P M Nigay
email: pmnigay@mines-albi.fr

T Cutard

A Nzihou

The impact of heat treatment on the microstructure of a clay ceramic and its thermal and mechanical properties

Keywords:

This paper presents the results of an experimental study on the microstructure, the thermal and the mechanical properties of a clay-based ceramic used in building applications. The X-ray tomography analysis showed a layered microstructure of clay with 200 µm sheets of porosity after the extrusion process. The gas release from the dehydration, dehydroxylation and decarbonation induced a 7 vol% formation of porosity during the heat treatment of the clay-based ceramic up to 850 °C. The porosity increase and the development of metakaolin led to a 38% decrease in the thermal conductivity. On the other hand, the Young's modulus of the clay-based ceramic was conserved due to the formation of smaller pores than the 200 µm sheets of porosity. The densification and the crystallization of amorphous phases also led to a 110% increase of the Young's modulus from 850 °C to 1050 °C. The Young's modulus of the clay-based ceramic was only decreased by the β→α quartz inversion of the cooling due to sand addition. Hence, this study provided a useful insight into how the microstructure of fired clay bricks can be specifically transformed by the porosity during the heat treatment to control the thermal and mechanical properties.

Introduction

Terracotta is a noble material used in the fields of sculpture, earthenware and construction since millennia [START_REF] Carty | Porcelain-raw materials, processing, phase evolution, and mechanical behavior[END_REF]. The expression changed over the years to describe the building materials such as fired clay bricks. Fired clay bricks are valuable materials due to their resistance and insulation performances. However, the combination of thermal and mechanical properties is the result of a long process.

The raw clay extracted from sedimentation pools is mixed with sand and water. The plastic mixture is shaped in the form of bricks using an extrusion process [START_REF] Krakowiak | Multitechnique investigation of extruded clay brick microstructure[END_REF]. The water is also removed by a drying process at temperatures up to 100 °C. The drying is promoted by the large particle size of the sand. The clay bricks experience afterwards a heat treatment at temperatures of 800 °C up to 1200 °C. In fact, the heat treatment temperatures are adapted to the purpose of the application [START_REF] Jordan | Influence of firing temperature and mineralogical composition on bending strength and porosity of ceramic tile bodies[END_REF]. The clay bricks are heat treated at temperatures around 800 °C to obtain high thermal properties. On the other hand, the fired clay bricks with high mechanical properties are obtained by heat treatment at temperatures around 1200 °C. This means that the fired clay-bricks present a trade-off between their thermal and mechanical properties. The trade-off comes to an issue since the fired clay bricks are used in building applications for resistance and insulation purposes at the same time [START_REF] Dondi | Recycling of industrial and urban wastes in brick productiona review[END_REF].

In recent years, many studies have shown the impact of the heat treatment temperature on the properties of fired clay bricks [START_REF] Müller | The influence of temper shape on the mechanical properties of archaeological ceramics[END_REF][START_REF] Ten | Thermal conductivity of traditional ceramics. Part I: influence of bulk density and firing temperature[END_REF]. The apparent trade-off between their thermal and mechanical properties was related to the microstructure. The clay bricks heat treated at around 800 °C presented a high percentage of porosity. The low thermal conductivity of air in these pores (0.026 W/m/K) leads to the acquisition of high thermal properties [START_REF] Smith | Thermal conductivity of porous materials[END_REF]. A recent study demonstrated that thermal properties are also enhanced by the anisotropy of the fired clay bricks [START_REF] Bourret | Anisotropy of thermal conductivity and elastic properties of extruded clay-based materials: evolution with thermal treatment[END_REF]. In fact, the heat conduction is highly limited through the plane of extrusion. The fired clay bricks present higher thermal properties in the direction of the thermal gradient of a wall. However, the literature indicates that the mechanical properties of the clay bricks are limited in the case of a heat treatment at around 800 °C. The low temperatures prevent from the transformation of clay minerals in thermally stable phases [START_REF] Andreev | Ceramic item deformation during firing: effects of composition and microstructure (review)[END_REF]. Moreover, the mechanical properties are restrained by the presence of pores acting as defects in the microstructure [START_REF] Kim | Clayware mechanical properties porosity dependent[END_REF].

According to the literature, the heat treatment of clay bricks at around 1200 °C leads to the acquisition of high mechanical properties. The sintering induces a consolidation of the clay bricks by a formation of necks between the particles [START_REF] Monteiro | Solid state sintering of red ceramics at lower temperatures[END_REF]. The mechanical properties are also enhanced by the densification of the fired clay bricks. In fact, the densification leads to an elimination of pores acting as defects in the microstructure [START_REF] Pialy | Effects of densification and mullitization on the evolution of the elastic properties of a clay-based material during firing[END_REF]. Some other studies attributed the acquisition of high mechanical properties to the transformation of the clay minerals into stable minerals such as mullite [START_REF] Freyburg | Influence of the clay type on the pore structure of structural ceramics[END_REF]. However, the literature has shown that the thermal properties of the fired clay bricks are limited in the case of a heat treatment at temperatures around 1200 °C. The elimination of pores reduces the proportion of air in the fired clay bricks, moreover, the thermal properties are limited by the progressive reduction of the anisotropy with the heat treatment [START_REF] Rhee | Porosity-thermal conductivity correlations for ceramic materials[END_REF].

The literature on the microstructure and the properties of the fired clay bricks with the heat treatment temperature is consistent. However, the research focuses on the characterization of the thermal and mechanical properties after heat treatment. The objective of this study is to investigate how the microstructure of fired clay bricks can be transformed during the heat treatment to control the thermal and mechanical properties. Hence, the study starts from the microstructure of a clay-based ceramic after the extrusion process. The experiments show the impact of the dehydration, dehydroxylation and decarbonation on the porosity of the clay-based ceramic during the heat treatment. The evolution of the porosity is also related to the evolution of the thermal and mechanical properties of the clay-based ceramic with the heat treatment temperature. The impact of the percentage of porosity on the thermal properties is separated from the effect of the solid skeleton using the Landauer's model. The impact of the solid skeleton is further confirmed by the quantities of mineral phases in relation with their properties obtained from the literature. Hence, this study shows how the microstructure of fired clay bricks can be specifically transformed by the porosity during the heat treatment to control the thermal and mechanical properties.

Materials and methods

Production of the clay-based ceramic

The clay-based ceramic investigated in this research was made of clay, sand and water. The clay was extracted as one batch in a clay quarry to ensure homogeneity. The clay was also ground at the laboratory with a 1 mm rolling. The clay-based ceramic mixture was prepared in a 70/30 wt% proportion introducing 8.4 kg of clay and 3.6 kg of sand in a kneading bowl. The 12 kg batch had a kneading for 10 min with an addition of water up to 17 wt%. The mixture was after extruded with an 8bars pressure to form a clay-based ceramic of 180×80×18 mm 3 . The clay-based ceramic experienced a drying at 25 °C, 65 °C and 105 °C for 24 h at each temperature in an electrical oven. The samples used for the experiments of this study were cut in the dried ceramic and rectified by polishing. The samples analyzed after heat treatment were heated in an electrical furnace (Nabertherm Controller C320) at different temperatures.

Determination of the composition

The inorganic elements of the clay, of the sand and of the clay-based ceramic mixture were measured with an ICP-AES instrument (Jobin Yvon Ultima 2). The materials were dissolved in a mixture of perchloric and hydrofluoric acid heated at 80 °C for 30 min. The solutions were diluted 10 times in purified water before the analyses. The concentrations of the inorganic elements obtained are reported in Table 1 as equivalent oxides. The aluminum and silicon oxides which are the main oxides of the clay minerals are presented separately. The other oxides include the calcium, magnesium and iron oxides. The distribution of oxides was placed on a ternary diagram of composition with common minerals of the terracotta. The position of the oxide equilibrium provided a prediction of the mineralogical transformations.

The mineralogical study was completed by X-ray diffraction (XRD) analyses using a Phillips Panalytical X′pert Pro MPD diffractometer.

The samples were analyzed as powders after heat treatment at 30 °C, 600 °C, 800 °C, 900 °C and 1000 °C with a 5 °C/min heating rate. Data were collected on a range of 2θ from 5°to 65°for each 0.017°with a CuKα radiation source (λ=1.5405 Å). The same automatic treatment of the background was applied to the scan of the different samples in the X′Pert HighScore Plus software. The identification of the minerals was carried out matching the ICDD database lines to the diffraction peaks. The content of minerals was measured by the RIR technique using the intensity of the diffraction peaks. A calibration was first performed using a 50/50 wt% mixture of the samples with an alumina reference. The quartz content was estimated comparing the intensity of its highest peak to the intensity of the highest peak of the reference. The intensity of the highest peak of the quartz phase was after compared to the intensity of the highest peak of other minerals in the samples.

Characterization of the microstructure

The heat flow and weight loss of the clay-based ceramic were measured as a function of the temperature using differential thermal analysis (DTA). The unfired samples were analyzed as 200 mg cylinders with a 5 mm diameter with a Setaram 92 instrument. Data were collected under air from 30 °C to 1100 °C at a heating rate of 5 °C/min. The deformation of the clay-based ceramic was also measured as a function of the temperature using thermo-mechanical analyses (TMA). The unfired samples were analyzed in the form of 200 mg cylinders with a 5 mm diameter with a Setsys 16/18 apparatus. Data were collected under air from 30 °C to 1100 °C at a heating rate of 5 °C/min. The same conditions were used for TMA and DTA to provide a consistent estimate of the bulk density [START_REF] Bailliez | The kinetics of surface area reduction during isothermal sintering of hydroxyapatite adsorbent[END_REF]. In fact, the bulk density of the clay-based ceramic (ρ (T) ) was calculated as a function of the temperature from Eq. (1) using the weight loss, the shrinkage, the initial weight (m 0 ), the initial length (l 0 ) and the initial radius (r 0 ) of the samples.

ρ = m (1 -Weight Loss ) πr l (1 -Shrinkage ) (T) 0 (T) 0 2 0 (T) 3 (1)
The transposition of the uniaxial deformation in a volume deformation supposes an isotropic behavior of the clay-based ceramic. The claybased ceramic that was investigated in this study presented an anisotropic behavior. However, the evolution of the bulk density was controlled by the evolution of the weight loss due to an insignificant evolution of the deformation. The evolution of the weight loss is independent of the direction and disposes of this isotropic condition.

