Pyrolytic conversion of organopolysiloxanes

Accepted Manuscript Title: Pyrolytic conversion of organopolysiloxanes Author: Djamila Hourlier Srisaran Venkatachalam Mohamed-Ramzi Ammar Yigal Blum PII: DOI: Reference:

S0165-2370(16)30618-0 http://dx.doi.org/doi:10.1016/j.jaap.2016.11.016 JAAP 3882

To appear in:

J. Anal. Appl. Pyrolysis

Received date: Revised date: Accepted date:

1-10-2016 22-11-2016 22-11-2016

Please cite this article as: Djamila Hourlier, Srisaran Venkatachalam, MohamedRamzi Ammar, Yigal Blum, Pyrolytic conversion of organopolysiloxanes, Journal of Analytical and Applied Pyrolysis http://dx.doi.org/10.1016/j.jaap.2016.11.016 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Pyrolytic conversion of organopolysiloxanes

Djamila Hourliera*, Srisaran Venkatachalama, Mohamed-Ramzi Ammarb, Yigal Blumc

a IEMN (CNRS, UMR 8520), Université Lille 1, Avenue Poincaré, CS60069, F-59652 Villeneuve d'Ascq Cedex, France. b CEMHTI, (CNRS, UPR 3079), Université Orléans, CS90055, F-45071 Orléans Cedex 2, France c Chemical Science and Technology Laboratory, SRI International, Menlo Park, CA, USA

Corresponding Author *E-mail: [email protected]

graphical abstract cross5%TMTVS linker polymer CH 3 Si

TMTVS

CH 3

O

O

Si

Si

O

CH 3

O

H

H3C H3C Si

Si O

H3C

O CH3 n

PHMS

Si CH 3

0.7 Intensity (A.U)

0.5

50µm

High carbon 1

Turbostratic carbon signature

0.3

1000

60% DVB polymer

DVB

Polymer-derived ceramics

Low carbon

crosslinker

CH3 Si CH3 CH3

D 1200

1400

0.996

G 1600

0.992 1800

Raman Shift (cm-1)

1200 1400 1600 1800 0.0 Raman Shift (cm-1)

2000

20µm 0.988

Highlights ➤TG/MS analysis provides an excellent analytical support for valuable information ➤ the chronology of events of the thermal degradation of organopolysiloxanes has been established. ➤ Divinylbenzene enables one to achieve desirable characteristics and controllable carbon contents. ➤ Raman mapping provides information on densities of defects in turbostratic carbon. Abstract This study extends previous work on in-situ generation of nanographene domains in silicon oxycarbide polymer-derived ceramic nanocomposites. The thermal conversion of cross-linked polyhydridomethylsiloxanes (PHMS) has been studied as a function of its initial organic structure and its content. Two cross-linking additives containing vinyl groups, tetramethyl-tetravinyl cyclotetrasiloxane (TMTVS) and divinylbenzene (DVB), have been reacted with the Si-H groups of the PHMS via Pt catalyzed hydrosilylation. These polymers gave a high ceramic yield of 7885 wt% upon pyrolysis at 1000°C-1400°C in inert atmosphere. The use of the aromatic crosslinker (DVB) induces significantly higher carbon content in a controlled fashion into the derived ceramics compared with the TMTVS. The chemistry involved in the in-situ evolution of lowdimensional graphene architectures embedded into silicon oxycarbide network has been studied in detail using thermal gravimetric, mass spectroscopy analyses (TG-MS) and Raman spectroscopy. The presence of aromatic functional groups (DVB) at the polymeric stage leads to graphene generation at lower temperatures than in the case of TMTVS cross-linker. Raman spectroscopy analysis of DVB cross-linked PHMS showed evidence of the beginning of formation of free sp2 carbon domains from 800°C. Raman mapping allows obtaining the chemical and the structural distribution of the resulting phases.

Keywords: polyorganosiloxanes, Raman spectroscopy, hydrosilylation reaction, graphene nanocomposites, silicon oxycarbide, TG-MS, polymer-derived ceramics (PDC), pyrolysis

1. Introduction Silicon-based polymers offer a large number of reactions by selecting the appropriate backbone units of the polymeric structure[1]. The most interesting modifications consist of organofunctional groups, either directly incorporated into the polymer backbone or grafted to the Si atoms. These polymers can be used for a broad spectrum of applications, as adhesion promoters for surface coating due to their outstanding water-repellent properties[2, 3] and as

2

additive incorporated in organic-based composites. They serve as high temperature elastomeric materials and function as selective membranes[4]. Another emerging field for silicon based polymers is “preceramic polymers”, in which the polymers are cross-linked and then thermally converted to silicon-based ceramics by pyrolysis. A large number of thermal conversions of organosiloxanes ceramic precursors have been studied and reported in the literature [1, 5, 6]. A common feature of polymer-derived ceramic route is the formation of amorphous metastable intermediate phases of SiCxOy in which an excess free carbon is embedded. It must be recognized that the thermal conversion of organically modified silicon precursors is the easiest way for preparing ternary and even multinary phases (Si/C/N/O/Al/Mg/B/Zr/Ti/etc.)[7, 8]. The excess carbon, which is not directly incorporated into the inorganic Si-O-C networks, results from the thermal decomposition of organic functional groups, incorporated in the polymeric precursor structure. The “free carbon” content depends on the molecular structure of the organic substituent group and its reactivity with neighboring functional groups present in the structure of precursor [9]. The free carbon, recognized for its turbostratic nature, can be either advantageous or detrimental, depending on the application for which the material will be used. In oxidizing atmosphere, carbon is the most vulnerable constituent [10], especially at low temperatures, where protective layer of silica cannot be formed, to inhibit the oxygen diffusion[10, 11]. In contrast, in inert atmospheres, the free carbon can be advantageous for the formation of SiC and other metal carbide phases via carbothermal reductions [12, 13]. It can also improve the thermal stability, by inhibiting and retarding transitions from the metastable states (oxycarbide phase) to the more stable phases (SiC and SiO2)[13, 14]. A great deal of attention has been directed in the past to lower the free carbon content in polymer derived ceramics. Such attempts were intensive in the case of polycarbosilane precursors for SiC fibers and matrices[15]. This ceramic fiber was intended to be used as reinforcement for thermostructural composites[16]. It should be noted that the first generation of SiC fibers (Nicalon™) from polycarbosilanes contained not only an excess carbon, but also SiOC phase. The oxygen incorporation is derived from the oxygen used during the curing step. Advances in chemical synthesis and processing have resulted, indeed, in fibers with reduced free carbon and oxygen content. New generations of stoichiometric SiC fibers (with Si/C=1) with no excess carbon and very low oxygen content are commercially available as Hi-Nicalon[17]. Knowing that the thermal conversion of organic groups to graphene domains is inevitable, the question arises: Can we take advantage of the in-situ formation of these nanodomains inside the SiOC, SiCN or similar matrices as functional materials for specific applications? A growing interest in carbon-rich silicon oxycarbide materials has been noticed for their potential as anode materials for rechargeable lithium-ion batteries[18, 19], or for temperature and pressure sensors [20, 21] that can be used in harsh environments. In our previous work, we have also demonstrated that the carbon-rich materials issued from pyrolysis of organopolysiloxanes exhibit a high absorption in terahertz domain[22]. Recently, the high carbon materials discussed in this article have been reported for their extremely high permittivity values, exceeding those of state of the art titanates[23]. Various methods were reported for tailoring chemical compositions of polymer-derived ceramics containing high carbon content. It is mostly done by incorporating phenyl-Si groups into the 3

