Vaporization study of SiC and SiC–2 mol% SiO2 powder mixtures. Grain morphology changes at high vapor pressures under pumping

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Journal of the European Ceramic Society 32 (2012) 4407–4417

Vaporization study of SiC and SiC–2 mol% SiO2 powder mixtures. Grain morphology changes at high vapor pressures under pumping G. Honstein, F. Baillet, C. Chatillon ∗ Science et Ingénierie des Matériaux et Procédés (SIMAP associé au CNRS-UMR 5466, UJF/INP-Grenoble), Domaine Universitaire, BP 75, 38402 Saint Martin d’Hères, France Received 3 June 2012; received in revised form 13 July 2012; accepted 16 July 2012 Available online 11 August 2012

Abstract The heat treatment of SiC powders, SiC–2 mol% SiO2 powder mixtures, and bimodal SiC powder mixtures has been studied with a quadrupole mass spectrometer linked by capillary tube to a special heat treatment reactor. Silica release was monitored on the CO(g) vaporized flow and the samples were analyzed by Raman spectroscopy and Scanning Electron Microscopy after the experiments. The present study showed that silica release by vaporization – first step in heating processes – is needed before any SiC growth process could start. The second step involving active SiC oxidation conditions by the remaining oxygen was conducive to the growth of “neck-like” connections between SiC grains and growth process was observed in the 1273–1600 K range. When the CO(g) release decreased as a result of higher temperatures or longer treatment times, carbon precipitation at the SiC surface was observed as the third step in the mass loss process. © 2012 Elsevier Ltd. All rights reserved. Keywords: Quadrupole mass spectrometry; SiC bimodal mixtures; SiC–SiO2 powder mixtures; High temperature vaporization; SiC growth

1. Introduction The growth of SiC crystals and the consolidation of SiC parts is always performed at high temperature and consequently the main processes that occur in this temperature range are related to matter transport by the gas phase as already stated by Knippenberg1 from growth process study in relation with structural characterizations of SiC polytypes, and Kreigesmann and Jodlauk2 from microscopic observations and shrinkage studies. A preliminary thermodynamic analysis of these transport processes has been performed by Honstein et al.,3,4 dealing first with intrinsic matter flows due to pure SiC vaporization, and second with SiC–SiO2 mixtures since silica is either present either as a native layer at the surface of SiC powder grains or used as an additive in the processes. This thermodynamic analysis showed that the growth of SiC crystals is possible only from carbon in the SiC–SiO2 –C 3-phases domain or from SiC vaporizing with a bare surface – i.e. without any silica layer – that is under active

∗

Corresponding author. Tel.: +33 0476 82 65 11; fax: +33 0476 82 67 67. E-mail addresses: [email protected], [email protected] (C. Chatillon). 0955-2219/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.jeurceramsoc.2012.07.018

oxidation conditions in relation with residual oxygen in the environment (atmosphere, containers). Under these conditions, the growth is in competition with an erosion process by vaporization as studied by Honstein et al.,4 and the best growth conditions are obtained for SiC crystals the non-stoichiometric composition of which being close to the SiC–C phase limit. However, erosion rate remains greater than growth rate and SiC cannot be obtained with the theoretical density but necessarily with a significant level of porosity. This thermodynamic study was followed by a Knudsen Cell Mass Spectrometric study in which Honstein et al.,5 observed the conditions of vaporization of either “pure” SiC powders and also pre-oxidized SiC powders in the low vapor pressure domain (<10 Pa). These experiments were performed as steady-state vaporizations under a very small controlled and known matter loss (the effusion process) and without any temperature gradients that could favor any growth processes by matter transport. Notwithstanding these restricted conditions, this experimental study confirmed that the growth of SiC was not possible as long as a silica layer was present at the surface of SiC grains. Finally the growth of pure SiC – observed as bridges or “necks” between grains – was possible under the only active oxidation conditions up to 1600–1700 K and the competitive erosion

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Fig. 1. Device for monitoring by quadrupole mass spectrometry the non-condensable gases extracted from the high temperature furnace via a capillary tube. TMP is for turbomolecular pump; PVP is for primary vacuum pump. The quadrupole mass spectrometer is a Pfeiffer QMG 422.

