Leila Karam

Julien Reboul
email: julien.reboul@upmc.fr

Sandra Casale

Pascale Massiani

Nissrine El

Porous nickel-alumina derived from metal-organic framework (MIL-53): a new approach to achieve active and stable catalysts in methane dry reforming

Keywords: alumina catalysts, dry reforming of methane, metal-organic framework

Nickel nanoparticles on

remaining high specific surface may also contribute to the excellent Ni resistance to sintering and in turn to the inhibition of carbon nanotubes formation during the reaction.

Introduction

In a world where both global warming and oil resources substitution have become major environmental issues, the reaction of dry reforming of methane (DRM) is taking great attention. This reaction targets the transformation of CH 4 and CO 2 , which are two abundant greenhouse gases, into a syngas with equimolar CO and H 2 composition perfectly suited for further production of hydrocarbons by Fischer-Tropsch synthesis [START_REF] Luyben | Control of parallel dry methane and steam methane reforming processes for Fischer-Tropsch syngas[END_REF][START_REF] Shi | Reforming with methane and new route of low-temperature methanol Synthesis[END_REF][START_REF] Navas-Anguita | Simulation and life cycle assessment of synthetic fuels produced via biogas dry reforming and Fischer-Tropsch synthesis[END_REF]. In spite of many studies dedicated to the topic, DRM has not reached yet the industrial level due to the lack of active and stable catalysts being also economically viable. At present, the most effective active phases in DRM are noble metals which have the disadvantage to be rare and expensive [START_REF] Hou | Production of synthesis gas via methane reforming with CO 2 on noble metals and small amount of noble-( Rh-) promoted Ni catalysts[END_REF][START_REF] Pakhare | A review of dry (CO 2 ) reforming of methane over noble metal catalysts[END_REF][START_REF] Bradford | CO 2 Reforming of CH 4[END_REF]. Hence, developing catalysts based on cheaper and more abundant active phases is strongly desirable. Amongst the family of transition metals that represents the best alternative, nickel is the most promising one [START_REF] Rostrup-Nielsen | CO 2 -reforming of methane over transition metals[END_REF]. However, most of the existing catalysts using common oxide supports as well conventional nickel deposition routes face two important drawbacks: firstly metal sintering, which is accentuated by the high temperature at which the reaction has to be carried out due to thermodynamic reasons [START_REF] Gao | A review of recent developments in hydrogen production via biogas dry reforming[END_REF][START_REF] Abdullah | Recent advances in dry reforming of methane over Nibased catalysts[END_REF]; secondly, important formation of carbon deposits during the course of the reaction, due to the occurrence of CH 4 decomposition as side reaction [START_REF] Jang | A review on dry reforming of methane in aspect of catalytic properties[END_REF][START_REF] Park | Promotional effect of Al addition on the Co/ZrO 2 catalyst for dry reforming of CH 4[END_REF] that leads to progressive activity loss by inhibiting the active sites and to plugging of the reactor when carbon nanotubes are formed in big amount [START_REF] Karam | Advantages of mesoporous silica based catalysts in methane reforming by CO 2 from kinetic perspective[END_REF][START_REF] Zhang | Coke-resistant Ni at SiO 2 catalyst for dry reforming of methane[END_REF]. In the last decade, numerous attempts like the use of a bimetallic catalyst [START_REF] Siang | Hydrogen production from CH 4 dry reforming over bimetallic Ni-Co/Al 2 O 3 catalyst[END_REF][START_REF] Németh | Hindered methane decomposition on a cokeresistant Ni-In/SiO 2 dry reforming catalyst[END_REF][START_REF] Horlyck | Elucidating the impact of Ni and Co loading on the selectivity of bimetallic NiCo catalysts for dry reforming of methane[END_REF] or altering the type of support [START_REF] Figueredo | A comparative study of dry reforming of methane over nickel catalysts supported on perovskite-type LaAlO 3 and commercial Α-Al 2 O 3[END_REF][START_REF] Park | Effect of supports on the performance of Co-based catalysts in methane dry reforming[END_REF] have been made to overcome these barriers and design more robust catalysts.

Recently, we showed that an efficient route to limit nickel sintering consists in using supports with high specific area and organized mesopores where the nickel nanoparticles can be occluded, leading not only to their stabilization against growing but also to their protection against carbon nanotube formation that is inhibited owing to steric constraints [START_REF] Hassan | Low temperature dry reforming of methane on rhodium and cobalt based catalysts : Active phase stabilization by confinement in mesoporous SBA-15[END_REF][START_REF] Karam | Tuning the properties of nickel nanoparticles inside SBA-15 mesopores for enhanced stability in methane reforming[END_REF][START_REF] Kaydouh | Effect of the order of Ni and Ce addition in SBA-15 on the activity in dry reforming of methane[END_REF][START_REF] Kaydouh | Highly active and stable Ni/SBA-15 catalysts prepared by a "two solvents" method for dry reforming of methane[END_REF]. Enhancing the interaction between the nickel active phase and the support was also found to play a determining role, especially when using ordered mesoporous alumina with high specific surface as support [START_REF] Tao | Methane reforming with carbon dioxide over mesoporous nickel-alumina composite catalyst[END_REF][START_REF] Lv | A highly dispersed nickel supported catalyst for dry reforming of methane[END_REF]. Furthermore, increased interaction and performances were reached on this type of support by preparing the catalysts via a one-pot method consisting in introducing the nickel active phase directly during the synthesis of the alumina matrix instead of adding it by conventional post-synthesis impregnation treatment [START_REF] Jabbour | Factors affecting the long-term stability of mesoporous nickel-based catalysts in combined steam and dry reforming of methane[END_REF][START_REF] Morris | Ordered mesoporous alumina-supported metal oxides[END_REF]. With such one-pot method combined to an Evaporation-Induced Self-Assembly (EISA) route, we recently synthesized well organized mesoporous Ni/alumina materials where nickel was integrated as isolated species in the walls. This led to a high nickel dispersion after reduction as well as after subsequent catalytic test in the harsh conditions of DRM [START_REF] Jabbour | Ordered mesoporous "one-pot" synthesized Ni-Mg (Ca)-Al 2 O 3 as effective and remarkably stable catalysts for combined steam and dry reforming of methane (CSDRM)[END_REF], with mean size of the metal nanoparticles in the spent catalysts much below those commonly reported in the literature [START_REF] Benrabaa | Physico-chemical properties and syngas production via dry reforming of methane over NiAl 2 O 4 catalyst[END_REF][START_REF] Shamskar | The influence of Ni loading on the activity and coke formation of ultrasound-assisted co-precipitated Ni-Al 2 O 3 nanocatalyst in dry reforming of methane[END_REF][START_REF] Zhang | Dry reforming of methane over Ni/MgO-Al 2 O 3 catalysts prepared by two-step hydrothermal method[END_REF][START_REF] Hou | Meso-porous Ni/Mg/Al catalysts for methane reforming with CO 2[END_REF]. The benefits of intimately mixing the active metal and the support precursors at the steps of material synthesis was similarly demonstrated in the literature in works where DRM catalysts were prepared via a solid-phase crystallization route consisting in first exchanging a mineral precursor (perovskites, hydrotalcites, pyrochlores or spinels) with nickel cations, then calcining it to obtain a nickel-containing oxide solid solution and finally reducing it to give the catalyst [START_REF] Takehira | Autothermal reforming of CH 4 over supported Ni catalysts prepared from Mg-Al hydrotalcite-like anionic clay[END_REF][START_REF] He | Co-Ni catalysts derived from hydrotalcite-like materials for hydrogen production by ethanol steam reforming[END_REF]. However, although active and stable, these last catalysts were poorly porous and had low specific surfaces, which is known to be detrimental to active site accessibility and therefore to catalytic performances.

Based on these previous studies, our purpose in the present work is to combine the advantages of high porosity, high surface area and intimate mixing between the nickel and support components to synthesize highly stable and active Ni/alumina catalysts for DRM. To this end, we chose to use an Alcontaining Metal-organic framework (MOF) as sacrificial precursor to prepare the catalysts. MOFs are highly porous hybrid materials made up by the regular assembly of metal clusters connected to one another's by multifunctional organic linkers. They are characterized by extremely high specific surfaces and pore volumes that make them of great interest as host materials as was already demonstrated in many applications such as gas storage, molecular separation or sensing [START_REF] Stock | Synthesis of Metal-Organic Frameworks (MOFs): Routes to Various MOF topologies, morphologies, and composites[END_REF]. In the field of catalysis, the high accessibility of their metal nodes and their reactivity have started to be exploited as well, even if stability issues in the conditions of catalysis still often remain to be solved. Besides, the scope of MOF application was recently broadened by using these hybrid materials as sacrificial templates to synthesize either carbonaceous or purely inorganic materials with unique textural properties depending on the thermal treatment used for their transformation [START_REF] Yuan | Worm-like mesoporous carbon synthesized from metalorganic coordination polymers for supercapacitors[END_REF][START_REF] Yan | Synthesis of mesoporous and tetragonal zirconia with inherited morphology from metal-organic frameworks[END_REF][START_REF] Zhao | Isobutane dehydrogenation over the mesoporous Cr 2 O 3 /Al 2 O 3 catalysts synthesized from a metal-organic framework MIL-101[END_REF][START_REF] Lü | MOF-templated synthesis of porous Co 3 O 4 concave nanocubes with high specific surface area and their gas sensing properties[END_REF].

In this contribution, the selected MOF is an aluminum-based MIL-53(Al) composed of aluminum hydroxyl chains connected by terephthalate organic linkers and running along one-dimensional channels. It gathers the attributes required to obtain a porous nickel-alumina solid solution after its calcination: on one hand, it possesses a high porosity and metal-cluster accessibility expected to favor the homogeneous dispersion of nickel species within the oxide structure; on the other hand, it contains organics whose decomposition enables the formation of porosity in the obtained alumina [START_REF] Liu | Mesoporous alumina nanosheets and nanorolls derived from topologically identical Al-based MOFs[END_REF]. An additional benefit is the occurrence in this structure of aluminum bridging hydroxy groups lying along the inorganic chains that can interact with the nickel precursors and hence act as anchorage sites.

Moreover, MIL-53(Al) is made from the assembly of relatively inexpensive components, namely hydroxy-aluminum species and terephthalate linkers, which can be extracted from natural or recycled feedstocks [START_REF] Li | Fengting Li The Metal-Organic Framework MIL-53(Al) Constructed from Multiple Metal Sources:Alumina, Aluminum Hydroxide, and Boehmite[END_REF][START_REF] Lo | Waste polyethylene terephthalate (PET) materials as sustainable precursors for the synthesis of nanoporous MOFs[END_REF] making this MOF-derived strategy economically and ecologically viable.