silicon based polymer[24]. For example, Scheffler et al. and others[25, 26] used the advantage of excess carbon to catalytically grow in-situ carbon nanotubes from hydrocarbon species released during the pyrolysis of phenyl containing polysiloxane (Sylres H44). For the preparation of nanoporous SiC-based materials and cross-linking purposes, many researchers incorporated a cross-linking agent containing vinyl groups to initiate the hydrosilylation reaction [27]. The vinyl groups also increase the carbon content of the PDC (polymer-derived ceramic) but not as much as an aromatic group. Blum et al. [28] have deliberately increased the carbon content by incorporating, at molecular level and in a controlled manner, divinyl benzene (DVB) to achieve a highly cross-linked polyhydridomethylsiloxane (PHMS) simultaneously with controlled increase of the carbon content. Their experiments have shown that the ceramic yield of the cross-linked polymers is even higher. A remarkable property of resulting materials is their oxidation stability in-spite of the high carbon content [11, 13]. However, volatile species from the pyrolysis of polymers cross-linked with DVB have not yet been studied in detail, which can illuminate differences in the mechanisms of the evolution of the graphenic phase when generated mostly from either methyl or the DVB groups. Previously, Maddocks et al.[29] attempted to analyze by thermogravimetry coupled with mass spectrometry the volatile species of cross-linked polycarbosilane with DVB. Unfortunately, oxygen contamination probably due to a leak in their system does not seem to be reliable. The authors even concluded their article by suggesting that further experiments in an oxygen-free atmosphere are required. The aim of the present work is to extend previous research on DVB cross-linked PHMS polymers by carefully investigating the gaseous species eliminated during pyrolysis and attempt to elucidate the reaction mechanisms responsible for the in-situ formation of the graphene in the SiOxCy ceramic host due to the presence of the DVB as a carbon precursor. 2. Materials and Methods Mixtures comprising a commercial polyhydromethylsiloxane (PHMS) and two types of crosslinkers, vinylmethyltetrasiloxane cyclomer (TMTVS: 5 wt% relative to PHMS weight) and divinylbenzene (DVB: 60 wt%) were processed according to the procedure previously reported [30]. A Pt divinyltetramethyldisiloxane catalyst in xylene (5 ppm relative to PHMS) was added to the mixture of PHMS and one of the two cross-linkers. The reactions were performed in a liquid phase with no solvents added (except for diluting the Pt catalyst), leading to a dense, cured polymeric, resin-like SiOC material (schema of Figure 1).The blends were stirred for about 20 min in dry environment, cast in a dish to form 2–3 mm thick specimens and cured at room temperature for 12 h. The curing can be accelerated at 40°C for PHMS-TMTVS or at 120°C for PHMS-DVB. The cured material is a hard and brittle rubbery material with very high level of transparency. Subsequently, the cured preceramic bulk specimen was pyrolyzed under flowing inert gas. For simplicity, hereafter in the text the gels will be referred to as PHMS for (PHMS-5 wt.%TMTVS), and as 60DVB for (PHMS-60 wt. % DVB). To understand what changes occur in cross-linked polymer during heat treatment the following techniques were used: thermogravimetry analysis coupled with mass spectrometry (TG/MS), Raman spectroscopy, and scanning electronic microscopy (SEM). Thermogravimetry analysis (TG) (Netzsch STA449F3 Jupiter apparatus) coupled with a quadrupole mass spectrometry (MS) (Aëolos QMS403D, 70 eV, electron impact) via a new heated capillary system, were applied to simulate the behavior of material during the annealing process and to determine the evolved gaseous species. Before each experiment, the TG system 4

was first evacuated and then flushed with ultrahigh purity helium for at least 30min before starting heating. The experiments were carried out under dynamic inert gas atmosphere (helium: 99.999 purity) at a flow rate of 80 cm3/min. The samples (20-400 mg) were heated in Al2O3 crucibles up to 1500◦C, with a heating rate of 10◦C/min. Raman spectroscopy measurements were performed using a Horiba Jobin-Yvon LabRam®HR micro-Raman system combined with a 473 nm laser diode as excitation source focused by a ×100 objective. The scattered light was collected by the same objective in backscattering configuration, dispersed by a holographic grating of 1800 grooves/mm and detected using a CCD camera. A minimum of five different regions per sample were analyzed. The microstructure characterization of the analyzed zones was accomplished using optical microscope equipped with a video camera, for viewing sample and record photographs. This is a very helpful, particularly, to verify the suitability of the selected parameters (exposure time and laser power) and avoid an additional contribution, which may arise from laser-induced sample heating. In order to obtain a complete picture of the chemical and structural distribution of various phases, Raman mapping has been performed. In this particular case, a laser excitation wavelength of 514.5 nm, a holographic grating of 600 grooves/mm, and a ×50 objective were used. The data were first processed by principal component analysis (PCA) giving an indication of independent components in the map. Then, a direct classical least squares (DCLS) method was used consisting in rigorously map reconstruction by a linear combination of spectra from the pure components contained in the sample. 3. Results 3.1Thermogravimetry analysis It is well known that some siloxane-based gels containing Si-H pendant units may have different adsorption/reaction in ambient air and moisture due to dehydrocoupling and oxidation reactions. Thus, it is essential to evaluate their air sensitivity, since it has a decisive influence on the thermal reaction mechanisms, and can affect the ceramic yields and evolved species. Monolith gels were compared with freshly ground powders and also aged powders. The state of matter is differentiated by the suffixes: FP (Freshly ground Powder), AP (one week Aged ground Powder), and M (Monolith). Preliminary tests with different sample masses and shapes were carried out, in order to adjust our experimental parameters and to overcome problems, such as the effects of heat transfer and initial mass sample, often encountered on thermal decomposition of materials. Under the same experimental conditions, our results confirmed a very good reproducibility and this is demonstrated in Figure 2a, which shows a direct overlay of duplicated experiments performed on the freshly ground powders of 60DVB. For monolithic gels, the thermograms indicate the thermal decomposition is independent of the initial mass, at least for masses ranging from 20 to 300 mg, as typically used in our experimental procedures (Figure 2a). As one might expect, the ground powders were generally more reactive than their monolithic counterparts (Figure 2b). The largest weight loss occurred with aged powders. This variation could be attributed to the presence of SiOH bonds issued from hydrolysis of Si-H. Cross-linked monolithic gel decomposition occurred in essentially two major stages. In the first temperature range between 150°C-600°C, 11% weight loss is recorded for 60DVB, compared to only 5% for PHMS. However, in the second stage, starting from 600°C to 1350°C, there is no 5

significant difference between the two gels; the weight loss is around 9% for both. For 60DVB, a third weight loss takes place after 30 min at 1350°C.At 1350°C, the ceramic yield from monolithic gels is relatively high, 86% for PHMS vs 78% for 60DVB.