The porosity of the clay-based ceramic (ε (T) ) was calculated as a function of the temperature from Eq. (2). The calculations were carried out using the bulk density of the clay-based ceramic (ρ (T) ) and the true density of the particles (ρ True ). The effect of the modifications in the mineral phases of the clay-based ceramic was included in the evolution of the bulk density with the temperature. The true density of the particles was then not depending on the heat treatment temperature. In fact, the true density was only measured by Helium pycnometer analyses after a heat treatment of the clay-based ceramic at the maximum temperature of 1100 °C and a fine grinding to eliminate the porosity. The true density of the particles was equal to 2.72 g/cm 3 . In this way, the percentage of porosity was determined in a continuous way at temperatures between 30 °C and 1100 °C.

ε = 1- ρ ρ (T) (T) True (2)
The microstructure of the clay-based ceramic was observed by Xray tomography using the Synchroton facility at ESFR Grenoble. The samples were analyzed as 200 mg cylinders after heat treatment at 950 °C. The micrograph of the cross-section shows the solid skeleton in light grey colours and the porosity in a black colour. The cross-section scans were also used to build a 3D representation of the porosity using imaging software. The 3D representation of the porosity shows the porosity of the clay-based ceramic in a light grey colour.

Determination of the thermal properties

The thermal properties of the clay-based ceramic were measured by a transitory plane source method (Hot Disk AB TPS 2500). Samples were analyzed as twins of 30×30×5 mm 3 using a nickel probe (3.189 mm radius). The thermal conductivity was calculated from the thermal response of the samples with the dimension and the power of the probe. The analyses were performed at room temperature after heat treatment at 30 °C, 200 °C, 400 °C, 600 °, 700 °C, 800 °C, 900 °C, 1000 °C and 1100 °C using a 5 °C/min heating rate. The values of the thermal conductivity after heat treatment at 250 °C, 450 °C, 650 °C and 850 °C were also calculated by linear regression analyses. In this way, the thermal conductivity could be compared to the percentage of porosity on the exact ranges of evolution with the temperature. The ranges were determined from the measurements of the weight loss, of the deformation, of the heat flow and of the percentage of porosity which are continuous with the temperature.

Determination of the mechanical properties

The mechanical properties of the clay-based ceramic were measured by a resonant frequency analyzer (IMCE HT650). The samples of 60×30×5 mm 3 were tested in a flexural mode. The Young's modulus was calculated from the fundamental resonance frequency and the samples dimensions. The analyses were performed with the temperature by introduction of the testing facility in an electrical furnace. The Young's modulus of the clay-based ceramic was then measured in a continuous way from 30 °C to 1050 °C.

Results and discussion

Mineralogical transformations

The oxide distributions of the clay, of the sand and of the clay-based ceramic mixture are displayed on the ternary diagram of composition in Fig. 1. The diagram also shows the minerals which are likely to form during the heat treatment. According to the literature, the clay minerals are transformed into an Al/Si spinel phase (Si 3 Al 4 O 12 ). The phase is expected to transform at high temperatures into a more stable phase of mullite (Si 2 Al 6 O 13 ). However, the presence of minerals such as calcium carbonates increases the concentration of other oxides than the aluminum and silicon oxides. The calcium carbonates induce a shift in the oxide equilibrium of the clay on the ternary diagram of composition. Hence, the clay minerals are more likely to transform at high temperatures into calcium-rich crystallites such as gehlenite (Ca 2 Al 2 SiO 7 ). The recombination with the calcium oxides would start at the interface of the clay minerals and calcium carbonates. However, the clay minerals could also transform into wollastonite (CaSiO 3 ) in order to reach the oxide equilibrium.

The sand added to the clay-based ceramic for its thermal stability show an oxide distribution close to the silica pit (SiO 2 ). However, a combination with the oxides from the clay would modify the oxide equilibrium of the sand. The minerals from the sand could transform into albite (NaAlSi 3 O 8 ), which is the sodium pole of plagioclase. Anorthite (CaAl 2 Si 2 O 8 ) which is the calcium pole of plagioclase could form at higher temperatures with a greater amount of energy. In fact, the oxide distribution of the clay-based ceramic mixture corresponds to the 70 wt% oxides of the clay and the 30 wt% oxides of the sand. The transformations of the clay-based ceramic are then unlikely to be the common transformations of the clay minerals. The energy supplied by the temperature should activate the transformations leading the clay and other minerals towards the global oxide equilibrium of the claybased ceramic mixture.

The X-ray diffraction (XRD) patterns of the clay-based ceramic are displayed after heat treatment at different temperatures in Fig. 2. The diffraction peaks show a presence of clay minerals after heat treatment at 30 °C. The clay-based ceramic is composed for 26 wt% of halloysite (Al 2 Si 2 O 5 (OH) 4 ) from the kaolinite-group. The illite-group minerals are present at 18 wt% with a regular illite (KAl 3 Si 3 O 10 (OH) 2 ) and an illite stratified by montmorillonite layers. Furthermore, the XRD pattern of the clay-based ceramic shows a 12 wt% content of calcium carbonates in the form of calcite (CaCO 3 ). The minerals from the sand are also present with 26 wt% of quartz (SiO 2 ) and 13 wt% of alkali feldspar (KAlSi 3 O 8 ).

The XRD pattern after heat treatment at 600 °C explains the transformations of the clay minerals during the heat treatment. The temperature increase leads to an elimination of the diffraction peak associated with halloysite (e.g., at 12°). In fact, the heat treatment induces a dehydroxylation of the kaolinite-group minerals from around 450 °C. The split of the hydroxyl groups that assured a cohesion between the tetrahedral and the octahedral sheets induces a collapse of the clay layers [START_REF] Aras | The change of phase composition in kaolinite and illite-rich clay-based ceramic bodies[END_REF]. Therefore, the dehydroxylation leads to an amorphization of halloysite into metakaolin (Al 2 Si 2 O 7 ) along with a water release as shown in Eq. (3).

s s g Al Si O (OH) ( ) → Al Si O ( ) + 2H O( ) 2 2 5 4 2 2 7 2 (3)
On the other hand, the clay minerals of the illite-group experience a dehydroxylation from 500 °C to 600 °C. The diffraction peaks associated with the illite minerals remain on the XRD patterns up to 800 °C. In fact, the dehydroxylation of illite starts by a condensation of hydroxyl groups into water molecules [START_REF] Gualtieri | Kinetics of illite dehydroxylation[END_REF]. The condensation of hydroxyl groups leaves an atom of oxygen between the tetrahedral and the octahedral sheets of the clay layers. This relocation avoids a collapse of the structure in contrast with kaolinite-group minerals. The dehydroxylation process only leads to a transformation of illite-group minerals in anhydrous illites (KAl 3 Si 3 O 11 ) along with a water release as

The transformations are related at higher temperatures to the clay minerals and the calcium carbonates. In fact, the temperature increase leads to a decarbonation of calcite from 650 °C to 850 °C. The diffraction peaks of calcite (e.g., at 29°) are eliminated after heat treatment at 900 °C. However, no diffraction peak associated with a formation of lime is observed on the XRD patterns. The calcium oxides of calcite are combined with the residual aluminum and silicon oxides of the dehydroxylation [START_REF] Peters | Mineralogical changes during firing of calcium-rich brick clays[END_REF]. Hence, the decarbonation transforms calcite and anhydrous illites into gehlenite (Ca 2 Al 2 SiO 7 ). The transformation induces a release of carbon dioxide (CO 2 ), amorphous silica (SiO 2 ) and alkalis (K 2 O) as shown in Eq. ( 5). (5)

The release of silicon and potassium oxides induces a convergence of the minerals towards the oxide equilibrium. The calcium oxides combine afterwards with the amorphous silica to form wollastonite (CaSiO 3 ) [START_REF] Cultrone | Carbonate and silicate phase reactions during ceramic firing[END_REF]. It explains the formation of wollastonite with a delay on the formation of gehlenite. The diffraction peaks associated with wollastonite are observed on the XRD patterns after heat treatment at 900 °C. This transformation occurs at the end of the decarbonation with a release of carbon dioxide as shown in Eq. ( 6). The clay and calcium carbonates are eventually combined after heat treatment at 900 °C. The 18 wt% of illite and 12 wt% of calcite produce 21 wt% of gehlenite and 13 wt% of wollastonite.

s s s g CaCO ( ) + SiO ( ) → CaSiO ( ) + CO ( ) 3 2 3 2 (6)
The composition of the clay-based ceramic is modified by the recrystallization of kaolinitic-residues at higher temperatures. The halloysite obtains a sufficient amount of energy to reach an organization after the dehydroxylation from 450 °C [START_REF] Sonuparlak | Spinel phase formation during the 980°C exothermic reaction in the kaolinite-to-mullite reaction series[END_REF]. The release of the last hydroxyl groups initiates a topotactic reaction at 950 °C. The transformation is observed on the XRD patterns with the diffraction peaks of the Al/Si spinel phase by 1000 °C. The 23 wt% of Al/Si spinel phase at this temperature matches with the 26 wt% of halloysite prior to the heat treatment. However, no diffraction peaks associated with mullite can be observed on the XRD pattern. This mineral is unlikely to form due to the specific equilibrium of oxides.

s s s 2Al Si O ( ) → Si Al O ( ) + SiO ( ) 2 2 7 3 4 12 2 (7)
The mineralogical transformations of the clay-based ceramic are related to the minerals from the sand at temperatures higher than 1000 °C. The diffraction peaks associated with albite (NaAlSi 3 O 8 ) and anorthite (CaAl 2 Si 2 O 8 ) appear on the XRD pattern as a result of the alkali feldspar degradation. The 3 wt% of albite and anorthite indicates a minor degradation of the 13 wt% of alkali feldspar at 1000 °C [START_REF] Krichen | Mineralogy and sintering mechanisms of a clay raw material from Kasrine[END_REF]. The clay-based ceramic is likely modified by a gas release associated with the mineralogical transformations. The clay-based ceramic is also modified by a solid-state sintering due to the absence of a liquid phase formation. In fact, the heat treatment provides some energy to the temperature-sensitive minerals of the clay-based ceramic mixture to reach the global equilibrium of oxides in a form of stable minerals.