The strategy followed for material synthesis and activation is depicted in Fig. 1. After synthesis and activation (solvent removal) of the parent MIL-53(Al), a cationic nickel precursor is impregnated within the pores (step a), then calcination is performed to obtain a porous Ni/Al oxide solid solution (step b) that is finally thermally treated under H 2 to form the reduced nickel nanoparticles that constitute the active phase for DRM. At all steps of the preparation, the materials are thoroughly characterized by complementary techniques (XRD, N 2 sorption, TG/MS and TEM) to understand the evolution of both the alumina-based support (texture and structure) and the supported nickel active phase (dispersion, location and state). In order to check the relevance of our strategy, and especially to evidence the benefit of occluding the nickel precursor within the MOF framework before its calcination, the performances of the obtained reduced catalyst are also compared to those of a nickel-alumina catalyst prepared by the impregnation of nickel within the pores of already calcined MIL-53(Al) as well as a more traditional nickel-alumina catalyst made by Ni impregnation of a commercial γ-alumina.

Experimental

Materials preparation

The parent Al-containing MIL-53(Al) material was synthesized according to a known microwavebased procedure reported to provide nanocrystallites of homogeneous distribution in shape and size [START_REF] Li | Fengting Li The Metal-Organic Framework MIL-53(Al) Constructed from Multiple Metal Sources:Alumina, Aluminum Hydroxide, and Boehmite[END_REF]. It consisted in mixing 1.21 g of AlCl 3 .6H 2 O and 0.42 g of benzene-1,4-di-carboxylic acid with 3 ml water and 5 ml of dimethylformamide (DMF), then transferring the mixture into a 50 ml reactor and heating it in a microwave oven for 30 minutes at 125 °C and with a power of 200 W. The solid was recovered by centrifugation and washed six times in 10 ml of DMF then in 10 ml of deionized water to ensure full removal of unwanted residues. After the last centrifugation, the white powder was dried for 24 h at 70 °C and heated in air at 220 °C for 72 h to remove the structuring agent and obtain the activated MIL sample. Nickel introduction was carried out by incipient wetness impregnation performed on the sample freshly dehydrated at 80 °C in an oven overnight. The procedure consisted in adding dropwise a Ni(NO 3 ) 2 .6H 2 O aqueous solution with a volume equal to the pore volume of the MOF support (as estimated from N 2 -sorption analysis, see below) and a Ni concentration established as to correspond to a Ni/Al atomic ratio in the sample of 0.07 (i.e. final Ni content in the dry catalyst of 5 wt%). The solid obtained after drying at room temperature for 24 h is hereafter denoted Ni/MIL. Part of it was kept as such to perform hydration-dehydration tests (see section 3.1) and the other part was submitted to calcination in air for 5 h at 500 °C (thin bed conditions, heating rate 0.5 °C.min -1 ) to give NiAl MIL .

For comparison purposes, two reference samples were also prepared from two distinct preformed alumina supports. The first one was obtained by placing the above MIL material in a ceramic plate and heating it in static air at 500 °C (heating rate 0.5 ° C.min -1 ) for 5 h to remove the organic linkers and lead to an alumina material (Al MIL ). The second support (Al COM ) was a conventional γ-alumina obtained by calcination at 500 °C in air of a commercial Boehmite (Sasol Plural, an alumina hydrate AlOOH).

After impregnation (5 wt% Ni, same conditions as above) followed by calcination in air at 500 °C, these two alumina supports gave samples Ni@Al MIL and Ni@Al COM , respectively.

Physicochemical characterization techniques

Elemental compositions (Ni and Al contents) were obtained by X-fluorescence (XRF) on a XEPOS spectrometer (Spectro Ametek) using a micropowder method and a calibration curve previously established from NiO and alumina mixtures with known chemical compositions.

Powder X-Ray diffractograms (PXRD) were registered on a PANalytical X'pert Powder instrument operated at 10 mA and 30 kV, using the Cu X-ray wavelength (λ=1.5405 nm). Crystalline phases were identified by referring to the international center for diffraction data (ICDD). The Scherer equation was applied to evaluate the average size (D in nm) of crystal domains, expressed as D=Kλ/βcosθ where K=0.9 is a constant, λ is the wavelength, β is the full width at half maximum (FWHM) of the diffraction peaks, and θ corresponds to peak position. N 2 -sorption adsorption-desorption isotherms were registered on an ASAP 2020 Micromeritics apparatus at -198 °C. Prior to analysis, the materials were degassed at 120 °C for 12h. BET (Brunauer-Emmett-Teller) surface areas were calculated within the relative pressure range 0.05-0.3. Total pore volumes were determined at saturation of the N 2 sorption. Micropore volumes were determined by the t-plot method and pore diameters were determined from the BJH (Barrett-Joyner-Halenda) model by using the desorption branch.

Reducibilities of the calcined materials were determined from temperature-programmed reduction (TPR) experiments carried out on an Autochem 2920 Micromeritics apparatus. Approximately 70-100 mg of the sample deposited on quartz wool in a U-shaped quartz tube was heated at a rate of 10 °C.min -1 from 25 °C to 900 °C under 5% H 2 /Ar gaseous flow (25 ml.min -1 ). The outlet gas flow was passed through an ice bath to trap the water formed during reduction (H 2(g) +NiO (s) →Ni 0 (s) +H 2 O (g) ) before

attaining the thermal conductivity detector (TCD). H 2 consumption was quantified by integrating the area of the TCD curve versus temperature and referring it to a calibration curve.

Thermogravimetric analysis (TGA), also expressed as differential thermogravimetric analyses (DTG), and coupled with differential thermal analysis (DTA), was performed on a TA SDT Q600 instrument working in horizontal mode. One of the two alumina crucibles was filled with about 25 mg of sample while the other (reference) crucible was kept empty. Both were located in the heating chamber and equipped with thermocouples to follow simultaneously weight and heat flow changes during heating, carried out in air flow (50 ml.min -1 ) from room temperature to 900 °C with a ramping rate of 5 °C.min -1 .

A capillary placed at the exit of the chamber and connected to a Fisons VG Thermolab Mass Spectrometer was also used to perform continuous mass spectrometry (MS) analysis of the exhaust gases.

Scanning Electron Microscopy (SEM) images were taken on a Hitachi SU-70 SEM-FEG microscope with an electron acceleration tension of 7 kV. After careful grinding, the samples were deposited on the analysis support and observed in mixed mode (70% of secondary electrons and 30% of retro-diffused Raman spectra were collected on a KAISER (RXN1) optical system equipped with a charge-coupled detector (CCD), a laser with λ = 785 nm (energy of 1.58 eV) and a microscope with an X50 long working distance (W.D. = 8.0mm) lens. The operation conditions were as follows: 10 mW laser power, 4 cm -1 resolution, 10 s acquisition time and a total of 30 accumulations per spectrum.

Catalytic measurements

Catalytic performances were studied in a fixed bed reactor (Hastelloy X, I.D.=9 mm, L=30.5 cm) situated inside a micro activity reference catalytic reactor (MAR, PID Eng and Tech Spain). Typically, 50 mg of catalyst were loaded into the reactor and reduced in situ under 5 vol% H 2 /Ar (30 ml.min -1 ) at 800 °C for 2 h (heating rate 5 °C.min -1 ). The temperature was then let to cool down to 650 °C, afterwards the reducing gas mixture was replaced by a flowing gas composed of an equimolar amount of CH 4 and CO 2 diluted in Ar (atomic ratio 5/5/90) and kept at atmospheric pressure and at a gas hourly space velocity (GHSV) of 72 L.g -1 .h -1 . This protocol was chosen based on TPR data and on previous studies [START_REF] Hassan | Low temperature dry reforming of methane on rhodium and cobalt based catalysts : Active phase stabilization by confinement in mesoporous SBA-15[END_REF][START_REF] Karam | Tuning the properties of nickel nanoparticles inside SBA-15 mesopores for enhanced stability in methane reforming[END_REF]. Catalytic performances were established from the gas composition at the exit of the reactor monitored with an INFICON micro GC along with a TCD and two parallel channels (Molecular sieve and Plot U). The data are reported after 0.5 h and 13 h of reaction (stability test) in terms of:

reactants conversions, estimated from the equation "X conversion (%) = (C in -C out ) / C in " , where C is the reactant (CH 4 or CO 2 ) and in and out represent the inlet and outlet concentrations of this gas, respectively, -H 2 /CO (mole ratio) = N H2 / N CO ", where N H2 and N CO are the molar concentrations of the formed H 2 and CO products.

Carbon balance was systematically verified during the tests and was always above 98%.

Results

Characteristics of the parent MIL-53, activated MIL and impregnated Ni/MIL materials.

After synthesis in microwave conditions, the parent MIL-53 sample is characterized by a X-ray diffractogram (Fig. 2a) similar to that reported in the literature for a MIL-53 material synthesized with the same microwave procedure (MW) [START_REF] Isaeva | Control of morphology and size of microporous framework MIL-53(Al) crystals by synthesis procedure[END_REF]. This validates the formation of the targeted MIL-53 framework that is known to display dynamic structure properties, implying pore size variations upon guest adsorption (so-called "breathing behavior") with an amplitude that depends on the size of the molecules adsorbed in the pores and on the strength of their interaction with the framework [START_REF] Loiseau | A rationale for the large breathing of the porous aluminum terephthalate (MIL-53) upon hydration[END_REF]. Such pores adaptation can explain the small differences observed between the PXRD pattern of the assynthesized MIL-53 prepared here by MW and that commonly reported after synthesis under conventional heating [START_REF] Isaeva | Control of morphology and size of microporous framework MIL-53(Al) crystals by synthesis procedure[END_REF], both methods inevitably leading to variations in the amounts of molecules (solvent and unreacted linkers) trapped within the pores due to different crystal formation rates.

After activation at 220 °C in air (to remove all adsorbates from the pores), the diffractogram changes to a completely different one (sample MIL, Fig. 2b), in accordance with the dynamic nature of the MIL-53 structure mentioned above. One can note however that the new PXRD pattern is not composed only of peaks belonging to the large pore form of MIL-53 (indicated by black triangles, typical of empty pores), but it also contains peaks related to the narrow pore form (indicated by black spheres) as well as reflections characteristic of an intermediate state between the open and closed forms (indicated by black diamonds). This multiphasic structural state is related to the thermal history of the sample as already described by others for MIL-53 [START_REF] Hsieh | MIL-53 frameworks in mixed-matrix membranes[END_REF] and can be taken as representative of the "activated form".

As shown in Fig. 2d, all the peaks corresponding to this activated form are still present after Ni impregnation (Ni/MIL), even if slightly shifted and with distinct relative intensities for some of them.

