M Samnouni

Nicolas Wichmann

X Wallart

Christophe Coinon

Sylvie Lepilliet

S Bollaert

InAlAs PHEMT

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

I. INTRODUCTION

Electron-Mobility Transistors (HEMTs) technology is the best candidate for low noise amplifier and Terahertz applications. Due to their superior high-frequency characteristics InP-based HEMTs are very promising for millimeter-wave applications and image sensor systems [START_REF] Sato | Compact receiver module for a 94 GHz band passive millimetre-wave imager[END_REF]. Currently, the record for fmax is 1.5 THz (fT = 610 GHz) [START_REF] Mei | First Demonstration of Amplification at 1 THz Using 25nm InP High Electron Mobility Transistor Process[END_REF] with pseudomorphic-HEMTs (PHEMTs) on InP substrate. In fact, for the last 10 years, the strategy used to improve electrical performance of HEMT was to reduce the gate length below 50 nm combined with epilayer engineering. This approach allows high cutoff frequencies and excellent transconductance with a record gm of 3.1 mS/µm [START_REF] Mei | First Demonstration of Amplification at 1 THz Using 25nm InP High Electron Mobility Transistor Process[END_REF] , recently published papers report high fmax while keeping longer gate length LG = 75 nm [START_REF] Takahashi | Maximum frequency of oscillation of 1.3 THz obtained by using an extended drain-side recess structure in 75-nm-gate InAlAs/InGaAs high-electron-mobility transistors[END_REF].

II. FABRICATION

InAlAs/InAs/InGaAs Pseudomorphic HEMT (PHEMT) epitaxial layers were grown on semi-insulating InP substrate by gas source molecular beam epitaxy (MBE). The layer structure (figure 1), from bottom to top, consists of an In0.52Al0.48As buffer layer, a bottom silicon-delta doped plane, an In0.52Al0.48As first spacer, an In0.53Ga0.47As/InAs/ In0.53Ga0.47As composite channel, an In0.52Al0.48As second spacer, a top silicon-delta doped plane, an In0.52Al0.48As barrier, an InP etch-stop layer, and finally a thick In0.53Ga0.47As cap layer. The epilayer structure was optimized for high electron density and mobility in the channel. To achieve high channel mobility, strained pure InAs layer was inserted in InGaAs channel. Room temperature mobility of µHall = 12000 cm 2 /V.s and a NHall = 2.65x10 12 cm -2 are obtained by Hall measurement. For PHEMT fabrication, electron beam lithography was used. The first step of the fabrication was mesa isolation by wet chemical etching down to the InAlAs buffer layer. Next, Ohmic contact was realized as follow: after defining the Drain-Source with 1.5µm distance by e-beam lithography, Ti/Pt/Au metals was evaporated.

The last process step concerns the asymmetric recess. The use of asymmetric recess will be motivated in the third part of this paper. For the recess, we need first to remove the InGaAs cap layer. Shinohara et al. [START_REF]Simple and high-precision asymmetric gate-recess process for ultrafast InP-based high electron mobility transistors[END_REF] reported one of the fabrication technique for an asymmetric gate recess. This technique based on three-layers resist. The difficulty of this technique is observed during the opening of the slits that requires several precisions for the sizing, wet and dry etching. Furthermore, the recessed area is not passivated. Takahashi et al. [START_REF] Takahashi | Enhancement of fmax to 910 GHz by Adopting Asymmetric Gate Recess and Double-Side-Doped Structure in 75-nm-Gate InAlAs/InGaAs HEMTs[END_REF] also reported a simple technique for an asymmetric gate recess, the process used based on one resist layer. After removing, the InGaAs cap layer the surface is passivated by Si3N4 film. Next, the gate electrode was deposited after definition of the gate footprint in Si3N4 by reactive ion etching RIE using SF6 gas. However, the etching of Si3N4 using SF6 have an effect on the electrical characteristics of the channel.

To verify the impact of the plasma etching, Hall measurements were performed after removing the Si3N4. The mobility and the electron density were degraded by 10% and 25% from 12000 cm 2 /V.s to 11000 cm 2 /V.s and from NHall = 2.65x10 12 cm -2 to NHall = 1.97x10 12 cm -2 respectively. Then, we used a different technique. After e-beam lithography defining the recess zone, the cap layer was removed by wet chemical etching. The second e-beam lithography was performed in bilayers resists to define the shape of the T-gate. Ti/Pt/Au metals was evaporated. Finally, the sample was loaded into PECVD chamber where 30 nm-thick Si3N4 were deposited to form a passivation layer and decreasing the sheet resistance in recessed zones. This technique allows the asymmetry of the recess.

III. RESULTS

In this part, we will focus on the RF characterization. Concerning the DC measurements, the drain current is around ID = 800 mA/mm and ID = 1000 mA/mm for the reference and optimized PHEMT respectively. In both HEMTs, the recess length on the source side LRS is 65 nm. The difference was in the recess on the drain side LRD (figure1). LRD is 175 nm and 225 nm for the referenced and optimized transistor respectively. The gate to source distance was reduced by 200 nm for the optimized transistor. The extrinsic peak transconductance gm of 1400 mS/mm was obtained at VDS = 0.8 V for both transistors. S-Parameters were measured up to 110 GHz. The extrapolation by (20 dB/decade) of the unilateral Mason's gain U gives fmax = 800 GHz for the reference PHEMT and fmax = 1.2 THz for the optimized PHEMT at VDS = 1 V (figure 2).

To understand the impact of increasing the drain side recess on fmax, four devices were characterized up to 110 GHz. The source side recess was fixed at ~ 65 nm and LRD varies from 70 nm to ~ 195 nm. The results are shown in figure 3. The symmetric devices, with LRD = LRS ≈ 65 nm had a fmax = 520 GHz. As LRD increased, the gain becomes more and more important thanks to the improvement of intrinsic parameters of the transistor also shown in figure 4. This indicates that the coupling between drain and gate is lower comparing to the gate source coupling [START_REF] Ballegeer | Physics and behavior of asymmetrically recessed InP-based MODFET's fabricated with an electron beam resist process[END_REF]. In order to extract the four noise parameters of the reference PHEMT, NF50 method [START_REF] Dambrine | A new method for on wafer noise measurement[END_REF] has been used. The results achieved at 94 GHz were NFmin of 1.8 dB (with associated gain of 11.6 dB) for an optimal gate voltage VGS_Opt = -0.3 V are shown in table 1.

IV. CONCLUSION

In conclusion, we have fabricated PHEMT operating above 800 GHz using an asymmetric gate recess with large gate length of 75 nm. This performance was obtained thanks to improvement of the epilayer structure with mobility channel of µHall = 12000 cm 2 /V.s. The use of an asymmetric recess improve the intrinsic parameters such as Cgs and Cgd capacitances. The output conductance gd was also improved when LRD increasing.

Fig. 1 .

 1 Fig. 1. Cross-section schematic of epitaxial layer structure.

Fig. 2 .Fig. 4 .Fig. 3 .

 243 Fig. 2. Measured U/|h21|² for reference and optimized PHEMT (LG = 75 nm and W=2x12µm)


Dichao Shi

Qifeng Yang

Christi Peterson

Anne-Félicie Lamic-Humblot

Jean-Sébastien Girardon

Anne Griboval-Constant

Moulay Lorenzo Stievano

Tahar Sougrati

Valérie Briois

Paul A J Bagot

Jean- Sébastien Girardon

Moulay T Sougrati

Robert Wojcieszak

Sébastien Paul

Dr Eric Marceau
email: eric.marceau@univ-lille.fr

Bimetallic Fe-Ni/SiO2 catalysts for furfural hydrogenation: Identification of the interplay between Fe and Ni during deposition-precipitation and thermal treatments

Keywords: bimetallic catalysts, furfural, X-ray absorption spectroscopy, Mössbauer spectroscopy, in situ measurements

published or not. The documents may come

Introduction

Noble metals have long been used in catalysis for the hydrogenation of organic molecules, but their natural abundance is low and, consequently, their cost is high and subject to large variations. Substitutes should be found among more abundant and cheaper transition metals. Fe is abundant, but it is also poorly active and poorly stable. In contrast, Ni is known for its activity, but it faces deactivation by carbon, and can favour hydrogenolysis or overhydrogenation side-reactions. However, associating Fe and Ni in supported bimetallic catalysts has been found to be beneficial in terms of activity and selectivity for reactions between H2 and ketones [1,2], aldehydes [2][3][4][5], esters [6,7], phenols [8] or nitriles [1,9]. Fe25Ni75/SiO2 (formulations are herein given in molar proportions) was first mentioned by Khumbar et al. to be more active than Ni/SiO2 in the hydrogenation of acetophenone and nitriles [1]. More recently, Yu et al. have reported that 5-hydroxymethyl furfural was almost quantitatively hydrogenated into the corresponding diol at 110°C with Fe-Ni catalysts, while the selectivity was lower than 80% with a monometallic Ni catalyst [4].

The structure of reduced Fe-Ni nanoparticles is strongly dependent on the Fe/Ni ratio.

Fe-rich nanoparticles exhibit a bcc structure while Ni-rich alloys belong to the fcc type [10].

Furthermore, as recently reviewed by Tomishige et al. [11], the particles sizes are often not uniform when the catalyst is prepared by co-impregnation, and can exceed 10 nm. The control of the Fe/Ni ratio in individual nanoparticles is also difficult. Monometallic nanoparticles, and bimetallic Fe-rich and Ni-rich nanoparticles, may co-exist in the catalyst [1,4,[12][13][14][15]. If the association between the two metals is not effectively controlled, it can be difficult to link the best catalytic results to precise Fe and Ni proportions in the catalyst formulation.

The thermal treatment of mixed Fe-Ni phases such as hydrotalcites appears to lead to well-defined systems [11,16,17]. In a few instances, deposition-precipitation has been implemented to yield oxide-supported Fe-Ni nanoparticles of controlled size. The decomposition of co-precipitated Fe and Ni cyanides allowed van de Loosdrecht et al. to attain particles sizes smaller than 4 nm on TiO2 [18]. Deposition-precipitation involving the decomposition of urea (DPU) was recently tested on alumina by Mutz et al. [17], resulting in a particle size distribution ranging from 2 to 7 nm for the Fe25Ni75 formulation, but the speciation of the metals in the first stages of the DPU process was not described.

DPU has been thoroughly investigated as a method to prepare well-dispersed monometallic Ni nanoparticles on silica, even at metal loadings as high as 40 wt.% [19][20][21][22].

Its extension to Fe-based bimetallic nanoparticles raises the question of the association between the two metals upon deposition and thermal treatments. We discuss here the choice of the synthesis parameters required to avoid early metal segregation during DPU, and we evidence the interplay existing between Fe and Ni-containing species along the preparation of a Fe-Ni/SiO2 catalyst. A combination of in situ characterization techniques (X-ray diffraction -XRD; X-ray absorption spectroscopy -XAS; Mössbauer spectroscopy -MössS) is associated to overcome the difficulty to distinguish the two poorly Z-contrasting metals, as conventional transmission electron microscopy or EXAFS alone could not differentiate between two elements with close numbers of electrons. Quick-X-ray absorption spectroscopy recorded simultaneously at the Ni and Fe K edges in in situ conditions will provide information on the evolution of the oxidation state and chemical environment of the two metals after DPU and during reduction, while Mössbauer spectroscopy will give valuable data on the speciation of ionic and metallic Fe at all stages of the preparation.