Impact of the mineralogical transformations on the microstructure

The evolutions of the weight loss and of the deformation of the claybased ceramic during the heat treatment are displayed in Fig. 3. The heat treatment induces a dehydration of the clay-based ceramic from 30 °C to 250 °C. The dehydration leads to a 2 wt% elimination of water remaining in the pores of clay particles from 30 °C to 100 °C. On the other hand, the 1 wt% elimination of water adsorbed on clay particles induces a shrinkage of the clay-based ceramic from 100 °C up to 250 °C. The temperature increase also leads to a dehydroxylation of the clay minerals. The dehydroxylation of the kaolinite-group mineral induces a 1 wt% elimination of constitution water at around 450 °C. Furthermore, a 1 wt% elimination of water is associated with the dehydroxylation of illites from 500 °C to 600 °C. The temperature increase leads to a α→β quartz inversion at around 600 °C [START_REF] Tarvornpanich | Microstructural evolution in clay-based ceramics II: ternary and quaternary mixtures of clay, flux, and quartz filler[END_REF]. The quartz inversion induces an expansion of the quartz particles during the heating of the clay-based ceramic. Furthermore, the shrinkage of the quartz particles results in a formation of defects during the β→α quartz inversion of the cooling. The heat treatment of the clay-based ceramic also leads to the decarbonation of the calcium carbonates. The decarbonation induces a 4 wt% elimination of calcite from 650 °C to 850 °C. On the same range of temperature, the clay-based ceramic experiences a solid-state sintering. The formation of necks between the particles leads to a shrinkage of the clay-based ceramic. However, the decarbonation limits the impact of the sintering to a consolidation of the clay-based ceramic up to 850 °C. The sintering only leads to a densification on a range of temperatures from 850 °C to 1100 °C. The evolution of the porosity during the heat treatment of the claybased ceramic is displayed in Fig. 4. The clay-based ceramic presents a porosity of 24 vol% after heat treatment at 30 °C. The air remaining in the mixture is trapped in the clay-based ceramic as pores during the extrusion process. Furthermore, the porosity is increased by the water release during the drying process. The heat treatment of the clay-based ceramic induces a 2 vol% formation of porosity during the dehydration. The porosity formation can be related to the release of water remaining in the pores of clay particles from 30 °C to 100 °C. Furthermore, the formation of porosity with the release of water absorbed on clay particles is limited by the shrinkage from 100 °C to 250 °C. The dehydroxylation of the clay minerals also leads to a 2 vol% formation of porosity between 450 °C and 600 °C. The porosity formation can be related to the release of constitution water from the kaolinite-group and illite-group minerals. The release of carbon dioxide associated with the decarbonation of calcite induces a 3 vol% formation of porosity from 650 °C to 850 °C. Therefore, the sintering does not reduce the porosity on this range of temperatures. The consolidation of the claybased ceramic consists in a coalescence of pores instead of an elimination [START_REF] Ducman | Frost resistance of clay roofing tiles: case study[END_REF]. The sintering only leads to a 2 vol% reduction of the porosity during the densification from 850 °C to 1100 °C. Hence, the clay-based ceramic conserves a 29 vol% of porosity at the highest heat treatment temperature of 1100 °C.

The microstructure of the clay-based ceramic after heat treatment at 950 °C is displayed in Fig. 5. The cross-section micrograph shows a heterogeneous microstructure of clay with pores, feldspar and quartz particles. The 200 µm pores orientated in a same direction result from the presence of air in the mixture during the extrusion process. In fact, the extrusion of process leads to a compression of the clay in the form of sheets. The air remaining in the mixture is then compressed between the sheets of clay in a form of porosity sheets. The set of porosity sheets also results in the formation of a layered microstructure, as shown on the 3D representation. Hence, the clay-based ceramic develop an anisotropic behavior as soon as the extrusion process. It is important to note that the layered microstructure is conserved after heat treatment of the clay-based ceramic at a temperature of 950 °C. The formation of pores during the mineralogical transformations does not eliminate the anisotropy. The anisotropic behavior associated with the porosity sheets is not either eliminated during the sintering. In fact, the sintering of the clay-based ceramic decreases the percentage of porosity between 850 °C and 1100 °C by an elimination of the smallest pores. The pores generated during the mineralogical transformations are then the smallest pores of the clay-based ceramic. The cross-section micrograph also shows the white particles of feldspar and angular particles of quartz. It can be noted a presence of large defects around the quartz particles. In fact, the defects are due to the larger shrinkage of quartz than clay particles during the β→α quartz inversion of the cooling. These defects associated with the quartz particles from the sand represent the largest defects in the microstructure with the porosity sheets of the extrusion process.

Impact of the microstructure evolution on the thermal properties

The evolution of the thermal conductivity with the heat treatment temperature is shown in Fig. 6. The clay-based ceramic presents a thermal conductivity of 1.07 W/m/K after heat treatment at 30 °C. This value represents a combination between the relatively high thermal conductivity of the solid skeleton and the relatively low thermal conductivity of air in the porosity sheets. The contribution of the solid skeleton can be determined from the Landauer's model [START_REF] Landauer | The electrical resistance of binary metallic mixtures[END_REF], reexpressed in Eq. ( 8) where λ is the measured thermal conductivity of the clay-based ceramic, λ s is the thermal conductivity of the solid skeleton, λ p is the thermal conductivity of air in the porosity, and v p is the pore volume fraction obtained from the percentage of porosity. It is important to note that the thermal conductivity of the solid skeleton is estimated at room temperature with the values of the pore volume fraction in temperature. This was only possible as the pore volume fraction of the clay-based ceramic at room temperature was equivalent to the pore volume fraction in temperature.

λ = 2λ -λλ (3v -1) λ + λ(2-3v ) s 2 p p p p (8)
According to the Landauer's model, the thermal conductivity of the solid skeleton is equal to 1.69 W/m/K after heat treatment of the claybased ceramic at 30 °C. On the other hand, the thermal conductivity of air is equal to 0.026 W/m/K. It is important to note that the impact of air on the thermal conductivity of the clay-based ceramic is enhanced in this direction of the thermal gradient of the walls by the anisotropic behavior of the porosity sheets.

The dehydration induces a 7% decrease of the thermal conductivity at temperatures between 30 °C and 250 °C. The thermal conductivity of the clay-based ceramic is decreased by the 2 vol% formation of porosity since the mineral phases are not transformed on this range of temperatures. In fact, the porosity formation reduces the proportion of solid skeleton with a relatively high thermal conductivity. The thermal conductivity of the clay-based ceramic is also decreased by the relatively low thermal conductivity of air in the newly-formed pores.

The dehydroxylation of the clay minerals also results in a 23% decrease of the thermal conductivity. The decrease of the thermal conductivity is related to the 2 vol% formation of porosity from 450 °C to 600 °C. According to the Landauer's model, the porosity formation represents a 4% decrease of the thermal conductivity. The thermal conductivity of the clay-based ceramic is then also decreased by a modification of the solid skeleton. The modification of the solid skeleton concerns the kaolinite-group mineral since the clay minerals from other groups only recombine with calcite at temperatures between 650 °C and 850 °C. In fact, the kaolinite-group mineral with a thermal conductivity approximated to the 2.80 W/m/K one of kaolinite [START_REF] Midttømme | Thermal conductivity of selected claystones and mudstones from England[END_REF], transform from 450 °C to 600 °C into metakaolin with a thermal conductivity of 0.39 W/m/K [START_REF] Michot | Thermal conductivity and specific heat of kaolinite: evolution with thermal treatment[END_REF]. The 26 wt% formation of metakaolin results in an extensive decrease of the thermal conductivity of the solid skeleton. According to the Landauer's model, the thermal conductivity of the solid skeleton is decreased from 1.55 W/m/K to 1.23 W/m/K. This modification of the solid skeleton represents a 20% decrease in the thermal conductivity of the clay-based ceramic.

The decarbonation of calcite does not significantly modify the thermal conductivity of the clay-based ceramic. The 3 vol% formation of porosity results, according to the Landauer's model, in a 6% decrease of the thermal conductivity from 650 °C to 850 °C. However, the impact of the porosity formation is prevented by the modification of the solid skeleton. The illite-group minerals and calcite, with thermal conductivities of 1.80 W/m/K and 3.40 W/m/K [START_REF] Midttømme | Thermal conductivity of selected claystones and mudstones from England[END_REF], combine on this range of temperatures into gehlenite and wollastonite, with thermal conductivities of 1.53 W/m/K and 4.03 W/m/K [START_REF] Horai | Thermal conductivity of rock-forming minerals[END_REF]. The 21 wt% formation of gehlenite and 13 wt% formation of wollastonite induces a significant increase in the thermal conductivity of the solid skeleton. According to the Landauer's model, the thermal conductivity of the solid skeleton is increased from 1.25 W/m/K to 1.33 W/m/K which represents a 6% increase in the thermal conductivity of the clay-based ceramic.

The sintering of the clay-based ceramic induces a 24% increase of the thermal conductivity from 850 °C to 1100 °C. The thermal conductivity of the clay-based ceramic is increased by the 2 vol% reduction of the porosity. The lower proportion of air represents, according to the Landauer's model, a 5% increase of the thermal conductivity. The thermal conductivity of the clay-based ceramic is also increased by the modification of the solid skeleton. In fact, the metakaolin with a thermal conductivity of 0.39 W/m/K [START_REF] Michot | Thermal conductivity and specific heat of kaolinite: evolution with thermal treatment[END_REF] recrystallizes on this range of temperature into an Al/Si spinel phase with a thermal conductivity of 0.56 W/m/K [START_REF] Michot | Thermal conductivity and specific heat of kaolinite: evolution with thermal treatment[END_REF]. The 23 wt% formation of Al/Si spinel leads to an extensive increase in the thermal conductivity of the solid skeleton. According to the Landauer's model, the thermal conductivity of the solid skeleton is increased from 1.33 W/m/K to 1.58 W/m/K. This modification of the solid skeleton represents a 19% increase in the thermal conductivity of the clay-based ceramic. Hence, the clay-based ceramic has a thermal conductivity of 0.87 W/m/K after heat treatment at 1100 °C.

Impact of the microstructure evolution on the mechanical properties

The evolution of the Young's modulus during the heat treatment of the clay-based ceramic is displayed in Fig. 7. The clay-based ceramic presents a Young's modulus of 7.5 GPa after heat treatment at 30 °C. The stiffness is obtained by formation of a solid skeleton during the extrusion process. However, the stiffness is also restrained by the 200 µm sheets of porosity acting as defects.

The heat treatment of the clay-based ceramic induces a 76% increase of the Young's modulus during the dehydration. The porosity formation could decrease the Young's modulus from 30 °C to 250 °C. However, the positive impact of the shrinkage exceeds the negative impact of the pores on this range of temperatures.

The dehydroxylation of the clay minerals does not induce any modification of the Young's modulus at temperatures between 450 °C and 600 °C. The Young's modulus of the clay-based ceramic is not increased on this range of temperatures due the amorphization of the kaolinite-group mineral as well as the absence of shrinkage. The Young's modulus is not decreased either by the newly-formed pores acting as defects. In fact, the dehydroxylation of the clay minerals results in a formation of small pores compared to the 200 µm porosity sheets of the extrusion process. This result suggests that the formation of smaller pores than the previous defects of the microstructure allows a conservation of the Young's modulus.

