Olivier Ledain

Jacques Sylvain
email: jacques@lcts.u-bordeaux1.fr

Laurence Maillé

Consolidation and conversion of carbon powders into TiC by Reactive Chemical Vapour Infiltration

Keywords: carbide, RCVI, hybrid process, powder consolidation, self-limitation

Ceramic samples were prepared using a hybrid process in which the ceramic powder route was combined with Reactive Chemical Vapour Infiltration, a new gas phase route. In this technique, a carbide growth occurs from the conversion of a carbonbearing powder and slows down with increase in carbide thickness due to solid-state diffusion limitation. This self-limitation is expected to allow a self-regulation of the growth between the interior and the surface of the sample and thus a better homogeneity. The chosen carbide was TiC obtained by reacting H 2 and TiCl 4 with a carbon powder introduced by slurry impregnation in felts. At 950°C, the carbide infiltration is homogeneous with depth but the powder is not consolidated, while at 1050°C the better conversion rate of carbon into carbide allows the consolidation but the infiltration is limited to only about 250 µm. The consolidation depth can be improved with time-temperature steps between both temperatures.

Introduction

Among the different methods used to fabricate ceramic matrix composites (CMCs), chemical vapour infiltration (CVI) allows carbide-based matrix to be obtained at moderate temperatures (900-1100°C) with high purity and without shrinkage in woven fibre preforms of different shapes and sizes [START_REF] Naslain | Design, preparation and properties of non-oxide CMCs for application in engines and nuclear reactors: an overview[END_REF]. However, CVI is a slow and so a time consuming technique and the final CMCs display significant residual porosity. The matrices comprise thus macropores whose shapes and singularities cause stress concentrations responsible for early crack initiations during mechanical loading [START_REF] Guillaumat | Probabilistic-statistical simulation of the non-linear mechanical behavior of a woven SiC/SiC composite[END_REF]. The ceramic powder route, for its part, allows the formation of the ceramic matrix inside the fibre preform by slurry impregnation [START_REF] Nakano | Carbon fiber reinforced silicon carbide composites[END_REF] [4] [START_REF] Naslain | Design, preparation and properties of non-oxide CMCs for application in engines and nuclear reactors: an overview[END_REF]. This method is fast and low cost and, unlike CVI, avoids the formation of large pores at yarn intersections and between the plies. But after impregnation, the green powder parts have to be consolidated by hot pressing and sintering. The high temperature and mechanical pressure required for carbide sintering can lead to severe fibre degradation. Because of the shrinkage that occurs during the sintering, this method is limited to the preparation of unidirectional CMCs. Furthermore, the sintering additives can affect the final properties of the material. The use of a hybrid process in which the slurry impregnation is followed by the gas phase route would overcome the limitations of each route used separately. However, a submicrometer powder compact is difficult to infiltrate by conventional CVI in large enough thickness due to a clogging effect. The gaseous precursor infiltrates in the fine porosity of the powder with difficulty. The carbide growth rate is lower at the interior of the powder than at the surface. This case is illustrated in Fig. 1. Because of the gaseous reactive species depletion within the powder, a pronounced deposition gradient is established around the particles and finally the seal-coating of the powder surface prevents further in depth infiltration. In this way, Fig. 2 shows a carbon powder compact after four hours of SiC infiltration by conventional CVI carried out in our laboratory. Despite the use of very low pressure (5 kPa) and temperature (800°C) conditions such as to allow low growth rate for maximum infiltration, the depth of SiC infiltration beneath the surface of the sample is only about 20 µm. A seal-coating is observed at the sample surface evidencing that the entrance of the material is closed and that deeper infiltration is not possible. Itoh et al. could overcome this difficulty by using pulsed CVI instead of conventional CVI [START_REF] Itoh | Pulsed chemical vapour infiltration of SiC to three-dimensional carbon fibre performs[END_REF]. They could prepare in their laboratory CMCs from fibre preforms partly filled with SiC fine powder and then infiltrated with silicon carbide by pulsed CVI at a temperature below 900°C and a total time per pulse of only 1.1-2.0 s. However, the pulsed method is not suited to CMC manufacture on a large scale because the periodic gas evacuation in a few seconds of large industrial CVI reactors does not seem feasible. The aim of this study is to develop a hybrid process that combines the advantage of slurry impregnation with a new gas phase route called reactive CVI (RCVI). RCVI is similar to reactive chemical vapour deposition (RCVD) but transferred to the infiltration and consolidation of porous media. In RCVD, the MC carbide coating is obtained by reacting a carbon-bearing substrate with the gas, which carries only the M element (being B, Hf, Ta, Si, Ti…) but not the C element like in classical CVD [START_REF] Takahashi | The chemical vapour deposition of titanium carbide coatings on iron[END_REF] [7] [8] [9] [10] [START_REF] Baklanova | Protective ceramic multilayer coatings for carbon fibers[END_REF]. Once the substrate surface in covered with a continuous film of carbide, the continuation of the growth involves carbon diffusion from the substrate to the surface through the growing carbide film. So the growth slows down when the carbide thickness increases due to the limitation of solid-state diffusion. The carbide growth involves the consumption and the conversion of the original substrate. When RCVI is applied to the infiltration of a porous carbon-bearing substrate like a powder compact, i.e. without carbon in the original gaseous phase introduced in the reactor, the same gradient phenomenon as in conventional CVI is observed at the beginning of the process due to gaseous precursor infiltration difficulty (Fig. 3.a). But, because the carbon solid-state diffusion is limited when the thickness of the carbide coating increases, the RCVI process is expected to be self-limiting with self-balancing growth rates between the interior and the surface of the powder (Fig. 3.b). Thus, a homogeneous infiltration and consolidation should be obtained within the whole material (Fig. 3.c). In this work, titanium carbide has been chosen to study this concept for the first time because a mixture of titanium tetrachloride and hydrogen reacts easily with carbon to form TiC by RCVD and because carbon can diffuse through this carbide that can be sub-stoichiometric [START_REF] Vincent | Effet des conditions expérimentales sur la CVD réactive de couches minces de TiC sur carbone[END_REF] [12] [START_REF] Ledain | Reactive Chemical Vapour Deposition of Titanium Carbide from H 2 -TiCl 4 Gas Mixture on Pyrocarbon: A Comprehensive Study[END_REF].

Experimental

Synthesis

Study materials, mainly consisting of green powder compacts, were prepared from RVG-2000 ® carbon felt with a thickness of 6 mm in which the powder was introduced by slurry impregnation. The advantage of using this felt is that it contains a small amount of carbon fibres (6-10% in volume) compared to the fibre preforms commonly used for the reinforcement of actual CMCs. This allows the easy preparation of the samples and allows the study to be focused on the RCVI of large green powder areas that are free of fibres while having a local porosity similar to that obtained in actual green CMCs prepared by slurry route. A stable aqueous suspension with a concentration of 34 vol. % of carbon powder (Luvomaxx N991 ® carbon black, d 50 # 600 nm) and with a wetting and dispersing additive (DisperBYK 181 ® ) was prepared. Then, the pieces of felt were impregnated under vacuum for one hour. The solvent was evaporated by heating at 100°C under vacuum for one hour. Two cycles of impregnation / evaporation were applied to enhance the filling in the porosity of the felt. The reactor used for the RCVI experiments was the same CVD device as Jouanny et al. [START_REF] Jouanny | Synthesis of TiC from porous carbon coating on Si-C-O (Nicalon) fibres by reactive chemical vapour deposition in pressure-pulsed mode or at atmospheric pressure[END_REF]. It was a silica tube with an inner diameter of 6 cm. The silica tube was placed inside a conventional electrical resistive furnace. The green study materials were placed in the ~10 cm long hot area of the reactor and heated at a temperature ranging from 950°C to 1050°C. The H 2 /TiCl 4 gaseous mixture was introduced in the reactor with a total pressure maintained at a constant value of 5 kPa, a total flow rate equal to 300 sccm (standard cm 3 min -1 ) and with a dilution ratio [H 2 ]/[TiCl 4 ] equal to 3. In the same experimental conditions, several study material samples were placed together in the reaction chamber. The influence of time on infiltration was studied by stopping regularly the experiment and removing a sample after a few hours.