by vaporization finally prevented any densification process as previously thermodynamically analyzed by Honstein et al.4 The present experimental paper explores the vaporization of larger samples in upper transient pressures range and under transient heat conditions in order to favor any matter transport between SiC grains by the escaping gas phase. Indeed, these temperature and concentration gradients are close to the real conditions in industrial furnaces where a carrier gas flow as well as temperature gradients due to thermal shields and cold walls cannot be avoided. In order to obtain some chemical references in the vaporization process, a quadrupole mass spectrometer is coupled with a vaporization reactor fitted with a capillary tubing to observe the CO(g) escaping from these large samples. As in the previous Knudsen Cell Mass Spectrometric study, the heat-treated powder samples were characterized before and after treatment by grain size analysis, X-ray diffraction, Raman spectroscopy and Scanning Electron Microscopy with Field Emission Gun (SEM-FEG) as already described by Honstein et al.5 2. Experimental set-up For partial vapor pressures higher than 10 Pa in a high temperature reactor, collisions occur in the sampling orifice or in the molecular beam generated in vacuum. Consequently, the Knudsen cell method as used previously by Honstein et al.,5 is no longer available to determine the vapor pressures in the reactor. For the experiments described here, with SiC containing SiO2 , and when condensation of SiO(g) vapor occurs at the reactor outlet, it is still possible to observe the remaining permanent CO(g) gas flow by introducing this flow into a spectrometer through a capillary tube. For this reason, a quadrupole mass spectrometer fitted with a capillary tube was used in order to attain a higher pressure range

in the transient vaporization regimes of the powders than when using an effusion method. The aim was to monitor the release of CO(g) during different incremented temperature plateaus or different temperature ramps and for sample masses at least 10 times greater than those in the previous study. Furthermore, measurements taken with a quadrupole mass spectrometer are faster and enable continuous monitoring of a vaporization process with greater gas release and faster mass loss compared to the Knudsen cell device described in the previous study. 2.1. Instrumentation description The experimental device contained a vacuum housing in stainless steel fitted with an inductive high frequency (HF) coil for heating the reactor, a HF generator, a primary vacuum pump for the reactor vessel and a Pfeiffer quadrupole mass spectrometer (QMG 422, 1–100 a.m.u. range) with its own pumping assembly. The reactor and mass spectrometer were connected by a stainless steel heated capillary tube. The pressure between the end of the capillary and the mass spectrometer input was controlled by a gas dosing valve fitted with its own pumping device. Fig. 1 shows the experimental arrangement and Fig. 2 shows the reactor details. The crucible, crucible cover and capillary connecting parts were made from graphite and the various thermal insulation shields from graphite wool. To prevent diameter reduction or clogging of the capillary due to vapor condensation (mainly SiO(g)), the crucible cover is machined as a cone and this cone is located in such a way that it is always cooler than other crucible zones due to the local temperature gradient at the outlet of the heating coil upper level. Black glass deposits were indeed observed systematically on the inner surface of the cone and in the connecting parts. The capillary tube is connected orthogonally to the top connection of the conical cover. The groove

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Table 1 Degussa Sipernat 350 SiO2 characteristics.

Fig. 2. 1: Crucible; 2: conical cover; 3: connection to the capillary; 4: graphite cylindrical ring for HF coupling; 5: graphite wool; 6: graphite spacer; 7: HF coil; 8: bottom plate; 9: holder; 10: grove for positioning; 11: capillary tube connection; 12: pyrometric sighting.

machined at the bottom of the crucible helps to locate it at the center of the coil and to stabilize it on the holder plate. Due to deposits and formation of layers on the crucible walls in the course of the experiments, each individual experiment was made comparable with the other experiments by using a new reactor assembly including a new crucible for each experiment. The sample is located in a graphite crucible inside the reactor and this crucible is heated by radiation from the reactor graphite walls. The experiments could be run with either a primary vacuum or a protective gas (Ar) atmosphere. Depending on the chosen temperature ramp, a characteristic vapor phase is established in the crucible by sample vaporization. Only the non-condensable gaseous species were sampled by the capillary, usually water vapor at the very beginning of heating and CO(g) throughout the heating process. • The samples were heated by induction heating. The work coil was coupled with a graphite susceptor wall that radiated the heat through the crucible to the sample. The power supply was obtained by a HF generator (12 kW). • The temperature is measured with a two-color pyrometer (IRCON Modline 5) by sighting through a porthole directly on the cone through a small hole machined in the insulation graphite wool. The pyrometer readings were initially calibrated to obtain the difference between the established temperature inside the crucible and the temperature measured during the experiments on the cone: an experiment without the SiC powder sample and a second pyrometer was performed: instead of the capillary, a second pyrometer was looking directly inside the crucible on its bottom surface. A constant temperature of Tcone +38 K (constant gradient) was measured inside the closed crucible in the 1150–1650 K range. • The gases from the reactor enter the ionization chamber through a 1 m long capillary tube of 0.15 mm inner diameter. The pressure in the reactor can be monitored within the 104 to 1.2 × 105 Pa range. In the present experiments,