The decarbonation of the calcium carbonates does not result in any modification of the Young's modulus at temperatures between 650 °C and 850 °C. The Young's modulus of the clay-based ceramic is increased by the formation of stable minerals such as gehlenite and wollastonite. The consolidation of the clay-based ceramic also leads to an increase of the Young's modulus on this range of temperatures. But, the positive impact of the calcium-rich crystallites along with the formation of necks between the clay particles is prevented by the coalescence of pores. This increase in the pore size of the clay-based ceramic enhances the negative impact of the 3 vol% formation of porosity on the Young's modulus.

The sintering only increases the Young's modulus after the decarbonation of the calcium carbonates. The densification of the claybased ceramic along with the crystallization of metakaolin into an Al/Si spinel phase induces a 110% increase of the Young's modulus at temperatures between 850 °C and 1050 °C. Hence, the clay-based ceramic presents a Young's modulus of 27.9 GPa at the highest heat treatment temperature.

However, the Young's modulus is decreased by 41% during the cooling of the clay-based ceramic. The decrease in the Young's modulus can be related to the β→α quartz inversion at around 600 °C. In fact, the transition leads to a formation of large defects around the quartz particles from the sand. This means that a formation of defects as large as the 200 µm sheets of porosity have a predominant impact on the Young's modulus. Hence, the clay-based ceramic presents a Young's modulus of 16.4 GPa after the cooling.

The sand is added in the clay-based ceramic mixture to promote a fast elimination of water during the drying process. However, a small reduction in the amount of sand could reduce the formation of defects at the quartz inversion. The clay-based ceramic could present a higher Young's modulus after the cooling.

Implications

The implications of this work are significant in the field of building materials. The building materials such as fired clay bricks are used for their thermal and mechanical properties. The literature has shown a major impact of the percentage of porosity. The insulating behavior of air in the pores provided high thermal properties to the fired clay bricks. Nevertheless, the mechanical properties were limited by the pores acting as defects in the microstructure. The fired clay bricks presented a trade-off between their thermal and mechanical properties.

The thermal properties of a clay-based ceramic were improved with a conservation of the mechanical properties in this study. The improvement was based on a formation of smaller pores than the defects of the microstructure. In fact, the pores generated during the heat treatment were smaller than large porosity sheets formed during the extrusion process. The porosity increase along with the development of metakaolin led to a 38% improvement of the thermal properties. The mechanical properties of the clay-based ceramic were only reduced during the β→α quartz inversion of the cooling. In fact, the decrease in the mechanical properties was caused by the quartz particles from the sand. The large grain size of the sand promotes a fast elimination of water during the drying process. But, a small reduction in the amount of sand could improve the mechanical properties of the fired clay bricks.

This research also provided information for a larger improvement of the properties using organic additives. The combustion of organic additives in fired clay bricks should lead to a porosity formation during the heat treatment. Hence, the porosity increase would improve the thermal properties of the fired clay bricks. The mechanical properties could also be conserved by addition of smaller particles than the defects of the microstructure. This information could widen the range of organic additives used in fired clay bricks to many organic wastes.

Conclusions

This study investigated the impact of heat treatment on the microstructure of a clay-based ceramic and its thermal and mechanical properties. The heat treatment of the clay-based ceramic induced a transformation of the temperature-sensitive minerals towards the global equilibrium of oxides. The clay mineral of the kaolinite-group transformed into amorphous metakaolin to recrystallize in Al/Si spinel phase. Furthermore, the illite-group minerals and calcite combined into calcium-rich crystallites such as gehlenite and wollastonite. The gas release associated with these transformations induced a porosity formation in the clay-based ceramic. The thermal properties of the clay-based ceramic were then improved by the low thermal conductivity of air in the newly-formed pores after a heat treatment at 850 °C. The improvement of the thermal properties was also due to the formation of metakaolin which decreased the thermal conductivity of the solid skeleton. On the other hand, the mechanical properties of the clay-based ceramic were conserved due to the formation of smaller pores than the defects of the microstructure. The densification of the clay-based ceramic along with the crystallization of amorphous phases also resulted in an improvement of the mechanical properties. The mechanical properties of the clay-based ceramic were only decreased during the β→α quartz inversion of the cooling as a result of the sand addition. Hence, this study provided a useful insight into how the microstructure of fired clay bricks can be specifically transformed by the porosity during the heat treatment to control the thermal and mechanical properties.

Fig. 1 .

 1 Fig. 1. Ternary diagram of composition with the clay, the sand and the clay-based ceramic mixture.

  KAl Si O (OH) ( ) → KAl Si O ( ) + H O( )

Fig. 2 .Fig. 3 .

 23 Fig. 2. X-ray diffraction patterns of the clay-based ceramic after heat treatment at 30 °C, 600 °C, 800 °C, 900 °C and 1000 °C.

Fig. 4 .

 4 Fig. 4. Evolution of the porosity and of the heat flow of the clay-based ceramic during the heat treatment.

Fig. 5 .Fig. 6 .Fig. 7 .

 567 Fig. 5. X-ray tomography micrograph of the cross-section (a) and 3D representation of the porosity (b) of the clay-based ceramic after heat treatment at 950 °C.

Table 1

 1 Oxide distribution of the main inorganic elements of the clay, of the sand and of the claybased ceramic mixture.

	Sample		Distribution of oxides (wt%)
		Al 2 O 3	SiO 2	Others
	Clay	16.8	44.8	38.4
	Sand	16.8	62.4	20.8
	Mixture	16.8	50.1	33.1

Acknowledgements

The authors would like to acknowledge the TERREAL Company for financial support and scientific contribution to this study. We also thank the staff of the ICA and RAPSODEE Centers at Ecole des Mines d′Albi for assistance with the research. Mr. Daniel Davidson and Mr. Wesley Awe are also thanked for proofreading this manuscript.

Funding

This study was funded by Agence Nationale de la Recherche et de la Technologie (CIFRE N°2011/1193).

Conflict of interest

The authors declare that they have any conflict of interest.


Thierry Cutard
email: thierry.cutard@mines-albi.fr

Nicolas Donval
email: bnicolas.donval@mines-albi.fr

Aurélien Mazzoni
email: caurelien.mazzoni@mines-albi.fr

Claire Michel
email: claire.michel@mines-albi.fr

Fabien Nazaret
email: enazaret@aurock.fr

Thermomechanical characterization of monolithic refractory castables

Keywords: refractory castable, mechanical behavior, microstructure, damage, creep, high temperature and experiments

Introduction

Refractory monolithic castables are characterized by a complex thermomechanical behavior. This is mainly due to their heterogeneity and to the large contrast which exists between the properties of their various constituents. This complexity deals with an asymmetric behavior in tension and in compression, with a high sensitivity of the thermomechanical behavior to the material thermal history and to the considered temperature level, with the various microstructural transformations that take place in such materials and with the non-linear behavior of monolithic castables. Complementary studies have been conducted on the thermomechanical behavior of such materials. Their results allow proposing a synthetic diagram to describe the evolutions of their behavior as a function of the considered temperature level. Such a description is given in figure 1 where three temperature domains have been defined. In the low temperature domain, the behavior is mainly a damage elastic one whereas the behavior is elasto-viscoplastic in the high temperature domain. These two domains are separated by a transition domain whom boundaries can vary in the 800-1000°C temperature range, depending on the material formulation.

Because of the complexity of the monolithic castable thermomechanical behavior, complementary tests must be carried on for its better characterization and for its better understanding. In the present paper, results of various mechanical tests are considered. They deal with tension, compression and bending tests. On the one hand, in some cases, the digital image correlation method has been coupled to mechanical tests in order to obtain displacement or strain fields, which allow completing the behavior curves obtained on test samples. DIC results are particularly helpful in following the development of damage processes at room temperature. On the other hand, such macroscopic observations have been completed with results of microscopic ones.

Optical and scanning electron microscopies allow linking microstructural mechanisms and thermomechanical behaviors in a wide temperature range. As indicated in figure 1, the behavior of monolithic castables remains a quasi-brittle one in a wide temperature range. Fiber reinforcement is one way to modify highly the monolithic castable behavior. That is why some results of studies conducted on short fiber reinforced refractory monolithic castables are considered too.

Devices to characterize refractory castables

Figure 2 presents some complementary test types for characterizing the thermomechanical behavior of monolithic castables. Because of the asymmetric behavior in tension and in compression of such materials, performing both tension and compression tests is the most appropriate way. But such tests remain difficult to be carried on, particularly tension tests [START_REF] Kakroudi | Effect of thermal treatment on damage mechanical behavior of refractory castables : Comparison between bauxite and andalusite aggregates[END_REF][START_REF] Nazaret | Influence of the mechanical behaviour specificities of damaged refractory castables on the Young's modulus determination[END_REF]. Performing three or four points bending tests is easier but one have to remember two main difficulties of bending tests: one deals with the difficulty to measure the bending deflection of the sample without introducing the behavior of loading contacts in the measurements [START_REF] Nazaret | Influence of the mechanical behaviour specificities of damaged refractory castables on the Young's modulus determination[END_REF], the other one deals with the difficulty to identify parameters of behaviors laws. This last point is due to the heterogeneous stress state of bending samples, to the asymmetric behavior of castables and to stress redistribution phenomena due to damage.

-a --b --c - Figure 3 gives some details on complementary test types that allow obtaining complementary results compared to mechanical test results. Considering a macroscopic scale, the DIC method is helpful to follow the development of damage processes at room temperature and during a tensile test. The diffuse damage step, the localization and the macrocraking initiation phenomena can be studied by considering strain or displacement fields. Links with the behavior curve can be done. At a microscopic scale, scanning electron microscopy observations are really useful to understand and to establish relationships between microstructural and thermomechanical changes. This is particularly true when considering the results of in-situ high temperature environmental scanning microscopy observations. Furthermore, dilatometry remains a method that allows identifying the temperature levels and ranges where microstructural changes take place. The different test types considered here (figures 2 and 3) have been described in details in previous papers [START_REF] Marzagui | Etude de deux bétons réfractaires silico-alumineux : microstructures et comportements thermomécaniques en traction et en flexion[END_REF][START_REF] Nazaret | Influence of the mechanical behaviour specificities of damaged refractory castables on the Young's modulus determination[END_REF][START_REF] Nazaret | Thermomechanical behaviour of a fibre-reinforced refractory concrete: tests and analysis[END_REF][START_REF] Marzagui | Characterization of microstructural evolutions in refractory castables by in situ high temperature ESEM[END_REF][START_REF] Robert | Use of Digital Image Stereo-Correlation (DISC) to characterize the mechanical behavior of refractory castable[END_REF].