Experimental

Catalysts preparation

A Fe50Ni50/SiO2 catalyst (in which 50/50 represents the relative atomic proportions of the two metals in the initial solution) was prepared using a modified DPU procedure originally developed for Ni/SiO2 catalysts [19][20][21][22][START_REF] Hermans | Preparation of Catalysts II[END_REF]. Sipernat-50 silica (Degussa; surface area, 400 m 2 g -1 , pore volume, 1.4 cm 3 g -1 ) was used as a support. Metal salts used were iron (II) sulfate heptahydrate (Alfa Aesar, 98%), nickel (II) nitrate hexahydrate (Merck, 99%) and nickel (II) sulfate hexahydrate (Aldrich, 99%). Other reactants were nitric acid (Prolabo, 70%), sulphuric acid (Alfa Aesar, 93-98%) and urea (Alfa Aesar, 99-100%). All reactants were used without further purification. 150 mL of aqueous solution of 0.84 mol L -1 urea, 0.02 mol L -1 H + introduced via a strong acid, 0.14 mol L -1 Ni(II) salt, 0.14 mol L -1 Fe(II) salt, were mixed with 1.14 g of silica in a three-necked flask allowing Ar bubbling. The suspension was continuously stirred and slowly heated to 80°C, resulting in the decomposition of urea, increase of the pH to 5.3, and deposition-precipitation of metallic compounds onto the support. The suspension was kept at 80°C for 22 h. After cooling to room temperature (RT), the solid was recovered by centrifugation, washed three times with distilled water and dried at RT. The choice of the reactants, their order of introduction, and the conditions of washing, drying and storage are discussed in section 3.1. The catalyst was reduced by heating in a flow of 10% H2/He (100 mL min -1 ) at 5°C min -1 up to 700°C (with a plateau of 2 h), without intermediate calcination.

A Ni100/SiO2 catalyst was prepared as a reference following the same procedure (Ni(II) concentration in solution: 0.28 mol L -1 ). It was verified by transmission electron microscopy that, in line with formerly reported results on this system [21,22], the sizes and distribution on the support of metal particles were similar to those of Fe50Ni50/SiO2. Two reference systems providing other Fe/Ni proportions (0.21 mol L -1 Fe(II) sulfate, 0.07 mol L -1 Ni(II) sulfate; 0.07 mol L -1 Fe(II) sulfate, 0.21 mol L -1 Ni(II) sulfate), and a reference Co-Ni/SiO2 system prepared from 0.14 mol L -1 Ni(II) nitrate hexahydrate and 0.14 mol L -1 Co(II) nitrate hexahydrate (Sigma Aldrich, 98%), were prepared for the sake of comparison, for the characterization of the dried systems (section 3.1) and to support a discussion on metal contents (section 3.2), respectively.

Characterization

Fe, Ni and Si elemental analysis was performed by inductively coupled plasma-optic emission spectroscopy using a 720-ES ICP-OES spectrometer (Agilent) with axial viewing and simultaneous CCD detection. N2 physisorption isotherms at liquid nitrogen temperature were carried out on a TriStar II Plus gas adsorption analyzer (Micromeritics) after activation at 150 °C in vacuum overnight. The specific surface area was evaluated with the Brunauer-Emmett-Teller (BET) model over the range P/P0 = 0.05-0.30. The total pore volume was measured from the volume of N2 adsorbed at P/P0 = 0.95. Temperature-programmed reduction (TPR) was followed using an Autochem apparatus (Micromeritics) under H2 (5%)/Ar (flow rate: 50 mL min -1 ; ramp rate: 7.5 °C min -1 ). Attenuated Total Reflection Infra-Red (ATR-IR) spectra were recorded using a Nicolet Magna iS50 spectrometer equipped with a diamond crystal and a DTGS detector (10 scans; resolution: 4 cm -1 ).

X-ray diffraction patterns (XRD) were recorded in ambient conditions using the Cu Kα radiation (l = 1.5418 Å; 40 kV, 30 mA) on a Siemens D5000 diffractometer, with a 0.05° scan step and 2 s time step. In situ measurements were conducted under H2 (3%)/N2 on a D8 Bruker diffractometer equipped with an Anton Paar XRK900 reactive chamber. Data were acquired in the Bragg-Brentano geometry, using the Cu Kα radiation (40kV, 40mA).

Diffractograms were recorded every 25 ˚C from 50 to 700 ˚C. A peak from the holder at 2q = 27.2° was used for calibration between different experiments. The average size of metal nanoparticles was evaluated by the Laue-Scherrer equation applied to the (111) reflection. 57 Fe Mössbauer spectra (MössS) were measured with a source of 57 Co embedded in a rhodium metal matrix. Measurements were performed in the transmission mode, with the source at room temperature. The measured sample, in the form of a self-supported pellet, was mounted in an in situ oven working at controlled temperatures up to 700°C under diluted H2 flow. The gas flow was maintained during cooling to 25°C for recording of the spectrum at RT. Low temperature spectra (-268 °C) were measured after transferring the sample pellet to a helium gas flow cryostat in an Ar-filled glove box (O2, H2O < 1 ppm), strictly avoiding any contact with ambient air. The spectrometer was operated with a triangular velocity waveform, and a gas-filled proportional counter was used for the detection of the gamma rays. Velocity calibration was performed with an a-Fe foil. The isomer shift δ, the electric quadrupole splitting D, the magnetic hyperfine field BHF, the linewidth at half maximum G and the relative absorption area of the different components were determined by fitting the spectra with the appropriate combinations of Lorentzian profiles. Time-resolved XAS data were measured in the transmission mode at the ROCK beamline [START_REF] Briois | [END_REF] using the SOLEIL home-made Quick-EXAFS monochromator equipped with a Si (111) channel-cut crystal and described in ref. [25]. The channel-cut oscillates with a ± 1.95° amplitude around 14.6 ° and with a frequency of 2 Hz, allowing simultaneous recording of the Fe (7112 eV) and Ni (8333 eV) K-edge spectra in 250 ms. Successive spectra collected with upward Bragg angles during the in situ monitoring of catalysts were merged to improve the signal-to-noise ratio, leading to a time resolution of 5 s per spectrum effectively analyzed.

Ionization chambers provided by Ohyo Koken Kogyo Co. Ltd. filled with a 2:1 mixture of N2 and He were used for measurements. The beam size was 2.5 x 0.5 mm. Normalization of spectra was performed by using the Python normal_gui graphical interface developed at SOLEIL for the fast handling of Quick-XAS data. The proportions of the different Fe and Ni species were determined by a multivariate data analysis, the MCR-ALS method [26][27][28].

MCR-ALS GUI 2.0 developed by the Tauler group [29], and freely available as a Matlab toolbox, was used for MCR-ALS minimization. The EXAFS signal extraction and Fourier transformation of the EXAFS spectra were done using the Athena graphical interface software [30]. EXAFS fitting of distances, coordination numbers and Debye-Waller factors was performed with the Artemis interface to IFeFFIT using least-squares refinements [31], firstly checked on NiO, Ni metallic foil and Fe2O3 references. Fourier-transformed EXAFS signals are presented without phase correction. The temperature-programmed reduction of the Fe50Ni50/SiO2 sample was performed using a dedicated cell [32] connected to the gas-feeding system installed on the ROCK beamline. The cavity of the sample holder (2 mm thick) was filled with the powder catalyst (20 mg) pre-mixed with boron nitride (BN, 10 mg). The in situ catalyst reduction was performed by heating the cell from RT to 500 °C under a H2 (5%)/He flow (50 mL min -1 ) with a heating rate of 5°C min -1 , followed by 2 h of isothermal treatment.

TEM observations were made at University of Oxford, using a JEOL JEM-2100 microscope fitted with a LaB6 source (JEM-2100) for bright field observation, and on a JEM-3000F FEGTEM microscope (JEM-3000) fitted with a spherical aberration STEM probe corrector and a Centurion EDX detector equipped with a super atmospheric thin window (SATW) detector and an INCA processor, for elemental analysis. E-beam size was selected to 1.2 nm for EDX point scan measurements on individual particles.

Catalytic tests

Fe50Ni50/SiO2 and Ni100/SiO2 catalysts were tested in the hydrogenation of furfural using a Screening Pressure Reactor (SPR) system from Freeslate equipped with 24 parallel stainless steel batch reactors of 6 mL each. 10 mg of catalysts, reduced at 700°C (section 2.1), were placed after transfer in ambient air into each of the 24 reactors and exposed in situ to pure H2 at 400°C (1 bar, 2 h) prior to reaction. 2 mL solutions of furfural (0.20, 0.56 or 1.00 mmol) in isopropanol were then injected under the controlled atmosphere of a glove box.

Tests were performed at 150°C for 2 h under 20 bar H2 and a continuous stirring of 700 rpm.

Reaction products were analyzed using Shimadzu GC-2010 Plus and GC-FID-MS-QP2010 Ultra EI gas chromatographs, equipped with a FID detector, a ZB-5MS column (30 m × 0.25 mm × 0.25 μm), and for the latter, with a QP2010 Ultra mass spectrometer. It was verified by quantifying acetone that in these experimental conditions, hydrogen transfer from isopropanol to furfural did not occur to a significant extent (less than 3% of converted furfural).

Results and discussion

Choice of the experimental parameters for the synthetic procedure and characterization of dried catalysts

Given the solubility products of iron and nickel hydroxides (pKs Fe(OH)3 = 38.8; Fe(OH)2 = 15.1; Ni(OH)2 = 15.2 at 25°C), Fe 2+ ions are expected to deposit in the same pH range as Ni 2+ ions, and Fe 3+ ions at much lower pH. Fe(II) sulfate was thus selected as the iron precursor salt for DPU, in preference to Fe(III) nitrate. In spite of this, initial attempts of DPU with nickel nitrate as precursor salt and nitric acid as acidifier led to the oxidation of In order to eliminate oxidizing species, nickel (II) sulfate and sulphuric acid were instead used as reactants, but the preparation still suffered from a lack of reproducibility. Diffraction peaks of Fe(III)-containing layered hydroxysulfate phases [33][34][35] were observed after drying, with ulterior formation of large Fe-Ni particles upon reduction (Fig. S1, Supplementary Material).

It was next chosen to degas with Ar the suspension containing silica, sulphuric acid, nickel (II) sulfate and urea for 2 h, and to introduce Fe (II) sulfate only after degassing. Ar bubbling was maintained during the whole heating and cooling steps. In these conditions, deposition took place at pH 5.3 and the suspension remained green. Centrifugation and washing were also carried out with degassed water.

