Martin Bondesgaard

Aref Mamakhel

Jacob Becker

Hidetaka Kasai

Gilles Philippot

Martin Bremholm

Bo B Iversen

Gilles

Supercritical flow synthesis of Pt1-xRux nanoparticles: comparative phase diagram study of nanostructure versus bulk

come

Introduction

Recent decades have shown growing demands for noble metals in a variety of applications such as catalysis, electronics, optical sensing, and photothermal therapy. 1 Among these are platinum (Pt), which has exhibited increasing costs with consequent incentives for reducing consumption e.g. through design of effective, durable yet Pt-deficient catalysts. 1 To this end, an obvious mean is the use of nanoparticles. Several factors affect the performance of a catalyst including particle size (surface area), crystallinity, morphology, crystal structure, degree of agglomeration, surface chemistry and elemental composition by alloying or atomic substitution. Great efforts have been put into gaining control of these parameters in order to optimize catalyst performance. [START_REF] Quan | High-Index Faceted Noble Metal Nanocrystals[END_REF][START_REF] Xia | Shape-Controlled Synthesis of Colloidal Metal Nanocrystals: Thermodynamic versus Kinetic Products[END_REF] In this regard, alloying Pt with cheaper elements has the explicit advantage of further decreasing raw material costs. Among alloys the Pt-Ru system has attracted attention due to excellent catalytic properties [START_REF] Chakraborty | Electrochemical impedance spectroscopy study of methanol oxidation on nanoparticulate PtRu direct methanol fuel cell anodes: Kinetics and performance evaluation[END_REF][START_REF] Kardimi | Synthesis and characterization of carbon nanotubes decorated with Pt and PtRu nanoparticles and assessment of their electrocatalytic performance[END_REF][START_REF] Peng | Designer platinum nanoparticles: Control of shape, composition in alloy, nanostructure and electrocatalytic property[END_REF][START_REF] Yandieva | Alumina-platinum catalyst in the reductive dehydration of ethanol and diethyl ether to alkanes[END_REF] for many applications, e.g. as catalysts in fuel cells. [START_REF] Chakraborty | Electrochemical impedance spectroscopy study of methanol oxidation on nanoparticulate PtRu direct methanol fuel cell anodes: Kinetics and performance evaluation[END_REF][START_REF] An | Synthesis of PtRu/carbon nanotube composites in supercritical fluid and their application as an electrocatalyst for direct methanol fuel cells[END_REF] Here, the overall driving factor is not only the inclusion of comparatively cheap Ru, 1 but also its ability to greatly diminish the detrimental effect of Pt being poisoned by CO. [START_REF] Stolbov | High CO tolerance of Pt/Ru nanocatalyst: Insight from first principles calculations[END_REF] A wide variety of controlled methods to synthesize nanoparticles have been developed, and the literature on Pt-systems is large and ever expanding. Among the typical methods for preparing Ptbased catalysts are solvothermal, [START_REF] Mi | Pulsed-Flow Near-Critical and Supercritical Synthesis of Carbon-Supported Platinum Nanoparticles and In Situ X-ray Diffraction Study of Their Formation and Growth[END_REF][START_REF] Mi | The formation mechanism of bimetallic PtRu alloy nanoparticles in solvothermal synthesis[END_REF] polyol reduction, [START_REF] Chen | Polyol Synthesis of Platinum Nanostructures: Control of Morphology through the Manipulation of Reduction Kinetics[END_REF] chemical vapor deposition, [START_REF] Garcia | Chemical Vapor Deposition of Iridium, Platinum, Rhodium and Palladium[END_REF][START_REF] Jeon | Selective Chemical Vapor Deposition of Platinum and Palladium Directed by Monolayers Patterned Using Microcontact Printing[END_REF] wet impregnation, [START_REF] Kim | Shaped platinum nanoparticles directly synthesized inside mesoporous silica supports[END_REF][START_REF] Zhang | A Generic Wet Impregnation Method for Preparing Substrate-Supported Platinum Group Metal and Alloy Nanoparticles with Controlled Particle Morphology[END_REF] aqueous solution with different reducing agents (NaBH4, hydrazine etc.), [START_REF] Peng | Designer platinum nanoparticles: Control of shape, composition in alloy, nanostructure and electrocatalytic property[END_REF] flame aerosol, [START_REF] Chakraborty | Mixed Phase Pt-Ru Catalyst for Direct Methanol Fuel Cell Anode by Flame Aerosol Synthesis[END_REF] electromagnetic radiation assisted reduction [START_REF] Pal | Microwave-assisted synthesis of platinum nanoparticles and their catalytic degradation of methyl violet in aqueous solution[END_REF][START_REF] Abedini | A review on radiation-induced nucleation and growth of colloidal metallic nanoparticles[END_REF] etc. More than 20 years ago Adschiri and coworkers developed a method to synthesize nanoparticles in supercritical fluids (SCFs). [START_REF] Adschiri | Rapid and Continuous Hydrothermal Crystallization of Metal Oxide Particles in Supercritical Water[END_REF] They have a range of unique properties, e.g. gas-like transport properties combined with liquid-like solvation capability and density, which can be effectively manipulated by changes in pressure and temperature. This makes SCFs very useful as a synthesis medium for nanoparticles. Indeed, a wide range of complex nanomaterials have been prepared during the past decades. [START_REF] Adschiri | Rapid and Continuous Hydrothermal Crystallization of Metal Oxide Particles in Supercritical Water[END_REF][START_REF] Adschiri | Green materials synthesis with supercritical water[END_REF][START_REF] Aymonier | Review of supercritical fluids in inorganic materials science[END_REF][START_REF] Bozbag | Synthesis of nanostructured materials using supercritical CO2: Part II. Chemical transformations[END_REF][START_REF] Cansell | Review on materials science and supercritical fluids[END_REF][START_REF] Douglas | Crystal growth of gallium nitride in supercritical ammonia[END_REF][START_REF] Kimura | Synthesis of platinum nanoparticles in high-temperatures and high-pressures fluids[END_REF][START_REF] Knez | Industrial applications of supercritical fluids: A review[END_REF][START_REF] Saito | Research activities on supercritical fluid science and technology in Japan -A review[END_REF][START_REF] Sanli | Synthesis of nanostructured materials using supercritical CO2: Part I. Physical transformations[END_REF][START_REF] Zhang | Applications of supercritical carbon dioxide in materials processing and synthesis[END_REF] In the case of the Pt1-xRux system, most of the previous solvothermal synthesis studies were done for the pure metal or for a limited number of alloy compositions. In the majority of studies, various types of batch syntheses were used. [START_REF] Kardimi | Synthesis and characterization of carbon nanotubes decorated with Pt and PtRu nanoparticles and assessment of their electrocatalytic performance[END_REF][START_REF] An | Synthesis of PtRu/carbon nanotube composites in supercritical fluid and their application as an electrocatalyst for direct methanol fuel cells[END_REF][START_REF] Chen | Crystal growth of platinum-ruthenium bimetallic nanocrystallites and their methanol electrooxidation activity[END_REF][START_REF] Yang | Ultrasonically treated multi-walled carbon nanotubes (MWCNTs) as PtRu catalyst supports for methanol electrooxidation[END_REF][START_REF] Antolini | Formation of carbon supported PtRu alloys: an XRD analysis[END_REF][START_REF] Albers | Characterisation of carbon supported platinum-ruthenium fuel cell catalysts of different degree of alloying[END_REF] Two studies report on solvothermal flow-synthesis of Pt, PtRu and Ru at different temperatures. Mi et al. reports the formation and growth mechanism of PtRu nanoparticles. [START_REF] Bondesgaard | Guide to by-products formed in organic solvents under solvothermal conditions[END_REF] Most continuous solvothermal studies are for pure Pt-nanoparticles, [START_REF] Mi | Pulsed-Flow Near-Critical and Supercritical Synthesis of Carbon-Supported Platinum Nanoparticles and In Situ X-ray Diffraction Study of Their Formation and Growth[END_REF][START_REF] Kimura | Synthesis of platinum nanoparticles in high-temperatures and high-pressures fluids[END_REF][START_REF] Harada | Synthesis of colloidal particles of poly(2-vinylpyridine)coated palladium and platinum in organic solutions under the high temperatures and high pressures[END_REF][START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF] again including Mi et al. employing pulsed-flow supplemented by in-situ studies of the formation and growth of Pt. Mi et al. also reported on the formation of pure Ru nanoparticles [START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF] using pulsed flow synthesis and including details on their formation and growth mechanism based on different precursor materials. To the best of our knowledge, only two other continuous-flow synthesis methods have been reported for this system, continuous flame aerosol synthesis [START_REF] Chakraborty | Mixed Phase Pt-Ru Catalyst for Direct Methanol Fuel Cell Anode by Flame Aerosol Synthesis[END_REF] and a continuous thermal CO-reduction method. [START_REF] Alyami | Tailoring ruthenium exposure to enhance the performance of fcc platinum@ruthenium core-shell electrocatalysts in the oxygen evolution reaction[END_REF] Pt1-xRux nanoparticles have been reported to exhibit interesting phase behavior that deviates from the bulk phase diagram. [START_REF] Bondesgaard | Guide to by-products formed in organic solvents under solvothermal conditions[END_REF][START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF] The hcp phase is the thermodynamically stable phase while fcc is metastable. [START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF] Recently, synthesis of phase pure fcc Ru was reported. [START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF][START_REF] Kusada | Discovery of Face-Centered-Cubic Ruthenium Nanoparticles: Facile Size-Controlled Synthesis Using the Chemical Reduction Method[END_REF] Here, we present the first systematic investigation of flow-synthesized Pt1-xRux nanoparticles across the entire compositional range and investigate the particle sizes, microstructures, morphologies and allotropism.

Experimental section

Materials

The solvents toluene (99.9%) and absolute ethanol were purchased from Sigma-Aldrich and VWR respectively. The metal precursors Pt(acac)2 (acac = acetylacetonate, 98%) and Ru(acac)3 (≥97%) were purchased from Strem Chemicals and Sigma-Aldrich, respectively. All chemicals were used as received.

Synthesis

The syntheses were carried out in a flow-reactor shown schematically in Figure 1. [START_REF] Hald | Supercritical synthesis of complex nano materials[END_REF][START_REF] Mi | Glucose-assisted continuous flow synthesis of Bi2Te3 nanoparticles in supercritical/near-critical water[END_REF] The metal precursors were dissolved in an 82:18 vol% mixture of ethanol-toluene. The metal concentration was 10 mM. A total of 11 syntheses were carried out with composition increments of 10 mol% from Pt0.0Ru1.0 to Pt1.0Ru0.0. A 200 mL precursor solution was loaded into an injector vessel [START_REF] Houlberg | Hydrothermal Synthesis and in Situ Powder X-ray Diffraction Study of Bismuth-Substituted Ceria Nanoparticles[END_REF] in order to prolong the lifespan of the pumps. The flow-rates of precursor and solvent (absolute ethanol) were 7 mL/min and 15.6 mL/min, respectively. The pressure was controlled by a proportional relief-valve (PRV) set to 200 bar. The solvent stream was preheated to 450°C and subsequently mixed with the ambient temperature precursor stream in a T-piece mixing geometry. The mixed stream passed through a vertical reactor also heated to 450°C, and the residence time in the reactor falls in the range 6-11 s as calculated from reactor volume, flow-rate and fluid density obtained at either Tmax (450 o C) or Tmixing (318 o C), respectively. Here, the solvent density [START_REF] Golubev | Experimental study of the density of aliphatic alcohols at various temperatures and pressures[END_REF] is assumed equal to that of pure ethanol. The product stream was cooled to ambient temperature before collection of the products.

