+B4C+C domain of the Si–B–C phase diagram, graphitic carbonand crystalline B4C are both absent in the as-deposited Si–B–C ceramics.
These studies clearly demonstrate the metastable character of thesehighly disordered Si–B–C ceramics and suggest that it would be relatedto the transient viscous behavior observed at high temperature. A com-parable case has already been reported for ceramics belonging to thequaternary Si–B–C–N system [16–20]. As-sputter deposited Si–B–C–Nfilms are usually amorphous. Some new phases appear and crystallizeafter thermal annealing in inert atmosphere, depending on the compo-sition of the coating [16,17]. In comparison, many studies have beendedicated to the thermal stability of polymer-derived bulk Si–B–C–Nceramics [18–20]. Except polycyclic carbon domains, these materialsare amorphous and their structure remains unchanged up to at least1400 °C. The crystallization of the SiC and Si3N4 phaseswas investigatedas a function of the annealing conditions in Ar or N2 [18,19] and thesestructural changes were found to be highly related to the shrinkageand the decrease of the creep rate observed at high temperature [19,20].
A similar phenomenon in the Si–B–C layers of a multilayer matrixcould be detrimental to the performances of CMCs in use, especially athigh temperature, for long durations and under mechanical loading.Such conditions could be met, for instance, in the hot parts of aero-nautic engines (e.g. primary flaps).
To better understand and predict the behavior of Si–B–C ceramicsat high temperature, Si–B–C coatings of various compositions havebeen deposited from specific CVD conditions. Their structure hasbeen studied at short and long range (i.e. at the atomic and a fewunit cell scale, respectively), in the as-deposited state and after ther-mal annealing. A kinetic study of the formation of the various phaseshas been carried out. This work is a contribution to the building of astructural model for the pristine and annealed material and of themechanisms responsible for the changes of the structure and the me-chanical properties at high temperature.
2. Experimental
Si–B–C films were deposited from a BCl3–MTS–H2 gas mixture at900–1000 °C and a total pressure of 5 kPa. Experimental details canbe found in [15], though the geometry of the reaction chamber was dif-ferent in this case: 100 mm inner diameter and an isothermal zone(+/−5 °C) length of approximately 100 mm. In the BCl3–MTS–H2 sys-tem, the initial composition of the gas mixture is usually defined by thetwo parameters α=xH2/xMTS and β=xMTS/xBCl3 (where xi are the ini-tial mole numbers) [4,15]. The total flow rate of the gas mixture rangedfrom 400 to 500 standard cm3 per minute (sccm). The gas compositiondomain explored was in the range 8≤α≤25 and 0.2≤β≤3.3.
Different types of substrates were used depending on the character-ization technique and the investigated property. For surface analyses,the coatings were deposited on small plates of single crystal Si wafers
(500 μm-thick, 100-orientation and polished, fromNeyco) or high puri-ty polycrystalline β-SiC wafers synthesized by CVD (500 μm-thick,polished, Fe content: 35 ppb, fromTechnical Glass Company). For a spe-cific purpose (see paragraph below) the coatings were also depositedon rayon-based carbon fiber tows (non-commercial samples, fromSnecma Propulsion Solide). For bulk analyses requiring a few tenths ofmgof pure powder, small blocks of open cell reticulated vitreous carbonfoams (experimental samples of a 250 μm cell size, from Commissariatà l'Energie Atomique) were used as substrates. The infiltrated foamswere ground and oxidized in dry air at 430 °C until the carbon substratewas completely removed. The powder sample was subsequentlycleaned in hot water to remove the superficial oxide layer.
Standard heat treatments were carried out with a radio frequency-heating furnace, under a high purity Ar flow (105 Pa), at 1300 °C during2 h. A high vacuum (Pb10−3 Pa) resistive-heating devicewas designedand used specifically for the kinetic study, to allow fast heating(20 °C s−1) and cooling (30 °C s−1) rates of small and elongated spec-imens. The specimens in this case were short lengths (5 cm) of carbonfiber tows coated with Si–B–C layers. The temperature was controlledwith a bichromatic pyrometer in both cases.