Drying at RT in ambient air caused a rapid change of colour of the solid from green to ochre, evidencing the oxidation of Fe(II) to Fe(III). Such oxidation is prone to happen to Fe(II)-containing layered hydroxides [36][37][38]; it has been shown that the presence of Ni(II) ions inhibits their further transformation to Fe3O4 [39,40]. Drying the solid under Ar flow in the three-necked flask, and storage under Ar, preserved the green colour longer, but oxidation of iron to Fe(III) was nevertheless seen to develop over a period of six weeks. Using XRD and XAS, it was checked by reducing dried solids either kept under Ar or aged in air that, unlike Fe oxidation during DPU, Fe oxidation upon drying and storage had no detrimental effect on the formation, nature and size of the reduced metal nanoparticles.

The XRD pattern of all dried solids exhibits a broad diffraction peak at about 22˚, characteristic of amorphous siliceous solids, a broad asymmetrical line at 2q = 34˚ and a narrow symmetrical one at 2q = 60˚ (Fig. S1). As previously shown for Ni/SiO2 catalysts prepared by DPU, the latter two peaks can be assigned to structures based on brucitic layers, either M II (OH)2 turbostratic hydroxides or ill-crystallized phyllosilicates [19].

MössS confirms the exclusive presence of octahedral Fe(III) ions in dried

Fe50Ni50/SiO2 after exposure to air (Fig. 1a, Table 1). The signal recorded at RT is asymmetric and must be fitted with at least two different quadrupole split contributions, as is often done for ions present in a continuum of environments. Given the relatively narrow range of isomer shifts and quadrupole splittings for octahedral Fe(III), it is virtually impossible to distinguish between Fe(III) oxyhydroxides and silicates [41]. However, the absence of a sextet allows discarding the presence of well-crystallized goethite FeOOH, which is expected to exhibit magnetic order at room temperature [42,43]. Moreover, the relatively low isomer shift of the narrow doublet, close to that measured for octahedrally coordinated iron in mixed metal Ni-Fe ferrites, suggests the formation of mixed-metal species [44]. To test this hypothesis, Mössbauer spectra were recorded on dried solids prepared following the same protocol but with two other Fe and Ni nominal contents (Fig. S2, Table S1). The significant changes in the quadrupole splittings, as well as the slight decrease of the average isomer shift when changing the metal contents, suggests that Fe(III) ions are co-deposited with Ni(II) ions.

Finally, electric quadrupole splittings are noticeably larger here than in well-defined oxyhydroxides and silicates, especially that of the second doublet. This is usually interpreted as the sign of distorsions in the chemical environment: distorted octahedra, surface sites, proximity to tetrahedral sites or to divalent cations, such as Ni 2+ , in hydroxides or silicates [39,42,[45][46][47].

According to the analysis of EXAFS oscillations, the local environment of both Fe 3+ and Ni 2+ ions is consistent with the octahedral sites of brucitic layers, common to the structures of divalent hydroxides and phyllosilicates [19] (Fig. 2, Fig. S3, Table 2). The first shell is satisfactorily fitted using 6 oxygen atoms (R = 1.99 Å for Fe 3+ , R = 2.05 Å for Ni 2+ ), probably in a more distorted symmetry in the case of Fe(III) as shown by the higher value of the Debye-Waller factor. The best fits for the second shell of neighbours involve 6 Ni atoms at R = 3.11 Å, characteristic of edge-sharing octahedra (it must be noted though that, due to the closeness of scattering phase and amplitude functions for Fe and Ni, the nature of the second neighbours as backscatterers cannot be unambiguously discriminated by EXAFS). If existing, iron environments involving two very different Fe-Fe distances above 3 Å, as encountered in FeOOH phases [46,48,49], are not predominant.

Evolution and composition of the phases deposited during DPU

The nature of the deposited phases and the metal contents were investigated as a function of DPU duration, from 1 to 22 h. While the X-ray diffractograms did not exhibit major changes besides the gradual broadening of the large halo associated to silica (Fig. S4), the evolution of ATR-IR bands (Fig. 3a) revealed that, following the pattern reported for monometallic Ni/SiO2 catalysts [19], the brucitic M II (OH)2 phase formed during the first 4 h of DPU evolved into an ill-crystallized 1:1 phyllosilicate. This is proved by ( 1 ICP measurements indicate that, despite the equimolarity between Fe and Ni in solution, nickel is the main metal initially depositing during the first hours (Fig. 3b). After a moderate drop between 1 and 4 h, the nickel content increases by a two-fold between 4 and 8 h, when the phyllosilicate phase forms. In contrast, the Fe content steadily increases throughout the process. As a result, a significant excess of nickel with respect to iron is present after 22 h DPU in the dried solid, compared with the atomic proportions in solution:

38 Fe at % and 62 Ni at % are measured by ICP on the solid, corresponding to 42% of the amount of Fe and to 67% of the amount of Ni initially introduced in solution.

This excess of nickel does not originate from its early introduction in the DPU solution, as it was also observed when nickel and iron sulfates were added together in solution after Ar degassing. For the sake of comparison, a reference system was prepared from nickel and cobalt (II) nitrates (nominal proportions: 50 Co at % and 50 Ni at %), without degassing and separate introduction of the reactants. The atomic proportions given by EDX measurements (34 Co at %, 66 Ni at %) prove that nickel tends to react with silica preferentially, compared with the other metal.

Finally, no IR band assigned to isocyanates or carbonates was observed in the spectra.

A shoulder at 607 cm -1 could be assigned to sulfate ions (n4 vibration mode [50,51]), but it was verified by ICP and XPS that sulfate ions were not detected on the dried solid.

In summary, the early precipitation of Ni(OH)2 onto silica during the first hours of DPU is followed by the formation of a poorly organized 1:1 phyllosilicate, in which, after drying, Fe 3+ and Ni 2+ ions occupy the octahedral layer. The (Fe+Ni)/Si ratio of 1.3 after 22 h of DPU is consistent with a 1:1 phyllosilicate stoichiometry (M/Si ratio varying between 1 and 1.5 for exclusively trivalent and divalent cations, respectively). In the absence of carbonates or sulfates, charge compensation can be ensured by deprotonation of the layers, and formation of O 2-or OH -anions formed upon oxidation of iron during drying and storage.

FeOOH as a separate phase is not predominant, but the existence of minor domains cannot be excluded as MössS evidences a distribution of environments around Fe 3+ ions.

Reduction process

The temperature-reduction profile of Fe50Ni50/SiO2 (after storage in ambient air and drying at 60°C to eliminate adsorbed water) evidenced three main reduction stages (Fig. S5): a minor one between 270 and 350°C, and two major ones between 350 and 500°C, and between 500 and 700°C, respectively. In situ X-ray absorption spectroscopy and Mössbauer spectroscopy were used to link these events to the reduction of Fe and Ni-containing species.

On the basis of the TPR profile, 700°C was selected as the reduction temperature at which both metals were expected to be quantitatively reduced.

The reduction of air-stored Fe50Ni50/SiO2 was monitored simultaneously at the Ni and Fe K edges by in situ XAS up to 500°C, the upper temperature limit allowed by the set-up (Fig. 4a andb). In line with the presence of a set of isosbestic points on the series of spectra recorded at the Ni K edge, the quantitative analysis by MCR-ALS procedure (whose principle is recalled in the Supplementary Material) requires two spectral components only (Fig. 4c, Fig. S6), that recall the experimental spectrum of Ni 2+ ions in the initial phyllosilicate, and that of metallic fcc Ni as confirmed by the EXAFS analysis of the final state (Fig. 2b andc, Fig. S3b andc, Table 2). The reduction of Ni 2+ ions starts above 350°C, with an inflection point at 400°C. Nickel reduction is complete after a plateau of 1 h at 500°C. The absence of an intermediate dehydrated phase in the analysis is justified by the higher temperature range at which nickel phyllosilicate would decompose to NiO (500-600°C) [22].

The progress of iron reduction, deduced from the spectra recorded at the Fe K edge, is more complex to analyse. The lack of isosbestic points indicates that more than 2 components are involved in the reduction, and the EXAFS signal of the final state could not be satisfactorily fitted with a limited set of parameters. It was thus chosen in a first approach to follow the decrease in energy of the Fe edge position, correlated to a decrease of the Fe oxidation state (Fig. 4d). The position in energy of the edge was interpreted with respect to reference compounds: hematite Fe2O3, Fe(II) sulfate, Fe foil. Fe 3+ ions originally present inside the layered structure underwent an accelerated reduction to Fe 2+ in the 275-350 °C temperature range. A more gradual decrease of the edge value above 350°C, starting when Ni reduces to the metallic state, can be interpreted as a progressive reduction of Fe 2+ iron to Fe metal triggered by Ni reduction. A comparison between the final XANES spectrum and the spectrum of reference metallic Fe shows that the reduction of Fe 2+ is not complete at the end of the plateau at 500°C. The MCR-ALS analysis of Fe reduction provides five independent spectral components (Supp. Mater., Fig. S7 andS8). The white lines of the first two components are identical, and the EXAFS oscillations are found at the same energies, though damped for the second component. The third component appears between 275 and 358°C. As seen above, the lower edge energy is associated to the Fe(II) oxidation state. The EXAFS oscillations are similar to those of the first component, though shifted to lower energies. The Fourier transforms of the first three components resemble each other, apart from the decrease of intensity of the peaks due to a temperature-induced increase of the Debye-Waller parameter.

The Fourier transform of the third component also differs from those recorded on oxides such as Fe3O4. This suggests that the set of the first three components is related to the same phase, i.e., the Fe-containing phyllosilicate, the second component reflecting mostly an increase of disorder with temperature (data related to the first two components are provided in Fig. S9 and Table S2), and the third component a reduced Fe 2+ state of iron ions in the layered phase.

The first event observed on the TPR profile is thus linked to the reduction of Fe 3+ to Fe 2+ .

The third component starts disappearing when nickel reduction starts, while the fourth component appears exactly in parallel with the formation of metallic nickel. The position in energy of the edge of the fourth component lets one assume that it is associated to Fe 2+ ions, though in a strongly altered environment: the shape of the first peak in the Fourier transform, which seems to contain both Fe-O and Fe-Fe (or Fe-Ni) contributions, suggests both the existence of an oxidic phase, and a proximity to metal particles. The last component, whose weight increases gradually during the plateau at 500°C, gives a well-defined Fourier transform identical to that of metallic Ni, and is interpreted as Fe(0) atoms located in the fcc environment of Ni-containing nanoparticles. In view of this last result, the fourth component may be understood as a Fe(II)-containing phase external to the Fe-Ni particles, from which Fe reduces and diffuses into the bimetallic nanoparticles upon increasing temperature and time.

These processes and the concomitant reduction of Ni 2+ to Ni(0) explain the second reduction stage on the TPR profile.

The evolution of the X-ray diffractograms of Fe50Ni50/SiO2 during reduction under H2 is presented as a function of temperature on Figure 5. Only the diffraction peaks of the phyllosilicate phase are visible up to 350°C. In accordance with XAS, three diffraction peaks characteristic of fcc metal nanoparticles ((111), ( 200) and (220) reflections) are detected above 350°C and their intensity increases when temperature increases up to 700°C. No peak related to bcc phases is observed (in particular the (200) reflection expected around 2q = 65.0°). The average size of the crystalline domains increases from 3 to 5 nm from 350 to 700°C. The diffraction peaks attributed to the phyllosilicate phase disappear above 500°C.

