Cristina Nita

Julien Fullenwarth

Laure Monconduit

Loïc Vidal

Camélia Matei Ghimbeu
email: camelia.ghimbeu@uha.fr

Influence of carbon characteristics on Sb/Carbon nanocomposites formation and performances in Na-ion batteries

Keywords: Sb, hard carbon, mesoporous carbon, alloy, Na-ion batteries

The impact of carbon characteristics (structure, porosity and surface chemistry) on the Sb/Carbon composites formation and electrochemical performances in Na-ion batteries is investigated herein. Mesoporous and non-porous hard carbon heated at different temperatures were used to support Sb nanoparticles prepared via a ball-milling/thermal reduction synthesis approach. The mesoporous carbon presenting higher surface area and oxygen functional groups induce the formation of small Sb nanoparticles (< 5 nm) well dispersed in the carbon network. For the hard carbon, the increase in the annealing temperature from 1100 to 1500°C leads to more organized structure but lower porosity and surface functionalities having as consequence the increase of Sb nanoparticle size and the particle agglomeration. The best electrochemical performances vs. Na were achieved for hard carbon/Sb composite treated at 1100°C which combines several advantages like dispersed Sb nanoparticle, low surface area and good conductivity (360 mAh g -1 after 200 cycles). Regarding the long-term capacity retention, for all materials a drastic decay was observed starting to 300 cycles and associated to the Sb particle size agglomeration and growth during repeated charge/discharge cycles and amorphization of Sb particles as demonstrated by post-mortem analyses.

Introduction

Lithium-ion batteries (LIBs) are widely used for portable electronic devices such as phones, laptops, cameras and more recently emerged for large-scale energy storage applications for electrical vehicles and intermittent solar and wind energy sources. However, the increase demand for Li-ion batteries in an economical and geographical context where the Li sources become more and more scarce and localized only in few areas of the world requires tremendous efforts to find alternative cost-effective solutions for current LIBs.

In the recent years, sodium-ion batteries (NIBs) have been brought into the spotlight due to the rich Na resources in the Earth's crust which may offer a low-cost energy system for near future. Despite the rather similar working principles of LIBs and NIBs, Na + is not able to insert into graphite anode due to its larger size, steric and thermodynamic limitations [START_REF] Dahbi | Negative electrodes for Naion batteries[END_REF][START_REF] Ellis | Sodium and sodium-ion energy storage batteries[END_REF][START_REF] Kim | Electrode Materials for Rechargeable Sodium-Ion Batteries: Potential Alternatives to Current Lithium-Ion Batteries[END_REF], leading to poor performances. Therefore, many studies were focused on hard carbon materials [START_REF] Saurel | From Charge Storage Mechanism to Performance: A Roadmap toward High Specific Energy Sodium-Ion Batteries through Carbon Anode Optimization[END_REF][START_REF] Wu | Surface-Driven Energy Storage Behavior of Dual-Heteroatoms Functionalized Carbon Material[END_REF][START_REF] Hou | Carbon Quantum Dots and Their Derivative 3D Porous Carbon Frameworks for Sodium-Ion Batteries with Ultralong Cycle Life[END_REF] having disordered structure and larger interlayer distance compared to graphite. Various types of hard carbons issue from the pyrolysis of monomers, polymers, biomass and natural resources have been proposed and capacities up to 300 mAh g -1 could be achieved in some cases via a storage mechanism involving adsorption and intercalation of Na + in the carbon porosity and nanographitic domains [START_REF] Ghimbeu | Insights on the Na+ ion storage mechanism in hard carbon: Discrimination between the porosity, surface functional groups and defects[END_REF][START_REF] Kang | Update on anode materials for Na-ion batteries[END_REF][START_REF] Ponrouch | High capacity hard carbon anodes for sodium ion batteries in additive free electrolyte[END_REF]. In order to achieve higher gravimetric and volumetric capacities, anode materials based on alloying materials have been studied as well as anodes for Na-ion batteries. Their charge storage mechanism relies in the formation of alloys with sodium during the charging of the anode leading to much higher and attractive capacities than carbon. Among the reported alloying materials, antimony, tin, and phosphorus are considered promising for Na-ion batteries thanks to their appealing theoretical capacities [10;11]. Antimony (Sb) exhibits a high theoretical capacity (660 mAh g -1 ) upon full sodiation to Na3Sb alloy, good kinetic performance and high volumetric energy density compared to carbon materials. However, Sb anodes suffer from a major inconvenient related to its large volume changes during sodiation/desodiation which induces particle growth and agglomeration, pulverization, loss of electrical contacts and thus performance degradation.

Such large volume changes of alloy materials may limit their applications for NIBs. Lately, several solutions have been proposed to overcome this problem including the nanostructuration of the active material, the use of a second metal or the use of carbon materials in order to buffer the Sb volume expansion.

The nanostructuration by synthesizing nanoparticles, nanocrystals or nanorods can facilitate the electrolyte ion diffusion ensuring faster kinetics and high-rate performance [START_REF] He | Monodisperse Antimony Nanocrystals for High-Rate Li-ion and Na-ion Battery Anodes: Nano versus Bulk[END_REF][START_REF] Liang | Large-scale highly ordered Sb nanorod array anodes with high capacity and rate capability for sodium-ion batteries[END_REF][START_REF] Zhang | Spherical nano-Sb@C composite as a high-rate and ultra-stable anode material for sodium-ion batteries[END_REF]. The use of metals such Al, Fe, Cu, Sn [START_REF] Xu | Atomic scale, amorphous FeOx/carbon nanofiber anodes for Li-ion and Na-ion batteries[END_REF][START_REF] Xie | Two-dimensional layered compound based anode materials for lithium-ion batteries and sodium-ion batteries[END_REF][START_REF] Wu | Graphene-based nano-materials for lithium-sulfur battery and sodium-ion battery[END_REF][START_REF] Wang | Recent developments in electrode materials for sodium-ion batteries[END_REF][START_REF] Li | CuS Microspheres as High-Performance Anode Material for Na-ion Batteries[END_REF][START_REF] Darwiche | Facile synthesis and long cycle life of SnSb as negative electrode material for Na-ion batteries[END_REF][START_REF] Choi | Carbon embedded SnSb composite tailored by carbothermal reduction process as high performance anode for sodium-ion batteries[END_REF] to minimize the volume changes and to prevent aggregation during cycling has been also reported, however, except Sn, most of these metals are electrochemically inactive and would scarify the capacity in some extent. Taking into account this aspect, several works have been focused on carbon materials which are able not only to buffer the active material volume expansion and to avoid agglomeration but also to provide electrical conduction and electrolyte diffusion pathways. A literature table gathering the main forms of Sb and Sb/C composites, their synthesis methods along with their electrochemical performances is provided in Table 1. Darwiche et al. used commercially available bulk Sb electrode and mixed with vapor ground carbon fibers and carbon black conductive additives which resulted in high performance at a high current rate (550 mAh g -1 at 4C, 160 cycles) [START_REF] Darwiche | Better Cycling Performances of Bulk Sb in Na-Ion Batteries Compared to Li-Ion Systems: An Unexpected Electrochemical Mechanism[END_REF]. Electrospinning was used to anchor Sb nanoparticles on self-sustained carbon nanofibers films derived from polyacrilonitrile (PAN) polymer [START_REF] Li | Antimony-Carbon-Graphene Fibrous Composite as Freestanding Anode Materials for Sodium-ion Batteries[END_REF][START_REF] Zhu | One-dimensional coaxial Sb and carbon fibers with enhanced electrochemical performance for sodium-ion batteries[END_REF][START_REF] Zhu | Electrospun Sb/C Fibers for a Stable and Fast Sodium-Ion Battery Anode[END_REF] and the capacities were comprised between 274 and 650 mAh g -1 . Liao et al. [START_REF] Liao | Ultrafine Sb nanoparticles embedded in an amorphous carbon matrix for high-performance sodium ion anode materials[END_REF] reported a hydrothermal approach to obtain Sb particles in an amorphous carbon matrix which exhibited high capacity of 622 mAh g -1 , at 100 mA g -1 current rate after 100 cycles. Other approaches such as sol-gel, impregnation of a carbon with a Sb salt followed by chemical reduction, polymerization in the presence of a Sb salt have been explored and yielded distinct materials with capacities ranging between 241 and 554 mAh g -1 (see Table 1) [14;27-30]. Some simple approach consisting in mixing of a carbon with metallic antimony by ball milling has been demonstrated as well. Zhou et al. [START_REF] Zhou | Wet milled synthesis of an Sb/MWCNT nanocomposite for improved sodium storage[END_REF] used multiwall carbon nanotubes (MWCNTs) and Sb powder and obtained uniformly mixture of Sb NPs on MWCNTs by a wet milling approach.

The materials exhibited a high capacity of 502 mAh g -1 at 100 mA g -1 current rate and 120 cycles. Qian et al. [START_REF] Qian | High capacity Na-storage and superior cyclability of nanocomposite Sb/C anode for Na-ion batteries[END_REF] prepared Sb/C nanocomposites by mechanical milling of commercial Sb powder with a carbon black (Super P) and demonstrated a high capacity of 610 mAh g -1 at 2 Ag -1 current rate over 100 cycles. Ramireddy et al. [START_REF] Ramireddy | Size and Composition Effects in Sb-Carbon Nanocomposites for Sodium-Ion Batteries[END_REF] prepared Sb/C composites in a similar manner by using graphite as carbon support. The particle size could be tuned from 1 to 50 nm by adjusting the carbon and Sb weight ratio in the mixture allowing to obtain high a stable capacity at high current rates. Ji et al. [START_REF] Wu | Dual Functions of Potassium Antimony(III)-Tartrate in Tuning Antimony/Carbon Composites for Long-Life Na-Ion Batteries[END_REF] proposed a self-thermal-reduction approach of potassium antimony(III)-tartrate to form Sb/C rod composite which exhibited low capacity and fading after ~ 70 cycles, i.e., 150 mAh g -1 at 100 mA g -1 for 200 cycles, due to the material volume expansion and fragmentation. Such limitation could be overcome by using potassium chloride during the synthesis which leaded to dual carbon/Sb nanoparticle composite (Sb-O-C/C) with high capacity (500 mAh g -1 ) and good capacity retention for more than 200 cycles (Table 1).

