Nadia Soulmi

Ana-Gabriela Porras-Gutierrez

Natalia E Mordvinova

Oleg I Lebedev

Cécile Rizzi

Juliette Sirieix- Plénet

Frédéric Lantelme

Henri Groult

Damien Dambournet

Laurent Gaillon

Sn(TFSI) 2 as Suitable Salt For the Electrodeposition of Nanostructured Cu 6 Sn 5 -Sn Composite obtained on Cu electrode in Ionic Liquid

The preparation of binder and carbon-free electrodes is of great interest owing to higher energy density. In this scope, electrodeposition is a suitable method that can be easily scale up providing that suitable chemical reactants are used. In this work, we used highly soluble Sn precursor based on TFSI counter ion dissolved in [EMIm + ][TFSI -] ionic liquid. The use of similar anionic groups in both Sn precursor and solvent allowed to avoid impurities typically encoureted when using Sn chloride precursor. The resulting solution was characterized by cyclic voltametry using either inert (Mo) and reactive substrate (Cu). In both cases, the electrodeposition occured in a diffusion controlled process. In the case of Cu, however, a Cu-Sn alloy that is η-Cu6Sn5 phase was identified. A FIB cross section revealed that Cu and Sn interdiffused and that no epitaxial growth occured. Prolonging the electrodeposition time favored the formation of β-Sn because Cu could no longer diffuse within the deposited structure. This resulted in a mixed β-Sn/η-Cu6Sn5 composite whose proportions depended on the deposition duration time. The deposited samples were directly assembled in lithium coin cells to characterize the capacity and cyclability of binder and carbonfree electrodes. Increasing the mass of deposited sample degraded the capacity and reversility of the system which was explained by a lower ability to accommodate the volume variation occurring during the electrochemical lithiation/delithiation process.

Introduction

Sn-based materials are challenging the practical widely used graphite anode materials for Li-ion batteries because of their higher lithium storage capacity (up to 991 mAh g -1 for a Li4.4Sn composition). However, due to dramatic volume expansion (up to approximately 300 %) during the charge/discharge cycling, Sn-based electrode provides only poor cycle life caused by its mechanical disintegration. 1,2 Downscaling the size of the material is one approach of choice to overcome this issue by buffering these expansion/contraction processes. 3,4 However to reach higher conductivity between the grains without loss of mechanical properties, several strategies associated to electrode formulation, embedding and processing have been used. 5 Sn/C based composites electrode, with high binder content (to maintain adhesion to the current collector and cohesion within the composite) is currently required to rise good electrochemical performances. 6,7 Likewise considerable interest have also been given to intermetallic Sn-based materials, where an inactive metal M vs. Li + /Li was used to modulate the volume change (e.g. M =Cu, Ni, Co, Fe, etc.). [8][9][10][11] Instead of composite electrode, the use of pure metallic Sn electrode could be employed by combining nanostructuration with a thin film approach. [12][13] Such electrodes can be obtained more easily by an electrodeposition process that uses an electric field to reduce ionic species onto a specific substrate. Typically, the deposition of metallic species can be performed directly on a collector of current with the advantages that tuning the experimental conditions (i.e. solvent, electrode substrate, applied potential or current density, metallic salt precursor, additives, working temperature, etc.) permit to modulate the morphology of the deposit. 14 Henceforth, electrodeposition in ionic liquids (ILs) media is of great interest owing to their wide electrochemical potential window stability, high deposition current efficiency and furthermore, compared to conventional aqueous electrolyte the deposit does not suffer from H2 evolution and metal oxide formation. 15,16 It has already been demonstrated that Sn can be electrodeposited from various ILs electrolytic solutions, 17 among which imidazolium 18 and pyrrolidinium 19 based ionic liquids solutions are most often used. Electrodeposition of Sn has also been achieved from ILs mixture with AlCl3, 20,21 or ZnCl2, 22 or even in deep eutectic solvent. 23 Even if experimental deposition conditions modulate the morphology, the thickness, and the nanostructuration of the deposited layer, in this kind of synthesis, they are also dramatically influenced by the choice of the anion/cation combination of the IL 19,[24][25][26][27] Indeed the physicochemical properties of the solvent (conductivity, solubility, viscosity) impact the deposition potentials the and the electrodeposition mechanism. In the case of Sn deposit, some studies have already underline the influence of the anion 24,[26][27][28][29] through viscosity and stabilization of the precursor tin species from one part. From another part, the influence of the ionic liquid cation through viscosity and likely its different adsorption /interactions with the substrate or the nuclei was also demonstrated [26][27][28] . Different complexes of Sn 2+ with various ILs have been identified by infrared spectroscopy, also underlining that both anion and cation of the IL play a role in the deposition process. 26 One of the main drawback when using ILs for Sn electrodeposition is the low solubility of the generally used halide precursor salt (SnCl2). 30 Moreover, the anhydrous salt should be privileged rather than its stable hydrated form (i.e. SnCl2.2H2O) since water is undesired in the electrolytic media. Al-Salman et al. 28 showed that the liquid (but very hygroscopic) solution of SnCl4 could also be used. Otherwise, when using SnCl2, due to is strong donor ligand properties [DCA - ] (dicyanamide) anion based ILs were demonstrated to provide highest solubility of this metal salt precursor, [31][32][33] and therefore has been thoroughly studied for Sn electrodeposition. [DCA -] anion is generally associated to [EMIm + ] (1-ethyl-3methylimidazolium) cation because this IL has a low viscosity and high conductivity. 26,27,31 Therefore, ILs with [TFSI -] (bis(trifluoromethanesulfonyl) imide) anion are generally preferred because they are known for their low hydrophilicity and hence low water uptake compared to [DCA -] anion, low viscosity and good conductivity. But through the fact that [TFSI -] is a non-coordinating anion, such ILs suffer from low solubility of halide salts. In one of our previous study, to overcome this issue, we had used anodic electrodissolution of a metallic Sn wire to introduce the desired amount of Sn 2+ in the solution. 27 One of the additional benefits of this method was that neither halides nor other counter ions were introduced in the electroplating bath. Therefore, the electrogenerated Sn 2+ cations could only complex with the anion of the IL used.

In the present study, [EMIm + ][TFSI -] IL was used to prepare solutions incorporating Sn(TFSI)2 salt, as plating solutions to electrodeposit Sn nanostructures. The salt Sn(TFSI)2 was chosen because of its good solubility in ILs composed with the same counter-anion. Sn nanostructures obtained directly on the current collector (active copper substrate), when varying electrodeposition conditions (duration, potential, concentration of Sn precursor) were shown to influence the electrochemical performances (charge/discharge capacity and cycling stability) of such a synthesized binder-free negative electrode.

Experimental

Chemicals.

1-Ethyl-3-methylimidazolium bis(trifluoromethanesulfonyl) imide ([EMIm + ][TFSI -]), 99.5% and the electrolyte 1M lithium hexafluorophosphate in ethylene carbonate : ethyl methyl carbonate (1:1 vol.%) with 2%wt. fluoroethylene carbonate were obtained from Solvionic. Sn(II) bis(trifluoromethyl sulfonyl)imide (Sn(TFSI)2) was obtained from Alfa Aesar, and the acetonitrile (ACN), HPLC grade 99.9 %, was obtained from Sigma-Aldrich.

All the chemical were used as received and stored in glove box filled with argon (H2O and O2 content < 1 ppm).

Preparation and electrochemical characterization of Sn electrodes.

All electrochemical measurements were performed in the glove box filled with argon (H2O and O2 content < 1 ppm). A three-electrodes electrochemical cell controlled by a potentiostat (PAR 273A) was used. Cu (A = 0.785 cm²) or Mo (A = 0.503 cm²) foils were used as working electrode (they were cleaned with 1M H2SO4 solution, rinsed with water and dryed prior to uses). The counter electrode was a Pt wire and a Sn bar was used as reference electrode.

Cyclic voltammetry (CV) curves were recorded at different scanning rates (0.005 -0.200 Vs -1 ) in an electrochemical window between -0.6 and +0.4 V vs. Sn 2+ /Sn. The electrodeposition of Sn was performed potentiostatically at -0.3 V vs. Sn 2+ /Sn for a given duration. After the deposition, the samples were thoroughly cleaned in the glove box to avoid oxidation of the surface, with acetonitrile to remove all the IL present at the surface of the electrode. 34 Characterization Methods.