More importantly, new tiny peaks appear at unusual positions and they are visible both after drying in room atmosphere and at 200 °C (identified by stars, Fig. 2d and 2d', respectively). Such new reflections were to our knowledge never reported for MIL-53, strongly suggesting that nickel cations are inserted within the pores where they provoke structural deformation by interacting with the framework as in the case of adsorbed organic guest molecules [START_REF] Horcajada | Flexible Porous Metal-Organic Frameworks for a Controlled Drug Delivery[END_REF]. In order to confirm the origin of these new reflections, a series of experiments was done with the aim to demonstrate that they appear only in the presence of Ni and do not simply come from different hydration states generating cell volumes and symmetry changes in response to the above-mentioned "breathing behavior". We prepared to this end a nickel free hydrated MIL sample to be compared to freshly impregnated Ni/MIL. This was done by immersing the freshly activated MIL powder into a volume of pure water equivalent to the one used for Ni impregnation. PXRD patterns were then taken, after submitting both rehydrated MIL and impregnated Ni/MIL to the same successive drying treatments consisting first in natural drying in room atmosphere (Fig. 2c and2d) then in drying at 200 °C for 24h (Fig. 2c' and 2d'). Whatever the treatment, the additional tiny peaks (indicated by black stars) were present for Ni/MIL (Figures 2d and2d') but not for MIL (Fig. 2c and2c'). Together with the limited intensity of these peaks that fits well the limited Ni content in the samples, this systematic difference is in line with a local framework distortion due to interaction with nickel cations occluded within the pores.

Such distortion also agrees with the small shift to higher angle of the ( 011) reflection (at 2θ = 8.7°) of the large-pore phase pattern, which is particularly disturbed after nickel impregnation (Fig. 2d).

The presence of nickel within the pores is further supported by the strong decrease of the saturation volume of N 2 adsorbed at low relative pressure (P/P 0 = 0.2) in Ni/MIL (Fig. 3c) compared to both activated MIL (Fig. 3a) and rehydrated MIL (Fig. 3b), which corroborates a pore hindrance linked to the presence of Ni 2+ . Accordingly, the microporous volume drops from 0.40 cm 3 .g -1 (activated MIL) to 0.09 cm 3 .g -1 upon nickel impregnation (Ni/MIL) and the BET specific surface simultaneously diminishes from 1130 m 2 .g -1 (MIL) to 300 m 2 .g -1 (Ni/MIL). By comparison, the pore volume is much less affected upon MIL rehydration (0.32 cm 3 .g -1 ), even if a small decrease is also noted, revealing some loss of structural integrity after water adsorption as well, probably at the origin of the PXRD peak broadening seen for both rehydrated (MIL) and impregnated (Ni/MIL) series of samples ((011) reflection at 2θ = 8.7° mainly, Fig. 2). Noteworthy, the N 2 adsorption-desorption isotherms of all materials (MIL, rehydrated MIL and Ni/MIL, Fig. 3a-c) also display a marked step at relative pressures above P/P 0 = 0.9 characteristic of a substantial intercrystalline meso-macro porosity occurring between aggregated nanocrystals. A last evidence of the intimate mixing between the nickel precursor species and the porous network in Ni/MIL is given by thermogravimetric analyses that show a weigh loss of about 15 % between 200 °C and 420 °C for this sample (Fig. 4Ac) but no loss in this temperature range for hydrated MIL (Fig. 4Ab).

As confirmed by the simultaneous MS detection of NO (mass 30) in the exhaust gas (Fig. 4B, bottom, blue curve), this weight loss corresponds to the thermal decomposition of nitrates introduced together with nickel during the impregnation step. Notably, the oxidizing nitrogen oxide species such as NO formed within the pores of Ni/MIL during heating induce a lowering by more than 60 °C of the temperature at which the carbonaceous (terephthalate) linkers degradation occurs (important weight loss between 400-550 °C, Fig. 4c, top) compared to both freshly activated and rehydrated MIL (process shifted to 500-650 °C for both of them, Fig. 4a and4b). The DTG profiles (Fig. 4B, middle curves) reveal that the degradation takes place in two steps, as previously reported for the amino-teraphthalate MIL-125(Ti) [START_REF] Hu | Surfactant-assisted synthesis of hierarchical NH2-MIL-125 for the removal of organic dyes[END_REF]. The second step is accompanied by an evolution of CO 2 for both MIL and Ni/MIL and also of NO for Ni/MIL (see respective curves in Fig. 4B), and the process is strongly exothermic as seen from the DTA profiles (Fig. 4A (NO, NO 2 , etc.) [START_REF] Yuvaraj | Thermal decomposition of metal nitrates in air and hydrogen environments[END_REF] with simultaneous formation of a NiO-based phase (for the sake of neutrality preservation); (3) then, strongly adsorbed NO x oxidizers remaining within the pores in close proximity (or in contact) with the organic linkers promote their decomposition, decreasing the temperature at which it occurs [START_REF] Pourmortazavi | Effect of nitrate content on thermal decomposition of nitrocellulose[END_REF].

In view of these TG profiles, we decided to set the calcination temperature of Ni/MIL during sample preparation (to form NiAl MIL ) at 500 °C, a temperature high enough to degrade the linkers but guaranteeing as well the development of the porosity and the large surface area of the pure mineral final phase. Indeed, it was recently reported that a temperature close to the onset of the parent MOF calcination temperature is optimal for the preparation of inorganic structures with preserved morphology and dimensions [START_REF] Lü | MOF-templated synthesis of porous Co 3 O 4 concave nanocubes with high specific surface area and their gas sensing properties[END_REF], which is important for preserving an intercrystalline porosity and large surface area in the obtained catalyst. The temperatures of peak maxima in Fig. 4B are slightly higher than in Fig. 4A due to the delay for the exhaust gases to reach the MS analysis chamber. a: total surface area (S.A.), total pore volume (V tot ) and mean pore diameter (mean) calculated from the 3.2 Physicochemical properties of the calcined NiAl MIL and reference materials.

Fig. 5 and6 show typical SEM and TEM micrographs of the parent (MIL) and impregnated then calcined (NiAl MIL ) materials. In both scanning (Fig. 5a) and transmission (Fig. 6a) modes, Ni/MIL appears as slightly elongated aggregated nanocristallites with uniform sizes of approximately 20 nm in length. Such morphology is coherent with the high intercrystalline volume identified above from N 2 physisorption. After calcination at 500 °C (sample NiAl MIL , Fig. 5b and6b), the size and shape of the nanograins are globally preserved although they tend to become thinner, forming fused ill-defined and more defective alumina flakes that are actually composed of several randomly interwoven nanosheets (inset in Fig. 6b). This is in line with the N 2 isotherm of this sample that displays features typical of layered mesoporous materials, namely (i) a progressive increase of the adsorbed N 2 volume, observed here at P/P 0 higher than 0.75 and (ii) a slit-type hysteresis spread over P/P 0 values, ranging here from 0.75 to 1 (Fig. 3d) and revealing a relatively polydisperse size distribution of large mesopores. The BET specific surface of 239 m 2 .g -1 is still high, even if much smaller than in the parent MIL material (Table 1) due to the organic linkers removal, and the microporous volume is close to that in Ni/MIL (Fig. 3d and 3c, respectively). In order to evaluate the impact on the material properties of the adopted preparation procedure (i.e.

impregnating first the activated MIL with nickel then calcining it), a different order was applied by first calcining the activated MIL-53 (to obtain Al MIL ) then impregnating and calcining it at 500°C (to remove adsorbed nitrates and water molecules). This second sequence order, in which the impregnation is done on a preformed alumina support (here Al MIL ), resembles a conventional preparation route.

From PXRD, the structure of Al MIL is typical of an amorphous alumina (Fig. 7b*). It changes after Ni impregnation and calcination (sample Ni@Al MIL , Fig. 7b) to a crystalline phase very similar to that in NiAl MIL except for two reflections (2θ at 45.8° and 66.5°) that appear broadened and subtly shifted towards positions typical of γ-Al 2 O 3 (2θ = 46.2° and 67.0°). This suggests a coexistence of both phases, as also supported by the TPR profile that contains two peaks, (i) a main one at high temperature (with maximum at T = 820 °C, Fig. S1b) indicative of some inclusion of nickel within the oxide matrix as in NiAl MIL , and (ii) a smaller one centered at about 700 °C characteristic of nickel interacting less strongly with the alumina support (most probably nickel oxo-clusters on the external surface of the alumina grains). In term of morphology, the preformed alumina support appears as nanograins composed of both packed and interwoven nanosheets (sample Al MIL , Fig. 5c and6c) and this amorphous two-dimensional aspect is lost after nickel impregnation and calcination, turning to aggregates of tiny nanocrystals with ill-defined shapes (sample Ni@Al MIL , Fig. 5d and6d). Note that this final morphology significantly differs from the lamellar shape seen above for NiAl MIL . Moreover, the grains appear assembled, forming a polycrystalline structure with polymodal porosity, in line with the three well defined steps observed in the N 2 adsorption isotherm of this sample (Fig. S2a), typical of the coexistence of micro, meso and macropores. Therefore, the order of the "calcination / Ni 2+ impregnation" sequence, which affected the temperature of degradation of the MIL organic linkers (TG/SM data, Fig. 4), also impacts the structure, texture and reducibility of the final material. To sum up, impregnation of the nickel cations within the activated MIL-53 to position them along the aluminum chains of the framework before the calcination step (NiAl MIL ) allows for a somewhat better dispersion (Table 1) and stabilization of nickel through their insertion into interwoven 2-dimensional alumina nanosheets. Notably, another advantage of this sequence order is to require only one calcination step during material preparation instead of two for conventionally impregnated Ni@Al MIL . For the sake of comparison, a last material (Ni@Al COM ) was prepared by conventional nickel impregnation of a commercial γ-alumina (Al COM ) and tested as traditional reference catalyst. Both

Ni@Al COM and its support Al COM show four main diffraction peaks, which are typical of crystalline γ-Al 2 O 3 (Fig. 7c* and 7c, respectively). No evidence of spinel phase can be observed on the diffractogram of Ni@Al COM , although diffraction peaks corresponding to a small fraction of this phase might be masked by those of the γ-Al 2 O 3 phase. The TPR trace of this material displays two peaks, as for Ni@Al MIL , but centered at lower temperatures than previously (630 °C and 800 °C, Fig. S1c), indicating a weaker Nisupport interaction than in both MOF-based samples, especially NiAl MIL . Similarly to Ni@Al MIL , a fraction of the nickel in Ni@Al COM is probably existing in the material as amorphous NiO, as revealed by the TPR signal observed at relatively low temperature (T < 700°C). The significantly higher dispersion and stabilization of Ni in the latter sample can be explained as follows: the impregnation of nickel

Characteristics of the reduced catalysts

Previous to catalytic tests, all the prepared nickel-alumina materials were treated under a flowing H 2 atmosphere to obtain reduced Ni 0 , the active phase for DRM. The H 2 consumption during TPR was used to estimate the amount of reducible Ni 2+ in all nickel containing samples, considering the consumption of one H 2 molecule per Ni 2+ cation. For all samples, the Ni content is in accordance with that measured by X-fluorescence and close to the expected 5 wt% Ni, even if slightly lower, by less than 10%, for sample Ni 0 @Al MIL (Table 1).

