Joris Kadok

Nicolas Bost

Antoine Coulon

Mohamed-Ramzi Ramzi Ammar

Séverine Brassamin

Cécile Genevois

Auriane Etienne

Aurélien Canizares

Jacques Poirier
email: jacques.poirier@univ-orleans.fr

Etienne Auriane

Aurélien Canizarès

Inhibiting the sp2 carbon deposition by adjunction of sulphurous species in refractory ceramics subjected to CO and H2 reducing atmosphere

niveau recherche, publiés ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

2

The formation of carbon from carbon monoxide decomposition, known as the Boudouard reaction 2CO CO2 + C solid has been studied extensively [START_REF] Boudouard | Les phénomènes de combustion dans les foyers industriels[END_REF][START_REF] Turkdogan | Catalytic effect of iron on decomposition of carbon monoxide: I. Carbon deposition in H2-CO mixtures[END_REF][START_REF] Walker | Carbon Formation from Carbon Monoxide-Hydrogen Mixtures over Iron Catalysts. I. Properties of Carbon Formed., II. Rates of carbon formation[END_REF][START_REF] Baird | Carbon formation on iron and nickel foils by hydrocarbon pyrolysis-reactions at 700°C[END_REF][5][6][START_REF] Irving | Interaction of evaporated carbon with heated metal substrates[END_REF][START_REF] Robertson | Carbon formation from methane pyrolysis over some transition metal Surfaces -I. Nature and properties of the carbons formed[END_REF][START_REF] Banerjee | Pyrolytic Carbon Formation from Carbon Suboxide[END_REF][START_REF] Presland | Growth of single-crystal graphite by pyrolysis of acetylene over 30 metals[END_REF][START_REF] Ruston | The solid reaction products of the catalytic decomposition of carbon monoxide on iron at 550°C[END_REF][START_REF] Nolan | Catalytic disproportionation of CO in the absence of hydrogen: encapsulating shell carbon formation[END_REF][START_REF] Dai | Single-wall nanotubes produced by metal-catalyzed disproportionation of carbon monoxide[END_REF][START_REF] Krause | The predictability of CO resistance of refractory materials by state-of-the-art test methods-a technical and scientific approach[END_REF]. The reaction, preferentially catalysed by cementite formed from metallic and oxide particles [START_REF] Turkdogan | Catalytic effect of iron on decomposition of carbon monoxide: I. Carbon deposition in H2-CO mixtures[END_REF], allows a maximum yield of the sp 2 carbon formation occurring between 500 and 600° [START_REF] Walker | Carbon Formation from Carbon Monoxide-Hydrogen Mixtures over Iron Catalysts. I. Properties of Carbon Formed., II. Rates of carbon formation[END_REF]. P.L. Walker et al. have studied the effect of temperature and inlet carbon monoxide-hydrogen composition on the rate of carbon formation. These authors have shown that: "as the hydrogen content of the gas mixture is increased, the temperature at which the maximum rate of carbon deposition occurs also increases, this temperature ranging from 528 °C for a 99.2% CO-0.8% H2 mixture to 630 °C for a 80.1% CO-19.9% H2 mixture" [START_REF] Walker | Carbon Formation from Carbon Monoxide-Hydrogen Mixtures over Iron Catalysts. I. Properties of Carbon Formed., II. Rates of carbon formation[END_REF].

The FexOy particles are transformed in the presence of CO into Fe3C catalytic cementite by a carburation reaction, and the formation of carbon originates from a dissolution-precipitation mechanism. The carbon deposition can be explained by the formation of liquid iron/ Fe3C droplets of nanometric size which dissolve carbon until the saturation limit. Carbon is segregated on the liquid surface droplet [START_REF] Krause | The predictability of CO resistance of refractory materials by state-of-the-art test methods-a technical and scientific approach[END_REF]. Among the catalysts, iron and iron oxides but also cobalt and nickel are particularly active between 350 °C to 750 °C [START_REF] Baird | Carbon formation on iron and nickel foils by hydrocarbon pyrolysis-reactions at 700°C[END_REF].

Moreover, the addition of a tiny fraction of reducing gas H2 is known to drastically increases the rate and the amount of carbon produced by the Boudouard reaction [START_REF] Turkdogan | Catalytic effect of iron on decomposition of carbon monoxide: I. Carbon deposition in H2-CO mixtures[END_REF][START_REF] Walker | Carbon Formation from Carbon Monoxide-Hydrogen Mixtures over Iron Catalysts. I. Properties of Carbon Formed., II. Rates of carbon formation[END_REF].

For example, Walker et al. [START_REF] Walker | Carbon Formation from Carbon Monoxide-Hydrogen Mixtures over Iron Catalysts. I. Properties of Carbon Formed., II. Rates of carbon formation[END_REF] reported that 0.1 g of carbonyl iron catalyst in a 9 % H2-91 % CO gas mixture at 600°C allows the formation of 10 g of solid carbon after 350 minutes [START_REF] Walker | Carbon Formation from Carbon Monoxide-Hydrogen Mixtures over Iron Catalysts. I. Properties of Carbon Formed., II. Rates of carbon formation[END_REF].

Depending on the CO/H2 ratio, this reaction can produce not only an amorphous or flocculent carbon but also more structurally organised carbon [5][6][START_REF] Irving | Interaction of evaporated carbon with heated metal substrates[END_REF][START_REF] Robertson | Carbon formation from methane pyrolysis over some transition metal Surfaces -I. Nature and properties of the carbons formed[END_REF][START_REF] Banerjee | Pyrolytic Carbon Formation from Carbon Suboxide[END_REF][START_REF] Presland | Growth of single-crystal graphite by pyrolysis of acetylene over 30 metals[END_REF][START_REF] Ruston | The solid reaction products of the catalytic decomposition of carbon monoxide on iron at 550°C[END_REF][START_REF] Nolan | Catalytic disproportionation of CO in the absence of hydrogen: encapsulating shell carbon formation[END_REF][START_REF] Dai | Single-wall nanotubes produced by metal-catalyzed disproportionation of carbon monoxide[END_REF], such as nanotubes (CNT for Carbon Nano Tubes), nanofibres (CNF for Carbon Nano Fibres), platelets or shells.

Recently, Bost et al. have initiated fundamental studies [START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF][START_REF] Bost | Probing the structural organisation of sp2 carbons obtained by the Boudouard reaction using in situ Raman scattering in reducing conditions[END_REF] to understand the effect of various experimental parameters such as valence and particle size of iron oxide catalysts on the amount of carbon deposition in CO/H2 gas mixtures and to characterise the structural organisation of sp 2 carbon. The influence of the degree of oxidation (Fe, FeO, Fe3O4 and Fe2O3) and the particle size (nanometre to millimetre size) of the FexOy catalyst particles on the carbon deposition rate was investigated using thermogravimetric analysis and Raman spectroscopy with different CO/H2 gas ratios.

The results show that the valence of the iron is not the key point that governs the amount of carbon deposited, but the carbon formation depends on the specific surface area of the catalytic particles. When the gas is composed of 100 % CO, a well-organised sp 2 carbon in the shape of encapsulating polyaromatic shells is formed around large catalytic Fe3C particles.

Under these conditions, the smaller the initial particle size is, the higher the carbon formation rate is due to a higher specific area which induces a higher catalytic activity of Fe3C [START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF].

Hematite (ρ = 5.25 g/cm 3 ) is firstly reduced to maghemite (ρ= 4.86 g/cm 3 ) then to magnetite (ρ = 5,15 g/cm3) and finally to cementite (ρ = 7,4 g/cm3). Krause and Pötschke [START_REF] Krause | The predictability of CO resistance of refractory materials by state-of-the-art test methods-a technical and scientific approach[END_REF] suggest the primary recrystallization of hematite to magnetite causes an enlargement of the reactive surface. The reactions of hematite reduction which occur with volume changes are possible under CO at temperatures of about 350°C.

Under CO-H2 atmosphere, these phase transformations result in a dramatic increase in the specific surface areas of catalytic particles. The carbon growth rate is hence faster and favours the nucleation of polyaromatic carbon nanofibres on the resulting Fe3C particles [START_REF] Bost | Probing the structural organisation of sp2 carbons obtained by the Boudouard reaction using in situ Raman scattering in reducing conditions[END_REF].

Consequently, the surface of the iron oxides plays a major role in the kinetics of the carbon formation reaction.

These different types of carbon produced by the Boudouard reaction inspire a great interest because they exhibit multiple interesting properties, such as remarkably high Young's modulus and tensile strength for CNTs and CNFs [START_REF] Ruoff | Mechanical and thermal properties of carbon nanotubes[END_REF][START_REF] Yazdanbakhsh | Challenges and Benefits of Utilizing Carbon Nanofilaments in Cementitious Materials[END_REF]. Graphene nanoplatelets are known to reinforce the mechanical and electrical properties of composite materials [START_REF] Drzal | Conference Proceedings of High Performance Fillers[END_REF].

However, the formation of large quantities of sp 2 carbon can be undesirable in some specific situations. Refractory materials, containing iron oxides as impurities of raw materials and subjected to a reducing gas, are strongly degraded due to destructive carbon deposits present in the open porosity [START_REF] Krause | The predictability of CO resistance of refractory materials by state-of-the-art test methods-a technical and scientific approach[END_REF]. The CO resistance of refractories is usually improved by the choice of raw materials with a low content of iron oxide impurities and by the increase of sintering temperature. However, these solutions are not effective enough for new applications such as low-CO blast furnace process in which refractories are subjected to CO and H2 gas.