3.2 Mass Spectrometry analysis: For describing the pyrolytic degradation of gels, one needs to know both the initial structure of precursors as well as the volatile species evolved during the heat treatment. Mass spectrometry coupled with thermogravimetry was used to continuously monitor the evolving volatile products, which were carried out of the TGA furnace through a capillary heated to 150°C. The timedependent evolution of the gases is shown in Figures 3 and 4. Due to the presence of numerous different fragments, only ions of most importance to our discussion have been selected and displayed for providing a clear presentation of the MS profile curves. It should be remembered that the peaks observed in the mass spectra cannot be directly assigned to the effluent issued from thermal degradation of the heated product. It has been shown that further fragmentation recombination by electron impact ionization occurs in the ion chamber of the mass spectrometer. During heating from room temperature to 450°C, the PHMS exhibits a mass loss of ∼ 0.7% to 5% depending of the state of matter (Monolith: M or Fresh powder: FP). The detected gases (Figure 3) correspond to H2O (M/z=18) and silicon-based species, mainly (CH3)3Si- (m/z = 59, 73) ascribed to a terminal site of the base PHMS polymer. The presence of H2O leads to the hydrolysis of these terminal sites and to the cleavage of (CH3)3Si-OH (M/z= 75, 45), according reactions (1) and (2). CH3 H3C Si O Si CH3

+H2O

CH3 H3C Si O Si CH3 Si H

CH3 H3C Si OH +OH Si CH3 (M/z=75,59) CH3 H3C Si H + CH3 (M/z=59,75)

Si O Si

(1)

(2)

The H2 (M/z=2) and H2O (M/z=18) evolution at this stage is associated with dehydrocoupling reactions between two Si-OH, and between Si-H and Si-OH, according to Equations (3) and (4): Si OH + H O Si

Si O Si

+ H2O (M/z=18)

(3)

Si OH + H Si

Si O Si

+ H2 (M/z=2)

(4)

It should be noted that during the formation of the cured PHMS via the hydrosilylation, the Pt catalyst also performs as a sluggish dehydrocoupling catalyst slightly converting Si-H bonds to 6

Si-OH due to the presence of moisture in the reaction solution. The in-situ formation of H2O is due to condensation reactions between 2 Si-OH species. Between 450°C and 600°C a greater weight loss of about 2.3% is recorded. A range of silicon containing species starting with (CH3)SiH2- (M/z = 45, 44, 43, 31, 30) and then (CH3)2SiH- (M/z = 59, 58, 45, 44, 43) and (CH3)3Si- (M/z = 73, 59, 42) are detected in this region. The elimination of hydrocarbon species containing at least two carbons began around 500°C, peaked at 620°C for methane (M/z = 16, 15, 14, 12) and at 595°C for ethane (M/z = 28, 26), ethylene (M/z = 28, 30) and potentially acetylene (M/z = 26, 25). The elimination of the multicarbon species is completed close to 700°C. Above 700°C, only methane and H2 are detected. It should be noted that no aromatic species are evolved. Also important to notice, that the methane evolution has a bimodal characteristics peaks at around 620 and 710C. For PHMS, the carbothermal reduction starts at 1500°C after 10 min of holding time, CO (M/z= 28,12), CO2 or SiO (M/z= 44) are detected. The TG-MS analysis of the DVB-based polymer reveals the same species listed for PHMS in the same temperature range (Figure 4). However, the hydrolysis reaction and departure of H2O are less pronounced due to the lower concentration of Si-OH groups. Between 450°C and 600°C, the weight loss reached 11.21% corresponds to both siloxane products but also a large fraction of methane, ethylene, ethane and benzene-based volatile species such as styrene (M/z = 104, 78) and/or xylene (M/z =106, 91) or toluene (M/z =92, 91). The emission of light hydrocarbon species, such as CH4 (M/z= 16) with a well distinguished peak at 595°C, are very likely due to the breaking of Si-CH3, or -CH2-CH2- bonds in the polymer and are not derived from the aromatic structures. The aromatic rings remain unaffected and are not fragmented by heat but condense at higher temperatures through loss of hydrogen attached to the periphery of aromatic carbon, leading to larger polyaromatic molecules. Continuation of the temperature increase above 600°C produced a new emission of CH4 (M/z = 16, 15) peaked up at 700°C, whereas H2 (M/z = 2) progressively increases until 800°C, but still present even at 1400°C. Heating further, the carbothermal reduction starts around 1450°C, earlier than for PHMS, with mainly an escape of CO (M/z = 12, 28) and also SiO or CO2 (M/z = 44). The carbothermal reduction reactivity starts earlier and is slightly greater in 60DVB than in PHMS.

3.3 Scanning Electron Microscopy When the gels are subjected to thermal treatment, a gradual change in color was observed. The transparent gels, both PHMS and 60DVB, slowly become light yellow at 450°C and yellow at 500°C. A notable difference was found at 700°C, while the PHMS remains a translucent orange material, the 60DVB turns into opaque and dark brown. Above 800°C, all samples become black due to the generation of free carbon [30]. At 1450°C, the remaining materials lose additional weight and the evolution of CO is detected by mass spectrometry due to carbothermal reduction. The materials turn to bluish color with white spots. The SEM observation focuses only on materials treated at 1450°C in which SiC phase has been detected by Raman as shown hereafter. The microstructure of PHMS and 60DVB treated at 1450°C, is displayed in Figure 5 and 6, respectively. The microstructures of the materials present vermicular particles interconnected with a high porosity (left inset in Figure 5). Circumferential crack-like voids (appearing darker spots than surrounding white material) can be discerned, which may be created by local reaction yielding to 7

the gas evolution and a porous material that shrinks away from its surrounding as shown in the right micrograph of Figure 5. Similarly to PHMS, the 60DVB material at 1450°C, shows (Figure 6) vermicular, voids but no matter is pulled out of the cavities. These cavities have diameters ranging from 1-2 μm.