Themorphology of the coatings was characterized by scanning elec-tron microscopy (SEM, FEI, Quanta-400 FEG). The bulk chemical com-position (through thickness) of the coatings was measured from theflat specimens either by electron probe microanalysis (EPMA, CamecaSX100, with an accelerating voltage of 7 kV, leading to a probed depthof about 0.5 μm), when the coatings were sufficiently thick, or byin-depth Auger electron spectroscopy profiling (AES, VG Microlab310-F, accelerating voltage: 10 kV and beam current: 5 nA, the compo-sition being averaged from the surface and through a homogeneouslayer of few hundred microns), for thinner coatings. High purity SiC(of same quality as the β-SiC substrate) and hot-pressed B4C (purity:99.5%, from Goodfellow) were used as the Si, C and B standards.
The ordering at long range (at a distance of a fewunit cells)was stud-ied byX-ray diffraction (XRD, BrukerD8,λCu–Kα1=0.154056 nm), eitherin the Bragg–Brentano geometry θ–θ mode, for the powder specimens,or parallel beam/glancing angle, for coatings on C fiber tows. Transmis-sion electron microscopy (TEM, Philips, CM30ST, accelerating voltage300 kV) was used to identify the crystalline phases in the coatings at ahigh spatial resolution. The powder samples were observed withoutpreparation. The ordering at intermediate range (typically a fewinteratomic distances) was evaluated, through vibration properties, byRaman microspectroscopy (Horiba-Jobin Yvon, Labram HR, λHe–Ne=632.8 nm). The lateral resolution is approximately 1 μm and the thick-ness probed varies from a few tenths to several hundreds of nanometers,depending on the composition and the structure.
The Si and B atoms' local environmentwas investigated by solid-statenuclear magnetic resonance (NMR). The 11B MAS NMR spectra wererecorded at 11.75 T on a Bruker Avance 500wide-bore spectrometer op-erating at νL=128.28 MHz and using a Bruker 4 mm probe and a spin-ning frequency of the rotor of 14 kHz. The spectra were acquired using aspin-echo θ−τ−2θ pulse sequence, with θ=90°, to overcome problemsresulting from the probe signal. The τ delay was synchronized with thespinning frequency and a recycle delay of 1 s was used. The chemicalshifts were referenced to BF3(OEt)2 (δ=0 ppm). Single pulse 29Si MASexperiment was performed on a Bruker Avance 300 spectrometer(7.0 T) operating at 59.66 MHz using a 7 mm double resonance MASprobehead, a spinning frequency of the rotor of 5 kHz and a smallflip-angle (π/6). Chemical shifts were referenced to TMS.
The local structure of the B and C atoms was studied by X-rayphotoabsorption spectroscopy (or X-ray Absorption Near-Edge Struc-ture: XANES), near the B1s and C1s edges (at 188 eV and 284 eV re-spectively). The experiments were carried-out on the TEMPObeam-line at SOLEIL [21]. The spectra were recorded in ultra high vac-uum (Pb10−8 Pa) in the total electron yield mode. The XANES signalcorresponds to a depth of about 10 nm (i.e. the low-energy secondaryelectrons path). The experimental photon energies are not adjusted
Fig. 1. Bulk elemental composition (as measured by EPMA) of the A–D coatings, shownin the Si–B–C ternary diagram (at T=1400 K, from [23]).
180 C. Pallier et al. / Surface & Coatings Technology 215 (2013) 178–185
to standard values. The absolute error is about 0.1–0.2 eV but therelative precision is less than a few meV. The surface compositionand chemical bonding of the specimens were also investigated byX-ray photoelectron spectroscopy (XPS, extremeScienta SES 2002,excitation: 830.2 eV, electron escape: approx. 0.5–1.5 nm), on TEMPO[22], just before the XANES analyses.