Damage behavior of refractory monolithic castables in the low temperature domain

Case of non-reinforced refractory monolithic castables

For an andalusite based refractory concrete, examples of room temperature tensile test results are given in figure 4a. One have been performed on a 110°C dried sample whereas the other one was performed on a 700°C fired sample. As previously described [START_REF] Marzagui | Etude de deux bétons réfractaires silico-alumineux : microstructures et comportements thermomécaniques en traction et en flexion[END_REF][START_REF] Marzagui | Characterization of microstructural evolutions in refractory castables by in situ high temperature ESEM[END_REF], damage develops during the first thermal cycle at such a temperature, particularly during the cooling stage. An example of such a thermal damage is given in figure 5a. For the 110°C dried material, the tension behavior is elastic linear up to the maximum stress level. In the case of the 700°C fired material, the behavior curve is characterized by a very restricted linear domain at the test beginning. It moves then to a wide nonlinear domain, due to diffuse damage processes. This example allows observing the deformation capacity of such a material when the mechanical damage develops as a superposition to the thermal damage, due to firing. More than the non-linearity of the envelope curve, loading/unloading loops confirm the development of a diffuse damage. Such phenomena have already been observed in other refractory castables [START_REF] Kakroudi | Effect of thermal treatment on damage mechanical behavior of refractory castables : Comparison between bauxite and andalusite aggregates[END_REF][START_REF] Lemaistre | Etude des propriétés thermomécaniques de divers réfractaires[END_REF][START_REF] Marzagui | Etude de deux bétons réfractaires silico-alumineux : microstructures et comportements thermomécaniques en traction et en flexion[END_REF]. Furthermore, an increase of the permanent strain levels is observed too. It can be related to the unclosing of damage that has been created by the tensile loading as observed in figure 5b. In that case, damage deals with numerous aggregate/concrete debondings that are connected by matrix microcracks. They appear as being more opened when their orientation is perpendicular to the tensile loading direction. Strain fields, obtained from performing the DIC method during a room temperature test on the 700°C fired andalusite based monolithic castable, allow observing the development of the strain localization phenomenon and of the progressive initiation of a macrocrack that will lead to the sample failure (figure 6) [START_REF] Marzagui | Etude de deux bétons réfractaires silico-alumineux : microstructures et comportements thermomécaniques en traction et en flexion[END_REF]. Such a method makes possible to detect strain localization sites earlier to classical strain measurements, for example with contact extensometers. Results of in-situ ESEM observations are shown in figure 7. This micrograph sequence focuses on the evolutions of a microcrack that has appeared at 150°C and that can be related to dehydration processes A progressive microcrack opening occurs when increasing the temperature with an increase of the opening amplitude in the 900-1000°C temperature range [START_REF] Marzagui | Etude de deux bétons réfractaires silico-alumineux : microstructures et comportements thermomécaniques en traction et en flexion[END_REF][START_REF] Marzagui | Characterization of microstructural evolutions in refractory castables by in situ high temperature ESEM[END_REF]. Such observations really contribute to obtain a better understanding of damage mechanisms that occur in monolithics during heat treatments and due to their heterogeneity. The high temperature tension behavior can now be considered (figure 4b). Up to 700°C, similarities are observed between room and high temperature behaviors. This must be related to the fact that similar damage mechanisms act in both cases. Considering now the 900°C, the behavior curve underlines the transition that occurs at this temperature level. Compared to the 700°C behavior curve, the maximum stress level strongly decreases whereas the material strain capacity is considerably increased. Viscoplasticity is now present in the mechanical behavior of the monolithic castable and is due to the formation of local liquid phases; particularly in silica fume rich areas.

Case of fiber-reinforced refractory monolithic castables (FRRC)

Figure 8a and figure 8b show behavior curves in tension and in four points bending of 500°C fired samples of a geopolymer based FRRC [START_REF] Nazaret | Caractérisation et modélisation du comportement thermomécanique d'un béton réfractaire renforcé de fibres[END_REF]. Both test types were performed at 20°C and at 500°C. Before the peak stress, the non-linear part is due to the development of diffuse damage by microcracking processes. While the peak stress is exceeded, the behavior moves to a softening branch and finally to a constant load deformation domain. Such a mechanical behavior is characteristic of fiber reinforced concretes. In the case of direct tensile tests, the peak stress value decreases when considering the high temperature test results. Nevertheless, in the case of the 500°C testing temperature, the FRRC develops a high softening behavior and the decrease of the stress values in the post-peak domain is really low. Contrary to the tensile test where the maximum stress level decreases according to the testing temperature, in the case of the four points bending test, the peak stress remains stable for the different testing temperatures. As previously shown [START_REF] Marzagui | Characterization of microstructural evolutions in refractory castables by in situ high temperature ESEM[END_REF], this must be related to the to characteristics of the tension softening curve. Indeed, the softening slopes are stronger at high temperature during tensile tests. The tension behavior can be highly modified by increasing the strain capacity of the material. This can be done by extending the non-linear pre-peak domain where diffuse damage occurs. Another way consists in enhancing the load capacity of the material in the post peak domain.

Adding metallic fibers creates another heterogeneity degree in these materials. A metallic component, characterized by a high expansion coefficient compared to ceramic ones, is included in the microstructure. When subjecting a FRRC to thermal cycles, specific internal stress fields are induced. They result in specific damage mechanisms at fiber/concrete interfaces [START_REF] Cailleux | Study of a ceramic refractory reinforced with metallic fibres : from the microstructure to the mechanical behaviour[END_REF]. As shown in figure 9, they deal with radial matrix microcracking and with fiber/concrete debonding. The presented micrographs have been obtained after subjecting the FRRC to a first thermal cycle.

Heating again the FRRC will lead to a closing mechanism of fiber/concrete debondings and to the development of a fretting pressure at fiber/concrete interfaces Figure 10 presents some results of DIC measurements performed on FRRC samples: a four points bending one (Fig. 10a) and a technological one (Fig. 10b) [START_REF] Robert | Use of Digital Image Stereo-Correlation (DISC) to characterize the mechanical behavior of refractory castable[END_REF][START_REF] Nazaret | Caractérisation et modélisation du comportement thermomécanique d'un béton réfractaire renforcé de fibres[END_REF]. In both cases, strain fields illustrate the capacity of the reinforced material to develop stress redistribution phenomena. Even if one macrocrack is initiated, because of fiber bridging effects, the macrocrack propagation is stable and multicracking phenomena are able to develop.

Elasto-viscoplastic behavior of refractory monolithic castables in the high temperature domain

Case of non-reinforced refractory monolithic castables

Results of non-isothermal creep tests in four points bending are given on Figure 11a. Tests have been performed under a 2MPa stress level in the 20°C to 1200°C temperature range for a nonreinforced andalusite based castable (And-LCC). Under such conditions, creep starts to occur between 750°C and 800°C. This indicates that local liquid phases starts to have effects at such temperature levels. Complementary four points bending tests have been performed to quantify the effects of modifying the strain rate on the high temperature behavior. Figure 11b gives some results obtained at 900°C for 0,2mm/min and 0,02mm/min crosshead speeds [START_REF] Mazzoni | Thermomechanical behaviour of refractory concretes: effect of a mineral fibre reinforcement[END_REF]. For the And-LCC a high strain rate sensivity is observed. This confirms that a high viscoplasticity component characterizes its high temperature behavior. In-situ ESEM micrographs of figure 12 give some illustrations of drastic microstructural changes that take place in the 1000-1300°C temperature range for the considered And-LCC monolithic castable [START_REF] Marzagui | Etude de deux bétons réfractaires silico-alumineux : microstructures et comportements thermomécaniques en traction et en flexion[END_REF]. From 1100°C, a small amount of liquid phase appears and is accompanied by a progressive vanishing of the binder phase particles. The amount of liquid phase clearly increases at 1200°C where liquid phase bridges are observed. At 1400°C crack bridging is clearly observed. Such transformations have a large influence on the high temperature material behavior. At a given temperature, the thermomechanical behavior can highly depend on the monolithic thermal history. Figure 13 gives an example of such a dependence by comparing the 900°C four points bending behavior for two 900°C firing durations: 5 hours and 150 hours. After a long time firing (150h), the viscoplastic part of the behavior disappeared, due to crystallization phenomena that have taken place in the castable binder phase. At this temperature level, located in the transition domain, the behavior has moved from a highly non-linear one to a quite linear one.

Case of fiber-reinforced refractory monolithic castables

Figures 14a and14b show bending curves for a non-reinforced And-LCC monolithic castable and for a Nextel720 ® reinforced one (And-N7 grade) [START_REF] Mazzoni | Comportement thermomécanique d'un béton réfractaire : effet du renforcement par des fibres minerals[END_REF]. Tests have been performed at 700°C, 900°C, 1000°C and 1100°C. Reinforcing the And-LCC castable with Nextel720 ® fibers leads essentially to large modifications of the pre-peak behavior. They mainly deal with an increase of the width of the pre-peak linear domain and with an increase of the stress peak and deflection peak levels. Nextel720® fibers can also be considered as delaying diffuse damage phenomena and to limit the viscoplastic behavior of the concrete. Because of the high alumina content of Nextel720 ® fibers, the temperature range in which such phenomena can be obtained is largely increased. No benefit effect is observed in the post-peak domain. Due to the Nextel720 ® fiber reinforcement of the And-LCC refractory concrete, the increase of the peak stress levels for the 700°C, 900°C, 1000°C and 1100°C testing temperatures are respectively of 86,1%, 79%, 94,7% and 52,7%. On figure 15 are some macroscopic views of a non-reinforced and of a Nextel720 ® reinforced andalusite grades [START_REF] Mazzoni | Comportement thermomécanique d'un béton réfractaire : effet du renforcement par des fibres minerals[END_REF]. After testing at 900°C, results show that fibers allow delaying the initiation of a macrocrack and permit multicraking phenomena. Similar results are obtained at 1200°C concerning the effects of a fiber reinforcement. But even if macroscopic behaviors appear as being similar, deformation mechanisms and microscopic fiber effects are note the same. In the case of an interrupted 900°C bending test (figure 16), observing cross sections of the reinforced grade (And-N7) allows identifying damage which is made of a microcrack network highly influenced by the presence of fibers. Numerous microcrack deviations and bridging configurations can be seen. Debonding phenomena have extensively taken place both at matrix/aggregate interfaces and at matrix/fiber interfaces too. In the case of an interrupted 1200°C bending test (figure 17), observing cross sections of the reinforced grade (And-N7) allows identifying damage which is made of a microcrack network and of a high density of cavities in the castable matrix [START_REF] Mazzoni | Comportement thermomécanique d'un béton réfractaire : effet du renforcement par des fibres minerals[END_REF]. Viscous deformation phenomena have played a major role in the deformation mechanisms of the castable and are highly influenced by the presence of fibers. Microcrack deviations, microcrack bridging configurations can be seen again. Such results allow obtaining a better understanding of fiber effects on the high temperature strain mechanisms of monolithic castables.

