Tin–carbon composites as anodic material in Li-ion batteries obtained by copyrolysis of petroleum vacuum residue and SnO2

Carbon 45 (2007) 1396–1409 www.elsevier.com/locate/carbon

Tin–carbon composites as anodic material in Li-ion batteries obtained by copyrolysis of petroleum vacuum residue and SnO2 J.L. Tirado a, R. Santamarı´a b, G.F. Ortiz a, R. Mene´ndez b, P. Lavela a, J.M. Jime´nez-Mateos c,*, F.J. Go´mez Garcı´a c, A. Concheso b, R. Alca´ntara

a

a

Lab. Q. Inorga´nica, Ed. C3 1a, Campus de Rabanales, UCO, 14071 Co´rdoba, Spain INCAR (CSIC), Francisco Pintado Fe 26, La Corredoria, 33011 Asturias, Spain Centro Tecnolo´gico REPSOL YPF, Ctra. A-5 km 18, 28931 Mo´stoles, Madrid, Spain b

c

Received 2 November 2006; accepted 19 March 2007 Available online 2 April 2007

Abstract Composite materials with tin nanoparticles surrounded by a ‘‘muffling’’ carbon matrix are formed simultaneously by adding 20% SnO2 to a vacuum residue and following carbonisation between 700 C and 1000 C. The primary purpose of the carbonaceous material is the reduction of SnO2, giving rise to SnS and Sn as nanoparticles. The homogenous distribution of both components induces therefore a synergetic effect on the properties of the electrode material, not only from the electrochemical point of view but also from that mechanical. Thus, the carbon matrix hinders the agglomeration of Li–Sn alloys during long term cycling and, simultaneously, tin particles improve the conductivity of the material and increase the overall capacity as compared with the reference carbon. In addition, a CVD treatment increases the performance of the material. 119Sn Mo¨ssbauer and 7Li MAS NMR spectroscopies allow a detailed study of partially charged/discharged samples and, therefore, the phases, steps and mechanisms occurring during the electrochemical process.  2007 Elsevier Ltd. All rights reserved.

1. Introduction Carbonaceous materials are extensively used as commercial anode materials in Li-ion batteries, due to their suitable working potential and structural stability upon cycling [1]. However, some drawbacks arising from production costs and/or technical problems, including electrolyte reactivity and structural stability, are still there [2,3]. Besides, in the case of graphite, the capacity is limited to that of the stoichiometric phase LiC6, that is, 372 mA h/g [4]. On the other hand, disordered low temperature (LT) carbons (<1000 C), although combining advantageous production costs with capacities even higher than the theoretical value for graphite [5,6], show high undesirable irreversible capacity in the first discharge and high polarization

*

Corresponding author. Fax: +34 91 348 86 13. E-mail address: [email protected] (J.M. Jime´nez-Mateos).

0008-6223/$ - see front matter  2007 Elsevier Ltd. All rights reserved. doi:10.1016/j.carbon.2007.03.041

in the charge process, which are limiting factors for their practical use in commercial batteries. Alternative materials, as certain metals, could be indeed a better election for anodes in Li-ion batteries, but they are not exempt from drawbacks as well. For instance, the case of the proper lithium is very well known [7], and other lithium alloying metals suffer different problems during the complete cycle of charge–discharge. Thus, tin, which can accept lithium up to Li22Sn5, i.e., 992 mA h/g [7], would be apparently a very attractive candidate for such application, as its capacity is almost three times that of graphite, and higher indeed than that shown for some of the best LT carbon materials in the first cycle. However, this promising behaviour is not free of its ‘‘dark’’ side: the incorporation of Li to the alloy imply a huge increase of volume (up to 258%, [7]), giving rise to, on the one hand, to a hard mechanical stress in the system and, on the other hand, to a breaking up of the tin particles and, consequently, to electrical and electrochemical isolation. In short, the retention