3. Results and discussion
3.1. As-deposited materials
The bulk composition (as measured by EPMA) and the depositionrate of a selection of four coatings A–D (Table 1) is presented in theSi–B–C ternary diagram [23] in Fig. 1 (the oxygen concentration re-mains below the detection limit of 1–2 at.%). They are all located closeto the quasi-binary B4C–SiC line. The composition of specimen B isclearly within the three-phase domain “B4C+SiC+C” whereas that ofcoating D apparently belongs to the “BxCy (x/y>4)+SiC+Si” domain.Homogeneous chemical reactions play a significant role in hot wallthermal CVD. In the BCl3–MTS–H2 system, themain effective precursorsof B, C and Si are respectively HBCl2 (produced from the reaction be-tween BCl3 and H2), the CH3 radical and chlorosilanes. The latter twospecies, resulting from the MTS thermolysis, react independently togrow the coating. Within the 900–1000 °C domain, the heterogeneousreactions between CH3 and the chlorosilanes start to prevail over the re-action between CH3 and HBCl2 [24]. The deposition rate increases withthe temperature, while the silicon and carbon concentrations increaseand the boron content decreases [15,24]. A similar influence of the tem-perature can be evidenced by the present results (compare D with C inTable 1). In addition, the increase of β (xMTS/xBCl3) tends to lead to a de-crease of the boron concentration and an increase of the silicon and car-bon amounts [4,15]. The same effect holds true in the present case(compare Dwith B), in spite of the particularly high β value of 3.3. How-ever, the silicon concentration is much more affected by the T or βincrease than the carbon concentration (which is only slightly im-proved). This is probably due to the higher reactivity of chlorosilanescompared to hydrocarbons and illustrates the complexity of such achemical system.
The structural analyses discussed below focus on A, B and D coatings.The structural features of specimen C, as assessed by XRD and Ramanspectroscopy (not shown), were indeed found very similar to those of B.Except specimenD, the as-deposited ceramics consist essentially of disor-dered phases. The XRD patterns, in Fig. 2, only show extremely broaddiffraction features which can be attributed to nanocrystalline orsubnanocrystalline β-SiC. The analysis of the 111-plane diffraction peakusing the fundamental parameters approach led to apparent crystallitesizes less than 1 nm for coating A and about 1–2 nm for B and C (the for-mer value is of course highly imprecise due to such broad diffraction fea-tures). Specimen D is clearly distinct from the other coatings. The XRDpattern shows well-defined diffraction peaks attributed to the β-SiCphase and corresponding to an apparent crystallite size of 15 nm. TheRaman spectra of coatings A, B and D, shown in Fig. 3, do not reveal anyobvious SiC features, the SiC grains being probably too small and scarce,and the SiC Raman cross-section being particularly low [25]. There is noindication either of any crystallized boron carbide phase or graphitic car-bon, as suggested by the phase diagram. Only very broad features are
Table 1CVD conditions, bulk elemental composition (as measured by EPMA) and depositionrate of the Si–B–C coatings (P=5kPa).
visible around 400–700 cm−1 (the steep left side being due to theNotch filter cutoff) and 850–1350 cm−1, which are typical of anamorphous state. By comparison with the B4C reference spectrum, thewide 1050 cm−1 band could be attributed to breathing modes oficosahedron-like units [26,13,15]. The Raman spectra of coating Dshows several broad features at wavenumbers corresponding to boroncarbide (400–600, 700–800 and 1000–1150 cm−1) and SiC (600–1000 cm−1) [26,27,15]. Few other sharp bands also appear at 670 and730 cm−1. They have not been assigned yet, but could result from thepresence of B-rich crystalline boron carbide(s) (BxCy, with x/y>4).
Among the specimens described above, only two coatings and aB4C reference (B4C powder, 200 mesh, purity: 98%, from Sigma-Aldrich), were selected for the solid-state NMR analyses. The firstone A, is a boron-rich material whereas the second one B, has highersilicon and carbon-concentrations. The 11B MAS-NMR spectra of thetwo as-deposited Si–B–C coatings (Fig. 4) reveal two main boron en-vironments. The more intense feature, centered at about 0 ppm, issimilar to the peak observed for icosahedral sites in crystalline
Fig. 2. XRD patterns of the as-deposited and heat treated (HT) Si–B–C coatings(1300 °C/2 h/Ar).