d50 /␮m

Na2 O/wt%

Fe2 O3 /wt%

SO3 /wt%

4.5

0.8

0.03

0.2

the reactor vacuum pump provided a primary vacuum in the reactor housing. The pressure of the permanent gases from the vaporization of the sample is continuously reduced in the capillary until a value of 1 mbar is obtained at the gas dosing valve inlet. The gas flow from the furnace via the capillary is pumped downward by an independent pumping system. Most of this flow is thus eliminated from the mass spectrometer input. The flow rate of the capillary is 2 sccm (standard cubic cm per minute) for 1 bar input gases and the residence time is about 0.3–1 s depending on the gas viscosity. In order to compare experiments with each other, the total conductance of the sampling system must remain the same through the whole experiment series. For this reason the gas dosing valve was completely open in each experiment. The samples had a mass of about 15–18 g and the capacity of the crucibles was adapted to the expected gas flow during a high temperature experiment. 3. Mass spectrometric observations 3.1. Samples The 0.5 ␮m diameter (“pure”) SiC sieved powder from StGobain and the Degussa Sipermat 350 SiO2 were used. The purity of the SiC powders was analyzed and published in a previous paper by Honstein et al.5 The oxidized bimodal SiC mixture was made from ≈33 mol% 0.5 ␮m diameter and ≈67 mol% 10 ␮m diameter sieved grain size powders and heat treated in air at 500 ◦ C for about 30 h. The characteristics of the SiO2 powders are presented in Table 1. 3.2. Experiments performed The set of performed experiments is presented in Table 2. A first set of experiments was performed by setting the temperature manually according to different temperature plateaus and minimizing the intermediate temperature jump time, while the second set was performed using a temperature ramp. The first steps of the heating processes were almost the same, • 0.5 h from room temperature to 1053 K • 1 h from 1053 to 1273 K • 1 h from 1273 to 1373 K, except for the “200” sample The main species observed were effectively the noncondensable CO(g) (at mass 28 a.m.u.) and sometimes, at the very beginning of the experiments, a small contribution of water (at mass 18 a.m.u.) was observed but this peak quickly disappeared. Other vapors as predicted by thermodynamics (see

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Table 2 Quadrupole mass spectrometric measurements performed with capillary sampling. Experiment label

Sample

T plateau or T maximum/K

Treatment time or end conditions

16-10 19-10 26-10 06-11 08-11 06-12 23-10 13-11 28-11 25 50 100 200

0.5 ␮m diameter SiC powder

1600 1700 1500 and 1690 1500 and 1740 1500 1650 1500 and 1600 1529 1560 1473 1473 1473 1473

1h 1 h 10 min 2.5 h and 1.5 h 1.5 h and 2 h No CO(g) observed, 40 min p(CO) maximum 1.5 h and 1.5 h No CO(g) observed, 45 min p(CO) maximum 25 K/h, 1273 < T < 1473 K 50 K/h, 1273 < T < 1473 K 100 K/h, 1273 < T < 1473 K 200 K/h, 1273 < T < 1473 K

0.5 ␮m diameter SiC powder + 2 mol% SiO2 mixture

Oxidized SiC bimodal powder mixture

Honstein et al.5 ) condensed as black glasses in the top of the cover before being able to enter (and clog) the capillary tube. 3.3. Manually adjusted temperature plateaus In the first heating steps, the CO(g) flow was detected slightly at 1053 K, but more clearly when reaching 1273 K. The CO(g) peak increased considerably for each rapid temperature rise, then decreased quickly as soon as the temperature stabilized, as shown for instance in Fig. 3. The CO(g) peak flow decreased for successive temperature jumps as did the final CO(g) flow (and consequently its pressure in the reactor) at the end of the different temperature plateaus and this behavior was the same for both “pure” SiC powders and mixtures with silica. Taking into account the preceding thermodynamic study and the experimental mass spectrometric observations by Honstein et al.,5,6 the present behavior is related to the existence of a silica layer at the surface of the SiC grains. The CO(g) peak is associated with the vaporization of mainly CO(g) and SiO(g), with the SiO(g) condensing at the outlet of the reactor (conical cover) before it is able to enter the capillary. Thus the CO(g) gas flow measured by the mass spectrometer was the observed “tracer” of the total vaporization of the silica layer.