The collected product suspensions were centrifuged and washed with ethanol several times and finally dried at ambient conditions.

Characterization

High resolution synchrotron powder X-ray diffraction (S-PXRD) measurements were performed at beamline BL44B2 [START_REF] Kato | The RIKEN Materials Science Beamline at SPring-8: Towards Visualization of Electrostatic Interaction[END_REF] , SPring-8, Japan, using a wavelength of 0.50037 Å. The data were analyzed with the FullProf software suite [START_REF] Rodriguez-Carvajal | FULLPROF: A Program for Rietveld Refinement and Pattern Matching Analysis[END_REF] using both Rietveld and Le Bail refinements. The instrumental profile broadening was estimated by refinement of a model to data collected on a CeO2 standard sample (NIST SRM 674b) and included in the refinements (IRF file). The background was described by a 12th order polynomial. Peak profiles for the phase region x < 0.3 were modeled with the Thompson-Cox-Hastings pseudo-Voigt function; for x ≥ 0.3 when the second phase emerges, a Le Bail model was used. Trial refinements showed that a purely Lorentzian line profile (X and Y parameters) was adequate, since the Gaussian parameters refined to values of zero (within one standard deviation) when the instrumental profile broadening was included. The X and Y profile parameters are related to the crystallite strain and size, respectively. The crystallite size was determined using the Scherrer equation. The Rietveld refinements were also used to quantify the weight fraction of the observed phases. Transmission Electron Microscopy (TEM) was performed on a Philips CM20 instrument equipped with a LaB6 filament operated at 200 keV. High resolution (HR) TEM images and elemental mapping by energy-dispersive X-ray spectroscopy (EDX) measurements were carried out on a FEI "Talos" F200X analytical (S)TEM-microscope operated at 200 keV. The program ImageJ (FIJI) was used for statistical size-analysis of the nanoparticles and for the determination of d-spacing from HR-TEM images by Fast Fourier Transformation (FFT) of individual domains showing well-defined lattice fringes.

Results and discussion

Phase behavior

The measured synchrotron PXRD patterns are shown in Figure 2a. For x ≤ 0.2 a single fcc phase is observed, while for x ≥ 0.3 reflections indexed to the hcp phase become increasingly dominant. This is in contrast to the bulk phase diagram [START_REF] Hutchinson | Solubility Relationships in the Ruthenium-Platinum System Platinum Met[END_REF] wherein a single-phase fcc alloy exists for 0 < x < 0.62 and an hcp alloy for 0.8 < x < 1. In the intermediate range (0.62 < x < 0.8) a biphasic region exists. The PXRD data obtained in this study show a significant expansion of the biphasic region when moving from bulk to nanoparticles. The region of single phase hcp is completely eliminated, leaving only one monophasic region in the whole phase diagram (fcc at x ≤ 0.2). Similar behavior has been reported by Kusada et al. [START_REF] Kusada | Solid Solution Alloy Nanoparticles of Immiscible Pd and Ru Elements Neighboring on Rh: Changeover of the Thermodynamic Behavior for Hydrogen Storage and Enhanced CO-Oxidizing Ability[END_REF] on the Pd1-xRux system where a large biphasic region was found consisting of the fcc and hcp phases. Phase fractions were derived from Rietveld refinements in the biphasic region [START_REF] Bish | Quantitative phase analysis using the Rietveld method[END_REF][START_REF] Dinnebier | Powder diffraction: theory and practice[END_REF] and they are plotted against the composition in Figure 3a. The fraction of the hexagonal phase gradually increases with growing ruthenium content. A small discrepancy to this trend is seen in the step from x = 0.9 to x = 1.0, which may arise from the absence of Pt in the last sample. The growing hcp content is consistent with a transition towards the stoichiometric range where the hcp phase is thermodynamically favored.

The crystal strains extracted from the Rietveld refinements for both the fcc and hcp phases are plotted in the lower part of Figure 3a. The fcc phase shows an increasing strain from x = 0.0 to 0.4, where it becomes roughly constant from 0.4 to 0.6. This reflects the stress imposed on the fcc structure by the hcp-favoring ruthenium. From x = 0.6 to 0.8 the fcc strain abruptly decreases which correlates well with the phase segregation seen in the bulk phase diagram. [START_REF] Hutchinson | Solubility Relationships in the Ruthenium-Platinum System Platinum Met[END_REF] Generally, from x ≥ 0.5 the hcp phase exhibits the least strain of the two phases, a reflection of its thermodynamic stability a) b)

in elemental bulk Ru and the driver of its growing volume fraction. Meanwhile, at x ≥ 0.8 the strain in the receding fcc phase once more increases as now imposed by the dominant hcp phase.

Figure 2b shows the progressive peak shift toward higher scattering angle with increasing Ru content, confirming the existence of a solid solution. Unit cell parameters were extracted from the Rietveld refinements, and in Figure 3b these are plotted against composition for the fcc as well as the hcp phase. For fcc, the a-axis of the unit cell follows a decreasing, linear trend with increasing Ru content. This is in agreement with Vegard's law. Kusada et al. [START_REF] Kusada | Solid Solution Alloy Nanoparticles of Immiscible Pd and Ru Elements Neighboring on Rh: Changeover of the Thermodynamic Behavior for Hydrogen Storage and Enhanced CO-Oxidizing Ability[END_REF] reported a similar trend for the Pd1-xRux system where it was shown possible to alloy the two elements over the entire composition range. The same trend is reflected in the c-axis in the hcp phase, whereas the a-axis remains nearly constant. These data are plotted as a*1.633 (the c/a ratio of an ideal hcp structure) in order to clarify deviations from ideal behavior e.g. compression or expansion. The a-values for hcp are given in the supporting information table S.1. The fcc interatomic distance is found to be larger than distances along all axes in hcp for x < 0.9, which is seen as a "driver" of the hcp formation as it implies weaker bonding in the fcc structure compared with the hcp structure. The bond strength in hcp only shows significant changes in the cdirection. The compression can be illustrated from the c/a ratio, which starts from a value close to the ideal ratio (1.633) and then decreases toward close to the bulk ruthenium value at x = 1.0. This trend is in agreement with the observed trend in strain, which corroborates the transition from a metastable to a more thermodynamically favored state.

Nanoparticle sizes

The particle sizes observed in TEM and the crystallite sizes derived from PXRD are shown in Figure 4. Note that figures for TEM have been converted from being number weighted into volume weighted in order to offer direct comparison with PXRD (<D>PXRD = <D 4 >/<D 3 >). The coincidence between TEM and PXRD is striking. It is concluded that the nanoparticles are fully crystalline, as also confirmed by HR-TEM (presented below). Considering, the volume weighting of the sizes, it is not surprising that size estimates from PXRD follow the crystal phase with the largest particles (revealed by TEM) at any given composition. Generally, stepping up from x = 0 the particle size gradually decreases from approx. 10 nm to 5 nm at x = 0.5. In the region x = 0.4-0.6 it levels out and becomes more or less constant at higher values of x. The trend of decreasing particle size for alloys compared to pure metals has been previously reported. [START_REF] Chen | Crystal growth of platinum-ruthenium bimetallic nanocrystallites and their methanol electrooxidation activity[END_REF] A correlation was found by Ouyang et. al. [START_REF] Ouyang | Atomistic origin of lattice strain on stiffness of nanoparticles[END_REF] showing an increased strain with decreasing particle size. This could be a contributing factor to limit the particle growth in combination with additional effects discussed below. Another driving factor for the splitting into two phases can be found in the in-situ studies by Mi et al. The temperature-resolved study of pure Ru (200, 250, 300 and 400°C) identified a trend that the amount of fcc phase was declining as the temperature was increased. [START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF] The in situ PXRD data revealed that the phases form simultaneous. This is evidence that the presence of fcc Ru is kinetically driven and implies a lower energy barrier for formation of fcc compared to hcp Ru. In the in situ PXRD study on Pt0.5Ru0.5 [START_REF] Mi | The formation mechanism of bimetallic PtRu alloy nanoparticles in solvothermal synthesis[END_REF] 7a-d) where some particles are found to have oblong areas with different structures and corresponding variation in composition ranging from Ruand Pt-rich regions, respectively. More details on these particles are provided in section 3.3.

It is possible that the relationship between fcc and hcp is not explained by kinetics alone, i.e. they could be affected by thermodynamic factors as well. In particular, it is known that many metastable phases cannot exist above a certain critical particle size [START_REF] Schonecker | Metastable cubic and tetragonal phases of transition metals predicted by density-functional theory[END_REF][START_REF] Kitakami | Size effect on the crystal phase of cobalt fine particles[END_REF][START_REF] Rong | Phase transformations and phase stability in nanocrystalline materials[END_REF][START_REF] Waitz | Phase Transformations of Nanocrystalline Martensitic Materials[END_REF] due to a lower surface energy than that of thermodynamically favored (bulk) phases. Hence, metastability occurs when the particle becomes sufficiently small. An example is the study by Kitakami et al. [START_REF] Kitakami | Size effect on the crystal phase of cobalt fine particles[END_REF] on Co nanoparticles, where the otherwise metastable fcc phase was found to be stable (phase pure) under 20 nm while only the hcp phase was observed for sizes above 110 nm. For intermediate sizes, a mixture of the two phases was observed with a decreasing amount of the fcc phase as particle size increased. In the present study, the crystallite size of the hcp phase is persistently lower than that of the fcc phase up until x ≥ 0.9. If it is assumed that the fcc phase is thermodynamically stable in the region 0 < x < 0.62 (and the hcp phase metastable), this behavior could be correlated to the bulk phase diagram.

For the nanoparticle phase diagram in the region 0.62 < x < 0.80 both phases coexist; this would correspond to the bulk phase in the region 0.8 < x < 1.0. In the nanosystem, the above thermodynamical argument hinges on the assumption that the critical particle size "inverts" in the composition range 0.62 < x < 0.8, rendering the fcc metastable for x ≥ 0.8. However, literature provides no clear picture of why such an inversion would occur. The matter of a potential critical particle size in this nanostructured system, must therefore be a subject of further studies. Yet a proposed explanation for the presence of a maximal critical particle size can arise from a combination of the facts discussed previously. When formation of a metastable structure is controlled by kinetics, the temperature need to be kept under a certain threshold to form the structure of interest to avoid sufficient energy in the system to pass the larger energy barrier for the the thermodynamically favored structure. When the reaction temperature is kept below this threshold, there is often also insufficient energy to initiate the further particle growth by either Oswald ripening or coalescence. These conditions favor that the particles size remain small, and this also makes the surface energy more dominant term in determining phase stability (higher fraction of surface atoms). When particle formation takes place above the temperature threshold, then the balance between the conditions is shifted toward forming the thermodynamic favored phase either partly or fully as seen in the study by Mi et al. [START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF] This may explain why inversion of crystallite sizes is observed for the two phases.