boron carbide [28,29]. However, it is slightly broadened andup-shifted compared to the reference B4C spectrum (~0 ppm vs. −5 ppm), suggesting a less organized structure [29]. This result is inagreement with the Raman analyses: the boron carbide phase proba-bly includes highly disordered icosahedron-like units, with similarB-environments. An additional peak, however, not observed for B4C,appears at 55 ppm for the two Si–B–C specimens. Its intensity is sig-nificantly higher for specimen B (Fig. 4). A similar peak was reportedfor polymer-derived Si–B–C–N ceramics and attributed to trigonalBC3 environments among various substitutional BCxN3–x sites[30,31]. These trigonal sites are normally absent in the crystallineB4C. Only bicoordinated (in the linear chain) and hexacoordinated(in the icosahedrons) environments are present in the rhombohedralstructure. These particular intericosahedral sites, in place of the BC2 en-vironments in B4C (in the C–B–C linear chains), could be related to theexcess carbon. The higher coordination number of the B atoms wouldindeed explain the absence of a free carbon phase in the ceramic. The in-crease of the 55 ppm peak intensity with the Si-content suggests thatthe amount of these particular B-environments could also be related, in-directly (since B–Si bonds were not evidenced by any mean), to thepresence of Si atoms in the ceramic. The proportion of these BC3 sites
Fig. 4. 11B MAS NMR spectra of the as-deposited and heat treated (HT) Si–B–C coatings(1300 °C/2 h/Ar) (* indicates spinning sidebands).
could be for instance associated with the interface between the SiCnanocrystals and the amorphous boron carbide phase (a-BxC).
For both materials, the 29Si MAS NMR spectra (Fig. 5) show a verybroad band centered at about 5 ppm and corresponding to a SiC4 en-vironment in a highly disordered material [32]. The rather lowamount of silicon in the A specimen is responsible for the poor sig-nal/noise ratio of the spectrum. The very broad lineshape observedfor the A sample signal suggests a large distribution of chemical shiftsrelated to a distribution of bond lengths and angles around siliconatoms. These highly distorted environments probably arise from theouter Si atoms of the SiC clusters (at the SiC/a-BxC interface). Theircontribution to the NMR spectra is naturally more significant for thesmallest SiC clusters or even in case of Si atoms isolated in theboron carbide phase. The spectrum of the B specimen reveals atleast two distinct SiC4 components: the first, narrower, around −13 ppm, and the second, broader, at about 5 ppm. These two environ-ments could be related to the different contributions of the siliconatoms located, respectively, inside the SiC nanocrystals, or at theirboundary, i.e., at the interface with the amorphous boron-rich phase.
The B1s photoabsorption spectra of three as-deposited Si–B–C coat-ings and the B4C reference (the same standard material as that used forthe bulk elemental analyses) are shown in Fig. 6. As for B4C, the Si–B–Cspectra show no clear distinction between the near edge region(1s→π⁎ transition) and the high energy domain (σ⁎ states), indicatinga complex state of hybridization [33–35]. The energy edge is lower forthe Si–B–C spectra (188.0 eV) than that for B4C (189.1 eV). Whereasthe B4C spectrum shows a main intense peak at 191.4 eV, four well dis-tinct components are present in the Si–B–C spectra at 189.9, 191.4,192.5 and 194.1 eV. These spectra are very similar to those of a C-rich“BxC1−x” coating deposited by coevaporation of boron and carbonatoms [35]. Furthermore, they share some common features with thespectra of high purity B4C annealed at various temperatures [34].Jimenez [34] identified four consecutive components labeled as A–D, re-spectively located at 190.9, 191.7, 192.3 and 193.7 eV. The most intenseA peak was attributed to π⁎ states presumably arising from boron atomsin a linear C–B–C chain. The B peak was assigned to B-rich carbide andthe C and the D peaks respectively to native boron suboxide and B2O3.This analysis was acknowledged and completed by Caretti et al. [35],who described five components in the spectra of “BxC1−x” coatings ofdifferent stoichiometry, labeled as B0–B4, the B1–B4 peaks essentially cor-responding to the A–D environments of Jimenez et al. The original B0peak at 189.7 eV, which is not present in the B4C spectrum,was attribut-ed to a BC3 trigonal environment [36]. The presence of a very similarcomponent at 189.7 eV in the Si–B–C coating spectra supports the exis-tence of BC3 sites in the material, as already suggested by 11B
Fig. 6. B1s photoabsorption spectra of the as-deposited Si–B–C coatings. Fig. 7. C1s photoabsorption spectra of the as-deposited Si–B–C coatings.