The X-ray diffraction (XRD) was performed using a Rigaku Ultima IV X-ray diffractometer with a Cu Kα radiation (λ= 1.54059 Å) with a scan rate of 0.1°min -1 between 10° and 80° (2θ).

Scanning Electron Microscopy (SEM) measurements were performed with a SEM-FEG ZEISS ultra 55 equipped with or coupled to a Bruker QUANTAX EDS analysis system.

The microstructure of the Sn-based electrode was examined with transmission electron microscopy (TEM). Two type of TEM samples were prepared. The first one was obtained by scrapping the electrode surface and further depositing the material together with ethanol on holey Ni holey carbon grid. In order to investigate substrate/film interface and possible interdiffusion, cross-section samples were prepared by focused ion beam (FIB) techniques in Helios Nano Lab 660FIB. Electron diffraction (ED) and high resolution TEM (HRTEM) experiments were performed on Tecnay G2 30 UT microscope operated at 300 kV and having 0.17 nm point resolution. High angle annular dark field scanning TEM (HAADF-STEM) imaging and energy dispersive x-ray spectroscopy (EDX) elemental mapping were recorded on a double aberration corrected JEM ARM200F cold FEG electron microscope operated at 200 kV and equipped with large solid-angle CENTURIO EDX detector.

Electrochemical Characterization.

The electrochemical performances as negative electrode material in Li-ion battery of the as-prepared Sn-based electrodes were carried out with CR2032-type coin cells. The electrolyte was 1 M LiPF6 dissolved in a mixture of ethylene carbonate (EC) and ethyl methyl carbonate (EMC) (1:1 vol. %) with 2%wt. of fluoroethylene carbonate (FEC). The cells were assembled with lithium metal as the anode and were tested in the voltage range of 2.0 -0.05 V vs. Li + /Li at different current densities.

Results and Discussion

Cyclic voltammetry study using inert (Mo) and reactive (Cu) electrodes.

A solution containing Sn(TFSI)2 solubilized in [EMIm + ][TFSI -] was characterized to investigate the electrochemical behaviour using two different working electrodes that are copper and molybdenum. Cu is commonly used as a current collector for negative electrode materials. It is, therefore, an interesting substrate for Sn electrodeposition as it allows to prepare an electrode ready to be used in a lithium cell. However, it can form alloys with Sn at room temperature which makes it a reactive substrate/electrode. [35][36][37] On the other hand, Mo does not form alloys with Sn and is therefore a non-reactive substrate suitable to characterize the electrochemical behaviour of Sn 2+ in solution.

Cyclic voltammetry was used to characterize the electrochemical properties of a 25 mM solution of Sn(TFSI)2 in [EMIm + ][TFSI -]. Figure 1 presents the cyclic voltamogramms obtained using Mo and Cu as working electrodes at a scan rate of 0.05 Vs -1 starting from the open current potential (EOCP) toward cathodic and reversed anodic sides. The potential window was -0.6 to +0.4 V which is within the electrochemical stability window of [EMIm + ][TFSI -] ionic liquid (Supporting Information, Figure S1) The EOPC was around +0.2 V. In both electrode materials, the shape of the reduction peak is characteristic of an electrodeposited solid that is insoluble. 38 On Mo electrode, a cathodic peak was observed at -0.17 V vs. Sn 2+ /Sn and was assigned to the reduction of Sn 2+ to Sn 0 . On the reverse potential scan, Sn was re-oxidized at around +0.10 V. The current loop during the cathodic and the anodic scan indicated that the electrodeposition of Sn into Mo required an overpotential to initiate nucleation and growth, as already observed in diverse ILs. 19,27 On the Cu electrode, the reduction peak appeared at a potential of -0.08 V, which was higher than that observed in Mo electrode. Moreover, the loop feature observed in the case of Mo was absent here. This meant that Cu electrode did not require an over potential to initiate growth of Sn due the interaction between Sn and the substrate. After the first reduction peak, the current still decreased smoothly which might indicate the occurrence of other chemical reactions coupled to the primary electrochemical process and/or the formation of others composition of Sn-Cu alloys. The occurrence of these additional reactions was also suggested by the ratio of the total anodic to cathodic charge (Qa/Qc) which was lower than unity indicating that the deposited material was not totally electrochemically removed during the positive scan. Cyclic voltammetry was performed at different scan rates. On the Mo electrode (Fig. 2a), an increase of the scan rate led to an increase of the over-potential. In contrast, the oxidation peak remained relatively unchanged. Due to the overpotential electrocrystallisation, it was difficult to compare and to interpret the variation of ∆𝐸𝐸 𝑝𝑝 = 𝐸𝐸 𝑝𝑝𝑝𝑝 -𝐸𝐸 𝑝𝑝𝑝𝑝 in terms of reversibility or irreversibility of the reaction.

On the Cu electrode (Fig. 2b), the reduction peak potential did not vary significantly with an increasing scan rate. For all scan rates, the reduction current started at the equilibrium potential of the redox couple. The overvoltage at the zerocrossing, η 0 = -0.05𝑉𝑉 , of backward curves was close to its theoretical value predicted for reversible systems. 39,40 The CVs, overall, indicated an electrochemical behaviour largely influenced by the nature of the working electrode. However, a linear relationship between the cathodic peak current density and the square root of the scan rate �𝐼𝐼 𝑝𝑝𝑝𝑝 = 𝑓𝑓�𝑣𝑣 1 2 ⁄ ��, in Mo and Cu electrodes, indicated that the reduction of Sn 2+ were governed by diffusion-controlled processes 8,41,42 (Figure S2) as follows :

𝐼𝐼 𝑝𝑝𝑝𝑝 = -0.6104𝑛𝑛𝑛𝑛𝑛𝑛𝐶𝐶 0 * � 𝑛𝑛𝑛𝑛 𝑅𝑅𝑅𝑅 𝑣𝑣𝐷𝐷 𝑆𝑆𝑛𝑛 2+ (1)
where 𝐼𝐼 𝑝𝑝𝑝𝑝 is the current intensity of the cathodic peak, 𝐶𝐶 0 * the concentration of the electroactive species and 𝐷𝐷 𝑆𝑆𝑛𝑛 2+ the diffusion coefficient of the Sn 2+ species. It should be noted that both plots did not pass through the origin. The non-zero intercept may be assigned to an adsorption phenomena and some surface chemical coupled reactions. 29,43 The diffusion coefficient of Sn 2+ species, 𝐷𝐷 𝑆𝑆𝑛𝑛 2+ can be determinated according to equation ( 1) for a reversible (diffusion-controlled) deposition process on a massive electrode. Because Cu 27,29,30 probably due to higher viscosity for [EMIm + ][TFSI -] that induced a lower mobility of Sn 2+ . Moreover, because the coordination state of Sn 2+ is modified by the counter ion solvation force of the ionic liquid anion versus the counter-ion of Sn precursor, different solvation environments influencing Sn 2+ mobility are expected.

Electrodeposition and characterization of deposited solids.

To better understand the reactivity of Sn 2+ on the Cu electrode, a cathodic deposition was performed at -0.3 V, a potential located after the reduction peak, in the mass-transfer controlled region for different duration times (5 min, 30 min, 1h and 2h). In such experiments, the mass of deposited solid scales linearly with the time of the electrodeposition (Figure S3). In general, the use of ILs as electrolyte prevents the occurrence of side reactions and therefore allows high efficiency to be obtained. 15 Increasing the duration time led to an increasing amount of deposited solids as revealed by the colour changes of the substrate going from black for lower mass to bright metallic for higher deposited mass (Figure S4). Moreover, all samples showed good adhesion to the substrate.