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

of capacity, in spite of its promising theoretical expectative, is very poor, much less indeed than that of many LT carbon materials. A possible solution to avoid these drawbacks is to prepare tin in the form of nanoparticles effectively dispersed by a ‘‘muffing’’ matrix that may cushion the mechanical effects of the volume changes. Obviously, this matrix should be reasonably electrically conductive, ionically permeable, at least for Li+, and chemically compatible with tin, the Li–Sn alloys and the electrochemical system as a whole. A carbonaceous material can be the election. In this way, several routes have been proposed. For instance: deposition of tin nanoparticles, arising from the reduction of SnCl2 by NaH, on graphite, stabilized by t-BuOH [8], nanocapsules of amorphous carbon coated tin prepared by reduction of SnCl4 with NaBH4 and further hydrothermal treatment of the dispersion of such nanoparticules on a glucose solution [9], coprecipitation in a suspension of functionalized nanotubes of SnO2, which should be reduced afterwards [10], or infiltration of nanoSnO2 in the nanotubes [11], confinement of nanoparticles of SnCl4 (which is further hydrolyzed and reduced) in micropores of activated carbons [12], copyrolysis of polystyrene and tetrafeniltin previously milled together [13], or covering of hydrophobized Sn nanoparticles dispersed in water by a shell of resorcinol and formaldehyde, further polymerized and calcined to give rise a powder of a sort of ‘‘nanotin-nuts’’ [14]. In all these methods the Sn nanoparticles are formed apart from the matrix, in different steps. Indeed, in a recent work in which the composite material is prepared in an ‘‘one pot’’ route by reduction of SnCl4 by (t-BuONa)-activated NaH suspension in THF on a graphite [15], the carbonaceous supporting material (graphite) is already formed, previous to the Sn(IV) reduction, therefore giving rise to independent tin aggregates and deposition of Snnanoparticles on the graphite surface. However, in order to get a more intimate mixed composite material, it would be very convenient that the carbonaceous matrix and the tin nanoparticles were formed at the same time. Thus, if the Sn and carbon precursors are selected properly, both processes could be carried out simultaneously. In this sense, we have studied the capability of SnO2–petroleum vacuum residue system, according to our previous experience in similar systems [16]. 2. Experimental To obtain of the Sn–C composites, commercial SnO2 (20% w/w) was copyrolized (420 C, 8 bar) with a heavy petroleum vacuum residue in a stirring reactor for 135 min. The high pressure was needed in order to assure a high carbonisation degree, with a yield of 52%. At this stage, no metallic tin is formed, according to Ellingham’s diagrams, although some tin is partially reduced to SnS, as can be checked by X-ray diffraction. In Fig. 1, the image of a polished section, showing the optical texture of this material, is included, obtained by an polarized-light optical microscope Zeiss Axioplan equipped with a 1 k retarder plate. From this image it can be directly deduced that the tin compounds (brightest spots) is perfectly dispersed within the carbon matrix.

1397

Fig. 1. Optical microscopic image of the polished copyrolysed SnO2-VR sample, using polarized light. Note the high dispersion degree of the tin nanoparticles (bright spots).

Afterward, the tin-containing coke was carbonised for 6 h at different temperatures: 700 (VR700), 800 (VR800), 900 (VR900) and 1000 C (VR1000) in a dynamic (3.7 l/h) nitrogen atmosphere. For the sake of comparison, another coke was also prepared under the same conditions, but without addition of metal oxide, and carbonised at 750 C (VR0). In order to check the effect of the surface properties in the electrochemical behaviour of the material, the VR800 sample was subjected to carbon vapour deposition (CVD): the VR800 material was heated at 2 C/min up to 800 C under a dynamic atmosphere of nitrogen (3.7 l/h); when the setting temperature was reached, the atmosphere was changed to toluene carried by a stream of nitrogen of 80 l/h.; after 6 h of treatment, the atmosphere was again changed to nitrogen at 3.7 l/h and let cool down free. This sample was labelled as VR800-CVD. The chemical analysis were carried out in a LECO-CHNS-932 analyzer, which has the auxiliary graphite furnace VTF-900 for oxygen analysis, following the usual procedures. Scanning electron microscopy (SEM) observation was performed in a Philips-FEI XL30-CP microscope, equipped with a conventional SE (secondary electrons), a solid state (two sectors) polar BSE (backscattered electrons) and a X-ray EDAX Sapphire (ultrathin Be window) detectors, on crushed and polished sections of the samples. A detailed description of the experimental procedure for microanalysis is included in [16–18]. X-ray diffraction (XRD) patterns were recorded in a Siemens D-5000 apparatus provided with CuKa radiation and using 0.04 2h/1.2 s steps. The 119Sn Mo¨ssbauer spectra (MS) were recorded at room temperature with a mixed Ametek–Wissel constant-acceleration spectrometer in transmission mode. The source was b-119mSn in a matrix of BaSnO3, working at room temperature. The velocity scale was calibrated using the magnetic sextet of a high purity 57Fe foil absorber, with a source of 57Co–Rh. The isomer shift scale is zeroed respect to the centre of the peak of the Sn(IV) signal in the BaSnO3 at room temperature. Experimental data were fitted to Lorentzian lines using a least-square-based method.1 The goodness of the fit was controlled by the classical test of v2. 7 Li MAS-NMR spectra were recorded using an Avance 400WD solidstate spectrometer (155.4 MHz resonance frequency for 7Li) with magic angle spinning (MAS, 14 kHz). LiCl was taken as an external standard for line shifts measuring (0 ppm). The discharged electrodes were rinsed with propylene carbonate, dried in vacuum and put into a hermetic NMR holder.

1 Landry F, Schaaf P. WinISO – Windows Mo¨ssbauer Fitting 625 Programme, 1998. Private communication.

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

The electrochemical behaviour was studied by using two-electrode Swagelok type cells. Counter electrodes were 9 mm discs of lithium metal. The working electrode consisted of a mixture of 92% active material and 8% PVDF binder, supported on a cooper foil of the same diameter. A 1 M LiPF6 (EC:DEC = 1:1) electrolyte solution was supported in Whatman glass fibber discs. Both the assembly of test cells and the handling of discharged/charged electrodes were carried out in a Mbraun glove box under an argon atmosphere. Cycling experiments were carried out in a MacPile system at 10 mV/0.1 h scan rate, while electrodes partially discharged and charged for post-mortem studies by 119Sn Mo¨ssbauer and 7Li NMR spectroscopies were prepared at 10 mV/0.3 h. Electrochemical impedance spectra were recorded in an Autolab PGSTAT12 system. For this purpose, three-electrode lithium cells assembled with a coke sample as working electrode and lithium pellets as counter and reference electrodes were cycled. The test cell was allowed to relax in open circuit for at least 5 h to achieve a quasi-equilibrium system. An AC voltage signal of 5 mV was applied from 100 kHz to 2 mHz.