The unit cell parameter a is known to decrease linearly with increasing Ni content in bimetallic fcc Fe-Ni alloys [52], as the atomic radius of Fe (0.172 nm) is larger than that of Ni (0.162 nm). A linear Vegard correlation between a, measured at RT, and the Ni atomic content was drawn from the JCPDS files of reference Fe-Ni alloys (Fig. S10); an uncertainty of 1 x 10 -4 nm on the value of a leads to an uncertainty of 1 Ni at % in the composition of the diffraction domains. After correction from the temperature-induced shifts measured for monometallic Ni100/SiO2 in the same conditions (Fig. S11 andS12), the positions of the ( 111) and (200) reflections for Fe50Ni50/SiO2 give access to the values of a at RT after reduction at 500 and 700°C (0.3547 and 0.3557 nm, respectively) and, through the correlation established above, to an assessment of the composition of the crystalline core of the particles (Fe22Ni78 and Fe31Ni69, respectively). The latter value of a is confirmed after measurement back at RT under H2 (0.3555 nm, Fe29Ni71).

The 57 Fe Mössbauer spectrum of the sample reduced at 500°C (Fig. 1b, Table 1) can be interpreted in terms of three components. The isomer shift of the first two large quadrupole split components is typical of high-spin Fe(II) ions, in line with XAS results. The component with the larger quadrupole splitting is consistent with Fe(II) ions in silicates, since the hyperfine parameters correspond to their typical range [53]. The quadrupole splitting of the second component, however, is too small to correspond to a bulk silicate, and can be attributed to Fe(II) ions in extremely distorted octahedral sites, or to lower coordination numbers. The third contribution, covering 57% of Fe, consists of a singlet typical of Fe(0), with an isomer shift at d = 0.06 mm.s -1 slightly more positive than that of iron metal, as reported for Fe-Ni alloys (d = 0-0.06 mm.s -1 at 25°C, the value of d increasing with the Ni content [54][55][56]). Based on the overall composition determined by ICP (38 Fe at %, 62 Ni at %), the composition of the Fe-Ni particles at 500°C is evaluated as Fe22Ni78, in good agreement with that determined by in situ XRD. The absence of a sextet shows that the Fe-Ni nanoparticles are superparamagnetic at RT, thus exhibiting a size smaller than 10 nm [55,56].

The third reduction stage on the TPR profile can be explained by comparing the Mössbauer spectra recorded at 500 and 700°C. Reduction at 700°C (Fig. 1c andd, Table 1) produces a complete modification of the 57 Fe Mössbauer spectrum, which can now be fitted with only two spectral components, a singlet and a quadrupole doublet. Both components are often observed for reduced, superparamagnetic iron nanoparticles, and are usually attributed to bulk and surface iron atoms, respectively [57]. The singlet at d = 0.04 mm.s -1 corresponds to Fe(0) at the core of Fe-Ni nanoparticles. The decrease of d suggests an environment richer in Fe compared with the reduced state at 500°C [54]. Since the hyperfine parameters of the quadrupole doublet can be confused with those of Fe(III), the same sample was measured at -268°C after transfer in a helium-flow cryostat. The low-temperature spectrum is dominated by a magnetically split contribution, that can be fitted with at least two magnetic sextets, comprising 90% of total Fe. The isomer shift, once taken into account the second-order Doppler shift, and the magnetic field of these sextets are in the typical range of Fe(0) in Fe-Ni alloys [58]. The two hyperfine fields are close to those measured for fcc Fe-Ni alloys in which Ni or Fe is the major metal (29 and 33 T, respectively [59]). In addition to these two components, a minor quadrupole split component accounting for about 10% of the total resonance area is still present at the centre of the spectrum. At this stage, it can be interpreted as Fe(0) species not ordered magnetically, either located at the periphery of bimetallic particles, or in highly dispersed monometallic Fe nanoparticles.

Characterization of the metal nanoparticles by TEM and EDX after reduction at 500 and 700°C

After reduction at 700°C, the Fe50Ni50/SiO2 catalyst exhibited a specific surface area of 240 m 2 gcatalyst -1 , an apparently low value compared with pristine silica, but explained mainly by the high metal loading ((Fe+Ni) loading = 52 wt%) and by the deep transformation of silica during DPU. The pore volume was 0.6 cm 3 .g -1 .

The metal nanoparticles on Fe50Ni50/SiO2 were characterized by bright field TEM after reduction at 500 or 700°C, and exposure of the samples to ambient air (Fig. 6a andb). After reduction at 500°C, they appeared as evenly distributed over the siliceous matrix and exhibited an average diameter of 5.1 nm (standard deviation: 0.9 nm), with a high proportion of smaller particles (Fig. S13a). Agglomeration was not observed after reduction at 700°C, and the bimetallic nanoparticles were still evenly distributed on the silica substrate despite the high metal loading, with a separation distance of 5 nm on average. The size distribution histogram shows that the average size, 5.4 nm, has only slightly increased between 500 and 700°C (Fig. S13b). The standard deviation was 1.3 nm and the size distribution was more symmetrical, with a lower proportion of smaller particles.

The elemental distribution of Fe and Ni was examined by EDX mapping after reduction at 500°C (Fig. 6c). No metal was detected on bare zones of the support. The composition of 28 individual particles was measured, resulting in an average composition of Fe31Ni69, and a standard deviation of 8 at % (Tab. S3). Both elements were present in every particle and Fe was never the major metal. It was not possible to perform point scans along the diameter of particles, owing to their small size, or to evidence local metal segregation.

Moreover, the comparison with in situ measurements is not straightforward, as the distribution of the metals may be perturbed by metal oxidation upon exposure to air [17].

DPU applied to the Fe-Ni system thus yields bimetallic nanoparticles of controlled size, in the same range as for monometallic Ni/SiO2 catalysts [21], and resistant to sintering at high temperature. The reduction above 350°C of the nickel ions from the phyllosilicate phase drives the progressive destruction of the layered structure and the reduction of surrounding iron ions to the metallic state, to form Ni-rich Fe-containing fcc nanoparticles. At 500°C, about 40% Fe 2+ ions are still present in an oxidic environment, in close association with the metal nanoparticles. Their gradual reduction between 500 and 700°C results in the diffusion of more Fe atoms inside the Fe-Ni particles. The composition of the particles, consistent with the values presented in sections 3.1 and 3.2, pertains to a limited range, with a standard deviation of 8 at %. In the absence of monometallic and Fe-rich particles, both the detection of non-magnetically ordered Fe by MössS, and the different values of the magnetic hyperfine field, suggest that a gradient of Fe concentration exists in the bimetallic particles, as formerly postulated by Nagorny and Bubert [59]. Due to the late reduction of Fe, the outer shells surrounding the bimetallic Fe-Ni crystalline core would be enriched in Fe, in line with former interpretations from the literature based on spectroscopic data and on beta emission from radioactive Ni [14,[59][60][61].

Catalytic tests

Depending on the catalyst nature and reaction conditions, the hydrogenation of furfural yields a slate of products resulting from the hydrogenation of the aldehyde function Fe50Ni50/SiO2 and Ni100/SiO2 catalysts were pre-reduced at 700°C, and underwent a second stage of reduction at 400°C prior to reaction to counter the effect of transfer in air. A characteristic TEM bright field micrograph of Ni100/SiO2, and the size histogram associated to this catalyst, are presented on Fig. S14. The metal dispersion D (%) can be calculated for the two catalysts by applying 97.1/Ø (nm) [64] to each class of diameter Ø from the two histograms (Fig. 6b andS14), assuming that metal particles are hemispherical and that this calculation, usually applied to fcc Ni, can be extended to fcc Fe-Ni alloys. A value of D = 20% is obtained both for Fe50Ni50/SiO2 and Ni100/SiO2. The metal surface area exposed per gram of catalyst is evaluated as 70 m 2 gcat -1 , assuming that one Fe or Ni atom occupies 6.3 Å 2 [64].

In the given experimental conditions, both catalysts showed high activity in the hydrogenation of furfural at 150°C (conversion: 97-100%), whatever the starting amount of furfural in solution (Table 3). In contrast, yields were highly impacted by several factors. Monometallic Ni100/SiO2 was rather selective to THFFA and MTHF, the over-hydrogenated products, in good correlation with the literature in similar conditions of temperature and pressure [5,65,66]. The formation of MF and MTHF by hydrogenolysis was favoured when the initial amount of furfural was high. Etherification of the alcohols appeared prominent. As , ined very low carbonyl group rema the of hydrogenation the from FFA of yield he a result, t conversion be explained by its an which c to MF by hydrogenolysis.

Unlike monometallic Ni100/SiO2, the Fe50Ni50/SiO2 catalyst exhibited a high yield in FFA whatever the initial amount of furfural, and low yields of THFFA, MF and MTHF. The carbon balance was high, demonstrating that lower amounts of heavy products were formed than with the monometallic Ni system. The fraction of ether formed from FFA slightly increased with the amount of furfural, but remained minor except when the catalyst did not undergo a pre-reduction at 700°C and thus contained iron mainly in an ionic form. As has been reported in the literature, the hydrogenation of the carbonyl group reaction prevails on the hydrogenation of ring and hydrogenolysis reactions on Fe-Ni catalysts [4,5]. The role of Fe surface atoms in the suppression of Ni overhydrogenating and hydrogenolyzing properties is thus confirmed here.

Conclusions

The preparation of bimetallic Fe-Ni/SiO2 catalysts by deposition-precipitation with urea requires the use of a Fe(II) salt and the absence of oxidizing species such as nitrates or dissolved oxygen. A layered structure best described as an ill-crystallized Fe-containing Ni(II)

  Fe(II) to Fe(III) upon heating, as evidenced by the change of colour of the suspension from green to orange, and to the early, separate precipitation of well-crystallized goethite Fe III O(OH) at pH 4.8, from which large particles of metallic bcc Fe would later grow upon reduction (diffraction peaks at 2q = 44.8 and 65.0°, (110) and (200) reflections; size: 40 nm).

  ) the progressive growth of the asymmetrical doublet of bands at 1007 and 1083 cm -1 (υ(Si-O) stretching vibration in a 1:1 phyllosilicate); (2) the decrease of the weak asymmetrical υ(Si-OH) band originating from silica at 974 cm -1 ; (3) the shift and decrease of the intensity of the υ(Si-O) vibration band assigned to silica at ca. 800 cm -1 ; (4) the decrease of the d(OH) band at 640 cm -1 , characteristic of turbostratic M II (OH)2, between 2 and 4 h, and, (5) from 4 h, the increase of a broad band at 664 cm -1 assigned to a tetrahedral Si-O mode in the 1:1 phyllosilicate structure. The poor crystallinity of the phyllosilicate phase may explain the absence of detection of Si second-shell neighbours in the EXAFS fits.