Considering these above mentioned works, it is clear that the carbon materials play an important role in the stabilization of Sb nanoparticles during synthesis and cycling, in improving the electronic conductivity and in fine in improving the electrochemical performance of SIBs. Despite that, most reported works are devoted to new synthesis routes to obtain Sb/C nanocomposites by employing different carbon and Sb sources and no systematic studies concerning the influence of carbon support characteristics on the formation of Sb nanoparticles and their performances in SIBs have been reported to our knowledge. Moreover, long term cycling is rarely discussed although of prime importance for developing robust energy storage systems. Therefore, we report in this work a simple ball milling procedure involving carbon materials and antimony salt followed by thermal reduction to obtain Sb/C nanocomposites.

The influence of carbon texture, structure and surface chemistry on the Sb particle size/dispersion was evaluated by several techniques along with their electrochemical performances. Carbons possessing higher surface area and larger amount of functional groups lead to smaller and better dispersed Sb NPs and improved electrochemical performances. The carbon/Sb anodes tested for Na-ion batteries exhibits high reversible capacity after the first charge-discharge cycle up to 639 mAh g -1 and a stable capacity retention up to 200 cycles with a capacity of 360 mAh g -1 . Moreover, irreversible capacity was successfully reduced by using carbon support with low surface area and low surface chemistry. The capacity fading at long term cycling is discussed as well in terms of materials characteristics.

Materials and methods

Materials Synthesis

In a first step, the carbon support, i.e., a mesoporous carbon and a non-porous carbon called herein hard carbon was prepared by sol-gel polymerization. The mesoporous carbon was obtained as described previously [START_REF] Jahel | Exceptionally highly performing Na-ion battery anode using crystalline SnO2 nanoparticles confined in mesoporous carbon[END_REF], briefly phloroglucinol (the carbon source, 1.65g) and pluronic F-127 (the structure directing and pore generating agent) were dissolved in an acidic ethanol solution, followed by the glyoxal cross-linker addition (1.62 mL). The obtained polymer gel after 3 days aging is dried followed by thermal pyrolysis at 900°C under Ar. The preparation of the hard carbon support was the same as for the mesoporous carbon, the only difference is the absence of pluronic F-127 in the mixture. Three annealing temperatures were used, i.e., 1100, 1300 and 1500 o C and the as-prepared carbons denoted: MC@900-for the mesoporous carbon and HC@T-for the hard carbons, where T=1100, 1300 and 1500 o C.

In a second step, the C/Sb materials were prepared by using a simple ball milling of carbon supports with antimony acetate (Sb (CH3COO)3) for 1h, followed by a thermal reduction in presence of Ar:H2 gas mixture (1:1, 10L/h flow rate). The materials were heated at 250 o C for 1h to allow the melting of the precursor, followed by the increase of the temperature up to 500 o C for 1h (heating rate 5k/min) to decompose and reduce the precursor to metallic Sb. A schematic illustration of the synthesis pathway can be found in

Materials characterization

The crystalline structure of the materials was characterized by X-ray powder diffraction The amount of antimony in the composite was determined by thermogravimetric analysis (TGA) conducted on a TGA 851 (Mettler-Toledo) instrument by heating the C@Sb composite under air (100 ml/min), from 30 o C up to 900 o C with a 10 o C/min rate. The textural properties of the carbon/Sb materials were evaluated using N2 and CO2 sorption isotherms by using a Micromeritics ASAP 2420 set-up. The materials were outgassed at 200°C under vacuum for 12h prior to the analyses. The BET specific surface area (SSA) was calculated in the relative pressure range, P/P0, of 0.05-0.3.

Electrochemical characterization

The electrode was prepared by mixing the carbon/Sb materials (70%) with carbon black (10%) and vapor grown carbon fibers -VGCF-S (10%) as conductive additives, and carboxymethylcellulose -CMC (10%) as binder. A quantity of 0.4 ml distillated water was added over the materials and homogeneously mixed by ball-milling for 1h. A film with a thickness of 100µm was laid down on a copper foil, dried at room temperature for 4h and at 100 o C under vacuum for other 12h.

The electrochemical performances of carbon support and carbon/Sb composites as negative electrode for Na-ion batteries were studied in coin cells assembled in an argon filled glove box. The electrolyte was NaClO4, 1M propylene carbonate (PC) as solvent and vinyl carbonate (VC)(1%) and fluoroethylene carbonate (FEC) (5%) as additives, which have proven to have a positive effect on the cycling of Sb particles [START_REF] Darwiche | Better Cycling Performances of Bulk Sb in Na-Ion Batteries Compared to Li-Ion Systems: An Unexpected Electrochemical Mechanism[END_REF]. The galvanostatic charge/ discharge measurements were carried out at room temperature on Mac pile (Biologic SA and MTI) battery testing system at room temperature. The capacity is reported per gram of C/Sb composite.

Results and discussion

Carbon support characterization

The structure of carbon materials may impact the Na insertion but also their conductivity, therefore was analysed in detail by several complementary techniques. The XRD patterns for the carbon supports (Figure 2a) exhibit three broad diffraction peaks, corresponding to carbon with low degree of graphitization. A slight increase of the peak intensity could be observed for the hard carbons treated at temperature higher than 1100 o C. The TEM pictures (Figure 2c,d,e,f) revealed a disordered structure for both mesoporous and hard carbons, but graphitic layers short domains are more obvious for the hard carbon, especially for the one treated at 1500 o C. Figure 2b shows the Raman spectra of carbon support. Two typical bands, D (1363 cm -1 ) and G (1600 cm -1 ), are visible and usually employed for the assessment of the disorder in the graphitic structure [42;43]. For mesoporous carbon, the D band is more intense and broader than the G band while for hard carbons the increase of the temperature induces an increase in the intensity of G band, sign of a more ordered structure. The ratio of ID/IG is 1.29

for the mesoporous carbon and decreases with the thermal temperature from 0.94 for HC@1100 down to 0.74 for HC@1300 and 0.73 for HC@1500, respectively. The Raman results confirm the XRD and TEM analyses which showed an increase of the carbon organization by increasing the temperature. MC@900 (c), HC@1100 (d), HC@1300 (e), HC@1500 (f)

The assessment of the carbon surface chemistry is one of the important factors since is impacting the interactions with the metal precursor salt and the stability of the particles during the synthesis but also the interactions with the electrolyte. The composition and the surface chemistry of the carbon supports (<10 nm thickness) was determined by XPS measurements and gathered in the Table 2. In the case of mesoporous carbon (MC@900), the amount of carbon is 95.7 at% and that of oxygen 4.3 at%. The deconvolution of the C1s peak (Figure S1a, Supporting Information) reveled the existence of ether (-C-OR), carbonyl (-C=O) and carboxyl (-COOR) groups at 286.1, 287.2 and 288.4 eV, respectively. The major contributions come from C-OR, COOR and then C=O. Similar groups and trend in term of quantities was found for hard carbon as well (Figure S1b,c,d, Supporting Information).

Concerning the amounts of oxygen in hard carbon, they were found to be higher than for mesoporous carbon and only slightly decreasing with the increase in the annealing temperature. This observation is rather unexpected since the increase in the temperature generally leads to the decomposition of functional groups and to the decrease in the amount of oxygen. However, it should be reminded that this composition is only in the extreme surface of the materials and the composition it might be different in the bulk. For this reason, the TPD-MS technique was employed to study the materials oxygen functionalities. The carbon materials were heated under vacuum and the evolved gases (H2, CO, CO2 and H2O) recorded as a function heating temperature. The functional oxygenated groups decomposed in CO and CO2 gases and their release temperature provides insights on their nature. Typically, the CO2 is derived from the decomposition of acidic groups, i.e., carboxyl or/and anhydride groups while CO is obtained by decomposition of more basic groups such as, carbonyl, phenol, ether and quinone groups [START_REF] Figueiredo | Modification of the surface chemistry of activated carbons[END_REF].

For the mesoporous carbon, MC@900, a large CO desorption peak ranged between 200 and 900°C is recorded (Figure 3a) suggesting the existence of a large variety of oxygenated groups such as anhydrides, phenols, ethers and quinones [41;45;46]. The CO peak is more intense and narrower between 600 and 900°C, indication of more predominant basic groups in the material such as ethers and quinones. The CO2 desorption profile is less intense than CO one, and is mainly observed between 200-500°C suggesting carboxyl and anhydride group [41;45;46]. These observations are in good agreement with the XPS results. Water release is also observed (100 -600°C), which can be assigned to physisorbed water in carbon porosity at low temperature and to chemisorbed water or in-situ formed water at higher temperatures. In addition, an intense peak of hydrogen is evolved at temperatures higher than 800°C suggesting the cleavage of the C-H bonds and the carbon structural organization. In the case of HC@1100 carbon, rather similar CO profiles are observed while the CO2 profile is on the contrary much larger indicated a greater number of anhydrides and lactones (Figure 3b)

besides the carboxyl groups. As for the water profile, the low temperature peaks are intense probably due to the high amount of micropores in this material, as will be demonstrated latter on. It is worth to note that hydrogen evolution is less important than in the case of mesoporous carbon due to the higher synthesis temperature and better organized structure. For HC@1500 (Figure 3c), the peaks are less intense compared to the other materials sign of a poor surface chemistry.

By integration of CO and CO2 desorption peaks, the amounts of the oxygenated functional groups can be determined (Table 2, Figure 3d). As a general trend, the amounts of CO are higher than CO2 for all materials, in well agreement with the desorption profiles intensity. The highest amounts of CO and CO2 are obtained for mesoporous carbon while for hard carbon they are decreasing with the increase of the annealing temperature (Figure 3d). The overall COx (CO + CO2, Table 2) quantities are following the same trend, and they are rather different from the one determined by XPS, confirming that the oxygen groups in the surface are rather different than in the bulk. It is possible that the higher amounts of water in the hard carbon may induce higher oxygen amount in the materials when analyzed by XPS. In the case of mesoporous carbon, functional groups maybe placed inside the pores and probably not accessible to be measured by XPS, resulting in lower oxygen amounts [START_REF] Slesinski | Self-buffered pH at carbon surfaces in aqueous supercapacitors[END_REF]. MC@900 HC@1100 HC@1300 HC@1500 200 400 600 800 1000 0,0 The textural properties of the carbon support were evaluated by nitrogen adsorption/desoption isotherms (Figure S2, Supporting Information). For MC@900, the N2 type IV isotherms characteristic to mesoporous materials were observed, confirming the role of Pluronic F-127 in the formation of mesopores (Vmeso = 0.47 cm 3 g -1 ). The existence of some micropores is indicated, as well, by the presence of the points in the low relative pressure region (Vmicro = 0.17 cm 3 g -1 ) with a surface area of 412 m 2 g -1 . In case of HC@T, the isotherms correspond more to type I, but not many points are observed in the low-pressure region, indicating the character almost non-porous for these materials in line with the SSA below 4 m 2 g -1 (Table 2).