The PXRD patterns of the samples reduced at 800 °C are shown in Fig. 7 slightly smaller than the value estimated by HR-TEM after measuring the sizes of at least 500 nanoparticles (Table 1). Considering the TEM image of Ni 0 Al MIL shown in Fig. 9a and the nonsymmetrical shape of the histogram of particles sizes for this sample (Fig. 10), it can be assumed that this average size discrepancy comes from contrast issues making uneasy the detection of the smallest nanoparticles (from 2 to 4 nm) on the alumina grains, in spite of the good resolution of HR-TEM images. In this sample, all detected nickel nanoparticles are homogeneously distributed all over the support which still appears as interwoven alumina sheets (as in NiAl MIL before reduction). The preservation of the morphology after reduction is also attested by textural features, the N 2 physisorption isotherm of Ni 0 Al MIL still showing a slit-like hysteresis typical of a layered material composed of poorly organized interlayered pores (Fig. 3e). The amount of N 2 adsorbed at P/P 0 > 0.8 is however higher than that before reduction (Table 1), indicating the formation of additional large mesopores within the material. This increase of the total pore volume and mean pore size upon reduction is likely due to the The reduction of nickel impregnated on the two preformed alumina supports (Al MIL and Al COM ) led similarly to the appearance of homogeneously dispersed Ni 0 nanoparticles (Fig. 9b and9c, respectively), but the particles and their size distributions were systematically larger than on Ni 0 Al MIL (Table 1), especially in the case of Ni 0 Al COM (size = 22 nm) (Fig. 10). Hence, the homogeneous NiAl 2 O 4based spinel nanophase formation seen during the first steps of preparation of Ni 0 Al MIL and issued from initial nickel insertion within the pores of the MOF with high surface area before linker decomposition seems to play an important role towards metal stabilization with a positive effect that remains after partial extraction from the alumina matrix upon reduction.

Catalysis performances

Finally, the catalytic performances of Ni 0 Al MIL were tested in the reaction of dry reforming of methane and compared to those of Ni 0 @Al MIL and Ni@Al COM (Fig. 11 and Table 2). After in situ reduction for 2 h at 800 °C to reduce nickel, each sample was cooled down to 650 °C and subjected to reaction for at least 13 h and for 100 h for Ni 0 Al MIL (stability test conditions). Fig. 11A and 11B show that the CO 2 and CH 4 conversions on both MOF-derived catalysts (Ni 0 Al MIL and Ni 0 @Al MIL ) are high, close to thermodynamic equilibrium, and much higher (2 to 3 times) than those on Ni 0 @Al COM . Noteworthy, Ni 0 Al MIL is not only the most active catalyst but it is also very stable as its performance after the prolonged 100 h on stream remain nearly equal to that measured at the early stage of the reaction (Figure 11D). The selectivity to DRM on this sample is also remarkable, as attested by two features that indicate limited occurrence of side reactions: (i) the close CH 4 and CO 2 conversions that differ by only 7% (Fig. 11A andB) and (ii) the H 2 :CO ratio very close to 1 (Fig. 11C). A rather good selectivity to DRM is also observed on Ni 0 @Al MIL but with a H 2 :CO ratio higher than 1 (Fig. 11C) that suggests the occurrence of some CH 4 decomposition (CH 4 (g) → C(s) + 2H 2 (g)). The formation of carbon deposits on the surface of this material is indeed confirmed by the 5 % weight loss detected at T = 600°C when the spent Ni 0 @Al MIL is submitted to TGA analysis (Fig. 12A) as well as by the D and G bands visible by Raman spectroscopy (Fig. 12B) [START_REF] Jabbour | Ordered mesoporous "one-pot" synthesized Ni-Mg (Ca)-Al 2 O 3 as effective and remarkably stable catalysts for combined steam and dry reforming of methane (CSDRM)[END_REF] and by the observation by scanning electron microscopy of carbon nanotubes on the catalyst surface (Fig. 13b and13b'). Regarding the Raman signals, the two main bands (D and G) are typical of the doubly degenerated phonon mode of C atoms in the sp2 carbon networks with high degree of symmetry and of order (graphitic carbon, G-band, 1600 cm -1 ) and of a disordered structural mode of carbon species (D-band, 1300 cm -1 ). The intensity ratio between these two bands can be used as an indicator of the crystalline degree and of the presence of defects in the carbon species; the smaller the ratio (<1) the higher the structural order. Here, both Ni@Al MIL and Ni@Al COM have a ratio significantly higher than 1 (4.6 and 4.2 respectively) which indicates the disordered character of the carbon deposits, in line with the occurrence of carbon nanotubes observed on these samples by electron microscopy. Such deposits are similarly clearly identified on the spent Ni 0 @Al COM catalysts, with a TGA weight loss (7 %) even higher (Fig. 12A) in spite of the low activity of this catalyst (low conversion level) and with D and G bands still clearly displayed on the Raman spectrograph (Fig. 12Bc). In contrast, such deposits are not detected on the most active Ni 0 Al MIL material, even after the 100 h test, neither by TGA (Figure 12A), nor by Raman spectroscopy (Fig. 12B) or electron microscopy (Fig. 13a and13a'). Finally, the occurrence of side reactions on the reference Ni 0 @Al COM material is confirmed by the products H 2 :CO, which is initially below 1 and decreases with time (Fig. 11B), indicating some H These results demonstrate the superior performances of the catalysts prepared from MIL-53, especially of Ni 0 Al MIL over Ni 0 @Al COM . The superiority of Ni 0 Al MIL in term of conversion is in line with the better Ni 0 dispersion in this sample (Fig. 13a and Table 1). It moreover reveals, together with the remarkable selectivity of Ni 0 Al MIL to DRM all along the run (higher than on Ni 0 @Al MIL also prepared from MIL-53), a preservation of the active sites and hence a persistent strong nickel-support interaction even in the harsh conditions of reaction. This is seen by the high and comparable nickel dispersion both after reduction (Fig. 9a,a' and 10A) and after reaction (Fig. 10A and13a). In contrast, a significant sintering takes place in Ni@Al COM in which the mean nickel particle size increases by about 30% after test (Table 1 and Fig. 10C,D). The formation in Ni 0 Al MIL of the stable spinel nickel aluminate phase, favored by the initial proximity between the nickel cations and the aluminum carboxylate clusters within the pore of MIL-53 during the calcination step, is believed to be at the origin of the stability of the dispersed Ni 0 nanoparticles. It could result from the presence of a remaining nickel aluminate spinel nanophase at the interface between the support and the metal nanoparticles after reduction that might insure a strong anchorage of the nanoparticles onto the resulting alumina support. One can also speculate that the layered texture of the material, which is remarkably well preserved during the 100h of reaction (Fig. 13a), contributes to the stability of the active phase as well by isolating and enhancing the occlusion of the nanoparticles, each of them being caught between two alumina sheets. This nanoparticle confinement effect may also explain the absence of carbon production on this sample. This explanation is in line with previous data showing that the occlusion of nickel nanoparticles within the mesopores of alumina supports prevent the catalysis of carbon nanotube formation, which occurs only from the nickel nanoparticles located on the external surface of the support [START_REF] Jabbour | Ordered mesoporous "one-pot" synthesized Ni-Mg (Ca)-Al 2 O 3 as effective and remarkably stable catalysts for combined steam and dry reforming of methane (CSDRM)[END_REF]. It is worth adding that although the nickel nanoparticles reported here are not the smallest reported in the literature so far (the application of atomic layer deposition recently afforded nanoparticle size below 3 nm [START_REF] Li | Steam reforming of n-dodecane over mesoporous alumina supported nickel catalysts: Effects of metal-support interaction on nickel catalysts[END_REF]), this study demonstrates the potential of MOFs as sacrificial hosts for preparing highly dispersed nickel nanoparticles whose size is below those prepared by employing other type of oxides as stable nickel cation hosts, such as alkaline earth metal substituted MZr 1-x Ni x O 3 -δ perovskites [START_REF] Dama | Active and durable alkaline earth metal substituted perovskite catalysts for dry reforming of methane[END_REF] or the lanthanum zirconate pyrochlore La 2 Zr 2 -xNi x O 7 -δ [START_REF] Saché | Ni stabilised on inorganic complex structures: superior catalysts for chemical CO 2 recycling via dry reforming of methane[END_REF]. Noteworthy, the use of MOF results in materials with a relatively high surface area (S. A. = 200 m 2 .g -1 ) compared to that obtained with other types of purely inorganic sacrificial nickel host cited above (S.A. < 20 m 2 .g -1 ), which can be seen as an additional benefit of this strategy in view catalysis applications.

Conclusion

In this work we demonstrate the benefit of using the aluminum carboxylate MIL-53 as a sacrificial nickel precursor host to prepare a dry methane reforming catalyst composed of nickel nanoparticles dispersed within a porous γ-Al 2 O 3 lamellar phase. The higher surface area of this metal-organic framework compared to preformed porous alumina supports allows for a more efficient dispersion of the nickel precursors and, in turn, for a superior dispersion and uniformity of the nickel nanoparticles obtained after the MOF calcination and subsequent reduction steps. Consequently, the performances in terms of activity, selectivity and stability of this new catalyst for dry reforming of methane within the reaction conditions applied in this study are higher than those of comparative nickel@alumina catalysts prepared by conventional nickel impregnation of preformed alumina supports. Although the application of more severe reaction temperatures (e.g. temperatures higher than 800°C) and higher gas pressures would be necessary to evaluate the performances of this MIL-53-derived catalyst under more realistic DRM conditions, the results displayed in this report demonstrate that MOF can be considered as a promising hybrid alternative to the sacrificial inorganic and non-porous hosts (perovskite, hydrotalcites) used so far to stabilize nickel species towards the synthesis of catalysts based on nickel metal nanoparticles. In particular, employing MOFs as a sacrificial framework leads to materials with a higher surface area that those commonly encountered in the domain and whose interlamellar spaces may provide an additional stabilization to the confined nickel nanoparticles.

Considering the large variety of MOF compositions described so far, this approach is believed to be promising for the preparation of appealing catalysts applied for other types of greenhouse gas valorization reactions, such as methanation, as it will be soon demonstrated in a shortcoming contribution.

Supporting Information. TPR data, N 2 sorption isotherms of the samples Ni@Al MIL and Ni@Al COM as well as of the samples Al COM , Ni@Al COM and Ni 0 @Al COM . CH 4 , CO 2 conversions and molar H 2 /CO product ratio obtained after the test of dry reforming of methane during 100 h over in-situ reduced Ni 0 Al MIL catalyst.