(

  furfuryl alcohol, FFA), hydrogenation of the furan ring (tetrahydrofurfuryl alcohol, THFFA) and/or hydrogenolysis of the C-O bond (2-methylfuran, MF; 2-methyltetrahydrofuran, MTHF) (Scheme 1). In the present work, ethers resulting from reactions with the isopropanol solvent were also identified by GC/MS (2-(isopropoxymethyl)furan, iPrOMF, m/z 81, 97, 139; 2-(isopropoxymethyl)tetrahydrofuran, iPrOMTHF, m/z 71, 85, 101, 144). Etherification can proceed with FFA or THFFA as a reaction intermediate, or via the direct reductive reaction of furfural, and can be catalyzed by unreduced metal ions acting as Lewis acidic centers [62,63].

1: 1

 1 phyllosilicate is formed after 22 h of DPU, with a lower proportion of Fe compared with the composition of the DPU solution (32 Fe at %, 68 Ni at %, against 50 Fe at %, 50 Ni at %), due to a preferential reaction of Ni 2+ ions with silica. The oxidation of Fe 2+ to Fe 3+ upon drying and storage does not challenge the subsequent formation of bimetallic particles upon reduction. Reduction under H2 proceeds in several stages: (1) reduction of Fe 3+ ions to Fe 2+ above 275°C; (2) reduction of Ni 2+ to Ni(0) above 350°C, triggering the progressive reduction of Fe 2+ ions to the metallic state, and the insertion of a fraction of Fe atoms into the fcc structure of Ni-rich particles; (3) deeper reduction of Fe 2+ ions between 500 and 700°C, resulting both in the diffusion of Fe atoms in the Fe-Ni core of the fcc particles, and in the enrichment in Fe of their outer shells. After reduction at 700°C, the bimetallic particles are homogeneous in size (average size: 5.4 nm) despite the elevated reduction temperature, and exhibit a standard deviation of composition of 8 Fe at %. The enrichment in Fe of the surface of the Fe-Ni nanoparticles suppresses hydrogenolysis and overhydrogenation side-reactions observed for Ni/SiO2 in the hydrogenation of furfural, and provides a catalyst selective in furfuryl alcohol. Side-reactions of etherification with the isopropanol solvent may be linked to the presence of unreduced metal ions acting as Lewis acidic sites.

Acknowledgements

This work has benefited from the support of the CSC-Ecole Centrale de Lille PhD scholarship program (D. S.), of the REU exchange program from the National Science Foundation (C. P.), and of the French National Research Agency (ANR), through the NobleFreeCat project (ANR-17-CE07-0022). Chevreul Institute (FR 2638), Ministère de l'Enseignement Supérieur et de la Recherche, Région Hauts-de-France and FEDER are acknowledged for supporting and funding partially this work.


We have developed a compact pulsed-field THz magneto-spectrometer based on a THz molecular laser and heterodyne detection both at room-temperature. The recently developed continuous-wave THz laser uses mid-IR-pumped ammonia as active medium. The receiver is based on a subharmonic mixer pumped by a multiplication chain. A pulsed magnetic field up to 9T is supplied by discharging a capacitor in a small coil at room-temperature. We demonstrate here the use of this spectrometer by measuring the effective mass of electrons in an InAs/AlGaSb heterostructure at room-temperature.

I. INTRODUCTION

OST of THz magneto-spectrometers operate with water or cryogenically cooled coils, often in large scale facilities (https://emfl.eu). Without cooling only electromagnets are available but they provide low magnetic fields only up to 1 T. At 1 THz it signifies that only effective masses up to 0.028m0 are accessible. Here we implemented a table-top THz magneto-spectrometer that combines at roomtemperature (RT) a pulsed magnet up to 9 T during 10 ms, a stable THz laser source and a sensitive large bandwidth realtime THz detector.

II. EXPERIMENT

The compact pulsed-field magneto-spectrometer is based on a THz molecular gas laser and heterodyne detection (see Fig. 1). The laser is an optically-pumped terahertz laser (OPTL) based on mid-infrared optical pumping of a polar molecular gas with a quantum cascade laser (QCL). This new concept, developed by us, allows pumping molecules and transitions that are not reachable with CO2 pump lasers [START_REF] Pagies | Low-threshold terahertz molecular laser optically pumped by a quantum cascade laser[END_REF]. Here the laser operates on the (3,3) pure inversion line of NH3 at 1.073 THz. The output power is of the order of a few 100s of µW. As shown in Fig. 1, the experiment is performed in transmission through the sample placed in the coil by guiding the THz beam inside a stainless steel circular waveguide. The receiver is based on a GaAs Schottky diode subharmonic mixer (Virginia Diodes WR1.0SAX) and the local oscillator (LO) is generated by a microwave synthesizer and a multiplication chain. The intermediate frequency (IF) is amplified, filtered, rectified and again amplified. The instantaneous detection bandwidth is larger than 1 MHz (time resolution better than 1 µs). The sample sits at RT in a copper magnet coil which supplies a field up to 9 T in a bore of 13 mm during typically 10 ms. The sample is a modulation-doped heterostructure grown by MBE on a GaAs semi-insulating substrate. It consists of a 1 µm-thick GaSb buffer layer promoting strain relaxation via the introduction of an array of Lomer dislocations confined at the GaSb/GaAs interface, a bottom Al0.8Ga0.2Sb barrier, a 15 nm thick InAs channel, a top Al0.8Ga0.2Sb barrier with a Si d-doping plane and a Si-doped InAs cap layer. Hall effect measurements lead to mobility values of 33 800, 325 000 and 600 000 cm2 /V.s with corresponding electron sheet density of 1.8, 1.5 and 1.5×10 12 cm -2 at 300 K, 77 K and 2 K respectively.

III. RESULTS

Fig. 2 shows the magnetic field pulse and the measured THz transmitted signal through the sample versus time. Two dips due to the electron cyclotron absorption are measured as the field sweeps up and down at a value close to 1.7 T. We can deduce an electron mass value of (0.045±0.002)m0 close to the published values [START_REF] Larrabee | Application of terahertz quantum-cascade lasers to semiconductor cyclotron resonance[END_REF]. Magnetic fields up to 12 T and a variable temperature of the sample are also achievable by cooling the coil with liquid nitrogen. In the future we intend to exploit the advantages of the heterodyne down-conversion: full amplitude/phase signal measurement and a phase-locked loop to further increase the signal/ratio.

Terahertz pulsed-field magneto-spectrometer at room-temperature M Fig. 1. Pulsed-field magneto-spectrometer setup.

Fig. 2 .

 2 Fig. 2. Magnetic field and THz signal versus time (room temperature sample).

Institute of Electronics Microelectronics and Nanotechnology, UMR CNRS 8520, Lille University, France

Virginia Diodes Inc., Charlottesville, VA, USA

Laboratoire National des Champs Magnétiques Intenses, CNRS-UGA-UPS-INSA, Toulouse, France


D Bimberg

X Wallart

Novel Quantum Dot Based Memories with Many Days of Storage Time Last Steps towards the Holy Grail?

Keywords: quantum dots, QD-flash, tunneling, NVSRA I

The feasibility of the QD-Flash concept, its fast write and erase times, is demonstrated together with storage times of 4 days at room temperature. The storage time of holes in (InGa)Sb QDs embedded in a (AlGa)P matrix can be extended by growth modifications to 10 y. Tunneling structures were recently demonstrated to solve the trade-off conflict between storage time and erase time. A QD-NVSRAM is suggested to become the first commercial application.

INTRODUCTION

An important limitation of charge-based memories is the so called "voltage-time dilemma". The essence of this dilemma is to achieve 10 years in data retention by a barrier, that needs to be overcome during the write process. If typical barriers like pn-junctions in silicon or the conduction band offset of a dielectric like silicon dioxide to the silicon semiconductor are used, the consequence, is that a memory cell can either be nonvolatile or fast. In practical scenarios were dielectrics are used as barriers also the endurance will be limited for nonvolatile memories. This is the reason why there is such a huge performance gap between random access memories like DRAM or SRAM that are volatile and nonvolatile memories. QD Flash promises to solve this dilemma by using self-organized nanostructures in compound semiconductor systems, where the barrier height can be tuned much easier than in the systems used in silicon technology [1]. As a result, a memory cell that is both nonvolatile and possesses all attributes of a random-access memory (RAM) can be realized, meaning it will have a fast write speed and unlimited endurance. Such a memory cell would be ideal to complement or even replace DRAMs in the memory hierarchy and has been hailed more than 10 years ago as the "holy grail" of solid-state memories [2] II. QD-FLASH: PRINCIPLE OF FUNCTION QD-Flash is based on self-organized quantum dots (QDs) inserted in III-V compound semiconductor heterostructures. Using III-V materials allows for fine-tuning of the band offsets and barriers in the device, thus guaranteeing much greater flexibility than the Si/SiO2 material combination. Furthermore, QDs eliminate the need for high electric fields, since capture of charge carriers occurs thermally in a fast and efficient fashion. In a recent review [START_REF] Bimberg | QD-Flash: Towards non-volatoility in quantum dot based storage[END_REF], we presented the concept of the QD-Flash memory alongside with results obtained until 2016. The write time in QD-Flash is limited only by thermal capture into the QDs, which occurs in a few picoseconds at room temperature. The erase time depends on the localization energy of the QDs and on the external voltage applied. Appropriate tuning of these parameters can lower the erase time to tens of nanoseconds. The storage time depends on the localization energy and capture cross-section of the QDs. These parameters are determined by the choice of materials used to manufacture the heterostructure and the size of the QDs. Changing the heterostructures from GaAs-based ones to GaP-based ones, allowing growth on Si, and from type I to type II arrangement of the bands, thus increasing the confinement has led recently to an increase in the storage time from nanoseconds to 4 days at room temperature [START_REF] Bonato | [END_REF]. QDs in III-V heterostructures form a triangular barrier, the shape of which depends on the external voltage. As a consequence, its width can be changed at will, thus speeding up or slowing down the erase process. The triangular barrier also allows for much lower electric fields during the erase process, thus increasing the durability of the device. Finally, QDs in III-V semiconductors can be manufactured to store holes instead of electrons employing the type II band alignment. Since holes have a larger effective mass than electrons their storage time is much longer for the same localization energy.

III. QD-FLASH CELLs

A QD flash cell is a modulation-doped field effect transistor (MODFET), in which a layer of QDs has been embedded between the 2-dimensional hole gas (2DHG) and the gate. Fully functioning protypes have been manufactured already 8 years ago using InAs QDs embedded in GaAs or Al0.9Ga0.1As [6], thus demonstrating the feasibility of the QD-Flash concept.