3.4 Raman spectroscopy Figure 7 provides the assignment of the various bands appearing in Raman spectra of the two starting cross-linked polymers. All precursors feature the PHMS backbone with strong bands associated with C-H asymmetric and symmetric stretching of -CH3 group attached to Si (Si-CH3) at 2969 cm-1 and 2909 cm-1, respectively. Another important band located at 2169 cm-1 is characteristic of the Si-H stretching mode. Other weak bands appear at 1409 cm-1 (C-H asymmetric bending), 1260 cm-1 (C-H symmetric stretching), are characteristic of CH3 attached to silicon. Furthermore, the bands at 763 cm-1 and 705 cm-1 are attributed to Si-C stretching of three methyl groups bonded to silicon corresponding to end groups, while the band at 493 cm-1 is characteristic of symmetric stretching of Si-O-Si. The band at 911 cm-1 is related to the Si-O stretching of the Si-OH group resulting from dehydrocoupling of Si-H bonds with external moisture. Evidence for vinyl bands of the crosslinker TMTVS in the region 1660 cm-1 is missing, indicating either all vinyl groups reacted or because their content is too low to be detected. For the 60DVB-based gels, additional bands located at 3059 cm-1 (C-H stretching of aromatic ring) are due to Csp2-H bonds of DVB, a doublet around 1610 cm-1 and 1583 cm-1 belonging to C=C stretching within the aromatic ring, and an intense band at 1009 cm -1 characteristic of the phenyl group. The Raman bands of the C-H in plan vibration of aromatic 1-4 di-substituted at 1231 cm-1 and CH2 scissors vibration in aliphatic at 1456 cm-1 are usually very weak. C-H aliphatic groups may result from saturating the vinyl groups via hydrosilylation or polyaddition [31]. 3.4.1 Thermal evolution of PHMS: The spectrum of cured PHMS is compared with its respective spectra obtained after heating at different temperatures. When PHMS is heated up to 550°C, the band assigned to Si-H at 2160 cm-1 diminishes in intensity and becomes broader. At higher temperature, the bands assigned to C-H groups are still present, while the intensity of Si-H stretching band decreased and was found to exhibit three shoulders that are characteristic of the presence of Si-H groups having different environments (the inset of Figure 8). The band at 911 cm-1 assigned to Si-OH disappeared at 550°C, which is consistent with further condensation of silanols to siloxanes Si-O-Si. Between 700°C and 850°C, the Raman Spectra are featureless suggesting that the precursor almost fully converted from organosilicon to inorganic structure, since no characteristic bands of Si-H, or any hydrocarbon groups are apparently observed. However, it should be noted that the fluorescence intensity level is progressively high for materials heat-treated between 500°C to 850°C. As a competing process to Raman scattering, the fluorescence emission can overwhelm the whole range of interest due to the presence of even very low amounts of florescent species that the host material may contain at this temperature range. This obstacle would be surmounted (in some cases) by using the appropriate laser excitation wavelength which avoid the fluorescence emission. In the present case, several attempts were made using laser excitation wavelengths in the visible range (from 457 nm up to 785 nm) without achieving improved spectra.

8

The heating at 900°C of PHMS provides the first evidence for the formation of sp2 carbon signature, in which the basic structural units (BSU) are formed with small coherent domain diameter. Due to the resonant Raman scattering process in sp2 carbon, the first order Raman spectrum exhibits two intense but relatively broad bands corresponding to disordered turbostratic carbon at about 1595 cm-1 (G-band) and 1343 cm-1 (D-band) when recorded with the excitation laser wavelength (473 nm). The G band is due to the in-plane bond stretching of all pairs of sp2 carbon atoms (E2g symmetry). Basically, the G band of graphite is located at 1581cm-1 but due to phonon confinement effect, it shifts toward higher wavenumbers up to 1600 cm-1. The D band corresponds to the defect-activated in-plane transverse optical (iTO) phonons located in the vicinity of Dirac point K (A’1 symmetry) where the strongest electron-phonon coupling occurs. Other weak features also occur in the second order Raman spectrum corresponding to the overtones and combinations of Raman fundamental bands. The D and G bands are the typical signature of distorted nanographene domains. When the samples were heated from 900°C to 1350°C, the line-broadening of the two bands (D and G) gets narrower suggesting a better structural organization. From 1450°C, the G band shifts toward lower wavenumbers leading to the appearance of another defect-activated longitudinal optical (LO) phonon (D’ band) located near the Brillouin zone center[32]. The G band shift is also a signature of the increase of the coherent domain diameter. This change is accompanied with a slight variation in the second order Raman spectrum, mainly the appearance of a welldefined bands (2D and D+D’ bands) located at 2690 cm-1 and 2937 cm-1 respectively, conforming the ordering of the sp2 carbon lattice. At 1450°C and for longer residence time (150 min), the PHMS displays signals in the wavenumber region 650 - 1000 cm-1 corresponding to Si-C bonds, whereas all the sp2 carbon features seem to be disappeared. In fact, since the SiC formation results from the carbothermal reduction involving free carbon and silica sites, the turbostratic sp2 carbon is supposed to be almost consumed, this explains the disappearance of the carbon features in the individual Raman spectrum. However, since the material is chemically heterogeneous, Raman mapping may be of outstanding importance in providing a complete picture of the distribution of the different phases. Owing to the huge number of the collected spectra, a first data processing by Principal Component Analysis (PCA) was performed giving an indication of the various components that are present in the Raman spectra. The reconstruction of the maps has been achieved using a Direct Classical Least Squares (DCLS) method, which is based on a linear combination of the spectra from the various pure components obtained by PCA. In other words, the more the regions tend to red color, the more the corresponding Raman spectra match the Raman spectrum of the reference. Figure 9a and 9b present the reflected light micrograph and the SEM image obtained on the PHMS heat-treated at 1450°C for 150min, where Raman mapping has been performed. Figure 9c presents a magnified SEM image corresponding to the square of Figure 9b and showing the surface microstructure. The map scores of Raman mapping show the spatial distribution of each of three phases [(d) low crystalline silicon carbide, (e) high crystalline silicon carbide, (f) disordered sp2 carbon] and their corresponding reference spectra used. Raman mappings of figure 9d, 9e, 9f indicate that the whole matrix is homogeneously associated with low crystalline silicon carbide while high crystalline is mainly located within the area that appears in white color in the optical image. Raman mapping also shows that almost all sp2 carbon was consumed particularly around high crystalline silicon carbide. The remaining sp2 carbon is located inside the cavities in 9

an extremely low content since a high zoom level was needed to detect the color contrast in the corresponding Raman map. 3.4.2 Thermal evolution of 60DVB: The dominant spectral changes of the 60DVB material begin around 100°C lower than for PHMS. At 500°C, the Si-H bands completely disappeared, without any changes in their environments during the elimination as shown in the inset Figure 10. Only bands characteristic of phenyl group are still present. Between 600°C and 700°C, the Raman spectra are featureless, which reflects an inorganic state of the material. Annealing at 800°C leads to the appearance of the two characteristic bands of disordered graphene’s D band (1343 cm-1) and G band (1600 cm-1). The 60DVB derived material promotes the formation of sp2 carbon at 100°C below than any other organic groups directly attached to silicon (methyl, vinyl, or even phenyl)[14]. Moreover, this first carbon residue appears to be more structurally ordered, characterized by narrower bands compared to the PHMS derived material. This is due to the fact that 60DVB precursor already contains aromatic groups which are the basis of the graphene domain evolution. Further annealing yields to a progressive narrowing of the D and G bands. The G band is slightly shifted towards lower wavenumbers leading to the appearance of D’ band and also a well-defined 2D peak at 2690 cm-1 appears. The change of all these parameters indicates a better organization of the sp2 carbon domains. Heating the 60DVB-based material at 1500°C is accompanied by weight loss, evolution of CO as detected by TG/MS and a change in microstructure where several black spots surrounded by white rims can be observed (Figure 11a). These black spots are found to be holes, as illustrated in Figure 11 b and c. Raman spectra reveal the presence of two phases: silicon carbide in the form of white rim and sp2 carbon covering the whole matrix. This indicates that the carbothermal reduction is local and related to the presence of free sp2 carbon. Raman mapping confirm for 60DVB also the local formation of SiC appearing as white circles in the optical image (Figure 11d). Otherwise, the material is mainly composed of sp2 carbon. However, one of the interesting findings here is the ability of Raman mapping to reveal the presence of two types of carbon structures in the vicinity of the formed SiC. Basically, the D band intensity, in the Raman spectrum of sp2 carbon, is a commonly used as an indicator to quantify the defect density of carbon based material. Therefore, Raman mapping shows that relatively well organized carbon structure (Figure 11 e and h), characterized by low intensity of defect band D, is surprisingly found to be directly in contact with the produced silicon carbide (Figure 11 d and g), whereas relatively disordered carbon (Figure 11 f and i) structure appears to be located in the interface between the matrix Si/C/O and well organized carbon. 4. Discussion The thermal conversion of cross-linked polymers involves a complex series of chemical reactions and many volatile species are driven off. Ceramic yields of 88 wt% to 79 wt% at 1400°C were achieved, depending on the molecular structure of organosiloxane precursor, and the degree of cross-linking. Below 400°C, apart from the evaporation of a small fraction of oligosiloxanes from the polymer via rearrangement reactions, other species are also detected by mass spectrometry such as H2O and H2 suggesting the presence of residual Si-OH bonds that undergo hydrolysiscondensation and dehydrocoupling reactions. Further evidence for the presence of Si-OH is also given by 29Si-NMR [30](T2: C-O2Si-OH units at - 55.9 ppm) and Infra-red analyses not reported here. Between 450°C and 600°C organosilane molecules were detected by MS for both the PHMS and 60DVB materials. These species are attributed to redistribution reactions involving Si-O, Si-H, and Si-C bonds. These reactions have been described in detail in earlier publications 10