Table 1 shows a comparison with reported values (at 400 o C) for crystallite size and unit cell parameters at x = 0 and 1. [START_REF] Mi | Pulsed-Flow Near-Critical and Supercritical Synthesis of Carbon-Supported Platinum Nanoparticles and In Situ X-ray Diffraction Study of Their Formation and Growth[END_REF][START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF] For x = 0 at 400°C, [START_REF] Mi | Pulsed-Flow Near-Critical and Supercritical Synthesis of Carbon-Supported Platinum Nanoparticles and In Situ X-ray Diffraction Study of Their Formation and Growth[END_REF] the unit cell parameter found by Mi et al. are a bit below that presented here (450 o C), whereas the crystallite size obtained is significantly larger.

For Ru nanoparticles (x = 1) the unit cell parameters for the hcp phase are reported to slightly exceed the present data. In the studies by Mi et al., [START_REF] Mi | Pulsed-Flow Near-Critical and Supercritical Synthesis of Carbon-Supported Platinum Nanoparticles and In Situ X-ray Diffraction Study of Their Formation and Growth[END_REF][START_REF] Mi | The formation mechanism of bimetallic PtRu alloy nanoparticles in solvothermal synthesis[END_REF][START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF] the precursor Ru(acac)3 resulted in a mixture of fcc and hcp Ru at all temperatures (200-400°C) whereas for the RuCl3, pure hcp phase was consistently observed. This suggests that the formation of hcp Ru nanoparticles presumably progresses through a reactions path mediated by anions. This corresponds well to literature where halide ions (e.g chloride) are known to be templating agents in the synthesis of metal nanoparticles e.g. to favor certain crystal structures [START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF][START_REF] Kusada | Discovery of Face-Centered-Cubic Ruthenium Nanoparticles: Facile Size-Controlled Synthesis Using the Chemical Reduction Method[END_REF] or particle morphologies. [START_REF] Lohse | Anisotropic Noble Metal Nanocrystal Growth: The Role of Halides[END_REF] Clearly the choice of precursor is crucial for obtaining fcc Ru. This has also been suggested by Kusada et al., [START_REF] Kusada | Discovery of Face-Centered-Cubic Ruthenium Nanoparticles: Facile Size-Controlled Synthesis Using the Chemical Reduction Method[END_REF][START_REF] Kusada | A Route for Phase Control in Metal Nanoparticles: A Potential Strategy to Create Advanced Materials[END_REF] who tested the hypothesis by using different Ru precursors (chloride, acetylacetonate and acetate) in the same reducing environment. Here the chloride yielded pure hcp phase whereas both the acetylacetonate and the acetate yielded phase pure fcc. The study by Ye et al. [START_REF] Ye | Ru Nanoframes with an fcc Structure and Enhanced Catalytic Properties[END_REF] has shown that it is possible to obtain the Ru fcc structure from a chloride precursor, which apparently contradicts the previous statement. However, their synthesis is based on heterogenous nucleation using seeded growth (Ru is deposited on a Pd seed). It is therefore reasonable that a different formation mechanism is observed compared with the homogeneous nucleation used here. When comparing the present data with in situ studies it is seen that the precursor also influences crystallite size with chloride precursors yielding larger particles than acetylacetonate precursors 37 at x = 0 and x = 1.

acetylacetonate was reduced in ethylene glycol yielding phase pure hcp Ru whereas reduction by ethanol plus triethylene glycol resulted in pure fcc Ru. [START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF][START_REF] Kusada | Discovery of Face-Centered-Cubic Ruthenium Nanoparticles: Facile Size-Controlled Synthesis Using the Chemical Reduction Method[END_REF] Using ethanol with butanediol produced mixtures of the two crystal phases. [START_REF] Mi | Controlling Allotropism in Ruthenium Nanoparticles: A Pulsed-Flow Supercritical Synthesis and in Situ Synchrotron Xray Diffraction Study[END_REF][START_REF] Ye | Ru Nanoframes with an fcc Structure and Enhanced Catalytic Properties[END_REF] This likely hinges on the reported temperature-controlled reaction behavior of organic solvents at solvothermal condition. [START_REF] Bondesgaard | Guide to by-products formed in organic solvents under solvothermal conditions[END_REF]

Elemental composition and morphology

EDX mappings of individual nanoparticles were performed for three samples (x = 0.2, 0.5 and 0.8)

to investigate the elemental composition and distribution in the nanoparticles, Figure 5. Additional examples are presented in supporting information (Figure S1). The maps reveal an even distribution of the two elements for x = 0.2. For the two other samples the distribution in the particles is mostly uniform although small fluctuations are seen. These might explain the small deviations of unit cell parameters seen in Figure 3b in the biphasic region. A similar trend was seen in a study by Hills et al, [START_REF] Hills | The Size-Dependent Structural Phase Behaviors of Supported Bimetallic (Pt-Ru) Nanoparticles[END_REF] who reported broad difference in composition of Pt-Ru nanoparticles.

) c) d) e) f) g) h) i) j) k) l) (d111=2.2 Å)
or the hcp (101) plane (d101=2.1 Å). From this, the long flat end facets of the elongated particles could be determined to be {111} and {101} for fcc and hcp respectively.

The phase shift between hcp and fcc by a stacking fault requires only a small shift in stacking sequence. Similar kinds of staking faults have been reported for other types of compounds. [START_REF] Ahmad | A comprehensive study on the mechanism behind formation and depletion of Cu2ZnSnS4 (CZTS) phases[END_REF][START_REF] Mayence | Probing planar defects in nanoparticle superlattices by 3D small-angle electron diffraction tomography and real space imaging[END_REF][START_REF] Lagrow | Can Polymorphism be Used to form Branched Metal Nanostructures?[END_REF] The observed stacking faults indicate that only a small energy difference separates the two crystal structures incurring easy transformation from one structure to the other. Indeed, it has been shown by DFT calculations that the energy difference between the hcp and fcc structures of Ru is small. [START_REF] Schonecker | Metastable cubic and tetragonal phases of transition metals predicted by density-functional theory[END_REF] This explains why it is possible for the two phases to co-exist across such a wide composition range.

A similar effect can be found in a fcc structure, when a high concentration of twin planes are present, which break the ABCABC stacking order and mimics the ABABAB stacking order in hcp structure on the atomic level. [START_REF] Fanfair | Twin-Related Branching of Solution-Grown ZnSe Nanowires[END_REF] When a "chain" of elongated particles was analyzed by EDS line scan,

Conclusion

The relations between crystal phase and microstructural behavior has been established at the nanoscale for the PtxRu1-x system. Crystallite and particle size were both found to decrease as the Ru content increased. The phase behavior of the system was found to deviate considerably from the bulk phase diagram and at the nanoscale the biphasic region is expanded to cover most of the phase diagram. The phase fractions produced are greatly influenced by the choice of precursor, solvent composition and synthesis parameters. As for the influence of precursor, it appears to be mediated by surfactant effects with acetylacetonate directing the product towards a metastable fcc state. Ultimately, this study has demonstrated the feasibility of employing continuous-flow synthesis for up-scale production of catalytic Pt-Ru nanoparticles, and elucidated options for "tailoring", which are offered by controlled synthesis and alloying.

Figure 1 .

 1 Figure 1. Diagram showing the reactor used for continuous flow supercritical synthesis of Pt1-xRux nanoparticles.

Figure 2 .

 2 Figure 2. a) S-PXRD for the Pt1-xRux system with Bragg reflections for the fcc phase (diamonds) and the hcp phase (solid

Figure 3 .Figure 3c

 33c Figure 3. a) Weight fractions in the two phase region and (below) strain values for the two phases obtained from the Rietveld refinements. Dashed lines indicate the bulk phase diagram. b) Unit cell parameters extracted from Rietveld

Figure 4 .

 4 Figure 4. Average particles size from TEM (volume weighted) and the crystallite size for the two phases (cubic and hexagonal) from PXRD. Dashed lines indicate the bulk phase diagram.

Figure 5 .Figure 6 .

 56 Figure 5. a-d) EDX of Pt0.8Ru0.2; e-h) EDX of Pt0.5Ru0.5;i-l) EDX of Pt0.2Ru0.8.

  Figure7cd, compositional fluctuation was identified across the row of particles. This can be correlated to the effects discussed above about imperfect alloying. The difference in stacking order between the fcc and hcp structures influences the nature of the facets at the surface of the particles, which can lead to different morphologies. This can explain why the morphologies change when the the hcp phase emerges for x ≥ 0.3 as also demonstrated by Kambe et al.[START_REF] Kambe | Effects of crystal structure, size, shape and surface structural differences on photo-induced electron transport in TiO2 mesoporous electrodes[END_REF]

Figure 6 .Figure 7 .

 67 Figure 6. TEM images a) Pt1.0Ru0.0 b) Pt0.9Ru0.1 c) Pt0.8Ru0.2 d) Pt0.7Ru0.3 e) Pt0.6Ru0.4 f) Pt0.5Ru0.5 g) Pt0.4Ru0.6 h) Pt0.3Ru0.7 i) Pt0.2Ru0.8 j) Pt0.1Ru0.9 k) Pt0.0Ru1.0. All the dashed boxes shows particles with internal phase shift due to stacking faults.

  it was concluded that Pt nanoparticles are formed first followed by the formation of a surface layer of Ru on Pt nanoparticles which diffuse into the Pt nanoparticles to form the Pt0.5Ru0.5 alloy. Unfortunately this measurement is made at 200 o C, with no data existing for 400 o C. In the present study, done at 450 o C, the time scale of the syntheses is significantly shorter (6-11 seconds) than that observed for this diffusion process (2 min). Based on this, it appears that the high temperature allows completion of parallel reactions of Pt-Ru and pure Ru due to different reaction time for Pt and Ru formation. On the other hand the short reaction time means that diffusion cannot be completed leading to two distinct, separate phases crystallize even at low Ru contents. This is supported by STEM-EDS (Figure

Table 1 .

 1 Crystallite size and unit cell parameters

	Mi et al. 10, 37		This study
	Crystallite size	Unit cell parameters Crystallite size	Unit cell parameters
	x = 0.0 24.7 nm	3.903 Å	10.3 nm	3.915 Å
	x = 1.0 4.9/8.6 nm (fcc/hcp) 2.728/4.325 Å (a/c) 4.3/5.9 nm (fcc/hcp) 2.704/4.292 Å (a/c)

Acknowledgments

This work was supported by Center for Materials Crystallography (DNRF93), the project Fuel Scale Integrate (Innovation Foundation Denmark) and the Danish Agency for Science, Technology and Innovation (DANSCATT). The synchrotron experiments were performed at SPring-8 with the approval of RIKEN (Proposal No. 20150018 and 20160037). J. Chevalier is acknowledged for his kind technical assistance related to the HR-TEM imaging.

Supporting Information

The Supporting Information is available free of charge on the ACS Publications website.

Additional results from Rietveld refinements, EDX mapping of selected samples, TEM size distributions.