XRD measurements were performed on the deposited solids obtained at different duration times (Figure 3). Overall, the XRD patterns of the electrodeposited solids could be indexed with two phases exhibit Bragg peaks characteristic of β-Sn (space group: I41/amd) and η-Cu6Sn5 (space group: P63/mmc). Prolonging the duration time led to an intensity increase of the [200] and [101] lines of β-Sn. At t > 1h, β-Sn became the predominant phase suggesting that η-Cu6Sn5 was formed at the early stage of the deposition and then, once Cu and Sn could no longer interdiffuse, β-Sn started to form. This was confirmed by energy-dispersive X-ray (EDX) spectroscopy measurements and associated Sn/ Cu ratio that showed that increasing the deposition time led to an increasing and decreasing contents of Sn and Cu, respectively (Figure S5). To extend the comprehension on the formation of Cu-Sn alloy, a FIB cross section TEM sample of a Cu electrode was prepared on which a solid was electrodeposited during 30 min. HAADF-STEM image of the corresponding cross section is shown was obtained on single [001] oriented grain and allowed to identify the η-Cu6Sn5 phase which was further confirmed by corresponding FT pattern. In Figure 4c, high-resolution HAADF-STEM image of Cu/deposited solid interface did not evidence any epitaxial growth between the substrate and the deposited solid. This suggested that Cu and Sn interdiffused to form the η-Cu6Sn5 phase. EDX elemental mappings shown in Figure 4d revealed that Sn was present both at the surface of the Cu electrode and within larger grains, i.e., around 200 nm. The HAADF-STEM image and Cu K mapping clearly showed the occurrence of attachment point from which these large particles were growth. This highlighted the diffusion of Cu coming from the substrate.

The morphology at the particle level was investigated by HAADF-STEM imaging and EDX mapping on particles obtained by scrapping the surface of the Cu electrodes. Typical examples of TEM images along with EDX mapping of the particles observed are shown in Figure 5a. For all samples, we observed particles featuring two different sizes, small ones in the range of 3-15 nm and larger ones with sizes ranging from 100-300 nm after 5 min, 200-400 nm after 30 min, 300-1000 nm after 1h and 200-800 nm after 2h. EDX mapping performed on the samples obtained after 5 min and 2h of deposition, revealed the presence of both Sn and Cu at the particle level. In the case of the 5 min deposition, we observed generally a homogenous distribution of Sn and Cu, in agreement with the predominant formation of η-Cu6Sn5 phase. However, small areas of Sn-O phase could also be observed. In the case of the 2h deposition, we observed Cu and Sn separation. While Sn was present all over the particle, Cu was not (see SI Figure S6 for other examples), which was in agreement with the formation of β-Sn as shown by XRD analysis. The formation of two phases was also confirmed by EDX mapping performed on a cross section for a sample prepared for 2h (Figure 5b). Scanning electron microscopy was further used to visualize the upper scale of the particle's organization (Figure 6). After 5 min of deposition, the solid is composed of nanoparticles. After 30 min and 1h, a network of interconnected particles was observed. This interconnected organization disappears after 2h of deposition suggesting a different growth mechanism.

Electrochemical properties.

The electrodeposited solids on Cu substrate provided electrode materials free of carbon and binder additives. They were directly assembled in coin-type half cells to investigate their properties vs. Li + /Li.

The multi-step lithiation delithiation processes on Sn electrode 3,29,44 as on Sn based alloy electrode 29,[45][46][47][48] are already well documented, and it has already been concluded that CV profiles and peak potentials are in both cases characteristic of electrodes containing metallic tin. Especially the plateau of potential attributed to the lithiation of Cu6Sn5 arise in the same domain of potential where intermetallic tin-lithium phases prevails. The potentials of the peaks in charge are similar in both cases to Sn-Li de-alloying process. Comparison of CV curves obtained on electrodes with lower and higher amounts of tin (Figure S7) does not permit to distinguish the composition of the electrode.

Figure 7 presents the galvanostatic discharge/charge curves obtained on the series of electrodeposited samples. The mass of the active materials determined by the measured current during the electrodeposition process yielded 0.04 mg for 5 min, 0.085 mg for 30 min, 0.15 mg for 1h and 0.29 mg for 2h. Cells were cycled between 0.05 V and 2.0 V (vs. Li + /Li) at a current density of 100 mA g -1 .

During the first discharge, all the samples featured an irreversible plateau at around 1.5 V vs. Li + /Li. Such a plateau was also observed in other studies and was assigned to the decomposition of the electrolyte (EC degradation) and the formation of the solid electrolyte interface (SEI) layer. 49 We noted that the plateau length decreased with the increase of the deposition duration time. The occurrence of isolated nanoparticles obtained after 5 min might favour the decomposition of the electrolyte due to higher surface area.

During the first cycle, all samples presented a large irreversible capacity. The charge capacity decreased with the increase of the deposition duration time with the corresponding values obtained for the first charge of : 900 (5 min), 400 (30 min), 265 (1h) and 190 (2h) mA hg -1 . For all the samples, the subsequent discharge and charge curves showed similar shape suggesting that the first cycle allowed to activate the electrode.

The cells were subsequently left for prolonged cycles under 100 mA g -1 . Overall, the capacity slightly decreases upon cycling. The solid prepared for 5 min presents the highest capacity with 510 mAh g -1 measured after 50 cycles. As shown in Figure 8, the solids prepared at t>5min present lower capacities. To determine the rate performance of this electrode material, C rate was increased stepwise from C/10 to 5C and then decreased to C/10, and cycled for 10 cycles at each rate (Figure S8). The best performing electrode is the thinnest at any rate, which correlated well with the high increase of lithium storage capacity with decrease of Sn mass. At high rate of 5C, 33% of the total capacity is recovered for the thinnest deposit. It was observed that this recovery of the total capacity increase either because the mass of deposited Sn decrease (25% for the thickest deposit) or because of the improve retention capacity brought by Cu-Sn alloys.

The difference in electrochemical properties for the series of samples can be related to two interconnected factors that are the particles morphology and the composition. A less dense film will allow to better accommodate the volume expansion experienced during the insertion/de-insertion reactions of lithium. Moreover, as compare to β-Sn, Cu-Sn alloys benefit from a lower volume expansion upon reactions with lithium. The evolution of the morphology after cycling was compared between the deposited samples obtained after 5 min and 1h (Figure 9). The first sample featuring the less dense structure allowed to better accommodate the volume variation by preserving the porous structure. After cycling, the sample prepared for 1h became much more compact which probably induced a blocking effect toward the lower part of the electrode.

Conclusions

We investigated the electrodeposition of Sn using soluble Sn precursor based on TFSI -counter ion dissolved in [EMIm + ][TFSI -] ionic liquid. The electrochemical behaviour of the Sn-ionic liquid solution was characterized by cyclic voltammetry using either inert (Mo) or reactive substrate (Cu). The diffusion coefficient of Sn 2+ in Mo electrode was found to be 1.61x10 -8 cm 2 s -1 which is comparable to values obtained with other type of ionic liquids. The electrodeposition onto Cu led to the formation of a Cu-Sn alloy that was η-Cu6Sn5 phase as identified by X-ray and electron diffraction. By analysing a FIB cross section on electrodeposited sample, we emphasized the interdiffusion of Cu and Sn atoms without any evidence of epitaxial growth. When prolonging the electrodeposition time, Cu atoms could no longer reach the upper surface favouring the formation of β-Sn. Tuning the reaction duration times enabled to tune the proportions of β-Sn/η-Cu6Sn5. The electrochemical properties of the binder and carbon-free electrodes were investigated in lithium coin cells. Increasing the mass of deposited sample degraded the capacity and reversibility of the system which was explained by a lower ability to accommodate the volume variation occurring during the electrochemical lithiation/delithiation process. This work is likely to be relevant to other types of elements providing similar salts.

Figure 1 .

 1 Figure 1. Cyclic voltammetry of 25 mM Sn(TFSI)2 in [EMIm + ][TFSI -] on Cu and Mo electrodes at a scan rate of 0.05 Vs -1 .

Figure 2 .

 2 Figure 2. Evolution of the cyclic voltammograms of a 25 mM Sn(TFSI)2 in [EMIm + ][TFSI -] solution at different scan rates on a) Mo, and b) Cu electrodes.

Figure 3 .

 3 Figure 3. XRD patterns performed on electrodeposited samples obtained at different deposition times.

inFigure 4 .

 4 Figure 4. a) Low magnification HAADF-STEM image of FIB cross section of a Cu electrode on which a solid was electrodeposited during 30 min. b) High resolution [001] HAADF-STEM image of single grain showing η-Cu6Sn5 phase and corresponding FT pattern given above c) HAADF-STEM image showing the interface between the Cu electrode and electrodeposited solid marked in (a) by red square (arrows). d) EDX elemental mapping of Cu K (green), Sn L (red), O K (blue) and overlaid colour map.