3. Results and discussion

due to the sulphur organic compounds included in the vacuum residue [16]. SnO2, when detectable, shows very weak peaks. The Mo¨ssbauer spectra have been deconvoluted in three components, a singlet –Sn(0)– and two doublets –Sn(II) and Sn(IV)– [19] (see Fig. 3). The higher the temperature, systematically the lower the concentration of Sn(IV) is, which is reasonable as the reduction of Sn(IV), in instance by carbon, is favoured at higher temperatures. However, 1.005 1 0.995 Absorption

1398

0.99 0.985 0.98

700 °C 6 h

0.975

3.1. Compositional studies: XRD and MS

0.97 0.965

-2

0 Velocity (mm/s)

2

4

6

-2

0 Velocity (mm/s)

2

4

6

-2

0 2 Velocity (mm/s)

4

6

-2

0 2 Velocity (mm/s)

4

6

1

0.996 0.994 0.992 0.99

800 °C 6 h

0.988 0.986

Table 1 Chemical composition of carbonised samples

-6

-4

1.005

Elemental analysis (%)

C/H

C

H

N

S

O

54.21 56.70 56.30 57.07 90.65

0.73 0.43 0.23 0.10 1.18

0.83 0.71 0.60 0.49 1.45

5.01 4.27 4.87 4.68 5.08

1.54 0.94 0.21 0.12 1.64

6 11 20 48 6

C/O

47 80 357 634 55

1

C/S

29 35 31 33 18

Absorption

VR700 VR800 VR900 VR1000 VR0

-4

0.998

0.984

Sample

-6

1.002

Absorption

The chemical characteristics of these carbonised composites are shown in Table 1, showing the evolution of oxygen and hydrogen with the carbonisation temperature. The XRD pattern of the carbonised coke VR0 is that typical for a low temperature carbon material, with just characteristic highly broadened (0 0 l) and (h k 0) bands. On the other hand, peaks due to tin-containing phases can be observed in the X-ray diffractograms of the composite materials (Fig. 2), as metallic tin (b-Sn, JCPDS 4-673) and partially reduced sulphides (SnS, herzenbergita, JCPDS 39-354),

0.995 0.99 0.985

900 °C 6 h

0.98 0.975 0.97 -6

1200

I (a.u.)

1000

VR12 VR0

700ºC-6h

800ºC-6h

900ºC-6h

1000ºC-6h

Sn (04-0673)

SnO2 (41-1445)

SnS (39-0354)

1.005 1 Absorption

1400

800

0.995 0.99

600

0.985

400

0.98

200

0.975

0 10

20

30

40

50 °2 theta

60

70

80

-4

90

Fig. 2. X-ray diffractograms of the pyrolysed and carbonised composite samples. The bar diagrams of the identified phases (b-Sn, SnO2 and SnS) are shown, indicating the corresponding JCPDS files.

1000 °C 6 h

-6

-4

Fig. 3. Mo¨ssbauer spectra of the studied samples, including the deconvolution and fitting results. Sn(0) is a singlet (green), meanwhile Sn(IV) and Sn(II) are doublets (blue and red, respectively). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

that is not the case of Sn(II) and Sn(0), whose concentrations vary not linearly (see Fig. 4). In order to explain these results, rough thermodynamical calculations for the possible main reactions related to tin evolution during carbonisation have been calculated (Table 2). According to these data, the most-favoured compound is SnS, conversely the formation of SnO is rather unlikely. Therefore, it can be propose that during coking at 420 C only a partial reduction of SnO2 to SnS takes place, as tin cannot be reduced by carbon at so low temperature. When that material is carbonised at higher temperatures, the preferred reaction is still the partial reduction to SnS, but as the matrix is carbon, this is most available, and due to simple kinetic effects the reduction to metallic tin by carbon occurs significantly; however, the higher the carbonisation temperature, the higher the diffusion of fluid phases (Sn, H2S, CS2, etc.) is, thus allowing indeed that not only the partial reduction would be increased but part of the metallic tin formed would be re-oxidised to SnS. On the other hand, SnS cannot be reduced to Sn directly only by the action of carbon (Table 2), unless hydrogen or oxygen were involved or the temperature was much higher. 70

60

concentration (%)

50

40

30 Sn(0)

20

Sn(II) Sn(IV)

10

0 700

800

900

1000

carbonisation temperature (°C)

Fig. 4. Evolution of the concentration of the oxidation states of the tin versus the carbonisation temperature for the composite materials, according to the integration of the respective Mo¨ssbauer spectra.