Characterization

Porosity (

) and pore size distribution of the green study materials were determined by mercury intrusion porosimetry (AutoPore IV, Micromeritics). Crystalline phases present in the samples were determined by X-ray diffraction patterns obtained in a Bragg-Brentano configuration (D8 advance, Bruker). Polished cross sections of samples were observed by scanning electron microscopy (SEM-Quanta 400 FEG V2 microscope) using a backscattered electron (BSE) detector and with an accelerating voltage fixed at 10 kV. After RCVI of various durations, the local apparent volume fractions of grown TiC ( ), remaining carbon ( ) and porosity (P) were determined at different depths in the sample by image analysis from the BSE-SEM images using a specially designed Matlab program. The used BSE-SEM images corresponded to an observed area of 15 µm wide and 13 µm high and were 1024 pixels wide by 884 pixels high. In a BSE-SEM image, the bright areas correspond to TiC which recognition is straightforward. The dark areas correspond to either residual carbon or porosity. The difficulty was to sort through the carbon particles of generally rather circular shape and surrounded by a layer of TiC and the pores of a more complex shape. The method to determine the various contents consisted in binarizing the image into pixels of either a black or a white colour using a threshold level between the bright and dark levels. The volume fraction of TiC was then the number of white pixels out of the total number of pixels in the image. The binary image underwent then a morphological closing through dilation followed by erosion of the white areas. This step was to remove blurred image areas, discontinuities in the TiC layers and thresholding artefacts. Each black area was then considered to be a carbon particle or a pore depending on a morphological criterion not reported here, but based on the roundness and the size. The contours of the sorted areas were drawn in blue for parts recognized as carbon and in red for parts recognized as pores in the original image in order to visually estimate the relevance of the results (Fig. 4). The volume fractions of carbon or pores were then proportional to the number of pixels of each kind of area. The given values result from the averages for different used dilation/erosion radii. Knowing the original porosity of the carbon study materials ( ), it was then possible to determine the apparent volume conversion rate between the volume of grown carbide and the volume of consumed carbon (Rv) after RCVI versus the depth of infiltration according to equation ( 1).

(1)

The minimal value of Rv is obtained for nearly stoichiometric TiC and is the ratio between TiC molar volume and carbon molar volume. This minimal theoretical value is equal to about two [START_REF] Ledain | Reactive Chemical Vapour Deposition of Titanium Carbide from H 2 -TiCl 4 Gas Mixture on Pyrocarbon: A Comprehensive Study[END_REF].

A sample of the last series of RCVI experiments was observed by transmission electron microscopy (TEM, Philips, CM30ST, the LaB 6 source operating at 300 kV). The thin slice (∼100 μm in thickness) was cut from the polished sample embedded in resin molding that was first observed by SEM. It was submitted to Ar + ion beam milling (JEOL, Ion Slicer, EM-09100IS) until the electron transparency of the region of interest is obtained. Selected area electron diffraction (SAED) patterns were recorded (selected-area diameter: 1.8 µm) as well as direct images in both bright field and TiC (111) dark field modes.

Results and discussion

After a thermal treatment at 1100°C under vacuum and before RCVI, the compact powder of the study materials has a porosity = 0.50 (±0.04) for an average pore size of 165 nm as measured by mercury porosimetry. It is worth noting that this result is in good accordance with the porosity P = 0.45 -0.50 locally measured by image analysis in the deepest areas of a sample infiltrated with a minimum of TiC by RCVI. This value was used to calculate Rv according to equation [START_REF] Naslain | Design, preparation and properties of non-oxide CMCs for application in engines and nuclear reactors: an overview[END_REF]. The presence of TiC after RCVI is confirmed by the characteristic peaks of the carbide on the X-ray diffraction patterns (not shown here). First, RCVI was studied at a temperature of 950°C for 4 hours to 36 hours. The volume fraction of TiC and the depth of infiltration in the powder increase with time (Fig. 5). In the infiltrated area, tends to a plateau between 0.4 and 0.5 from 24 hours and up to a depth greater than 500 µm. The maximum value is reached more quickly near the surface than in depth. However, if there is no observed sealcoating as expected, the maximum of (0.52) is not located just at the end surface of the sample but at a depth of about 50 µm (Fig. 6.a&b). The reason is unclear, but this may be due to a good gas exchange by forced convection facilitated on the edge of the sample. The volume conversion rate (Rv) decreases with depth and is below the minimal theoretical value of two. This means that a part of the consumed carbon is not converted into carbide but is lost in the form of gaseous by-product species which are evacuated out from the reactor. Such a behaviour has already been pointed out by Ledain et al. in the RCVD of TiC on pyrocarbon [START_REF] Ledain | Reactive Chemical Vapour Deposition of Titanium Carbide from H 2 -TiCl 4 Gas Mixture on Pyrocarbon: A Comprehensive Study[END_REF]. The carbon lost is responsible for the rather high porosity measured in the samples and the poor particle coalescence by TiC as observed by SEM (Fig. 6) despite the quantity of introduced carbide. It allows in return a homogeneous in-depth infiltration. In Fig. 7, the apparent mean thicknesses of the carbide coatings around each particle are plotted with infiltration time for different depths. During the first period up to 8 hours, the travel speed of infiltration front is about 60 µm/h. Near the whole material surface, the carbide coating thickness is already about 100 nm in these first hours whereas almost no coating is observed in depth. In the second period between 8 to 18 hours, the travel speed of infiltration front slows down and is about 20 µm/h. The carbide coating suddenly appears in intermediate depths. In the third period between 18 to 36 hours, carbide appears at the depth of 500 µm. The travel speed of infiltration front further slows down to about 8 µm/h. The thickness of the carbide layers increases gradually and slowly up to a maximum of about 130 nm everywhere in the sample up to this depth of 500 µm. Hence, the self-balancing of the growth rates, which allows the homogenization of carbide thickness with depth, is evidenced here. A second round of RCVI experiments was carried out at a temperature of 1050°C for 4 hours to 24 hours. At this temperature, the volume fraction of TiC is higher and the porosity is lower than at 950°C near the surface. But there is nearly no improvement of the maximum of infiltration depth with time which remains limited to only about 250 µm (Fig. 8).