Carbon/Sb hybrid materials

Taking into account the results obtained in our previous study on the influence of Sn loading [START_REF] Nita | Understanding the Sn Loading Impact on the Performance of Mesoporous Carbon/Sn-Based Nanocomposites in Li-Ion Batteries[END_REF], we considered 60 wt.% amount of Sb, the appropriate loading to obtain a good compromise between the Sb loading and the Sn particle homogeneity in the material.

The XRD patterns of the carbon/Sb composites exhibit several peaks which were indexed to the Sb metallic phase (COD 9008575) (Figure 4a). Therefore, the implemented synthesis pathway allowed to decompose the antimony acetate precursor into antimony oxides and their further reduction under hydrogen and by carbo-reduction reactions results in the formation of metallic Sb, according to the following reactions:

Sb(CH3COO)3 => Sb2O3 + (C-COOH) + H2O+COx (1) Sb2O3 + H2 +C => Sb +H2O +COx (2)
Material SSA m The composition and thermal/chemical stability of carbon/Sb hybrid materials were determined by TGA in air (Figure 4b). The TGA curve for the mesoporous carbon presents a slight increase of the mass weight followed by two main mass weight losses observed at ~ 400 and 600 °C, associated to the oxidation of carbon (C+O2  CO2). Regarding the Sb/HC@T, the gain in mass before 500 o C is much more evident. In both cases, this gain could be associated to the oxidation of the Sb particles (Sb + O2 SbO2), in agreement with other works [START_REF] Liao | Ultrafine Sb nanoparticles embedded in an amorphous carbon matrix for high-performance sodium ion anode materials[END_REF], [START_REF] Yi | Enhanced Li-and Na-storage in Sb-Graphene nanocomposite anodes[END_REF]. In contrast with the mesoporous material, where this increase was not well visible (more likely because the particles are situated in the pores), for the hard carbon/Sb composites, the particles situated on the carbon surface determines firstly the oxidation of the metallic Sb, and then the carbon combustion. The derivatives for the TGA curves (not shown here) present mainly one exothermic peak for each material, but also one endothermic peak. 2 indicate that the synthesis route allowed preparing materials with high loading of Sb and close to the calculated ones. Insights on the structure and morphology of the different materials were assessed by STEM in dark field. The Sb particle size have a great influence on Na + diffusion pathways and on the storage mechanism as indicated in the introduction part. When using mesoporous carbon framework, the Sb nanoparticles are homogeneous distributed in the carbon framework (Figure 5a) and they have small size (< 5 nm) as better visualized in the STEM images provided in the Figure S4a, Supporting Information. . For Sb/HC@T the homogeneity seems to strongly depend on the annealing temperature of the carbon support. The Sb/HC@1100 material presents a homogeneous distribution of Sb nanoparticles (Figure 5b) with slightly higher size than for mesoporous carbon and some larger particles observed (10-20 nm) as well as seen in Figure S4b, Supporting Information. A closer look on a large particle by HRTEM allow to visualize the crystallite lattice fringes, however, on small particles it was difficult to observe it (Figure S5, Supporting Information). Increasing the annealing temperature of the carbon support to 1300 and 1500 o C, respectively, a significant increase in particle size is observed (up to 20 nm), along with a tendency to form agglomerates (up to 50 nm) (Figure 5c,d).

Figure 5. STEM images for Sb/MC@900 (a), Sb/HC@1100 (b), Sb/HC@1300 (c) and Sb/HC@1500 (d).

The precise determination of the particle size evolution is rather difficult to realize due to the high density of the particles. However, this behavior may be explained on one hand by the decrease of the amount of oxygen groups of carbon with the increase of the annealing temperature which decreases the hydrophilicity of carbon support and of the functional groups/ active sites serving as anchoring points for the antimony acetate precursor. On the other hand, an important role on the agglomeration of Sb particles could be played by the decrease of the pore volume and with the temperature, as will be seen in the following part.

The textural properties of the composite materials were evaluated by N2 physisorption technique at 77K (Figure 6) and the corresponding textural data are gathered in Table 2. The surface area and the pore size/volume are carbon characteristics which may strongly impact not only the composite synthesis but also its electrochemical performances. A high pore volume and surface area can be beneficial since it may allow a good diffusion of the electrolyte but also can accommodate the Sb nanoparticles during synthesis and electrochemical cycling. On the contrary, the high surface area may induce undesirable reactions of the support with the electrolyte. Therefore, such characteristics must be determined. Type IV isotherms characteristic to mesoporous materials were observed in the case of the Sb/MC@900. The textural properties, namely the specific surface area and the pore volume decreased after the Sb addition from 412 to 208 m 2 g -1 (Table 2) and respectively, from 0.64 to 0.34 cm 3 g -1 . For the Sb/HC@T materials, surprisingly, the Sb addition induces a slight increase of the textural properties (SSA and pore volumes) compared to the values for the HC@T support (Table 2). The Sb/HC@1300 and Sb/HC@1500 materials present an increase of the adsorbed nitrogen volume at high relative pressures (P/P0) indicating an inter-particle porosity due to the larger amount of Sb particles on the carbon surface. The pore size distribution (PSD) (Figure 6b) determined from the adsorption branch of N2 isotherm indicates the existence of both micro-(0.6 nm) and meso-pores (7.5 nm) for all composite materials but also pores larger than 20 nm corresponding rather to the inter-particle porosity.

Electrochemical performances

The electrochemical tests for carbon/Sb materials were performed at C/10 constant current rate in the voltage range of 0.01-1.5V (vs. Na/Na + ). Firstly, the carbon supports were tested in the same conditions as the composites materials for comparison purposes in order to evaluate their contribution to the capacity (Figure 7a). The carbon supports revealed a high irreversible capacity in the first cycles, representing 91-94% from the initial capacity no matter the material. However, it is worth to note that in the case of mesoporous carbon, MC@900, the irreversible capacity is significantly much higher than for the hard carbon, HC@1100 i.e., 738 mAh g -1 vs. 64 mAh g -1 (Figure 7c and Table S1). This may be related in large extent to the higher specific surface area but also to the richer surface chemistry of mesoporous carbon which induce more electrolyte decomposition and SEI layer formation [49;50]. Such higher irreversible capacity of the carbon support may negatively impact the irreversible capacity of the MC@Sb composite as well.

Regarding the reversible capacity, it was found to be dependent on the carbon support. The mesoporous support presents a higher reversible capacity (55 mAh g -1 ) after the 100 cycles compared to the hard carbon supports, where the reversible capacities decrease with the annealing temperature of the carbon from 13 mAh g -1 for HC@1100 to 4 mAh g -1 for HC@1500 (Figure 7a). This decrease of the capacity with the temperature could be associated with the higher hydrophobicity for the materials treated at high temperatures and the slight increase of the graphitization level of the carbon. The reversible capacity values are particularly very low for hard carbon materials since usually materials with similar specific surface area and structure delivers much higher capacity (> 200 mAh g -1 ) [START_REF] Beda | Hard carbons derived from green phenolic resins for Na-ion batteries[END_REF]. Such different behavior may come from the different electrochemical conditions used herein (binder, electrolyte, additives) which were optimized for Sb anodes [START_REF] Darwiche | Better Cycling Performances of Bulk Sb in Na-Ion Batteries Compared to Li-Ion Systems: An Unexpected Electrochemical Mechanism[END_REF] and which do to favor good performances for hard carbons. Considering the low reversible capacity of the pristine carbons, their contribution to the capacity of the C/Sb composite materials can be therefore neglected. For all carbons supports, the coulombic efficiency is quite low in the first cycle, but after 20 cycles it strongly increases to more than 90% and after 30 cycling reaches 100%.

Generally, for the C/Sb composite materials, the irreversible capacity was significantly reduced compared to the carbon supports (Figure 7b and Table 3). This can be understood by the lower amount of carbon contained in the composites (~ 50 %) which limits the irreversible reactions with the electrolyte. Moreover, the irreversibility of the composite materials which use hard carbon as support was successfully reduced compared to the mesoporous materials (48% for Sb/MC@900 and 25-29% for Sb/HC@T) (Figure 7e). This decrease can be related to the lower SSA in case of the hard carbon, which in contact with the electrolyte reduce the SEI formation. A slight contribution to the irreversible capacity could come also from the small quantity of Sb oxides covering the Sb nanoparticles a well-known phenomenon for nanoparticles in contact with air (see Raman spectra provided in Figure S6, Supporting Information). Concerning the reversible capacity in the first cycle, the Sb/MC@900 presents the highest specific capacity (639 mAh g -1 ), while for Sb/HC@T the values decreases with the annealing temperature from 535 mAh g -1 for Sb/HC@1100 to 419 mAh g -1 for Sb/HC@1500. It is worth noting that in most of the literature works regarding Sb/C composites, the cycling performances are presented only for a limited number of cycles (Table 1), which do not allow to assess the long-term behavior of such materials. We have therefore tested the materials for the long-term cycling, and during the 75 cycles a decrease of the capacity is observed for all materials. After the 250 th cycle, another strong decay is noticed and for 500 cycles the materials kept only about 8% from their initial capacity for Sb/MC@900, Sb/HC@1300 and Sb/HC@1500 and 17% for Sb/HC@1100, respectively. The most well performing material from cycle stability point of view is Sb/HC@1100 material which, moreover exhibits the highest reversible capacity. A comparison between the reversible capacities of Sb/HC@1100 and Sb/MC@900 is presented in Figure 7e, f. The better performances obtained for the hard carbon/Sb composite as support could be explained by a good compromise between the low surface area and the good dispersion of the Sb particles in the carbon framework. Both these parameters may impact the performances. The Sb particles size are small and well dispersed in both composites ensuring short electrolyte pathways and insertion of Na. However, the carbon supports are rather different in terms of porosity (specific surface area and pore size distribution). The mesoporosity of MC carbon allow better diffusion of the electrolyte but in the same time the associated high surface area induce high irreversible capacity and probably thick and unstable SEI layer which are detrimental during long cycle. On the contrary, the low surface area in the case of HC ensures a low irreversible capacity related to an efficient SEI.

The better structural organization of the HC carbon (see Figure 2) may also have a positive impact on the conductivity of Sb nanoparticles [START_REF] Wu | Binding low crystalline MoS2 nanoflakes on nitrogendoped carbon nanotube: towards high-rate lithium and sodium storage[END_REF]. Moreover, it has been shown [START_REF] Liu | A waste biomass derived hard carbon as a high-performance anode material for sodium-ion batteries[END_REF] that the increase of the annealing temperature of hard carbon improves its electronic conductivity.