[5, 9]. For example, the redistribution reaction between Si-O and Si-H according to Equation (5), yields organosilane (CH3)3-SiH and the formation of T units in the material. Si-O and Si-H bonds exchange involving terminal sites of PHMS ((CH3)3SiO) CH3 H3C Si O H3C

O

H Si O CH3

CH3 O H3C Si H + O Si O CH3 CH3 (CH3)3SiH

(5)

T

The presence of such silicon environments derived from these reactions, are strongly supported by the results of solid 29Si-NMR spectra[30] that clearly show the enhancement of T3 (CSiO3) units at – 65.8 ppm in parallel to the disappearance of DH units at – 33.7 ppm to – 36.8 ppm, when heating the two materials up to 600°C. As mentioned in the previous MS section, organosilane species cleaved out during the heating may in turn undergo additional series of rearrangements resulting of simple gaseous hybrids and light hydrocarbon molecules like ethylene and methane. The Si-H bond (at 2160 cm-1) detected by Raman spectroscopy of the 60DVB polymer diminishes and hardly observed at around 500°C due to the fluorescence and also its low concentration. In contrast the Si-H is still present at 600°C in the case of PHMS but has been decomposed into three sub-bands assigned to three different environments (Units DH at 2170 cm1 , units DH2 at 2125 cm-1, and units TH at 2235 cm-1). It is well known that the Si-H stretching absorption frequency is very sensitive to the atoms attached to the adjacent silicon [33]. In fact, the Si-H bonds can survive up to 700°C, as seen by infrared spectroscopy. The heated 60DVB also releases at this stage fragments derived from the divinylbenzene as detected by MS at M/z = 91, 78. The 13C NMR spectra of 60DVB[30], heated at 400°C, reveals a singular peak at -28 ppm, which is assigned to aliphatic carbons (-C-C-C- bonds). Such a signal was never seen in polymers containing vinyl and Si-H, which may react either by hydrosilylation and/or polyaddition of vinyl groups[24]. The presence of linear hydrocarbon chain containing three carbons may be originated from the activity of DVB. In the alkenylbenzene, in general, the double bond shows higher activity than the resonance-stabilized benzene ring. Knowing that silicon is electropositive in most of its compounds, when compared to carbon and oxygen, the exchange of the double bond of DVB and Si-CH3 may occur leading to the formation of new DH units, according to the equation (6): CH CH2 O Si CH3 O H

H C CH2 CH3 O Si H O

(6)

The influence of DVB cross-linker is quite significant. The remaining styrene groups are also capable of further cross-linking by reaction of the Si-H or by addition on another vinyl group. However, it seems that the unsaturated substituents easily polymerized by addition than by hydrosilylation reaction, suggesting that the Si-H bonds are still present in the material, in spite of high DVB content. The polyaddition involving Vinyl groups leads to the insertion of carbon into silicon network. These reactions are well documented in[24]. In the same temperature range, 11

nucleophile attack of residual Si-OH on Si-C and on Si-H occurred, according to Equation (7) and (8) Si CH3 + Si H

+

Si OH Si OH

Si O Si Si O Si

+ CH4

(7)

+ H2

(8)

These types of reaction are slightly shifted toward higher temperatures; starting from 435°C and is continuing up to 725°C. The main products released are hydrogen and light hydrocarbons such as CH4, C2H4 for PHMS, and CH4 only for 60DVB. This type of reactions correspond to the first peak of methane with a maximum situated at 575°C for 60DVB and at 620°C for PHMS. Between 650°C and 1000°C, the volatile degradation products consist mainly of CH4 and H2. In spite of the similarity of evolved gaseous products, there are, however, some differences in the gas evolution patterns. The 60DVB material exhibited one peak with a maximum release situated at 695°C whereas in PHMS, two peaks clearly appeared; one at 700°C and another one as a shoulder at 770°C In order to clarify the mechanisms and to establish the chronology of the various events, it is useful to compare other structures of gels, already analyzed in our previous work [5], with the present observations with PHMS and 60DVB. The selected gels are defined as follows: T for the gel issued from CH3Si(–OEt)3, DHTH 1/1 for the one issued from the same ratio of DH (CH3)HSi– (OEt)2 and TH triethoxysilane HSi(–OEt)3, TTH from a mixture of T and TH, and TQ from a mixture of CH3Si(–OEt)3 and Si(–OEt)4. The reactions involved in the synthesis and the characterization of gels were described elsewhere [5]. The formation of methane evolved during the heating of various structural units of gels is reported in Figure 12 for a comparative analysis. Contrary to the DHTH (1/1), the TQ and TTH gels are partially condensed and contain Si-OH bonds that react, according the equations (7) and (8) discussed above. These reactions take place below 650°C and are common to all gels containing Si-OH bonds: PHMS, 60DVB, TTH, and TQ. Above 650°C, a further release of CH4 and H2 occurs for all gels containing Si-H and Si-CH3, like in PHMS, 60DVB, TTH and DHTH (1/1). This suggests that the Si-H bonds take part in the reaction and initiate the scission of SiCH3. The formation of free-radicals by homolytic bond cleavage of Si-H, followed by hydrogen abstraction, or initiation of homolytic cleavage of Si-C and then radical combinations, have been well documented in the literature [5]. Thus it would be unnecessary to detail them all here. Above 750°C, another peak of CH4 appeared for all gels containing Si-CH3. This suggests that in the absence of Si-H in the structure of gels (TQ, T), the homolytic cleavage of Si-CH3 starts at higher temperatures, above 700°C, as evidenced by a well separated peak of methane issued from TQ and T gels, for which the maximum release is shifted at around 800°C. These homolytic cleavage reactions typically have an initiation, propagation, and a termination stage. Between 900°C and 1400°C, the pyrolysis gases are essentially hydrogen, issued from the external hydrogen attached to aromatic structures of free carbon phase. Above 1400°C, carbothermal reduction may occur, depending on the content of free carbon and its distribution within silicon oxycarbide phase. These heterogeneous reactions between solid phases (SiC xOy, SiO2 and free carbon), generate porosity and lead to the formation of SiC phase and departure of CO and SiO. Since the carbothermal reactions require the mobility of the elements it is performed 12

by gas phase interactions of SiO with the carbon phase, according to equation (9) and (10)[12, 14] SiO2(s) + 3C 2C + SiO(g)