*To whom correspondence should be addressed:

Julien.reboul@sorbonne-universite.fr

Figure 1 .

 1 Figure 1. Schematic view of the innovative MOF-based preparation route of the Ni-alumina catalyst: (a) impregnation of the activated MIL-53(Al) with an aqueous Ni(NO 3 ) 2 solution to obtain Ni/MIL followed by (b) calcination at 500 °C to form the porous NiAl MIL solid solution and (c) reduction under H 2 at 800 °C to give the Ni 0 Al MIL catalyst. The grey color is for carbon atoms, blue is for aluminum, red for oxygen and green for nickel.

  signals). A JEOL-JEM 2010 electron microscope operating at 200 keV (LaB 6 gun) was used for High-Resolution transmission electron microscopy (HR-TEM) experiments. The observations were done on ultrathin sections of the samples prepared by mixing a few milligrams of the powder with an EPON 812 embedding resin in a beam capsule, heating the mixture at 60 °C for 48h to provoke polymerization, then cutting the polymerized blocks with a diamond knife. The obtained slices (50-70 nm in thickness) were deposited on copper grids covered with a carbon membrane layer. Average sizes of nickel particles, when visible, were estimated using the ImageJ software and taking at least 500 particles into consideration. Combined STEM-HAADF and EDX Mapping was performed on a JEOL-JEM 2020 microscope to visualize further the Ni distribution in ultra-thin sections (prepared as above) of a selected sample.

Figure 2 .

 2 Figure 2. X-ray diffraction patterns of: (a) as synthesized MIL-53(Al), (b) activated MIL, (c) rehydrated MIL dried in room atmosphere or (c') at 200 °C, (d) impregnated Ni/MIL dried in room atmosphere or (d') at 200 °C. The dot vertical line emphasizes the position of the 011 reflection of the PXRD pattern corresponding to the MIL-53 large-pore phase. Triangles and spheres indicate the peak positions corresponding to the open and closed forms of the MIL-53 structure, respectively (as reported in [41]). Diamonds show unattributed reflections that may correspond to an intermediate state. Stars indicate peaks appearing after nickel impregnation.

Figure 3 .

 3 Figure 3. N 2 physisorption isotherms after degassing at 120 °C of: (a) freshly activated MIL (red spheres), (b) rehydrated MIL (blue triangles), (c) Ni/MIL (green squares), (d) NiAl MIL (black diamonds) and (e) Ni 0 Al MIL (brown spheres). Filled and empty symbols correspond to adsorption and desorption isotherms, respectively.

Figure 4 .

 4 Figure 4. (A) TGA (top curves), DTG (middle curves) and DTA (bottom curves) profiles for samples: (a) freshly activated MIL (green lines), (b) rehydrated MIL (red dotted lines) and (c) impregnated Ni/MIL (blue lines). (B) Mass spectrometry profiles corresponding to the detection of H 2 O (mass 18), NO (mass 30) and CO 2 (mass 44) for hydrated MIL (red dotted lines) and impregnated Ni/MIL (blue lines).

N 2

 2 sorption isotherms. b: temperature at the maximum of TPR peaks. c: Ni contents determined from the amount of H 2 consumed during TPR taking into account that the theoretical amount is 860 μmol.g -1 for the reduction of 5wt% Ni 2+ . In brackets: values determined by XRF. d: mean size values calculated from the histograms obtained from representative TEM images of each samples. In brackets: values estimated from PXRD by applying the Scherrer equation considering the Ni 0 reflection at 2θ =51.9°. Note that the particle sizes could not be identified by PXRD after test because of the lack of sufficient amount of samples. e: nanoparticle dispersion calculated from the formula D = 6×10 9 × (V Ni /A Ni )×(1/d) where d is the particle size (estimated from PXRD), V is the volume of Ni atom in the bulk metal and A is the area of the Ni surface atom [49][50].

Figure 5 .

 5 Figure 5. Representative SEM images of (a) MIL, (b) NiAl MIL , (c) Al MIL and (d) Ni@Al MIL .

Figure 6 .

 6 Figure 6. Representative HR-TEM images of (a) MIL, (b) NiAl MIL , (c) Al MIL and (d) Ni@Al MIL .

Figure 7 .

 7 Figure 7. Diffractograms of the Ni containing samples: (a) NiAl MIL , (a') Ni 0 Al MIL , (b) Ni@Al MIL , (b') Ni 0 @Al MIL , (c) Ni@Al COM , (c') Ni 0 @Al COM . The diffractograms of the preformed alumina supports (light grey lines) are shown for comparison: (b*) Al MIL and (c*) Al COM . Intensities are divided by 3 in patterns c*, c and c'.

  nitrate on the parent MIL-53 with especially high surface area enables the initial occlusion of the nickel precursors along the organic linkers and [Al(OH)O] n clusters decorating the pores, leading after the MOF calcination step (organic removal) to the formation of an Al 2 O 3 -NiAl 2 O 4 phase in which nickel and aluminum cations are intimately mixed in a homogenous oxide phase. In contrast, impregnation of a preformed alumina (obtained by calcination of either MIL-53 or boehmite) then impregnation and calcination leads to a mixture of alumina, NiAl 2 O 3 and NiO with an overall weaker interaction between nickel and the alumina matrix. Interestingly, the presence of nickel at the very first step of the preparation seems to direct the formation of the crystalline spinel NiAl 2 O 4 phase during calcination (giving NiAl MIL ), while no spinel phase is formed when starting from pure MIL-53 (giving amorphous Al MIL ).

Figure 8 .

 8 Figure 8. STEM-HAADF image of NiAl MIL (left) and EDX elemental mappings of the same area for Al (in green) and Ni (in yellow).

  and compared to those before reduction. The reduction of NiAl MIL to Ni 0 Al MIL leads to the formation of new diffraction peaks assignable to the (111), (200) and (220) planes of crystalline Ni 0 (at 2θ = 44.5°, 51.9° and 76°, respectively, Fig. 7a'). Bands attributable to an alumina-based phase are still present but they are slightly shifted (compared to those before reduction) towards positions characteristic of pure γ-Al 2 O 3 which attests of Ni extraction from the oxide matrix towards the pore surface. Application of the Scherrer equation to the Ni 0 reflection at 2θ = 51.9° leads to a mean Ni 0 nanoparticle size of 5.3 nm,

  exfoliation of the packed nanosheets through (i) the transformation of the NiAl 2 O 4 spinel phase into γ-Al 2 O 3 and (ii) the formation of the Ni 0 nanoparticles that may tend to further space apart the nanosheets one from another.

Figure 9 .

 9 Figure 9. HR-TEM images of the reduced samples: (a and a') Ni 0 Al MIL , (b and b') Ni 0 @Al MIL , (c and c') Ni 0 Al COM .

Figure 10 .

 10 Figure 10. Histograms established from at least 500 Ni 0 nanoparticles in typical HR-TEM images of (A) Ni 0 Al MIL , (B) Ni 0 @Al MIL and (C) Ni 0 @Al COM before and after 13h of catalytic test (black and white bares, respectively). Related mean sizes are shown in (D) for the catalysts before (black circles) and after (grey circles) test.

Figure 11 .

 11 Figure 11. Comparison of the catalytic performances in DRM of the Ni 0 Al MIL (red squares), Ni 0 @Al MIL (green spheres) and Ni 0 @Al COM (blue triangles): (A) CH 4 conversion, (B) CO 2 conversion, (C) H 2 :CO ratios measured during 13h of reaction and (D) CO 2 (filled squares) and CH 4 (empty squares) conversions as well as H 2 :CO ratios (red dash-dotted line) measured during 100h of reaction for Ni 0 Al MIL (the black dash-dotted line represent H 2 :CO = 1). Black continued and dashed lines are the calculated values at the thermodynamic equilibrium for the CH 4 and CO 2 conversions, respectively.

Figure 12 .

 12 Figure 12. (A) TGA profiles (a, b, c at the top) associated to MS signals corresponding to CO 2 (a', b', c' at the bottom) and (B) Raman spectroscopy profiles of the spent catalysts (a, a') Ni 0 Al MIL , (b, b') Ni 0 @Al MIL and (c, c') Ni 0 @Al COM .

Figure 13 .

 13 Figure 13. TEM (a,b) and SEM (a',b') images of spent NiAl MIL (a,a') and Ni@Al MIL (b,b'). The dashed oval and white arrows in (b) highlight the presence of carbon nanotubes.

  , bottom curves). Hence, all data indicate that the linkers degradation is displaced towards a significantly lower temperature after nickel impregnation, giving the following scenario during heating of Ni/MIL: (1) the temperature increase up to 150 °C leads to the dehydration of the hexahydrate nickel nitrate used as nickel precursor according to the equation Ni(NO 3 ) 2 .6H 2 O → Ni(NO 3 ) 2 + 6H 2 O, the theoretical mass loss of 59% for this event being close to the 50% measured by TG; (2) this is followed between 200 and 400 °C by nitrates decomposition into NO x

Table 1 .

 1 Textural characteristics, reduction temperature and nickel content.

	Sample names	Textural properties S.A. (m 2 .g -1 ) V tot (cm 3 .g -1 a )	Φ (nm) mean	Reduction Temperature (°C)	b	Ni content (wt%)	c	Ni 0 size (nm) nanoparticle d reduced spent	0 Ni nanoparticle dispersion (%) e
	MIL	1130	0.80	n.d
	NiAl MIL	239	0.82	10	830		5.1 (5.2)
	Ni 0 Al MIL	200	1.02	16					6.8 (5.3)	7.5	19
	Al MIL	318	1.19	6.2
	Ni@Al MIL	244	0.94	5.1	700/820		4.7 (4.6)
	Ni 0 @Al MIL	130	0.55	5.6					10.2 (8.0) 11.0	12
	Al COM	185	0.51	9.2
	Ni@Al COM	168	0.47	9.9	630/800		5 (5)
	Ni 0 @Al COM	125	0.43	10.3					22 (18)	29	5.6

  2 consumption through reverse water-gas shift (RWGS) reaction (CO 2

+ H 2 → CO + H 2 O) and/or direct reduction of CO 2 (CO 2 + 2H 2 ↔ C + 2H 2 O) that consumes a part of the produced H 2 .

Table 2 .

 2 Catalytic test results obtained with Ni 0 AlMIL and the comparative catalysts Ni 0 @Al MIL and Ni 0 @Al COM .

		Conversion (%)					Molar composition of main products	Coke content (%)
	Catalyst	t=0.5h		t=13h		t=100h		t=0.5h	t=13h	t=100h
											t=13h	t=100h
		XCH4	XCO2	XCH4	XCO2	XCH4	XCO2	H2/CO	H2/CO	H2/CO
	NiAlMIL	75	79	76	79	74	80	1.02	1.01	1.03	-	-
	Ni@AlMIL	71	79	70	77	n.p.	n.p.	1.12	1.09	n.p.	3	n.p.
	Ni@AlCOM	22	27	20	26	n.p.	n.p.	0.93	0.89	n.p.	3	n.p.