However, for the hard carbon supports annealed at higher temperature, i. Taking into account the as-presented evolution, we can conclude that there is no need to use sacrificial templates in order to create mesopores in the carbon framework. The advantage of the mesopores for improving the electrolyte diffusion and the accommodation of Sb particles it is contra balanced by its higher SSA which determines higher irreversibility.

Moreover, there is no reason to treat the carbon support at high temperature (1300 o C and 1500 o C, respectively) since the Sb agglomeration phenomena negatively impact the electrochemical performances. The coulombic efficiency of the C/Sb composite materials reaches fast 100% (after 10 cycles), which means that parasitic reactions take place because of the electrolyte degradation are limited after first cycles. That being said, the Sb/HC@1100 could be considered the optimal composite for the proposed applications.

Nevertheless, to better understand this fading behavior which is rarely discussed, Figures 8 presents the galvanostatic curves and derivative curves for Sb/MC@900 and Sb/HC@1100 for several specified cycles. For the Sb/MC@900 (Figure 8a,c), the first discharge profile shows a gradual voltage drop to 0.4 V, followed by a plateau and another strong decrease down to 0.01 V, delivering a discharge capacity of 639 mAh g -1 . In the same time, 55% of the inserted Na can be reversibly desodiated, giving a charge capacity of 335 mAh g -1 . The subsequent discharge curves are very different compared to the first discharge because of the formation of SEI on the surface of the anode. The differences in shape and potential between the first insertion and the following ones are similar to the classically observed ones for Sb system and likely attributed to the "electrochemical grinding" occurring in the first discharge. This phenomenon induces the active material particle size grinding leading to smaller and sometimes amorphous particles [START_REF] Darwiche | Facile synthesis and long cycle life of SnSb as negative electrode material for Na-ion batteries[END_REF], [START_REF] Mao | Active/Inactive Nanocomposites as Anodes for Li-Ion Batteries[END_REF]. The next curves are almost overlapped indicating the good reversibility of the chemical reactions in the electrode. For the discharge curve which corresponds to the 75 th cycle, a faster decrease of the potential is observed, corresponding to the significant decrease of the capacity which took place during the first 75 cycles.

The Figure 8b, d presents the galvanostatic and derivative curves for Sb/HC@1100 for selected cycles. The first discharge profile shows a rapid voltage drop to 0.8 V, followed by a gradual decrease down to ≈ 0.45V (Figure 8b). A plateau is observed at this value, followed by another strong decrease down to 0.01V, delivering a discharge capacity of 535 mAh g -1 . The next curves present the same voltage drop to 0.8 V, from here a slow continue decrease of the voltage is observed down to 0.45, followed by a strong decrease to 0.01V. The curves corresponding to the second and third cycles are overlapped indicating the reversibility of the chemical reactions, while the curve for the 75 th cycle presents faster decrease to 0.01V. The 3 cathodic peaks observed in the derivative curves (Figure 8d) situated between 0.38 and 0.72V correspond to the multi-step formation of amorphous NaxSb and crystalline Na3Sb alloys at the end of discharge [START_REF] Darwiche | Better Cycling Performances of Bulk Sb in Na-Ion Batteries Compared to Li-Ion Systems: An Unexpected Electrochemical Mechanism[END_REF] according to the following equations:

Sb + xNa + + xe -↔ NaxSb (3) NaxSb + (3-x) Na + + (3-x) e -↔ Na3Sb (4)
For the first 3 anodic cycles, an intense peak is seen at 0.78 V and a small oxidative peak at 0.90 V which splits in the 75 th cycle into two peaks located at 0.81 and 0.91 V characteristics to the desodiation of crystalline Na3Sb. At 300 cycles, only one peak is observed which is shifted at 0.94 V, is broader and less intense, and further vanished at 500 cycles. The same behavior is seen for Sb/MC@900, but in this case, the 0.94 V peak almost disappeared after the 300 cycles. The disappearing of the peaks can be related to the observed fading of the two materials which occurred earlier for Sb/MC@900 compared to Sb/HC@1100. The reason behind the disappearing of the electrochemical peaks may be the amorphisation of Sb particles as observed in the literature [START_REF] Darwiche | Better Cycling Performances of Bulk Sb in Na-Ion Batteries Compared to Li-Ion Systems: An Unexpected Electrochemical Mechanism[END_REF].

In order to better understand the changes which took place in the anode during the charge-discharge cycling, STEM analyses were performed on the post mortem electrodes (Figure 9). Two materials were selected which present very different physico-chemical properties in order to distinguish easier between the factors which could impact the electrochemical performances. So, the STEM images were taken for Sb/MC@900 (Figure 9a) and Sb/HC@1500 electrodes recovered after 500 cycles (Figure 9b). The post mortem Sb/MC@900 material revealed a good homogeneity of Sb particles in the carbon matrix, comparable with that of the material before cycling (Figure 5a), confirming the well confinement of the Sb particles in the mesoporous structure. However, few agglomerates of particles are formed during the cycling as well which are no more electrochemically active after several cycles and could be partially responsible for the decay of the reversible capacity after 75 cycles. Moreover, the decay after 300 cycles could be also associated to the amorphization on Sb nanoparticles. On the contrary, the post mortem Sb/HC@1500 electrode shows more agglomerated particles (Figure 9b) compared to Sb/MC@900 composite, the material being more prone to form agglomerates even before the cycling (higher particle sizes and agglomerates before cycling, Figure 5d). In addition, the benefits of the better dispersed particles for the Sb/MC@900 material competes with the weakness due to the high surface area of the mesoporous carbon, which interacts with the electrolyte and forms continuously the SEI layer, originating a significant decay of the capacity in cycling. For the Sb/HC@1500, the irreversible capacity is significantly lower, thanks to the lower textural properties of the carbon support, but due to the larger Sb particles which are partially agglomerated [START_REF] Jahel | Confined Ultrasmall SnO2 Particles in Micro/Mesoporous Carbon as an Extremely Long Cycle-Life Anode Material for Li-Ion Batteries[END_REF] the material loses in performances. Therefore, both mesoporous and hard carbon could be successfully used as support to prepare carbon/Sb anodes for Na-ion batteries, by making a balanced compromise between the amount and the distribution of the Sb particles in the carbon matrix, and the characteristics of carbon support. Well dispersed particles obtained at moderate Sb loading (~50%) and annealing temperatures of the carbon matrix could favor the diffusion of the electrolyte on the anode surface. On the other hand, the carbon porosity allows the particle confinement and accommodate the particles limiting their agglomeration during the cycling, but cannot totally prevent it, few aggregates being formed no matter the carbon support (more for hard carbon). However, in the same time this porosity could react with the electrolyte and form a SEI layer inducing a high irreversible capacity. Therefore, small and well dispersed Sb particles on a suitable carbon support is not sufficient to ensure long term cycling. The emplacement of the particles in the carbon network can be as well an important factor. Ji et al.

[34] succeed to avoid the fading of Sb/C composites by embedding carbon Sb nanoparticles in an amorphous carbon network. Such confinement of Sb particles in the carbon structure rather than in the carbon porosity or in the surface may better limit the particle growth and aggregation. Such hypothesis can be sustained by other two works [START_REF] Zhang | Spherical nano-Sb@C composite as a high-rate and ultra-stable anode material for sodium-ion batteries[END_REF], [START_REF] Zhu | Electrospun Sb/C Fibers for a Stable and Fast Sodium-Ion Battery Anode[END_REF], where the Sb particles are trapped in the carbon structure by simultaneous synthesis of Sb and carbon in one-pot synthesis providing long term cycling performances.

Conclusions

Carbon materials presenting different structure, porosity and surface chemistry were synthesized by eco-friendly sol-gel methods and used to support Sb nanoparticles by ball milling-thermal reduction approach. The influences of the presence/absence of mesopores and the annealing temperature on the physico-chemical and electrochemical properties of carbon/Sb hybrid materials was investigated. Small and well dispersed Sb nanoparticles were obtained for mesoporous carbon and hard carbon supports treated at low temperature. By increasing the annealing temperature from 1100°C to 1500°C, the carbon porosity and surface oxygen-based functionalities decreases inducing the Sb particle size increase and agglomeration. The electrochemical tests performed at a C/10 current rate in a voltage range of 0.01-1.5 V (Na/Na + ) prove a successfully reduction of the irreversible capacity of mesoporous carbon and hard carbon from ≈ 91-94% to 48% and ~ 25-29% by preparing Sb/MC@900 and Sb/HC@T hybrid materials. The reversible capacity reached 535 mAh g -1

for Sb/HC@1100 after the first cycle, while the other composites delivered lower capacities.

For the long term cycling, the Sb/HC@1100 composite combining both low surface area, small Sb nanoparticles and optimal structure exhibited the most stable behavior up to 300 cycles, while in the case of the other materials fading occurred in the first 75 cycles. In addition, for all materials a drastically capacity fading is observed at around 300 cycles with almost complete loss of the capacity. This behavior could be attributed to the Sb particle size agglomeration and amorphization during cycling.

Figure 1 .Figure 1 .

 11 Figure 1. Schematic representation of the synthesis protocol developed for carbon/Sb

Figure 2 .

 2 Figure 2. XRD patterns (a) and Raman spectra for carbon supports (b); TEM pictures for

Figure 3 .

 3 Figure 3. TPD-MS desorption profiles of mesoporous carbon, MC@900 (a) hard carbon,

Figure 4 .

 4 Figure 4. XRD patterns of carbon/Sb hybrid materials prepared using different carbon

  The peaks are shifted to higher temperature in function of the thermal treatment temperature of the carbon support. The presence of two peaks corresponding to the carbon combustion could indicate that the size of the Sb particles (small particles/aggregates) or the place where they are located (inside/outside of the pores) influence differently the carbon combustion. The real amount of Sb was calculated by taking into account the SbO2 quantity obtained after the TGA at 900 o C in air (FigureS3, Supporting Information) and the Sb/O atomic ratio. The values presented in Table

Figure 6 .

 6 Figure 6. N2 adsorption/ desorption isotherms (a) and pore size distribution (b) for carbon/Sb

Figure 7 .

 7 Figure 7. Electrochemical performances of the carbon support and carbon/Sb materials (voltage range 0.01-1.5V (Na + /Na), C/10 current rate). The specific discharge capacity vs. the

Figure 8 .