SiO(g) + 2CO (g) SiC(s) + CO (g)

(9) (10)

Thus, the weight loss during this reaction can be substantial. Nevertheless, this is not the case for both PHMS and 60DVB materials. Especially, it is negligible at 1500C for monolytic material that is still dense prior to the carbothermal reduction and has very low surface area, which can interact with the evolved gases. The volatiles can have a major effect on the material’s structure, and therefore on the end-use and properties of materials. In spite the high content of the extra carbon via DVB, thermal conversion leads to a slightly higher pyrolysis yield (21 wt% loss) at 1400°C, than for example phenyl directly attached to silicon (26 wt % loss), that we previously studied in our other work on thermal conversion of preceramic polymers[14]. This indicates that most of the DVB has been retained in the evolved ceramic structure. 5. Conclusion Thermogravimetry coupled with mass spectrometry analysis provides an excellent analytical support for valuable information about (i) evolution of volatile species during the pyrolysis, (ii) separation and elucidation of reactions involving different substituents, (iii) the presence of Si-OH, even in low concentration. This can be achieved by heating the sample to trigger off condensation reactions leading to the water loss (H2O), easily distinguished at M/z = 18. The major findings from this work are summarized as follows. (1) When DVB is incorporated into organosilicones frameworks, it is not easily fragmented via thermal cracking to yield volatile species and is converted to turbostratic carbon residues constituted of few graphene layers nanometrically embedded in the SiOC host. In a wider sense, the cleavage of DVB and other aromatic structures in the organosiloxane network is hindered. In contrast Straus and Madorsky[34] showed that the pyrolysis of polydivinylbenzene polymer losses nearly 84 wt% at 450°C, mainly in volatile species. At 960°C, the carbon yield is only 6% based on the weight of the original polymer. (2) The bonding environments surrounding Si atoms are initially Si-H, Si-CH3, Si-CxHy-Si, Si-O-Si and Si-OH, undergo redistribution reactions, thereby new environments around silicon atoms are formed, and at the same time silicon-based compounds are evolved. The rearrangement (metathesis) reactions proceed in the temperature range between 350°C to 530°C. (3) The nucleophile attack of Si-OH on Si-C and on Si-H giving off light hydrocarbon molecules CxHy and H2, respectively, occurred between 450°C to 650°C. (4) During progressive heating, Si-H homolytic cleavage generates free-radical such as active hydrogen atoms that initiate the Si-R cleavage bonds earlier in temperature (around 650°C); hence methane (when R = CH3) appears as main volatile product. Generally, in 13

the absence of Si-H bonds, the homolytic cleavage of Si-C proceeds at higher than 650°C, as seen in TQ and T polymers. (5) The presence of Si-OH and Si-H have a significant effect on the reactivity of Si-CH3, and thus on the insertion of carbon into silicon oxide network. (6) Raman mapping revealed the presence of different phases (SiC, free sp2 carbon). For the first time, two different carbon structures have been well identified (disordered and well organized carbon) in polymer derived ceramics. (7) The presence of high carbon content promotes the thermal reduction at high temperature. However, due to the dense nature of the polymer-derived material, the carbothermal reduction begins only at about 1500°C and it is inhibited even when the excess of the free carbon is significant in the 60DVB material. In the last 40 years, the development of materials from organosilicon polymers that can be used in harsh atmosphere and having desirable mechanical characteristics at high temperatures has led ceramists and chemists to develop and study mostly polymers that lead to low free carbon content in the pyrolyzed materials. Ironically, many of today’s PDC applications count on the presence of the in-situ developed graphene and turbostratic carbon domains. For example, absorbers for photonic and Li battery anodes beyond the capability of carbon-based anodes, made it necessary to develop carbon-rich SiOC. Graphene structures resulting from hydrocarbon groups have good absorbing properties in large range of frequencies from microwaves to infra-red. Moreover, contrary to common belief, and despite the high carbon content in DVB-derived materials, they are remarkably resistant to oxidation, compared to other carbon-based materials.

References [1] Monthioux M, Delverdier O. Thermal behavior of (organosilicon) polymer-derived ceramics. V: Main facts and trends. Journal of the European Ceramic Society. 1996;16(7):721-37. [2] Mahltig B, Fischer A. Inorganic/organic polymer coatings for textiles to realize water repellent and antimicrobial properties—A study with respect to textile comfort. Journal of Polymer Science Part B: Polymer Physics. 2010;48(14):1562-8. [3] Ma M, Hill RM. Superhydrophobic surfaces. Current Opinion in Colloid & Interface Science. 2006;11(4):193-202. [4] Collinson MM. Sol-Gel Strategies for the Preparation of Selective Materials for Chemical Analysis. Critical Reviews in Analytical Chemistry. 1999;29(4):289-311. [5] Bahloul-Hourlier D, Latournerie J, Dempsey P. Reaction pathways during the thermal conversion of polysiloxane precursors into oxycarbide ceramics. Journal of the European Ceramic Society. 2005;25(7):979-85. [6] Colombo P, Mera G, Riedel R, Sorarù GD. Polymer-Derived Ceramics: 40 Years of Research and Innovation in Advanced Ceramics. Journal of the American Ceramic Society. 2010;93(7):1805-37. [7] Homeny J, Nelson GG, Risbud SH. Oxycarbide Glasses in the Mg-Al-Si-O-C System. Journal of the American Ceramic Society. 1988;71(5):386-90. [8] Toma L, Fasel C, Lauterbach S, Kleebe H-J, Riedel R. Influence of nano-aluminum filler on the microstructure of SiOC ceramics. Journal of the European Ceramic Society. 2011;31(9):1779-89.