The charge state of the QDs is controlled by the gate voltage and read out is done using the 2DHG. The structure of QD-Flash is sketched in Fig. 1 Fig. 1: Sketch of a QD-Flash

The logic state "0" is realized in a QD-Flash when the QDs are not occupied by holes. Conversely, the logic state "1" is realized when holes are localized in the QDs. The band structure configuration of QD-Flash during the basic operations (store, write and erase) is depicted in Fig. 2. In the storage configuration (Fig. 2a), if holes are confined (logic state "1") they cannot escape the QDs because of their huge localization energy, which thus represents the emission barrier. If the QDs are not charged (logic state "0"), the holes present in the valence band are prevented from entering the QDs by the triangular barrier created by the band-bending of the Schottky gate contact, which thus represents the capture barrier. The storage time of QD-Flash is determined by the capture and emission rates of the QDs [6]. In this system, capture and emission are dominated by thermally assisted tunneling. Consequently, the storage time depends on the barrier height, temperature and applied voltage. In order to write a logic "1", a write pulse is applied to the gate. The pulse is chosen to bring the system into flat band configuration (Fig. 2b). The capture barrier is removed, and holes are thermally captured into the QDs. Thermal capture is a very fast process, with time constants of the order of the picoseconds at room temperature [6]. Conversely, erasing (i.e. writing the logic "0") is performed by applying an opposite bias to the gate. The erase pulse bends the valence band even further, thus making the triangular barrier thinner (Fig. 2c). This in turn allows the charge carrier to tunnel out of the QDs, effectively erasing the stored data. Erase times much below 1 ns have been predicted with an appropriate choice of applied electric field and employing tunnel junctions [7].

Charge carriers are captured to the QDs from the 2DHG, which is therefore depleted whenever the QDs are charged.

QD

IV. MATERIAL SCIENCE CHALLENGES

Much longer storage time are predicted for the GaSb/GaP system as compared to the InAs/GaAs QDs, due to larger hole localization energy in type II structures and the larger band gap discontinuity [1]. We have therefore investigated for the first time the quite challenging growth of GaSb QDs on (001) GaP substrates by means of Molecular Beam Epitaxy (MBE) [8]. QD growth was performed by deposition of 1.2 to 2 monolayers (ML) GaSb on GaP at a growth temperature ranging from 410 to 485°C and growth rates of 0.1 to 0.2 monolayers per second (ML/s). QD diameter, height and density have been systematically determined by atomic force microscopy and their dependence on the growth conditions has been interpreted in terms of scaling theory. Typical measured values lead to densities in the 10 10 /cm 2 range, diameter around a few tenths of nm and height between 3 and 6 nm, which make these QDs valuable candidates for the targeted memory applications (Figure 3). Similar QD layers as that depicted in figure 3 have been capped starting with a 7 nm thick GaP layer grown at 430°C at 0.2 ML/s. Then the growth has been interrupted for 4 minutes to increase the growth rate to 1 ML/s for the rest of the 900 nm thick capping layer with a progressive increase of the growth temperature from 430°C to 600°C. TEM analysis demonstrates that, upon capping, the QD layer morphology is strongly modified by a strong P/Sb exchange reaction leading to the formation of a 10 nm thick intermixed GaP/GaSb layers. In spite of the large mismatch between GaSb and GaP (11.8%), the QD layers are coherently grown on the substrates without any dislocations. Finally QD layers have been embedded in GaP and n + p-diodes have been fabricated. The localization energy, capture cross-section, and storage time of holes in the ground state of the quantum dots have been determined via deep-level transient spectroscopy. Their localization energy of 1.18(+/-0.01) eV agrees with the theoretical prediction of 1.4 eV once the observed interdiffusion is considered. As storage time of holes of 3.9 (+/-0.3) days at room temperature is determined, marking an improvement of 3 orders of magnitude from previous record numbers, see Fig. 4. Replacing the GaP barriers either by AlP or InAlP using either MBE or MOCVD will shift the hole storage time much beyond the 10y limit defined for Flash memories. Hole localization energies of 1.9.eV and 2.0 eV are expected for the material combinations GaSb/AlP and In0.5Ga0.5Sb/AlP, respectively, leading to hole lifetimes 10 orders of magnitude larger than presently observed [START_REF] Bimberg | QD-Flash: Towards non-volatoility in quantum dot based storage[END_REF].

V. OUTLOOK

Next material science work will focus on -demonstrating the beyond 10 y lifetime for AlP/ or InAlP/InGaSb QDs, -combining the tunnel junctions with the AlP/GaSb -fabricate a GaP-based fully operational flash cell similar to Fig. 1 and -employ previously developed concepts [8] for single q-bit emitters for localising precisely QDs, targeting single QD memories -start circuit design work for integration. A memory is more than 1 cell.

Thus in a few years' time Si-compatible quantum dot memories could catch the NVSRAM market. In the long run it could even become the non-volatile alternative for today´s DRAM technology the industry has been seeking for more than two decades now. .

  [START_REF] Mikolajick | Proceedings of NVMTS[END_REF]. THIS "HOLY GRAIL" WOULD OFFER UNIQUE NEW FUNCTIONALITIES ENSURING FUTURE PROGRESS OF MEMORY DEVELOPMENT AFTER THE END OF MOORE'S LAW.

Fig. 2 :

 2 Fig.2: Schematic representation of the valence band during memory operations in QD-Flash: (a) store, (b) write (writing "1"), and (c) erase (writing "0").

  charge carriers interfere with conduction in the 2DHG through Coulomb scattering. The combination of these two effects causes the conductivity of the 2DHG to fall when the QDs are occupied and to rise when they are unoccupied. Therefore, a simple measurement of the sourcedrain resistance serves as read-out of the storage state of the device.

Fig. 3 :

 3 Fig.3: 1µmx1µm AFM image of GaSb QDs after deposition of 1.2 ML GaSb on a (001) GaP surface at 450°C at a growth rate of 0.1 ML/s and with a Sb/Ga ratio of 2.5. The resulting density is around 2x10 10 /cm 2 .

Fig. 4 :

 4 Fig. 4: Carrier storage time versus hole localization energy as a function of capture cross section.

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Marisa Rico-Santacruz

Patricia García-Muñoz
email: garciamunoz@unistra.fr

Valérie Keller
email: nkeller@unistra.fr

Nelly Batail

Charlotte Pham

Didier Robert

Nicolas Keller

Valerie Keller

Alveolar TiO2-β-SiC photocatalytic composite foams with tunable properties for water treatment

Keywords: TiO2-β-SiC composite foam, Shape Memory Synthesis replica method, photocatalysis, adsorption properties, water treatment

INTRODUCTION

Wastewater treatment is still a concerning issue in most parts of the world, where the removal of inorganic and organic contaminants from wastewater remains nowadays a major concern, especially when some of the most hazardous and refractory targets such as phytosanitary or plant protection products are concerned. Thus extensive research has been driven by the need to establish tertiary treatments in the current wastewater treatment plants, while keeping in mind the necessity to improve the process sustainability when facing this challenge [1].

Photocatalysis is one of the advanced oxidation processes widely investigated in the frame of a sustainable development policy for mineralizing water contaminants or at least for degrading them into readily biodegradable compounds [2,3]. Powdery photocatalysts are usually used as nanoparticle suspensions which take advantage of both an efficient light distribution and a maximum exposed surface for optimizing the degradation activity.

Nevertheless, the necessity of implementing a costly, technologically-limited and timeconsuming photocatalyst recovery step such as nano-filtration for separating it from the treated water, as well as safety issues related to the handling of nanopowders, make the process less viable on large scales for a real application and with a less secure environment for handling and recycling.

Therefore approaches focussing on the immobilization of photocatalysts and especially on the design of macroscopic structures are of high interest for solving those issues. The need to minimise inherent drawbacks resulting from the catalyst immobilization led to investigate cellular monolithic solids as promising alternative to packed-bed photocatalytic materials.

Among them, multichannel materials such as honeycomb-shaped or square-channel structures enable operating at low pressure drop while increasing the reactor surface/volume ratio; however, they suffer from poor internal light irradiation, due to the complexity to provide light to the photocatalyst and to the reactor core [4]. Open-cell alveolar solid foams have recently received a strong interest in water treatment for taking advantage of an active static mixer effect inside the reactor while benefitting from better light transmission within the material [5][6][7][8][9][10]. In comparison to polymer, metallic and other ceramic foams such as eg. cordierite, alumina or zirconia [11][12][13], medium surface area alveolar β-SiC foams gained a strong interest thanks to a high density of oxygenated surface groups for anchoring large amounts of TiO2, as well as to a high thermal and chemical stability that allow severe reaction conditions to be applied on the foam and possibly deactivated photocataysts to be regenerated thermally or by soda washing.

Previous results evidenced the interest of using open-cell self-bonded TiO2/β-SiC alveolar foam structured photocatalysts for water treatment, prepared by immobilizing the TiO2 photocatalyst onto the medium surface area pre-shaped β-SiC alveolar foam support [5,14,15]. However, the implementation of a post-synthesis immobilization technique for stabilizing TiO2 onto the foam support opens the door to detrimental stability problems that could provoke a release of photocatalyst to the water. Indeed, the resistance of the supportcatalyst junction towards strains derived from particle to particle and particle-fluid mechanical interactions in the reactor environment, remains an essential quality criteria for a good support, in order to avoid detachment of photocatalyst particles from the support.

The aim of this work is the design of TiO2-β-SiC photocatalytic composite foams that can gather within a ready-to-use photocatalytic media, the TiO2 active phase and the β-SiC alveolar foam structure. As a result, there is no need to implement a post-synthesis immobilization process on the foam, so that such a TiO2 coating-free composite foam would also benefit from a better stability in water. They were elaborated through a synthesis approach derived from the Shape Memory Synthesis (SMS) replica method developed for preparing pure β-SiC alveolar foams. This replica method was already used for preparing TiO2-β-SiC composite powders, through a double carburization process involving the carboreduction of a precursor TiO2 into TiC together with the formation of β-SiC, with a subsequent selective oxidation of TiC into TiO2 at low temperature by taking advantage of the β-SiC resistance to oxidation [16]. In addition, we aimed at controlling the adsorption properties of the TiO2-β-SiC composite foams towards the pollutant in water by applying an appropriate final calcination treatment. Diuron (C9H10Cl2N2O, 3-(3,4-dichlorophenyl)-1,1-dimethylurea) has been taken as model substrate to degrade for evaluating the liquid phase photocatalytic activity of β-SiC-based composite foams under simulated solar light, since Diuron is one of the most commonly used contact herbicides from the family of substituted phenylureas, whereas photocatalysis has been reported to be efficient for degrading Diuron and mechanistic pathways have been proposed [2,17,18].

EXPERIMENTAL PART

Synthesis of TiO2-β-SiC photocatalytic composite foams by SMS replica method

Alveolar open-cell TiO2-β-SiC composite foams with medium surface area were developped at the SICAT company (Strasbourg, France) according to the Shape Memory Synthesis (SMS) method (Figure 1). The SMS replica method consists in the transformation of a pre-shaped polyurethane foam into the corresponding β-SiC carbide at 1360 °C via a 4 steps synthesis process [19]. First, impregnation/infiltration of a pre-shaped precursor polyurethane foam with a phenolic resin containing micronized (<20 µm) metallic silicon and carbon black powder was performed. Here, the classical SMS replica synthesis method implemented for obtaining bare β-SiC foams has been modified by incorporating controlled amounts of TiO2 powder (Crenox 8602, purity 91.9%, SBET > 250 m²/g, particle size= 50-1 µm and dp=5-10 nm) inside the slurry used for the infiltration/impregnation step. Then, the green body was dried overnight at 120 °C, followed by a reactive step under Ar atmosphere at 1360 °C. During the thermal treatment, the polyurethane was pyrolized giving almost no carbon yield, whereas the high carbon yield resin was pyrolized into amorphous carbon binding the carbon black, the silicon and the TiO2 powders together. Above 1000 °C, the carbon skeleton was subsequently attacked by SiO vapors formed by reaction between Si and the residual traces of oxygen, whereas CO was allowed to carburize the metallic silicon, yielding to the formation of β-SiC. The TiC was formed during SiC synthesis by reaction between the TiO2 and the carbon species. Finally, a thermal treatment in air was performed at temperatures between 400 °C and 700 °C for removing the residual unreacted carbon by combustion and for producing the final TiO2-β-SiC composite foams. Two materials with theoretical 10 and 20 wt. % of TiO2 in respect of the β-SiC matrix were prepared.