does not increase any more either beyond 16 hours. Like at 950°C, seal-coating is not observed and the maximum of (here ~0.8) and density is not located just at the end surface of the sample; it is located here at a depth of about 10 µm (Fig. 9.a). At this temperature, the volume conversion rate Rv is higher than the minimum theoretical value of two. This volume increase during the carbon conversion without loss allows the particle coalescence by TiC as evidenced by the SEM observations where the material appears very dense, up to a depth of about 100 µm (Fig. 9). At the same value of ≈ 0.4, the carbide phase coating the carbon particles forms a continuous matrix at 1050°C (Fig. 9.b&c) whereas nearly independent core/shell-type particles are observed at 950°C with only a few simple points of contact (Fig. 6). It is possible that, under the combined effect of confinement and high temperature, the carbon containing gaseous by-products resulting from particle attack react in the porosity with titanium precursor to form additional TiC. The increase of the carbon volume conversion efficiency enhances the dense area formation near the sample surface that seals fresh gas access toward the core of the sample preventing the good infiltration of the powder. At 950°C, the temperature would be too low to allow the reaction of these carbonaceous gaseous by-products which are then led straight out the sample. A third series of RCVI experiments was carried out by raising in steps the temperature from 950°C to 1050°C. The purpose was to combine the in-depth homogeneity of infiltration achieved at low temperature with the pore-filling and density obtained at high temperature. A first 16-hours step at 950°C allowed the infiltration of TiC in the study materials up to a depth of 500 µm with a maximum limited to ~0.5 near the surface (Fig. 10). This result is in accordance with the one previously obtained in the first series of experiments carried out at the constant temperature of 950°C. The following temperature steps were then divided into a series of six hours in order to get an overview of the intermediate states of infiltration. At 1000°C, an increase of both infiltration depth (with a travel speed of infiltration front of ~ 8 µm/h) and is observed. Finally, at 1050°C an increase of is still observed near the surface of the samples, at a depth of 25 to 100 µm. The SEM observations show that the final study material is well densified up to a depth of 200-300 µm. The TiC matrix exhibits a good continuity with porosity as low as 3% in the densest area (Fig. 11.a) and 18% 200 µm beneath the surface (Fig. 11.b). At this depth, the use of a unique temperature of 950 C results in a porosity higher than 30%. In this area, the good pore-filling is in accordance with a volume conversion rate Rv ≈ 2.2 close to, but greater than the minimal theoretical value, reflecting the efficient conversion rate of carbon into carbide. A TEM observation of the sample at an estimated depth of around 150 µm is presented Fig. 12. The former carbon particles have been transformed into polycrystalline carbide particles. Each carbide particle can be identified by its more or less circular contour surrounded by pores. When the diameter of the transformed particle is lower than ~600 nm, no carbon is observed in the core, evidencing a total conversion. But when the particle size is higher, carbon remains in the centre and a core/shell structure is observed as in the case at the top and lower left of Fig. 12.a. In that case, the thickness of the carbide coating (shell part) around the remaining carbon (core part) is around 300 nm. The SAED pattern of the area is typical of polycrystalline titanium carbide (Fig. 12.b). A (002) carbon ring is hardly visible and they are no diffraction spots other than those of TiC, evidencing that there is no free titanium. A few carbide crystal grains are highlighted in dark field mode (Fig. 12.c) without noticing preferential orientation or particular aspect ratio. At this depth in this sample, a part of which has been made at 950°C for 16 hours, the carbide crystal grain size ranges from 20 to 200 nm with an average of 80 nm. This average grain size corresponds to the coating thickness from which the carbide growth slows down dramatically after 3 hours of RCVI at this depth (Fig. 7). As long as the carbide coating thickness is lower than this average grain size, the carbide growth can be considered to occur from direct reaction between the gas phase and carbon; the gaseous molecules have direct access to the carbon substrate between the growing carbide grains. Beyond this thickness, the grain edges join together and further reaction is only possible through solid state diffusion. At 950°C the growth is stopped when the coating thickness reaches 130 nm (Fig. 7), i.e. hardly twice the average grain size, because of too difficult diffusion. As long as gaseous species can reach the sample area, a temperature rise up to 1050°C allows the reactivation of growth through solid state diffusion through carbide. The carbide coating can then be composed of a stack of several grains and reach up to ~300 nm in thickness. According to Van Loo et al. who studied different TiC-Ti, TiC-C and TiC 1-y -TiC 1-z diffusion couples, carbon is the only diffusing component in TiC [START_REF] Van Loo | Diffusion of carbon in TiC 1-y and ZrC 1-y[END_REF]. Consequently at the particle level, the carbide growth should be driven only by the outwards diffusion of carbon and not by the inwards diffusion of titanium. But, when a nonreciprocal mutual solid-state diffusion process through an interface of two materials takes place, "Kirkendall voids" are formed close to the interface that leads to the formation of hollow objects [START_REF] Fan | Formation of nanotubes and hollow nanoparticles based on Kirkendall and diffusion processes: a review[END_REF] [START_REF] Geon Dae Moon | Chemical transformations of nanostructured materials[END_REF]. Here, there is no evidence of hole or void in place of carbon in the centre of the particles and the particles are not hollow. The production of dense carbide particles means that titanium can diffuse inwards. The discrepancy with Van Loo et al.'s observations may result from a difference in the involved mechanism: grain boundary diffusion in very fine-grained carbide in the RCVI case that facilitates element mobility vs. bulk diffusion in coarse-grained carbide in the diffusion couple technique case. The junction between three particles can be observed in Fig. 12.d. The particle junction results in grain boundary formation in the manner of a sintering mechanism, but the material consolidation is limited by the remaining pores between particles.

Rv and measured at each step and at different depths are reported in Fig. 13.

Near the surface, Rv nearly reaches its maximal value at the end of the first step, whereas at a greater depth it gradually increases with time. Thus, in depth, carbon increasingly tends to convert into carbide instead of being evacuated from the sample as in the case of RCVD or RCVI at a constant low temperature of 950°C (Rv ≈ 1.4 between 150 and 300 µm deep after 36 h). Compared to RCVIs at a constant single temperature, whether at 950°C or 1050°C, increasing in steps the temperature helps to improve the material densification up to a depth of 300 µm.

Conclusion

At a low temperature of 950°C, RCVI in the carbon -titanium carbide system allows the infiltration of the powder up to a depth of ~700 µm, while the infiltration of SiC by conventional CVI is limited to a depth of only 20 µm in the same carbon powder at even lower temperature of 800°C. The thickness of the TiC coating around the particles is homogeneous thanks to the self limiting process with a plateau of about 130 nm. But the particle coalescence is poor and a part of the carbon is evacuated in the gaseous phase. This phenomenon, by limiting the swelling of the converted particles, allows the material to be homogenized with a gradual increase of the infiltration depth. At a single high temperature of 1050°C, the infiltration depth is limited to only about 250 µm. There is still no seal-coating at the end surface like in conventional CVI, but owing to too rapid growth and good volume conversion rate of carbon into carbide without carbon lost, a dense area is formed near the surface which limits the diffusion of the gaseous reactive species. There is no increase of the infiltration front depth with time. With the use of time-temperature steps from 950°C to 1050°C, the consolidation of the powder is improved up to a depth of 300 µm. However to fabricate several millimetres thick CMCs with this hybrid process, the RCVI method needs to be further improved. Optimization should include the increase in the first step at 950°C over 16 hours. A process modelling, considering the phenomena observed as a function of temperature and depth, would allow for that optimization. Finally, it would also be interesting to consider different chemical systems with other elements than carbon and titanium.

Acknowledgements

This work was supported by Herakles (Safran group) and the Atomic Energy Commission.


L Maillé

A Guette

R Pailler

Y Le Petitcorps

P Weisbecker

Elaboration of C/C Composites Based on the Infiltration of a Hydrocarbon Precursor in Supercritical State into the Preform

Keywords: supercritical fluid, Chemical Vapor Infiltration, C/C composites

A work based on the development of a new process for the fabrication of a C/C composite for aeronautic and aerospace industries, in order to reduce the infiltration duration of the carbon preform, is reported in this paper. The infiltration of a carbon fibrous preform by the mean of a hydrocarbon in the supercritical fluid state is done at high temperature in order to get a carbon matrix. The microstructure of the pyrocarbon coating is characterized by SEM and TEM. The experimental parameters (temperature, hydrocarbon pressure, residence time) are tuned in order to elaborate the carbon matrix. The best conditions lead to a rapid densification in-depth from the top of the preform. After several experiments, the densification has been improved.

Introduction

Ceramic Matrix Composites are developed for aeronautic and spatial applications [START_REF] Naslain | SiC-Matrix Composites: Non brittle Ceramics for Thermo-Structural Application[END_REF][START_REF] Naslain | Introduction aux matériaux composites. Tome 2 : Matrices métalliques et céramiques[END_REF] for their refractoriness and their low density. These composites are usually get by a long infiltrating step of a weaved preform by Chemical Vapor Infiltration (CVI) using a hydrocarbon precursor as the propane for example [START_REF] Naslain | SiC-Matrix Composites: Non brittle Ceramics for Thermo-Structural Application[END_REF]. However, the CVI-step is a rather long process which is a huge drawback. As a result, the manufacturing cost of these composites is very expensive. In order to reduce the cost, other processes can also be employed to get the matrix, as liquid phase routes with mineral charges [START_REF] Rietveld | A profile refinement method for nuclear and magnetic structures[END_REF][START_REF] Rodriguez-Carvajal | Recent advances in magnetic structure determination by neutron powder diffraction[END_REF][START_REF] Sambasivan | Phase chemistry in the systemthermochemical review with phasestability diagrams[END_REF][START_REF] Paulasto | Interfacial reactions in Ti/Si 3 N 4 and TiN/Si diffusion couples[END_REF][START_REF] Hillig | Making ceramic composites by melt infiltration[END_REF][START_REF] Ber | Ceramic matrix composites densification by active filler impregnation followed by PIP process[END_REF][START_REF] Magnant | Carbon fiber/reaction-bonded carbide matrix for composite materials -Manufacture and characterization[END_REF].