1399

3.2. Texture: microscopic observation As expected, the optical texture of the pyrolized and carbonised sample without tin is the typical mixture of flow domains and small mosaic structure [16], consequence of the two reaction pathways (reactivities) occurring typically during the coking of such feed [21,22]. However, the introduction of SnO2 changes completely this picture, avoiding the formation of flow domains, giving rise to a regular mosaic with thin and uniform texture and very small particles containing the metal compounds embedded in the carbonaceous matrix (see Fig. 1). The morphology of all the samples, crushed or polished, is quite similar when observed by SEM using BS, and, therefore, only those for VR800 are shown (Fig. 5). In general, the crushed samples (Fig. 5 top) are formed by irregular faceted bits of 50–100 lm of length, covered by a thin and loose powder of submicron particles, probably originated from small fragments detached during grinding. In addition, some shiny spheres of metallic tin (according to EDS spot analysis) can be observed as well. On the other hand, the matrix is not homogeneous, but at least three brightness levels can be observed. Considering that the contrast in BSE–SEM is completely depending on the average atomic weight in flat surfaces, these areas must be ascribed to different compositions: the darker one being that corresponding to the carbonaceous matrix, the brighter being a phase with a high concentration of tin (maybe embedded metallic tin), and the intermediate one probably being a lighter phase of tin (sulphide). The size of these moieties is around one micron. The BSE images of the polished samples give a clearer picture on the composition of these materials. Up to four different grey levels can be identified in all the samples (Fig. 5b). The brightest particles (‘‘d’’ in Fig. 5b) are metallic tin, according to the EDS spot analysis (the darker area in its ‘‘SW’’ part is an artefact originated by the decomposition of the resin used for mounting the sample); many other spherical (droplets) particles with the same brightness level (and, therefore, composition) can be observed around the larger particles as an exudation. At higher magnification particles a, b and c in Fig. 5 show a texture with a basic pattern: a carbonaceous matrix

Table 2 Approximate thermodynamic parameters of the reactions for tin possibly involved during the carbonisation, calculated from the entropies and the enthalpies of formation, considering that the approximate thermodynamic parameters for coke is that for graphite and for sulphur in coke that of CS2 (or thiophene, that is quite close) Reaction

DH at 25 C (kcal/mol)

DS at 25 C (cal/mol)

DG (kcal/mol) At 700 C

2 SnO « Sn + SnO2 SnO2 + C « Sn + CO2 SnO2 + 1/2C « SnO + 1/2CO2 SnO2 + 1/2CS2 + 1/2C « SnS + CO2 SnO + 1/2CS2 « SnS + 1/2CO2 Sn + 1/2CS2 « SnS + 1/2C

2 45 23 1 11 32

All the data have been taken from Weast [20].

2 50 26 33 7 16

0 3 2 34 18 17

At 800 C 0 8 4 37 19 15

At 900 C 1 13 7 40 19 13

At 1000 C 1 18 10 43 20 12

1400

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

Fig. 5. BSE–SEM images of the crushed (top) and polished section (bottom) of the VR800 sample. Selected areas of interest are shown at higher magnification.

(dark background, which should not be confused with porous) in which dendrites or particles of tin compounds (brighter features) are included. In particle ‘‘a’’ most of the non carbonaceous material is dark grey, which should be tin sulphide and maybe some remaining oxide, besides some bright spots that should be metallic tin. On the other hand, the texture of particle ‘‘b’’ is more mosaic type (indicative of a growing of the crystalline materials) with distinguishable particles of medium grey; the content of sulphur and oxygen in this area is similar to that in particle ‘‘a’’, but the concentration of tin is higher, indicating a higher degree of reduction. Finally, the carbon content in ‘‘c’’ is quite low, correspondingly to the lower amount of dark (carbonaceous) matrix that can be observed, as well as that of sulphur, which is almost negligible; therefore, this particle should be mainly very dispersed tin, in the form of very small particles (clearly in the submicron scale), maintained agglomerated by the remaining carbonaceous matrix and sulphide.

It should be remembered that at any of the carbonisation temperatures, the tin formed has to be in liquid state (mp = 231.93 C [20]), retained within the porous system of the carbonaceous material, which acts as a ‘‘sponge’’. Nevertheless, some liquid tin exudates to the surface of the particles, rising up the droplets that can be observed on them, growing up, coalescing and building up the bigger drops or particles shown in Fig. 5. The higher the temperature, the lower the viscosity of tin, the higher the conversion of tin oxide to metallic tin and sulphide, and the higher the crystallinity of the tin compounds are. These entire effects give rise to the different types of particles observed. Thus, the proportion of the different types of areas is not the same in all the samples, being VR1000 the most homogeneous and that with more drops of tin. The compositional maps determined by EDS (Fig. 6) corroborate those conclusions. Oxygen is quite uniformly distributed and with very low concentration. Sulphur and tin show some inhomogeneity, although not very impor-

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

1401

Fig. 6. Quantitative EDS maps, after ZAF correction, of the VR700 sample for the elements C, O, S (using K lines) and Sn (using L line), with a resolution of 200 · 256 pixels (around 0.35 lm). A individual spectrum is recorded for 1 live second for each pixel. The operational conditions were 20 kV and 0.3 bars of pressure. The BS image is also included for reference purposes.