The literature on the mechanism of metal dusting corrosion [START_REF] Grabke | Thermodynamics, mechanisms and kinetics of metal dusting[END_REF][START_REF] Zeng | Relationship between the Growth of Carbon Nanofilaments and Metal Dusting Corrosion[END_REF][START_REF] Chu | Mechanisms of Metal Dusting[END_REF], even though it is indirectly concerned with this issue, points to possible avenues of research to prevent the formation of carbon. In a CO-H2 atmosphere at approximately 600°C, the metal dusting phenomenon causes large deposits of fine graphitized carbon particles [START_REF] Grabke | On the mechanism of catastrophic carburization: metal dusting[END_REF] along with the catastrophic corrosion of Fe-, Co-and Ni-based metals and alloys. The metal dusting is explained by the growth of a cementite layer at the surface of the metal by carbon transfer from CO(g) followed by graphite formation on the cementite. The decomposition of the cementite leads to the release of fine metal particles that will act as catalysts for further carbon formation [START_REF] Schneider | Effect of H2S on 31 formation and decomposition of Fe3C and Fe5C2 under metal dusting conditions[END_REF]. Hydrogen increases the carbon formation rate [START_REF] Ermakova | Decomposition of Methane over Iron Catalysts at the Range of Moderate Temperatures: The Influence of Structure of the Catalytic Systems and the Reaction Conditions on the Yield of Carbon and Morphology of Carbon Filaments[END_REF]. The addition of a sulphurous gas under metal dusting conditions is well known to prevent the usual deleterious effect caused by this phenomenon on metals [START_REF] Schneider | Effect of H2S on 31 formation and decomposition of Fe3C and Fe5C2 under metal dusting conditions[END_REF], [START_REF] Berry | Influence of impurities and role of iron carbides in deposition of carbon from CO[END_REF][START_REF] Schneider | Computer Coupling of Phase Diagrams and[END_REF][START_REF] Schneider | Effect of H2S on metal dusting of iron[END_REF].

Schneider et al., for instance, carefully studied the influence of H2S(g) on iron under similar conditions, i.e., between 500 °C and 700 °C in a CO-H2-H2O-H2S atmosphere [START_REF] Schneider | Effect of H2S on 31 formation and decomposition of Fe3C and Fe5C2 under metal dusting conditions[END_REF], [START_REF] Schneider | Computer Coupling of Phase Diagrams and[END_REF], [START_REF] Schneider | Effect of H2S on metal dusting of iron[END_REF].

These authors reveal that under such conditions, the adsorption of elemental sulphur occurs via the reaction: H2S = H2 + S adsorbed Karcher et al. also showed that SO2 gas inhibits the carbon deposition on an iron and steel surface [START_REF] Karcher | Inhibition of carbon deposition on iron and steel surfaces[END_REF]. Troilit or pyrrhotite crystallize at the surface of cementite. A layer of FeS is formed at the catalyst surface of the metal and acts as a H2-CO barrier which prevents the nucleation and the growth of carbon.

The aim of the present paper is to investigate the inhibition of the sp 2 carbon formation in refractory ceramics from CO decomposition catalysed by iron oxides under the most favourable conditions of the reaction, i.e., in the presence of H2 at 600 °C, and to describe in detail the mechanisms of carbon inhibition.

Only few research studies focus on the inhibition of the sp 2 carbon deposition in refractory ceramics submitted to CO and H2 reducing atmosphere [START_REF] Walker | Carbon Formation from Carbon Monoxide-Hydrogen Mixtures over Iron Catalysts. I. Properties of Carbon Formed., II. Rates of carbon formation[END_REF], [START_REF] Berry | Influence of impurities and role of iron carbides in deposition of carbon from CO[END_REF], [START_REF] Karcher | Inhibition of carbon deposition on iron and steel surfaces[END_REF]. The effect of sulphur on the carbon deposition produced by the Boudouard reaction, catalysed by cementite in presence of H2 gas, was investigated under laboratory conditions. First, the materials and the experimental means are presented and the Boudouard reaction with or without S inhibitors is described in detail. The morphological and structural evolution of iron carbide and carbon nanoparticles is studied. The effect of the addition of different concentrations of solid sulphur S(s) and sulphur dioxide SO2 (g) that inhibits the carbon formation is quantified and a mechanism of carbon inhibition by the sulphur is proposed.

Finally, this study paves the way to an effective solution to inhibit the sp 2 carbon formation by the Boudouard reaction in the refractories by poisoning the Fe3C catalyst with sulphur.

Materials and methods

Materials

Iron present in refractories have two origins. The natural iron impurities of the raw materials are hematite, especially if the raw materials have been heat treated. In addition, the impurities from crushing and granulation may contain mainly iron and a smaller amount of Ni, Co and W. With regard to the damage of the refractories, hematite plays an essential role. So, Fe2O3 iron oxide particles were used to carry out laboratory trials. Fe2O3 powder with a particle size of 2.5 µm was purchased from the Sigma Aldrich Company (Table 1). The specific surface of Fe2O3 particles was determined for information from the N2 adsorption isotherms at 77 K at a relative pressure (P/P0) range of 10 -6 to 1 atm. (Autosorb-1MP, Quantachrome) by applying the BET equation. It should be noted that the initial geometry of Fe2O3 particles is not an essential experimental value since this geometry changes drastically when the reactive atmosphere (CO, H2) is present. Solid sulphur and sulphur dioxide are used as inhibitors of the carbon deposition via the Boudouard reaction, as detailed in Table 1. The reducing gas mixture (71 % CO, 3 % CO2, 11 % H2, 15 % N2), hereinafter referred to as CO/H2 gas, was similar to the CO-H2 recycled gas mixture in future low-CO blast furnace industrial processes, in the context of Ultra Low CO2 Steelmaking (ULCOS) initiative. This composition of gas was previously used in the research work published in the papers [START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF][START_REF] Bost | Probing the structural organisation of sp2 carbons obtained by the Boudouard reaction using in situ Raman scattering in reducing conditions[END_REF] A low cement castable (Refracast LCC-80R/TS -Refratechnic) has been used to validate inhibition of carbon deposition by gaseous sulphur.

Methods

To study the effect of sulphur on the formation of carbon produced by the Boudouard reaction, Fe2O3 powders were mixed with S. The boiling point of sulphur being 444.6 °C, under the test conditions (600°C) the sulphur is gaseous.

Laboratory tests were also carried out with SO2 (g) to understand the effect of gaseous sulphur on the formation of carbon produced by the Boudouard reaction. SO2 (g) was mixed continuously with the reducing CO/H2 gas. The SO2 (g) quantities were controlled using valves and gas flow meters of great accuracy. H2S (g) was not used for two reasons: firstly, to have only one hydrogen source (H2) for a better control, and second, because SO2 (g) is a less toxic gas than H2S (g).

Preliminary trials were performed in a horizontal tubular furnace, composed of an alumina tube (Pyrox®/Eurotherm 902P). In order to quantify the carbon growth rate, these macroscopic measurements of carbon formation were completed by thermogravimetric analysis measurements. Thermogravimetric analyses were carried out on Fe2O3 samples exposed to a CO/H2 gas mixture at 600 °C. Measurements of Fe2O3 mass variation were performed using a thermobalance from NETZSCH STA 409C/CD coupled with a QMS 403/5

Skimmer. The reproducibility of the TGA instrument, given by NETZSCH, is ± 0.1 %. Fe2O3

samples of approximately 30 mg were placed in an alumina crucible and were heated at 600 °C with a rate of 10 °C•min -1 under 100 ml•min -1 of an Ar inert gas prior to being exposed to a reducing CO/H2 gas (100 ml•min -1 ).

The structural organisation of samples and the kinetics of carbon formation were determined by X-ray diffraction and Raman spectroscopy.

The crystalline phases were analysed with powder X-ray diffraction (PXRD) using a Bruker D8 A25 diffractometer (Bragg-Brentano geometry θ-θ) equipped with a LynxEye XE detector and using the Cu Kα radiation (λ = 1.5418 Å). Each scan was recorded in a 2θ range from 20 to 60° with a step size of 0.016 at 10 s per step.

Raman spectroscopy was performed using a T64000 Jobin-Yvon multichannel spectrometer equipped with an Ar-Kr laser source, aBX41 Olympus microscope and a liquid nitrogencooled CCD detector. The spectra were collected in a backscattering geometry under a microscope (×20 long working distance objective) using the 514.5 nm wavelength (2.41 eV; 3 mW of laser power), a 600 groves•nm -1 grating giving a spectral resolution of 3 cm -1 in the 1000-2000 cm -1 wavenumber range. Heating the sample was achieved using a TS1500

Linkam device (Tadworth, U.K.) [START_REF] Ruoff | Mechanical and thermal properties of carbon nanotubes[END_REF]. The TS1500 heating stage was connected to a pressurized gas cylinder using inox tubes. The same temperature program was used for Raman spectroscopy and for TGA (10°C•min -1 until 600°C). The carbon deposit was analysed, in situ, at high temperature for 240 minutes, with a Raman acquisition time of one minute.

After the experiments, the samples were also characterised by microscopic observations.

Scanning electron microscopy (SEM) was carried out using a Hitachi S4500 FEG on polished sections. High resolution transmission electron microscopy (HRTEM) was also used to characterise the nanostructure of the different samples. HRTEM images were collected using a Philips CM20 microscope. Elemental compositions and STEM images were carried out on a JEOL ARM200F (JEOL Ltd.) cold FEG operating at 120 kV and equipped with a TEM/STEM double spherical aberration (Cs) correctors. Scanning transmission electron microscopy -high angle annular dark field (STEM-HAADF) images were acquired in the angular range of 50-180 mrad with an 8 cm camera length and a 0.1 nm probe size. Elemental composition line scans were performed by STEM-EDS using a JEOL EDS (Energy Dispersive Spectroscopy) system and a 0.13 nm probe size. Samples were prepared by dispersing a small amount of the material in ethanol using an ultrasound bath. A drop of the solution was deposited onto a holey carbon-coated copper grid.

Results

3.1.Effect of sulphur on the amount of carbon produced by the Boudouard reaction

This section focusses on preventing the formation of carbon from CO decomposition catalysed by iron oxides, by adding sulphur, in a solid or gaseous form. Laboratory experiments were carried out to understand the relevant parameters taking place in the carbon formation and to examine the effect of the S addition on the C deposition by the Boudouard reaction. The experiments were carried out under optimal conditions for the formation of carbon, i.e., with Fe2O3 at 600 °C under a gas mixture (71 % CO, 3 % CO2, 11 % H2, 15 % N2).