Elodie Boursier
email: elodie.boursier@u-bordeaux.fr

Marja Giedre

Jean-Christophe Archipovaite

Stephane Delagnes

Guilmot Petit

Philippe Ernotte

Patricia Lassonde

Benoit Segonds

Yannick Boulanger

François Petit

Légaré

Marija Giedre

Stéphane Delagnes

Patricia Lassonde

Benoît Segonds

Dmitry Légaré

Eric Roshchupkin

Cormier

Giedre Marija Archipovaite

Jean-Christophe Delagnes

Stephane Petit

Guilmot Ernotte

Philippe Lassonde

Patricia Segonds

Benoit Boulanger

Yannick Petit

François Legare

Dmitry Roshchupkin

Eric Cormier

Study of middle infrared difference frequency generation using a femtosecond laser source in LGT

Many applications such as linear and multidimensional infrared spectroscopy require the generation of widely tunable femtosecond pulses with high average power covering the atmospheric transmission window (3 -5 µm). Furthermore, long wavelength few-cycle pulses are of prime importance in strong field physics especially when carrier-envelope phase effects are investigated. Since too few laser materials with sufficient bandwidth in this spectral domain has been developed yet, the only alternative is phase-matched difference frequency generation (DFG) and optical parametric amplification (OPA) in nonlinear crystals pumped by powerful high beam quality laser sources. Both DFG and OPA covering the SWIR (1.4 -3 µm) and part of the MIR (3 -5 µm) bands have been successfully demonstrated in widespread nonlinear crystals such as -BaB2O4 (BBO), KTiOAsO4 (KTA), and LiNbO3 (LNB) [1][2][3][4][5][6][7][8][9][10][11] associated with commercial femtosecond pump laser sources based on Yb 3+ -or Nd 3+ -doped materials (~ 1 µm) or Ti:Sa laser (~ 0.8 µm). Recently, we identified Langatate La3Ga5. 5Ta0.5O14 (LGT), initially known for its piezoelectric properties, as another promising nonlinear crystal with a transparency window ranging from 0.3 up to 6 µm [12,13]. Czochralski growth method may provide centimeter-size Langatate crystals with very good optical quality and homogeneity.

LGT belongs to the trigonal 32 point group and is a positive uniaxial crystal. The Sellmeier coefficients retrieved by the sphere method are reported in [13] with a nonlinear coefficient d11 = 2.4 ± 0.4 pm/V at λ = 0.67 µm. Optical damage appeared at 4.3 GW/cm 2 , i.e. 21.7 J/cm², when illuminated by a 1.064 µm (10 Hz, 5 ns) Nd:YAG laser with a 220 µm beam diameter [13]. Calculations showed that a LGT slab cut along a phase-matching (PM) direction (𝜃 𝑃𝑀 = 56° and 𝜑 𝑃𝑀 = 90°) and pumped at 0.967 µm can generate broadband pulses spanning from 1.5 to 3.5 µm [13]. In this Letter, we report to the best of our knowledge, on the first study of middle infrared difference frequency generation in LGT crystal using ultrashort femtosecond pulses in the femtosecond regime. The investigations carried out at low intensity in this Letter highlight the potentiality of LGT as a novel nonlinear crystal for DFG and optical parametric amplification in the middle infrared. Two kinds of sources provided the incoming pump and signal beams. One was an amplified continuum generator pumped around 1 µm by using a high-repetition-rate and high-averagepower laser based on Yb 3+ -doped material. The second source was a TOPAS ® OPA pumped by a Ti:Sa laser. The tunability of the generated wavelength was measured in the same rotating LGT slab, from three experimental configurations. We also determined the optical damage threshold and the nonlinear refractive index 𝑛 2 of LGT.

Our calculations, using Sellmeier coefficients from ref. [13], predict that phase-matched Type-I DFG/OPA (𝜆 𝑝 𝑜 , 𝜆 𝑠 𝑒 , 𝜆 𝑖 𝑒 ) can be generated in a LGT slab cut at 𝜃 PM = 58.3° in the 𝑦-𝑧 plane (𝜑 𝑃𝑀 = 90°) by combining two incoming beams collinearly. It generates a beam continuously tunable while rotating the slab. Here 𝜆 𝑝 and 𝜆 𝑠 are the incoming pump and signal wavelengths, respectively, while 𝜆 𝑖 is the corresponding idler wavelength, with the convention 𝜆 𝑝 < 𝜆 𝑠 ≤ 𝜆 𝑖 . Then during the DFG/OPA process, the beam at 𝜆 𝑠 exhibits an amplification (OPA) while the beam at 𝜆 𝑖 is generated (DFG). The superscripts 𝑜 and 𝑒 stand for the ordinary and extraordinary waves, respectively. In order to validate our calculations and to evaluate the full potential of LGT, three different experimental configurations (a), (b) and (c) were implemented, as shown in Fig. 1. In (a), a broadband pump at 𝜆 𝑝 = 0.7 µm is mixed with a narrowband signal at 𝜆 𝑠 = 1.03 µm. In (b), a narrowband pump at 𝜆 𝑝 = 1.03 µm is mixed with a broadband signal at 𝜆 𝑠 = 1.5 µm. Finally in (c), a broadband (80 nm) pump at 𝜆 𝑝 = 1.3 µm and a broadband (135 nm) signal at 𝜆 𝑠 = 1.9 µm are mixed together. In (a) and (b), we used a homemade Yb 3+ -doped fiber chirpedpulse amplification (FCPA) source based on a rod-type large mode area fibers operating at 100 kHz. The 350 fs (FWHM) laser pulses emitted at λ = 1.03 µm are here considered as narrowband (~6 nm), and the measured energy was about 60 µJ at the input of the LGT sample. One part of the 1.03 µm beam is smoothly focused in an undoped YAG (Y3Al5O12) rod in order to generate a broadband continuum (0.6 -1.8 µm). Out of this continuum, we can either select a broadband (90 nm FWHM) visible beam centered at 0.7 µm using (a), or a broadband (436 nm FWHM) IR spectrum centered at 1.5 µm using (b). Consequently, the weak visible beam was first amplified by OPA in a -BBO crystal pumped at 0.515 µm (SHG of 1.03 µm) in order to reach about 60 nJ. The 1.5 µm beam energy was less than 5 nJ, and the beam sizes at 1/e² were about 180 µm for both beams in case of (a) and (b). For (c), we used a commercial 2.5 kHz repetition rate and 35 fs pulse width TOPAS ® OPA pumped by λ = 0.78 µm of a Ti:Sa laser. The incoming beams size at 1/e² was 180 µm with 1 µJ energy level. Depending on the central wavelength, the spectra of the incoming pump and signal beams are recorded with an Ocean Optics HR2000 spectrometer (0.4 -1.1 µm,  = 1 nm) and a NIRQuest (0.9 -2.5 µm,  = 7 nm). The idler spectra are measured with a Mozza ® Fastlite™ monochromator (1 -5 µm,  = 5 cm -1 ) in (a) and (b), and a homemade scanning monochromator (1.5 -5 µm,  = 2 nm) in (c). Figures 2a, 3a and4a show spectra of the incoming beams and of the phase-matched idler beam by rotating LGT. Note that for sake of clarity, the measured amplified signal spectra in configuration (b) are not shown. The LGT slab is cut at the predicted optimal angle 𝜃 PM = 58.3° in the 𝑦-𝑧 plane, and controlled using X-Rays in backscattering geometry. In order to limit the group-velocity mismatch between the pump and signal beams, a slab thickness of 𝐿 = 3.84 mm was used in the three experiments. It was polished to optical quality and mounted on a rotating stage. The incident angle was measured by using the incident-to-reflected beam angle method. This technique results in ±1° error on the measured angle. Fig. 2. Experimental data for the generated beam at 𝜆 𝑖 (red) using a broadband pump at λp = 0.7 µm (blue) and a narrowband signal at λs = 1.03 µm (green) as incoming beams in a LGT slab using configuration (a). a) Incoming and idler beams normalized intensities, and b) internal angles are plotted as a function of the wavelength. Blue circles are the associated selected pump wavelengths, and lines are calculations from ref. [13].

The phase-matching conditions are recovered by determining the wavelength of the maximum value of all recorded spectra. Then we plotted the tuning curves in LGT based on the internal angle as a function of the phase-matching wavelength 𝜆 𝑖 (see Figs. 2b, 3b and4b) and 𝜆 𝑠 (see Fig. 3b). Due to the weak signal gain in configurations (a) and (c), it was not possible to observe any significant spectral modification with the spectrometers. However in (c), the corresponding tuning curve was calculated using energy conservation between the measured idler wavelength and the pump set at λp=1.32 µm, which corresponds to the peak value of the pump spectrum (see dots in Fig. 4b). Based on the same principle, we also determined in configuration (a) the pump wavelengths from the broadband spectrum shown in Fig. 2a that were combined with the narrowband signal, in order to generate the phase-matched idler wavelengths shown in Fig. 2b. Calculations based on ref. [13] agree very well with our experimental data within the precision of our measurements in the three experimental configurations. For the sake of clarity, in the case of broadband pump and signal (configuration (c)), we have calculated the phase-matching curves (see Fig. 4b) corresponding to three pump wavelengths: 𝜆 𝑝2 = 1.32 µm (maximum of the pump spectrum), 𝜆 𝑝1 = 1.18 µm, and 𝜆 𝑝3 = 1.41 µm (wavelengths corresponding to 50 % of pump maximum spectral intensity). All recorded data fit within the internal phase-matching curves. Note that the measured spectral bandwidth of the idler waves matches with the simulated curve after considering all the incoming experimental parameters, i.e. pulse duration, beam sizes, pulse energy, and sample characteristics. Thinner crystal would certainly provide broader spectra at all wavelengths. In Fig. 5 we compare the spectral range of the generated wavelengths in LGT with that obtained in -BaB2O4 (BBO), KTiOAsO4 (KTA) and Periodically-Poled LiNbO3 (PPLN) from refs. [1][2][3][4][5][6][7][8][9][10][11]. The transparency range of each crystal is shown with the dashed white bars and the measured spectral range of the generated wavelengths through the blue and red colored bars. The later corresponds to cases for which the nonlinear crystals are pumped by two femtosecond commercial laser sources around either 0.8 µm (blue bars) or 1 µm (red bars). Figure 5 points out that the measured spectral ranges of LGT are comparable to that of PPLN when pumped around 1 µm. In comparison BBO and KTA are much more limited in this aspect. Lines are calculations after ref. [13] for three selected quasimonochromatic pump wavelengths 𝜆 𝑝1 = 1.18 µm, 𝜆 𝑝2 = 1.32 µm and 𝜆 𝑝3 = 1.41 µm from the pump spectral width.