Conclusion

Results considered in the present paper have shown the complexity of the high temperature mechanical behavior of monolithic refractory castables. Such a complexity comes from the high heterogeneity of non-reinforced and of reinforced materials. From a mechanical point of view, the complexity deals with the non-linearity of the behavior in a wide temperature range. Origins of this non-linearity depend on the temperature since microscopic mechanisms that are involved change between low and high temperature levels. From an elastic damage character in a low temperature domain, the behavior moves to an elasto-visco plastic one at high temperature. A transition domain has to be defined to relate the progressive transition between these two behavior types. In both cases, the behavior can drastically change with the material thermal history because of large microstructural changes. Another component of the complexity deals with the asymmetric behavior between tension and compression. Combining complementary mechanical test types is essential to obtain a well appropriate characterization and knowledge of a refractory castable behavior. Considering results of methodologies that allow obtaining details on microstructural evolutions in the material is of main importance too. Microstructural evolutions are complex too. They can deal with phase changes, densification, and damage phenomena.

Fig. 1 .

 1 Fig. 1. Main evolutions of microstructures and thermomechanical behavior of monolithic castables at high temperature [1 to 7].

Fig. 2 .

 2 Fig. 2. Complementary test types for characterizing the thermomechanical behavior of monolithic castables: a) tension, b) compression, c) bending.

Fig. 3 .

 3 Fig. 3. Helpful test types to complete mechanical tests when studying monolithic castables: a) digital image correlation (DIC), b) in-situ high temperature environmental scanning microscopy (in-situ HT-ESEM), c) dilatometry.

  temperature behavior: effect of a 700°C firing.

  temperature behavior: effect of the testing temperature.

Fig. 4 :

 4 Fig. 4 : Mechanical behavior in tension of an andalusite based monolithic castable

Fig. 6 :

 6 Fig. 6: Strain fields determined by DIC at different steps of a tensile test (andalusite based material)

Fig. 7 :

 7 Fig.7: Evolutions of a microcrack during the first heating in an andalusite based monolithic castable (high temperature in-situ ESEM micrographs).

Fig. 8 :

 8 Macroscopic behavior of a geopolymer based FRRC after a 500°C firing and as a function of the testing temperature (comparison between the room temperature and the 500°C behaviors): a) in the case of a direct tensile test, b) in the case of four points bending tests.

Fig. 9 :

 9 Fig. 9: Evolutions of fiber/refractory concrete interfaces of a geopolymer based refractory concrete reinforced with short metallic fibers, after a 500°C firing.

Fig. 10 :

 10 Fig. 10: Room temperature strain fields determined by DIC on a geopolymer based refractory concrete reinforced with short metallic fibers. Illustration of stress redistribution phenomena: a) on a four points bending sample, b) on a technological sample.

Fig. 11 :

 11 Fig. 11: Four points bending behaviors at high temperature: a) Non-isothermal creep test results for a non-reinforced And-LCC castable and for Nextel 720 ® (And-N7) and Nextel 610 ® (And-N6) reinforced grades, under a 2MPa stress level. b) Strain velocity effect at 900°C both for a non-reinforced And-LCC grade and for a Nextel 720 ® reinforced grade.

Fig. 12 :

 12 Fig. 12: High temperature microstructural evolutions during the first heating of an andalusite based monolithic castable (high temperature in-situ ESEM micrographs).

Fig. 13 :

 13 Fig. 13: Effect of the firing time on the four points bending behavior at high temperature: a) Four points behavior of a non-reinforced And-LCC castable at 900°C, b) Illustration of crystallization phenomena that occur for long firing times (here 150h) at 900°C.

Fig. 14 :

 14 Fig. 14: Four points bending curves at various temperatures (700°C, 900°C, 1000°C and 1100°C): a) for a non-reinforced grade (And-LCC), b) for a Nextel 720 ® reinforced grade (And-N7).

Figure 15 - 1 Figure 15 - 2 Fig. 15 :

 15115215 Figure 15-1 Figure 15-2 Fig. 15: Macroscopic views of four points bending samples (a: lateral faces, b: zoom on lower faces, in tension) for a non-reinforced grade (And-LCC) and for a Nextel 720 ® reinforced grade (And-N7): -after 900°C tests at a crosshead speed of 0,02mm/min, -after non-isothermal creep tests up to 1200°C.

Fig. 16 :

 16 Fig. 16: Views of diffuse damage developed during a 900°C four points bending test (interrupted before the peak stress) in an andalusite based castable reinforced with 2 vol. % of Nextel720 ® fibers. White arrows indicate microcrack deviation phenomena, close to N7 fibers.

Fig. 17 :

 17 Fig. 17: Views of damage developed during a 1200°C four points bending test (interrupted at the peak stress) in an andalusite based castable reinforced with 2 vol. % of Nextel720 ® fibers. White arrows indicate microcrack deviation phenomena and cavitation limitation phenomena, close to N7 fibers.


T Kagnaya

C Boher

L Lambert

M Lazard

T Cutard

Wear mechanisms of WC-Co cutting tools from high-speed tribological tests

Keywords:

In machining processes, surface roughness and dimensional accuracy of machined parts depend on tool wear. Cemented carbide based cutting tools remain widely used in machining processes for their wear resistance. The aim of this paper is to study wear mechanisms of a WC-Co cutting tool grade under tribological conditions. Dry friction experiments are carried out on a high-speed pin-on-disc tribometer considering WC-Co pins against steel discs made of an AISI 1045 grade. Furthermore, two type-K thermocouples are embedded in WC-Co pins in order to estimate the contact temperature during sliding tests. A large sliding speed range is considered: from 60 up to 600 m/min. Results deal with wear rate versus tribological conditions, evolution of friction coefficient and temperature versus sliding speed and evolution of wear mechanisms. WC-Co tribological pins exhibit different wear mechanisms: abrasion, adhesion, transgranular WC micro-cracking and WC/WC debonding. Circulation of debris in the friction contact depends on sliding speed and on test duration. The evolution of surface temperature versus friction coefficient is studied too. Furthermore, a particular attention is paid on the relationship between the thermal energy in the pin and the mechanical energy in the contact. The thermal energy is calculated from the temperature values of K-thermocouples and the mechanical energy dissipated in the pin is calculated from friction coefficient evolutions. Relationships are established through modelling approaches.

Introduction

WC-Co cemented carbides are extensively used as cutting tools in machining processes for their excellent wear resistance. In machining process, cutting tools are generally subjected to severe mechanical and thermal conditions. Such conditions influence considerably cutting tool wear. The study of the cutting tool wear still remains a challenge in cutting process because it depends on the temperature level and on friction conditions at tool/chip and tool/workpiece interfaces. During machining process, conditions at these interfaces are not fully understood. Several studies [START_REF] Yang | Pin-on-disc wear testing of tungsten carbide with a new moving pin technique[END_REF][START_REF] Yang | Determination of the wear coefficient of tungsten carbide by a turning operation[END_REF][START_REF] Yang | Wear coefficient of tungsten carbide against hot-work tool steel disc with two different pin settings[END_REF][START_REF] Zemzemi | Analyse du frottement aux interfaces pièce-outil-copeau au cours de l'usinage d'un acier 42CrMo4[END_REF] were performed in laboratory tribological conditions (pin/disc friction conditions) to fully understand wear mechanisms at tool/chip and tool/workpiece interfaces. To study wear of machining tools, Yang [START_REF] Yang | Pin-on-disc wear testing of tungsten carbide with a new moving pin technique[END_REF][START_REF] Yang | Determination of the wear coefficient of tungsten carbide by a turning operation[END_REF] developed a moving pin technique for pin-on-disc wear testing using the whole disc surface area. Using three different disc steel grades, he showed that the wear coefficient obtained with the moving pin tribometer is more consistent than the one obtained by classic friction test (stationary pin test) for the same friction conditions. The highest wear rates obtained in the case of the moving pin are explained because the pin always slides on a new surface. Moreover, mechanical characteristics of the disc remain constant. In addition, the comparison of the wear coefficient between tribological pins and insert tips used in turning operation showed a difference of one order of magnitude. It is suggested that the high turning temperature at the tool-chip interface may have lowered the hardness of the work material during the turning operation to give the lower wear coefficient values. To take into account the temperature effect in tribological tests, Yang [START_REF] Yang | Wear coefficient of tungsten carbide against hot-work tool steel disc with two different pin settings[END_REF] investigated wear coefficient of tungsten carbide on a high temperature pin-on-disc tribometer. To simulate the turning operation, the vertical alignment of the pin is modified to achieve different contact surfaces.

Whatever test temperatures, the pin wear coefficient is always more important than insert tips. For high temperature friction tests, they observed that the most involved wear mechanism is adhesive wear without any deep grooves and with the presence of debris. At room temperature abrasive wear is predominant. This observation showed that the temperature plays an important role in wear mechanisms. Jianxin et al. [START_REF] Jianxin | Friction and wear behaviors of the PVD ZrN coated carbide in sliding wear tests and in machining processes[END_REF] studied friction and wear behavior of WC tools with a PVD ZrN coating in sliding wear (20-100 N and 0.2-1 m/s) and in machining processes. With the sliding wear tests they showed that the coated carbides offer the best abrasive resistance and the same conclusion is obtained for cutting operations.

Therefore, some authors have used tribological tests to measure wear coefficient and to understand wear mechanisms of carbide tools. Nevertheless, machining processes have some differences with tests of sliding wear (contact area, chip formation, chip plastic deformation, chip geometry, heat generated during cutting, temperature level, etc.). Indeed, pin-on-disc friction tests can be viewed as the continuous friction between flank-face and workpiece. This comparison is less convincing for the instantaneous friction between rake face and chip during cutting process. On the other hand, during cutting processes, the generated microchips can be embedded between surfaces in friction and can play a part as debris like in tribological tests. Moreover, the temperature of the interface is a parameter that affects the mechanical and thermophysical properties of cemented carbides [START_REF] Acchar | Strength degradation of a tungsten carbide-cobalt composite at elevated temperatures[END_REF][START_REF] Östberg | Effect of TaC on plastic deformation of WC-Co and Ti(C,N)-WC-Co[END_REF]. The aim of this paper is to study the relationship between the temperature generated by friction and wear of WC-Co tool grade. The tribological tests are performed with a pin-on-disc tribometer at a very high sliding speed (10 m/s). The study focuses on SEM analysis of worn surfaces of the pins, calculations of wear parameters and temperature measurements. An analysis of the temperature evolutions versus friction coefficient and wear mechanisms at different conditions are investigated.