Preliminary experiments

In order to ascertain the beneficial effect of sulphur addition on the inhibition of carbon deposition, preliminary macroscopic tests were carried out on a larger scale of observation.

Fe2O3 powders were mixed with pure sulphur. The samples (1 g) were placed in an alumina basket and then heated to 600°C under air at a rate of 10°C•min -1 prior to being exposed to the CO/H2 gas with a 28 L•h -1 flow rate, for one hour. EDS analysis reveals also the presence of FeS with a concentration gradient of S about 4 mm thick (Figure d,e). These preliminary tests show that sulphur and sulphur dioxide inhibit carbon deposition. In contact with SO2 (g), FeS is formed on the surface of the Fe2O3 particles and the carbon deposit is not formed. SO2 gas seems to have a delayed effect on the carbon formation.

Sulphur obviously poisons cementite, which no longer acts as a catalyst. As a result, CO is more stable than C + CO2, as it should be under the given ambient conditions.

The next step is to quantify the effect of inhibiting sulphur on the carbon deposition produced by the Boudouard reaction.

Thermogravimetry analysis

TGA experiments were performed to quantify the effect of inhibiting sulphur on carbon growth rate as a function of the S amount in the various solid or gaseous mixtures.

TGA experiments with solid sulphur

Fe2O3 powder was mixed with sulphur in different amounts: 1, 5, 10 and 30 wt.% S. Figure 4 shows the mass variation of samples in time (full lines) obtained from thermogravimetric measurements as well as the variation in the temperature (dashed lines) as a function of time.

The TGA results of mixtures (Fe2O3+S) are compared to an experiment without S addition. Several chemical reactions with gain and loss of mass occur simultaneously. During heating at a rate of 10°C.min -1 , under an inert atmosphere (100 ml•min -1 Ar), Fe2O3 + S mixtures have a mass loss associated with sulphur vaporisation at 440 °C, proportional to the amount of the introduced S (Figure 4). The mass balance of the TGA experiments performed with S + Fe2O3 samples at the temperature of sulphur vaporisation (440°C) and prior to the injection of CO/H2 gas (600°C) is presented in the table 2. Almost all the sulphur volatilises prior the injection of the CO/H2 gas. The amount of sulphur adsorbed on the surface of the Fe2O3 particles is extremely small. The FeS layer is formed at the volatilization temperature of sulphur, prior the CO / H2 injection. Higher concentrations of solid precursor sulphur are not beneficial for the formation of the FeS layer. The stability over time of FeS will be discussed in the paragraph: "Retarding effect of sulphur on the carbon deposition".

After 60 minutes, the injection of Ar gas is stopped and CO/H2 reducing gas is introduced continuously (100 ml•min -1 ). A weight loss approximately of 25 wt.% is observed due to the chemical reduction of the iron oxides by hydrogen and carbon monoxide (the theoretical weight loss is 30.06 %), as previously explained in literature [START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF]. The Fe2O3 reference sample shows a very important carbon deposit. The samples containing sulphur have a significantly lower weight gain, similar for all the samples (approximately 12.5 wt.%) that stabilizes after 400 minutes.

TGA experiments with gaseous sulphur compounds

A continuous introduction of gaseous sulphur dioxide SO2 in CO / H2 mixture has many advantages compared to a solid sulphur addition. Indeed, the solid sulphur volatilizes at 440 °C and the amount of sulphur adsorbed on the surface of the particles is extremely low. The SO2 (g) leads to a more sustainable amount of sulphur and thus more efficient inhibition of carbon formation in time.

The experiments were carried out with Fe2O3 at 600 °C under a mixture of CO/H2 reducing gas containing SO2 (g). The gas flow rate was 100 ml•min -1 .

A mass balance in the following experimental conditions (180 min at 600 °C under a mixture of CO / H2 gas containing 1000 ppm of SO2 (g) with a gas flow rate of 100 ml.min -1 ) show that ratio S injected / Fe2O3 is equal to 3.4 wt. %.

Figure 5 shows the TGA results (mass variation as a function of time) for SO2 (g) content equal to 1, 10, 50, 100 and 1000 ppm, and the results are compared to an experiment without SO2 (g) addition. Once the gas mixture was injected, all the samples underwent a similar mass loss for approximately 20 minutes, due to the Fe2O3 chemical reduction into pure iron prior to its carburation in iron carbide [START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF]. For the samples exposed to a gas mixture containing 1 to 10 ppm of SO2 (g), a dramatic mass gain is observed due to carbon formation, roughly similar to what could be found without SO2 (g). Table 3 indicates the linear carbon growth rate in mg•h -1 for different experimental conditions. SO2 content (1 and 10 ppm) is obviously not sufficient to lower significantly the carbon formation. When 50 ppm of SO2 was injected, a lower mass gain was observed, and the carbon growth rate was decreased by a factor of 5 in comparison with the experiment without SO2. For 100 ppm and 1000 ppm of SO2, the carbon growth rate values were decreased by a factor of 70. However, a slightly higher mass gain is observed for the sample with 100 ppm of SO2 in comparison with 1000 ppm.

Retarding effect of sulphur on the carbon deposition

To determine if the initial introduction of SO2 gas has a retarding effect on the carbon formation produced by the Boudouard reaction, a TGA experiment was carried out in two steps:

1. The Fe2O3 sample was exposed to the N2 gas (100 ml•min -1 ) containing 1000 ppm of SO2 (g), over 3 hours at 600 °C. In that case, the sulphur is adsorbed on the surface of the iron oxide sample.

2. Then, the composition of the gas is modified. The initial gas, N2 containing 1000 ppm of SO2 (g), is replaced by the CO/H2 gas (100 ml•min -1 ) without SO2(g) additions for 3 hours in addition at 600°C.

The resulting curve is shown in Figure 5 (bold dashed blue curve, superimposed on the other curves, on the same graph). A retarding effect of sulphur is observed. The carbon growth rate for this experiment is 16.4 mg•h -1 , which is significantly lower than for the experiment without SO2 (21.6 mg•h -1 ) but higher than for the experiment under CO/H2 reducing gas containing 1000 ppm SO2 (0.3 mg•h -1 ), which shows a significant inhibition but not a total prevention of the carbon formation. Moreover, the temporary exposition to 1000 ppm of SO2 had a lesser effect (a few hours) compared to a continuous exposition of SO2 (g) with a lower content (100 ppm). The chemisorption of SO2(g), reduced to S on the surface of Fe2O3

particles is probably at the origin of the prevention of the carbon formation. Otherwise, the elemental physisorbed sulphur would evaporate immediately because its boiling temperature (444 °C) is below the temperature of experiment. The sulphur species, which are retained in the deeper sections of the Fe2O3 catalyst, diffuse slowly up to the surface could thus explain the observed effect. At the microscopic scale, this mechanism is complex. It will be described in detail in a subsequent publication. The affinity of sulphur to iron being higher than the affinity of oxygen to sulphur, a sulphur rich layer formed at the surface of the initial Fe2O3 particles. The chemical reaction Fe + S FeS is an exothermic reaction: ΔH°FeS = -102 KJ/mole. FeS limits the nucleation and growth of the sp 2 carbon nanotubes (nucleation which requires the formation of liquid / iron / iron carbide droplets). Sulphur effectively delays carbon deposition but does not suppress it. FeS has a limited effect in time (a few hours) and is therefore not stable. The reason for this instability is an important point that is discussed in paragraph 4.

So, the mechanism of sulphur chemisorption does not appear to be sustainable and hence requires a continuous SO2 (g) exposition.

From TGA experiments, it is shown that 1% Ssolid or 100 ppm SO2 (g) are sufficient to inhibit carbon formation. The structural and morphological evolution of the samples is presented in the following paragraph.

3.2.Structural evolution of samples and kinetics of carbon formation

The evolution of crystalline phases of each sample, after the TGA experiments, was characterised by powder X-ray diffraction. PXRD patterns of the samples exposed to 0, 10, 50, 100 and 1000 ppm of SO2 under TGA are presented in Figure 6.a. 4) Å (Pearson symbol mS28) [START_REF] Leineweber | Crystal structure determination of Hägg carbide, x-Fe5C2 by first-principles calculations and Rietveld refinement[END_REF]. The crystal structure of these two compounds is similar in the sense that they can be described by a stacking of identical Fe and Fe-C atomic planes. However, the two compounds differ in the stacking order and orientations of those planes. The presence of these two phases in the PXRD patterns is in accordance with the literature. Schneider et al. reported that the increasing contents of H2S for pure iron samples under metal dusting conditions (CO + H2 at 500 °C) favours the formation of this Fe5C2 compound [START_REF] Schneider | Effect of H2S on 31 formation and decomposition of Fe3C and Fe5C2 under metal dusting conditions[END_REF]. On every PXRD pattern, the (002) carbon reflexion is broad shaped, indicating an imperfectly crystallized carbon. Its intensity tends to decrease with the increase of the SO2 content. These results are consistent with TGA experiments where the carbon formation was decreasing with the increasing SO2 (g) content. Iron sulphide phases are not detected on PXRD diffractograms.

In addition to X-ray diffraction, Raman spectroscopy was performed to study the structural evolution of carbon deposits, for iron oxide samples exposed to 0, 10, 50, 100 and 1000 ppm of SO2 (Figure 6.b)

Raman spectroscopy is a very sensitive characterisation method that reveals the presence of carbon as well as its structural state from perfectly crystalline to amorphous [START_REF] Ammar | Characterizing various types of defects in nuclear graphite using Raman scattering: Heat treatment, ion irradiation and polishing[END_REF][START_REF] Ammar | How to obtain a reliable structural characterization of polished graphitized carbons by Raman micro spectroscopy[END_REF].

The Raman spectra of the sp 2 carbon contain three main bands that range from 1000 to 2000 cm -1 :

-G band, located at approximately 1580 cm -1 when recorded at room temperature, that corresponds to the in-plane bond stretching of carbon atoms (E2g symmetry)

-Two defect bands (D: ≃ 1250-1400 cm -1 and D': ≃ 1600-1630 cm -1 ) originating from a peculiar interaction between electrons and phonons due to defects in the crystal structure [START_REF] Maslova | Determination of crystallite size in polished graphitized carbon by Raman spectroscopy[END_REF].