Figure 5 .

 5 Figure 5. a) HAADF-STEM image of the particles obtained by scrapping the electrode surface obtained on the electrodeposited solids for 5 min and corresponding EDX elemental mapping of Cu K(green), Sn L(red), O K(blue) and overlaid colour image. b) HAADF-STEM image of FIB sample of a Cu electrode on which a solid was electrodeposited during 2h and corresponding EDX elemental mapping of Cu K(green),Sn L(red), O K(blue) and overlaid colour image .

Figure 6 .

 6 Figure 6. Scanning electron micrographs of the electrodeposited samples prepared at different duration time.

Figure 7 .

 7 Figure 7. Galvanostatic discharge/charge curves obtained for the first three cycles on the series of electrodeposited samples. a) 5 min, b) 30 min, c) 1h and d) 2h. The cells were cycled under 100 mA g -1 between 0.05 and 2.0 V vs. Li + /Li.

Figure 8 .

 8 Figure 8. Long term cycling of the electrodeposited samples. The cells were cycled under 100 mA g -1 between 0.05 and 2.0 V vs. Li + /Li.The evolution of the morphology after cycling was compared between the deposited samples obtained after 5 min and 1h (Figure9). The first sample featuring the less dense structure allowed to better accommodate the volume variation by preserving the porous structure. After cycling, the sample prepared for 1h became much more compact which probably induced a blocking effect toward the lower part of the electrode.

Figure 9 .

 9 Figure 9. Scanning electron micrographs of electrodeposited samples (after 5 min and 1h) on Cu substrate before and after 50 cycles vs. Li + /Li.

Acknowledgements

O.L acknowledges the "Agence Nationale de la Recherche (ANR)" in the framework of the "Investissements d'Avenir" with the reference "ANR-11-EQPX-0020" (GENESIS EQUIPEX) for the financial support that allowed the acquisition of the FIB system (HELIOS Nanolab 600 from FEI) for TEM sample preparations.

Conflicts of interest

There are no conflicts to declare.

Supporting Information

Sn(TFSI)2 as Suitable Salt For the Electrodeposition of Nanostructured Cu6Sn5-Sn Composite obtained on Cu electrode in Ionic Liquid

Nadia Soulmi a , Ana-Gabriela Porras-Gutierrez* a , Mordvinova Natalia b , Oleg Lebedev b , Cécile Rizzi a,c , Juliette Sirieix-Plénet a,c , Henri Groult a,c , Damien Dambournet a,c and Laurent Gaillon* a,c


Ségolène Palussière

Jérémy Cure

Andrea Nicollet

Pierre Fau

Katia Fajerwerg

Myrtil L Kahn

Alain Estève

Carole Rossi

The role of alkylamine in the stabilization of CuO nanoparticles as a determinant of the Al/CuO redox reaction

We report on a new strategy to synthesize Al/CuO nanothermites from commercial Al and ultra-small chemically synthesized CuO nanoparticles coated with alkylamine ligands. These usual ligands stabilize the CuO and prevent them from aggregation, with the goal to enhance the interfacial contact between Al and CuO particles. Using a variety of characterization techniques, including microscopy, spectroscopy, mass spectrometry and calorimetry (ATG/DSC), the structural and chemical evolution of CuO nanoparticles stabilized with alkylamine ligands are analyzed upon heating. This enables us to depict the main decomposition processes taking place at the CuO surface at low temperature (< 500 °C): the ligands fragment into organic species accompanied with H2O and CO2 release, which promotes the CuO reduction into Cu2O and further Cu. We quantitatively discuss these chemical processes highlighting for the first time the crucial importance of synthesis conditions that control the chemical purity of the organic ligands (octylamine molecules and derivatives such as carbamate and ammonium ions) on the nanothermite performances. From these findings, an effective method to overcome the ligands-induced CuO degradation at low temperature is proposed and Al/CuO nanothermite reaction is analyzed, in terms of onset temperature and energy released. We produce original structures composed of aluminium nanoparticles embedded in CuO grainy matrixes exhibiting onset temperature ~200 °C below usual Al/CuO onsets, having specific combustion profiles depending on synthesis conditions, while preserving the total amount of energy released.

Introduction

Nanothermites have received increasing interest over the two last decades because of their high volumetric energy densities (up to 16 kJ/cm 3 ), adiabatic flame temperature (> 2600 °C), and high reaction (burn) rates. They feature an exothermic redox reaction involving aluminum (Al) and a metal oxide, which results in a stable product after reaction. Different nanothermites with specific shapes (nanowires, 1-3 core-shell, 4,5 nanoparticles (NPs) mixture, 6-13 nanofoils 14 ….), length scale (10 to few 100 nm), have been actively investigated for a wide range of applications including explosive additives, propellant rate modifiers, 15 welding, 16 and also for micro initiation and environmentally clean primers. 17 Altogether, this research effort has experimentally demonstrated that a variety of different combustion effects can be obtained by manipulating the reactive system (Al and oxide) at the microscale, which cannot be achieved with bulk materials only. Typically, maximizing Al and oxidizer interfacial contact is paramount for achieving optimum reaction kinetics (combustion efficiency and ignitability), and numerous approaches have been reported in the literature to direct and optimize the nanoparticles assembly. [START_REF] Séverac | [END_REF] It can be done in inorganic, aqueous or organic solutions, using aerosols or sol-gel methods. 18- 20 Kim et al. proposed a method in which the Al and Fe2O3 self-assembly is controlled by the electrostatic forces which exist between charged aerosol particles. 21 Another alternative is to use complementary DNA strands to control the structure of the nanothermites by guiding the self-assembly of the nanoparticles to enhance the energy release. 22-24 A recent approach demonstrates that Al and oxide nanoparticles assembled on functionalized graphene sheets form highly reactive macrostructures. 25 In this work, a novel stabilizing strategy has been developed to synthesize Al/CuO nanothermites from Al and ultra-small CuO NPs functionalized with alkylamine ligands.

These usual ligands stabilize the CuO NPs and prevent them from aggregation, which can be beneficial to enhance the coating of Al particles with smaller CuO NPs. Thus synthesized CuO NPs are then mixed in hexane with commercial Al NPs to form the nanothermite composite following an assembly process. The resulting nanothermites exhibit a weak Al+CuO redox reaction preceded by several low temperature exotherms that are investigated in the present study. The reduction of CuO NPs capped with alkylamine ligands into an intermediate Cu2O phase is observed at ~250 °C before further reduction to metallic copper at ~ 450 °C i.e. prior to the thermite reaction onset.

This contrasts with what is typically observed in bulk CuO in both isothermal and ramping reduction conditions, which requires high temperature to get reduced (527 and 877 °C for CuO Cu2O and Cu2O Cu, respectively). This work provides a detailed understanding of the early CuO reduction process due to the oxidation of the surrounding ligands before presenting an effective method to overcome it. After the chemical synthesis, the CuO are annealed at 500 °C (ramping 5 °C.min -1 ) under oxygen atmosphere to eliminate organics and re-oxidize the reduced CuO. These annealed CuO nanoparticles mixed with commercial Al exhibit lower onset temperature compared to nanothermites produced from CuO commercial nanoparticles while releasing the same total energy.

Materials and methods

Chemicals and reagents

All chemicals products are ordered from Sigma Aldrich, excepted copper monochloride (Acros Organics). Commercial aluminum (Al, 99.9+%) and copper oxide (CuO, 99%) nanoparticles are supplied by US research nanomaterials (Texas, USA) and have an average particle size of 96 ± 43 nm and 31 ± 15 nm, respectively. The particles are largely spherical. The active Al mass ratio determined by thermogravimetric analysis (TGA) is 69 %. 26 A solution of methyllithium (1.6 M in Et2O, 0.02 mol, 1 eq.) is added dropwise by syringe into a solution of 1,3-diisopropylcarbodiimide (0.02 mol) in 30 mL of Et2O at -30°C. The solution is stirred for 4 h at room temperature. The resulting solution is then transferred to a copper monochloride (1.98 g, 0.02 mol) solution in 25 mL of dried Et2O. The solution is stirred for 18 h under argon atmosphere. The solvent is removed under reduced pressure, and the resulting solid is extracted with dried pentane (75 mL). The pentane-extracted solid is filtered through a glass frit to afford a pale yellow solution.