In this paper, we explore an original and new densification process in order to elaborate a Carbon/Carbon (C/C) composite: the Chemical Supercritical Fluid Infiltration. A supercritical fluid is a substance at a temperature and pressure above its thermodynamic critical point. Its properties are simultaneously those of gases and liquids [START_REF] Trequesser | Supercritical fluid deposition of compositionally uniform yttria stabilized zirconia films[END_REF]. Indeed, it has the unique ability to diffuse through porous solids like a gas, and to dissolve materials like a solvent. This combination of the solvent ability and the high diffusivity should enable a rapid infiltration and a good densification of the porous preform.

Materials and experimental procedures

Two types of materials are used: (i) graphite pipes (length 6 cm, diameter 0.6 cm) are used as a substrate for kinetic studies of the deposition rate, and (ii) 3D porous carbon fibrous preforms (length 2 cm, diameter 1.5 cm) for infiltration study. The initial open porosity of the preforms is around 77%.

A cold wall reactor schematic represented in Fig. 1, is used at high pressure, in batch flow. It is made of inconel alloys. The graphite pipe is heated, using the Joule effect. The temperature is measured by a thermocouple located inside the graphite pipe protected by an alumina tube.

The carbon films or matrix are elaborated using the methane as the hydrocarbon precursor.

The kinetic of the carbon deposit can be studied by modifying the values of the residence time of the methane. The other parameters (initial pressure, initial temperature) are kept constant.

The temperature selected for each samples is 950°C, with an initial pressure of 50 bars, in these conditions the methane is a supercritical fluid (T critical = -82.6°C and P critical = 46 bar). For kinetic studies of the deposition rate, thin films are performed on the graphite pipes from 2.5 to 20 minutes of deposition. The duration selected (cf. Table 1) for one step infiltration of the carbon matrix texture 3 D, is fixed to 60 minutes.

Another experiment with several steps of infiltration, described table 1, is used to improve the matrix densification, with an initial pressure of 50 bar. The polished cross sections of thin carbon films or matrix are observed with an optical microscope, (Reichert-Jung MEF3).

Results and discussion

Pyrocarbon coating on a graphite pipe

With the variation of the residence times, it is possible to reach an important thickness of pyrocarbon: from 3µm to 30 µm for 2.5-5 minutes and 10-20 minutes respectively. It confirms the interest to work using a supercritical fluid, by comparison, several hours are necessary to reach the same thickness with the CVI-process [START_REF]Design -fabrication and application of thermostructural composites (TSC) like C/C -C/SiC and SiC/SiC composites[END_REF].

The thickness of the pyrocarbon coating varies significantly along the pipe due to a presence of a temperature gradient (Fig. 2). The temperature is a particularly significant and sensitive parameter.

The texture of the pyrocarbon coating is similar to the one elaborated with CVI-process (Fig. 3). Pyrocarbon thin film is characterized by a cone-like growth structure [START_REF] Bourrat | Regenerative laminar pyrocarbon[END_REF].

In order to demonstrate the feasibility to perform a pyrocarbon thin film with a hydrocarbon precursor in supercritical state, carbon preforms have been densified with the same protocol.

Infiltration of the pyrocarbon matrix into a fibrous preform

The influence of the hydrocarbon precursor residence time during the matrix infiltration is studied. The temperature and the initial pressures are fixed at 950°C and 50 bars respectively.

These first experiments show a good densification (infiltrated of 0.5 minute depth of the preform) for a short residence time (10 minutes). However, as it could be observed on the sample performed with residence time of 60 min, the densification is not complete (Fig. 4).

Various pores are present between some fibers even on the surface of the preform.

The cross sections observations show a densification thickness of 1.5 mm in the depth of the preform with a residence time of 60 minutes (against 0.5 mm with 10 minutes). Consequently, with this experiment, the densification is significantly improved to the first one, but a lot of porosities are presents in the matrix. However, the pyrocarbon thin film is homogeneous in thickness and adherent to the fiber.

In order to improve the densification, other experiments are performed, using the semi continuous feeding reactor (table 1). Five steps of infiltration have been done, equivalent to 42 min for the residence time. The change in the temperature between the step 3 and 4 was done in order to introduce the pyrocarbon at low temperature inside the preform, and to finish the densification at high temperature, to fill the porosities.

The matrix densification using the semi continuous feeding of the reactor has been improved. After these five steps, the gain mass of the sample is increased to 90% by this way. Whereas it varies from 41% and 43% for one batch of 10 minutes and 60 minutes respectively (table 1).

It is possible to compare the cross sections of the figure 3 and4, and to observe an improvement of the infiltration with the semi continuous feeding reactor. A good densification thickness of 1.5 mm in the depth of the preform is observed. A pyrocarbon coating is observed around each fiber until 2.5 mm in-depth of the preform (figure 5 a-b).

TEM characterizations confirm the presence of pyrocarbon, and an anisotropy structure (Fig. 6).

Conclusion

The objective of this study was to demonstrate that a pyrocarbon coating could be deposited around each carbon fibres within a 3D carbon preform with a methane precursor in its supercritical state. A device has been set up for this purpose alloying to work under a pressure of 50-100 bar at 950-1050°C for less than one hour of infiltration. According to the experimental parameters chosen the infiltration depth of the preform could reach 1.5 mm and a pyrocarbon coating is observed around each fiber until 2.5 mm in-depth. The results were interesting, however other experiments have to be conducted in order to get a full infiltrated preforms if necessary.

LIST OF FIGURE AND TABLE CAPTIONS

FIGURESFigure 1 .

 1 FIGURES

Figure 2 .

 2 Figure 2. Influence of the gradient temperature of the graphite pipe on the pyrocarbon thickness.

Figure 3 .

 3 Figure 3. Cross section of a pyrocarbon layer deposited on the carbon pipe (pyrocarbon elaborated with 20 minutes of residential time).

Figure 4 .

 4 Figure 4. Cross sections of the infiltrated of C/C composite with a residence time of 60 min, at different scale.

Figure 5 .

 5 Figure 5. Cross section of sample realised with the five steps described in table 1.

Figure 6 .

 6 Figure 6. TEM characterizations.

TABLETable 1 .

 1 Infiltration conditions of the 3 D carbon preforms.

TABLE 1 (

 1 *) For each step, the vessel is evacuated and then feed with the CH 4 precursor at a constant pressure of 50 bar.

	samples	temperatures (°C)	durations (min)	conditions	Mass gain
	A	950	60	One step infiltration	43 %
				Semi-continuous
				feeding
	B	950	10	Step 1(*)	41 %
		950	10	Step 2
		950	10	Step 3
		1050	10	Step 4
		1050	2	Step 5	90 %

Acknowledgment

The authors would like to thank S. Couthures for the fabrication of the experimental process and assistance in testing.


L Amarsid

F Radjai

SURFACE WAVE GENERATION BY SEABED COLLAPSE SIMULATED BY THE COUPLED LATTICE BOLTZMANN / DISCRETE ELEMENT METHOD

The aim of this work is to model the generation and propagation of surface waves as a consequence of submarine landslides. The numerical approach presented in this paper allows for the simulation of both underwater granular flows at the grain scale and free surface wave propagation. The discrete element method (DEM), extensively used for dry granular avalanches, is coupled with the Lattice Boltzmann Method (LBM) for the integration of Navier-Stokes equations. The fluid-grain coupling is introduced at the grain scale by imposing the no-slip condition for the fluid at the grain surface and by integrating at each time step the resulting hydrodynamic forces exerted by the fluid on each grain. The so-called "mass tracking method" is used to introduce a free surface in the model.