182 C. Pallier et al. / Surface & Coatings Technology 215 (2013) 178–185
MAS-NMR analyses. These environments, however, are not the onlyboron environments in the material (contrary to hexagonal sp2 struc-tures such as h-BN or boron-substituted carbons). The similarity of thespectra with the reference B4C spectrum suggests that they are likelymixed with other sites of higher coordination number, probably locatedin icosahedron-like sites. The three other peaks, matching the B1, B3 andB4 components, can be tentatively assigned to various tricoordinatedsites, BCxO3−x (0≤x≤2), the oxygen atoms arising from superficialcontamination from the ambient air [35]. The increase of the “BC3”peak intensity with the silicon concentration in the coatings is worthyof note. Even if the volume (and the elemental composition) consideredis different in both cases, this result is in agreement with the MAS11B-NMR analyses and, particularly, to the intensity changes of the55 ppm peak.
The C1s photoabsorption spectra of two as-deposited coatings areshown in Fig. 7. As for the B1s spectra, the C1s spectra reveal a complexhybridization state compared with graphite (sp2) or diamond (sp3)[34]. All the Si–B–C spectra show a peak at 284.9 eV and an edge atabout 286.5 eV, which both resemble the B4C reference features.The former feature is also close to the π⁎ transition of graphite ob-served at a slightly higher energy (285.2 eV). The SiC reference spec-trum (recorded from a polycrystalline β-SiC wafer of the same qualityas the SiC substrates) displays a main peak at 285.0 eV, also very closeto the 284.9 eV peak in B4C. The distinction between the two refer-ence carbides is difficult due to such a similarity in the C1s features.The 284.9 eV peak is less intense and broader in the Si–B–C and SiCspectra than in the B4C spectrum. The high energy σ⁎ region (from287 to 300 eV) better differentiates the two coatings. The Si-rich Dspectrum indeed shares with the SiC reference a similar shape ofthe σ⁎ region.
3.2. Heat-treated materials
As in Section 3.1, the results presented below refer only to A, B andD specimens. The structure of the heat-treated C and D specimenswas indeed found very similar from the XRD and Raman spectroscopydata (not shown for C). The crystallization behavior of A, B and C willbe investigated in more detail in Section 3.3.
As expected from the Si–B–C phase equilibrium [23], the bulk ele-mental composition of the coatings was not significantly affected bythe 1300 °C–2 h treatment (not shown). On the other hand, the dif-ferent compositions of the coatings resulted in notably different
structural changes after the heat treatment. The XRD pattern(Fig. 2) and the Raman spectrum (Fig. 3) of the annealed sample Aboth reveal the presence of rhombohedral B4C (evidenced on theRaman spectrum by the two sharp peaks at 485 and 535 cm−1, dueto the linear C–B–C chain vibrations, and the icosohedral breathingmodes at about 1080 cm−1) [26]. A SiC crystal growth can also be ob-served on the XRD pattern, from less than 1 nm to about 3 nm(Fig. 2). Moreover, two bands appear after annealing on the Ramanspectra at 1600 and 1350 cm−1 (Fig. 3). These features are referredto as, respectively, the G peak, corresponding to the phonon of E2gsymmetry expected for graphite, and the D peak, assigned to structur-al defects in the hexagonal layers [37]. Both bands are broad, indicat-ing a disordered form of polycyclic (sp2) carbon [38].
After annealing, the XRD analysis of the B specimen reveals an in-crease of the SiC crystallite size from 1–2 nm to 5–6 nm. No otherphase is detected by XRD, as shown in Fig. 2. In this case, the Ramanspectrum does not show the typical B4C features but only the two in-tense D and G bands associated to the disordered free carbon phase(Fig. 3). The SiC features are still absent on the spectrum due to themuch higher Raman cross-section of carbon, compared to SiC. TheTEM analysis of the heat-treated B sample (Fig. 8) confirms the signif-icant growth of the β-SiC grains as well as the formation of polycycliccarbon, as a few stacks of highly curved layers, both phases being ho-mogeneously distributed in the material. Crystalline boron carbide isnot detected either by XRD, Raman spectroscopy or TEM.