The ID/IG intensity ratio of the D and G bands and the G band full width at half maximum (FWHM) are commonly used to characterise the disorder of carbon-based materials [START_REF] Ammar | Characterizing various types of defects in nuclear graphite using Raman scattering: Heat treatment, ion irradiation and polishing[END_REF], [START_REF] Pimenta | Studying disorder in graphite-based systems by Raman spectroscopy[END_REF][START_REF] Nakamizo | Raman spectra of ground natural graphite[END_REF][START_REF] Negri | Resonance Raman contribution to the D band of carbon materials: modeling defects with quantum chemistry[END_REF][START_REF] Castiglioni | Raman spectroscopy of polyconjugated molecules and materials: confinement effect in one and two dimensions[END_REF][START_REF] Donazzi | Annular reactor testing and Raman surface characterization in the CPO of methane and propyleneAppl[END_REF].

In a previous publication [START_REF] Bost | Probing the structural organisation of sp2 carbons obtained by the Boudouard reaction using in situ Raman scattering in reducing conditions[END_REF], Bost et al. have shown that the Raman spectrum of sp 2 carbon produced by the Boudouard reaction catalysed by iron oxide, in a pure CO atmosphere without hydrogen, exhibits only a sharp G band. The absence of the defect-induced D band suggests a high degree of crystallinity with important coherent domain diameters of the resulting carbon, similar to what could be found for graphite. In contrast, the sp 2 carbon produced in a CO/H2 gas mixture turns out to be in the form of nanofibres, and the corresponding Raman spectrum exhibits a pronounced D band activated by the boundaries that constitute these nanofibres.

When 0, 10 and 50 ppm of SO2 is used, Raman spectra (Figure 6.b), characterised by a pronounced D band, display similar structural organisation of the nanofibres produced.

However, the Raman spectra of the samples exposed to 100 and particularly 1000 ppm of SO2 exhibit some differences. The FWHM of the G band is relatively wider for these samples, particularly the sample exposed to 1000 ppm of SO2. This is a signature of the production of more disordered carbons or nanofibres with lower coherent domain diameters. Basically, ex situ Raman spectroscopy cannot be used here to provide quantitative information about the amount of the sp 2 carbon produced owing to the absorbent character of this material to the visible light and therefore the penetration depth of the excitation laser used. In fact, the band located at approximately 1310 cm -1 and appears as a shoulder of the D band is attributed to magnon scattering mode and its presence is an indirect evidence of the small amount of the produced carbon, especially for the samples exposed to 1000 ppm of SO2. PXRD is completely silent about this where no (002) carbon reflection is observed, which shows the high sensitivity of Raman spectroscopy in probing sp 2 carbon materials and the complementarity of the characterisation techniques used.

In situ Raman spectroscopy, at 600°C, has also been carried out to determine the kinetics of the carbon formation (Figure 7). Approximately 1 mg of Fe2O3 was first heated to 600 °C under N2 prior to being exposed to the reducing gas mixture with or without the SO2 containing gas (≃ 24 L•h -1 ). Raman spectra were acquired every minute during the experiments.

The intensity of the G band (in arbitrary units) versus time is plotted for the in situ Raman experiments without SO2 and with 1000 ppm of SO2, as shown in Figure 7.

Considering the fact that the G band is proportional to the amount of the carbon exposed to the excitation laser, two pieces of information can then be extracted here.

1. The introduction of 1000 ppm SO2 in the gas mixture clearly has a delaying effect in the formation of carbon since this latter starts to be probed after more than 1 hour (Figure 7.b), while the detection of carbon was immediate in the absence of SO2 (Figure 7-a).

2. The formation of carbon without SO2 is so important that the intensity IG starts to decrease after 90 minutes (Figure 7.a) due to the laser defocusing. In Figure 7.b, the intensity IG remains relatively constant without decreasing and is also evidence of a tiny amount of carbon without defocusing the excitation laser. Considering all the experimental results (TGA, XRD and Raman spectroscopy), we can conclude that with the addition of increasing SO2 quantities in the gas mixture (> 100 ppm), the carbon formation is significantly lowered and delayed, according to a mechanism that will be proposed in the discussion section. The effect of the Fe3C catalyst is inhibited. However, the carbon formation is not completely prevented.

3.3.Morphological and elemental characterization of samples

To better understand the effect of sulphur on the inhibition of the carbon produced by the Boudouard reaction catalysed by Fe2O3, the samples were observed after TGA experiments by scanning electron microscopy (SEM) and high-resolution transmission electron microscopy (HRTEM).

Figure 8 shows typical SEM micrographs of the samples obtained in back scattering electron mode, before (a) and after (b and c) the thermobalance experiment using CO/H2 gas without SO2(g) addition. Before the experiment (Fig. 8-a), the Fe2O3 particles are approximately 5 μm in diameter with a very angular aspect. After experiment (Fig. 8-b andc), the sample exhibits a dense population of particles, appearing brighter and measuring less than 0.5 µm in diameter. Considering the initial size of the Fe2O3 particles (2 -5 µm), the particles were decreased in size during the experiments by the phase transformations of hematite (chemical reduction of Fe2O3 into Fe3C with volumes changes and mineral recrystallizations), as previously explained [START_REF] Krause | The predictability of CO resistance of refractory materials by state-of-the-art test methods-a technical and scientific approach[END_REF][START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF].

EDS analysis systematically reveals the large presence of iron and carbon elements. The resulting material turns out to be a mixture of sp 2 carbon and iron carbide particles, as revealed by PXRD. This mixture contains a large quantity of sp 2 carbon, according to TGA and Raman spectroscopy results. The sp 2 carbon is in the form of nanofibres observed as dark phase in SEM images. Figure 9 shows typical SEM micrographs of the samples after the thermobalance experiment using CO/H2 gas with SO2(g) addition.

The sample exposed to 1000 ppm of SO2 (g) exhibits a quite different morphology. The particles sizes do not seem to have been reduced or have been reduced to a lesser extent. EDS analysis revealed very weak traces of sulphur, but its presence is not sufficient to precisely determine its composition and its location. As already asserted with Raman spectroscopy, a pure carbon phase, which should be present in this sample, is not observed on the micrographs because of its very small quantity. PXRD patterns revealed the presence of iron carbide particles such as Fe3C and/or Fe5C2 particles (Fig. 6)

Figure 9-Scanning electron micrographs of the samples after TGA experiments using CO/H2 gas with 1000 ppm SO2 (g). Images were taken at 15 kV in back-scattered electron mode (BSE)

TEM micrographs of the samples not exposed to SO2 show the presence of iron carbide (cementite) particles with a reduced size compared to the initial particle size (Figure 10.a).

The rounded shape of these particles may suggest a molten state, as previously explained by Krause and Pötschke [START_REF] Krause | The predictability of CO resistance of refractory materials by state-of-the-art test methods-a technical and scientific approach[END_REF]: "a prerequisite of carbon growth is the formation of liquid iron/carbide raw droplets which absorb carbon until the saturation limit; otherwise carbon deposit will not take place". A part of the carbon forms nanofibres growing from Fe3C For STEM-HAADF images, the intensity is proportional to the density of the sample and the atomic number of the element (proportional to Z 1.7 ). Consequently, and contrary to the previous TEM images, the iron carbide particles appear with a brighter contrast than the carbon particles. The inner layer is very thin (0.5 -1 nm) and seems to be composed of lighter elements compared to the outer layer (3 -5 nm), which appears brighter.

Atomic resolution STEM-HAADF images show that the particle and the outer layer are crystallized but without an apparent orientation relationship (see Figure 11.d). Such observation is, however, not as clear for the inner layer where it is not possible to distinguish the atoms. Occasionally, the inner layer does not thoroughly surround the iron carbide. In such a case, the outer layer is absent (see Figure 11.e). To determine the evolution of the chemical composition, a 20-points EDS line scan has been performed with a 0.13 nm probe size through the three different domains: iron carbide particle, inner layer and outer layer (see Figure 12).

However, due to the spherical shape of the iron carbide particle, both layers, inner and outer, present all around this particle contribute slightly to modifying the composition of the iron carbide particle measured. Likewise, the composition found for the inner layer is largely altered by the presence of the outer layer due to its subnanometre-size. The compositions given at the boundaries between the two layers should be taken with the greatest care.

According to the result of the analysis, iron is present in the three domains. Iron is more concentrated in the particle, associated with carbon. The ratio (% at) of approximately 3:1 between these two elements indicates that the particle is composed of iron carbides, Fe3C

and/or Fe5C2, as expected. The outer layer composition is not biased by the other domains and seems to contain the three elements: iron, oxygen and carbon, which would suggest the presence of an iron oxide along with an iron carbide and/or free carbon. The composition of the inner layer is more delicate to interpret due to its very thinness (0.5 -1 nm) and the large contribution of the outer layer co-probing. Nevertheless, this inner layer is the domain richest in sulphur with a rather high content of iron, which suggests the formation of an iron sulphide such as FeS, as asserted by the preliminary experiments (see paragraph 3.1.1) and a previous study [START_REF] Schneider | Effect of H2S on 31 formation and decomposition of Fe3C and Fe5C2 under metal dusting conditions[END_REF].

Discussion

From the experimental results and the literature, a mechanism that governs the inhibition of carbon deposition by the Boudouard reaction catalysed by cementite is proposed.

Under CO/H2 gas without sulphur, the iron oxide particles are chemically reduced by hydrogen and CO. The iron oxide particles are converted to Fe3C iron carbide. Liquid iron/ carbides droplets are formed which dissolve carbon until the limit of carbon saturation and favour the nucleation and the growth of nanocarbon fibres [START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF]. Figure 13.a summarizes the mechanism of carbon formation by the Boudouard reaction, catalysed by cementite.