Acknowledgments. The authors sincerely acknowledge the ERANET EU-FP7 initiative, the national ANR (France) and CNRS-L (Lebanon) agencies for their financial support through the SOL-CARE (Energy-065) project (JC-ENERGY-2014 first call). The University of Balamand Research Council is also thanked for the BIRG 02/2016 contribution.


Peter Dunne
email: peter.dunne@ipcms.unistra.fr

Ciaran Fowley
email: c.fowley@hzdr.de

Gregor Hlawacek

Jinu Kurian

Gwenael Atcheson

Silviu Colis

Niclas Teichert

Bohdan Kundys

Munuswamy Venkatesan

Jürgen Lindner

Alina Maria Deac

Thomas M Hermans

J M D Coey

Bernard Doudin

Helium Ion Microscopy for Reduced Spin Orbit Torque Switching Currents

Keywords: spintronics, spin orbit torque switching, nanomagnetism, ion beam irradiation MAIN

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

adjacent heavy metal (HM) layers. In contrast to spin-transfer-torque MRAMs [START_REF] Perrissin | A Highly Thermally Stable Sub-20 Nm Magnetic Random-Access Memory Based on Perpendicular Shape Anisotropy[END_REF] , SOT devices require less demanding 3D fabrication, greatly simplifying their production. Moreover, their planar nature allows for easy visualisation and easier exploitation of the stray magnetic fields in applications such as reprogrammable magnetic domain configurations for spin-wave logic [START_REF] Wagner | Magnetic Domain Walls as Reconfigurable Spin-Wave Nanochannels[END_REF] .

The present drive to reduce switching currents and energy consumption is constrained by the selection of materials necessary to switch a magnetic domain via SOT. Similarly, real devices require multi-step nanofabrication to define memory storage cells. Here we present a new method to both reduce the switching currents, and to avoid multi-step lithography by using He + irradiation to locally modify the material properties in a single layer. Consequently, we can reduce the critical current density in a macroscopic device by almost an order of magnitude. Furthermore, we demonstrate how in-situ electrical measurement during irradiation allows precise control over the manipulation of the perpendicular magnetic anisotropy, capturing its evolution as a function of ion dose, making possible high-precision control and reproducibility of the process.

Light ion irradiation can lead to structural reorganization of stable material phases while preserving overall atomic composition with little to no topological damage [START_REF] Fassbender | Tailoring Magnetism by Light-Ion Irradiation[END_REF] . This allows for the reduction of PMA in thin magnetic multilayers [START_REF] Hoshi | Differences between Sputtering Methods in The Formation of Amorphous Magnetic Alloy Films[END_REF][START_REF] Chappert | Planar Patterned Magnetic Media Obtained by Ion Irradiation[END_REF] , which is caused by the intermixing of layers and increase of interfacial roughness [START_REF] Fassbender | Tailoring Magnetism by Light-Ion Irradiation[END_REF] . Furthermore, He + and Ar + broad-beam irradiation has been used to tailor SOT driven domain-wall dynamics [START_REF] Zhao | Enhancing Domain Wall Velocity through Interface Intermixing in W-CoFeB-MgO Films with Perpendicular Anisotropy[END_REF] , improve the spectral linewidth of SOT nanooscillators [START_REF] Jiang | Reduced Spin Torque Nano-Oscillator Linewidth Using He + Irradiation[END_REF] , and reduce the critical current density to switch the magnetisation of extended thin films via SOT [START_REF] Yun | Lowering Critical Current Density for Spin-Orbit Torque Induced Magnetization Switching by Ion Irradiation[END_REF] . Patterned magnetic anisotropy using broad beam irradiation has been demonstrated using shadow masks, but the smallest patterned volumes are mostly determined by the lithography process rather than the ion-beam interaction volume [START_REF] Chappert | Planar Patterned Magnetic Media Obtained by Ion Irradiation[END_REF] . On the contrary, focused helium ion beams allow precise control of perpendicular magnetic anisotropy [START_REF] Franken | Precise Control of Domain Wall Injection and Pinning Using Helium and Gallium Focused Ion Beams[END_REF][START_REF] Fowley | Local Modification of Magnetic Anisotropy and Ion Milling of Co/Pt Multilayers Using a He + Ion Beam Microscope[END_REF] and magnetisation [START_REF] Röder | Direct Depth-and Lateral-Imaging of Nanoscale Magnets Generated by Ion Impact[END_REF] on the order of size of the collision cascade (< ~10 nm) [START_REF] Hlawacek | Helium Ion Microscopy[END_REF] . With helium ion microscope (HIM) based He + irradiation, the lateral dimensions of the magnetic structures are limited only by the collision cascade [START_REF] Hlawacek | Helium Ion Microscopy[END_REF] due to the minimal beam-size of only 0.5 nm. SQUID magnetometry on unpatterned films with the substrate orientated parallel and perpendicular to the applied field indicated robust PMA (Fig. 1a), as expected for a sub-nm Co film [START_REF] Carcia | Perpendicular Magnetic Anisotropy in Pd/Co and Pt/Co Thin-film Layered Structures[END_REF] , with an effective anisotropy field of 1.28 T, and a saturation magnetisation µ0Ms of 1.02 T.

Samples of

Anomalous Hall effect measurements confirmed that the PMA remains intact on the patterned Hall bars (Figs. 1b,c). Here, we found square hysteresis loops and coercivities of ~20 mT (Fig. 1b) for magnetic fields applied perpendicular to the film along z, and effective magnetic anisotropies of well over 1 T (Fig. 1c) for hard axis measurements (along x). At positive saturation (+mz), the anomalous Hall resistance is 0.59 Ω.

Current-induced (Ix) switching measurements under a bias field µ0Hx = 150 mT (Fig 1d) yield a critical current density, Jc, of 6.0 MA cm -2 for full magnetisation reversal in the microfabricated structure. To minimise thermal effects while sweeping the current, Ix was pulsed for a duration of 2 ms, with a delay of 1 s between subsequent pulses, i.e., a duty cycle of 0.2%. Similar to the field driven switching in Fig. 1b, there is only one step and the saturated state has a resistance of 0.57

Ω.

The patterned devices were locally irradiated with helium ions with a 30 kV acceleration voltage using an Orion Nanofab helium ion microscope system, while in real time, we monitored the evolution of the anomalous Hall resistance in-situ. This electrical characterisation system consists of four Kleindieck MM3A micromanipulators which allow three-dimensional positioning of the contacting needles in a high vacuum environment (see Fig. 2a). A Keithley 2400 source meter was used to apply current and read the voltage drop corresponding to the longitudinal and transverse resistances. A current of 1 mA applied along the x-axis is used to probe the evolution of the anomalous Hall resistance close to zero applied magnetic field. The combination of the HIM with the possibility for in-situ control of the magnetisation state via Hall resistance measurements is unique in that it allows unprecedented spatial resolution, high flexibility in the patterning by avoiding high resolution but low yield lithography steps, resulting in a high turn around for explorative experiments. This compliments alternative approaches based on laser illumination with a 1.5 µm lateral resolution [START_REF] Cao | Deterministic Magnetization Switching Using Lateral Spin-Orbit Torque[END_REF] , ion broad beam irradiation which requires thick photoresist masks that limit the achievable spatial resolution [START_REF] Yun | Lowering Critical Current Density for Spin-Orbit Torque Induced Magnetization Switching by Ion Irradiation[END_REF] , electric field control of anisotropy [START_REF] Lee | Fast Magneto-Ionic Switching of Interface Anisotropy Using Yttria-Stabilized Zirconia Gate Oxide[END_REF] , or even oxidation which is CMOS compatible but inflexible and not ideal for fundamental studies [START_REF] Baek | Complementary Logic Operation Based on Electric-Field Controlled Spin-Orbit Torques[END_REF] .

During irradiation, the anomalous Hall resistance steadily decreases (Fig. 2b) until a critical dose of 35 ions nm -2 is reached, and a sharp switching is observed. The sharp fall in ΔR is complete at 37 ions nm -2 . The total fall of resistance (0.52 Ω) is close to the full deflection observed in the exsitu measurements shown in Fig. 1, marked as a grey dashed line, and indicates a reduced effective anisotropy field in the irradiated zone (HK2) compared to that of the extended structure (HK1) (Fig. 2b). The layer intermixing introduced by irradiation gives control over Keff, and HK2 thus becomes an experimentally tuneable parameter. Although the irradiations were carried out over the complete area of the Hall cross, as shown in Fig. 2c, the close proximity of the unirradiated areas allow for residual information from those regions to be gathered in the Hall loops, leading to the offset of ~0.1 Ω in Fig. 2b.

We selected four doses along this irradiation curve to demonstrate the ease of applying this technique to the reduction of critical currents for spin orbit torque switching. Figure 2d shows the corresponding anomalous Hall response versus µ0Hx for 1, 20, 30, and 50 ions nm -2 . Aside from the reduction in magnetic anisotropy field, evidenced by the saturation of RAHE at lower in-plane applied fields, several steps are observed in the anomalous Hall loops for irradiation doses ≥ 30 ions nm -2 . The additional steps correspond to the unchanged anisotropy in the unirradiated adjacent regions. This offset is highly dependent on the accuracy of overlaying the irradiation box with the junction area, and the residual Hall resistance is never precisely the same for each junction.

However, the multi-step switching, as seen in Fig, 2d, is always a result of the reversal of neighbouring regions adjacent to the Hall cross.

Assuming coherent rotation of the magnetisation, the normalised first quadrant magnetisation, 𝑀 " # , (Fig. 2e) can be fitted with [START_REF] Fowley | Direct Measurement of the Magnetic Anisotropy Field in Mn-Ga and Mn-Co-Ga Heusler Films[END_REF] 𝑀 " # = % 1 -( 𝜇 * 𝐻 , 𝜇 * 𝐻 - .

/ (1) which yields the magnetic anisotropy field, µ0HK, as a function of irradiation dose (Fig. 2f). We observe a continuous decrease in the anisotropy field, from 1.32 T for as-deposited films, to only 0.12 T for samples irradiated beyond the critical dose of 35 ions nm -2 . The saturation magnetisation, Ms, (Fig. 2f) was determined from the relative change in the saturation Hall resistance with ion dose, when the magnetic field is applied out-of-plane, along z (not shown here).

Ms is reduced at higher ion doses, due to alloying of the Co/HM interfaces, but is less sensitive than the magnetic anisotropy field, HK, only reducing to 86% of the initial value compared to 33%

for the anisotropy field after a dose of 30 ions nm -2 .