Fig. 3. Decimal logarithm graph of the as-measured CO+ ionic intensity with corresponding temperature (experiment 06-11) during the heat treatment of a SiC–2 mol% SiO2 powder mixture as a function of experiment duration.

The heat treatment of different samples – “pure” SiC, SiC–2 mol% SiO2 mixture and oxidized bimodal SiC mixture – are displayed in Fig. 4. The different measured ionic intensities were normalized by reference, via the mass spectrometric sensitivity ratios, to the background peak at mass 28 (N2 ) as monitored before heating, and related to the total measured pressure in the mass spectrometer housing (ionization gauge and total intensity in the mass spectrometer). The height of the CO(g) peak of “pure” SiC powder is very low compared to the two other samples, while the CO(g) peak for the oxidized bimodal SiC mixture is the largest. The peak for the SiC–SiO2 mixture indicated that the mixture takes longer to release its silica content. The peak area is proportional to the initial silica content – the smallest area being for “pure” SiC – while the shape of the peak is related to the way silica was evaporated. So it would appear that the oxidized bimodal SiC powder mixture contained more than 2 mol% SiO2 . However, the lower release rate of CO(g) for the SiC–2 mol% SiO2 mixture is worthy of note and indicates that the silica grains served as an oxidation tank for the SiC grains (maintaining for a longer time the silica layer at the surface of SiC). As a result,

Fig. 4. Ionic intensity of CO+ during the vaporization of “pure” SiC powder, SiC–2 mol% SiO2 powder mixture and pre-oxidized SiC bimodal powder mixture as a function of time. The sensitivities of the different experiments were normalized using the nitrogen peak and total pressure measured in the vacuum background before starting the heating process.

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Fig. 6. Breakdown of the experimental curve obtained for the CO(g) release intensity of pre-oxidized SiC bimodal mixtures showing the three contributions: fine grain vaporization, coarse grain vaporization and residual CO(g) from a tank source.

• A shoulder appeared at the decrease in CO(g) release, and this shoulder may be due either to thermal diffusion inside the sample or to some chemical change of the sample.

Fig. 5. Heat treatment of pre-oxidized SiC bimodal mixtures: (a) ionic intensities of CO+ normalized to their maxima as a function of the heating time for different temperature ramps (200, 100, 50 and 25 K/h) from 1273 to 1473 K and (b) zoom on the shoulder for the slow temperature ramps.

the mean oxygen potential in the reactor decreases at a slower rate. A slight bump was also observed in the rapid decrease of the peak, probably indicative of different regimes for the CO(g) release or different vaporization processes. 3.4. Heating with temperature ramps The CO(g) peaks measured with fixed temperature ramps applied to pre-oxidized SiC bimodal mixtures are displayed in Fig. 5 for different temperature rise rates between 1273 and 1473 K. The shapes of the peaks were different for the different rates, and the characteristic temperatures that showed regime changes are presented in Table 3. A residual amount of CO(g) – sufficient to be detected – was observed for all samples even 1.5 h after the beginning of heat treatment at 1273 K. The observation of release regime changes can be related to the heating rates by two different features: • The maximum release occurred at the peak for the vaporization rates of the mixtures: the faster the temperature ramp, the greater the initial silica loss at the peak.

As already illustrated in Fig. 3, the CO(g) flow increase occurred systematically up to temperature stabilization at the 1473 K plateau and this effect can be attributed to the only vaporization of the sample, i.e. of SiC with its SiO2 layer. The CO(g) release then decreased rapidly indicating that the silica layers started to disappear at least for SiC grains close to the surface of the powder bed and a shoulder was observed that became more pronounced for the slow temperature ramps. The influence of temperature ramps on sample vaporization behavior was caused by the heat zone expansion and temperature distribution inside the powder bed which had intentionally been made relatively large in this reactor. When heating a low thermal conductivity sample from the walls in contact with the coupling coil, the heat zone moves from the rim (at the crucible walls) to the core of the sample – and similarly toward the powder bed surface – according to a smaller temperature gradient for slow heating ramps. It was then possible to observe more clearly the distribution of vapors between different grains. As in previous experimental work by Honstein et al.,5 a lower net vaporization rate for larger grains was observed, presently larger grains reach the same vaporization flow at a different temperature compared to small grains. As a result, CO(g) release may take longer for large grains compared to small grains. The observed CO(g) peak was broken down as shown in Fig. 6 taking into account two different vaporization rates, as well as a background component. Fig. 6 clearly shows the two peaks due to the release of CO(g) by the bimodal grain distribution and a third peak (noted “other”) that probably corresponds to the residual CO(g) due to the active oxidation regime after silica release, as noted in the previous experimental observations by Honstein et al.5