Only limited changes are observed after annealing the D specimen.A slight β-SiC grain growth is observed by XRD (from 15 to 19 nm),while no other phases are revealed. This is confirmed by the Ramananalyses, which show almost no change in the spectrum, but slightlysharper features for the annealed specimen (Fig. 3).
The nature of the various phases in the annealed coatings andtheir organization at long and intermediate range (as evidenced byXRD and Raman spectroscopy, respectively), is in good agreementwith the expected ternary isothermal section of the Si–B–C phase di-agram (Fig. 1). Specimen A contains only a small excess of carbon andvery low amounts of SiC. Although it is probably not totally crystal-lized after annealing, rhombohedral B4C is the major crystallinephase in this coating. Specimen B is apparently more crystallizedand, as expected, includes more SiC and free carbon than A. CoatingD is atypical among the Si–B–C coatings studied so far [4,7–15], as itcontains a major part of SiC and, as expected, no excess carbon,even after annealing. The presence of free silicon is not observed
Fig. 8. SiC TEM dark field image (111-plane) (a), selected area electron diffraction pattern (b) and high resolution image (lattice fringes) (c) of heat treated B coating (1300 °C/2 h/Ar).
(contrary to what is suggested by the phase diagram) and the boroncarbide phase is too scarce for its structure to be clearly identified.
The 11B MAS NMR spectra of the A and B specimens (Fig. 4) show astrong decrease of the intensity of the 55 ppm band after annealing(it almost disappears in the A sample). The main band at ~0 ppm isshifted to −5 ppm for the two coatings, in accordance with the stan-dard rhombohedral B4C spectrum. The vanishing of the trigonal BC3sites, in favor of the more typical B-environments in boron carbides,is obviously related to the conversion of the a-BxC phase into crystal-line B4C. The NMR data confirm the XRD and Raman analyses, at leastfor specimen A. Yet, a few BC3 sites are still present in the B coatingafter annealing. As suggested above, they could be related to the larg-er amount of SiC/a-BxC interfaces in this material.
A very narrow peak is observed at about−16 ppm in the 29Si MASNMR spectra of the heat-treated A and B specimens (Fig. 5), corre-sponding to the β-SiC phase [33,39]. Such a strong decrease of thelinewidth and the slight downshift observed after annealing both re-sult from the SiC grain growth and the formation of a more homoge-neous SiC4 environment. The peak position and the shouldersobserved at low chemical shifts indicate the presence of smallamounts of other polytypes, likely as stacking faults [40].
The heat-treated coatings were also analyzed by photoabsorption atboth B1s and C1s edges. Both spectra reveal strong π⁎ and σ⁎ transitionstypical of a sp2 hybridization of the C and the B atoms (not shown). TheXPS analyses (not shown) revealed that the surface composition(trough about 1 nm) was strongly altered by the 1300 °C–2 h treat-ment. A few attempts were made to heat the coatings under ultrahigh vacuum, directly in the TEMPO line's preparation chamber. Thesurface changes were less significant in this case, but TEM analyses re-vealed, however, the presence of a 10 nm-thick anisotropic layer essen-tially made of carbon. This thin outer layer probably results from theremoval at high temperature of the boron and silicon oxide species, ini-tially present at the surface of the as-deposited specimens (due to con-tamination) and to the silicon sublimation promoted by high vacuum.The depth probed by XANES being of the order of 10 nm, thephotoabsorption spectra are of course greatly influenced by this decom-position layer. Obviously, these results cannot be related to the bulkanalyses (e.g. EPMA, XRD or NMR) in this particular case. Comparative-ly, Raman analyses, performed on the same annealed specimens, clearlyshow boron carbide features indicating a limited contribution of theC-rich surface layer due to a significantly larger probed depth. For this
reason, XRD and Raman spectroscopy were preferred to follow moreprecisely the structure changes, versus temperature and time.