 8 Figure 8. The galvanostatic (a, b) and derivatives (c, d) curves for Sb/MC@900 and Sb/HC@1100 materials for the specified cycles (voltage range 0.01-1.5V (Na + /Na), C/10

Figure 9 .

 9 Figure 9. STEM images after 350 cycles for Sb/MC@900 (a) and Sb/HC@1500 (b).

Table 1 .

 1 Literature review on Sb/C hybrid materials synthesis methods and performances for Na-ion batteries.

	Materials	Synthesis method	Electrochemical performances vs. Na	Refs.
			Current	Crev	Cycle
			density	(mAh g -1 )	number
			(mA g -1 )
	Bulk Sb	Commercial	4C	550	160	[22]
	Sb/C	Sol-gel with citric acid	4C	115	50	[35]
	Sb/C	Sol-gel with chitosan	500	403	100	[36]
	Sb/C	Sol-gel	100	423	200	[27]
	Sb@N-doped C	Sol-gel	50	305	60	[30]
	sheets
	Sb/CNFs	Impregnation/chemical	100	542	100	[37]
		reduction
	Sb/graphenes	Chemical reduction	100	241	10	[29]
	Sb/graphene/CNFs		100	274	100	[23]
	Sb/CNFs	Electrospinning	200	600	80	[24]
	Sb/CNFs		100	650	300	[25]
	Sb nanorode arrays	Template	1000	574	250	[13]
		electrodeposition
	Sb/porous carbon	Salt template	100	456	500	[38]
	Sb nanocrystals	One-pot colloidal	660	580	100	[14]
	Sb/amorphous C	Hydrothermal route	100	622	110	[26]
	Sb/MWCNTs	Wet milling	200	400	100	[31]
	Sb/C black	Mechanical milling	2000	610	100	[32]
	Sb/C graphite	Mechanical milling	100	420	100	[33]
	Sb/C yolk-shell	Polymerization	50	554	200	[39]
	Sb/N-doped C	Ionic liquids	100	440	100	[28]
	NanoSb/C	Aerosol spray pyrolysis	100	385	500	[12]
	Sb/C roads	Self-thermal-reduction	100	150	200	[34]
	Sb-O-C/C	Thermal-reduction	100	500	200	[34]

Table 2 .

 2 Properties of the mesoporous and hard carbon/Sb materials. Specific surface area (SSA) determined by BET method, chemical composition assessed by XPS, COx functional groups obtained by TPD-MS and Sb amount determined by TGA analyses.

Table 3 .

 3 Electrochemical performances for mesoporous and hard carbon/Sb materials.

	Material	C 1st	C 1st	C irrev	C rev 2 nd	Crev 100 th	Crev 300 th	Crev 500 th
		discharge	charge	(mAh g -	cycle	cycle	cycle (mAh	cycle (mAh
				1 )	(mAh g -1 )	(mAh g -1 )	g -1 )	g -1 )
	Sb/MC@900	639	335	48%	354	224	117	49
	Sb/HC@1100	535	381	29%	398	377	285	92
	Sb/HC@1300	462	347	25%	362	271	95	37
	Sb/HC@1500	416	311	25%	323	253	114	28

Acknowledgements

The authors acknowledge the financial support of this work from University of Haute-Alsace, France and European Union's Horizon 2020 Program (project NAIADES, call: LCE10-2014, Contract no. 646433). Jean-Marc Le Meins, Samar Hajjar-Garreau and Bénédicte Réty are acknowledged for the help provided with the XRD, XPS and TPD-MS analyses.

Data availability

The raw/processed data required to reproduce these findings cannot be shared at this time due to legal or ethical reasons.

Competing financial interests

The authors declare no competing financial interest

Appendix A. Supporting information

Supplementary data associated with this article can be found in the online version at doi:XX.XXXX/XXX.


Cristina Nita

Julien Fullenwarth

Laure Monconduit

Jean-Marc Le Meins

Julien Parmentier

Moulay Tahar Sougrati

Camélia Matei

carbon/Sn based nanocomposites in Li-ion batteries

Keywords: carbon/Sn hybrids, mesoporous carbon, tin-based nanoparticles, Li-ion batteries

Understanding the Sn loading impact on the performances of mesoporous

Introduction

During the last years, the interest for environmentally friendly processes and alternative sources dedicated to energy generation and storage has significantly increased. For electrical vehicles and portable electronics applications, the Li-ion batteries have been considered the most important power sources due to their high energy density, operating voltages and lightweight design [1] . However, the fast development of the technology demands high performing batteries, but in the same time safety and low-cost.

A Li-ion battery contains several components, each of these contributing in different ways to the improvement of the battery performances, the anode inducing a significant impact [2] . Nowadays, graphite is the most used anode in Li-ion batteries, with a theoretical capacity limited to 372 mAh g -1 , because of the limited Li storage capacity (one Li atom for six carbon atoms) [3] . In order to improve the capacity as compared to graphite an increasing interest was focused on other anode materials. For such purposes, the Sn or Sn-based materials have been considered among the best candidates, especially due their higher theoretical capacity (781 mAh g -1 for SnO2 [START_REF] Courtel | Nanotechnology for Lithium-Ion Batteries[END_REF] and 990 mAh g -1 for Sn [START_REF] Wang | [END_REF] ). However, such metals suffer an important volume expansion during lithiation step and therefore, particles agglomeration and electrode pulverization occurs along with the loss of electrochemical performances. These drawbacks could be overcome by downsizing the Sn particle size [6] and by using different carbon materials (carbon nanotubes [7] , [8] , carbon fibers [9] , carbon nanospheres [10] , graphenes [11] , graphite [12] , mesoporous carbon [3] , [13] ) as supports for Sn-based hybrid materials. This allow to alleviate the macroscopic volume expansion of the composite material and, in the same time, increase its electrical conductivity and provide efficient pathways for electrolyte diffusion as demonstrated by several works [14] , [15] , [16] , [17] .

Wu et al. [18] reported the preparation of SnO2-CNTs using glucose through a hydrothermal approach, followed by a carbonization at 500 o C in nitrogen atmosphere. The materials present a discharge capacity of 400 mAh g -1 at current density of 100 mAg -1 , after 50 cycles, slightly higher than that of graphite, but lower than the theoretical capacity of the composite because of the solid-electrolyte interphase (SEI) formation. Cui et al. [17] reported the preparation of Sn/onion-like carbon nanocapsules by arc discharge method with a reversible capacity of 585 mAh g -1 after 100 cycles at very low current rate, 100 mA g -1 , while Han et al. [16] succeeded to increase the capacity up to 902 mAh g -1 (at 100 mA g -1 after 1000 cycles) by introducing nitrogen (N) atoms in the carbon framework. The high capacity and the stable cycling could be promoted by the very low current rate used. Another limitation is the complexity of the synthesis procedure which involves specific pressure and temperature conditions. Guo et al. [10] prepared Sn nanoparticles embedded in N-doped hollow nanosphere which delivered 500 mAh g -1 at 5C after 500 cycles. Nano size tin particles embedded in porous multichannel carbon microtubes were synthesized by electrospinning and provided a capacity of 350 mAh g -1 , at 2 C rate after 600 cycles [19] . Liu et al. [20] repported capacities up to 520 mAh g -1 at 500 mA g -1 , after 1000 cycles, for Sn/N-doped carbon nanospherers prepared by arc discharge method. Jahel et al. [21] , [3] pointed out that the confinement of electrochemically active SnO2 nanoparticles (ca. 2 nm in diameter) in the pores of a mesoporous carbon is the key parameter to avoid the particle growth and agglomeration, affording high capacity and long stable cycling life (up to 2000 cycles). The improvement of the C/SnO2 anode materials performances was achieved, with a reversible charge capacity up to 443 mAh g -1 at 1C current rate after 2000 cycles. However, the irreversible capacity of the first charge-discharge cycle was still high, the later one remaining still an important problem of such materials and is mainly related to several factors as enumerate next: (i) solid electrolyte formation due to the reactions of the electrolyte with the high surface area of the porous carbon support, ii) the electrolyte consumption due to the reduction of tin oxide species to Sn metallic phase, according to the reaction described in [3] and iii) the small amount of active phase (40 wt.% SnO2, corresponding to 31 wt.% Sn) in the composite, and therefore, the high amount of carbon (with a low electrochemical activity after the first discharge).

In order to overcome these drawbacks, the present study aims to systematically investigate the influence of Sn-based nanoparticles loading on the irreversible capacity, the reversible capacity and the cycle stability of carbon/Sn hybrids. As far as we know and highlighted by the literature review gathered in Table 1, no systematic study have been performed until now for Sn/C hybrids to evaluate how the Sn/SnO2 loadings can influence the formation of Sn/C hybrid materials (particle size, surface chemistry and porosity) and further impact their electrochemical performances in Li-ion batteries. In addition, the contribution of each component (carbon and Sn) to the irreversible/reversible capacity and the long term cycling is not addressed either to our knowledge.The improvement of irreversible and reversible capacity in the first cycles compared to the commercial graphite anode by adding high amounts of Sn, as well as the capacity fading behavior with further cycling is discussed in the light of material characteristics (nanoparticle dispersion/confinement in the carbon mesopores, amount of SnO2 present in the nanoparticles and specific surface area).

Results and Discussion

Material characterization

Mesoporous carbon materials with embedded Sn nanoparticles were synthesized and characterized in order to study the influence of Sn loading on the C/Sn hybrid material characteristics and performances. The idea of using carbons with tunable pore size was motivated by our previous work [3] , [21] where we have shown that the confinement of small SnO2 nanoparticles inside the pores is the key parameter to avoid NPs agglomeration and to ensure long and stable cycle life in the battery.

The XRD pattern of mesoporous carbon exhibits three diffraction peaks, corresponding to the hexagonal graphite structure (Figure S1a, Supporting Information). The presence of broad peaks indicates a low degree of graphitization. SAXS patterns (Figure S1b, Supporting Information) display an intense peak at 0.97 o (2 theta) sign of the presence of a mesoporous structure which is confirmed by TEM images reveling a wormlike structure with mesopores having relatively uniform size (Figure S1c, Supporting Information).