The optical damage threshold was studied by using a 500 µm thick slab cut at 𝜃 = 9° in the 𝑦-𝑧 plane where we found no phasematching condition [13]. We used a sub-15 fs pulse duration, 100 Hz repetition rate Ti:Sa laser (800 nm) and the same experimental setup as in ref [14]. The energy was very low, i.e. around 1 µJ, and the beam was tightly focused at the surface of the sample using an off-axis parabola mirror. Accurate calibration and measurement give a 10 µm ± 5 % beam waist diameter [14]. The LGT slab sits on a 3-axis motorized translation stage in order to control the optimal focusing position. We adopted a so-called 1-on-1 protocol, i.e. one shot per site and ten tested sites per considered fluence, while maintaining a normal incidence and sufficient spacing to avoid crosstalk. The energy, and thus the fluence, were the only varying parameters. A preliminary calibration was achieved with fused silica as a reference material. Then, the LGT crystal was processed to determine the damage probability by increasing the fluence continuously. A fluence of 425 mJ/cm² still led to 0 % damage while the probability reached 80 % at 500 mJ/cm², setting then the optical damage threshold. This result corroborates the optical damage threshold measured in nanosecond regime [13] if corrected for the pulse duration difference [15]. Finally, we directly measured the third-order nonlinear electric susceptibility 𝜒 𝑥𝑥𝑥𝑥 (3) of LGT at 0.343 µm from third-harmonic generation performed at the surface of a 820 µm thick slab cut at (θ = 57°, φ = 90°). For this purpose, we used an Amplitude Systemes t-Pulse ® oscillator (390 fs, λω = 1.030 µm, 9 MHz) with a linear polarization oriented parallel to x-axis of LGT, and implemented the setup described in ref [16]. We took into account the Fresnel losses, the third-harmonic transmission, the nonlinear interaction length and the refractive indices corrections at λω and λ3ω = λω/3 = 0.343 µm [16]. We studied a fused silica slab in the same conditions, to use 𝜒 𝑥𝑥𝑥𝑥 𝑆𝑖𝑂 2 (0.343 µm) = 2.48×10 -22 m².V -2 from [17] as a reference. We found for LGT that 𝜒 𝑥𝑥𝑥𝑥 (0.343 µm) = 18 𝜒 𝑥𝑥𝑥𝑥 𝑆𝑖𝑂 2 (0.343 µm) = (4.50.5)10 -21 m².V -2 . The nonlinear refractive index 𝑛 2 is linked to 𝜒 𝑥𝑥𝑥𝑥 (3) by the relation 𝒏 𝟐 (λ 3ω ) = 𝟑𝝌 𝒙𝒙𝒙𝒙 (λ 3ω )

𝟒𝒏 𝟎 (λ 3ω )𝜺 𝟎 𝒄 where n0 stands for the ordinary refractive index of

LGT, 𝒄 is the light speed and 𝜺 𝟎 the vacuum dielectric constant. Then we found for LGT that 𝒏 𝟐 (0.343 µm) = (6.3±0.7) × 10 -19 m 2 .W -1 , which is about ten times that of BBO and three times that of KTA or PPLN at 0.780 µm [13,18]. It means that LGT may be associated with stronger self-focusing due to Kerr effect and then lead to potential damage. However, it is worth to mention that LGT can be grown in the largest sizes and the greatest homogeinity [13] if compared to KTA for which inhomogeneity is a real issue.

In conclusion, we performed the optical characterization of LGT in the femtosecond regime. Three experimental configurations of DFG/OPA were studied and led to the angular tuned generation of wavelengths ranging between 1.2 and 4.7 µm. These results show that LGT stands as a serious competitor of BBO, KTA and PPLN for parametric generation in the atmospheric transmission window.

In particular the high optical damage thresholds of LGT we estimated may favor this crystal compared to well-known nonlinear crystals. Finally we directly measured the third-order nonlinear electric susceptibility of LGT and determined the corresponding value of the nonlinear refractive index 𝑛 2 of LGT. Further investigations will study the experimental parametric gain and evaluate the potential of LGT for large aperture high beam quality OPA in the femtosecond regime.

Fig. 1 .

 1 Fig. 1. Schematic drawing of the three experimental configurations (a), (b) and (c), using the same LGT crystal cut at (𝜃 𝑃𝑀 = 58.3°, 𝜑 𝑃𝑀 = 90°).

Fig. 3 .

 3 Fig. 3. Experimental data of the idler (red) and signal (green), using a narrowband pump 𝜆 𝑝 = 1.03 µm (blue) and a broadband signal 𝜆 𝑠 = 1.5 µm (green) as incoming beams. a) Incoming and idler generated normalized intensities, and b) internal angles are plotted as a function of the wavelength in LGT slab. Lines are calculations after ref. [13]. Figs. 2b, 3b and 4b show that the same LGT slab cut at 𝜃 PM = 58.3° is able to generate or amplify wavelengths between 1.4 and 4.7 µm by tilting the crystal in the 𝑦-𝑧 plane over 21° with different pump sources. Depending on the pump source, a continuous tunability is demonstrated from 2.1 to 2.6 µm, 1.4 to 3.6 µm, and 3.4 to 4.7 µm.Calculations based on ref.[13] agree very well with our experimental data within the precision of our measurements in the three experimental configurations. For the sake of clarity, in the case of broadband pump and signal (configuration (c)), we have calculated the phase-matching curves (see Fig.4b) corresponding to three pump wavelengths: 𝜆 𝑝2 = 1.32 µm (maximum of the pump spectrum), 𝜆 𝑝1 = 1.18 µm, and 𝜆 𝑝3 = 1.41 µm (wavelengths corresponding to 50 % of pump maximum spectral intensity). All recorded data fit within the internal phase-matching curves. Note that the measured spectral bandwidth of the idler waves matches with the simulated curve after considering all the incoming experimental parameters, i.e. pulse duration, beam sizes, pulse energy, and sample characteristics. Thinner crystal would certainly provide broader spectra at all wavelengths. In Fig.5we compare the spectral range of the generated wavelengths in LGT with that obtained in -BaB2O4 (BBO), KTiOAsO4 (KTA) and Periodically-Poled LiNbO3 (PPLN) from refs.[1][2][3][4][5][6][7][8][9][10][11]. The transparency range of each crystal is shown with the dashed white bars and the measured spectral range of the

Fig. 4 .

 4 Fig. 4. Experimental data of the idler (red) using a broadband pump 𝜆 𝑝 = 1.3 µm (blue) and a broadband signal 𝜆 𝑠 = 1.9 µm (green) as incoming beams. a) Incoming and idler normalized intensities and b) internal angles plotted as a function of the wavelength in LGT slab.Lines are calculations after ref.[13] for three selected quasimonochromatic pump wavelengths 𝜆 𝑝1 = 1.18 µm, 𝜆 𝑝2 = 1.32 µm and 𝜆 𝑝3 = 1.41 µm from the pump spectral width.

Fig. 5 .

 5 Fig. 5. Transparency ranges (dashed white bars) and measured spectral ranges generated in BBO, KTA, PPLN and LGT when pumped around 0.8 µm (blue bars) and around 1 µm (red bars).

Fig. 6

 6 plots the damage fluence as a function of the nonlinear parameter for widespread nonlinear crystals in the nanosecond regime. LGT is represented by the plain (ns) and light (fs) red disks. The different chemical families of nonlinear crystals are distinguished by different colors.

Fig. 6 .

 6 Fig. 6. Damage threshold as a function of the highest nonlinear coefficient 𝑑 𝑖𝑗 for Alkali Metal Oxides (AMO3), Metal Semiconductor Diphosphide (MSP2), II-III/V-VI = II/VI and III/V semiconductors, Li III-IV = LiGaS2, LiInS2 and LiInSe2 crystals in the nanosecond regime.

Acknowledgments. The authors thank Raphaël Clady and Olivier

Uteza for their work on the experimental setup, and Jérôme Debray for the orientation and cutting of LGT slabs.

Funding. This project received funding from the ANR (ANR-10-IDEX-03-02), the ANR-ASTRID (ANR-15-ASTR-0005), ELI-ALPS GOP-1.1.1-12/B-2012-0001, Laserlab-Europe EU-H2020 654148, FEMTO and CMDO+ CNRS Networks.

Full references:


Driss Kenfaui

Sophie Guillemet-Fritsch

Zarel Valdez-Nava

Lionel Laudebat

Christophe Tenailleau

Pascal Dufour

Marie-Laure Locatelli

Tuning electrical conductivity in AlN-based ceramics by incorporating graphene therein

Aluminum nitride (AlN) is the material of choice for high power modulus substrates. However, due to the strong covalent bonds in the AlN system and the low diffusion coefficient of its constituent elements, additives have to be used as sintering aids, typically Y2O3, to achieve nearly full material density. The thermal and electrical conductions can be also fostered by adding CaF2 and carbon particles, respectively. In this paper, we report on tuning the electrical conduction within AlN-based composites developped by incorporating graphene nanoplatelets (GNP) into the AlN-based ceramics comprising the Y2O3 and CaF2 additives. GNP were produced by exfoliation of graphite in isopropyl alcohol. These composites were prepared by Spark Plasma Sintering (SPS) through a new configuration called "multiple preparation" with a view to fabricating three samples per SPS cycle, revealing potential time and energy savings. A magnitude of 10 order of magnitude higher as compared to the commercial AlN was obtained for these composites by adjusting the amount of GNP as well as the way of exfoliation to produce them. This work paves the way for the development of novel multifunctional ceramic composites for Power Electronics applications.

Introduction

In Power Electronics (PE), the substrate has to perform concomitantly the dual functions of electrical insulation and thermal conduction. It also ensures the mechanical reliability of the arrangement of heterogeneous materials constituting the power module. Owing to their physical and mechanical properties, the Al2O3, AlN and Si3N4 ceramics are the substrates that are capable of fulfilling these functions in a PE module.

However, only Al2O3 and AlN ceramics are selected for high power applications because of their higher breakdown voltage and thermal conductivity, [START_REF] Chasserio | Ceramic substrates for hightemperature electronic integration[END_REF] with a preference for the AIN material which displays the best compromise of both properties. Si3N4 ceramics are less suitable due to higher manufacturing cost.

These ceramics currently fulfill the multiple functions of the substrate in the traditional Sibased PE systems. They will nonetheless undergo stronger electrical field in new SiC-based systems operating under higher voltages. Their functional and mechanical properties must therefore be controlled and/or improved locally and/or in terms of volume in order to cope with more severe stresses arising both from operation (a rise in voltage and temperature) and integration. In other words, a new generation of substrates should be developed to take advantage of the high potential offered by large gap semiconductors. Thus, AlN ceramics turn out to be the best-suited existing substrates to face those challenges 2- 4 due to their excellent aforesaid attributes. [START_REF] Xiong | Microstructural effects on the transmittance of translucent AlN ceramics by SPS[END_REF][START_REF] Song | Heating mechanism of spark plasma sintering[END_REF] The production of dense AlN ceramics by conventional sintering requires however high temperature levels associated with long dwell time. Indeed, these materials weakly densify under the sole action of heat due to the strong covalent bonds in the AlN system and the low diffusion coefficient of their constituent elements. [START_REF] Yu | Microstructural characterization and microstructural effects on the thermal conductivity of AlN (Y2O3) ceramics[END_REF][START_REF] González | Development of AlN and TiB2 composites with Nb2O5, Y2O3 and ZrO2 as sintering aids[END_REF][START_REF] He | Study on microstructure and thermal conductivity of spark plasma sintering AlN ceramics[END_REF] Long sintering time and high temperature levels not only cause an increase in the cost of production but also promote the grain growth of which leads to the degradation of the mechanical properties of the sintered ceramics. Spark Plasma Sintering (SPS), well-known for its ability to accelerate the kinetics of sintering, [START_REF] Xiong | Microstructural effects on the transmittance of translucent AlN ceramics by SPS[END_REF] is today the most widely used method for producing highly dense AlN ceramics in very short times. [START_REF] Yu | Microstructural characterization and microstructural effects on the thermal conductivity of AlN (Y2O3) ceramics[END_REF][START_REF] González | Development of AlN and TiB2 composites with Nb2O5, Y2O3 and ZrO2 as sintering aids[END_REF][START_REF] Song | Heating mechanism of spark plasma sintering[END_REF] Nevertheless, the addition of a limited fraction of sintering additives (rare earths and/or alkaline earth oxides) is essential to promote the densification of AlN materials by creating a liquid phase, [START_REF] González | Development of AlN and TiB2 composites with Nb2O5, Y2O3 and ZrO2 as sintering aids[END_REF][START_REF] He | Study on microstructure and thermal conductivity of spark plasma sintering AlN ceramics[END_REF][START_REF] Song | Heating mechanism of spark plasma sintering[END_REF] and to improve their thermal conductivity by reducing the oxygen impurities present therein. [START_REF] Qiao | Effect of Y2O3 on low temperature sintering and thermal conductivity of AlN ceramics[END_REF] AlN particles are usually covered with a thin layer of Al2O3 oxide (1-2%) which reacts with the sintering additives (e.g. Yttrium oxide Y2O3) and AlN powder. This results in the formation of oxides (yttrium aluminate garnet Y3Al5O12 (YAG), yttrium-aluminum perovskite YAlO3 (YAP), monoclinic yttrium-aluminum Y4Al2O9 (YAM), etc.), which fosters sintering in liquid phase of the AlN material during the sintering process.