Both the HK and Ms play an important role in current-induced SOT switching. Under a macrospin approximation with the condition 𝐻 , ≪ 𝐻 -, where 𝐻 , is a magnetic bias field applied in-plane, the critical current density, Jc, required to switch the magnetisation direction of a magnetic state (in SI units) is 29

𝐽 5 = 2e𝜇 * 𝑀 8 𝑡 : ℏ𝜃 => ( 𝐻 - 2 - 𝐻 , √2 . ( 2
)
where e is the electronic charge, ℏ is the reduced Planck constant, tf is the magnetic free-layer thickness, and 𝜃 => is the spin Hall angle. However, when the applied field is comparable to the anisotropy, as we find for increased ion doses, the critical current density can instead be written as 29

𝐽 5 = 2𝑒𝜇 * 𝑀 8 𝑡 : ℏ𝜃 => A % 𝐻 - / 32 C8 + 20𝑏 / -𝑏 H -𝑏(8 + 𝑏 / ) I / JK (3)
where

𝑏 = > L > M
. The effect of ion dose on Jc is initially minimal for doses up to 20 ions nm -2 (Figs.

2f, 3a, & Table 1), with 𝐽 5 N reducing from -6.0 to -5.0 MA cm -2 (for µ0Hx = 150 mT). For 30 ions nm -2 , close to the critical dose, the critical current density is, however, reduced by almost one order of magnitude to -800 kA cm -2 . At larger doses, 50 ions nm -2 , we only observe the switching in the peripheral regions adjacent to the Hall cross.

Nonetheless, reducing the critical current density is not the sole criterion for improvement of SOT switching of magnet states. Additionally, we define a SOT switching efficiency, 𝜂, being the ratio of the magnetic energy barrier (𝐾 Q:: ) per unit area (in the magnetic free layer tf) versus the electrical power (P) per unit area (in the entire stack thickness d) required to switch the magnetic state, in units of seconds:

𝜂 = 𝐾 Q:: 𝑡 : 𝑃𝑑 = 𝜇 * 𝑀 8 (𝐻 --𝐻 , )𝑡 : 2𝐽 5 / 𝜌𝑑 ( 4
)
where 𝜌 is the total resistivity of the device. The choice of units is deliberate; the minimum pulse length for deterministic SOT switching is not always reported, but if it is known, then 𝜂 can be divided by this duration to determine the dimensionless efficiency. Either way, a larger value of 𝜂 implies a more stable magnetic element that requires less power to switch. Defined in this way, the switching efficiency of different material systems can be compared regardless of the pulse length, so long as the bias fields are similar. Furthermore, increasing N, the number of Pt/Co/W repeats, will not impact the efficiency if the relative amount of magnetic material remains the same, i.e. tf /d is constant. If the bias fields are not similar, substituting Jc from Eq. ( 2) into ( 4)

under the condition 𝐻 , ≪ 𝐻 -, leads to an intrinsic SOT efficiency for PMA systems:

𝜂 V = ℏ / 𝜃 => / 2𝜇 * 𝑀 8 𝐻 -𝑒 / 𝜌𝑑𝑡 : (5)
The advantage of Eq. ( 5) is that it depends solely on intrinsic material properties, and not on the experimental conditions, thus making it a more transparent comparator between different reported systems, particularly when the experimental current pulse may be much longer than the intrinsic switching time. Its drawback is that it is based on a macrospin approximation, which cannot envisage a multidomain response, and therefore overestimates the critical current density by a factor 10 -100.

Nevertheless, Eq. 5 provides a theoretical guideline for optimum parameters when considering material systems. This expression illustrates that to achieve large switching efficiencies, a minimal amount of conductive material must be used, while retaining a large spin Hall angle and a small anisotropy field. We find that neither the calculated spin Hall angle nor resistivity change significantly upon irradiation, with the former remaining ~0.4, and the latter ranging from 215 -235 µΩ.cm. Similar to Jc, we find there is negligible change in the intrinsic switching efficiency, 𝜂 V , for irradiation doses < 20 ion nm -2 , which remains close to 2 fs (Table 1), increasing to a maximum of 4.2 fs at 30 ions nm -2 . In contrast, the experimental efficiency, 𝜂, is initially ~7.1 ps, and drops at 20 ions nm -2 to 4.3 ps, before then increasing to its maximum of 86 ps at 30 ions nm -2 , an order of magnitude improvement upon the virgin state. The discrepancy between 𝜂 and 𝜂 V is due to the breakdown of the macrospin approximation used to calculate the latter. In all cases, the switching efficiency compares favourably to the literature values listed in Table 2, where efficiencies are of order 0.01 -10 ps, with a maximum of 69 ps found for metallic Ta/Pt/Co/W/Ta multilayers [START_REF] Zhang | Current-Induced Magnetization Switching in Pt/Co/W and Pt/Co/W0. 82Pt0.18 Structures with Perpendicular Magnetic Anisotropy[END_REF] , and 21 ps for insulating barium ferrite on Pt [START_REF] Li | Spin-Orbit Torque-Assisted Switching in Magnetic Insulator Thin Films with Perpendicular Magnetic Anisotropy[END_REF] , while zero-field switching devices [START_REF] Lee | Fast Magneto-Ionic Switching of Interface Anisotropy Using Yttria-Stabilized Zirconia Gate Oxide[END_REF] , present much lower efficiencies (0.74 ps).

Returning to the SOT switching for 30 ions nm -2 , shown in Fig. 3a, we sketch the magnetic states at points 1, 2, 3 and 4 in Fig. 3b. As the interior of the Hall cross can be addressed separately from the rest of the device, this forms a two-state system that can be addressed electrically as shown in Fig. 3a, or magnetically as shown in Figs. 2c and3c,d. Electrical switching involves applying a bias in-plane field (here 150 mT), and pulsing currents in a range of ± 3 mA to ensure switching of only the irradiated region. This leads to a switch between states 1 and 2, or 3 and 4 in (Fig. 3a,b). Otherwise pulsing currents of order ± 7.5 mA switches the magnetisation of the entire device between states 1 and 3 (Fig. 3a,b). Magneto-optic Kerr effect (MOKE) imaging under magnetic fields normal to the device (µ0Hz) confirms that the entire irradiated region switches independently from the rest of the device, Fig. 3c. Individual hysteresis loops for the irradiated (30 ions nm -2 ) and non-irradiated regions (as-dep1, as-dep2) are shown in Fig. 3d. The irradiated and unirradiated regions exhibit coercivities which are well spaced in magnetic field, µ0Hc = 5 mT and µ0Hc = 16 mT, respectively. The reader is directed to the supplementary video 1 to view the full magnetisation response during a field sweep.

Although there is an equivalence between electrical and magnetic switching, electrical SOT switching is more robust to variability in lithographic processing of the Hall bars than magnetic switching. This is because the former depends on the perpendicular magnetic anisotropy of a device, compared to coercivity for the latter. We find less than 1% variation in the magnetic anisotropy for the non-irradiated devices, µ0HK =1.32 ± 0.01 T, as the anisotropy is determined by the deposition process, which in bulk, non-patterned, thin films is 1.28 ± 0.05 T. In contrast, we find a 26% variation of coercivity in the same devices, 14 ± 4 mT, reflecting the strong dependence of the coercivity on the lithographic processing and slight variations in dimensions of the Hall bars. Since we address the anisotropy directly with ion irradiation, this illustrates the advantageous precision and robustness of focussed ion beam irradiation for magnetic element patterning compared to material removal.

In order to assess the suitability of the device as a memory storage element, we analyse the effect of He + irradiation on the thermal stability of the created magnetic bits. As expected, the reduction in 𝐽 5 does come at the cost of thermal stability, defined by a parameter Δ = 𝐾 Q:: 𝑉/𝑘 \ 𝑇, where 𝑉 is the magnetic element volume 𝑘 \ is Boltzmann's constant and 𝑇 is temperature. However, even for a 1 nm thick, 32 nm sided element, i.e. a 1024 nm 3 volume, ∆ = 64 for doses as high as 20 ions nm -2 , more than the 10-year stability criterion of 60 used in magnetic recording [START_REF] Evans | Thermally Induced Error: Density Limit for Magnetic Data Storage[END_REF] . Close to the critical dose, at 30 ions nm -2 , ∆ drops to 38-the equivalent of 95% data retention for 16 days [START_REF] Weller | Thermal Effect Limits in Ultrahigh-Density Magnetic Recording[END_REF] but unlike data storage, long-term stability is not critical to all applications. In fact, weakly bistable devices can be suitable for spin Hall oscillators, where the strength and modulation of spin wave propagation depends on the magnetic anisotropy [START_REF] Fulara | Spin-Orbit Torque-Driven Propagating Spin Waves[END_REF] . An easy way to improve the thermal stability would be to add N Pt/Co/W repeats to the system. ∆ will increase linearly with N, while not impacting η (Eq. 4), as stated previously. Furthermore, the reduction in local magnetic anisotropy allows for minimal losses during magnetisation switching. This level of control highlights the effectiveness of this method to tailor the switching current density in PMA multilayers. Although only two steps are shown in Fig. 3, one can also write switching current spatial gradients of almost any value directly. This allows for complex tailoring of not only field free memory elements [START_REF] Yu | Switching of Perpendicular Magnetization by Spin-Orbit Torques in the Absence of External Magnetic Fields[END_REF][START_REF] Lau | Spin-Orbit Torque Switching without an External Field Using Interlayer Exchange Coupling[END_REF][START_REF] Cai | Electric Field Control of Deterministic Current-Induced Magnetization Switching in a Hybrid Ferromagnetic/Ferroelectric Structure[END_REF][START_REF] Sheng | Adjustable Current-Induced Magnetization Switching Utilizing Interlayer Exchange Coupling[END_REF] , but even permits the control of spin-wave waveguide locations coupled with on-and-off spin-wave transmission control [START_REF] Wagner | Magnetic Domain Walls as Reconfigurable Spin-Wave Nanochannels[END_REF][START_REF] Sluka | Emission and Propagation of 1D and 2D Spin Waves with Nanoscale Wavelengths in Anisotropic Spin Textures[END_REF] . Such switching current spatial gradients can be used to tailor the synaptic-like response from PMA memory elements [START_REF] Cao | Tuning a Binary Ferromagnet into a Multistate Synapse with Spin-Orbit-Torque-Induced Plasticity[END_REF] when applied to neuromorphic computing or even logic elements [START_REF] Baek | Complementary Logic Operation Based on Electric-Field Controlled Spin-Orbit Torques[END_REF] . Extremely low thermal stability, would mean metastable 0 and 1 states which can be applied to probabilistic computing [START_REF] Borders | Integer Factorization Using Stochastic Magnetic Tunnel Junctions[END_REF] or random number generation [START_REF] Lee | Design of High-Throughput and Low-Power True Random Number Generator Utilizing Perpendicularly Magnetized Voltage-Controlled Magnetic Tunnel Junction[END_REF] .