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[9] Mutin PH. Control of the Composition and Structure of Silicon Oxycarbide and Oxynitride Glasses Derived from Polysiloxane Precursors. Journal of Sol-Gel Science and Technology. 1999;14(1):27-38. [10] Brewer CM, Bujalski DR, Parent VE, Su K, Zank GA. Insights into the Oxidation Chemistry of SiOC Ceramics Derived from Silsesquioxanes. Journal of Sol-Gel Science and Technology. 1999;14(1):49-68. [11] Modena S, Sorarù GD, Blum Y, Raj R. Passive Oxidation of an Effluent System: The Case of Polymer-Derived SiCO. Journal of the American Ceramic Society. 2005;88(2):339-45. [12] Weimer AW, Nilsen KJ, Cochran GA, Roach RP. Kinetics of carbothermal reduction synthesis of beta silicon carbide. AIChE Journal. 1993;39(3):493-503. [13] Vomiero A, Modena S, Soraru GD, Raj R, Blum Y, Della Mea G. Investigation on the oxidation process of SiCO glasses by the means of non-Rutherford backscattering spectrometry. Nuclear Instruments and Methods in Physics Research Section B: Beam Interactions with Materials and Atoms. 2003;211(3):401-7. [14] Latournerie J, Dempsey P, Hourlier-Bahloul D, Bonnet J-P. Silicon Oxycarbide Glasses: Part 1— Thermochemical Stability. Journal of the American Ceramic Society. 2006;89(5):1485-91. [15] Yajima S, Hasegawa Y, Okamura K, Matsuzawa T. Development of high tensile strength silicon carbide fibre using an organosilicon polymer precursor. Nature. 1978;273(5663):525-7. [16] Naslain R. Design, preparation and properties of non-oxide CMCs for application in engines and nuclear reactors: an overview. Composites Science and Technology. 2004;64(2):155-70. [17] Chollon G, Pailler R, Naslain R, Laanani F, Monthioux M, Olry P. Thermal stability of a PCSderived SiC fibre with a low oxygen content (Hi-Nicalon). Journal of Materials Science. 1997;32(2):32747. [18] David L, Bhandavat R, Barrera U, Singh G. Silicon oxycarbide glass-graphene composite paper electrode for long-cycle lithium-ion batteries. Nat Commun. 2016;7. [19] Kaspar J, Graczyk-Zajac M, Riedel R. Carbon-rich SiOC anodes for lithium-ion batteries: Part II. Role of thermal cross-linking. Solid State Ionics. 2012;225:527-31. [20] Riedel R, Toma L, Janssen E, Nuffer J, Melz T, Hanselka H. Piezoresistive Effect in SiOC Ceramics for Integrated Pressure Sensors. Journal of the American Ceramic Society. 2009;93(4):920-4. [21] Roth F, Guillon O, Ionescu E, Nicoloso N, Schmerbauch C, Riedel R. Piezoresistive Ceramics for High-Temperature Force and Pressure Sensing. Sensors and Measuring Systems 2014; 17 ITG/GMA Symposium; Proceedings of; 2014 3-4 June 2014; 2014. p. 1-4. [22] Hourlier D, Venkatachalam S, Boyaval C, Ducournau G, Blum YD, Lampin J-F. Hybrid organicinorganic nanocomposites for terahertz applications. European Materials Research Society Spring Meeting, E-MRS Spring 2014, Symposium I - Solution processing and properties of functional oxide thin films and nanostructures. Lille, France 2014. [23] Blum YD, Chan WK, Hodges JW, Hui DK, Krishnamurthy S, Macqueen DB, et al. High permittivity nanocomposites for electronic devices. Google Patents. [24] Latournerie J. Ceramiques nanocomposites SiCO : Synthese, caracterisation et stabiblite thermique: Unversite de Limoges 2002. [25] Scheffler M, Greil P, Berger A, Pippel E, Woltersdorf J. Nickel-catalyzed in situ formation of carbon nanotubes and turbostratic carbon in polymer-derived ceramics. Materials Chemistry and Physics. 2004;84(1):131-9. [26] Segatelli MG, Pires ATN, Yoshida IVP. Synthesis and structural characterization of carbon-rich SiCxOy derived from a Ni-containing hybrid polymer. Journal of the European Ceramic Society. 2008;28(11):2247-57. [27] Schmidt WR, Interrante LV, Doremus RH, Trout TK, Marchetti PS, Maciel GE. Pyrolysis chemistry of an organometallic precursor to silicon carbide. Chemistry of Materials. 1991;3(2):257-67. [28] Kleebe H-J, Blum YD. SiOC ceramic with high excess free carbon. Journal of the European Ceramic Society. 2008;28(5):1037-42.

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[29] Maddocks AR, Hook JM, Stender H, Harris AT. Heterogeneously catalysed crosslinking of polycarbosilane with divinylbenzene. Journal of Materials Science. 2008;43(8):2666-74. [30] Kleebe H-J, Gregori G, Babonneau F, Blum YD, MacQueen DB, Masse S. Evolution of C-rich SiOC ceramics. International Journal of Materials Research. 2006 2016/07/22;97(6):699-709. [31] Marciniec B. Hydrosilylation Polymerisation. In: Marciniec B, ed. Hydrosilylation: A Comprehensive Review on Recent Advances. Dordrecht: Springer Netherlands 2009:191-214. [32] Ammar MR, Galy N, Rouzaud JN, Toulhoat N, Vaudey CE, Simon P, et al. Characterizing various types of defects in nuclear graphite using Raman scattering: Heat treatment, ion irradiation and polishing. Carbon. 2015;95:364-73. [33] Baraton M-I, Besnainou S. SiH surface species on a SiC nanosized powder: ab initio and FT-IR studies of their behaviour under oxidation. Molecular Engineering. 1996;6(4):327-46. [34] Straus Sidney MLS. Thermal stability of polydivinylbenzene and of copolymers of styrene with divnylbenzene and with trivinylbenzene. Journal of Research of the National Bureau of Standards -A Physics and Chemistry. 1961;65A:243.

16

H

H3C H3C H 3C Si

CH3

Si O

CH3

CH3

Si

Si

CH3

Si O

H3C

CH 3

O

O

Si

n

PHMS

O O

H 3C

Si

CH 3

DVB

Pt Catalyst

Pt Catalyst

TMTVS

O H 3C

Si

H

H 3C Si

Si

Si

O

PHMS crosslinked

CH 3

O

H3C Si O Si

O

O

O Si

H3C

O

Si

O

O

Si

O

CH 3 Si

Si

H3C

O

Si

H

O

H3C

H

O

H3C

O H3C

O

Si

H CH 3 Si

O

O

Si

H3C

O

H3C

O

H 3C

O

O

H

Si

H 3C

O Si

Si

H 3C O

Si

Si

H 3C

CH3

Si

CH 3 H O

O

H3C

H3C

O H

Si

Si

O

H

Si

Si

O

O

O

H

Si

H 3C

H 3C

CH3

O

O

H 3C H

Si

H 3C

O

H

O

Si

Si

Si

O

Si

H

O

Si

O

H 3C

H CH3 O

O

O

O

Si

CH 3

CH3

Si

Si

H 3C

H 3C

O

O

CH3 Si

CH3

O

H

H3C

O

H

O Si

O Si

O

Si

Si

CH 3

O

Si

H 3C

Si

Si

O CH3

CH 3

O

O Si

H 3C

Si

CH 3

H 3C

CH3

O

H2O Pt Catalyst

PHMS-DVB crosslinked

Si CH3

Figure 1. Schematic representation of the two PHMS-derived polymers Temperature (°C)

Temperature (°C)

150 300 450 600 750 900 1050 1200 1350 dwell

150 300 450 600 750 900 1050 1200 1350 dwell

-10

-10 -20.3%

-20 -30

-29.5% 60DVB FP 28mg 60DVB FP 46mg 60DVB M 46mg 60DVB M 309mg

-40 -50

-14.2%

Mass loss (%)

b) 0

Mass loss (%)

a) 0

-16.2% -20.3%

-20 -30

60DVB FP 46mg 60DVB AP 42mg 60DVB M 46mg PHMS FP 43mg PHMS M 43mg

-40 -50 -60

-60 0

20

40

60

-29.5%

80

100

Time (min)

120

140

160

0

20

40

60

80

100

Time (min)

120

-59.3% 140 160

Figure 2. Thermogravimetric analyses of samples in helium at 10°C/min a) effect of sample size and initial mass of Monolithic 60DVB. b) Comparison of PHMS and 60DVB; AP: Aged Powder, FP: Fresh Powder, M: Monolith 17

PHMS,He,374mg

Temperature (°C) 100

300

500

700

900

1100

1300

1500

dwell 0

a)

Intensity (A.U.)