For C/Sn hybrid materials, the XRD patterns exhibit several intense peaks which were all indexed to the tetragonal Sn phase (pdf 00-004-0673) (Figure 1). A small quantity of SnO2 is also formed as evidenced by the small and large (211) peak at 51 o (2theta), noted ♦ [19] .This oxide is probably present as a layer formed around the Sn nanoparticles and it could be formed during the "exposure" of the sample in air, since it is well known that smaller the nanoparticles are, easier their oxidization is [25] but also can be due to incomplete reduction of the Sn-based precursor. The nature and amount of such oxide layer is worth to be investigated in more detail since it is electrochemically active in the battery and contributes to the active mass of the samples. If the structure of the materials was not changed by varying Sn loading, the size of the Sn crystallites seems to be significant influenced. Particle size information was extracted by two complementary experiments: TEM and powder XRD. For this last one, the crystallite sizes which are related to size of domains over which diffraction is coherent (later called "apparent size") were extracted from the FWHM (full width at half maximum) of individual diffraction line profile via WinPlotr [26] , DIFFRAC.EVA [27] and then using the Scherrer formula [28] . Main objective is to obtain a trend for the same family of compounds. Instrumental resolution function of the diffractometer was taken into account for the calculation because of the rather low FWHM values but let us keep in mind the limitation of the method according to the different hypothesis it involves (Gaussian peak profiles, size distribution not too large, etc…). The selected Sn reflection for this study is (200) which is the most intense, not too affected by asymmetric profile and with no peak overlap. Anisotropic line broadening is not observed for all these hybrid materials. This study allows to see the variation of the apparent size of the Sn crystallites as a function of the Sn loading (Table 2). For materials with different Sn loading, the apparent crystallite size increases with the Sn content from 67 nm for 20 wt.% of Sn to 148 nm for 80 wt.% of Sn. This tendency is in agreement with Oswald ripening mechanisms and in line with our previous work [29] , [30] .

Insights in the structure and morphology of the different materials were assessed by STEM (Figure 2). All the composites display small particles, uniformly distributed in the carbon matrix, but also bigger particles are present as aggregates. The HRTEM images allow to estimate a size of primary particles of about ~5nm (Figure S2e, Supporting Informations). The homogeneous distribution of the smallest particles suggest that the wet impregnation/reduction protocols allow a good diffusion of Sn 4+ and BH4 -species in the porous carbon framework. Few aggregates with diameters ranging between 50 to 200 nm can be distinguished, especially for materials with high Sn loading (70%Sn@C and 80%Sn@C). For the larger particles, their size increases with the loading as shown previously [29] , [30] and also observed by XRD microstructural analysis (Table 2). The large size of the agglomerated particles justifies also the high and very sharp XRD peaks (Figure 1).

In addition to TEM, the energy dispersive X-ray (EDX) mapping of the C/Sn-based materials confirms the uniform dispersion of Sn in the carbon support as well as the presence of aggregates (Figure S2, Supporting Information). The most encountered problem, and rarely evaluated in the literature was related to the quantification of the Sn particles loading in the carbon framework and usually this problem is not addressed in many works. By considering the XRD results, it was demonstrated that the materials contains carbon, metallic Sn and SnO2. The composition and thermal/chemical stability of C/Sn-based hybrid materials were determined by TGA in air (Figure S3a, Supporting Information), which allows the combustion of carbon and determination of final SnO2 quantity, coming both from the initial SnO2 content, but also from the oxidation of metallic Sn during TGA under air (Table 2). The amounts of SnO2 obtained varies between 14 and 89 wt.%. However, only the TGA results do not allow to discriminate between the initial Sn and SnO2 in the hybrid materials. Therefore, complementary experiments to determine Sn/SnO2 ratios of the materials were performed by Mössbauer spectroscopy (Figure 3, Table 2). Spectra obtained with different Sn loadings and carbon supports were fitted with 2 unresolved doublets corresponding to SnO2 (green) and metallic tin (blue). SnO2 has an isomer shift of 0.0-0.01 mm/s and a quadrupole splitting of 0.51-0.57 mm/s, while tetragonal Sn (β-Sn) has an isomer shift of 2.51-2.53 mm s -1 and a quadrupole splitting of 0.14-0.26 mm s -1

. The results confirm the presence of both compounds: Sn and SnO2. The Sn/SnO2 ratio measured by Mössbauer spectroscopy increases in the increase of Sn loading in the composite (ratio from 31.6/68.4 to 60.5/39.5 Sn/SnO2 for Sn loading ranging from 40% to 60% wt, Table 2). For high loadings (samples with 70 and 80 wt.% Sn), increase of the Sn loading does not induce a significant evolution of the Sn/SnO2 ratios.

The amount of formed SnO2 may be explained by two factors, i.e, the particle size and particle location in the carbon network.On one hand, the Sn crystallites size increases from 67 nm to 148 nm (Table 2) when increasing the loading from 20 to 80% and the oxidation tendency generally increases with the decrease of the particle/crystallite size [25] , [31] . However, starting with 60% of loading, a slight increase of SnO2 amount is observed (Table 2) and this may be probably related to the higher amount of agglomerated Sn particles situated on the carbon surface, therefore, not confined in carbon pores (Figure 2) and more exposed to oxidation.

For 20%Sn@C sample, the metallic Sn quantity is so low (limit of detection) that it can hardly be detected by Mössbauer, especially because of the Lamb-Mössbauer factor which is ten times smaller for Sn than for SnO2 [32] . So, for the future calculations and experiments it will be assumed that the 20%Sn@C sample contains only SnO2.

Based on Mössbauer results which provided the Sn to SnO2 ratio in each composite, the amounts of Sn and SnO2 could be determined taking into consideration the TGA results and further used to correctly asses the electrochemical performances (Table 2). The composition and chemical oxidation state of Sn on the surface of the materials (<10 nm thickness) were determined by XPS measurements (Figure S3b, Supporting Information and Table 3). Thus, the C1s, O1s and Sn3d peaks were detected. For all materials, two peaks are observed on the Sn 3d spectra and the separation binding energy between them is always 8.4 eV, which is in good agreement with the energy splitting of Sn and SnO2 [21] . The shift of the peak positions for 20%Sn@C material could be induced by the high amount of SnO2, confirmed also by Mössbauer spectroscopy (Table 2), which is less conductive than the metallic Sn. The Sn amount increases by increasing the loading from 26.7 wt.% for 20%Sn@C to 72.1 wt.% for 80%Sn@C, while the atomic ratio O/Sn is comprised between 1.66 and 1.9 which confirms the presence of a SnO2 layer in the surface of the material (Table 3).

The textural properties of the carbon and composite materials (specific surface area, porous volume and pore size) were evaluated by N2 physisorption technique (77K). A mixture of type I/IV isotherms characteristic to micro/mesoporous materials [3] were observed in all cases (Figure 4a). The low relative pressure up-takes evidence the presence of micropores (size < 2 nm). The microporous volume (Vmicro) of the carbon support is about 0.28 cm 3 g -1 , with a high BET specific surface area (SSA) of 665 m 2 g -1 (Table 4). The total porous volume (VT) for carbon is 0.97 cm 3 g -1 , indicating a mesoporosity size of 0.69 cm 3 g -1 which could allow the accomodation of the Sn-based nanoparticles. The Sn addition induces a decrease of the N2 adsorbed volume proportional with the Sn loading (Figure 4a). The evolution of the porous volume and SSA with the Sn loading are presented in the Figure 4c, d andTable 4. The Vmicro decreases from 0.21 cm 3 g -1 for 20%Sn@C to 0.07 cm 3 g -1 for 80%Sn@C, while the SSA go down to 181 m 2 g -1

. The VT presents the same decreasing trend from 0.79 cm 3 g -1 for 20%Sn@C to 0.17 cm 3 g -1 for 80%Sn@C, inducing the same behavior for the Vmeso values. As can be noticed (Fig. 4c), the mesoporous volume decrease more than microporous volume when increasing the Sn loading, suggesting a more preferential location of the particles in larger pore. This is rather related to the size of the Sn particle ~5 nm which can be better accommodated in moesopores having sizes of about 8 nm as will be discuss latter than in micropore having size < 1nm. The pore size distribution (PSD) was obtained from the adsorption branch of the N2 isotherms using 2D NLDFT model. The results show the presence of two types of pores: micropores ( 0.8 nm) and mesopores ( 8 nm) (Figure 4b). In contrast with the SSA and porous volumes, the pore size distribution before and after Sn/SnO2 impregnation remains constant no matter the Sn loading. The presence of both types of pore are important, since they allow the diffusion of the electrolyte through the material and particularly the larger pores may accommodate the Sn particles and avoid their expansion/agglomeration during cycling.

Therefore, the incorporation of the Sn/SnO2 nanoparticles induces a decrease in the surface area and porous volume. This may be explained mainly by two factors: i) higher density of metallic Sn-based particles than carbon, ii) confinement of the particles in the carbon pores (pore blockage with the particles). To discriminate between these two effects, the textural values were calculated per gram of carbon based on Sn and SnO2 contents determined in Table 2. For low loadings ranging from 20 to 40%, the SSA and porous volume calculated per gram of carbon are similar to that of pristine carbon (C), while for high loadings i.e., 60 to 80 %, surprisingly these values gradually increases (Figure 4d). Two hypotheses may be proposed for this behavior: modification of carbon during the Sn 4+ infiltration/reduction/drying steps or the existence of some porosity induced by Sn aggregates.

To check the first hypothesis, the composite was washed with HCl to remove the Sn/SnO2 particles followed by several analytical characterizations to confirm the efficiency of such treatment (Figure S4, Supporting Information). XRD pattern of the material after washing (60%Sn/C@HCl) reveals the absence of Sn peaks and the presence of specific phase of carbon with low degree of graphitization. In addition, the TGA analysis in air indicates as well the complete removal of Sn-based nanoparticles since no SnO2 residue was observed after thermal treatment, the curve of 60%Sn/C@HCl is overlapped to that of carbon one. The N2 adsorption/desorption isotherms, present similar profiles to the pristine carbon support, suggesting that the carbon texture was not either modified by the synthesis procedure. Therefore, all these characterization point out that the carbon matrix texture, structure and thermal stability is not affected by the Sn nanoparticles synthesis procedure. However, for other metallic systems confined in microporous carbon, the modification of carbon porosity (specific surface area and pore size) could be observed due to the synthesis procedure [29] .

The second hypothesis assumes some porosity coming from the small size of Sn particles or between the Sn-particles. The Sn/SnO2 can be considered non-porous, with no intrinsic porosity. However, due to their nanometric size, the SSA may increase and an inter-particle porosity cannot be excluded as well. This porosity increases due to the fact that, by increasing the particle loading, more Sn particles are formed outside the carbon network and they may contribute to the SSA increase. In the same time, a part of particles can be confined in the pores of carbon leading to lower SSA and pore volume.