In summary, a He + microscope has been used to perform mask-less light-ion irradiation on perpendicular magnetic anisotropy stacks for local control of the SOT switching properties. In-situ electrical measurements of anisotropy reduction provides real-time control of the anisotropy, allowing accurate determination of the critical dose needed to reach a desired, up to the transition to in-plane magnetic anisotropy. The corresponding reduction in critical current densities for SOT switching are nearly linear with decreasing anisotropy. Irradiation at doses just below the critical dose allow us to achieve almost an order of magnitude control over the critical current required for robust switching, and DC switching current densities as low as 800 kA cm -2 are achieved, which we believe is the lowest to date. Moreover, the switching efficiency of 86 ps is larger than the previous state of the art. It illustrates the unique advantages of this approach, and opens the door to preferential switching of predetermined areas of a device, limited only by the nm-resolution of the He + ion beam microscopy, while preserving the flat topography of the initial magnetic stack. The properties are: effective anisotropy Keff, thermal stability parameter Δ, negative and positive critical currents 𝐼 5 N , 𝐼 5 _ , the corresponding experimental critical current densities 𝐽 5 N , 𝐽 5 _ , the theoretical critical current density 𝐽 5 `a determined using Eq. ( 3) with µ0Hx = 150 mT, and an effective spin hall angle measured to be 𝜃 => = 0.4, the switching efficiency 𝜂 calculated from Eq. ( 4), and intrinsic efficiency 𝜂 V using Eq. ( 5). Note that for 50 ions nm -2 , only switching of adjacent unirradiated regions was detected. [START_REF] Sinova | Spin Hall Effects[END_REF] , and for multilayers with adjacent layers of opposite 𝜃 => , the magnitude of the larger one was chosen; c) calculated for µ0Hx = 0; d) when not reported, calculated using bulk resistivity values from [START_REF] Haynes | CRC Handbook of Chemistry and Physics[END_REF] , and then corrected using the M-S model [START_REF] Mayadas | Electrical-Resistivity Model for Polycrystalline Films: The Case of Arbitrary Reflection at External Surfaces[END_REF] ; e) topological insulator used for spin injection; f) ferromagnetic insulator; g) zero field switching device, moving between a remnant state mz = 0, and a downward pointing -mz state.

DETERMINING MAGNETIC ANISOTROPY FROM ANOMALOUS HALL EFFECT

If the magnetisation, M, of a perpendicularly magnetised thin film were to rotate coherently under an in-plane external magnetic field H x , then its magnetic free energy can be described using the Stoner-Wohlfarth model as [1]

E = K U sin 2 θ -µ 0 M H x cos (θ -π/2) (1)
where K U is the uniaxial anisotropy, and θ is the angle between the magnetisation and the normal to the surface. Minimising the free energy with respect to θ results in

θ = sin -1 µ 0 M H x 2K U (2)
The anomalous Hall effect is sensitive to the perpendicular component of the magnetisation, giving a signal proportional to M z

M z = M cos θ = M cos sin -1 µ 0 M H x 2K u = M 1 - µ 0 M H x 2K u 2 (3)
For the normalised data used here, the equation is recast as

M z = 1 - B x B K 2 (4)
where M z is the normalised z component of the magnetisation, B x = µ 0 H x and B K = µ 0 H K are the applied field and anisotropy field in units of T.

Normalised AHE data for a sample after irradiation (20 ions/nm 2 ) are shown in Fig. 1. The in-plane component, M x was assumed to be M x = 1 -M 2 z , also shown in Fig. 1. The data was fitted using Eq. 4 from 0 to 300 mT, but above ≈430 mT the magnetisation deviates due to the formation of domains. The anisotropy field, H K , was extrapolated from the coherent region to M z = 0, M x = 1, was found to be µ 0 H K = 611 mT. From unpatterned film SQUID magnetometry, we found the saturation magnetisation to be 0.83 MA m -1 , and noting the following relation

µ 0 H K = 2K ef f M s (5)
where M s is the thin film saturation magnetisation, the effective anisotropy K ef f is 0.51 MJ m -3 .

M z M x M z f i t M x f i t M z , M x µ 0 H ( T )
FIG. 1. Normalised AHE data for increasing in-plane applied magnetic field (Mz, square scatter plot), with its corresponding fit (solid line), and extrapolated in-plane magnetisation (Mz).

COERCIVITY AND ANISOTROPY DISTRIBUTION

The anomalous Hall resistance of 8 non-irradiated samples was measured for µ 0 H z and µ 0 H x (Fig. 2). A significantly larger distribution of coercivities (14 ± 4 mT) was observed compared to the variation of anisotropy (1.32 ± 0.01 T).

THICKNESS-DEPENDENT CORRECTION OF RESISTIVITIES

The resistivity of metal thin films is known to increase significantly for thicknesses below 10 nm [2]. As the experimental resistivity is not always reported in the literature for SOT devices, we take bulk resistivity values and apply a consistent correction to the resistivity of each layer. Assuming grain-boundary dominated scattering, the thin film resistivity can be written using the M-S model as [3]:

ρ f = ρ 0 f α (6)
where ρ f is the corrected thin film resistivity and ρ 0 is the bulk resistivity with

f α = 1 - 3 2 α + 3α 2 -3α 3 ln 1 + 1 α (7)
and

α = λR coef D(1 -R coef ) ( 8
)
where λ is the electron mean free path without grain boundaries (typically 10 -40 nm), R coef is the electron grain boundary reflection coefficient (typically 0.4 -0.6), and D is the grain size. We fitted the M-S model to the measured resistivity of our complete stack, ρ = 214 µΩ.cm, resulting in λ = 25 nm, R coef = 0.611, and where we fix the grain size to the equal to the thickness of the layer, D = d. We then applied this model to determine the resistivity of the unreported literature stacks.

  Ta(5)/Pt(2)/Co(1.0)/W(1.5)/Pt(1.5)/Ta(1.5) were prepared by magnetron sputtering on Si wafers, and patterned into 10 µm Hall bar structures by a combination of UV lithography, ion milling and lift-off. An asymmetric stack of Pt, Co, and W was chosen, as the opposite spin Hall angles of the Pt-Co and Co-W interfaces are known to maximise SOT switching efficiency 23 .

Figure 1 .

 1 Figure 1. a) SQUID-VSM magnetometry parallel (Mx) and perpendicular (Mz) to the substrate.

Figure 2 .

 2 Figure 2. a) False coloured helium ion microscopy image of the in-situ contacting of a 10 µm Hall

Figure 3 .

 3 Figure 3. a) Current driven SOT switching curves at the same doses as Fig. 2 under a bias field of

4 FIG. 2 .

 42 FIG. 2. Anomalous Hall resistance of virgin devices for a) µ0Hz and b) µ0Hx. Panel b) shows the normalised first quadrant.

Table 1 .

 1 Irradiation dose dependency of the thin film magnetic properties.

	Dose	Keff	Δ	𝜌	𝐼 5 N	𝐼 5 _	𝐽 5 N	𝐽 5 _	𝐽 5 `a	𝜂	𝜂 V
	ions nm -2	kJ m -3		µΩ.cm	mA	mA	MA cm -2 MA cm -2	MA cm -2	ps	fs
	0	537	133	214	-7.5	7.5	-6.0	6.0	459	7.1	1.2
	1	521	129	214	-7.5	7.5	-6.0	6.0	442	6.9	1.2
	20	257	64	219	-6.2	7.5	-5.0	6.0	184	4.3	2.5
	30	153	39	224	-1.0	3.3	-0.8	2.7	85	86	4.2
	50	38	9	235	-	-	-	-	-	-	-

Table 2 .

 2 Comparing Switching Efficiencies. then the smaller of the two values is taken; b) when not reported, literature values of 𝜃 => were taken from

	Stack	Jc a)	Bx 𝜃 => b) 𝛥 c)	P d)	𝜂	𝜂 V Ref.
		MA cm -2 mT			PW m -3 ps	fs
	Pt 3/Co 0.6/AlOx 1.6	78	47.5 0.07	74	2330	0.02 0.09	10
	Ta 3/Pt 5/Co 0.6/Cr 2/Ta 5	2.70	20	0.41 106	0.89	28 2.02	43
	Ta 3/Pt 5/Co 0.8/W 3/TaOx 2	2.30	5	0.244 136	0.42	69 0.78	30
	Ta 5/CoFeB 1.1/MgO 2/Ta 2	3.09	20	0.08	56	1.91	14 0.19	44
	Bi2Se3 7.4 e) /CoTb 4.6/SiNx 3	3.0	100	0.16	73	2.48	6.2 0.23	45
	Pt 5/BaFe12O19 3 f)	5.7	123	0.07 232	6.26	18 0.04	31
	Pt 4/Co 0.4/Ni 0.2/Co 0.4 /Pt 2 g)	7.1	0	0.07	13	10.2	0.74 0.72 37
	a) If 𝐽 5 -≠ 𝐽 5 _,

Acknowledgements

We thank Fabien Chevrier for technical support, the staff of the STnano nanofabrication facility.

The nanofabrication facilities (NanoFaRo) at the Ion Beam Center at the HZDR are also gratefully acknowledged.

ABBREVIATIONS

HM, Heavy metal; MRAM, magnetic random-access memory; PMA, perpendicular magnetic anisotropy; SOT, spin orbit torque

Data and Code Availability Source data for figures 1-3 and any other data that support the findings of this study, and the Python scripts used to process the data, are available on the Zenodo data repository: https://zenodo.org/communities/mami-h2020/

ASSOCIATED CONTENT Supporting Information.

The following files are available free of charge.

Supporting video 1, "MOKE_C_30ions.avi":

Ex-situ normalised magnetisation of a 10 µm wide Hall bar under a magnetic field sweep from -50 mT to +50 mT and then back to -50mT using a MOKE microscope. The central 10 x 10 µm 2 area was irradiated with a dose of 30 ions/nm 2 . Supporting text, "DEPNA_single_layer_SI.pdf" consists of:

1. Derivation of equation ( 1) using a macrospin approximation, and an example fit to the sample irradiated with 20 ions nm -2 .

2. Anomalous Hall resistance plots for 8 non-irradiated samples for µ0Hz and µ0Hx.

Discussion of the Mayadas-Shatzkes model used for applying thickness dependent corrections to bulk AUTHOR INFORMATION

Corresponding Authors *peter.dunne@ipcms.unistra.fr *c.fowley@hzdr.de

Author Contributions

The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

Funding Sources

We acknowledge the support of the Labex NIE 11-LABX-0058_NIE within the Investissement

Notes

The authors declare no competing financial interest.