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Table 3 Observed temperature regime changes for CO(g) release during the capillary mass spectrometric experiments performed during the heating of pre-oxidized bimodal SiC mixtures with different temperature ramps between 1273 and 1473 K. Sample and temperature ramp/(K/h)

Start of release/K

Maximal release/K

End of release/K

Shoulder/K

200 100 50 25

1273 1273 1273 1273

1347 1361 1410 1422

1446 1461 1473 1473

1367 1367 1430 1450

3.5. Conclusion The CO(g) release was observed as soon as the samples reached 1273 K, and this CO(g) release is directly related initially to silica vaporization and then to silica depletion of the samples. The gradual temperature increase has different effects: • A CO(g) peak due to primary vaporization of SiC and SiO2 in contact at the interface SiC grain/SiO2 layer. • At each temperature rise step, the SiO2 reserve remaining in the still unheated sample regions, due to temperature gradients, is mobilized. • In the final temperature rise step the CO(g) is observed to decrease and this proved that silica has been completely released. • The remaining oxygen (stored in the crucible material or furnace) contributed to active oxidation of the SiC grains. Time is also an important factor for the outgoing oxygen flow that leads to silica elimination. 4. Characterization results The heat treated powders in the furnace coupled to the quadrupole mass spectrometer through the capillary tube were analyzed by Raman spectroscopy and FEG-SEM as already explained in the previous mass spectrometric study by Honstein et al.,5 and the original powders were those analyzed in this previous study. 4.1. Powders heat treated at different temperature plateaus 4.1.1. Raman spectroscopy Raman spectroscopy is usually used to identify the proportion of different SiC polytypes (4H, 15R, 6H etc.) when using the RAMAN diffraction pattern. With powders, the fluorescence of the photons emitted to eliminate the stored energy of the incident light in the grains caused a regular increase of the background line as a function of the frequency in the detected spectrum range contrarily to the observed flat and low level of the background line of a spectrum for a solid sample in which the incident energy could be easily dissipated into the bulk solid and into the holder. For powder grains, on the other hand, the contacts built between the grains could not produce the same effect. In case of any connections growth between the grains during the heat treatment, the background noise decreased. It is this effect that had been primarily retained in the present

Fig. 7. Raman spectra of “pure” SiC powder (experiment 16-10), SiC–2 mol% SiO2 powder mixture (experiment 26-10) and pre-oxidized bimodal SiC powder mixture (experiment 23-10) samples characterized after heat treatment using the capillary mass spectrometer.

study and not the probable variation in the different polytype proportions that were sometimes observed. Fig. 7 displays the Raman spectra for the “pure” SiC powder, the SiC–2 mol% SiO2 powders mixture and the oxidized bimodal SiC powder with zooms in different frequency regions. From the observations made, the following conclusions can be drawn: • The background level of the oxidized bimodal SiC powder mixture in the low frequency region is higher than the other two. This is an indication that silica amorphous material remains in this mixture. • The SiC–2 mol% SiO2 powder mixture in the high frequency region did not show any fluorescence while the other two powders showed an increase in fluorescence level. This may be indicative of interconnections built between the grains in the SiC–2 mol% SiO2 mixture. • All samples showed carbon content with peaks in the 1500–1700 cm−1 region. The heating treatment could have been extended far beyond the active oxidation regime, resulting in carbon precipitation. • In the low-frequency region, the RAMAN peaks for “pure” SiC and the mixture SiC–2 mol% SiO2 were similar. These peaks corresponded to the 6H-SiC polytype, except the two peaks for 4H (195 cm−1 ) and 15R (256 cm−1 ) for oxidized powders. • In the high frequency region, the shoulder at 789 cm−1 for the only SiC–2 mol% SiO2 sample corresponds theoretically