tant and not always coincident. This means that, at the same time, there exist areas enriched in metallic tin, tin sulphide and (probably organic) sulphide. For example, this is very clear in the lower left corner of the image in Fig. 6, where a droplet of metallic tin can be observed between two particles, in which borders there is a accumulation of tin sulphide and maybe metallic tin. On the other hand, in the top left part of the main particle, a spot of high concentration of sulphur is observed, without an increase of the tin concentration. 3.3. Electrochemical behaviour of lithium cells The discharge and charge branches for the first two cycles for the carbonised material without tin oxide (VR0), VR800 and pure SnO2 are displayed in Fig. 7, in order to be compared. The behaviour of the VR0 is that well-known for LT carbons [16]. Likewise, the electrochemical curves for SnO2 show in the first discharge an extended plateau at 0.9 V, corresponding to the reduction of Sn(IV) to Sn (1.5 Ah/g). After that, some overlapping bands down to 0 V are observed, due to the expected formation of the successive intermetallic compounds between Li and Sn, as Li2Sn, LiSn, Li7Sn2 and Li22Sn5 [23]. Obviously, in charging and further cycles only the reversible formation of the LixSny species is detected, as the oxide is not formed any longer. On the other hand, the profiles for the electrochemical cycling of the VR800 composite can be considered as the contribution of both tin products and carbonaceous material. Thus, electrochemically formed intermetallic tin compounds leads to an increase of the discharge capacity of ca. 186 mA h/g. The irreversibility in the first discharge is still present, due to the carbonaceous part of the material and the remaining oxidised tin species, mainly sulphide. In this way, the broadened band corresponding to the SEI forma-

tion is clearly observed below 1.0 V. Moreover, the set of reversible peaks assigned to the formation of the Li–Sn alloys are detected both on charge and discharge, although the signals differ in shape and intensity from those shown by pure SnO2, probably due to the surface reactivity of tin compound particles when the alloy is formed. The characteristic peak of SnO2 reduction is not observed, as it was reduced during carbonisation, as already indicated before. The effect of carbonisation temperature on the electrochemical behaviour of the composite material can be deduced from Fig. 8. Although the I versus V curves show a similar pattern in all cases, the V versus capacity curves evidence a decreasing polarization with the carbonisation temperature, which can be inferred from the diminution of the hysteresis between charge and discharge branches. Also, on increasing the carbonisation temperature, the overall discharge capacity decreases, especially from 700 C to 900 C. Both effects can be ascribed to the increase of the structural ordering in and between the layers of the carbonaceous phase during carbonisation (and the consequently more suitable path for lithium migration) and to the simultaneous condensation of these graphene layers, with the consequent elimination of edges and defects, specially free radicals, which act as ‘‘anchor’’ sites for Li [24]. On the other hand, Fig. 9 shows the potentiostatic cycling (10 mV/0.1 h, 0.002-1 V; [25]) for up to cycle 25 for the composite (VR700–VR1000) and reference (VR0) samples, SnO2 and the additional sample prepared from VR800 by CVD (VR800–CVD). After the first discharge, there are not remarkable differences neither a clear tendency with the carbonisation temperature. It would be indicative that carbonisation mainly acts on removing surface sites where lithium is irreversibly trapped. The tin reduction upon carbonisation also contributes to decrease significantly the irreversible capacity recorded for the

1402

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409 Voltage (V)

Capacity (mAh/g) 0

200

400

600

0.0

0.5

1.0

1.5

2.0

2.5

2.5

200

VR0 100

1.5 0 1.0

I (mA/g)

Voltage (V)

2.0

-100

0.5

750ºC - 6 hours VR0

0.0 2.5

200

VR10 Voltage (V)

1.5

0

1.0

I (mA/g)

100

2.0

-100 0.5

800ºC - 6VR800 hours

0.0 0

200

400

600

0.0

0.5

Capacity (mAh/g)

1.0

1.5

2.0

2.5

Voltage (V) 200

2.5

SnO2 0

1.5 -200 1.0

I (mA/g)

Voltage (V)

2.0

-400

0.5

SnO2

0.0

-600 0

500

1000

1500

Capacity (mAh/g)

0.0

0.5

1.0

1.5

2.0

2.5

Voltage (V)

Fig. 7. Discharge/charge (V versus capacity and I – mass normalized – versus V) curves (first two cycles) for VR0 (top), VR800 (middle) and pure SnO2 (bottom).

original SnO2 upon the first discharge. For subsequent cycles, capacity values are kept higher than those of the theoretical capacity of graphite. Particularly, VR800 is the sample showing the best retention. The initial irreversible capacity and the slight fading of the reversible capacity with cycling are typical of carbonaceous materials, and they are intimately related, specially the first effect, with the irreversible accumulation of Li in the SEI, which is in turn largely affected by the surface of the material. For these reasons, we prepared a new sample by CVD. As it can be checked in Fig. 9, this treatment improves the cycling behaviour of the material, especially for longer periods, getting a more stable cycling and an increase of the specific capacity in more than 20 mA h/g. 3.4. Impedance spectroscopy The decrease of the charge/discharge hysteresis when the carbonisation temperature is increased (Fig. 8) may be related to the kinetic response of the electrode material. To evidence this point, the impedance spectra were recorded for two selected samples, VR700 and VR900, as

well as for VR0 as reference for elucidating the compositing effect. The profile of the spectra is characteristic of disordered carbons. Fig. 10 shows the Nyquist plots for the selected samples. The relevant features are two depressed semicircles at high and intermediate frequencies, and a straight line in the low frequency extreme. The experimental data were fitted using an equivalent circuit reported elsewhere [16] (Fig. 11). R1, R2 and R3 are ascribed to the resistance imposed to lithium migration through electrolyte solution, electrolyte interface (SEI) and insertion reaction into the bulk, respectively. The semicircle depression is usually assigned to the presence of inhomogeneities in the electrode material, caused by roughness, porosity and/or polycrystalline state, which hinder the frequency dispersion in the interface. A constant phase element (Q) takes into account the depression of the semicircles [26]. A Warburg component (W) is introduced to describe the system response to the Li+ diffusion through the intercalated carbon. The introduction of tin particles in the carbon matrix notoriously improves the conductivity of the composite, leading to a significant diminution of both R2 and R3 resis-