With SO2(g) additions, under CO/H2 gas, the carbon deposition by the Boudouard reaction is strongly retarded. The effect of sulphur could be explained by a multistep mechanism.

-Reduction of SO2 to H2S by CO/H2 gas For high H2/SO2 ratios, the two species H2 and SO2 are known to react together to form H2S.

Tai Feng et al. have studied the reduction of SO2(g) to elemental sulphur with H2 and mixed H2/CO gas [START_REF] Feng | Reduction of SO2 to elemental sulfur with H2 and mixed H2/CO gas in an activated carbon bed[END_REF]. Thermodynamic equilibrium calculations using the equilibrium module of the thermodynamic software Factsage 6.4 and experiments were carried out on a H2-CO-SO2 system and the temperature effects and H2/SO2 ratio were evaluated.

When the H2/SO2 ratio was greater than three, which is the case in our experiments, the thermodynamic calculations and the laboratory experiments showed that SO2 is completely reduced at 600°C to H2S. The elemental sulphur is completely absent from the reduction products. The reaction is described as follows:

3 H2 + SO2 = 2H2O + H2S (1)
-Chemisorption of sulphur Sulphur chemisorption and carbide formation occur simultaneously. Sulphur has a highly active surface. It tends to be strongly absorbed or segregated at the surfaces of catalytic iron and Fe3C particles. The reaction is described as:

H2S = H2 + Sads (2)
The reaction 2 leads to considerable coverage with adsorbed sulphur at low sulphur activities as [START_REF] Grabke | Role of sulphur in carburization, carbide formation and metal dusting of iron[END_REF].

As = Ks . pH2S/pH2 , with Ks: equilibrium constant

The micrographs (Fig. 11 and Fig. 12) show a very thin FeS layer. Laboratory experiments indicate that this FeS layer is formed immediately when the gas SO2 is injected.

-

Formation of Fe3C and Fe5C2 Hägg iron carbides

As shown in Figures 6 and10, for the samples exposed to a low SO2 content (0-10-50 ppm), PXRD patterns and TEM images revealed the large presence of Fe3C but also the formation of sp2 carbon. In those cases, Fe5C2 iron carbides are not formed.

For an exposition at higher contents (50 -100 -1000 ppm), the carbon formation strongly decreases and the presence of Fe5C2 is observed.

The formation of the Hägg carbide (71.4 at. % Fe) would be explained by the partial depletion of iron in the cementite phase (75 at. % Fe) according to the reaction:

2 Fe3C + Sadsorbed Fe5C2 + FeS
In a paper about the influence of H2S on metal dusting [START_REF] Schneider | Influence of H2S on metal dusting[END_REF], Schneider et al. also reported the same experimental results. These authors observed the formation of cementite Fe3C at low contents of H2S and the formation of the Hägg carbide Fe5C2 at higher contents of H2S.

-Inhibition of the carbon deposition catalysed by Fe3C (from the reduction of Fe2O3) in presence of H2.

The literature about metal dusting corrosion indicates that the addition of sulphur compounds (H2S) in the atmosphere inhibits metal dusting and prevents the cementite decomposition [START_REF] Grabke | On the mechanism of catastrophic carburization: metal dusting[END_REF]. The formation of the FeS layer limits the decomposition of the metastable cementite: TEM-EDS analysis revealed the formation of an iron sulphide inner layer (0.5-1 nm), surrounding the iron carbide nanoparticles (Fig. 11 and Fig. 12). The iron provided to form the inner FeS layer is coming from the cementite (Fe3C).

The protective layer of FeS has another beneficial effect. As shown previously [START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF], the nanometre size and the specific surface of catalytic particles promote the growth of carbon fibres.

of the iron carbide specific area, as shown in Figure 9.

In principle, sp 2 carbon nucleation and growth should be prevented by this monolayer FeS on the catalytic surface. But the FeS layer is not stable, S adsorption is a dynamic phenomenon and there are always some defects and vacancies in the FeS monolayer which will lead to a slow continued carbon formation. Thus, sulphur strongly retards but not fully suppress carbon formation. This explains a linear carbon growth rate of 0.3 mg.h -1 for 1000 ppm SO2 (see Fig. 5 and table 3) and the outer layer surrounded the iron carbide particles (Figures 11 and12). The STEM-EDS analysis revealed no sulphur but a significant amount of carbon, iron and oxygen in the outer layer. The mechanism explaining the formation of this layer remains an open question that it would be worth to study in detail for a publication in a subsequent paper.

Practical application to CO/H2 resistant refractories

The retarding effect of sulphur on the carbon deposition produced by the Boudouard reaction by gaseous sulphur has been validated by an industrial refractory material (low cement castable: Refracast LCC-80R/TS -Refratechnic). Its chemical composition is given in Table 4. This castable, doped with 5 wt.% of iron oxide particles to enhance the carbon deposit, was dried and pre-fired in air at 900 °C, for 8 hours. Two experiments were carried out:

-Experiment 1. The samples were heated to 600°C under Ar, at a rate of 10°C min -1 prior to being exposed to the CO/H2 gas with a 28 L.h-1 flow rate, for 3 hours.

-Experiment 2. The samples were heated to 600°C under air prior to being exposed to the CO/H2 gas containing 100 ppm of SO2 (g).

The results are presented in figure 14. The samples exposed to CO / H2 gas are much damaged and an important carbon deposit is observed. The samples exposed to CO / H2 containing 100 ppm of SO2 (g) are not damaged.

There is no significant carbon deposit.

This work reveals that the addition of sulphurous gas (100ppm) in CO / H2 gas in an effective solution and an efficient inhibitor to limit the carbon formation and the damage of refractories. For a specific industrial application, it will be essential to ensure that -The processed materials are not contaminated with sulphur -Sulphur release is compatible with environmental problems -The temperatures of the refractory lining do not allow the condensation of sulphuric acid in colder parts. This is the case with the future low CO-blast furnaces.

The addition of SO2 into CO-H2 gas is very weak (100 ppm)

The processed materials (elaboration of pig iron) allow low contamination with sulphur.

The blast furnace gas is purified or its reuse (valorization as energetic gas) is possible.

It's obvious as well, that the addition of a sulphur carrier to the refractory raw material mix is not a practical solution, because the FeS-layer is not stable, as described in this paper.

Conclusions

The inhibition of the carbon formed by the Boudouard reaction and catalysed by Fe3C has been investigated under the most favourable conditions of the reaction, i.e., in the presence of H2 at 600 °C. This work reveals that sulphur, in the solid and gaseous states, is an efficient inhibitor to limit the carbon formation. From 100 ppm of SO2(g), the carbon growth rate decreases by a factor of 70 compared to an experiment without inhibitor, but the reaction is not completely prevented. In a CO/H2 mixture at 600 °C, iron oxide particles are chemically reduced to Fe3C and their sizes decrease. With the addition of sulphur such as 100 to 1000 ppm of SO2(g), the size of iron oxide particles (Fe2O3) decreases to a lesser extent but are still reduced to cementite (Fe3C) and to Hägg iron carbide (Fe5C2). The sulphur intake has contributed to the formation of a very thin FeS layer (0.5 -1 nm) around the iron carbide particles, preventing the catalytic effect of the Boudouard reaction.

Sulphur has a retarding effect. However, a temporary exposition of SO2 (1000 ppm) has a lesser effect compared to a continuous exposition of SO2 (100 ppm).

This study associated with the previous work [START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF], [START_REF] Bost | Probing the structural organisation of sp2 carbons obtained by the Boudouard reaction using in situ Raman scattering in reducing conditions[END_REF], shows that is possible to control the structure, size and organisation of the sp 2 carbons produced by the Boudouard reaction, catalysed by iron oxides.

-In the presence of pure CO gas, a carbon with a high structural organisation and an important coherent domain diameter is produced in the form of shells around the catalyst particles; -In the presence of a CO and H2 mixture, a carbon with a nanofibre morphology is favoured;

-In the presence of CO and H2 mixture with an addition of S, the catalytic effect of the iron oxide particles is prevented.

This study paves the way to an effective solution to limit the sp 2 carbon deposition by the Boudouard reaction catalysed by iron oxides in the presence of hydrogen by poisoning the Fe3C catalyst with sulphur.

Fig. 1

 1 Fig. 1 shows the degradation of refractory castable samples after a CO-test (determination of resistance to carbon monoxide), according to the European standard EN ISO 12676. Carbon deposits on iron oxide spots and thereby disintegrates the refractory material.

Figure 1 .

 1 Figure 1. Deposits of carbon and degradation of refractory castable samples after laboratory CO-test (by courtesy of ECREF).

Figure 2

 2 Figure2shows a photograph of the samples. The 100 wt.% Fe2O3 sample shows a significant carbon deposition with a morphology of "popcorn" in accordance with the literature[START_REF] Bost | The catalytic effect of iron oxides on the formation of Nano-carbon by the Boudouard reaction in Refractories[END_REF].The samples containing sulphur have only a change in colour (Figure2 c-d). The black colour is caused by the reduction of hematite to cementite. It can be shown that carbon formation is inhibited if solid sulphur is part of the starting composition.