A lot of work has therefore been reported on the heterogeneous ceramics (AlN + Y2O3), developed by SPS method, exhibiting a strongly densified microstructure with compounds (YAG, YAP, YAM and/or other oxides) concentrated at the grain boundaries. He et al. [START_REF] He | Study on microstructure and thermal conductivity of spark plasma sintering AlN ceramics[END_REF] reported almost complete densification (> 99%) of AlN + 3 wt% Y2O3 ceramics sintered by SPS under a controlled nitrogen atmosphere at 1800°C under an uniaxial pressure of 40 MPa for 10 min.

These materials present a microstructure in which the oxide YAlO3 is incorporated into the grain boundaries of the AlN matrix. However, this secondary phase appears thick and partially covers the AlN grains, which is prohibitive for the thermal conductivity of these materials due to the diffusion of phonons at the AlN / YAlO3 interfaces.

It has also been reported that the addition of calcium fluoride CaF2 (melting temperature of 1418°C) to the AlN + 3 wt% Y2O3 mixture leads to the formation of an additional oxide Ca3Al2O6 and other products which evaporate during SPS sintering, reducing then the amount of secondary phases at the grain boundaries. [START_REF] He | Study on microstructure and thermal conductivity of spark plasma sintering AlN ceramics[END_REF] The reduction in the Y2O3 fraction in favor of CaF2 lead to a microstructural with the secondary phases occupying only the junctions and concentrating at the triple points between the grains, thereby enhancing the thermal conduction in AlN ceramics. Moreover, Nakano et al. [START_REF] Nakano | Grain boundary phase in AlN ceramics fired under reducing N2 atmosphere with carbon[END_REF] reported that the thermal conduction in an AlN ceramic depends both on the different constituent phases, and on the distribution of secondary phases within the microstructure. Hence, the overall thermal conductivity of the heterogeneous ceramic could be increased if the microstructure includes reduced amount of secondary phases located preferentially at the triple points between the grains and evenly distributed in the AlN matrix. We therefore expect that the thermal conductivity of AlN + 1% m Y2O3 + 2% m CaF2 ceramics will be higher as compared to the AlN + 3% m Y2O3 and AlN ones.

On the other hand, little work has been reported on the electrical properties of AlN-based ceramics, especially the electrical conductivity (s). Indeed, it is only in the last few years that the graphene nanoplatelets (GNP) have been incorporated in the AlN microstructure in order to amend its s. [START_REF] Simsek | The effect of graphene nanoplatelets on the thermal and electrical properties of aluminum nitride ceramics[END_REF][START_REF] Yun | Mechanical, electrical, and thermal properties of graphene nanosheet/aluminum nitride composites[END_REF] The choice of this additive is driven by the advantages of a large specific surface area and high charge carrier mobility of the graphene sheets, but also by its exceptional mechanical properties [START_REF] Lee | Measurement of the elastic properties and intrinsic strength of monolayer graphene[END_REF] which give GNP a high resistance to irreversible damage at high temperature under mechanical stress such as during SPS processing, thereby distinguishing them from carbon nanotubes (CNT). [START_REF] Fan | Preparation and electrical properties of graphene nanosheet/Al2O3 composites[END_REF][START_REF] Ramirez | Graphene nanoplatelet/silicon nitride composites with high electrical conductivity[END_REF] CNT have indeed been reported to be reactive with oxides and nitride-based materials at high temperature, making it difficult to maintain their structure during SPS sintering. [START_REF] Kusunose | Production of a grain boundary phase as conducting pathway in insulating AlN ceramics[END_REF] GNP also exhibit thermal conductivity (k~ 6-10 Wm -1 K -1 ) higher with two orders of magnitude compared to reduced graphene oxide (rGO). [START_REF] Román-Manso | Thermal conductivity of silicon carbide composites with highly oriented graphene nanoplatelets[END_REF][START_REF] Renteria | Strongly anisotropic thermal conductivity of free-standing reduced graphene oxide films annealed at high temperature[END_REF] Note that rGO is obtained by the reduction of the graphene oxide (GO) during sintering of AlN. [START_REF] Xia | Mechanical and thermal properties of reduced graphene oxide reinforced aluminum nitride ceramic composites[END_REF] The most frequently techniques used for producing GNP are: [START_REF] Reina | Layer area, few-layer graphene films on arbitrary substrates by chemical vapor deposition[END_REF][START_REF] Guermoune | Chemical vapor deposition synthesis of graphene on copper with methanol, ethanol, and propanol precursors[END_REF][START_REF] Coleman | Liquid-Phase Exfoliation of Nanotubes and Graphene[END_REF] i) Physical exfoliation of graphite using adhesive tape, which leads to the production of graphene in small area, but with the presence of residues duct tape on it.

ii) Epitaxy of graphene from silicon carbide (SiC), or the oriented growth, that requires the use of a high vacuum, a source of SiC as well as very high heating temperatures (~ 1300°C) to evaporate the silicon atoms, then the carbon atoms will rearrange themselves to form layers of graphene.

iii) Heterogeneous catalysis using chemical vapor deposition (CVD) which allows forforming graphene on a large surface, on copper or nickel supports. This method also involves the use of high temperatures and low pressures, hence a high cost. iv) Exfoliation of graphite in liquid phase. [START_REF] Coleman | Liquid-Phase Exfoliation of Nanotubes and Graphene[END_REF] Simsek et al. [START_REF] Simsek | The effect of graphene nanoplatelets on the thermal and electrical properties of aluminum nitride ceramics[END_REF] employed the last technique (iv) to produce GNP by exfoliating the graphite in isopropyl alcohol using the conventional ultrasound. The obtained GNP were then incorporated within based-AlN matrix to elaborate ceramic composites (AlN + 2.9 wt% Y2O3 + x vol% GNP) by SPS at 1700 -1750°C under a uni-axial pressure of 50 MPa for 5 min under a controlled nitrogen atmosphere. An almost full composite density (> 99%) was achieved for the volume fractions x ≤ 5 vol% of GNP, but it dropped slightly to 97.8% for x = 10 vol%. The evolution of s with x revealed that s can be increased from s = 10 -14 S.cm -1 for a ceramic with no GNP up to 1 S.cm -1 for the composite containing 10 vol% of GNP.

In the light of the state of the art drawn up above, we decided to develop the ceramic composites AlN + 1 wt% Y2O3 + 2 wt% CaF2 containing a GNP fraction of 2.5 and 10 vol% by SPS. The GNP were produced by exfoliation of graphite in isopropyl alcohol using two exfoliation methods: conventional sonication (ultrasonic bath) and sonotrode. These composites were fabricated under the same SPS conditions in order to assess the contribution of the selected additives. They were particulary elaborated using "multiple preparation configuration" with the aim of fabricating three samples per SPS cycle. It enabled to have the number of samples needed to perform different material characterizations and to assess the degree of reproducibility. Their microstrcutres features and electrical conductivity were compared with those of the ceramics AlN, AlN + 3 wt% Y2O3 and AlN + 1 wt% Y2O3 + 2 wt% CaF2. Note that the reference and the description of the samples are summarized in Table 1.

Experimental section

Powder preparation and shaping

The AlN powder (Atochem, France), with a median grain size d50 = 1.9 μm, and the sintering additives Y2O3 (purity: 99.9%, Shinetsu, Japan) and CaF2 (purity: 99.7%, Alfa Aesar, Germany) were selected as the starting precurssors. The respective fractions of Y2O3 and CaF2 were set at 1 and 2 wt% with a goal of reaching nearly full densification concomitantly with improving the thermal conductivity of AlN-based materials.

To tune the electrical conduction, GNP were chosen for their aforementioned advantages.

Two methods of exfoliation were tested to produce them: ultra sonic bath and sonotrode.

The first batch of GNP were produced by exfoliation of 2.5 vol.% of graphite (median lateral particle size: 15 µm, Skyspring Nanomaterials, Inc. -USA) in 20 ml of isopropyl alcohol using a classical ultra sonic bath. The choice of this liquid dispersant is governed by its evaporation temperature (82.6°C) which is much lower than that of (NMP) N-methyl-2-pyrrolidone (202°C) or (DMF) N, N-dimethylformamide (153°C), solvents known for their greater ability to stabilize the dispersion of graphene sheets produced by exfoliation of graphite. [START_REF] Coleman | Liquid-Phase Exfoliation of Nanotubes and Graphene[END_REF] Indeed, such a low evaporation temperature is required to be able to separate more easily the solvent from the final powder mixture. This separation is carried out by evaporating the solvent using the rotary evaporator as describted later in this section.

However, the GNP, produced by the conventional sonication (ultrasonic bath), are still stacks comprising a very high number of graphene sheets. However, the effects of graphene at the electronic, and therefore electric, level are more prominent in the quantum domain, that is, when the thickness of the graphene layers is close to the mean free path of the electron. In practice, below ten layers of graphene, the electrical effects significantly increase.

We therefore decided to use the sonotrode, a probe sonicator (Sonics Materials VCX-750, Fisher Scientific, France) fitted with a high-frequency ultrasonic generator, to sonicate a similar liquid solution comprising 2.5 vol.% of graphite in 20 ml of Isopropylic alcohol. The sonicator was set to a power of 150 W with an amplitude of 40% and a pulsed mode of the cycle 4s ON (sonication) / 6s OFF (pause). The aim was to provide the required energy to exceed that of Van der Waals type bonding between the graphene sheets for the purpose of exfoliating the graphite. The pause time allows to limit the rise in temperature of the solution and therefore to slow its evaporation . The solution volume was then kept constant throughout the sonication time.

The solution was sonicated for 2 h; the total time of the sonication experiment was 5 h taking into account the total pause time. It was thereafter left to settle with a view to evidencing the graphite exfoliation. Figure 1 shows an image taken 15 h after the solution sample was sonicated. The volume Va (black part of the container) is visibly higher than that of the solution before sonication, suggesting exfoliation of the graphite. This result was also obtained for two similar solutions previously subjected to effective sonication times of 0.5 and 1 h. The apparent volumes Va, obtained after these solutions were left to settle for 15 h, were compared. The longer the sonication time, the better the exfoliation. The solution, with the highest volume, was thereafter subjected to 3 additional sonication tests of 2 h each. The volume Va continued to increase indicating a more exfoliation of the graphite.