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409 Capacity (mAh/g) 0

200

400

600

1403

Voltage (V) 800 0.0

0.5

1.0

1.5

2.0

2.5

200

2.5

Voltage (V)

1.5

0

1.0

I (mA/g)

100

2.0

-100

700ºC - 6 hours

0.0 2.5

200

2.0

100

1.5

0

1.0

I (mA/g)

Voltage (V)

0.5

-100

800ºC - 6 hours

0.0 2.5

200

2.0

100

1.5

0

1.0

I (mA/g)

Voltage (V)

0.5

-100

900ºC - 6 hours

0.0 2.5

200

2.0

100

1.5

0

1.0

I (mA/g)

Voltage (V)

0.5

-100 0.5

1000ºC - 6 hours

0.0 0

200

400

600

800 0.0

0.5

Capacity (mAh/g)

1.0

1.5

2.0

2.5

Voltage (V)

Fig. 8. Discharge/charge (V versus F and I versus V) curves for the composite materials carbonised at different temperatures (700, 800, 900 and 1000 C) for 6 h. The intensity has been mass normalized.

1300 VR700

1200

VR800

1100

VR900

capacity (mAh/g)

1000

VR1000

900

VR800-CVD

800

VR0

700

SnO2

600

VR0+Sn

500 400 300 200 100 0 1

2

3

4

5

6

7

8

9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 cycle

Fig. 9. Potentiostatic cycling (10 mV/0.1 h, 0.002-1 V) for the composite Sn–C samples studied, including that without tin (VR0) and compared to SnO2 and the VR800 sample treated by carbon deposition of toluene (VR800-CV; see text) and a mechanical mixture of 80% of VR0 and 20% of SnO2 (VR0 + Sn).

1404

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

of the tin-less coke (VR0), although its evolution is linear as well, the slope is clearly higher. In the case of R3, the picture is not too different, except for higher absolute values of the resistance and a much more clear difference between the slopes of VR700 and VR900, higher for the first material. Besides, while for R2 in the composite materials the ordinate in the origin can be assumed to be 0, this is not the case for R3 in any of the studied materials. 3.5. Spectroscopic characterisation of partially discharged materials For a better understanding of the mechanism involved in the lithium reaction with the Sn/C composite, several electrodes prepared using VR700 were partially discharged and charged along the first cycle (marked points in Fig. 13) and examined by XRD, 119Sn Mo¨ssbauer and 7Li NMR. 3.5.1. XRD The electrochemical process induces a progressive loss of crystallinity in the material (Fig. 14). Although a num-

Fig. 10. Nyquist plots for VR700, VR900 and VR0 samples after the first and tenth cycle.

2.8

Q2

Q1

2.4

R1

g

R2

2.0

Voltage (V)

R3 W1

Fig. 11. Analogue circuit for the fitting of the impedance response.

1.6

1.2

f a

0.8

e

tance values at the end of the first discharge. Also, the effect of the carbonisation temperature is reflected by a decrease of electrode internal resistances when the treatment temperature increases. According to the graphic in Fig. 12, it can be check that R2 values, although generally low, evolve practically in a linear form, maybe with a slightly bigger slope for the lower temperature sample (VR700), which can be indicative of the positive effect of compositing on forming a stable passivation film. In contrast, for the R2

d 0

150 1

300

450

600

Capacity (m (mAh/g)

Fig. 13. Discharge–charge curve of a cell prepared from VR700. For the indicated points by circles, the cell was open and the carbonaceous material analysed (see text).

0.6

VR0(750)6 VR0

VR0750(6) VR0

700ºC-6h VR700

700ºC-6h VR700

900ºC-6h VR900

0.5

⋅

VR900 900ºC-6h

⋅

0.4

0.15

R3 (Ohm.g)

R2 (Ohm.g)

c

0.0

0.25

0.2

b

0.4

0.1

0.3 0.2

0.05

0.1

0

0 0

2

4

6 Cycles

8

10

12

0

2

4

6 Cycles

8

10

Fig. 12. Evolution of R2 and R3 (see Fig. 11) according to cycling for VR0, VR700 and VR900.

12

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

+

+

Intensity / a.u.