Experiments and materials

High-speed tribometer

Friction tests are carried out with a very high-speed pin-on-disc tribometer in dry conditions. The disc is 170 mm in diameter and 10 mm in thickness. The disc has a continuous rotating movement and the rotating speed can vary from 30 to 3000 rpm. Linear sliding speed depends on radius of the wear track and can range from 0.05 to 30 m/s (Fig. 1a). The pin has a cylindrical shape (diameter 10 mm) with a truncated conic and with a flat circular surface of 2 mm in diameter (contact surface). Before any test, the pin and the disc are cleaned using dry air and the pin is cleaned in an ultrasonic tank using acetone and ethanol. The normal loading is carried out by dead weights. The tangential force (friction) is measured using a strain gauge located in the friction plane. Two type-K thermocouples are embedded in the pin to measure the temperature evolution during friction. The thermocouple holes are 0.6 mm in diameter and the thermocouple diameter is 0.5 mm. The first (TC1) and the second (TC2) thermocouples are embedded respectively at 1 mm and at 3 mm from the contact surface of the pin (Fig. 1b andc). The distance between the two thermocouples is always 2 mm. To be sure that the thermocouples are well embedded, a silver glue is used.

The temperatures (TC1 and TC2) of the pin and the friction forces are directly recorded during the test with a software developed on LABVIEW. The data sampling frequency for temperatures and for friction forces is 10 Hz.

Materials

The disc steel grade is AISI 1045. It is a medium carbon content steel with a perlito-ferritic microstructure. The cemented carbide pin is composed of tungsten carbide grains and of a cobalt binder (WC-6.5 wt.% Co grade G1 K20). The grain shape is polygonal and their largest size is in the range of 350 nm to 2 m (Fig. 2). The WC-Co hardness is 1550 HV 10 and the density is 14.90 g/cm 3 .

Set-up conditions

Several friction tests are carried out under various sliding speeds to reproduce wear mechanisms as observed at the surface of machining tools. In comparison of tribological machining wear mechanisms, criterion that are considered deal with similarities between microstructural evolutions of the contact surfaces. The frame of the tribometer is able to bear very high-speed tests but it is designed for normal loadings less than 100 N. For this reason, the used normal load is preferentially 60 N. For one sliding speed and different durations, tests are performed to measure the evolution of the pin wear. The test conditions are listed in Table 1.

From the experimental values (friction force, pin height loss and thermocouple temperatures), four physical quantities are determined to compare and explain our results. They deal with the pin wear volume, the pin wear rate, the mechanical cumulated energy dissipated by friction in the contact and the heat energy cumulated in the pin. The pin wear volume (WV) is calculated from the diameter deviation of the unworn (D 1 ) and worn (D 2 ) surface (Eq. ( 1)). The wear volume (WR) is obtained considering Eq. ( 2).

WV = 24 (D 2 -D 1 )(D 2 1 + D 2 2 + D 1 D 2 ) (1) WR = WV F N L (2)
where WV (mm 3 ) is the pin wear volume, WR (mm 3 /(N m)) the pin wear rate, D 1 and D 2 (mm) respectively the mean value of an initial diameter and of a final diameter of the pin contact surface, F N (N) is a normal load, and L (m) is the sliding distance.

The mechanical cumulated energy dissipated by friction in the contact (E MC ) (Eq. ( 3)) is directly calculated from the friction curve versus test duration (cumulative energy). The energy dissipated in the contact is the work of the friction force. For each time interval t, the dissipated energy E can be calculated with Eq. ( 3). Considering the mean value of the friction coefficient and assuming a constant sliding speed and normal load, Eq. ( 4) can be used:

E = F N V t 0 (t) dt = F N V ¯ i t (3)
E MC = E = n i=0 F N V ¯ i t (4)
where E MC (J) is a mechanical cumulated energy dissipated by friction in the contact, V (m/s) is linear sliding speed, t is increment of time (s), n is total number of data measures (t f / t), t f is friction test duration (s), F N is friction force, and ¯ i is the mean value of Coulomb friction coefficient during the time interval t.

The mechanical cumulated energy dissipated by friction only in the pin (E MP ) is a fraction of the mechanical cumulated energy dissipated by friction in the contact (E MC ) (Eq. ( 5)). It depends on the partition coefficient (ˇ) calculated with the Peclet number. The partition coefficient is expressed versus linear sliding speed (V) in Eq. ( 6) (Appendix A).

E MP = ˇEMC (5) ˇ= 1.66V -0.385 (6)
The heat energy cumulated in the pin (E TP ) measured from the deviation of the two temperatures (TC1 and TC2) is given in Eq. ( 7), considering a unidirectional thermal conduction with spherical coordinates (Appendix B).

E TP = S p h n i=0 T i t (7)
where E TP (J) is the heat energy cumulated in the pin (J), p (W/(m • C)) the thermal conductivity of the pin, S the considered surface (m 2 ), h (m) the distance between the two thermocouples, and

T(t) = TC1(t) -TC2(t) ( • C
) is the deviation between temperature of TC1 and TC2.

Analyses

Optical microscopy, and scanning electronic microscopy (SEM) and energy dispersive spectrometry (EDS) are used to analyze WC-Co pin surfaces and pin cross-sections. For the last ones, the surface cut is made in a direction parallel to the friction direction (AA ′ orientation) or perpendicular to the friction direction (BB ′ orientation). The topography of the wear surface of the pins is carried out with a confocal microscope.

Results and discussion

Evolution of the friction coefficient

For each test, a mean value of friction coefficient is calculated from the friction coefficient curve. For a normal load of 60 N, the mean friction coefficient versus sliding speed is plotted in Fig. 3. The friction coefficient decreases when the speed increases (0.63-0.4). Two phenomena can be considered to explain this linear decrease: (i) the debris circulation into the contact and (ii) the increase of the contact surface temperature. In the first case, debris circulation induces that more the sliding speed is high more the pin is worn. The WC-Co particles are smashed to powder in the contact and participate to the friction as a third-body. The pin wear will be discussed in the next subsection. Furthermore, the increase of the contact temperature leads to damage the mechanical properties of the two antagonists, particularly the disc ones.

The dependence between the evolution of the friction and the thermocouple temperature one is also very high (Fig. 4). If the temperature level is directly dependent on the mechanical energy dissipated in the contact (thus on the product between contact pressure and sliding speed), the temperature evolutions are resulting from the friction and consequently from the circulation of debris in the contact. In this approach, the influence of the pin height loss is not taken into account even if it could influence the measured temperatures. Considering that in our tests, the maximum pin height loss is equal to 0.13 mm after 1 h of test, and considering the quasisteady state of the temperature evolution, it could be assumed that the level of the temperature is slightly influenced by this parameter.

Evolution of the pin wear volume

For the different sliding speeds and test durations, the plot of the wear volume of WC-Co pins versus the mechanical energy into the contact (sliding distance × normal load) is considered as linear and verifies the Archard law (Fig. 5). The pin wear rate (WR) is equal to 2.44E-7 (mm 3 /(N m)).

The plot of the wear volume of WC-Co pins versus the mechanical energy dissipated in the pin (E MP ) is also considered as linear (Fig. 6). In this case, the curve slope is equal to 4.16E-6 (mm 3 /(N m)). Even if the partition coefficient (ˇ) is not a constant value, the same trend is obtained with the mechanical energy dissipated by friction only on the contact (E MP ).

For the pin, whatever the different sliding speeds for a normal loading of 60 N, the curve trends are linear. It means that the kinetic of pin debris generation is constant with a wear rate of 2.44E-7 (mm 3 /(N m)). If the kinetic is the same, the linear trends can also indicate that the mechanisms of wear loss can be considered as being the same. Several authors [START_REF] Huq | Expressing wear rate in sliding contacts based on dissipated energy[END_REF] already studied the linear dependence between dissipated energy and wear loss.

Pin wear mechanisms 3.3.1. Macroscopic scale considering friction and thermocouple temperature evolutions

The evolutions of the friction coefficient and of the thermocouple temperatures are given for a 600 m/min sliding speed, a normal load of 60 N and a test duration of 1 h (Figs. 7 and8).

These figures show that similarities exist between temperature and friction coefficient evolutions versus time. The friction curve presents several stages with more or less large fluctuations. The friction coefficient as well as the thermocouple temperatures reaches a steady state after a short-term period (200 s). Decrease of the friction coefficient during this stage (stage 1) is related to a mechanical accommodation between the antagonists and to a "polishing process" of the WC-Co pin surface. At high sliding speed, during the first moments of friction, shocks between asperities of the antagonists induce sparks and an increase of the mean temperature [START_REF] Gomes | The effect of sliding speed and temperature on the tribological behaviour of carbon-carbon composites[END_REF]. During tests, sparks were clearly observed at the pin incoming. For a speed of 600 m/min, the temperature of the first thermocouple TC1 is higher than 300 • C after 3 s of friction and reach 400 • C after 25 s (Fig. 8). The very fast increase of temperature during this stage results from the break of initial asperities at the first time of the friction between the two antagonists. After this stage, the friction coefficient is constant as well as the mean temperature (stage 2). Observing such a steady state means that the wear mechanism of the pin is constant during the test duration and that the boundary conditions in terms of heat convection exchanges around the pin are quite a constant.

Curves are generally composed of regular fluctuations with a frequency less than 1 Hz (Figs. 7 and8). They probably result from the tribometer frame vibrations and of large random fluctuations due to the debris evolution in the contact.

Large variations of the friction coefficient and of the temperature are noticed after stage 2 (other tests confirm these evolutions). Authors [START_REF] Paris | High-speed tribological behaviour of a carbon/silicon-carbide composite[END_REF] previously observed such a phenomenon. The very hard WC debris contributes as a third-body in the friction contact to form deep scratches parallel to the friction direction at the pin contact surface (Fig. 9). And circulation of hard particles into the contact could explain these large variations of the friction coeffi- cient evolution. It is assumed that circulating debris are essentially made of WC-Co because steel ones (iron debris) are preferentially transferred at the pin surface.