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to the 6H, at 938 cm−1 to the 15R polytype. However, the different polytypes usually present in this region overlap and this made it impossible to study these fluctuations. • Comparing the low and high frequency regions, it is clear that the SiC–2 mol% SiO2 sample lacks the 15R SiC type. From the previous thermodynamic analysis and experimental work by Honstein et al.,3–5 the observed variations in CO(g) flow during the heat treatment procedure showed that the different samples were first de-oxidized (silica release) then maintained in the active oxidation steps the duration of which were different depending on the sample. These steps correspond to different and varying Si/C gas phase ratios and consequently could produce the growth of different polytypes as already proposed qualitatively by Knippenberg.1 4.1.2. FEG-SEM analysis The three above powder samples were analyzed by FEGSEM and the results are presented in Fig. 8. The agglomerates of the “pure” SiC powder sample observed in Fig. 8a had already been observed in the initial powder. No grain-to-grain connections were found, thus in agreement with the high level of the Raman fluorescence observed for this sample. The temperature and duration of the heat treatment at 1600 K were not sufficient to promote the growth of “neck-like” bridges. The best connections were observed to form for the SiC–2 mol% SiO2 powder after the 26-10 experiment (Fig. 8b) in agreement with the lack of fluorescence level: this sample was heat-treated at 1500 K for 2 h and at 1690 K for 1.5 h. In addition, surface energy minimization was observed with the formation of facets. This sample in fact remained in the active oxidation regime for a longer time – as revealed by the lower CO(g) flow decrease in Fig. 4 – and this situation with bare SiC surface favored the growth process as already explained by Honstein et al.4 For this sample containing initially silica grains the growth process may have probably started during the 1500 K plateau, and this was followed by a competition stage, with surface diffusion reorganization as well as etching by vaporization at 1690 K. Fig. 8c shows that the fine grain fraction of the oxidized bimodal SiC sample was still present, and some interconnections were built between small and large grains. The heat treatment of this sample (experiment 23-10) was the same as for the SiC–2 mol% SiO2 powder, but the few connections built were not sufficient to prevent Raman fluorescence. However, the silica content of this sample was probably higher than for the SiC–2 mol% SiO2 powder when looking at the total release of CO(g) in the mass spectrometric experiment. Indeed the silica existing as an only layer did not impose a longer time in the active oxidation regime as shown in Fig. 4 than independent silica grains. The analysis of different treatments of the samples would appear to confirm these growth mechanisms. Fig. 9a shows the SiC–2 mol% SiO2 powder after heat treatment at 1500 K for 40 min and Fig. 9b the oxidized bimodal SiC mixture after heat treatment at a maximum temperature of 1529 K for 45 min. The

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grains of the SiC–2 mol% SiO2 powder have clearly built up necks between each other. Furthermore, they have facets which indicate either surface diffusion to minimize the energy or etching by vaporization. Fig. 9b shows that the oxidized bimodal SiC mixture did not build any necks, but contained only rounded grains. As these two samples were treated similarly and the experiment was stopped once all CO(g) had disappeared – i.e. after all the silica has been released – the reason for the difference in growth process must be due to the distribution of the initial silica in the sample: layers on SiC or independent grains. Fig. 10 shows the same samples, SiC–2 mol% SiO2 powder (experiment 06-12) and oxidized bimodal SiC mixture (experiment 28-11) after heating was stopped just after the maximum CO(g) peak was obtained. This occurred at 1650 K for the SiC–2 mol% SiO2 powder, and 1560 K for the oxidized bimodal SiC mixture. The SiC–2 mol% SiO2 powder presented necks and facets, while the oxidized bimodal SiC mixture did not show any change when compared to the initial samples. The 90 K difference in temperature may explain the presence or absence of observed growth (depending on the growth rate) but this was probably not the main reason since the two samples reached the same vaporization state at the observed CO(g) peak. The main reason is probably that, in the oxidized mixture, the larger grains do not lose their silica layer at the peak summit due to lower vapor pressure as determined in the previous Knudsen Cell Mass Spectrometric experiment by Honstein et al.,5 and consequently the loss of silica occured more continuously and more slowly. For the SiC–2 mol% SiO2 powder the small SiC grains (or part of them) lose first their silica layer due to higher vapor pressure – and the growth process could start under the active oxidation conditions maintained by vaporization of the residual silica grains. Remind that the sample is a large sample and it is not necessarily homogeneous in terms of composition and temperature: grains close to the surface become depleted first with silica. Fig. 11 shows the SiC–2 mol% SiO2 powder (experiment 0611) treated at 1500 K for 1.5 h, then 1740 K for 2 h. At the end of the heat treatment the CO(g) flow decreased below the detection limit. Fig. 11a shows a group of interconnected grains (left) with necks and facets, while the other grains look different, well interconnected but with numerous facets. The Electron Back Scattering (EBS) image (chemical analysis) in Fig. 11b shows that the first group remains as pure SiC grains while the second group is carbon covered. The connections were probably built at different times, and the second group produced carbon by vaporization under excessively low oxygen pressure conditions (very low CO pressures) as already explained for pure SiC by Honstein et al.3,4 The two groups probably came from two different locations in the reactor crucible. 4.1.3. Conclusions Initial SiO2 content was necessary to prepare the conditions for SiC growth in the form of interconnections between grains, but as long as the SiC remained covered by a silica layer the growth process could not start. The primary heating step enhanced the cleaning of the SiC surface by vaporizing the silica layer, and this step prepared the next active oxidation