Synthesis of N,N'-diisopropylacetamidinate copper [Cu( i Pr-Me-AMD)]2:

A saturated solution is formed by removing part of the pentane under reduced pressure. The powder is recrystallized in 5-10 mL of pentane and stored at 7 °C for 3 days. Colourless crystals [Cu( i Pr-Me-AMD)]2 are obtained after removing the pentane under reduced pressure, and stored under inert atmosphere at 7 °C (30 %).

Synthesis of CuO nanoparticles by wet chemistry

Two batches of CuO NPs are produced following an organometallic route described previously from a N,N'-diisopropylacetamidinate copper precursor and octylamine (OA) as a stabilizing agent. The reaction is complete after 16 h as confirmed by NMR analyses reported in a precedent work. 27 CuO NPs grown either by hydrolysis and/or oxidation in the ambient air (named CuO_air thereafter), or by oxidation under O2 atmosphere (named CuO_O2 thereafter). For CuO_air, the precursor [Cu( i Pr-Me-AMD)]2 (0.24 mmol, 1 eq.) is mixed with octylamine (1.95 mmol, 8 eq.) under argon atmosphere, and then exposed to ambient air at room temperature. This leads to the reaction of OA with H2O and CO2 molecules, resulting in the formation of carbamate/ammonium species (RNCO2 -/RNH3 + ). For CuO_O2, the as prepared mixture [Cu( i Pr-Me-AMD)]2 / OA contained in a flask, is placed into a reactor and purged by a O2 line at 0.3 bar, giving a black/brown liquid after 1 h. A summary of the materials used in this work is provided in Table 1.

Table 1. Database of Al and CuO nanoparticles used in the experiments:

Preparation of Al/CuO nanothermites

To prepare the nanothermites, chemically synthesized CuO NPs (CuO_air and CuO_O2) are first washed by centrifugation (8700 rpm, room temperature, 10 min) with ethanol and acetone in order to remove excess of ligands that are not coordinated onto the nanoparticles surface. The mass losses for CuO_air and CuO_O2 correspond to 72 and 87 % of their initial mass, respectively. Then, they are mixed with commercial Al nanoparticles following the procedure depicted in Figure 1 Al/CuO_air, Al/CuO_O2 and Al/CuO_com thereafter. Al/CuO_com constitutes the reference system. All nanothermite mixtures are sonicated in a hexane suspension for a total of 10 min in an ultra-sonic bath before being dried at 70 °C under vacuum.

Functionalization of commercial CuO NPs (CuO_com)

To check the reduction effects of AO ligands on CuO_com, a batch of CuO_com nanoparticles (500 mg, 6.29 mmol) are functionalized with OA (18.3 mg, 0.14 mmol) by mixing them and stirred for 1 h in the ambient air. The mixture is centrifuged (8700 rpm, RT, 10 min) and CuO_com@OA nanoparticles are retrieved before chemical and thermal analysis.

Characterization of CuO and Al NPs

The morphology, particle size, agglomeration are observed by TEM (Transmission Electron Microscopy) using a JEOL JSM operating at 100 kV. Analyses of the images are performed using ImageJ. The thermal decomposition of CuO NPs are characterized under a ramping heating profile @ 10 °C.min -1 by Thermogravimetric/Differential Scanning Calorimetry (TGA/DSC) using a Mettler Toledo Thermogravimetric analyzer. Experiments are performed with ~5 mg of CuO nanopowders in platinum crucibles over a temperature ranging from ambient to 500 °C.

The chemical composition of the materials is characterized using a SEIFERT XRD 3000 TT X-Ray Diffractometer (XRD) with Cu-Ka radiation ( = 1.54059 Å) fitted with a diffracted-beam graphite monochromator. The diffraction angle (2) is scanned from 10 to 80°. The decomposition of CuO particles is investigated using MS (mass spectroscopy) (Pfeiffer Omnistar 1-200 uma), with temperature range from ambient temperature to 500 °C with a heating rate at 10 °C.min -1 . All measurements are performed under Ar flowing gas (99.999%).

Characterization of nanothermites Al/CuO

The thermal properties of nanothermites are characterized by DSC using a NETZSCH DSC 404 F3 pegasus device equipped with a DSC-Cp sensor type S. ~5 mg of nanothermites are analyzed in a platinum crucible over a temperature ranging from ambient to 1000 °C with a heating rate of 10 °C• min -1 .

Additionally, a SEM (Scanning Electron Microscopy) Helios 600i instrument coupled with an EDX (Energy Dispersive X-ray spectrometry) spectroscopy system to conduct elementary analysis is used to image nanothermites. For the observations, samples are prepared by depositing and evaporating one droplet of the nanoparticles solution onto a silicon wafer.

Results and discussion

Chemical and structural characterization of CuO and Al NPs

Representative TEM images of the synthesized CuO NPs (CuO_air and To further investigate the interactions between CuO particle surfaces and ligands, ATR-FTIR absorbance spectroscopy is performed and results are shown in Figure 3b.

The stretching modes of the CuO bonds in CuO_com are clearly detected at low wavenumbers, 590 cm -1 , 523 cm -1 and 475 cm -1 . However, as for XRD analyses, some differences are found between synthesized CuO NPs in air and controlled O2

atmosphere. The low intensity of signal between 3000 and 700 cm -1 obtained for CuO_O2 is clearly explained by the fact that only a small quantity of ligands remains absorbed on CuO surface after washing. Nevertheless, a weak peak can be detected around ~2918 cm -1 corresponding to C-H vibrational modes, which confirms the presence of ligands around CuO particles.

Focusing on CuO_air, in addition to C-H modes from the alkyl groups (CH2) at 2953, 2918 and 2851 cm -1 , characteristic peaks at 3379, 3320, 3141 cm -1 and 1574 cm -1 are associated to N-H and N-H2 and C-O2 modes. Those bands are characteristic to the carbamate/ammonium species (R-NH3 + + R-NHCO2 -), formed when OA is exposed to carbon dioxide and water present in ambient air. Supported by Bacsik et al. 33 who investigate properly the band positions of chemical functions of the carbamates ions pairs via IR spectroscopy, by heating and looking at the desorption of carbon dioxide, we conclude that specific peaks between 1600 and 700 cm -1 are attributed to the ammonium and carboxylate groups of the ammonium carbamate species at different stretching symmetry. This reaction does not occur in the case of CuO_O2 synthesis, i.e. where the atmosphere does not contain neither carbon dioxide nor water molecules. These ion pairs may bind to CuO nanoparticles with a higher stabilization than neutral amine terminations being largely eliminated through washing, thus playing a dominant role in the ligand shell stabilization. 32 In summary, although in both synthesis routes, ligands include OA and amidine resulting from the decomposition of the copper (I) acetamidinate precursor, they react differently with the CuO surface depending on the environmental atmosphere.

Thermal properties of Al/CuO nanothermites

The sonication of the Al and CuO NPs in hexane leads to homogeneous mixtures that are observed in the TEM photographs (Figure 4). The Al/CuO_com mixtures (Figure 4a) contain aggregates with spread homogeneous mixture of spherical Al and CuO NPs. Furthermore, Al/CuO_air mixtures (Figure 4b) show a heterogeneous distribution unlike Al/CuO_O2 ones (Figure 4c) where small CuO NPs cover Al NPs, pointing to the intimate interfacial contact between Al and CuO NPs. Therefore the amount of small CuO NPs around Al NPs is greater than for the Al/CuO_com mixture. The results of DSC experiments are presented in Figure 5. Al/CuO_com (Figure 5a) exhibits two major exothermic peaks and one endotherm after 500 °C, whereas Al/CuO_air and Al/CuO_O2 (Figure 5b andc) show weaker exothermics peaks at similar temperatures, typically associated with the thermite reaction between CuO and Al nanoparticles. 17 The temperatures associated with these exothermic and endothermic reactions, as well as the heat released between 450 and 900 °C are summarized in Table 2. Following the first exothermic peak (i) at 578 °C, which corresponds to the reaction of solid Al with CuO, the small endotherm (k) is attributed to the melting of Al at 658 °C. A second exothermic peak (j), caused by the reaction of liquid Al with solid Cu2O and some CuO, occurs at 783 °C. During this reaction, the amorphous alumina transforms into  Al2O3, [START_REF] Abdallah | [END_REF] which explains the sharp decrease in the DSC. However, we still can identify and interpret some thermal responses on the DSC curves of Al/CuO_air and Al/CuO_O2. The most significant one is the intense endotherm seen in the Al/CuO_air system at 189 °C that is not detectable in the Al/CuO_O2 system. In addition, two exotherms are observed slightly above 200 °C and the heat released between 450 and 900 °C for Al/CuO_O2 is higher than that for Al/CuO_air (see Table 2). To summarize, we clearly observe three different thermal responses depending on the CuO NPs used, highlighting the role of synthesis protocol and ligands nature in the hindering of the thermite reaction. Al/CuO_O2. Indexed peaks are described in Table 2.