Figure 2 -

 2 Figure 2-Photography of samples before and after a treatment under a gas flux of 28 L•h -1 of a CO/H2 reducing gas, at 600°C, for one hour (Pyrox®/Eurotherm 902P furnace). a, b -100 % Fe2O3; c, d-70 % Fe2O3 + 30 % S

Figure 3 -

 3 Figure 3-Scanning electron micrograph of Fe2O3 bulk sample after experiment a) and b) Overview and detail of the sample c),d),e),f) EDS analysis of S

Figure 4 -

 4 Figure 4-TGA curves of the samples exposed to the CO/H2 reducing gas mixture at 600 °C -Mixture of Fe2O3 powders and pure S (1, 5, 10 and 30 wt.%) The grey area shows the heating stage to 600°C under an inert gas (Ar)

Figure 5 -

 5 Figure5-TGA curves of the Fe2O3 samples exposed to the reducing gas mixture containing small fractions of SO2 (g) at 600 °C (continuous curves). The results from the retarding effect of SO2 (g) experiment (see paragraph Retarding effect of sulphur on the carbon deposition) are also given for comparison (bold dashed curve). The grey area shows the heating stage to 600°C under an inert gas(Ar)

Figure 6 -

 6 Figure 6-a) Powder X-ray diffraction patterns of the samples obtained after TGA experiments under 0, 10, 50, 100 and 1000 ppm of SO2. Fe3C and sp 2 carbon are found in each pattern whereas the Fe5C2 carbide is found only under 100 and 1000 ppm of SO2. The 0, 10 and 50 ppm patterns are normalized with the highest Bragg reflection of the Fe3C phase (031) and the 100 and 1000 ppm patterns with the highest Bragg reflection of the Fe5C2 phase (510). b) Raman spectra of the carbon obtained after TGA experiments under 0, 10, 50, 100 and 1000 ppm of SO2. These spectra contain three main bands: G band and the defect-induced D and D' bands * Fe2O3 magnon peak From 0 to 50 ppm of SO2, PXRD diffractograms exhibit reflections that can be all indexed with sp 2 carbon (C -JCDPS 01-089-8487) and cementite (Fe3C -JCDPS file 01-079-4902). For 100 and 1000 ppm of SO2, the Hägg carbide (Fe5C2-JCDPS 01-076-3842) is also observed. Compared to the cementite which crystallizes in an orthorhombic lattice with parameters a = 5.0807(2) Å, b = 6.753(1) Å and c = 5.5152(3) Å (Pearson symbol oP16) [30], the Hägg carbide crystallizes in a monoclinic lattice with parameters a = 11.56953(7) Å, b = 5.45252(3) Å, c = 5.06043(3) Å and β = 97.7308(4) Å (Pearson symbol mS28) [31]. The

Figure 7 -

 7 Figure 7-Intensity of the G band versus time for the in situ Raman experiments (at a temperature of 600°C, under CO/H2 gas) a-without SO2, b-with 1000 ppm of SO2.

Figure 8 -

 8 Figure 8-Scanning electron micrographs of the samples. The powder was deposited on a substrate. a) Initial Fe2O3 sample prior the thermobalance experiment, at 20 kV. b) and c) Sample after the thermobalance experiment using CO/H2 gas without SO2(g) addition Images were taken at 15 kV in back-scattered electron mode (BSE)

  particles. These carbon fibres are composed of polyaromatic layers surrounding an amorphous core (see Figures 10.b and 10.c). Other Fe3C particles without filamentous carbon are enclosed by a graphitized carbon shell (not shown on this micrograph).

Figure 10 -

 10 Figure 10-Transmission electron micrographs of the sample after TGA experiments without SO2. a), b), c) Overview and details of the sample d) HRTEM micrograph showing details of the nanofibre. e) Fast Fourier Transform (FFT) from the graphitized carbon located on the external layer (left panel) and from the amorphous core of the nanofibre (right panel).

Figure 11 -

 11 Figure 11-Micrographs of the sample after TGA experiments exposed to 1000 ppm of SO2 (g). a), c) TEM image representative of the sample. b), d), e) high-resolution HAADF-STEM images of a particle surrounded by a two-layer shell.

Figure 12 -

 12 Figure 12-a) STEM-HAADF micrograph and b) composition profile (% at) through the iron carbide particle, the inner layer and the outer layer.

Fe3C 3 Fe

 3 + C and the carbon nucleation. The adsorbed big sulphur S 2-ions block the reaction sites for the carburization reactions and the carbon transfer can occur only on sulfur free areas (1θs), where θS is the fractional occupancy of the S adsorption sites.

Figure 13 -

 13 b summarizes the mechanism of carbon inhibition produced the Boudouard reaction by adjunction of sulphurous species.

Figure 13 -

 13 Figure 13-The formation of carbon from the Boudouard reaction, under CO/H2 gas, catalysed by iron oxide particles a) without sulphur addition and b) with sulphur addition

Figure 14 .

 14 Figure 14. Photography of refractory samples before and after exposure to CO/H2 and CO/H2 containing 100 ppm of SO2(g)

Table 1 :

 1 Origin and purity of the materials used

		Materials	Origin	Purity (wt%)	Particle size (µm)	Specific Surface (m 2 /g)
	Catalysts	Fe2O3	Sigma Aldrich	>99	2-5 a	8.0
	Reducing gas	CO/H2	Air Liquide	>99.9
	Sulphur dioxide Solid sulphur	SO2 (g) S	Air Liquide Alfa Aesar	>99.9 >99,5	≤ 10 b
	a Data as provided by the Sigma-Aldrich Company,

b

Prepared by grinding in an agate mortar.

Table 2 -

 2 Mass balance of the TGA experiments

	Initial S	Mass loss: Δm (%)	Mass loss Δm (%)
	(wt%)	at 440°C -44 min	at 600°C -60 min
		at S vaporisation temperature	at injection of CO/H2 gas
	1	0,92	1.03
	5	4.55	5.36
	10	9.25	10.11
	30	27.97	29.99

Table 3 -

 3 Linear carbon growth rate for TGA experiments. The initial sample masses were approximately 30 mg of Fe2O3.

	SO2 Content	Carbon Growth Rate
	(ppm)	(mg•h -1 )
	0	21.6
	1	21.0
	10	19.5
	1000	0.3

*The carbon growth rate was not determined for 50 and 100 ppm SO2 content, as the curves were not linear.

Table 4 -

 4 Chemical composition of the low cement castable a

	Oxides	wt. %
	Al2O3	81
	SiO2	16
	CaO	2.5
	Fe2O3	0.5

a Data as provided by Refratechnic company

Acknowledgements

This research was supported by the French National Agency: ADEME (TGRBF project).

The authors acknowledge the ICMN laboratory (Orléans, France) for the CM20 TEM access and the GENESIS platform for the TEM JEOL ARM 200F access. GENESIS is supported by the Région Haute-Normandie, the Métropole Rouen Normandie, the CNRS via LABEX EMC and the French National Research Agency as a part of the programme "Investissements d'avenir" with the reference ANR-11-EQPX-0020.

Annie Richard and Dr Ahmed Bachar are also acknowledged for the SEM/EDS analysis and for preliminary tests.


Gabriel Kerbart

Christelle Harnois

Christelle Bilot

Sylvain Marinel

Pressure-Assisted Microwave Sintering: A Rapid Process to Sinter Submicron sized Grained MgAl2O4 Transparent Ceramics

Keywords: Microwave, Uniaxial pressure, Sintering, Transparent Ceramics, Spinel

This work focuses on the development of an original process based on a 2.45 GHz singlemode microwave cavity equipped with a uniaxial press, to sinter transparent spinel MgAl2O4 ceramic in air. The samples were conventionally pre-sintered to a density of 90% TD before microwave sintering to the final stage of densification. The influence of thermomechanical cycle on the material properties was investigated. Transmittance, grain size distribution, hardness and fracture toughness of the samples were measured and correlated to the microstructure. This new sintering process has allowed obtaining transparent samples with sub micrometric grain size and high mechanical properties, with relatively short times and low temperature. These first results can be compared to some obtained by SPS or HIP. The technical input of this method is that all the process is here conducted in air atmosphere.

Introduction

Transparent polycrystalline ceramics have attracted much interest from the scientific community but also from industry due to their numerous potential applications in laser host, jewelry, spacecraft windows or IR windows and dome for military applications [START_REF] Wang | Transparent Ceramics: Processing, Materials and Applications[END_REF]. These ceramics are characterized by their attractive thermo-mechanical properties from room to high temperature (>1000°C), their intrinsic transparency in the visible-IR range, and the low cost of the raw materials [START_REF] Petit | Drying step optimization to obtain large-size transparent magnesium-aluminate spinel samples[END_REF]. However, the transparency can only be achieved for highly pure and fully dense materials. In the case of the magnesium aluminate spinel (MgAl2O4), hereafter termed spinel, this high density can only be achieved with high sintering temperature coupled with mechanical pressure. Techniques such as High Isostatic Pressure (HIP) and Spark Plasma Sintering (SPS) are the most effective in that way.

In the literature, the best results obtained by HIP require the use of a pre-treatment such as conventional sintering or hot forging [START_REF] Krell | Advanced Spinel and Sub-Um Al2O3 for Transparent Amor Applications[END_REF][START_REF] Sutorik | The Production of Transparent MgAl2O4 Ceramic Using Calcined Powder Mixtures of Mg(OH)2 and c-Al2O3 or AlOOH[END_REF]. For instance, with a conventional pre-sintering step followed by a HIP sintering at 1400°C, Krell et al. [START_REF] Krell | Advanced Spinel and Sub-Um Al2O3 for Transparent Amor Applications[END_REF] obtained transparent spinel, exhibiting 81-85% RIT at 640 nm, with a grain size of 400-600 nm, hardness of 15 GPa and fracture toughness of 2 MPa.m 1/2 . A similar grain size has been obtained by Goldstein et al. [START_REF] Goldstein | Transparent Polycrystalline MgAl2O4 Spinel with Submicron Grains, by Low Temperature Sintering[END_REF] on spinel sintered at 1320°C -170MPa. They measured a transmittance value of 75% RIT at 550 nm with similar mechanical properties.

Transparent spinel has been obtained using SPS at 1300°C-80 MPa by Morita et al. [START_REF] Morita | Spark-Plasma-Sintering Condition Optimization for Producing Transparent MgAl2O4 Spinel Polycrystal[END_REF]. They reported a grain size of 450 nm, a RIT value of 47% at 550 nm, and a hardness of 15GPa. Bonnefont et al. [START_REF] Bonnefont | Fine-Grained Transparent MgAl2O4 Spinel Obtained by Spark Plasma Sintering of Commercially Available Nanopowders[END_REF] communicated the production of a transparent ceramic with a mean grain size of 275 nm, exhibiting 74% RIT at 550 nm and 15 GPa hardness using SPS technique 1300 °C-72 MPa. Sokol et al. [START_REF] Sokol | Nano-Structured MgAl 2 O 4 Spinel Consolidated by High Pressure Spark Plasma Sintering (HPSPS)[END_REF] obtained a nanosized grains transparent spinel, with a mean grain size of 50 nm, by HPSPS at 1000 MPa and 1000°C, presenting >80% RIT at 550 nm, 18 GPa hardness, and 2.66 MPa.m 1/2 fracture toughness.