Microscopic scale considering SEM observations and EDS analyses

SEM observations of the pin deal with contact surface examinations and with cross-section ones. The wear modes observed are micro-plastic deformation and micro-cracks initiation of WC grains. Compared to the WC phase, the binder phase is very soft. Under high shear stress levels, the cobalt binder phase can flow at the microscopic scale and induces direct contacts between WC grains. As seen in Fig. 10, the polygonal grain shape induces stress concentration and micro-cracks initiation [START_REF] Pirso | Friction and wear behaviour of cemented carbides[END_REF][START_REF] Pirso | Friction and dry sliding wear behaviour of cermets[END_REF]. The weak toughness of WC (about 10 MPa √ m) could explain the fast propagation of the cracks through the grains. Plastic deformation of WC grains is also observed (Fig. 11). In addition, under huge local shear stresses, WC fragments are removed from the matrix and act as abrasive debris in the contact to contribute to the scratches formation. Grains fragmentation and grains debonding are observed in Fig. 12.

For 600 m/min, abrasive scratches and iron oxide transfer are observed at the pin contact surfaces (Fig. 13). The scratches are generated by the debris flow through the friction contact. Oxygen is always detected in the iron-rich transfer made of iron oxide. Thus, disc wear is involved in transfer debris at the pin surface. The transfer layer thickness is generally very low (2-3 m) (Fig. 14) even if in some cross-sections the thickness could reach 10 m. The transfer is essentially composed of iron oxide (Fig. 14). Below 600 m/min, no macroscopic transfer is observed at the pin contact surface. Furthermore, a lot of small iron oxide transfer areas are dispersed through the material (Fig. 15, points C and D), at the spaces between WC grains (grain boundaries, triple points, etc.).

Finally, the macroscopic layers generated by transfer mechanisms of the iron disc are essentially detected for test durations higher than 17 min. Therefore, it can be concluded that occurrence of transfer mechanisms take place at the same time that the large fluctuations and the low increase of values on the friction curves (after the stage 2) and on the thermocouple ones (Figs. 7 and8). Therefore, adhesion could be a wear mechanism that induces more energy dissipation than abrasion or groove formation.

Consideration of energies dissipated in pin wear

The wear mechanisms of pins are depending on the energy dissipated in the contact by friction (E MC ). The temperature increase in the pin is directly linked to the large part of mechanical energy dissipated in the contact. We make assumption that the whole dissipated energy in the pin (E MP ) is transformed in heat, leading to an increase of the thermocouple temperature. So, energy consumed for plastic deformation of the sub-contact layers and energy of the secondary processes (e.g. acoustic emission, oxidation, failure, etc.) are neglected.

In this subsection, we are studying both the level of the dissipated power and of the dissipated energy, leading to pin wear mechanisms. As previously written, two energy and power quantities can be calculated from measurements: the mechanical energy dissipated in the contact (E MC ) by measuring the friction coefficient (Eq. ( 4)) and the thermal energy dissipated in the pin (E TP ) (Eq. ( 7)) by measuring the temperature deviation between the two thermocouples TC1 and TC2. From E MC and using the partition coefficient, the mechanical energy dissipated in the pin (E MP ) is calculated (Eq. ( 5)). Powers corresponding to these energies are considered as being equal to the ratio "dissipated energy/total test duration" (e.g. P TP = Power and E TP = Energy). The power values are instanta-neous quantities and not cumulated ones. In power values, the test duration is not taken into account.

For very high sliding speeds (≥300 m/min), the thermal power dissipated (P TP ) in the pin and the thermocouple temperature (TC1 measured after 15 min of test duration) are slightly speed depending (Figs. 16 and17). For low sliding speeds (<300 m/min), the thermal power dissipated in the pin and the thermocouple temperatures are strongly speed depending. These two levels of thermal power dissipated in the pin must be related to the debris circulation in the tribo-contact. For P TP high levels, scratches and iron transfer at the pin contact surface are observed after a short time test (15 min). Temperatures are also very high. For P TP low levels, even after 2 h of test, no scratches and no transfer at the pin contact surface are observed and the temperature remains low. Only the contact surface of the pin is polished.

The thermal energy dissipated in the pin (E TP ) plotted versus the mechanical one (E MP ) has a linear dependence with a slope coefficient equal to 1.13 (Fig. 18). That means that 88.5% of E MP is

Conclusions

Friction and wear of a WC-Co cemented carbide are studied on a very high sliding-speed tribometer. Tests have been carried out with a constant normal load and various sliding speeds and test durations. Two thermocouples have been embedded in the pin to quantify the thermal cumulated energy dissipated in the pin. Wear parameters and temperature measurements have been performed to better understand the wear mechanisms of the WC-Co pins.

The friction coefficient decreases when the sliding speed and the pin temperature increase. The friction coefficient and the temperature curves always have the same evolution versus test duration. The wear volume of the pin is directly depending on the mechanical energy dissipated in the pin. At low sliding speeds, the wear mechanisms of the pin deal with plastic deformation and microcracking of WC grains, fragmentation and debonding of WC grains and polishing of the pin contact surface. At high sliding speeds, a supplementary wear mechanism is observed. It deals with transfer of iron oxide. When iron transfer appears, the friction coefficient and temperature curves versus test duration are very chaotic. The thermal power dissipated in the pin depends on the sliding speed. 88.5% of the mechanical energy dissipated in the pin are transformed into thermal energy. The last percents of this energy are lost in convection and wear phenomena. It seems that adhesion (iron transfer) plays an important part in this energy because fluctuations of friction coefficient and temperature, versus time, are very high when adhesion appears.

Appendix A

To determine the part of the dissipated energy in the pin, it is supposed that all the mechanical energy dissipated in the contact is converted into heat. The generated heat is distributed between the pin and the disc. The calculation of dissipated energy in the pin (E MP ) is based on the determination of a heat partition coefficient (ˇ) and total dissipated energy (E MC ). This coefficient is related to experimental conditions. For high-speed sliding tests, it depends on the dimensionless thermal number Pe (Peclet number) and can be determined following Eq. (A1).

Appendix B

The thermal energy dissipated in the pin (E TP ) is calculated using Eq. [START_REF] Östberg | Effect of TaC on plastic deformation of WC-Co and Ti(C,N)-WC-Co[END_REF]. It is based on some assumptions. Numerical simulation investigations performed to determine pin contact surface temperature [START_REF] Kagnaya | Friction at high sliding speed of WC/Co pin versus steel disc: estimation of the contact temperature[END_REF] showed that isothermal surfaces have a spherical form (due to the conical shape of the pin). We consider that for long test times, the temperature reaches a steady state. In this case, heat conduction is the principal heat transfer mode. Considering a hollow sphere, if a flux is applied on the interior surface of the sphere and if it is considered that, there is not loss of energy, the same quantity of flux is fully transmitted to the outside surface of the sphere (Fig. B1). On the one hand, we consider that the angle between points of thermocouples implantation is very small. In this way, it is supposed that points of measurement are located on a same radius. Based on these assumptions, the heat flux density can be calculated using Eq. (B1). On the other hand, it is considered that the heat flux density calculated using Eq. (B1) is applied on the surface (S) of the spherical cap passing by thermocouple TC1. So the total thermal energy dissipated in the pin (E TP ) during the test is calculated using the following equation (B2).

Fig. 1 .

 1 Fig. 1. (a) Global view of the high-speed tribometer; (b) tribological pin in contact with the disc; (c) pin design and thermocouples location.

Fig. 2 .

 2 Fig. 2. Microstructure of the WC-Co grade and WC grain size.

Fig. 3 .

 3 Fig. 3. Friction coefficient versus sliding speed (normal load: 60 N).

Fig. 4 .

 4 Fig. 4. Mean friction coefficient versus mean temperature of thermocouple TC1 (V = 60-100-300-400-600 m/min; FN = 60 N, test durations are given in Table1).

Fig. 5 .

 5 Fig. 5. Pin wear volume (WV) versus mechanical energy given in the contact (sliding distance × normal load) (normal load: 60 N).

Fig. 6 .

 6 Fig. 6. Pin wear volume (WV) versus mechanical energy dissipated by friction in the pin (EMP) (normal load: 60 N).

Fig. 7 .

 7 Fig. 7. Friction coefficient evolution versus time (V = 600 m/min; normal load: 60 N).

Fig. 8 .

 8 Fig. 8. Temperature evolution versus time (V = 600 m/min; normal load: 60 N).

Fig. 9 .

 9 Fig. 9. Pin surface topography (V = 600 m/min; P = 60 N; t = 60 min).

Fig. 10 .

 10 Fig. 10. SEM micrograph of the pin cross-section (AA ′ orientation) (600 m/min -60 N-1 h). Observation of WC micro-cracks due to shear stresses.

Fig. 11 .

 11 Fig. 11. SEM micrograph of the pin cross-section (600 m/min-60 N-2 h). Observation of WC plastic deformation due to shear stresses and to the polygonal shape of WC grains.

Fig. 12 .

 12 Fig. 12. SEM micrograph of the pin cross-section (600 m/min-60 N-17 min). Observation of WC debonding and of WC fragmentation.

Fig. 13 .

 13 Fig. 13. SEM micrograph of the pin contact surface (600 m/min-60 N-27 min). Observation of iron oxide transfer at the pin contact surface.

Fig. 14 .

 14 Fig. 14. SEM micrograph of the pin cross-section (600 m/min-60 N-1 h). Observation of the thickness of the iron oxide transfer (EDS analysis: iron oxide).

Fig. 15 .

 15 Fig. 15. SEM micrographs of the pin contact surface (600 m/min-60 N-1 h): (a) global view, (b) point A and (c) point B.

Fig. 16 .

 16 Fig. 16. Thermal power dissipated in the pin versus sliding speed (FN = 60 N).

Fig. 17 .

 17 Fig. 17. Mean temperature of the TC1 thermocouple versus sliding speed (FN = 60 N; time = 15 min).

Fig. 18 .

 18 Fig. 18. Mechanical energy dissipated in the pin versus thermal energy dissipated in the pin.

  where V (m/s) is the sliding speed, ˛p (m 2 /s) the thermal diffusivity of the pin (˛p = p / p c p ), p (W/(m • C)) the thermal conductivity of the pin and d (W/(m • C)) the thermal conductivity of the disc.

S = 2 r 2 ( 1 -

 21 cos Â) (B3)where  is the half angle and r (m) is the radius of the sphere passing by thermocouple TC1.

Fig. B1 .

 B1 Fig. B1. Conical pin configuration and spherical cap.

Table 1

 1 Test conditions.

		Load, 60 N
	Sliding speed (m/min)	60	100	300	400	600
	Test durations (min)	120	15	15	5, 15, 60	17, 27, 60, 120

Table