I. INTRODUCTION

The generation of tsunami is a major devastating hazard associated with submarine landslides occurring on the continental shelf [START_REF] Mason | Submarine landslides: processes, triggers and hazard protection[END_REF]. The key mechanisms that control the nature of such tsunamis are not yet clearly established and seem to involve various parameters such as the volume of the landslide, its velocity and initial acceleration, its thickness, as well as the depth of water on which depends the propagation speed [START_REF] Lallemand | Lattice Boltzmann method for moving boundaries[END_REF].

Idealised two-dimensional numerical and laboratory models of tsunamis generated by submarine landslides have recently been proposed [START_REF] Staron | The spreding of a granular mass: role of grain properties and intial conditions[END_REF][7] [START_REF] Bouzidi | Momentum transfer of a Boltzmann-lattice fluid with boundaries[END_REF] but they all assume a rigid block slide and do not account for the granular scale. A two-layer continuous model has been developed to represent water flow on a granular bed that is considered as a shallow fluid of greater density as water [START_REF] Spinewine | Solution properties and approximate Riemann solvers for two layer shallow flow models[END_REF]. The two fluids differ in terms of rheology, compressibility, viscosity, and potential for mixing. Although successful in accounting for general phenomenology in a short computation time, such models, defined at a large scale, are limited in describing all features of seabed granular flows. A pending research issue thus remains the parameterization of the interactions between the water phase and the sediment phase. Owing to the number of flow variables involved and measurement imprecision, estimating such parameters from laboratory experiments remains difficult.

Contrarily to submarine landslides, aerial granular landslides have been extensively studied during the past two decades by means of experiments and numerical simulations [10][15] but the generalization of the results in the presence of a fluid remains unexplored to the present date. The onset and propagation of seabed instabilities involve strong coupling between water and the granular phase via water pressure gradients in the pore space or dilatancy of the granular phase [START_REF] Legros | The mobility of long-runout landslides[END_REF][3] [START_REF] Cassar | Submarine granular flows down inclined planes[END_REF]. Different approaches have been developed to take into account fluid-grain coupling [START_REF] Mcnamara | Grains and gas flow: Molecular dynamics with hydrodynamic interactions[END_REF][27] [START_REF] Bouzidi | Momentum transfer of a Boltzmann-lattice fluid with boundaries[END_REF]. A methodology has been recently proposed to enable predictive simulations of underwater granular flows at the grain scale, but without accounting for free surface and possible wave generation [START_REF] Mansouri | A 3D DEM-LBM approach for the assessment of the quick condition for sands[END_REF][5] [START_REF] Topin | Micro-rheology of dense particulate flow: application to immersed avalanches[END_REF].

The present work is based on this methodology in which the Discrete Element Method (DEM), extensively used to model dry avalanches [START_REF] Staron | The spreding of a granular mass: role of grain properties and intial conditions[END_REF], is coupled with the Lattice Boltzmann Method (LBM) that enables the introduction of a water phase and the calculation of the hydrodynamic forces on the grains. The so-called "mass tracking method" is used to introduce a free surface in the model [START_REF] Thürey | Stable free surface flows with the lattice Boltzmann method on adaptively coarsened grids[END_REF].

This paper is structured as follows. The methodology is presented in section 2: The LBM is briefly introduced for 2D water flow with free surface modelling, and the coupling methodology with DEM is presented. Section 3 provides three validation tests and results: (i) the 2D model is first validated against 1D shallow water model for free surface wave propagation, (ii) the falling speed of a grain is compared to analytical results and (iii) the first results of the collapse of a grain column underwater are given.

II. METHODOLOGY

LBM approach for modelling free surface flow in two vertical dimensions

The LBM consists in applying the Boltzmann equation to the fluid, which is modelled as a collection of fluid particles moving on a fixed rectangular grid. To model free surface flows in two dimensions, the so-called D2Q9 model (2 dimensions, 9 velocities) is used along with the "mass tracking" approach (Figure1) [START_REF] Thürey | Stable free surface flows with the lattice Boltzmann method on adaptively coarsened grids[END_REF]. . At each time step, the density and velocity are transported from one cell to its 8 neighbouring cells (i=1,…,8) or remain within the cell (i=0). To each of the nine possible directions i, are associated a local density distribution function f i , and a velocity u i such that

            i i i i i t f t t f t , , , , x u x u x x   (1)
with u i =u e i where u=x/t.

At each time step, two main operations are performed on flow particles: collision and streaming. These are briefly recalled below (see e.g. [START_REF] Wachs | A DEM-DLM/FD method for direct numerical simulation of particulate flows: Sedimentation of polygonal isometric particles in a Newtonian fluid with collisions[END_REF] or [START_REF] Sukop | Lattice Boltzmann Modeling[END_REF] for more details).

Streaming:

The particles with local density

  t f i ,
x located at position x at time t, move with velocity u i to the point x+u i t during a time step t:

    t f t t t f i i i , , x u x      (2)

Collision:

It is assumed that the particles with density

  t f i ,
x entering simultaneously the same cell will collide with each other, yielding a new distribution function

  t f i , out
x defined by the BGK model [START_REF] Wachs | A DEM-DLM/FD method for direct numerical simulation of particulate flows: Sedimentation of polygonal isometric particles in a Newtonian fluid with collisions[END_REF]:

            , , 1 , , eq out u x x x i i i i f t f t f t f    (3)
where  is a dimensionless relaxation time and the eq i f are the equilibrium distribution functions given by

                 u u u e u e u 2 2 4 2 eq 2 3 2 9 3 1 , c c c w f i i i i   (4)
where the w i are density weights (w 0 =4/9; w i =1/9 for i=1…4 and w i =1/36 for i=5…8).

Boltzmann equation:

Equations ( 2) and ( 3) are combined to give the Boltzmann equation that includes a supplementary term proportional to the external force F: [START_REF] Delenne | Onset of Immersed Granular Avalanches by DEM-LBM Approach[END_REF] where c s is the so-called sound speed of the Boltzmann network, given by .

          F v c t w f t f t f t t t f i s i i i i i i 2 in eq in , , 1 , ,            u x x u x

Treatment of the free surface:

The free surface is taken into account by introducing 3 types of cells: fluid, empty and interface (Figure 1). Interface cells are partially filled cells that separate the air (empty cells) from fluid. The mass tracking approach [START_REF] Thürey | Stable free surface flows with the lattice Boltzmann method on adaptively coarsened grids[END_REF] consists in tracking the movement of the fluid interface by the calculation of the mass exchange m i between interface cells and their neighbours. This flux can be directly computed from the f i movement at the streaming step.

For each given direction i, three configurations can occur according the type of the neighbour cell (located at x+u i t):

-Type E (empty cell). The mass exchange is null:

  0 ,     t t m i x (6)
-Type F (totally filled cell). The mass exchange is simply:

     t f t t f t t m i i i i , , , ~x u x x        (7)
where ~i represents the opposite of direction i. The first (resp. second) density function is the amount of fluid that is entering (resp. leaving) the interface cell.

-Type I (interface cell). The mass exchange between two interface cells must account for the quantity of fluid contained in the cells (fluid fraction ε = m/):

            2 , , , , , ~t t t t f t t f t t m i i i i i x u x x u x x             (8)
A more precise computation of the mass exchange can be done for this third case according to the type of the two cells neighbours. It is not presented here for the sake of conciseness but can be found [START_REF] Thürey | Interactive free surface fluids with the Lattice Boltzmann Method[END_REF]. The mass exchange is computed for each neighbouring cell and the mass is thus updated to: [START_REF] Ladd | Numerical simulation of particular suspensions via a discretized Boltzmann equation, Part 2, Numerical results[END_REF] Note that, at the end of this step, some interface cells can be emptied or filled. The cell type is thus changed accordingly and the lacking or exceeding mass is reallocated to neighbouring cells.

               8 1 , , , i i t t m t t t m x x x

Example:

Figure 2 shows an example of free surface simulation. Results obtained for the classical dam-break problem are in very good agreement with those obtained from experimental test cases such as proposed in [START_REF] Duarte | Experimental study on dam-break waves for silted-up reservoirs[END_REF]. For the sake of comparison with [START_REF] Duarte | Experimental study on dam-break waves for silted-up reservoirs[END_REF] the classical non dimensional variable 0 / h g t T  is used, where h 0 is the initial water depth in left side of the dam.