However, those methods involve specific atmosphere conditions during sintering (argon pressure or vacuum in contact with graphite, for respectively HIP and SPS). Otherwise, SPS samples may be contaminated by carbon, which is not favorable to obtain high transparency [START_REF] Morita | Fabrication of Transparent MgAl2O4 Spinel Polycrystal by Spark Plasma Sintering Processing[END_REF][START_REF] Zhu | Fabrication of Transparent MgAl 2 O 4 from Commercial Nanopowders by Hot-Pressing without Sintering Additive[END_REF]. Therefore, both methods may lead to transparent ceramics with a slight coloration, due to partial reduction [START_REF] Krell | Transparent Compact Ceramics: Inherent Physical Issues[END_REF][START_REF] Bernard-Granger | Influence of Graphite Contamination on the Optical Properties of Transparent Spinel Obtained by Spark Plasma Sintering[END_REF].

Over the years, microwave energy has emerged as an efficient technology for the synthesis and sintering of oxide ceramics, such as spinel [START_REF] Macaigne | Microwave Sintering of Pure and TiO2 Doped MgAl2O4 Ceramic Using Calibrated, Contactless in-Situ Dilatometry[END_REF], BaTiO3 [START_REF]Latest Global Developments in Microwave Materials Processing[END_REF], CaCu3Ti4O12 [START_REF] Riquet | Direct and Hybrid Microwave Solid State Synthesis of CaCu3Ti4O12 Ceramic: Microstructures and Dielectric Properties[END_REF] etc. This process offers a greener alternative to conventional heating, with a faster heating rate and less energy consumption than infrared radiation heating, while working under air. Furthermore, since the penetration depth of the microwave electrical field is a few centimeters in a lot of dielectric materials, it allows a volumetric heating, resulting in an efficient heating technology with a homogenous temperature distribution within the sample [START_REF] Oghbaei | Microwave versus Conventional Sintering: A Review of Fundamentals, Advantages and Applications[END_REF]. However, as previously evoked, removing all the porosity not only requires the diffusion activation at high temperature (natural sintering) but also high pressure. This is the reason why an original dielectric microwave heating process equipped with a uniaxial press has been developed. This paper aims at presenting this process and the results obtained for the sintering of spinel under uniaxial pressure. A 2.45 GHz single-mode cavity was used and specifically designed to be able to apply a uniaxial pressure during sintering. The overall system will be described and the microstructure, the mechanical and optical properties of the final sintered products will be correlated to the processing conditions.

Experimental procedure 2.1. Hot press microwave furnace

Microwaves are produced by a 2.45 GHz microwave generator (Sairem GMP 20KSM) and are transmitted to the resonant cavity through a normalized TE10 rectangular waveguide, ended by a TE10m cavity (m=3 or 4). The cavity is tuned to TE103 with the coupling iris, so that the heating is provided by the interactions of the material with the E-field (E-field mode). TE103 resonance conditions are achieved by using a coupling iris and a movable short circuit piston, while the impedance matching is tuned by a waveguide hybrid junction with two short circuit pistons (magic T). The microwave generator can deliver a maximum power of 2 kW. The overall system is shown on figure 1. The temperature is measured by an infrared pyrometer (OPTRIS CTM-3SF75H3-C3) pointing on the center of the cavity from its backside, where the sample will be located. This heating system is coupled with a uniaxial press positioned above the cavity. The force is applied by the upper steel punch onto the sample placed between two alumina cylinders to avoid direct contact between the hot sintering parts and the water-cooled punches. This system can deliver a maximum load of 1.6 tons. A minimum value of 8 MPa is necessary at the beginning of the cycle to assure punch and samples contact.

The displacement of the punch, the applied pressure and the internal measured temperature are recorded as a function of time, allowing drawing in-situ dilatometric curves. The system can be controlled either manually (the operator chooses the delivered power and applied pressure), or in an automatic mode. In this case the thermo-mechanical cycle follows a pre-defined program of temperature-delivered power and applied pressure. In each case, tuning of resonance, by shortcircuit punch position, and adaptation of the impedance are done manually.

To limit thermal losses, an aluminosilicate insulation box surrounds the sample. Two SiC plates (>99.9 % purity, from Umicore, Bagnolet, France), acting as susceptors, are disposed on each side of the sample and perpendicular to the microwaves propagation direction to initiate the heating. It is important to mention that the sintering is carried out in static air. Furthermore no mold was used in these experiments for two reasons (1) it is quite difficult to find a suitable material for molds, able to work in air, at high temperature and pressure in a microwave chamber, and presenting no (or slight) reactivity with the material. (2) It has been shown that spinel exhibits a pseudo-plastic behavior at high temperature [START_REF] Beclini | MICROSTRUCTURAL SUPERPLASTIC DEFORMATION IN MgO • A1203 SPINEL[END_REF] from which we can take advantage for the densification [START_REF] Coble | Diffusion Models for Hot Pressing with Surface Energy and Pressure Effects as Driving Forces[END_REF].

Using the as-defined assembly, the sample heating is provided by both thermal flux coming from the susceptors and dielectric microwave heating. This is the so-called hybrid heating mode. Images are extracted from the thermal camera and are displayed on figure 2. Those images show that (at the early stage of the process) the two SiC plates are first heated by microwaves (Fig. 2.a) and transfer the heat to the sample by thermal radiation. Above approximately 1000°C, the dielectric losses of the spinel increase and it starts to heat by itself. On the photography, it can be seen that the spinel compact preferentially glows (fig. 2b), showing that the latter is being subjected to a hybrid heating process.

Sample Preparation

The starting material used in this study is a spinel MgAl2O4 provided by Solcera (Evreux, France), slip casted according to a proprietary method with S30CR powder from Baikowski (Poissy, France) (>99% purity, impurities are S 700 ppm, Na 41 ppm, Ca 15 ppm, Si 12 ppm, and Fe 6 ppm). The green samples are shaped in the form of small discs of 10 mm in diameter and 5 mm in thickness. The density was measured in ethanol by Archimedes method. The green samples were found to be close to 60% of theoretical density.

The green discs were polished in order to have two parallel faces. Then, a first heating is performed to sinter the material through the first stages of densification. The 8 MPa minimum starting value imposed by the hot press system, and the fact that no mold is used, imply that discs present initially a sufficient mechanical resistance for the upcoming microwave/press treatment. Samples were thus pre-sintered using a conventional furnace at 1400°C for 12 minutes, in air, with a heating rate of 10°C/min. This short heating cycle is sufficient to activate the densification mechanism, while limiting grain growth and avoiding intra-grain porosity entrapment. After this first densification step, the samples show density close to 90% TD, and their dimensions are about 8 mm of diameter, and 4 mm in thickness. Thereafter a final step of densification is carried out with microwave heating under a uniaxial pressure in order to remove all porosities and obtain transparency.

Temperature Calibration

As already mentioned, the temperature is measured by an infrared pyrometer pointing towards the sample surface, located in the center of the cavity. The low detection limit of this pyrometer is 250°C. No detail is given upon the value(s) of the emissivity coefficient used by the system. Using insitu dilatometric curves, it has been experimentally observed that the shrinkage starts at a displayed temperature of ~1100°C, for an applied pressure of 25 MPa (Fig. 4.a.). To evaluate the accuracy of the temperature measurement around that threshold point, a calibration test has been performed in the sintering conditions. Germanium oxide, which has a melting point of 1115°C, was chosen as a reference, following the procedure developed by Macaigne et al. [START_REF] Macaigne | Microwave Sintering of Pure and TiO2 Doped MgAl2O4 Ceramic Using Calibrated, Contactless in-Situ Dilatometry[END_REF]. This experiment was carried out by compacting GeO2 powder on the side of a spinel disc, notched to hold the powder in place. The sample was then positioned at the center of the cavity as described earlier, with GeO2 powder facing a CCD camera at the front side of the machine. This camera is used to record the time at which the melting occurs. At the same time, the infrared pyrometer measures the temperature at the back a b side surface of the MgAl2O4 sample. The transverse view of the cavity is represented in Fig. 3. with both the CCD camera, and the pyrometer. Assuming the sample heats homogeneously and that the germanium oxide at the front side is at the same temperature as the spinel (or very close), the measured temperature can be compared to the expected melting temperature of GeO2.

The observed melting temperature was 1117°C. As the theoretical melting point of germanium oxide is 1115°C, this experiment confirms a good estimation of the temperature value, at least in the range which is relevant for this study.

Characterizations

The transmittance is measured using a Perkin Elmer lambda 1050 spectrophotometer, ranging from 1800 nm to 300 nm, with a detector change at 860 nm. This modification of detectors creates a measurement artifact, which is corrected using the visible detection as reference. The light is produced by a deuterium tungsten source. The measurement consists in comparing the intensity of two 3 mm wide square lasers, passing through two 2 mm wide circular holes. In a first time, an analysis with the sample support, but no sample is carried out to calibrate the machine. Then, the samples are placed on the side of the support, without moving it out of the machine. The measurement is performed within 5 to 6 minutes.