A reference β-SiC foam supported TiO2 catalyst has been prepared via the classical deposition method consisting in dipping the foam 35 times into a 15 g/L TiO2-P25 (Evonik) ethanolic suspension, with intermediate rinsing steps with ethanol solution, drying at 100 °C and final calcination treatment at 380°C. Prior deposition, the bare β-SiC foam was decarboned for 2 h at 700 °C in air for removing residual unreacted carbon species by combustion.

Evaluation of the photocatalytic efficiency

The experiments were carried out in an artificial weathering in fast-action instrument Suntest XLS+ (Atlas Material Testing Technology BV, Gelnhausen, Germany) equipped with a Xenon arc lamp of 1700 W adjustable power. The equipment has a Solar ID65 filter to limit the UV radiation at 320 nm, simulating solar exposition according to ICH Q1B guidelines.

The runs were performed adjusting the lamp power to 250 W m -2 with a simulated solar emission within 300-800 nm, which corresponds to an average solar radiation in a summer day in southern Europe. The reaction volume was 100 mL and the starting concentration of diuron (3-(3,4-dichlorophenyl)-1,1-dimethyl-urea) 10 mg L -1 . At each time interval, 1 mL of solution has been sampled and then filtered through 0.20 µm porosity filter to remove the photocatalyst powder if any, before the concentration of Diuron has been determined by UV-Visible spectrophotometry (Cary 100 scan Varian) by monitoring the disappearance of the main absorption peak at λ=248 nm, and Total Organic Carbon measurements were performed using a Shimadzu TOC-L analyzer to determine the organic carbon load. Short-chain acids and chloride anion were analyzed by ion chromatography (Metrohm 790 IC) using a conductivity detector.

Characterisation techniques

The titania content in the catalysts were determined by chemical analysis after a microwave-assisted acidic dissolution in aqua regia at 185°C under autogenic pressure.

Inductively coupled plasma optical emission spectroscopy (ICP-OES) was carried out on an

Optima 7000DV spectrometer (Perkin Elmer) at the Analysis Platform of IPHC-Strasbourg, France.

X-ray diffraction (XRD) patterns were recorded on a D8 Advance Bruker diffractometer in a θ/θ mode and using the Kα1 radiation of a Cu anticathode (λ = 1.5406 Å). The anatase and rutile mean crystallite sizes were calculated from the Scherrer equation applied with the usual assumption of spherical crystallites.

The surface area measurements have been carried out on a Micrometrics Tristar 3000 using N2 as adsorbent at -196 °C with a prior outgassing at 200 °C overnight to desorb the impurities or moisture. The Brunauer-Emmett-Teller (BET) specific surface area and the total pore volume have been calculated from the N2 adsorption isotherm.

Scanning electron microscopy (SEM) was performed in secondary electron mode on a JEOL JSM-6700 F FEG microscope.

Thermogravimetric analysis (TGA) was carried out with a 20% (v/v) O2/N2 mixture at a flow rate of 40 mL/min at a heating rate of 10 °C/min from 25°C to 900°C using a Q 5000TA thermoanalyzer.

RESULTS AND DISCUSSION

Characterization of the TiO2-β-SiC composite foams

Figure 2 shows optical images of the pre-shaped alveolar polyurethane precursor, as well as of the bare β-SiC foam and of the corresponding TiO2-β-SiC composite foams obtained through the SMS replica method. It shows that both the alveolar open-cell structure and the cell size of the foam were maintained during the SMS process going from the pre-shaped polyurethane foam to the TiO2-containing β-SiC composite foam. It is worth noting that despite the introduction of the TiO2 powder into the precursor slurry used for infiltrating the pre-shaped polyurethane foam, the SMS replica synthesis method still reflected the fact that the original macrostructural features of the alveolar foams were retained after synthesis. Thus, TiO2-β-SiC composite foams can be manufactured with adjustable shapes suitable for being adapted to the reactor geometry, and they have been prepared as 5 cm (d.) x 2 cm (h.) disks.

The SEM images shown in Figure 2 evidences the strong heterogeneity of the TiO2-β-SiC composite foam surface compared to the surface of the bare β-SiC foam.

The TiO2-β-SiC composite foams and the reference β-SiC foam exhibited globally similar macroscopic features in terms of characteristic parameters (cell size at 5200 ± 300 µm, window size at 2500 ± 150 µm and strut diameter at 550 ± 75 µm), with as a result a similar open porosity of about 0.91-0.94 ± 0.05, as well as a similar geometrical surface area within the 540-620 m -1 range. One can therefore assume that the TiO2-β-SiC composites show a similar light transmission profile through the foams than the reference foam [5].

The main physico-chemical properties of the TiO2-β-SiC composite foams are shown in Table 1. Figure 3 shows the XRD patterns obtained on the composite foams calcined at temperatures ranging from 400 °C to 700 °C. In addition to the diffraction lines attributed to the β-SiC polymorph in fcc structure (JCPDS card , that are similar to those of the reference β-SiC foam, diffraction peaks corresponding to anatase TiO2 (JCPDS card were observed for all the composite foams, with the characteristic most intense peak at 25.3° corresponding to the diffraction of (101) planes. Formation of rutile TiO2 phase was also observed at 700°C in the composite foams with 10 wt. % of titania, and already at 500 °C in those containing 20 wt. % of titania, with the characteristic most intense peak at 25.1° corresponding to the diffraction of (101) planes (JCPDS card 21-1276). The intensity of the diffraction peaks corresponding to both titania phases increased with increasing the calcination temperature and the titania content in the precursor slurry, as it was also confirmed by Raman analysis (not shown). However, it is worth noting that the rutile content remained limited at about 10%. This contrasted with similar works performed on TiO2-β-SiC composite powders, for which TiO2 rutile contents up to 70% were observed for a calcination temperature of 600 °C. This was attributed on one hand to the continuous aspect of the β-SiCbased foams that is known to improve heat transfer properties through the whole foam, while intrinsically β-SiC can act as a thermal regulator to disperse the heat issued from the TiC oxidation thanks to its high thermal conductivity. For the same reasons, the TiO2 mean crystallite size was not strongly impacted by the calcination temperature increase, with a mean size around 10-12 nm.

The complete oxidation of TiC into TiO2 at 400°C with no formation of suboxides was in agreement with the works of Shimada [20] and of Sheng et al. [21], that evidenced that the exothermic and one-dimensional diffusion process for the oxidation of TiC by oxygen already occurred at 350°C. Further, similar works performed on TiO2-β-SiC composite powders obtained via this replica method showed that the chemical environment of the Ti 4+ (Ti-O) surface species in TiO2 network was not influenced by the calcination temperature within the studied range [16].

In the case of the highly loaded composite foam (ie. with 20 % of TiO2), the formation of crystallized silica was observed, with the appearance of an additional peak at 22.0° characteristic of the diffraction of (101) planes of silica (JCPDS card 65-0466). Whereas the influence of the silicon oxidation kinetics cannot be ruled out, this can result from an overall deficit in reductive carbon compared to the silicon species within the precursors. Also, considering the stoichiometry difference between TiO2+3CàTiC+2CO and TiO2+Si+CàSiO2+TiC overall reactions, SiO2 could result from the existence in the infiltrated precursor foams of both carbon-poor and carbon-enriched areas, due to unhomogeneity either of the foam coating or within the infiltration slurry.

The foams with a theoretical content of 10 wt. % and 20 wt. % displayed a real TiO2 content of 16 ± 2 wt. % and 24 ± 4 wt. %, obtained by ICP-OES. Considering the analysis depth in SEM, EDS analysis reported in Fig. 2 evidenced the presence of TiO2 inside the external layer of the composite foams, measured at 15 and 26 at. % of Ti for the F10-700 and F20-700 foams, with 46% and 29% at. % of Si, respectively. This was associated to the modification of the foam surface, observed by SEM, and this evidenced that TiO2 was not homogeneously dispersed within the β-SiC matrix bulk. We proposed that migration of titanium from the bulk of the foam to its surface probably occurred during the high temperature synthesis process.

No significant change was observed in terms of N2 sorption isotherm and hysteresis profiles compared to the reference β-SiC foam, meaning that the main shape of pores was not affected, and the composite foams remained made of aggregates/agglomerates of particles forming slit shaped pores with non-uniform size and/or shape particles (Fig. SI1) [22]. By contrast, the TiO2-β-SiC composite foams displayed a higher surface area within the 34-65 m 2 /g range when compared to that of the reference β-SiC foam at 25 m 2 /g. This was assigned on one hand to the residual presence of non-reacted high surface area TiO2 nanoparticles or to some carbon excess that can bring micropores, whereas one cannot rule out the possible formation of voids and cracks during the TiC to TiO2 oxidation, as reported in the case of TiO2-β-SiC composite powder systems and mixed Zr-β-SiC materials [23,24]. The decrease in surface area observed for both TiO2 contents when increasing the calcination temperature from 400 °C to 700 °C could result from the combustion of residual unreacted carbon from the β-SiC foam synthesis.

TGA experiments were performed for determining the amount of residual unreacted carbon in the TiO2-β-SiC foams as a function of the calcination temperature (Figure 4). The profiles showed that a final thermal treatment in air at 700°C was necessary for removing completely the residual unreacted carbon by combustion, while the TiO2-β-SiC composite foams calcined at 400°C and 500°C contained about 7-8 wt.% and 5 wt.% of residual unreacted carbon, respectively, whatever the TiO2 content within the composite (Table 1)

Dark adsorption experiments

Figure 4b depicts the evolution of the diuron concentration during the dark adsorption period on the TiO2-β-SiC composite foams calcined at temperatures ranging from 400 °C to 700 °C. It shows that the final calcination temperature for obtaining TiO2 from TiC strongly influenced the adsorption behavior of the foams towards the diuron molecule. On both composite foams calcined at 400 °C, a complete adsorption of diuron was observed within 60 min and 150 min of dark adsorption period, while 85-90% diuron adsorption was achieved on composite foams calcined at 500°C. By contrast, calcining the composite at 700 °C suppressed the strong adsorption on the foams whatever their TiO2 content, the diuron concentration remaining constant for 3 h. This adsorption behavior of the composite foams was in agreement with the TGA profiles, the highest the amount of residual unreacted carbon, the highest the adsorption profile of the foams. The highest the amount of unreacted carbon in the composite foam, the highest the surface area and the highest the adsorption profile. Therefore the amount of residual un-combusted carbon issued from the pre-shaped PU precursor foam and from the pyrolized resin, could be tuned from 0 to 8 wt.% by the final calcination temperature (400-700 °C range), to tune the adsorption properties of the TiO2-β-SiC composite foams.