Table 2. Detail of reactions associated to the thermal responses of nanothermites plotted in Figure 5. Results are consistent with DSC traces presented earlier (Figure 4b andc). The main difference between both sets of NPs concerns the CuO_air endotherm peak, noted (l), which is followed by two exotherms (m) and (n) also visible for CuO_O2, while being less intense.

Whatever the synthesized CuO conditions, the significant mass loss asserts that decomposition reactions are taking place at the CuO particle surface. Mass spectrometry analyses of both synthesized CuO NPs upon heating (@ 10°C.min -1 ) under argon atmosphere reveal that mass losses are accompanied with the release of different species. Figure 7 shows MS spectra for both CuO_air and CuO_O2 at ambient temperature (AT), 100, 175, 235, 280, 341, 382 and 400 °C. A typical commercial CuO background spectrum for the mass spectrometer consists of a large peak at mass to charge ratio (m/z) 18 (H2O), smaller peaks at m/z 28 (N2) and m/z 32 (O2). 30 Detailed mass spectra, included all detected m/z species, during the CuO_air and CuO_O2 ramping are given in supplementary file, Figure S1 and Table S1. For both CuO_air and CuO_O2, the major decomposition products include m/z 18 (H2O) and m/z 44 (CO2). These observations can be better understood by quantifying the kinetics of desorption plotted in supplementary file (Figure S2). Importantly, H2O molecules start to be released at room temperature for CuO_air and in larger amounts compared to CuO_O2, whereas the kinetics of H2O molecules release is similar for both CuO (see Figure S2). This is consistent and matches well the endotherm peak (I) seen in DSC analyses (Figure 6b). This is also consistent with the chemical nature of CuO NPs grown under air, where OA ligands are potentially decomposed into carbamate/ammonium ion pairs (RNCO2 -/RNH3 + ) due to the reaction with ambient H2O and CO2, as suggested also from FTIR results (Figure 3b). In such a scenario, these ions should increase the polarity of the surface, facilitating the adsorption of water molecules, despite the presence of hydrophobic alkyl chains. On the contrary, the The early reduction of CuO in the case of ultra-small particles has been evidenced by Pike et al. who observed the reduction of CuO nanoparticles being 4-16 nm in width at 250°C using a ramping heating profile, in CO2 atmosphere. 31 Altogether these results also question on the effect of CuO nanoparticles size on their stability in temperature, i.e. ability to withstand a higher temperature before their reduction into Cu2O and then to metallic copper nanoparticles. To address this question, larger particles, CuO_com, are functionalized with OA (named CuO_com@OA) and characterized upon ramping at 10°C.min -1 in Ar atmosphere (See supporting information S5). Figure S5 displays FTIR and XRD characterization of CuO_com@OA showing alkyl chain of OA at the stretching mode in the area of 3000 cm -1 , proving that bonding is realized between OA and commercial CuO. XRD analysis of annealed CuO_com@OA (Figure 5b) demonstrates that at 500 °C, CuO phase is partially reduced into Cu2O and Cu phases. The DSC diagram of CuO_com@OA mixed with Al (Figure S5c) exhibits a heat release of 0.79 kJ.g -1 between 450 and 900 °C.

Altogether, this confirms the role of ligands in the reduction process of CuO nanoparticles but also that, in the case of bigger particles, CuO phase is still present after 500 °C, allowing the reaction with Al. Therefore, CuO nanoparticles size has to be taken into account with precision to avoid early reduction.

Optimization process: atmosphere controlled thermal treatment of CuO nanoparticles

To overcome the early reduction of ultra-small CuO nanoparticles, the synthesized CuO nanoparticles are annealed at 500 °C under oxidizing environment to burn the organic coating without reducing the CuO and enhance their size. In Figure 9 Annealed CuO nanoparticles grow during the heating treatment that could be the effect of nanoparticles sintering. From XRD spectra (Figure 9), crystallite size is evaluated at 33 ± 6 nm and 41 ± 8 nm for annealed CuO_air and CuO_O2, respectively.

After mixing with Al commercial NPs, SEM pictures in Figure 10 nanothermites. Indexed peaks are described in Table 3.

Finally, the thermal performances of both nanothermites obtained by mixing Al NPs with annealed CuO NPs are characterized by DSC and diagrams are given in Figure 10. Two exotherms (i and j) are observed for the Al/annealed CuO_O2 mixture (Figure 10e) whereas only one (i) is present for the Al/annealed CuO_air mixture (Figure 10f).

For both nanothermites, the reaction onset occurs earlier than for the Al/CuO_com nanothermite (Figure 5). A strong endotherm appears at 650 °C (k) for Al/annealed CuO_air mixture (a) attributed to Al-Cu eutectic. The first and second exothermal peaks (i and j) are associated to the redox reaction. The total heat released for both nanothermites presented in Table 3 remains competitive with Al/CuO_com (Table 2).

Importantly a major part of this total energy is released prior Al melting: 63 % and 48 % for Al/CuO_air and Al/CuO_O2 compared to 33% for Al/CuO_com. This is an interesting result for application requiring low temperature reactive materials.

Conclusion

Figure 1 .

 1 Figure 1. Schematic of the synthesis of CuO nanoparticles following an organometallic

  CuO_O2), CuO_com and Al NPs are shown in Figure2, with the addition of size distribution following a Gaussian fit. The CuO_air objects are characterized as polymorphous shapes and mean diameter of 6.8 ± 2.2 nm, whereas CuO_O2 particles exhibit smaller size with a mean diameter of 3.8 ± 0.7 nm and more isotropic morphologies. On the contrary, CuO_com displays large particles with an extended size distribution from 15 to 120 nm and mean diameter of 31 ± 15 nm. In the same way, Al NPs feature spherical shapes with a mean diameter of 96 ± 43 nm. The synthesis of CuO NPs allows an exquisite control on getting smaller size and monodispersity down to the nanoscale, due to the slow and homogeneous chemical interplay between ligands (OA and/or carbamate) and growing CuO particles.27

Figure 2 .

 2 Figure 2. TEM images of (a) CuO_air, (b) CuO_O2, (c) CuO_com (d) and Al NPs. The

Figure 3 shows

 3 Figure3shows XRD diagrams corresponding to CuO_air, CuO_O2 and compared with

Figure 3 .

 3 Figure 3. X-ray diffraction patterns (a) and FTIR spectra (b) obtained on CuO_air,

Figure 4 .

 4 Figure 4. TEM images of the different Al/CuO mixtures prepared following the

Figure 5 .

 5 Figure 5. DSC curves obtained for (a) Al/CuO_com, (b) Al/CuO_air and (c)

  of ligands with CuO hindering the Al/CuO redox reactionTo get into further details on the reaction of CuO nanoparticles with their surrounding chemical cues, neat CuO_air and CuO_O2 NPs are chemically analyzed upon ramping from ambient to 500 °C. TGA curves (Figure6a) show that commercial powders, CuO_com, feature no mass loss whereas CuO_air and CuO_O2 lose a total of 58 % and 29 % of their initial mass, respectively. For both synthesized CuO NPs, the mass loss occurs in three consecutive steps. The first one occurs at ~200 °C and the mass loss is much pronounced for CuO_air (~37 %) compared to CuO_O2 (~ 4 %). The two other loss steps (at 285 °C and 412 °C, respectively) show similar behaviour for both sets of NPs. In parallel to TGA, DSC measurements are also recorded for the three batches of CuO NPs (see Figure6b): while no peak is observed for CuO_com in the range of ambient -500 °C, both synthesized CuO NPs exhibit different reactions and associated thermal signatures at temperatures lower than the expected Al-O reactions.