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Fig. 8. FEG-SEM images of: (a) “pure” SiC powder sample after 16-10 experiment (Tmax = 1599 K), (b) SiC–2 mol% SiO2 mixture sample after 26-10 experiment (Tmax = 1700 K), (c) pre-oxidized bimodal SiC mixture sample after 23-10 experiment (Tmax = 1600 K).

step by accumulation of the necessary residual oxygen through the SiO(g) and CO(g) species in the reactor. The active oxidation regime was established for a longer time when starting with independent silica grains, i.e. mixtures of SiC–SiO2 . This was better than an equivalent pre-oxidation of the SiC grains. After the active oxidation regime, due to the final release of oxygen-containing species by SiO(g) and CO(g), the oxygen potential decreased and carbon precipitation at the SiC surface was observed. In this final step the vaporization is really occurring for pure SiC i.e. in the Si–C binary system with total lack of oxygen. At that time the vaporization becomes non-congruent as already analyzed by Honstein et al.3 : the vapor phase is

richer with silicon than with carbon and this last component precipitates at the surface. 4.2. Powders heat treated with different temperature ramps 4.2.1. Raman spectroscopy Fig. 12 compares the same oxidized bimodal SiC powder mixture heated in the same temperature interval with different temperature ramps between 1273 and 1473 K. In all samples the carbon peak appeared at values > 1500 cm−1 . The fluorescence levels were different but no regular trends were observed. Zooming the low frequency region, it can be seen that the 4H

Fig. 9. FEG-SEM images of: (a) SiC–2 mol% SiO2 mixture sample after 08-11 experiment (Tmax = 1500 K for 40 min) and (b) pre-oxidized bimodal SiC mixture sample after 13-11 experiment (Tmax = 1529 K for 45 min).

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Fig. 10. FEG-SEM images of: (a) SiC–2 mol% SiO2 mixture sample after 6-12 experiment (Tmax = 1650 K) and (b) pre-oxidized bimodal SiC mixture sample after 28-11 experiment (Tmax = 1560 K).

Fig. 11. FEG-SEM images of a SiC–2 mol% SiO2 mixture sample after 06-11 experiment (Thold = 1500 K for 1.5 h and 1740 K during 2 h): (a) normal imaging mode and (b) back scattered electron imaging mode detection showing the graphite layer on the surface of the group of grains on the right.

Fig. 12. Zooms from the Raman spectra of the pre-oxidized bimodal SiC powder mixture sample from experiments performed with different temperature ramps.

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Fig. 13. FEG-SEM images of pre-oxidized bimodal SiC powder mixture samples from experiments performed with different temperature ramps up to 1473 K: (a) 200 K/h, (b) 100 K/h, (c) 50 K/h, and (d) 25 K/h.

SiC polytype was no longer present for the lowest temperature ramp, i.e., 25 K/h. As this sample remained for a longer time in the 1273–1473 K range, the slower change in the gas phase Si/C ratio might have favored the formation of other polytypes detrimental to the 4H type. 4.2.2. FEG-SEM observations Fig. 13 shows microscopic observations. Connections between small grains as well as between small and coarse grains were observed together with facets but no significant differences due to the temperature ramp. Connections of similar quality were obtained in experiment 23-10 performed with two temperature plateaus at 1500 K and 1600 K for 1.5 h, a time comparable with the present ramp times going from 0.5 to 4 h but with a maximum temperature of 1473 K. Graphite was not observed at the surface of the grains meaning that growth occurred mainly in the active oxidation regime. 4.2.3. Characterization conclusion The effect of different temperature ramps in the 1273–1473 K range and for the two successive regimes – silica release and active oxidation of SiC – did not influence the growth of connections (necks) between the grains of a bimodal SiC mixture. The only observed effect is the lack of 4H SiC type growth for the slower temperature ramp (25 K/h). 5. General conclusion The present work was undertaken following the preceding work performed by Knudsen cell mass spectrometry with the