Grain modelling: sample construction and DEM computation

The Discrete Element Method is used to build a densely-packed sample of polydisperse spherical grains. A set of grain diameters according to the cumulative beta distribution is obtained with the statistical model proposed by [START_REF] Topin | Micro-rheology of dense particulate flow: application to immersed avalanches[END_REF]. The beta distribution has the advantage to be bounded on both sides and capable of representing double-curved distributions similar to the soil grain-size distributions encountered in practice. Once the diameters are generated, the grains are first placed on a regular grid in a rectangular column as a dilute sample and then compacted between two plates using the procedure proposed in [START_REF] Schafer | Force schemes in simulations of granular materials[END_REF].

The force F between two grains is defined by F=(N,T) where the normal part N is due to the contacts between particles and the tangential part T stems from friction. A linear-elastic approximation is used to represent the force law between a pair (i,j) of contacting particles:

    otherwise 0 0 f o r 2           ij ij n n ij N n mk k N     ( 10
)
where  is the gap or the overlap between the two contacting particles i and j and   its derivative with respect to time, k n is the normal stiffness, m=m i m j /(m i +m j ) is the reduced mass and

  1 , 0  
is a damping parameter which controls energy dissipation due to inelastic collision. The friction is computed with the classical Coulomb law expressed as a nonlinear relation between friction force T and the sliding velocity.

     n N v T f t     , m i n
where  is the tangential viscosity parameter,  f is the friction coefficient and v t the tangential velocity at contact.

Coupled LBM/DEM model for fluid/solid interactions

Solid particles (grains) previously generated are introduced in the fluid domain of LBM using the same discretization scheme. Since in two dimensions the poral space does not percolate through the sample, for the fluid computation, we reduce the diameters of the grains in order to have a minimum of two nodes of fluid at contact between grains. A new type of cell is thus used: G for grain cell (see figure 3). As for free surface, the interaction between fluid and solid particles appears only at the interface. Grains are considered as moving boundaries on which the non-slip condition is imposed to the fluid using the standard bounce back rule. The hydrodynamic forces acting on grains are computed using the momentum exchange method proposed by [START_REF] Ladd | Numerical simulation of particular suspensions via a discretized Boltzmann equation, Part 2, Numerical results[END_REF]. The fluid particles that are bounced back on grain boundaries transmit forces to the grain proportionally to their momentum change [START_REF] Bouzidi | Momentum transfer of a Boltzmann-lattice fluid with boundaries[END_REF] [START_REF] Lallemand | Lattice Boltzmann method for moving boundaries[END_REF]. The implementation of this method is straightforward since particles momenta are known at each time step. The total hydrodynamic forces exerted on a solid particle are obtained by summing up the forces on all boundary nodes of that particle. Finally, these forces are used to compute the motion of grains.

Example:

Figure 4 shows the velocity of a grain in freefall in comparison with theoretical results given by the second Newton's law and a friction law of the form f=-ku 2 valid for an inertial regime i.e. for a Reynolds number greater than one, which is the case here.

III. NUMERICAL EXPERIMENT AND RESULTS

For the sake of conciseness and clarity, only qualitative results are presented here. A more precise physical interpretation as well as parameter study will be presented in detail in further work. . This test has been performed with five different column heights and the same water depth. The ratio of the water depth over the initial water depth at T=0.08 has been reported in Figure 5 for each test case as well as the maximum and minimum water depth at T=0.08 as a function of the initial granular column height. As could be expected, the amplitude of the generated wave increases with the granular column height. It can also be seen in Figure 6 (left) that the extreme values are not located at the same abscissa. Further studies will allow for a more precise parameter analysis to be conducted.

IV. CONCLUSION

The presented work gives the basis for the simulation of fluid/grain interactions in presence of free surface. The methodology has first been validated for simple cases. The dam-break test case (without grain) gives similar results to those experimentally obtained in the literature. The falling velocity of a unique grain in free fall has been successfully compared to theoretical laws. As a first parametric analysis, we assessed the influence of the initial column height on the amplitude of the wave generated at the free surface. Other parameters will also be studied such as the aspect ratio of the grain column, grain distribution characteristics, etc. An interesting problem is also to investigate the triggering of the collapse using a cohesive model for the contact of granular materials such as proposed in [START_REF] Delenne | Mechanical behavior and Failure of cohesive granular materials[END_REF].

The comparison of the proposed methodology with experimental results obtained in [START_REF] Duarte | Experimental study on dam-break waves for silted-up reservoirs[END_REF] or [START_REF] Zech | Dam-break induced sediment movement: Experimental approaches and numerical modeling[END_REF] for dam-break upon sediments is also one of the forthcoming goals of this work.

Figure 1 .

 1 Figure 1. Left: D2Q9 scheme and Right: discretization of the flow domain according to cells types. The flow domain is discretized in square cells of linear size x, to which are associated a position vector x, a density   t , x 

Figure 2 .

 2 Figure 2. Dam-break on dry bed. Screen capture taken at T=0, T=0.88, and T=1.57 (these simulations are in good agreement with experimental results given in figure 3 of [7]).

Figure 3 .

 3 Figure 3. Grains and fluid discretization (cells types).

Figure 4 .

 4 Figure 4. Validation of the grain fall velocity in water. The theoretical curve is obtained with a friction law of the form f=-ku 2 s.

Figure 5 .

 5 Figure 5. Collapse of a grain column. Up: ratio column height/water depth = 0.508 (spl05). Down: ratio=0.915 (spl09). Screen captures taken at T=0, T=0.04 and T=0.08.

Figure 6 .

 6 Figure 6. Left: ratio of water depth h at T=0.08 over the initial water depth h0 for the different samples. Right: minimum and maximum water depth as a function of the initial granular column height (all values being normalized by h0)


L Maillé

M A Dourges

S Le Ber

R Pailler

CERAMIC MATRIX COMPOSITES MANUFACTURED BY MULTISTEP DENSIFICATION OF Si-O-C FIBRE PREFORM

Keywords: CMC, active filler, slurry impregnation, flexural strength

Ceramic matrix composites have been manufactured at temperatures below 1100°C by a multistep process which consists in (i) a polymer impregnation of Si-C-O fibre preform and pyrolysis (PIP) in order to consolidate the preform ; (ii) an impregnation with a slurry of submicronic TiSi 2 powder, (iii) a heat treatment under nitrogen (N 2 ) in order to reduce the porosity thanks to the volume expansion associated with the nitridation of the active filler ; and (iv) a final densification by PIP. Various liquid polymeric precursors were investigated to perform this final densification. The morphology, the chemical composition and the mechanical properties of the composites were characterized. Based on these characterizations, a preceramic polymer precursor was chosen in order to produce a ceramic matrix composite.