3.3. Crystallization kinetics
The influence of the annealing temperature and time on the struc-tural changes was more precisely studied by performing heat treat-ments on the Si–B–C-coated carbon fiber bundles. The annealedspecimens were respectively examined by Raman spectroscopy tofollow the formation of free carbon and B4C (e.g., for specimen A,see Fig. 9), and XRD to evaluate the growth of the SiC crystals. The in-tensities of the D band at 1350 cm−1 (attributed to the free carbonphase) and the C–B–C chain characteristic peak at 485 cm−1 (dueto rhombohedral B4C) on the Raman spectra (Fig. 9) were plotted ver-sus the heat treatment temperature and time (Figs. 10 and 11). Thisinformation is only qualitative as the Raman intensities may also fluc-tuate with the surface state of the specimens. The apparent SiC crys-tallite size was evaluated from the 111-diffraction peak (Figs. 12and 13).
The D specimen was not further examined as the as-depositedmaterial is already crystalline and its structure is unaffected afterannealing at 1300 °C–2 h. The behavior of the C coating which con-tains higher Si and C amounts than A and B, while being also essen-tially amorphous, was also investigated in this section.
For a dwell time of 15 min, the free carbon phase appears first at1200 °C for the boron-rich (A) coating (Fig. 9), and 1100 °C for the sili-con and carbon-rich (C) coating (Fig. 10). The amount of free carbonkeeps on increasing with temperature (as shown by the intensitychanges of D andG bands) up to 1300 °C, for all the coatings. The reorga-nization of the a-BxC phase, confirmed by the sharpening of the broad1050 cm−1 band (Fig. 9), starts only at 1250–1300 °C. For theboron-rich coating A, the crystallization of rhombohedral B4C is clearlyevidenced beyond 1300 °C, on the XRD pattern (not shown) and theRaman spectra, from the typical vibrations of the C–B–C chains and ico-sahedrons, respectively at 485, 535 and 1080 cm−1 (Fig. 9). In parallelwith the formation of free carbon and B4C, a coarsening of the β-SiCnanocrystals is observed by XRD when T increases (XRD pattern notshown). The apparent crystallite size increases progressively up to1400 °C (for t=15 min) for coating C, whereas it rises rapidly beyond1250 °C for coating A and stabilizes at a similar level at 1400 °C (Fig. 12).
Fig. 9. Raman spectra of A coating as a function of the annealing temperature (t=15 min/high vacuum).
Fig. 11. Intensity of theD band (free carbon) and theC–B–C chain characteristic peak (B4C),as derived from the Raman spectra, as a function of the annealing time (T=1300 °C/highvacuum).
184 C. Pallier et al. / Surface & Coatings Technology 215 (2013) 178–185
On the other hand, the annealing time at a given temperature hasonly little influence on the final material. Free carbon is the onlyphase appearing in small amounts after a few minutes at 1200 °C,while a slight SiC grain growth is detected (not shown). No furtherchanges in the structure are observed for durations up to 60 min.They are indeed more significant at 1300 °C. Free carbon traces firstappear almost instantly (for a dwell time t=0 min, i.e., after heatingat 1300 °C and immediate cooling) and the amount increases rapidlyand stabilizes during the first 15 min of dwell time (Fig. 11). The B4Cphase is formed shortly afterwards and the structure remains rela-tively stable between t=30 and 60 min. A rapid growth of the SiCgrains is also observed during the first 15–30 min at 1300 °C(Fig. 13). As for free carbon and B4C, the SiC crystallite size remainsapproximately constant between 30 and 60 min. The final grain sizedepends on the composition of the coating. It is about 5 nm for theB and C coatings and 7.5 nm for A. In spite of their different composi-tion, the structural evolution of specimens B and C versus tempera-ture and time is relatively similar. The pre-existing SiC nanocrystals(1–2 nm) in these two as-deposited coatings probably impede the re-organization of the a-BxC phase and the subsequent crystallization of
Fig. 10. Intensity of the D band (free carbon) and the C–B–C chain characteristic peak(B4C), as derived from the Raman spectra, as a function of the annealing temperature(t=15 min/high vacuum).