Figure 6 .

 6 Figure 6. (a) TGA and (b) DSC curves of CuO_air, CuO_O2 and CuO_com NPs under

Figure 7 .

 7 Figure 7. MS responses of (a) CuO_air and (b) CuO_O2 plotted at various

CuO_O2

  NPs exhibit less affinity with water molecules, indicating that the hydrophobic contribution of alkyl chains dominates. The wider range of temperatures in which water molecules are desorbed in CuO_air (see Figure S2) is another illustration of its surface/ligands chemical complexity, characterized by different adsorption configurations associated with the presence of ligands of different nature and organization (OA, carbamate, ammonium …). At 175 °C, CO2 start to be released abruptly in CuO_air system which is earlier than what is observed for CuO_O2 (280 °C). Increasing temperature up to 235 and 280 °C, NO species (m/z 30) are detected in CuO_air, which indicates that CuO_air are surrounded by more ligands and chemical species than CuO_ O2. To summarize MS experiments, upon ramping heating, the ligands decomposed into fragments (H2O, R-CO2, R-NH3 …) with a release of CO2 and H2O occuring at ~270-280 °C for both CuO NP systems. The exothermic peaks (m) and (n) in Figure 6b can therefore be attributed to the oxidation of some released species (R-CO2, R-NH3 …) with oxygen from the CuO surface. XRD diagrams (Figure 8) of CuO_air (a) and CuO_O2 (b) taken at 200 and 500 °C clearly show the CuO_air decomposition into cuprite (Cu2O, major phase) at 200 °C and, to a lesser extent, copper. At 200 °C, the decomposition of CuO_O2 is less pronounced, as cuprite phase is present but still represents a minor phase. At 500 °C, for both samples, copper structure is the main phase with the presence of a small quantity of Cu2O. Early reduction of CuO is confirmed by FTIR analysis given in supporting information, Figure S3: cuprite phonon modes at 616 cm -1 is evidenced at 200 °C for both samples, whereas they disappear at 500 °C when copper is formed. These effects are even stronger for CuO_air with the formation of copper at 200 °C, due to the larger quantity of ligands and water at the CuO surface. XRD analysis of CuO_com annealed at 200 and 500 °C under argon is presented in Figure S4. The initial CuO structure doesn't change until 500 °C, pointing to the crucial role of the ligands in the CuO reduction process below 500 °C observed for CuO_air and CuO_O2.

Figure 8 .

 8 Figure 8. X-ray diffraction patterns of (a) CuO_air and (b) CuO_O2 taken at 200 and

  , XRD diagrams show CuO_air and CuO_O2 annealed under oxygen at 200 and 500 °C. Surprisingly, at 200 °C, CuO_air reduces mainly into copper and Cu2O whereas CuO_O2 partly remains as CuO structure, but with a strong reduction component into Cu2O. The CuO reduction process appears inevitable due to the preference of ligands to extract oxygen atoms from the CuO surface instead of reacting with oxygen molecules from the atmosphere. Since ligands are coordinated to the surface, oxygens from CuO surface are more available for the oxidation compared to those in the atmosphere. However, at 500 °C both CuO_air and CuO_O2 get re-oxidized into a tenorite structure with a remaining minor cuprite phase for the CuO_air system.

Figure 9 .

 9 Figure 9. X-ray diffraction patterns of (a) CuO _air and (b) CuO_O2 taken at 200 and

  Figure S6). CuO are organized as extended domains of sintered particles, as shown

Figure 10 .

 10 Figure 10. SEM images of (a) and (b) Al/annealed CuO_air mixtures, (c) zoom on

Figure 11 .

 11 Figure 11. (a) Onset temperatures and (b) energy released prior 500 °C (b) determined

  . Three batches of Al rich nanothermites (Al to CuO mass ratio of 1.2) are then prepared, which are named

	Material	Source	Mean diameter (nm)
	Al (purity 69 %)	US research nanomaterials	96 ± 43 nm (TEM)
	CuO_com	US research nanomaterials	31 ± 15 nm (TEM)
	CuO_air	US research nanomaterials	6.8 ± 2.2 nm (TEM)
	CuO_O2	US research nanomaterials	3.8 ± 0.7 nm (TEM)

Table 3 .

 3 Thermal reactions associated to the Al/CuO mixtures in the Figure10

kJ.g -1 ) (450-900 °C)

	Indexed peaks	Peak temperature	Peak temperature	Reactions
		Al/annealed	Al/annealed CuO_O2	associated
		CuO_air (°C)	(°C)
	(i)	627	553	Al + O
	(j)	657	653	Al2Cu melting ; Al
				melting
	(k)	/	725	Al + O
	Heat released (0.81	0.93

Acknowledgements

We would like to thank Dr. Pierre Alphonse, Dr. Laure Vendier and Jean-François Meunier for their assistance and expertise in the TGA/DSC measurements, Xray diffraction analyses and mass spectroscopy measurements respectively. We also thank Severine Vivies for her collaboration in producing the nanothermites and their characterization. We thank the European Commission and Region Occitanie for their FEDER support (THERMIE grant) having funded the DSC equipment. We also thank the European Research Council as well as the university Fédérale de Toulouse for its financial support through the PyroSafe project (grant number 832889) and "chaire d'attractivité" MUSE, respectively.

Copper oxide nanoparticles of 6 nm in average size were synthetized following an original organometallic route before being mixed with Al nanoparticles to form a nanothermite composite. As synthesized, we found that ligands coordinate to CuO nanoparticles leading to their early reduction that is significantly different from that of usual bulk CuO used in nanothermite composite. Notably the synthesized nanoparticle reduce to an intermediate Cu2O phase at ~250°C before further reduction to metallic copper at ~450 °C. Thus, energetic performances of the nanothermites prepared from synthesized CuO show unexpected exothermic reactions below 500 °C. Using a variety of characterization techniques, including microscopy, spectroscopy, mass spectrometry and calorimetry, the structural and chemical evolution of CuO nanoparticles stabilized with ligands (alkylamine) were analyzed upon heating enabling us to unravel that, at low temperature (~200 °C), ligands fragment into organic species accompanied with important release of H2O and CO2 molecules. This process promotes the CuO reduction into Cu2O and further Cu. Finally, we proposed an effective method to overcome the CuO degradation. After the synthesis, the CuO nanoparticles coated with the ligands are annealed in oxidizing atmosphere enabling the oxidation of ligands with the CuO reduction at low temperature, to finally re-oxidize the reduced CuO. The CuO nanoparticles increase in size while keeping roughly the same shape, probably being a matrix of sintered nanoparticles. Finally the thermites produced by mixing thus CuO particles with commercial Al ones exhibit earlier onset temperature than those produced from commercial CuO nanoparticles with similar heat of reaction.

Competting financial interests

The authors declare no competing financial interests.

Corresponding author

Correspondance to: Carole Rossi, rossi@laas.fr


Pedro González Losada
email: pedro.gonzalezlosada@esiee.fr

Lionel Rousseau
email: lionel.rousseau@esiee.fr

Marjorie Grzeskowiak
email: marjorie.grzeskowiak@univ-mlv.fr

Julie Degardin
email: julie.degardin@inserm.fr

Sarah Berhanu
email: sarah.berhanu@inserm.fr

Elisabeth Dubus
email: elisabeth.dubus@inserm.fr

Serge Picaud
email: serge.picaud@inserm.fr

Gaëlle Lissorgues
email: gaelle.lissorgues@esiee.fr

Protuberant Electrode Structures for New Retinal Implants

Keywords: microfabrication, retinal prostheses, implantable device, 3D electrode, retinal stimulation

Focal stimulation of neural tissue is a challenge in retinal prosthetic devices in which the aim is to improve the spatial resolution of the stimulation and therefore increase the electrode density of the prosthetic devices. Our work intends to create a new implant able to enhance the focalisation of the stimulation signal through protuberant electrodes. These electrodes are micro fabricated on a soft polyimide substrate using classical metal electrodeposition techniques. Before proceeding with fabrication a FEM model of the electrode's current density was done to select the best-performing structures and geometries in terms of local stimulation. Based on these models, several prototypes were fabricated and implanted in vivo into a rat's eye to verify the adaptation to the retina tissue.