INTRODUCTION

The processing of Ceramic Matrix Composites (CMC), using impregnation and pyrolysis of preceramic polymer precursor (PIP) offers advantages over others processing routes (chemical vapour infiltration, sol-gel route, reactive melt infiltration or several other techniques [ 1 -5 ] ) ; particularly, PIP is based on relatively simple technologies. In order to be competitive on the civil aeronautics market, low cost CMC processing routes must be developed [START_REF] Gadow | Mechanical properties of ceramic matrix composites with siloxane matrix and liquid phase coated carbon fiber reinforcement[END_REF][START_REF] Suttor | Fiber-reinforced ceramic-matrix composites with a polysiloxane/boron-derived matrix[END_REF][START_REF] Azim | Development of coating for temperature sensing[END_REF][START_REF] Woiton | Self-assembled microstructured polymeric and ceramic surfaces[END_REF] . PIP of SiC fibre preforms could be attractive but due to an inherent shrinkage obtained after pyrolysis of the polymer, larges pores are formed in the composite and multiple cycles of PIP must be performed to increase the densification. Complementary methods of densification such as slurry impregnation with filler can be performed before the PIP to obtain an effective process with a low price / performance ratio. However, the repetition of numerous impregnation and pyrolysis cycles is then necessary to obtain a dense material [START_REF] Jones | Polymer derived ceramic matrix composites[END_REF] . As it has been suggested by P. Greil in the case of ceramic preparation, this problem could be overcome with the use of active fillers, which react during pyrolysis under reactive atmosphere to form oxides, carbides or nitrides leading to a significant volume expansion that can decrease the matrix porosity [START_REF] Greil | Active-filler-controlled pyrolysis of preceramic polymers[END_REF][START_REF] Greil | Near Net Shape Manufacturing of Polymer Derived Ceramics[END_REF][START_REF] Greil | Near net shape manufacturing of ceramics[END_REF][START_REF] Greil | Pyrolysis of Active and Passive Filler-loaded Preceramic Polymers, Handbook of Advanced Ceramics[END_REF][START_REF] Erny | Microstructure development of oxycarbide composites during active-filler-controlled polymer pyrolysis[END_REF][START_REF] Akkas | Silicon oxycarbide-based composites produced from pyrolysis of polysiloxanes with active Ti filler[END_REF][START_REF] Xie | Active filler (aluminum-aluminum nitride) controlled polycarbosilane pyrolysis[END_REF][START_REF] Ade | Electroconductive ceramic composites with low-to-zero shrinkage during sintering[END_REF][START_REF] Torrey | Composite polymer derived ceramic system for oxidizing environments[END_REF][START_REF] Vojtech | High temperature oxidation of titanium-silicon alloys[END_REF] . Among various expansion agents, titanium disilicide powder (TiSi 2 ) is identified as an interesting active filler [20- 23] . Under nitrogen atmosphere, the nitridation of sub-micronic TiSi 2 powder starts around 1000°C, and leads to the formation of TiN (d = 5.43 g/cm 3 ) and Si 3 N 4 (d = 3.19 g/cm [START_REF] Rosso | Ceramic and metal matrix composites: routes and properties[END_REF] ) with a large volume increase even if the reaction is incomplete. This study explores the processing of ceramic matrix composites by slurry impregnation of submicronic active filler, subsequent nitridation and liquid polymer impregnation and pyrolysis with liquid polymeric precursors. First, pellets prepared from sub-micronic nitrided TiSi 2 powder were impregnated with different precursors and heat treated to explore the behaviour of the polymers. According to the experimental results on monolithic matrices, liquid polymeric precursors were chosen to manufacture the fibre reinforced ceramic matrix composites. Finally, the mechanical behaviour (3-point bending tests) of CMCs is characterized in order to select the best polymer precursor.

MATERIALS AND EXPERIMENTAL PROCEDURE

Materials

A high purity micrometer-sized TiSi 2 powder (C-54 stable phase, 99.95% in purity, ~ 45 µm, Neyco) was used during this work. This raw powder was milled with a planetary ball mill (Retsch PM200). Composites were fabricated from 2D fibre preforms (~2mm thickness) made of woven Nicalon fibres (Nippon Carbon Co) covered with a PyC interphase. The fiber preforms are first consolidated by one PIP cycle with a phenyl-containing polysiloxane (resin A) ; after this consolidation step the porosity is around 48%. All the preforms used were purchased in the consolidated state. The four resins used as preceramic polymers for the densification of pellets are presented on table 1 The first step is the milling of TiSi 2 commercial powders, which are not fine enough to be impregnated into the fibre preforms. They were milled with a planetary ball mill (Retsch PM200). Tungsten carbide (WC, d=15.6g/cm 3 ) accessories were chosen in order to perform high energy millings: a 50 mL hermetic jar. 18 g of raw powder and 305 g of WC balls (3 mm in diameter) were loaded with 15 mL of isopropanol in the WC jar. The millings were performed at a jar rotational speed of 960 rpm during 3 hours. In order to limit the temperature raise in the jar, the milling time was split in eighteen 10 minutes effective milling periods interrupted by 20 minutes cooling periods. The milled powder was collected by evaporating the solvent at 100°C under vacuum (grain size diameter d 50 : 0.5 µm; d 90 : 1.1 µm).

Processing of composite pellets: TiSi 2 powder was compression moulded to make pellets. Pellets are nitrited 1 h with a stream of N 2 at a temperature of 1100°C. They were used to study the behaviour of the preceramic precursor during impregnation and pyrolysis in contact with the solid nitrided TiSi 2 . The pellets were immersed in the liquid polymeric precursors (alone or in solution) in a beaker. The beaker was placed in a vessel, which was evacuated and heated at 60°C in order to reduce the viscosity of the polymer. After one hour of impregnation, the samples were removed from the beaker and put in a furnace. The pellets were heated under argon atmosphere at a ramping rate of 20°C/min up to 200°C, maintained at this temperature for 2 h for curing ; the temperature was then raised to 900°C for pyrolysis at a rate of 10°C/min, and maintained at this temperature for 1 h.

Processing of fibre-reinforced composites: Ceramic matrix composites were prepared with 2D Nicalon fibres. These SiC-based fibres are unstable at high temperatures because of the silicon oxycarbide phase they contain, which decomposes beyond 1100 °C ; the manufacturing of CMC is therefore limited to this maximum temperature. Active filler was impregnated in the preform using slurry impregnation method. The active filler powders were mechanically mixed with ethanol in order to obtain slurry. The addition of poly-ethylene imine (PEI) enabled to stabilize the colloidal suspension. A concentration of 15%V of active filler powder could be obtained with an optimum concentration of 2.5 mg of PEI per square meter of powder. The fibre preforms were immersed in a beaker of suspension and the impregnation was performed under vacuum for one hour. The samples were then removed from the suspension, and the solvent was evaporated by heating at 100°C under vacuum for one hour. The composites impregnated with the active filler were then nitrided under a flow of nitrogen gas at a ramping rate of 10°C/min up to 1100°C, and maintained at this temperature for 5 h. The next step was the PIP cycle with the preceramic polymers presented table 1. The PIP cycles were carried out with the same experimental set-up as for the pellets preparation.

Characterizations methods

The pore size distribution and porosity of the materials were determined by mercury porosimetry (Micromeritics Autopore IV 9500). The microstructure of CMC was observed by scanning electron microscopy (SEM). Experiments were carried out using a Quanta 400 FEG microscope.

3-points bending strength of test pieces 50 mm (length) x 10 mm (width) x 2 mm (thickness) in size was measured with a universal testing machine (Instron 5860) at a cross-head speed of 0.5 mm/min with a span of 30 mm, at room temperature. Two specimens were performed for each experimental condition.

The ultimate flexural strength (σ R ) was calculated according to the equation ( 1): σ R = 3 F L / 2 w t² (1) where F is the maximal applied load, L the support span (50 mm), w the width (~ 10 mm) and t the thickness (~ 2 mm) of the specimen.

RESULTS AND DISCUSSION

PIP of precursors in pellets

In order to select the preceramic polymers, composite pellets of ceramic were prepared using nitrided TiSi 2 cylinders and PIP precursors listed on table 1. The study of the behavior of resins A to D is presented in this paper. As it is shown in table 2, the impregnations were performed in solutions of 50wt% of xylene in the case of resin B, of 70wt% of methyl ethyl ketone (MEK) in the case the resin C and without solvent for the resin A and the resin D. Table 2 presents the composition and porosity obtained for the four composites after one PIP cycle. The ceramic proportion is higher for the materials prepared with the precursors used without solvent. Such result could be expected if we take into account the conditions of impregnation. A large amount of solvent is added to obtain a suitable viscosity for resins B and C. Even if the yield of transformation is high, particularly in the case of resin C, the final proportion of polymer residue in the pellet is low, due to the elimination of the solvent. The best results were obtained with resin A and the resin D. Mercury intrusion porosimetry was used to investigate the microstructure of materials. The pellets A and B fabricated with resin A and the resin D offer a lower porosity (about 8%) than the others obtained from resin B and C (20 and 28% respectively).