B4C (the B4C phase is indeed not observed on the XRD pattern after a1300 °C–2 h annealing, see Fig. 2). But the ordering of the differentphases reciprocally affects each other's; the presence of theintergranular B4C and free carbon phases that appeared afterannealing, indeed probably explains the limited SiC crystal growthat 1300 °C of B and C (from 1–2 nm when as-deposited, to 5 nm),compared to A (from less than 1 nm, to 7.5 nm, see Fig. 13). On theother hand, such a higher SiC growth rate of the boron-rich material(A) at 1300 °C could be due to the fact that both SiC and B4C phasescrystallize simultaneously (Figs. 2, 9, 13). The B and C atoms motionassociated with the a-BxC conversion into B4C probably promotesthe Si atoms diffusion required for such a large SiC grain growth.
4. Conclusions
The solid deposited by CVD from a BCl3–MTS–H2 mixture is madeof a quasi-amorphous boron carbide matrix (a-BxC) surroundingβ-SiC subnano or nanocrystals, except for high deposition tempera-ture and MTS concentration in the gas phase (e.g. at 1000 °C and for
Fig. 12. SiC apparent crystallite size (111-plane) of the Si–B–C coatings as a function ofthe annealing temperature (T=1300 °C/high vacuum).
β=3.3), where it consistsmainly of polycrystallineβ-SiC. The SiC clustersize decreaseswhen the boron content increases. The a-BxC phase showsonly a very short-range ordering. It includes icosahedron-like units, pre-sumably highly disordered compared to those in rhombohedral B4C. Theintericosahedral B sites include trigonal BC3 environments which arenormally absent in crystalline boron carbide(s). The relative stability ofsuch disordered a-BxC structure, at temperatures approaching 1000 °C,could be related to the high proportion of SiC/a-BxC interfacial atoms inthematerial. The singular B local environment probably explains the ab-sence of free carbon, contrary to what suggests the ternary Si–B–C dia-gram for the majority of the coatings.
If one excludes the most crystalline materials deposited at hightemperature and MTS flow rate, the annealing of the various Si–B–Ccoatings at increasing temperature and time induces the formationof increasing amounts of disordered polycyclic carbon (within the1100≤T≤1300 °C domain), the crystallization of the a-BxC phaseinto B4C (1250≤T≤1300 °C) and the coarsening of the SiC nanocrystals(1200≤T≤1400 °C). These structural changes are achieved in less than30 min, the final structure being apparently frozen in a state whichstrongly depends on the chemical composition (i.e. on the B and the Siconcentrations) and the temperature. In the coatings containing a rela-tively high amount of SiC, the SiC nanocrystals (1–2 nm) probablyimpede the formation of large B4C crystals during annealing. Viceversa, the presence of the intergranular B4C and free carbon phaseslimits the SiC crystal growth at 1300 °C. On the other hand, in theboron-rich coatings containing only small SiC clusters, both the SiCand B4C phases crystallize simultaneously and lead to larger crystals at1300 °C. The B and C atoms motion involved in the crystallization ofthe boron carbide phase likely promotes the SiC grain growth.
At the atomic scale, the phase rearrangement results in (i) the forma-tion of CC3 (hexagonal) sites in the carbon layers, which are initially ab-sent in the material, (ii) a strong decrease of the proportion of trigonalBC3 sites, in favor of the more conventional inter- (C–B–C chains) andintra-icosahedral sites of rhombohedral B4C, and (iii) more regular tetra-hedral SiC4 sites involved in largerβ-SiC nanocrystals. Yet, a few BC3 sitesapparently remain after crystallization in the Si-rich coatings, possiblydue to the SiC/B4C interfacial bonding.
We are currently investigating the high temperature tensile be-havior of Si–B–C-coated monofilaments at high temperatures. The
consequences of the structural changes on the thermo-mechanicalbehavior of the Si–B–C coatings will be the matter of a forthcomingpaper.
Acknowledgments
The authors are indebted to CNRS and Snecma Propulsion Solide(Groupe Safran), for allowing a grant to C. Pallier. They are gratefulto Synchrotron SOLEIL for supporting the X-ray photoabsorption ex-periments on TEMPO (proposal no. 20110047). They also acknowl-edge A. Delcamp, from SPS, for valuable discussions.
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