aim of enhancing the growth processes of necks between SiC grains using higher transient vaporization pressures as well as temperature ramps and gradients in conjunction with larger samples. For these conditions some composition gradients existed in the samples in terms of silica local contents. The primary step in the different heat treatments of “pure” SiC, SiC–2 mol% SiO2 or bimodal SiC pre-oxidized powders was the release of silica to ensure that the SiC grain surfaces became clean. The two main species involved in this release were SiO(g) and CO(g). In the present work where the vaporization flow was sampled by capillary tube, SiO(g) was first condensed at the exit of the vaporization reactor and the only gaseous species that could be observed was CO(g) serving as a tracer of the vaporization reaction steps. The different steps in the CO(g) release were identified: – first the vaporization of silica, then – second the transitory state of active oxidation of SiC and – third and final step the carbon precipitation at the SiC surface for very low CO(g) pressures – mainly for values lower than our detection limit. The growth of SiC – observed as necks connecting the different SiC grains – occurred only under active oxidation conditions just after the removal of the silica layers that normally exist on the SiC grain surface as native or as resulting from a preceding passive oxidation step due to the presence of independent silica powder added initially in the SiC powder. The building of necks between grains was observed sooner for small grains than for coarse grains because the silica release was systematically slower for coarse grains in agreement with the partial pressures as determined in the preceding Knudsen Cell Mass Spectrometric experiment.

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The distribution of silica in the sample either as a layer – native or created – at the SiC surface or as independent grains in a SiC–SiO2 mixture influenced the duration of the next active oxidation step and consequently the time available for the SiC growth process by vaporization and condensation processes under this regime. In case of large samples, and for a limited vaporizing open surface, the bottom part of the powder bed probably serves as a silica reservoir that imposes active oxidation conditions to the silica depleted bed surface layer. The available time for growth is also related to the crossing of the active oxidation window depending on the oxygen release rate (or oxygen partial pressure decrease in the powder bed). As the active oxidation window is crossed with different rates of decrease in oxygen potential due to the dynamics of nonclosed industrial heating furnaces or laboratory cells or reactors, the SiC growth process is usually performed under variable silicon to carbon potentials – or Si/C gas phase composition ratios – and different SiC polytypes could be grown successively. This change in the chemical environment during growth could explain the variations in polytypes proportions observed in the present experimental work. The last step is the carbon precipitation at the surface of the SiC grains (or growth of graphite crystals) due to too low oxygen pressures (or CO(g) pressures), which also signals the beginning of pure SiC etching or erosion. The observation of facets is

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probably related to the approach of this erosion regime since it was observed to occur systematically with neck building. Acknowledgments The authors would like to thank the St-Gobain CREE Research Centre at Cavaillon (France) for sponsoring the present study. References 1. Knippenberg WF. Growth phenomena in silicon carbide. Eindhoven, Netherlands: Philips Gloeilampenfabrieken; 1963. p. 161–274. 2. Kreigesmann J, Jodlauk J. Characterizing the consolidation of bimodally distributed fine grained silicon carbide powders, cfi/Ber. DKG 2002;79: 37–44. 3. Honstein G, Chatillon C, Baillet F. Thermodynamic approach to the vaporization and growth phenomena of SiC ceramics. I. SiC and SiC–SiO2 mixtures under neutral conditions. J Eur Ceram Soc 2012;32:1117–35. 4. Honstein G, Chatillon C, Baillet F. Thermodynamic approach to the vaporization and growth phenomena of SiC ceramics. II. The SiC surface under oxidative conditions. J Eur Ceram Soc 2012;32:1137–47. 5. Honstein G, Chatillon C, Baillet F. Vaporization study of mixed SiC powders. Partial pressures and grain morphology changes under vacuum conditions. J Eur Ceram Soc 2012;32:3851–60. 6. Honstein G, Chatillon C, Baillet F. High temperature mass spectrometric study of the vaporization behaviour of SiC–SiO2 system. J Alloys Compd 2007;452:85–8.