Electrochemical performances

From the physico-chemical analyses on the C/Sn-based materials, we observed that the modification of the Sn loading produces changes regarding the size and dispersion of the Snbased particles, their tendency to form agglomerates by increasing the Sn loading being highlighted. The Sn nanoparticles were found to be partially oxidized, and such oxide may play a role on the irreversible capacity. The textural properties such as SSA and porous volumes decreased with increasing of the Sn loading. All these characteristics may influence on the Li-ion batteries performances will be discussed in the following part.

The electrochemical tests for carbon and carbon/Sn-based materials were performed at constant (C/5) and variable current rate (between C/10 and 3C) in the voltage range of 0.01-1.5 V (vs Li/Li + ). Figure 5 presents the galvanostatic and derivative curves for 60%Sn@C for the specified cycles. The first discharge profile show a rapid voltage drop to 1.5 V, followed by a gradually decrease down to 0.01 V, delivering a discharge capacity of 1480 mAh g -1 . The subsequent discharge curves are very different compared to the first one because of the formation of SEI on the surface of the anode during the first discharge. In the same time, the following curves are almost overlapped, indicating the good reversibility of the chemical reactions in the electrode and a good cycling stability. The initial charge capacity is 646 mAh g -1 , corresponding to a coulombic efficiency of 89%.

The derivative curves (Figure 5b) for the first discharge present three peaks badly defined at 0.63, 0.0.50 and 0.38 V, the first one being shifted to 0.42 starting for the second cycle. During charge three better defined derivative peaks are at 0.60, 0.72 and 0.80 V, in good agreement with those assigned to the reversible lithiation of Sn into Li Li2Sn5, LiSn and Li2.5Sn, respectively (Eq. 2) [2] .

Material

BET SSA m 2 g -1 VT cm 3 g -1 Vmicro cm 3 g -1 Vmeso cm Note that no clear peak appears around 1 V in discharge which would have been ascribed to the conversion reaction (Eq. 1) [33] , [2] and this may be due to the nanometric size of Sn particles.

After 50 cycles the peaks are no more visible.

SnO2 + 4Li + + 4e -→ Sn + 2Li2O (Eq. 1)

Sn + xLi + + xe -↔ LixSn (Eq. 2)
To evidence the contribution of the carbon support on the electrochemical results of C/Sn hybrid materials, the carbon support, was tested in the same conditions as the composite materials (Figure 6). The reversible capacity of carbon support after 100 cycles was 204 mAh g -1 , but a high irreversible capacity of 776 mAh g -1 was observed during the first chargedischarge cycle, representing 78% from the first charge capacity (Table S1a, Supporting Informations). The coulombic efficiency is quite low in the first cycle (≈16%), but from the second one, it strongly increases to 80% and after 25 cycling reaching 100% (Figure 6a). The high irreversible capacity and the low columbic efficiency in the first cycles are related to the high surface area of carbon materials which induces the electrolyte decomposition with the formation of a solid electrolyte interphase.

The electrochemical performances of the hybrid materials were evaluated at constant and variable current rate and compared with their theoretical capacities (Figure 6). The theoretical capacity for each material was calculated by taking into account the real amount of each compound (Sn/SnO2/C) (Table 2), with the theoretical capacities for Sn and SnO2 (990 mAh g -1 [START_REF] Wang | [END_REF] and 781 mAh g -1 [START_REF] Courtel | Nanotechnology for Lithium-Ion Batteries[END_REF] , respectively) and the experimental reversible capacity obtained for the carbon which was used as support.

The irreversible capacity of C@Sn composite materials is even higher than the carbon one, suggesting a contribution of Sn-based nanoparticles to the irreversible capacity (Figure 6b). In order to determine the impact of Sn-particles on the irreversible capacity, we calculated its contribution by taking into account the experimental irreversible capacities obtained for the carbons support.

As can be noticed a decrease of the irreversible capacity with the decrease of the carbon content in the C@Sn composites is clearly observed. The 20%Sn@C presents an irreversible capacity of 1323 mAh g -1 , from which 667 mAh g -1 due to the carbon while by increasing the Sn loading to 80%, the irreversible capacity was reduced to 653 mAh g -1 (43%). These values are lower than those of Sn powder, where an irreversible capacity around 1350 mAh g -1 representing 93% from the initial capacity was reported in the literature [34] . The significant decrease of the irreversible capacity for 80%Sn/C compare to that of Sn powder (93%) and, in the same time, to the irreversible capacity of the carbon support (78%) indicates the importance of carbon matrix. This decrease tendency of irreversible capacity may be related to the decrease of the specific surface area of the C@Sn materials, as demonstrated by the rather linear correlation of irreversible capacity with the specific surface area (Figure 6c). The presence of a smaller fraction of porous carbon limits the electrolyte decomposition and the solid electrolyte interphase (SEI) formation. Despite the strong decrease of irreversible capacity by decreasing the carbon fraction and the specific surface area, an important irreversible capacity is still coming from the Sncontribution. This is often assigned to the electrolyte consumption due to the conversion reaction of SnO2 into metallic Sn and Li2O (Eq. 1). If we represent the evolution of the irreversible capacity versus the oxygen amount in the Sn/SnO2 particles (Figure 6d), one can be noticed that the irreversible capacity decreases with the increase of the oxygen amount in the materials, therefore, with the increase of the quantity of SnO2. At a first glance, this result seems rather unexpected since a higher amount of SnO2 in the composites is supposed to induce a higher irreversible capacity due to the reaction (1). However, some authors found also that the presence of SnO2 in Sn/SnO2 materials is beneficial to limit the SEI formation, to improve the CE and also to provide long term cycling stability [35] , [36] , [37] . Therefore, the conversion reaction is not the only factor which may affect the irreversible capacity of Sn/SnO2 materials. The electrolyte decomposition on the surface of small and reactive Sn/SnO2 nanoparticles, with the formation of SEI layer or the modification of the particle size/morphology during Li + insertion may account as well for the observed irreversible capacity [38] , [39] , [35] .

Concerning the reversible capacity (Figure 6e), this was also significantly improved by increasing the Sn loading from 397 mAh g -1 for 20 wt.% Sn to 871 mAh g -1 for 80 wt.% Sn, respectively. The values are even higher than the theoretical capacities for these materials (Table S1b, Supporting Information) due probably to some synergetic effects between the Sn nanoparticles and the carbon matrix. For high loadings (60, 70 and 80 wt.%) after a drastically decrease of the capacity during the first 20 cycles (≈ 40%), stable capacities of 524, 430, 326 mAh g -1 is maintained up to 80 th cycle, respectively. For lower Sn loadings, after the first cycles, the capacity is directly extremely low (~300 mAh g -1 ), but not so far from the theoretical values (285 and 340 for 20 wt.% Sn and 40% Sn respectively). Another decrease of the capacity occurs after 100 cycles for the higher loadings, which finally present almost the same capacity (≈ 300 mAh g -1 ) after 350 cycles (Figure 6e). The fading of capacity for the higher loadings could be a consequence of the large amount of agglomerated particles situated outside of the carbon pores, confirmed by STEM (Figure 2). These particles were not properly in contact with the conductive carbon network and non-confined as well in corbon pores to limit their volume expansion/strain during lithiation/delithiation. STEM analysis after cycling were performed for materials with 60 and 80 wt.% of Sn to see how the Sn-based particles were affected during the cycling (Figure 7). Compared to the materials before cycling (Figure 2b,c), the STEM images at this stage indicate that the 60%Sn@C composite better conserves the particles size and dispersion (Figure 7a), while for 80%Sn@C material the agglomerated particles are more present than before cycling (Figure 7b). These post-mortem investigations confirm that a moderate amount of Sn in the porous carbon frameworks is required to balance the particle aggregation, the material conductivity and electrolyte diffusion.

The evolution of the capacity for long cycling could indicate a similar maximal electrochemical potential no matter the Sn NPs loading. Therefore, to choose the most promising material which could be used as anode, we should consider the irreversible capacity, as well as the reversible capacity but not only during the first cycles which presents the highest value for the material with 80 wt.% of Sn, but also the long-term cycling stability behavior, which seems to be especially, by the dispersion of the particles in the carbon matrix.

The coulombic efficiency of the C/Sn composite materials reaches fast 100% (after 10 cycles), which means that no parasitic reactions take place because of the electrolyte degradation is disappeared after first cycles. Figure 6f shows the cycling performances of the C/Sn-based materials at different current rates. The reversible discharge capacities of the composites strongly depend on the C rate and thus, for Sn/C composite the values decrease to 578, 466, 314, 249 and 210 mAh g -1 at C/10, C/5, C, 2C and 3C, respectively, and successfully regain up to 90% from the initial capacity, when the current rate is set back to its initial value.

Conclusions

Mesoporous carbon/Sn-based materials with different Sn loading were prepared by a facile and fast route based on an impregnation/reduction bottom-up approach. For all loadings, a mixture of Sn/SnO2 was formed and several complementary techniques (Mössbauer spectroscopy, XPS and TGA) were employed to carefully asses their amount and ratio in the carbon matrix to correctly evaluate their impact on the electrochemical performances. At low Sn loadings (< 60%) the Sn NPs are very homogenous dispersed in the carbon network with particle size around 5 nm and only few aggregates observed. The increase of the Sn loading induces the formation of larger Sn particles and many Sn agglomerates (~ 200 nm), which are not embedded in the carbon porosity, as reveled by XRD and TEM. Moreover, the increase of the Sn loading in the carbon framework induces a significant decrease in the textural properties (SSA and pore volume) which could be assigned partly to the Sn density but also to the confinement of Sn particles into the carbon pores. It was evidenced also that the Sn nanoparticles may present a non negligible inter-grain porosity.

The increase of the Sn loading leads to a decrease of the irreversible capacity during the first cycle, which was mainly related to the decrease of the SSA. Improvement of the reversible capacity is achieved as well; however, the evolution of the reversible capacity for long cycling demonstrated significant fading for high loadings triggered by Sn NPs placed outside the carbon pores which undergo severe agglomeration due to volume expansion upon lithiation. Therefore, similar electrochemical performances are reached regardless the Sn NPs loading after 350 cycles. Our founding suggests that a high Sn loading combined to small and dispersed particle size in the carbon framework may be the key factor to achieve high reversible capacity and long term cycling.

Experimental Section

Material synthesis

The mesoporous carbon support was prepared by a softtemplate phase separation route as previously reported in our work [21] , [3] , involving the dissolution of phloroglucinol (the carbon source, 1.65g) and Pluronic F-127 (pore generation agent, 3.27 g) in a solution of ethanol (81 mL) containing small quantity of hydrochloric acid (0.6 mL), followed by the glyoxal addition (1.62 mL). The latter one is used as cross-linker to form a tridimensional phenolic resin with phloroglucinol able to selfassemble with the triblockpolymer template Pluronic F-127 via H-bonding giving rise to mesophase nano-assemblies. Thermal annealing at 900°C of organic polymer phase allows to obtain the mesoporous carbon support.