Transmittance measurement has been normalized to the same thickness (2 mm) to allow a comparison between samples. The following relations were used Eq.1-2 [13]:

2 2 ² 1 * 1 1 * 2 1 1
Where IT2 is in-line transmission for the corrected thickness (d2 = 2 mm), and IT1 is in-line transmission measured with real thickness d1. Refraction index n is given by Sellmeier relation:

² 1 1,8939 -0,09942 3,0755 -15,826 2
The determination of the grain size and the microstructures were obtained using a scanning electron microscope Zeiss supra 55, coupled with an EDAX EBSD detector. The sample surfaces were polished down to 1 µm with a diamond suspension, and 0.25 µm with colloidal silica. Since the spinel is insulating, the samples were glued to the support with a silver paint, which helped to evacuate the surface charges. Carbon coating was avoided in order to prevent pattern deterioration since EBSD characterization scans only the first 40 nm in depth [START_REF] Dingley | Progressive Steps in the Development of Electron Backscatter Diffraction and Orientation Imaging Microscopy: EBSD AND OIM[END_REF]. The analysis was made with the sample tilted to 70° under the electron beam towards the EBSD detector, to maximize the data collection. The acquisition was made using the TSL OIM data Collection 6 software, and the results were interpreted using TSL OIM Analysis 6. The step between each measurement point was set to 50 nm, to have at least half a dozen points for each grain. The grain sizes were determined by an automatic intercept method from the software, using 40 horizontal lines.

Hardness and fracture toughness were evaluated using a micro-hardness tester Matsuzawa MMT-X7. A first set of ten indentations was realized with a weight of 1 kgf during 13 seconds on the surface of each sample to compare with results from the literature. The hardness is calculated by the Clemex software and is given in Vickers Hardness (HV). The obtained values were then averaged to get a better precision on the global hardness of the sample. The uncertainties have been calculated using the maximum and the minimum measured values on each sample.

For the fracture toughness, a similar protocol was followed. The half width of the indentation is then measured, as well as the length of the crack, with the same software that was used previously for the hardness. The crack pattern has been determined to be Palmqvist type by Sokol et al. [START_REF] Krell | Advanced Spinel and Sub-Um Al2O3 for Transparent Amor Applications[END_REF]. The value of the tenacity has been calculated using the following equation (Eq.3. [START_REF] Krell | Advanced Spinel and Sub-Um Al2O3 for Transparent Amor Applications[END_REF]) :

1c 0.06 * !"# 4$ 3
Where Hv the hardness in GPa, P the load in Newton (here equal to 9.807 N), and L the crack length in m.

Results and discussion

Sintering process

Using Spark Plasma Sintering on spinel, it has been shown that intermediate values of pressure (60 MPa < P < 80 MPa) and temperatures (1300°C < T < 1600°C) must be used to get transparent materials [START_REF] Rubat Du Merac | Fifty Years of Research and Development Coming to Fruition; Unraveling the Complex Interactions during Processing of Transparent Magnesium Aluminate (MgAl 2 O 4 ) Spinel[END_REF]. By analogy with this sintering technique under uniaxial pressure, similar experimental parameters were chosen as starting conditions here. Tests with pressure of 50 MPa and temperature up to 1100°C-1150°C were conducted. However, samples had been broken under such pressure, probably because no mold was used. The pressure was then gradually decreased through different tests down to 30 MPa. Moreover, to avoid cracking formation in the sample, it was necessary to apply it step by step and at a high temperature, around 900°C, and to remove the pressure at a similar temperature. This method allows the production of crack-free sintered spinel.

Through the observation of the dilatometric curves (Fig. 4.a), for pressure around 20 to 25 MPa, shrinkage starts at around 1100°C. This temperature corresponds to a 750 W delivered power for the assembly and sample size used in this experiment. It must be noticed that using those conditions of temperature and pressure, a creeping mechanism can take place, which does not contribute here to the densification of the ceramics and may induce crack formation, due to cavitation. Therefore, the dwell times were kept short enough to limit this side effect. An optimal dwell time has been empirically determined by calculating the theoretical shrinkage to obtain 100% TD (assuming only axial shrinkage and no radial deformation). According to these criteria, we found an optimal protocol to get crack-free dense sample. The best chosen conditions are as follows: a dwell time of 8 min at 750 W and 30 MPa (Fig. A thin white crown (less dense than in the center) can sometimes be visible on the periphery of the samples. The density gradient can be understood by two phenomena that are inherent to the technique. The first is the sample quite fast cooling from its surface by thermal radiation due to an imperfect thermal isolation. The second is the lack of radial constraints implying a pressure gradient from the center of the disk to its periphery. The impact of the dwell time on the physical properties of the spinel has been studied through three experiments performed with dwell times of 5, 8, and 10 min. After the experiment, the density of each sample was measured and found to be 100% of theoretical density. However, for a similar starting density, the total shrinkage is slightly higher for samples with longer dwell times (Fig. 5.).

Characterization 3.2.1. Transmittance

During the measurement, two interfaces between spinel and air are crossed by the light beam. Those two interfaces, with abrupt change in refraction index, lead to a partial reflection of the laser (6-7% at each interface). This phenomenon implies a theoretical limit to the transmittance of 87% of light intensity without anti-reflect coating [21].

Figure 6 shows the value of the transmittance as a function of wavelength for the different samples. An artifact of measurement can be seen around 900 nm, due to the detector change during the analysis. It appears that an increase of the dwell time first improves the quality of the optical properties, with a higher transmittance in the IR domain for the 8 minutes dwell samples. Transmittance as high as 60 % at 600 nm, 77% at 1000 nm and 83% at 1600 nm have been obtained with this new method. Those transmittance results are similar to spinel sintered by SPS, but are still inferior to the results obtained by HIP densification in the visible range [START_REF] Rubat Du Merac | Fifty Years of Research and Development Coming to Fruition; Unraveling the Complex Interactions during Processing of Transparent Magnesium Aluminate (MgAl 2 O 4 ) Spinel[END_REF]. It seems that there is an optimum sintering duration, since a dwell time of 10 min reduces the transmittance by 5 to 10 % for 2 mm thick samples, as can be seen on Fig. 6. This optimum might be explained by the sintering method used in this experiment. Grains sliding mechanism takes place through the creeping phenomena, in resulting the pores formation by cavitation for longer dwell times [START_REF]Latest Global Developments in Microwave Materials Processing[END_REF].

Even if the measurement of density by Archimedes method in ethanol gave a density of ~ 100%, there is still some residual porosity in the samples. This remaining porosity can explain the decrease of transmittance in the visible range. It has been shown that with 0.05% of pores 100 nm wide, transmittance is reduced by half at 600 nm [21]. Furthermore, with 0.05% of porosity, the diameter of the pores must be less than 20 nm to not affect the in-line transmittance [21]. However, it has not been possible to measures either the amount of porosity nor the sizes of the pores in this work.

Grain size

To control if the heating has been homogeneous, measurements were done at different areas from the middle of the discs to the edge. No significant difference in microstructure or in grain size has been observed in function of the position, confirming that the volume heating has occurred and that the samples are homogeneous.

A typical result of EBSD analysis on the spinel samples studied here is shown in Fig. 7. with an auto grain and an IQ with a measurement done on the sample which was subject to a 5 min dwell time. The mean grain size measured for the samples with 5 and 8 min dwell time is 0.45 and 0.5 ± 0.05 µm, and for the sample with 10 min, the mean grain size is 0.55 ± 0.05 µm. This process seems consistent for the production of transparent spinel with sub-micron grain size. The short stage at high temperature combined with pressure is likely responsible for this observed limited grain growth. The submicron grain size obtained here should have a good impact on the mechanical properties of this material. Hence, the hardness and fracture toughness were characterized on these three samples.

Hardness and fracture toughness

Figure 8 shows the results of mechanical characterizations. For comparison, dash lines figuring the results from the literature in the same measurement conditions have been drawn [START_REF] Krell | Advanced Spinel and Sub-Um Al2O3 for Transparent Amor Applications[END_REF][START_REF] Goldstein | Transparent Polycrystalline MgAl2O4 Spinel with Submicron Grains, by Low Temperature Sintering[END_REF].

Hardness values are ranging from 17 to 15 GPa depending on the dwell time (Fig. 8.). Hardness decreases with longer dwell time, which is consistent with the greater grain sizes observed. Fracture toughness is almost constant whatever the dwell time, and remains close to values reported in the literature on HIPed samples [START_REF] Krell | Advanced Spinel and Sub-Um Al2O3 for Transparent Amor Applications[END_REF][START_REF] Goldstein | Transparent Polycrystalline MgAl2O4 Spinel with Submicron Grains, by Low Temperature Sintering[END_REF]. Overall, mechanical properties are similar to those measured for HIPed samples. However, long sintering time presents two major drawbacks. The first is obviously the grain growth, and the second is the creation of porosity due to grains sliding, which alters the optical properties.

Conclusion

Sub-micrometric transparent polycristalline spinel MgAl2O4 has been obtained in a short thermal process (8 minutes dwell time at the high temperature stage) using a new microwave pressureassisted system. Inspired from sinter-forging and adapted to the microwave furnace, this process was successfully used to produce samples with high transmittance values and high mechanical properties due to sub-micrometric grain size. This new technique allows the production of transparent spinel in air, with a transmittance up to 83 % in the IR range.

Fig. 1 .

 1 Fig. 1.Scheme top view of the 2.45 GHz microwave furnace

Fig. 2 .

 2 Fig. 2. Photography of a spinel sample showing radiative heating below 1000 °C from the SiC susceptors (a) and hybrid microwave heating over 1000 °C (b)

Fig. 3 .

 3 Fig. 3. Scheme of the experimental set-up for the temperature calibration test

  4.b).

Fig. 4 .

 4 Fig. 4. a) In-situ dilatometric curve b) Temperature and pressure cycle for a 5 min dwell experiment

Fig. 5 .

 5 Fig. 5. Dilatometric curves as a function of time for corresponding dwell times, with corresponding samples in front of the CRISMAT logo. Samples are sorted from shorter to longer dwell time in descending order.

Fig. 6 .

 6 Fig. 6.Transmittance as a function of wavelength normalized for 2 mm thick samples: Impact of the dwell time on the quality of the transparency

Fig. 7 .

 7 Fig. 7. a) Auto grain image generated with EBSD data of a transparent spinel b) IQ image generated by OIM

Fig. 8 .

 8 Fig. 8. Hardness and fracture toughness depending on dwell time compared to HIPed references

Acknowledgement

The authors would like to thank Solcera for the samples, and the Labex for the financial support. The authors are also thankful to Jerome Lecourt and Jean-Pierre Burnouf for their precious help in conducting the experiments.