(g) (f) (e) (d) (c)

**

***

(b) (a)

15

20

25

30

35

40

45

50

Angle / 2 Theta

Fig. 14. X-ray diffractograms of the materials partially charged/discharged prepared using the VR700 coke, corresponding to the points marked in Fig. 13.

ber of crystalline forms of these alloys have been reported [27], the nanosized Li–Sn intermetallic particles formed upon discharging preclude the clear discern of all their reflections as expected for high crystallinity products. The pattern of the electrode discharged at 0.3 V (b in Fig. 14) shows additional small reflections around 27.5 and 37.5 (2h) (marked as *) which can be correlated with low lithium content tin alloys as Li2Sn5 and LiSn. Deeper discharges beneath 0.2 V involved the disappearance of these phases and the development of a new broadened reflection at around 38, whose intensity continuously increases up to the end of the discharge at 0 V. This new signal could be assigned to Li22Sn5. However, peaks of b-Sn, which have been progressively fading as expected, are still observed at the end of the first discharge. This fact can be ascribed to either a kinetic effect during the discharge of the electrodes or the existence of less accessible tin particles isolated into the carbon matrix. The highly broadened bands, especially (0 0 2), of the carbonaceous material are really difficult to discern, and tentatively could be pointed out when the lithium content is the highest, that is, when the degree of amorphization of tin compounds is also high. On charging, the low degree of crystallinity is preserved, although the reversibility of the electrochemical reaction is deduced from the loss of Li22Sn5 reflection and the partial recovering of intensity for b-Sn signals. 3.5.2. 7Li NMR The 7Li NMR spectroscopy has been extensively used to describe the formation of Li–Sn compound upon electrochemical reaction. Idota et al. [28] identified large chemical shifts into SnO to an extent close to those of the Li–Sn alloys. Others authors have confirmed the occurrence of Li–Sn compounds electrochemically produced by compar-

0.8 V

0.3 V

Discharge

+ +

0.1 V

0.0 V

0.5 V Charge

*

ison of their spectra with those of chemically obtained phases. However, the electrochemical route yields very different lithium local environment when maximum Li/Sn ratio is reached at the end of the discharge, what points out to a structurally different phases [29]. Further studies on TCO and SnO demonstrated a more complex mechanism in which the participation of oxygen must not be ruled out [30]. For VR700 discharged at 0.8 V, an unique signal is observed in the spectrum, located around 0 ppm (Fig. 15). This signal can be actually decomposed in two Lorentzian components with distinct linewidth, corresponding to highly ionic lithium bonds. Taking into account that SnS and SnO2 are reduced at 1.4 and 0.9 V, respectively, the broadened Lorentzian signal can be associated to amorphous lithium sulphide and some possible remaining oxide, while the narrower signal probably arises from lithium ions from LiPF6 not efficiently rinsed and/or the alkaline metal ion consumed on the formation of a passivating film on the carbon particle surface [31]. On increasing the lithium content, new bands appear, including signals at ca. 75 ppm (almost negligible), 41 ppm (most intense), 30 ppm (medium) and 16 ppm (low) for the electrode discharged at 0.3 V (Fig. 15). Previous works have shown how the bands ascribable to Li–Sn compound shifts toward low field values when the lithium content increases

Arbitrary units

+ Sn LiSn Li22Sn5

1405

1.0 V

2.1 V

80

60

40

20

0

-20

-40

Chemical shift (ppm)

Fig. 15. 7Li NMR spectra of the partially charged/discharged cells using VR700 material.

1406

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

during Li/SnO cell discharging [30]. In fact, a non-linear variation of the shift versus Li/Sn ratio is reflected by an initial decrease of its values, which is monitored on increasing the lithium content from LiSn2 (75 ppm) to Li7Sn2 (7 ppm). Then, an abrupt increase is observed, leading to a shift value around 100 ppm for Li22Sn5 [32]. From these data, the sig-

nals around 41, 30 and 16 ppm can be ascribed to LiSn, Li7Sn3 and Li5Sn2, respectively. Moreover, new signals at 1.5 and 5 ppm appear near 0 V. The signal centred at 1.5 ppm is assigned to Li2S (and/or Li2O), notoriously increased in the first stages of the discharge. The signal located at 5 ppm continuously increases and is also slightly

1.001

% Absorption

1 0.999 0.989 0.997 Discharged at 0.8 V

0.996

1.001

1.001

1

1

0.999

0.999 % Absorption

% Absorption

0.995 -6

0.989 0.997 0.996 0.995 0.994 -6

Charged at 0.5 V

6

-2

0 2 Velocity (mm/s)

4

6

-2

0 2 Velocity (mm/s)

4

6

-2

0 2 Velocity (mm/s)

4

6

Discharged at 0.3 V

0.995 0.994

-4

-2

0 2 Velocity (mm/s)

4

-6

6

1

1

0.999 % Absorption

0.999 % Absorption

4

0.996

1.001

0.998 0.997 0.996 0.995

0.991 -6

0 2 Velocity (mm/s)

0.997

1.002

0.992

-2

0.989

1.001

0.994 0.993

-4

-4

0.989 0.997 0.996 0.995 0.994

Charged at 1.0 V

Discharged at 0.1 V

0.993 0.992

-4

-2

0 2 Velocity (mm/s)

4

-6

6

-4

1.002 1.0005 1 1 0.9995

% Absorption

% Absorption

0.998 0.999 0.9985 0.998

Charged at 2.1 V

0.992

0.988

0.9965 -6

0.994

0.99

0.9975 0.997

0.996

-4

Fig. 16.

-2

117

0 2 Velocity (mm/s)

4

6

0.986 -6

Discharged at 0V

-4

Sn Mo¨ssbauer spectra of partially discharged/charged cells built from VR700.