On the other hand, the powder mixture AlN + x1 wt% Y2O3 + x2 wt% CaF2 was homogenized in absolute alcohol using a mechanical mixer adjusted to a rotation of 200 rpm for 1 h. The powder was then heated at 80°C for 15 h to remove the solvent.

Then, this powder was introduced with GNP dispersed in isopropyl alcohol into a glass flask fixed in a rotary evaporator. The flask was subjected to a rotation of 70 rpm for 2 h to homogenize the resulting solution. It was heated up to 70°C under vacuum to reduce the partial pressure of the solvent. The produced vapor was condensed after being pipped to the cooling system, and the liquid solvent was finally conveyed to a recovery tank.

The resulted powder mixure was shaped in the SPS apparatus (Dr. Sinter 2080 device from Sumitomo Coal Mining -Fuji Electronic Industrial, Saitama, Japan). It was first loaded into a graphite mold of 20 mm in diameter. A graphite sheet was previously placed on the internal wall of the mold in order to avoid any reaction between it and the material, on the one hand, and to facilitate the demolding of the sample after sintering, on the other hand. The mold was then transferred to the SPS and heated up to 1800°C for 10 min, at a rate of 100°C/min, under uniaxial pressure of 50 MPa and nitrogen atmosphere. Such experimental conditions for SPS shaping were seleced from the literature [START_REF] He | Study on microstructure and thermal conductivity of spark plasma sintering AlN ceramics[END_REF][START_REF] Simsek | The effect of graphene nanoplatelets on the thermal and electrical properties of aluminum nitride ceramics[END_REF] with the aim of achieving complete densification and limiting grain growth.

It is worth noting here that in addition to the simple preparation, i.e. one pellet per SPS cycle (Figure 2a), we also performed the multiple preparation of samples that involves the elaboration of three pellets per SPS cycle. The required quantity of powder for each pellet was placed between two graphite discs in the matrix as illustrated in Figure 2b, which allows then the preparation of three samples by applying only a single SPS cycle. Figures 2c-d show, for instance, the microstructures of the AY samples elaborated using simple preparation and multiple one at top, middle and bottom positoins. They evidence almost similar microstrctural features with no significant difference between the samples. Note that the AY microstructure is described later in detail ( § 3.2.1).

The multiple preparation, proposed for the first time to our knowledge, has the merits of producing, in a very short time, the number of samples required to carry out different material characterizations, on the one hand, and to evaluate the degree of reproducibility of the results obtained, on the other hand. This testifies to the relevance of such approach in terms of not only the higher production of the sintered material but also the triple saving of time and energy.

Characterization techniques

After being sonicated for 8 h as mentioned above, the solution (2.5 vol.% of graphite in 10 ml of isopropyl alcohol) was sonicated for an additional time of 1 h. Immediately after this operation, a few drops loaded with exfoliated graphite were poured onto a perforated carbon grid. It was then transferred into the transmission electron microscope (TEM) for observation (HT-7700 Hitachi 120 kV). The accelerating electron voltage was set at 80 kV to reduce the damage that could be induced by the electron beam on the graphene while maintaining the image resolution of the particle fringes in order to distinguish the number of platelets and the local thickness on the folded edges.

The microstructure of the elaborated AlN-based materials was examined by means of the Scanning Electron Mircoscope (SEM JEOL JSM6510LV). Their chemical composition was determined by using X-ray diffractometer (Bruker D4).

The dielectric properties were probed at 20°C under low voltage (3 VPP) in the 10 -1 -10 6 Hz frequency range for the prepared samples by the mean of a broadband impedance analyzer (Novocontrol Alpha-A). The SPS-sintered pellets of 20 mm in diameter were initially mirrorpolished and annealed at 150°C during 5 h to remove all traces of water. A 80 nm-thick gold electrode was then deposited by sputtering on both faces of each sample. The metalized pellet was finally placed into the measuring cell between two conducting electrodes to form a parallel capacitor.

Results and discussion

Observation of GNP

TEM characterization, performed for the solution sonicated for 9 h, revealed the presence of a large number of two-dimensional and quasi-transparent platelets (Figure 3a) depicting nano-scale thicknesses and lateral sizes in the range of 600 -2000 nm, similar to those reported elsewhere [START_REF] Coleman | Liquid-Phase Exfoliation of Nanotubes and Graphene[END_REF] for monolayers of graphene obtained by exfoliation of graphite in the solvent NMP (N-methyl-2-pyrrolidone). Generally, the observed GNP have straight edges and corners with distnguished angles allowing to perceive their nano-scale thickness, and even to forecast the very small number of monolayers. TEM images, presented in Figure 3a andb, are similar to those of graphene monolayers reported elsewhere. [START_REF] Coleman | Liquid-Phase Exfoliation of Nanotubes and Graphene[END_REF][START_REF] Paton | Scalable production of large quantities of defect-free few-layer graphene by shear exfoliation in liquids[END_REF][22] We have also noted the presence of less transparent and thicker platelets as shown in Figure 3b. These platelets therefore include a higher number of monolayers.

Microstrcture and chemical compositon of the SPS-sintered materials

Ceramics A, AY and AYC

Figure 4a depicts the SEM image of the cross section (fractured surface) obtained by cleavage of the AlN material (sample A) produced by SPS under the aforementioned conditions. It shows a homogeneous microstructure with a density higher than 98% refelecting the relevance of the selected sintering conditions. Such microstructure comprises micro-sized grains with no significant growth occurred during the sintering, due to the very short time of the SPS processing, and no secondary phase that can be visibly distinguished at the grain boundaries. However, X-ray diffraction analyzes (Figure 4b) revealed the presence of traces of Al9O3N7, Al8O3N6 and Al10O8N3 oxides, very likely due to the reaction of oxygen impurities with AlN during sintering.

The SEM observation of the AY ceramic (Figure 4c) shows the typical microstructure of AlN materials produced by SPS with sintering additives: the secondary phases (clear contrast) are located at the grain boundaries of the AlN matrix (dark contrast). A highly dense microstructure exceeding 99% with very limited grain growth is noted as well. Unlike the sample A (Figure 4a), the sintering is carried out in the liquid phase formed following the reaction of Y2O3 with AlN and Al2O3 present on the surface of the grains. The secondary phase, produced in liquid form during sintering, therefore flows at the grain boundaries and solidifies during cooling, thus contributing to densifying the material. Chemical analyzes indicate that its composition is close to that of YAP as shown by the XRD analysis given in Figure 4d.

However, the addition of 3 wt% Y2O3 led to the formation of a relatively high amount of YAP which partially covers the AlN grains. This might promote phonon scattering at the YAP / AlN interfaces and consequently degrade the thermal conduction in the material. These results are in agreement with those reported by He et al. [START_REF] He | Study on microstructure and thermal conductivity of spark plasma sintering AlN ceramics[END_REF] The micrograph of the SPS-sintered AYC ceramic is presented in Figure 4e. It evidences a homogeneous microstructure with hexagonal grains having rounded edges. The density of this material is comparable to that noted for the AY one. Nevertheless, the AYC ceramic includes much smaller amount of secondary phases at the grain boundaries, present along the junction and concentrated at the triple points between the grains. This is attributed not only to the lower fraction of the added Y2O3 but also to the formation of secondary phases which evaporate during sintering as evidenced by the traces in clear contrast observed on the 'spacers' at the end of the SPS process. The X-Ray analysis shows the presence of YAP as well, coexisting however with an additional oxide Ca3Al2O6. This based-AlN matrix was used for fabricating the composites presented later. Its recorded X-Ray diffractogram is then similar to that given in Figure 6c.

The reduction of secondary phases within this matrix can be explained as flowes: since the decomposition temperature of CaF2 (1418°C) is lower than the used dwell temperature (1800°C), the following chemical reactions took place during the treatment of the AlN + 1 wt%Y2O3 + 2 wt%CaF2 powder mixure by SPS: As the evaporation temperature of the AlF3, Ca3N2 and Ca11N8 products is lower than 1800°C, they therefore evaporated during the SPS sintering.

Composite AlN + 1 wt% Y2O3 + 2 wt% CaF2 + 2.5 vol.% GNP (AYC2.5G)

Figure 5 shows the SEM images of the fractured surfaces obtained by cleavage of the composite AYC2.5G elaborated by SPS. The GNP incorporated in this material were previously produced by exfoliation of graphite in isopropyl alcohol using conventional sonication (ultrasonic bath) for 1 h as reported by Simsek et al. [START_REF] Simsek | The effect of graphene nanoplatelets on the thermal and electrical properties of aluminum nitride ceramics[END_REF] The microstructure seems to be not disturbed by the presence of GNP (Figure 5a) indicated by yellow arrows. It displays a density (99%) as high as that recorded for ceramics free of GNP (Figures 4), regardless of the used configuration (simple or multiple) shown in Figure 2. GNP appear to be distributed throughout the AlN matrix. However, we cannot confirm the uniformity of this distribution from SEM observations because graphene monolayers, or GNP including a very small number of graphene sheets, are very difficult to observe using SEM. High-magnification microstructural examination reveals that the in-plane of the GNP platelets tends to be perpendicular to the axis of the SPS pressing axis (Figure 5b). We also observe that GNP are inserted between the grains and follow the contours of the latter, producing ripples without fracture, thus testifying of the high mechanical resistance of graphene at high temperature under loading. We note, however, the presence of several thick GNP stacks containing a very large number of graphene layers (Figure 5c), which indicates the insufficiency of conventional sonication to produce GNP comprising a reduced number of layers. As in the case of the AYC ceramic (Figure 4e), the secondary phases, produced in reduced quantity, are concentrated at the triple points between the hexagonal grains of the AlN. These phases, illustrated in clear contrast, are indicated by red arrows (Figure 5b). The results of the local EDS microanalyses, given in the inset of this Figure, indicated the presence of the elements Y, Al, O and Ca, supporting then those of the characterization by X-Ray diffraction which revealed the presence of YAP and Ca3Al2O6.

Composite AlN + 1 wt%Y2O3 + 2 wt%CaF2 + 2.5 vol.% GNP (AYC2.5GSO)

Figure 6a shows the fractured surface of the ceramic composite AYC2.5GSO fabricated by SPS.

In this case, the incorporated GNP were obtained by exfoliation of graphite in isopropyl alcohol after sonication for 9 h using the sonotrode. Unlike the previous composite sample (Figure 5a), the microstructure is nearly full densified (99%) whatever the used configuration (Figure 2), but revealed little GNP, which bodes well for the presence of graphene monolayers and GNP containing a very small number of layers which are difficult to observe using SEM. This result is consistent with the TEM observations and highlights much better exfoliation of graphite by sonotrode. Examination at high magnification of the observed GNP (Figure 6b) also showed that they include a smaller number of layers as compared with the previous ceramic (Figure 5c). They are introduced between the grains and also follow the contours of the latter by exposing waves without fracture.

The local EDS microanalyses combined with the characterization by X-ray diffraction (Figure 6c) showed the presence of the oxides YAP and Ca3Al2O6 concentrated at the grain boundaries (Inset of Figure 6a).