Figure 1. Pore size distribution of the pellet before and after PIP with different preceramic polymers

Figure 1 shows the pore size distribution of the pellets, after one PIP cycle, for the four materials prepared. The size of the pores is lower than 1 µm. Table 2 indicates a similar total pore volume of about 8% for pellets A and D. But in figure 1, we can observe that the size range of the pores is different, around 100 nm for the pellet A and around 10 nm for pellet D. Difference of initial viscosity of the two precursors can explain this observation. Based on these results, resins A and D were selected to prepare fibre-reinforced ceramic matrix composites.

Preparation and characterization of Ceramic Matrix Composites 3.2.1. Composites processing

Three types of composites were produced using the following steps: infiltration of the milled TiSi 2 powder (18 ± 4 %w), nitridation and one PIP cycle with the resin A and D. The composition of the composite is given in detail in table 3. Proportion of pre-ceramic residue is higher in CMC2, as observed in the pellets.

Characterizations of composites

The composites were characterized by mercury intrusion porosimetry. The mechanical properties were also determined using 3-point bending tests. The influence of (i) the PIP precursor and (ii) the repetition of the impregnation and pyrolysis cycles were studied. Table 4 summarizes the results. Figure 2 exhibits typical load-displacement curves for all composites. The very first (slightly) non-linear part of the stress-displacement curves may be related to the progressive adjustment of the samples. The first linear part of the curve shows an elastic behaviour of the material; it is followed by a cracking stage until the sample fails. The strengths are distributed between 114 and 163 MPa with only one PIP cycle of resin A and resin D respectively; and reach 215 MPa with three PIP cycles using the resin D. As it was expected, the values of strengths depend particularly on the number of PIP cycles performed. This point can be confirmed regarding the final porosity of CMC measured as 15.5% for CMC2 and 7.3% for CMC3. These results are interesting because the same range of flexural strength values is reported in literature for CMCs elaborated by liquid routes, but in some cases authors choose to multiply PIP cycles (> 6) [START_REF] Ziegler | Fiber-reinforced composites with polymer-derived matrix: processing, matrix formation and properties[END_REF][START_REF] Dong | Microstructural evolution and mechanical performances of SiC/SiC composites by polymer impregnation/microwave pyrolysis (PIMP) process[END_REF][START_REF] Cordoba | Nitriding study of titanium silicide intermetallics obtained by mechanical alloying[END_REF][START_REF] Lu | Fabrication of a NicalonTM fiber/Si3N4-based ceramic-matrix composite by the polymer pyrolysis method[END_REF][START_REF] Shin | Low-temperature processing of ceramic woven fabric/ceramic matrix composites[END_REF] , even if they introduce passive or active fillers [START_REF] Nakano | Fabrication and characterization of three-dimensional carbon fiber reinforced silicon carbide and silicon nitride composites[END_REF][START_REF] Stantschev | Long fibre reinforced ceramics with active fillers and a modified intra-matrix bond based on the LPI process[END_REF] . We can notice that using resin D to perform the PIP improves significantly the flexural properties compared to the use of resin A, although the porosity range is the same. To understand this result, microstructure was explored and particularly the repartition of the ceramic residue within the intra-bundle and inter-bundle parts of the composites.

Microstructure of composites

Figure 3 presents CMC1 composite micrographs. This observation shows that the powder is impregnated into the inter-bundle macropores. The TiSi 2 nitridation created TiN, Si 3 N 4 and free Si phases [START_REF] Maillé | Study of the nitridation process of TiSi2 powder[END_REF] . Preforms used were consolidated with the resin A, therefore it is not always possible to determine if the observed ceramic residue is the results of the consolidation by the final PIP cycle. However, the following remarks can be made: grain wetting with the resin A appears poor and the ceramic residue does not adhere to grains. Figure 4 presents a SEM micrograph of CMC2, prepared with one PIP cycle using resin D. Figure 3 and 4 allow us to appreciate the differences between the behaviour of resins A and D during thermal treatment. Indeed, the resin shrinkage due to the pyrolysis of resin A produces a dense ceramic residue without cracks ; the ceramic residue is not much adherent to fibres and to nitrided powder (figure 3). On the contrary, the pyrolysis of resin D produced a residue which Figure 5 shows the distribution of the different constituents inside the composite CMC3. Consolidation residue is particularly localised within the bundle. The nitrided powder fills the inter-bundle macropores. This powder is consolidated by the PIP cycles. Some cracks due to polymer shrinkage can also be observed. This SEM micrograph confirms the low porosity of the composite.

CONCLUSION

CMC were produced by a relatively simple process combining three different methods of densification. Using impregnation of sub-micronic active filler and nitridation with a volume expansion, part of the porosity was filled. The subsequent PIP cycles with a suitable pre-ceramic polymer allowed to produce a composite with mechanical properties in the same range than others CMCs obtained by more complex or more time-consuming processes. Choosing this liquid route and Nicalon fibre preforms enabled to develop a low-cost process performed at temperatures below 1100°C. After a study of various ceramic precursors, a resin was selected. This pre-ceramic polymer displayed the best properties to elaborate these composites: a low viscosity at ambient temperature and a good wetting of the nitrided TiSi 2 powder. The flexural properties should be improved significantly if the number of PIP cycles was increased.

Figure 3 .

 3 Figure 3. SEM micrograph of polished cross section of a composite (Nicalon preform) elaborated with one impregnation of TiSi 2 powder, nitriding 5 h at 1100°C and with one PIP cycle using resin A.

  to the nitrided powder and fibre as it was revealed by the cracks observed.

Figure 4 .

 4 Figure 4. SEM micrograph of polished cross section of a composite (2D Nicalon preform) elaborated with one impregnation of TiSi 2 powder, nitrided 5 h at 1100°C and with one PIP cycle using (CMC2).

Figure 5 .

 5 Figure 5. SEM micrograph of cross section of CMC3.

  .

	Table1. Ceramic precursor characteristics
	Ceramic precursor	Composition	Viscosity
	Resin A	Methylphenylvinylhydrogen polysiloxane resin	900 mPa.s
	Resin B	Alkylsilicone resin + xylène	6-12 mm 2 /s 50%w xylène
	Resin C	Methyl silicone resin + methyl ethyl ketone	powder 50%w methyl ethyl ketone
	Resin D	Polymethoxy methylsiloxane / methyl silsesquioxane resin	23-30 mPa.s
	2.2. Processing of composites
	Powder milling:

Table 2 .

 2 Results of impregnation and pyrolysis of ceramic precursor on pellets of TiSi 2 powder

	Samples	Precursor	Solvent	precursor proportion	Impregnation weight gain	Ceramic yield (900°C/ Ar)	Proportion (weight) of polymer residue	Porosity
	Pellet A	Resin A	-	100%	27%	73%	21%	8.5%
	Pellet B	Resin B	xylène 50wt%	50%	22%	65%	15.5%	20%
	Pellet C	Resin C	MEK 70wt%	30%	16%	87%	14.3%	28%
	Pellet D	Resin D	-	100%	35%	65%	26%	8%

Table 3 .

 3 Fibre-reinforced composites compositions

	Sample Powder	Thermal treatment	Polymer used for the final PIP	Fraction of nitrided powder (Weight %)	Fraction of ceramic residue (Weight %)
	CMC1	TiSi 2	5 h, 1100°C, N 2	Resin A 1 PIP cycle	20	18
	CMC2	TiSi 2	5 h, 1100°C, N 2	Resin D 1 PIP cycle	16	17.3
	CMC3	TiSi 2	5 h, 1100°C, N 2	Resin D 3 PIP cycles	14	26

Table 4 .

 4 Flexural strengths of the CMCs prepared under various conditions (average of 2 specimens).

				Flexural
	Sample	Polymer	Porosity (%)	strength
				(MPa)
	CMC1	Resin A 1 PIP cycle	14.0	114 (± 20)
	CMC2	Resin D 1 PIP cycle	15.5	163 (± 20)
	CMC3	Resin D 3 PIP cycles	7.3	215 (± 25)

ACKNOWLEDGEMENTS

This work was supported by Herakles (Safran Group) and by the French national project NaCoMat.