Introduction

Work already carried out in retinal implants shows that several geometric modifications can be made. On the one hand, the use of a ground plane around the electrodes instead of a far return electrode keeps the stimulation signal more localised on the electrode area [START_REF] Joucla | Improved focalization of electrical microstimulation using microelectrode arrays: A modeling study[END_REF]. On the other hand, the insertion of the electrode in a cavity of few micrometres surrounded by a ground plane creates 3D shapes which represents an improvement [START_REF] Djilas | Three-dimensional electrode arrays for retinal prostheses: Modeling, geometry optimization and experimental validation[END_REF]. An important aspect is the interaction between the electrode and the neural tissue: a closer interaction will entail a better focalisation. To obtain such effect, 3D mushroom or pillar electrodes have been used in Micro Electrode Arrays to improve the interaction between cultivated cells and the electronic device [START_REF] Huys | Novel concepts for improved communication between nerve cells and silicon electronic devices[END_REF].

Our work tries to integrate both elements in a soft implant to study two main aspects: the performance of the electrical stimulation and the inner integration of the implant with the tissue in vivo. We carried out our research in three phases. First we developed a FEM model of the electrode with different geometries to evaluate its current delivering performance. Then, taking into account these results and using the planar implant fabrication technology previously developed in our laboratory, we added new fabrication steps to build the protuberant electrodes. Finally, in order to verify the adaptation of the retinal tissue to these protuberant shapes, the devices were implanted in sub-retinal position into P23H blind rats with a very thin retina, similar to the one that could be found in the targeted pathologies.

Materials and Methods

FEM Simulation

A finite element model (FEM) has been developed using COMSOL. The geometry consists of a single protuberant metallic electrode embedded in an insulating substrate and surrounded by a liquid environment with a conductivity which represents the physiological environment. The interface between the insulating substrate and the liquid environment is modelled using insulating boundaries described by Equation [START_REF] Joucla | Improved focalization of electrical microstimulation using microelectrode arrays: A modeling study[END_REF]. For the electrode-electrolyte interface, robin boundary condition described by Equation ( 2) is used as previously defined by Joucla et al. [START_REF] Joucla | Improved focalization of electrical microstimulation using microelectrode arrays: A modeling study[END_REF]. The conductivity values for the different materials are considered homogeneous and isotropic, as summarised in Table 1.

Two types of structures are modelled for comparison: a flat electrode in a cavity surrounded by a ground plane and a single protuberant electrode in a cavity surrounded by a ground plane, with the same cavity depth.

σ∇V n = 0, ( 1
)
σ∇V n = g(𝑉 𝑚𝑒𝑡𝑎𝑙 -𝑉), (2)

Microfabrication

As mentioned in the introduction, the fabrication technology is based on the planar implant technology to which new steps haven added in order to put in the 3D shaped wells the protuberant electrodes. The size of the hexagonal cavities is 100 μm and the diameter of the protuberances is 12 um. We designed different cavities in the same implant: some of them empty, and the others with 1, 2, 3 or 4 protuberant structures. The fabrication process for the soft implants on a silicon wafer is the following: After the deposition of a sacrificial layer composed of titanium (100 nm) and aluminium (500 nm), 10 μm of polyimide (PI 2611) are spin coated. Layers of titanium (100 nm) and gold (500 nm) are deposited by sputtering and patterned by photolithography to define electrodes, tracks and pads. Finally, SU8 2002 (MicroChem) is spin coated to form a 2 μm passivation layer. Using photolithography, openings are created on the pads and electrodes. Then, 500 nm of aluminium is sputtered to define the shape of the implant by lithography. To etch the polymer, we use reactive ion plasma composed of a mix of Ar and 02 at 120 Watt. To peel off the implants, the aluminium layer is dissolved by electro erosion

In our case, we do not peel off the implants and continue the fabrication to produce 3D structures. On the whole wafer, we deposit a thin layer of titanium (50 nm) and 150 nm of copper corresponding to the seed layer. A photoresist is deposited to form the mould for protuberant electrodes. The wafer is placed in a copper solution and a current of 100 mA is applied to create copper protuberances. The photoresist is cleaned in solvent and a thick photoresist (15 μm) is used to perform the counter ground plane. Once the photolithography step is completed, the wafer is placed for a second time in the copper electrodeposition solution in the same growth conditions. Wafers are cleaned and the initial seed layer is etched. A photoresist protection is applied before peeling off the implants. After cleaning, a thin layer of 2 μm of parylene C is used to encapsulate the whole copper structure (as Copper is not biocompatible). For the next generation of implants, the 3D structures will be built through gold electrodeposition (for biocompatibility).

The implantation of the devices is done in P23H/HO rats at the age of 9 months and only in one eye. The placement of the implant on the sub-retinal position is checked periodically during the implantation time by means of eye fundus observation (Micron III) and OCT images. After 19 weeks the rats are sacrificed, the retina is explanted and cell labelling is performed.

Results

FEM Simulation

The cross sectional view of the current density is shown in Figure 1 for the planar geometry (a) and protuberant one (b). In the planar geometry a border effect is identified both in the border of the electrode and the ground plane: most of the current penetrates into the liquid in this part of the electrode. In the protuberant geometry, we appreciate a redistribution of the current density as the border effect still exists but part of the current flows to the media through the protuberance.

Microfabrication

In order to characterise the electrodeposition time for the protuberant electrodes and the ground plane, several tests were done on wafers where only the electrodes or the ground plane were patterned. The time and the thickness of the metallic structures were measured using a Wykko profilometer to obtain the relation between the electrodeposition time and the electrodeposed metal thickness. These values were used for the electrodeposition of the structures on the soft implants and the results are summarised in Table 2. The surface of the electrodes was examined using a scanning electron microscope (ZEISS NEON40) to see the general aspect of the implant (Figure 2a) and the aspect of the individual electrodes (Figure 2b).

Implantation

The digital endoscope (Figure 3a) and OCT images taken during the whole implantation period can show that the implant is properly placed in the sub-retinal position. A confocal analysis of the explanted tissue allows us to check the adaptation of the retina to the new implant geometry. The images obtained (Figure 3b) do not show a glial reaction of the tissue around the implant and the cells can descent into the cavities whether there are protuberances or not.

Conclusions

Focal stimulation is a challenge for neural prostheses and especially for retinal ones. Improvement of the stimulation process can be achieved by modifying several parameters as the type of stimulation signal, the materials used for the prostheses fabrication or their geometry. We explore here this last option to try to improve the interface between the electrode and the tissue. For this purpose we first developed a FEM model of the electrode that we could micro fabricate and which gave us an idea of the best structures to explore. After micro fabrication, we implanted them in a PH23 rat's eye to examine the adaptation of the retinal tissue to this new 3D geometry. Images obtained using a confocal microscope and specific labelling showed that the retina is able to adapt to the shape of this new implant and that bipolar cells are present in the cavities.

Figure 1 .

 1 Figure 1. Comparison of the current density delivered by a flat electrode and a protuberant electrode in a two level configuration: ground plane and flat electrode plane are not at the same level-same colour scale.

Figure 2 .

 2 Figure 2. SEM image of the implant (a) General view of the implant head on soft substrate; (b) Detail of the configuration consisting of 4 protuberant electrodes embedded in the well formed by the ground plane.

Figure 3 .

 3 Figure 3. (a) Verification of the correct placing of the implant in sub-retinal position two weeks after implantation: Eye fundus of P23H rats with the implant obtained with a Micron III; (b) Confocal image of the retinal tissue and implant, where blue cells represent the bipolar cells after 19 weeks.

Table 1 .

 1 Parameters used for the simulation.

	Material	Conductivity (S/m)	Relative Permittivity
	Physiological liquid	1.47 × 10 -2	81
	Insulating material	1.31 × 10 -18	3.4
	Conductive material	45.6 × 10 6	-

Table 2 .

 2 Relation between the electrodeposition time and the thickness achieved.

		Electrode	Ground Plane
	Wafer	Time [min]	Thickness [um]	Time [min]	Thickness [um]
	W1	17	3.25	57	23.00
	W2	17	3.20	50	22.00

Acknowledgments: Work performed within the ANR grant project NEUROMEDDLE.

Conflicts of Interest:

The authors declare no conflict of interest.