The C/Sn hybrid materials with different Sn loadings (x%Sn@C, where x=20/40/60/70/80) were prepared by liquid impregnation of the porous carbon network (C) with a SnCl4•5H2O ethanol solution under stirring for 30 min, followed by a chemical reduction process with NaBH4 solution in water in order to reduce the SnCl4•5H2O to Sn (0.6g NaBH4/1g SnCl4•5H2O). The reducing agent was eliminated by washing the material with distilled water at room temperature. The final material was dried at 80°C for 12h.

Materials characterization

The crystalline structure of the materials was characterized by X-ray powder diffraction (XRD) technique, by using a powder diffractometer D8 ADVANCE A25 from Bruker in Bragg-Brentano reflection geometry θ -θ. This diffractometer is equipped with the LynxEye XE-T high resolution energy dispersive 1-D detector (Cu Kα1,2), leading to ultra-fast X-ray diffraction measurements. The X-ray powder diffraction data uses proprietary Bruker format which are converted into usable one via PowDLL 2.71 [26] . The Sn crystallite size was determined using the Scherrer formula and the FWHM (Full Width at Half Maximum) was extracted via DIFFRAC.EVA [27] (CuKα2 contribution previously removed) and also checked via WinPLOTR profile fitting procedure [START_REF] Roisnel | Materials Science Forum[END_REF] . The material surface morphology/structure was investigated with a JEOL ARM-200F transmission electron microscope working at 200 kV. EDX mapping was obtained with a JED 2300 detector. Thermogravimetric analyses (TGA) were conducted on a TGA 851 (Mettler-Toledo) thermogravimeter by heating the samples under air (100 ml/min), from 30 o C up to 900 o C with a 10 o C/min heating rate.

119 Sn Mössbauer spectra were collected with a constant acceleration spectrometer using a Ca 119m SnO3 source in the transmission geometry. They were fit by combinations of Lorentzian lines. The absorber used for the determination of the hyperfine parameters contained 120 mg cm 2 of the Sn composite. The velocity scale is calibrated with a 57 Co source and a pure iron foil absorber. The isomer shifts are given relative to BaSnO3 standard at room temperature. The absorption areas were further used to calculate the quantities of Sn and SnO2 loaded in the carbon/Sn hybrid materials. Lamb-Mössbauer factors of 0.5 and 0.05 have been assumed for the conversion of absorption areas for SnO2 and Sn, respectively [32] . X-ray photoelectron spectroscopy (XPS) was performed with a VG Scienta SES 200-2 spectrometer equipped with a monochromatized Al Ka X-ray source (1486.6 eV) and a hemispherical analyzer. The pass energy was 100 eV.

The textural properties of the carbon and carbon/Sn hybrid materials were evaluated using N2 adsorption/ desorption isotherms measured with a Micromeritics ASAP 2420 device at 77K. The samples were out-gassed in vacuum at 80 o C for 12 h before the adsorption analysis. The specific surface area (SSA) was calculated from the linear plot in the relative pressure range of 0.01-0.05. The micropore volume (Vmicro, pore size < 2 nm) was obtained using the Dubinin-Radushkevich (DR) equation in the relative pressure interval (P/P0) from 10 -4 to 10 -2 . The total pore volume (VT) was determined from the amount of nitrogen adsorbed at a relative pressure of 0.95. The mesoporous volume (Vmeso, 2nm < pore size <50 nm) was determined by subtracting the microporous volume form the total pore volume. The pore size distributions (PSD) were determined from the adsorption branch of nitrogen isotherms using the 2D NLDFT heterogeneous surface model for carbon materials implemented in SAIEUS (Micromeritics) [START_REF] Jagiello | [END_REF] .

Electrochemical characterization

The electrode preparation was made by mixing in a jar the synthesized materials (70%) with carbon black (10%) and vapor grown carbon fibers -VGCF-S (10%) as conductive additive, and carboxymethylcellulose -CMC (10%) as binder. A quantity of 0.4 ml distillated water was added over the materials and homogeneously mixed by ball-milling for 1h. The film with a thickness of 100µm was laid down on a copper foil, dried at room temperature for 4h and at 100 o C under vacuum for other 12h. The electrochemical results were calculated based on the active mass, which was considered the mass of the carbon/Sn/SnO2 composite. The corresponding loading of the active mass in the electrode was comprised between 0.9 and 1.9 mg/cm 2 .

The electrochemical performances of carbon and carbon/Snbased materials as negative electrode for Li-ion batteries were studied in coin cells assembled in an argon filled glove box. The electrolyte used for the coin cell preparation was LiPF6 with ethylene carbonate (EC), propylene carbonate (PC), dimethyl carbonate (DMC) (1:1:3) as solvents and vinyl carbonate (VC)(1%) and fluoro ethylene carbonate (FEC)(5%) as additives. The galvanostatic charge and discharge measurements were carried out at room temperature on Neware battery testing system (BTS3000) at room temperature. The capacity is reported per gram of C/Sn composite.

  Nita, Dr. J-M Le Meins, Dr. J Parmentier, Dr. C.

Figure 1 .

 1 Figure 1. XRD patterns for C/Sn hybrid materials with different Sn loading (x% Sn@C, x=20/40/60/70/80).

Figure 2 .

 2 Figure 2. STEM images for carbon/Sn-based hybrid materials having different loadings corresponding to 20 wt.% Sn (a), 60 wt.% Sn (b); 80 wt.% Sn (c).

Figure 3 .

 3 Figure 3. Room temperature Mössbauer spectra for carbon/Sn-based materials having different Sn loadings x%Sn@C (x=20/40/60/70/80).

Table 4 .

 4 Textural values of carbon support and x% Sn@C materials (x=20/40/60/70/80).

Figure 4 .

 4 Figure 4. Textural properties for carbon and carbon/hybrid materials. N2 adsorption/desorption isotherms (77K) (a) and pore size distribution (b). The evolution of the porous volume (c) and SSA (d) with the Sn loading.

Figure 5 .

 5 Figure 5. Galvanostatic (a) and derivative curves (b) for carbon/Sn composites 60%Sn@C at C/5 current rate in the voltage range of 0.01-1.5 V (vs. Li/Li).

Figure 6 .

 6 Figure 6. Specific discharge capacity and coulombic efficiency for the carbon and carbon/Sn materials as a function of cycle number at constant current rates C/5 C rate in the voltage range of 0.01-1.5 V (vs. Li/Li + ) (a). The irreversible capacity after the first discharge (b). The correlation between the irreversible capacity and BET specific surface area (c) and oxygen amount in Sn/SnO2 nanoparticles as determined by XPS (d); The reversible capacity after 2 cycles, 40 cycles and 350 cycles, respectively (e). Specific discharge capacity as a function of cycle number at variable current rates between C/10 and 3C (f).

Figure 7 .

 7 Figure 7. STEM images after 350 cycles for 60%Sn@C (a) and 80%Sn@C (b).

Table 1 .

 1 Carbon/Sn-based materials for lithium-ion battery.

	Carbonaceous
	/Sn-

based material Preparation method of composite Electrochemical performances vs. Li Refs Curren t density (mAg -1 ) Rever sible capac ity (mAh g -1 ) Cycle s numb er

							Sn/carbon	Acid-etching	30	413	100	[15]
							nanotubes	treatment
							composite
							Sn NPs/Carbon	Direct current	100	600	10	[14]
							nanotubes	(DC) arc-
							nanocapsules	discharge
								method
							SnO2	Solvothermal	200	497	300	[23]
							nanocrystals/	method
							carbon
							nanotubes
							SnO2/carbon	Solvothermal	100	709	100	[7]
							nanotubes	method
							SnO2/Carbon	Hydrothermal	100	400	50	[18]
							nanotubes	approach
							Sn/N-doped	Arc-discharge	500	520	1000	[20]
							carbon	method
							nanospherers
							Sn NPs/	Electrospinnin	2C	350	600	[19]
							carbon	g,
							microtubes	carbonization
							Sn particles/	Impregnation/	0.1 C	740	160	[9]
							Carbon fibers	heat treatment
							Sn/graphite	Electrocatalysi	C/15	433	35	[12]
								s
							Sn NPs	Hydrothermal	100	662	100	[11]
							/carbon/	process
							graphene
							Sn/carbon	Sol-gel	20	588	200	[5]
							nanocomposite	carbothermal
								reduction
							Sn/porous		100	788	300	[24]
							carbon
							Sn/	Sol-gel	20	400	15	[13]
	Onion-like	Arc-discharge	100	585	100	[17]	mesoporous	polymerization
	carbon NCs/Sn	method					carbon
	Sn NPs	Arc-discharge	100	902	1000	[16]	ous carbon SnO2/Mesopor	Treatments Heat Impregnation	1 C	511	500	[3]
	nanocapsules/	method							200	367	2000
	N-doped onion-
	like carbon
	Sn	Chemical	100	420-	30	[8]
	nanowires/Car	vapor		630
	bon nanotubes	deposition
		(CVD)
	Sn/Carbon	Hydrothermal	100	400	50	[22]
	core-shell	method, Heat
		treatment
	Sn	Solvothermal	0.2 C	1070	200	[10]
	nanoparticles/N	method
	-doped hollow		5 C	500	500
	nanosphere

Table 2 .

 2 Crystallite size determined by XRD and quantification of Sn and SnO2 amounts contained in the Sn/C materials, based on Mössbauer spectroscopy and TGA analysis.

Table 3 .

 3 The chemical composition of the carbon/Sn-based materials obtained by XPS analysis.

	Material	Composition, wt %	Repartition of oxygen, wt%
		C	Sn	O	O-(C)	O-(Sn)	O-(Sn)/Sn
	20%Sn@C	61.8	26.7	11.5	4.7	6.8	1.90
	40%Sn@C
		45.3	41.2	13.5	3.1	10.4	1.88
	60%Sn@C
		11.8	69.9	18.2	2.6	15.6	1.66
	80%Sn@C
		8.8	72.1	19.2	2.7	16.5	1.69

Acknowledgements

The authors gratefully acknowledge the financial support of this work from University of Haute-Alsace, France through the Ph.D. scholarship of Cristina Nita. We thank Loic Vidal and Samar Hajjar for STEM and XPS analyses performed via the IS2M Mulhouse technical platform.