Electrical conductivity of the elaborated AIN-based materials

Figure 7a presents a comparaison of the frequency dependence of the electrical conductivities s (f) measured at 20°C for the AlN material prepared by SPS (sample A) and for the commercial ceramic from Impak company. The commercial sample appears to be more resistive, as is evident from the s shift towards lower values in the 0.5 -5 x 10 5 Hz frequency range. At 0.1 Hz, s is lower with an order of magnitude as compared to the A ceramic prepared by SPS (Table 2). This difference is probably tied to the oxides detected therein (Figure 4a) . It underlines the necessity to incorporate sintering additives in the AIN matrix so as to remove oxygen impurities during sintering, thereby avoiding the formation of these oxides and, subsequently, decreasing s.

Figure 7b shows the s (f) curves recorded at 20°C for the composites AlN + 1 wt% Y2O3 + 2 wt% CaF2 + 2.5 vol.% GNP produced by SPS using single (AYC2.5G) and multiple (AYC2.5G-H and AYC2.5G-M) configurations. For comparison, the curve s(f) recorded for the ceramic comprising a fraction of GNP as high as 10 vol.% (AYC10G) is also plotted. The GNP incorporated in these samples were previously produced by exfoliation of graphite in isopropyl alcohol using conventional sonication (ultrasonic bath) for 1 h. The curves s (f) of the A ceramic and of the commercial one are also given. The conductivity s of the sample AYC10G is independent on the frequency (~ 6.5 x 10 -2 S.cm -1 ), which approves the choice of GNP to modify s within AlN ceramics. The samples comprising 2.5 vol.% GNP, although shaped under the same conditions in the SPS apparatus, show visibly dispersed s values and their electrical behaviors differ from each other. Indeed, the sample AYC2.5G presents a monotonic decrease in s from 1 x 10 6 to 10 -1 Hz, corresponding to a capacitive behavior that is similar to that noted for the commercial AIN material and for the A ceramic (Figure 7a). The sample AYC2.5G-Mexhibits a similar behavior in the 2 x 10 3 -10 6 Hz range, and its s clearly becomes less frequency dependent in the 0.1 -2x 10 3 Hz range by showing a plateau, characteristic of a resistive behavior. It is also observed for the sample AYC2.5G-H over the entire explored frequency range. At 0.1 Hz, the s values recorded for these last three samples are 1.7 x 10 -12 ; 2 x 10 -9 and 1.5 x 10 -6 S.cm -1 (Table 2), respectively. Such dispersion, in terms of s as well as of the electrical behaviors, evidences that the electrical effect of GNP is different within the last three samples. This can be attributed to either a non-uniform distribution of GNP in the prepared mixture AlN + 1 wt% Y2O3 + 2 wt% CaF2 + 2.5 vol% GNP, or the difference in the quality of the exfoliated GNP contained in each sample.

The curve s(f) recorded at 20°C for the composites AYC2.5SO produced by SPS using simple (AYC2.5GSO) and multiple (AYC2.5GSO-H, AYC2.5GSO-M and AYC2.5GSO-L) configurations are given in Figure 7c. GNP incorporated in these samples were obtained by exfoliation of graphite in isopropyl alcohol upon sonication for 9 h using the sonotrode.

The curves s (f) of the A ceramic and of the commercial sample are also plotted for comparison. We note that the composites show frequency-insensitive electrical conductivity, indicating resistive behavior over the entire explored frequency range. The samples, prepared following multiple configuration, display almost identical values, revealing a remarkably high degree of reproducibility. They indicate a s value close to 10 -4 S.cm -1 , emphazing an increase of approximately 10 orders of magnitude as compared to the commercial sample and of the AlN ceramic produced by SPS. In addition, this value is markedly higher than the s values noted for the previous ceramics (Figure 7b), comprising the same fraction of GNP obtained by conventional sonication of graphite. This increase in s is very likely related to the presence of a high number of graphene monolayers with a higher mobility of charge carriers, which induces more marked electrical effects in the AlN ceramic. It may also result from a better distribution of GNP in the composites. The AYC2.5GSO sample, produced by single configuration, exhibits similar electrical behavior and a s value of 10 -5 S.cm -1 , close to that of the AYC2.5G-H, AYC2.5G-M, AYC2.5G-L (low position) samples. These results are consistent with the MET and SEM observations, and present, to our knowledge, a degree of reproducibility never reported so far for AlN-based composites containing conductive additives such as GNP. However, the values obtained (10 -4 -10 -5 S.cm -1 ) are much higher than the value reported by Simsek et al. [START_REF] Simsek | The effect of graphene nanoplatelets on the thermal and electrical properties of aluminum nitride ceramics[END_REF] for the same fraction of GNP. This difference can be ascribed to better exfoliation of the graphite by means of the sonotrode under the optimized conditions, and probably by an improved distribution of graphene in our elaborated composites as well.

Figure 7d shows the s(f) curves recorded at different temperatures (20-195°C) for the composite AYC2.5GSO. Whatever the temperature level, the s of this composite does not depend on the frequency, revealing then a resistive behavior over the whole explored frequency (10 -1 -10 6 Hz) and temperature (20-195°C) ranges. s increases as the temperature raised and reaches 2.82 S.cm -1 at 195°C. Finally, since the GNP particles are two-dimensional, s was otherwise monitored on a bar cut from the SPS-sintered composite AYC2.5GSO in the directions parallel (σ ∥ ) and perpendicular (σ⊥) to the applied pressing axis (Inset of Figure 7e). Figure 7e presents the recorded conductivities σ ∥ and σ⊥ versus frequency at 20°C for this composite. It evidences that the GNP entailed an anisotropy therein with a ratio (σ⊥ / σ ∥ ) = 4.4, stemming from the fact that they were oriented, under the effect of uniaxial pressure during the SPS sintering, in the direction perpendicular to the pressing axis. Indeed, since GNP particles possess the superior σ in their in-planes over that of their out-planes, they induce σ⊥ larger than σ∥ within the composite AYC2.5GSO.

Conclusions

Graphene nanoplatelets were incorporated in AlN-based ceramics comprising sintering additives of Y2O3 and CaF2. SPS process combined to the multiple preparation configuration allowed to produce full densified (>99%) ceramic composites with three samples per SPS cycle, enabling the assemenet of reproductibility and significant time and energy savings.

Their electrical conductivity was demonstrated to be tuned by acting on the exfoliation process. More efficient exfoliation leads to increased number of graphene monolayers with higher charge carrier mobility. The electrical conductivity reaches 10 -4 S.cm -1 for composite containing 2.5 vol.% NGP, i.e. 10 order of magnitude higher as compared to the commercial AlN ceramics. The fine tuning of the electrical conductivity through graphene incorporation is suitted for many other systems. is plotted. GNP incorporated in these samples were produced by exfoliation of graphite in isopropyl alcohol using conventional sonication for 1 h. (c) Composites AlN + 1 wt% Y2O3 + 2 wt% CaF2 + 2.5 vol.% GNP produced by SPS using simple (AYC2.5GSO) and multiple (AYC2.5GSO-H, AYC2.5GSO-M, AYC2.5GSO-L) configurations. GNP incorporated in these samples were obtained by exfoliation of graphite in isopropyl alcohol upon sonication for 9 h using the sonotrode. Note that for comparison, the curves s (f) recorded for the A ceramic and the commercial one (from Impak company) are also plotted in Figure 7b andc

Figure caption

Figure 1 .

 1 Figure 1. (a) The prepared solution, containing 2.5 vol.% of graphite powder in 20 ml of isopropyl alcohol, before sonication. (b) Images taken immediately after the sonication of this solution for 2 h by using the sonotrode and (c) after the sonicated solution was left to settle for 15 h. The volume Va (black part of the container) is visibly higher than that of the solution before sonication, indicating the exfoliation of the graphite.

Figure 2 .

 2 Figure 2. Schematics illustrating the two SPS configurations used for elaborating AlN-based ceramics: (a) single and (b) multiple preparations. Images of the corresponding single pellet and the stack of three pellets upon SPS treatment are given for each configuration. SEM images of the fracture surfaces obtained by cleavage of the AlN + 3 wt% Y2O3 (AY) ceramics, produced in nitrogen atmosphere by SPS at 1800°C under 50 MPa for 10 min, using (c) simple preparation and multiple preparation configuration in (d) top, (e) middle and (f) bottom positions.

Figure 3 .

 3 Figure 3. TEM micrographs of GNP deposited over the TEM grid immediately upon 9hsonication of a solution containing 2.5 vol.% of graphite in 20 ml of isopropyl alcohol. They show (a) quasi-transparent two-dimensional flakes, and (d) less transparent and thicker flakes comprising a higher number of single-layers.

Figure 4 .

 4 Figure 4. SEM images of the fracture surfaces obtained by cleavage of the ceramics produced in nitrogen atmosphere by SPS at 1800°C under 50 MPa for 10 min: (a) AlN (A) ceramic and (b) its corresponding XRD pattern, (c) AlN + 3 wt% Y2O3 (AY) ceramic and (d) its corresponding XRD pattern, and (e) AlN + 3 wt%Y2O3 + CaF2 (AYC) ceramic.

Figure 5 .

 5 Figure 5. SEM images of the fractured surface of the ceramic AlN + 1 wt% Y2O3 + 2 wt% CaF2 + 2.5 vol.% GNP (AYC2.5G) produced by SPS, with the GNP produced using the conventional sonication (ultrasonic bath). (a) An overall view showing the distribution of GNP

Figure 6 .

 6 Figure 6. (a) SEM image of the fracture surface of the composite AlN + 1 wt% Y2O3 + 2 wt% CaF2 + 2.5 vol.% GNP (AYC2.5GSO) produced by SPS, with the GNP produced using the the sonotrode. Little GNP observed in the microstructure suggesting the presence of graphene monolayers and GNP containing a very small number of layers invisible by using SEM. Inset: secondary phases YAP and Ca3Al2O6 concentrated at grain boundaries. (b) Observation at higher magnification of GNP inserted between the grains of the AlN matrix. (c) The XRD pattern recorded for this composite.

Figure 7 .

 7 Figure 7. Frequency dependence of the electrical conductivity s(f) measured at 20°C for the elaborated AIN-based materials. (a) AlN ceramic fabricated by SPS (sample A) and the commercial AlN one (from Impak company). (b) Composite AlN + 1 wt% Y2O3 + 2 wt% CaF2 + 2.5 vol.% GNP (AYC2.5G) produced by SPS using simple (AYC2.5G) and multiple (AYC2.5G-H (High position) and AYC2.5G-M (middle position)) configurations. For comparison, the curves s(f) recorded for a composite comprising 10 vol.% of GNP (AYC10G)

  . (d) s (f) curves recorded at different temperature levels (20-350°C) for the composite AYC2.5GSO. (e) The electrical conductivities σ∥ and σ⊥ versus frequency measured at 20°C for the AYC2.5GSO composite in the directions parallel (σ∥) and perpendicular (σ⊥) to the applied pressing axis as illustrated in Inset.

FIGURES Figure 1

 1 FIGURES

Figure 6 26

 6 Figure 6

Figure 7

 7 Figure 7

Table caption Table 1 .

 caption1 References of samples prepared and measured in this work. -US* : Ultra Sonic bath

Table 2 .

 2 The electrical conductivity values recorded at 0.1 Hz and 20°C for the studied samples.

Table 2

 2

Acknowledgements

The authors thank the National Research Fondation (ANR, ASTRID program) for funding.

Surfactant-free single-layer graphene in water, Nat. Chem., 2017, 9, 347-352.