Photocatalytic runs

Preliminary tests performed on the TiO2-β-SiC composite foams calcined at 700 °C did not show any significant activity in the degradation of diuron, only a very small kinetic rate constant of 0,27x10 -3 min -1 being observed for the composite with 20 wt.% of TiO2 (Fig. 5a,b). By contrast, a strong TOC release was observed for both composite foams. This TOC profile was attributed to the release and subsequent mineralization of carbon-containing residues issued from the foam synthesis process remaining adsorbed at the catalyst surface and potentially blocking the TiO2 active sites, so that no significant conversion of diuron was observed. Figure 5c shows that those carboned species can be removed by implementing a photocatalytic pre-cleaning step of the composite foams in Milli-Q water for 5 h under UVA at 60 mW/cm 2 prior the evaluation of the degradation activity.

After a first dark adsorption period with a similar adsorption profile than that observed on uncleaned foams (not shown), the diuron degradation was performed under solar irradiation on the TiO2-β-SiC composite foams (Fig. 5 b,c). Whatever the adsorption properties of the composite foam, ie. whatever the remaining diuron concentration to be degraded after the adsorption step -10 mg/L or 1-1.5 mg/L for foams calcined at 700°C and 500°C respectively-, complete removal of the diuron was achieved under irradiation. The higher the TiO2 content, the shorter the necessary reaction time. Complete photocatalytic degradation of diuron was obtained on the composite foams calcined at 700 °C within 120 min and 90 min of reaction for a content of 10 wt.% and 20 wt.% respectively, corresponding to a kinetic rate constant of 48x10 -3 min -1 and 53x10 -3 min -1 .

It was worth noting that the TiO2-β-SiC composite foams calcined at 700°C outperformed the reference TiO2/β-SiC foams prepared via impregnation and stabilization of TiO2 on conventionnal β-SiC foams in terms of diuron degradation, with an apparent kinetic constant of 11.2x10 -3 min -1 . This evidenced that most of the TiO2 phase formed during the synthesis of the TiO2-β-SiC composite foams were located in the external layer of the foam rather than within the bulk of the foam matrix -and in a higher amount in the case of the F20-700 foam-, so that TiO2 was available for photonic activation. This was in agreement with the EDS-SEM results. Although we could not rule out that the presence of residual unreacted carbon might have an influence on the charge transfer property, we proposed that this high apparent kinetic constant could result from the establishment of TiO2-β-SiC heterojunctions at the nanometric scale, with the coupling of the p-type semiconductive β-SiC crystallites of the foam support with the n-type TiO2 crystallites [25][26][27][28]. Currently, the semiconductor coupling is a wellinvestigated strategy for reducing the recombination rate of charge carriers [29,30]. The intrinsic electric field induced at the heterojunction interface with matching band potentials enhances the transportation of photogenerated charge carriers between the composite material and promotes the separation of the photoexcited electron-hole pairs in both semi-conductors [31]. The positive role of heterojunctions between SiC and TiO2 has been proposed in few works on nano-TiO2 coated micrometer-sized SiC particles to account for a better photocatalytic efficiency [32,33]. While such a heterojunction is unlikely to allow boosting the efficiency in conventionnal β-SiC foam supported TiO2 photocatalysts due to the low TiO2/β-SiC heterojunction density, especially when a thick micrometric TiO2 coating is applied, we can hypothesize that the nanometric size of both TiO2 and β-SiC crystallites in intimate contact inside the nanocomposites is a key-parameter for increasing the heterojunction density in the material, and consequently the charge carrier separation.

As far as composite foams with enhanced adsorption properties were concerned (after calcination at 500°C), the higher the adsorption capacity, the shorter the reaction time necessary for removing the diuron from the water. However, the kinetic rate constants were not compared to those achieved on the composite foams calcined at 700°C, due to far different initial concentrations of diuron to be degraded.

By contrast, the TiO2-β-SiC composite foams did not outperform the conventional TiO2/β-SiC foam in terms of diuron mineralization, with a kinetic rate constant for TOC removal of 0.58x10 -3 mg/Lmin -1 and 1.0x10 -3 mg/Lmin -1 being observed for the F10-700 and F20-700 composite foams, respectively, to be compared to 2.0x10 -3 mg/Lmin -1 for the reference foam (Fig. 5b). The smaller mineralization rates obtained on TiO2-β-SiC composite foams while they outperformed the reference in terms of diuron degradation were attributed to a large amount of generated intermediates that may be adsorbed on the catalyst surface provoking its partial deactivation. A similar opposite behaviour was also reported by Ohtani and co-workers, who have shown that hydrothermally-synthesized TiO2 nanoparticles displayed smaller TOC removal rates than TiO2-P25 although a higher reactivity toward the main pollutant was achieved. They assigned this behavior to the formation of carbonaceous deposits at the catalyst surface [34].

The degradation rates of the composites indicated the scarce mineralization of the effluents reached after 4 h. Those values followed a rise trend at higher temperatures and when higher amounts of surface titania were contained in the catalysts. But, the TOC conversion reached did not exceed the 30% for the best case (F20-700).

The oxidation process was further evidenced on the composite foams by the detection of the short-chain organic acids (oxalic acid as the most representative, formic and acetic as trace amounts) and of chloride ions released to the solution (Table 2). The inorganic chloride concentration detected in the solution was quite low (between 5 and 10 %), while the chloride atomic balance was almost closed in most of the cases by Cl species found to adsorbed on the composite foam surface as chloride ion or organochlorides as evidenced by inorganic chromatography analysis with a previous leaching step .

The identification of many metabolites of diuron by Solis et al. [35] allowed a multipathway degradation route to be proposed, in which the first step consists in the attack of the alkyl function or in the substitution of chlorine from the aromatic ring. After several hydroxylations, the process led to the formation of acetic, formic and oxalic acid as last reaction intermediates. In our case, the detection of chorine and of oxalic acid at the end of the reaction fully supported a similar oxidation mechanism on the TiO2-β-SiC composite foams. This route is in good agreement with the earlier well-known works of Malato et al. on the photocatalytic degradation of diuron [2].

Previous works also showed that the diuron degradation route was not affected by the association of TiO2 to fibers and silica binder, or to Pt nanoparticles [17,36]. Further, using the photocatalytic degradation of phenol, 4-chlorophenol and 2,4-dichlorophenoxyacetic acid in water as model reactions, Matos et al. evidenced that the reaction mechanism was not affected by the association of TiO2 with carboned adsorbents, similar reaction intermediates being observed when compared to the neat TiO2 (even if the product distribution might differ in some cases) [37,38]. Therefore we proposed that the diuron photodegardation mechanism was similar on the TiO2-β-SiC composite foams than on the TiO2 powders and on the β-SiC foams supported TiO2 photocatalysts.

Finally, the stability of the composite foam materials was evaluated by performing several sequential runs using the most photoactive photocatalytic foam (F20-700). A fairly good reusability of the catalyst was evidenced by the total diuron elimination, with apparent kinetic rate constants for diuron degradation calculated in average as 52x10 -3 ± 48x10 -3 min -1 (Fig. 5d). It must be noted that no release of titanium ions or titanium oxide nanoparticles to the solution was observed while performing reusability tests, as determined by ICP-OES analysis of the solution after treatment.

Conclusion

Alveolar TiO2-β-SiC photocatalytic composite foams incorporating different titania contents (16 and 24 wt.%) have been synthesized through a SMS replica method in which TiO2 powder has been mixed to the infiltration slurry used for transforming the pre-shaped polymer foam into its corresponding carbide, with a final selective oxidation of the TiC carbide into small TiO2 crystallites. They are gathering within a ready-to-use material the macroscopic alveolar structure provided by the β-SiC foam and the active TiO2 photocatalyst exposed in the external layer of the foam rather than dispersed within the matrix, so that no postsynthesis deposition or synthesis of the TiO2 active phase onto the foam support is necessary.

The adsorption behaviour of the TiO2-β-SiC composite foams towards the diuron pollutant in water was tuned by monitoring the amount of residual carbon in the material, obtained by performing a controlled final calcination step between 400°C and 700°C. The adsorption profile of the composites was as high as the calcination temperature was low, so that the composite foams can benefit from both adsorption and photocatalytic properties, whereas a calcination at 700°C drastically diminished the diuron adsorption. The photocatalytic activity was evidenced by the complete diuron removal from water and the reduction of TOC, together with the detection of short-chain acidic organic intermediates and of chloride anion.

  The TiO2-β-SiC composite foams calcined at 700°C outperformed a reference β-SiC supported foam TiO2 catalyst in terms of diuron degradation, while the reference foam remained more efficient in terms of mineralization rate. The stability of the composites was evaluated by performing several reuses of the photocatalytic foam. development (EFRE/FEDER) is acknowledged for partially supporting the work, in the frame of the project PHOTOPUR, within the Program Interreg-V Rhin Supérieur and the Sciences Offensive Call.

Figure 1 .Figure 2 .

 12 Figure 1. Schematic diagram of the SMS replica synthesis of the TiO2-β-SiC composite foams.

Figure 3 .

 3 Figure 3. XRD of the TiO2-β-SiC composite foams containing theoretical amount of

Figure 4 .Figure 5 .

 45 Figure 4. (a) TGA profiles of the TiO2-β-SiC composite foams. The small weight gain

Table 1 .

 1 Main physico-chemical properties of the TiO2-β-SiC composite foams. data averaged on three composite foams obtained by ICP-OES b the residual unreacted carbon content was determined by TGA analysis by evaluating the weight loss observed between 500°C and 800°C. c measured by water displacement. d derived from the model developed by Edouard and coll. using the open porosity and the mean strut diameter[39,40].

	β-SiC	F10-400	F20-400	F10-500	F20-500	F10-700	F20-700

a

Table 2 .

 2 Cl -ion release, oxalic acid concentration and TOC conversion observed after 4 h of reaction with the TiO2-β-SiC composite foams. remaining adsorbed after test in the composite foam (as organochloride or Cl -ion) has been determined by ionic chromatography, and expressed in respect of the initial amount of Cl in the diuron

	Catalyst	% Cl -release a	% Cladsorbed b	[Oxalic acid]	[Acetic acid]	[Formic acid]	XTOC (%) after 4h
				(mg/L)	(mg/L)	(mg/L)
	F10-400	5	90	2.4	1.8	0.7	8
	F20-400	10	40	3.0	-	-	12
	F10-500	7	83	1.8	-	-	16
	F20-500	5	90	2.9	-	0.4	19
	F10-700	5	33	0.5	0.2	-	25
	F20-700	8	85	0.7	0.3	0.01	30

a the Cl -ion release is expressed as molar percentage in respect of the initial amount of Cl in the diuron. b the Cl

Acknowledgement

The authors thank the French Agence Nationale de la Recherche (ANR) for funding the Project COPHOTOFE under reference ANR-13-ECOT-0010.The European Fund for regional