-5 -10

Mass loss (%)

M/z=26 M/z=28 M/z=44 M/z=59 M/z=73 M/z=75

M/z=2 M/z=12 M/z=15 M/z=16 M/z=18

-15 -20 -25

0

20

40

60

80

100

120

140

Time (min) 250

300

350

M/z=2 M/z=18 M/z=29 M/z=43 M/z=45 M/z=59 M/z=75 M/z=73 M/z=30 M/z=28 M/z=26

450

400

59

b)

500

c)

550

600

43

650 M/z=2 M/z=26 M/z=28 M/z=29 M/z=30 M/z=31 M/z=42 M/z=43 M/z=44 M/z=45 M/z=58 M/z=59 M/z=73

2

44 45 43

73

Intensity (A.U)

1

Temperature (°C)

Temperature (°C) 200

-30

160

PHMS,He,374mg

59 58

Intensity (A.U)

75 18 2

45 28 26

15 14

73 75

26

31 42 30 29

M/z=75

28

18

30 15

20

25

30

35

40

40

45

50

55

60

65

Time (min)

Time (min)

Figure 3. TG/MS analysis for gases evolved during thermal degradation of monolithic PHMS in inert atmosphere helium, (a) represents a whole temperature range, (b) m/z of volatile species detected in the temperature range 185°C - 435°C, and (c) between 435°C and 685°C

60DVB, He,366mg

60DVB, He,366mg

Temperature (°C) 300

500

700

900

1100

1300

1500

a)

dwell

Intensity (A.U)

250 300 350 400 450 500 550 600

0

b)

91 92

-10 -15 -20

20

40

60

80

100

Time (min)

120

140

160

He 60DVB

18

-30

28 30

M/z=2 M/z=26 M/z=28 M/z=30 M/z=43 M/z=44 M/z=45 M/z=59 M/z=73 M/z=75 M/z=91 M/z=92 M/z=104 M/z=106

59 106 104 44 78

73 45

-25

0

2 26

-5

Mass loss (%)

M/z=2 M/z=12 M/z=15 M/z=16 M/z=18 M/z=26 M/z=28 M/z=44 M/z=59 M/z=73 M/z=75 M/z=91

Intensity (A.U)

100

Temperature (°C)

75

M/z=78

43

20 7

25

30

35

40

45

Time (min)

50

55

Figure 4. TG/MS analysis for the gases evolved during thermal degradation of monolithic 60DVB in inert atmosphere helium (a) represents a whole temperature range, (b) m/z of volatile species detected in the temperature range 235°C - 635°C

60

65

7b

18

Figure 5. SEM micrographs of PHMS pyrolyzed at 1450°C in Helium (Left) surface view of the sample; top inset shows higher magnification of area between circular pits, (right) higher magnification of circumferential crack-like voids

as CH(Si-CH3)

C-H, ring

s

 CH (Si-CH3)

Si-H

ring C=C,

CH2, scis

C-H, ring

33

3

sSi-O-Si

 Si-O (Si-OH)  C=C, ring breath

 Si-C(Si-(CH ) )

 Si-C (Si-CH )

Raman Intensity (A.U.)

Figure 6. SEM micrographs of 60DVB pyrolyzed at 1450°C in Helium: (Left) top view of the sample; top inset shows higher magnification of area between circular pits, (right) higher magnification of cavities.

60DVB

19  CH3(Si-CH3)

s

sCH3(Si-CH3)

PHMS

400

800

1200

1600

2000

-1 Raman Shift (cm )

2400 2800

3200

Figure 7. Raman spectra acquired from the two starting cross-linked polymers

b)

700°C

2173

Raman Intensity (A.U.)

2125

SiC

2235 600°C 550°C 450°C PHMS

3

2000

600°C

3 3

2100

SiC

2200

2300

2400

1

550°C

1

450°C

1

SiC

Raman Intensity(A.U.)

a)

1450°C,150',white spot 1450°C,150',black region

D G 1500°C 1450°C,30' 1350°C 1000°C 900°C

3

400

800

1200

850°C 800°C

1

PHMS

1600

2000

-1 Raman Shift (cm )

2400 2800

3200 400

800

1200

1600

2000

2400

2800

3200

-1 Raman Shift (cm )

Figure 8. Raman spectra acquired from PHMS after heating at different temperatures

a

b

C

20µm

e

Low Crystalline SiC

h

f

High Crystalline SiC

Disordered sp2 Carbon

i Intensity (A.U)

Intensity (A. U)

g

Intensity (A.U)

d

20

Figure 9. (a) Optical micrograph and (b) SEM image of the PHMS heat-treated at 1450°C for 150 min where Raman mapping has been performed. (c) magnified SEM image corresponding to the square shown in b. Raman map scores and their corresponding reference spectra showing the spatial distribution of the various components: (d, g) low crystalline SiC, (e, h) high crystalline SiC and (f, i) disordered sp2 carbon. o

670 C

b)

o

600 C o 550 C

Raman Intensity (A.U)

300°C 60DVB o

500 C

2000 2100 2200 2300 2400

x1

A

o

400 C

3

3

G D'

2D

SiC

400°C

3

D SiC

x1

x1

o

300 C

Raman Intensity (A.U)

a)

D+D'

1500°C, white region

1500°C, black region

1450°C 1350°C 1250°C 900°C

3

400

x1

60DVB

800

1200

1600

2000

-1 Raman Shift (cm )

2400 2800

3200

800°C

400

800

1200

1600

2000

Raman Shift (cm-1)

2400

2800

3200

Figure 10. Raman spectra acquired from 60DVB after heating at different temperatures

21

b

e

g Intensity (A.U)

High crystalline SiC

700

800

900

Raman Shift (cm-1)

1000

f

h

well organized sp2 C

Intensity (A.U)

d

c

1000

1200

1400

1600

1800

Raman Shift (cm-1)

2000

i

disordered sp2 C

Intensity (A.U)

a

1000

1200

1400

1600

1800

2000

Raman Shift (cm-1)

Figure 11: (a) Optical micrograph and (b) SEM image of the 60DVB heat-treated at 1500°C where Raman mapping has been performed. c) magnified SEM image corresponding to the square shown in b. Raman map scores and their corresponding reference spectra showing the spatial distribution of the various components: (d,g) High crystalline SiC, (e,h) well organized sp2 carbon and (f,i) disordered sp2 carbon.

22

Intensity (A.U.)

Nucleophilic Free-radical attack of Si-OH mechanisms by Si-CH3 homolytic homolytic Redistribution cleavage cleavage reactions of Si-C/C-H of Si-H Fragmentation of Silanes

PHMS 60DVB H H D T (1/1) H TT TQ T

100

200

300

400

500

600

700

800

900 1000

Temperature (°C)

Figure 12. Evolution of methane in various Si-based precursors

23