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

shifted to low field values when the discharge progresses, behaviour quite similar to that observed in others disordered carbons [33]. Further discharge to 0.1 V promotes the loss of intensity of the Li/Sn alloys with low lithium contents, while the band located at ca. 15 ppm increases. Signals at 7 ppm and/or 100 ppm due to the presence of Li7Sn2 and/or Li22Sn5 were absent. At 0 V, the peaks corresponding to Li–Sn phases strongly decrease in intensity as compared to that attributed to the Li reaction with the carbonaceous material. Therefore, carbon matrix is involved not only in the stabilisation of the Li–Sn intermetallic compounds formed upon electrochemical lithium reaction, but also contributes to capacity by its own lithium reaction in the whole range of cell potential down to 0 V. On charging, the reversible reaction is observed for the carbon material, while the presence of signals ascribed to Li–Sn compounds and ionic lithium compounds were still detected even at cell potentials as higher as 2.1 V. This result seems to confirm that charging at cut off voltages higher than 1 V does not contribute to the reversibility of the process. 3.5.3. 119Sn Mo¨ssbauer spectroscopy Complementary information about the changes in local environment of tin atoms and their oxidation states along the first cycles can be obtained using 119Sn Mo¨ssbauer spectroscopy. Fig. 16 shows the spectra recorded at different points of the first cycle. All of them are characterized by a broadened and asymmetric profile due to the overlapping of several tin species. Li–Sn alloys have been chemically synthesized by other authors and further analysed by this technique to describe the tin local environment according the structure of these intermetallic compounds [27]. For samples obtained by an electrochemical procedure, it is difficult to achieve pure rather than a mixture, as has been evidenced by the 7Li NMR spectra. The complexity of the profiles is enhanced if we consider the possible interactions of Li–Sn alloy particles with the oxygen atoms surrounding them [34]. For this reason, the fitting of our spectra was carried out taking into account the hyperfine parameters previously reported for pure phases [27], allowing the contribution to vary. For an electrode discharged at 0.82 V, with a lithium content equivalent to 74 mA h/g, the first observed effect is the sharply decrease of the splitted band previously ascribed to SnS. This reduction process agrees well with the electrochemical profile (Fig. 13) because this electrochemical reaction is recorded at ca. 1.4 V. Moreover, no change seems to be detected for b-Sn. Unexpectedly, the SnO2 signal, centred around 0 ppm, seems to increase in intensity. Oxidation phenomena arising by an erroneous handling of the partially discharged electrode were discarded after repeating the experience and taking special care on the manipulation to obtain reproducible results. We can speculate about internal oxidation reactions undergoing between freshly reduced tin atoms in close contact

1407

with high reactive SEI products, because both species are generated at similar cell potentials along the cell discharge. In the spectrum of the electrode discharged at 0.3 V (324 mA h/g) the centroid is clearly shifted to lower isomer shift values, due to the transformation of metal tin into Li– Sn alloys. Bearing in mind the results of the well resolved bands in NMR spectra, the MS results indicate the complete disappearance of Sn(II) and a minor contribution of b-tin. This result clearly points out the non-crystalline character of these intermetallic compounds, as they are barely discernible in XRD patterns. The lack or crystallinity can be regarded as a beneficial factor for an enhanced electrochemical behaviour, as the important volume difference found between crystalline intermetallics and tin, can be minimized in highly dispersed non-crystalline or nanosized materials. Further discharge to 0 V involves continuous changes in spectra profiles and centroid shifting to lower isomer shift. The deconvolution of the profiles (Table 3) agrees well with the successive appearance of new Li–Sn alloys with higher content in the alkali metal. Also, the contribution of b-Sn was almost negligible as compared to the intermetallic phases, despite of the observation in XRD pattern. As expected, the general trend of the charging process evidences the reversible extraction of lithium from the alloys along tin oxidation (Fig. 16 and Table 3). Diluted Li–Sn alloys are retrieved and even metallic tin is recovered at 1.0 V. However, only minor amounts of SnO2 and SnS were produced when the electrode was taken to 2.1 V. Moreover, a significant contribution of alloys remained after the oxidation, indicating a poor electron retrieval over

Table 3 Phases detected by Mo¨ssbauer spectroscopy and their concentration calculated from the deconvolution of the spectra shown in Fig. 16 corresponding to partially charge/discharge VR700 samples

Discharge

Sample (d/c v)

Phase

C (%)

0.8

SnO2 Sn SnS Sn LiSn Li7Sn3 LiSn Li7Sn3 Li5Sn2 Li7Sn3 Li5Sn2

21.1 76.9 2.0 14.3 46.2 39.5 6.0 64.8 29.2 34.9 65.1

LiSn Li7Sn3 Li5Sn2 Sn LiSn Li7Sn3 Sn LiSn Li7Sn3 SnS SnO2

9.0 40.3 50.7 20.4 66.8 17.5 32.1 28.0 21.7 9.1 9.1

0.3

0.1

0 Charge

0.5

1

2.1

1408

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409

1 V. In fact, the electrochemical curves show an abrupt increase of the cell potential at the end of the charge in which no significant electrochemical processes undergoes. This fact must be assigned to kinetic factors which are revealed by the high polarization of the cycles for the sample carbonised at 700 C. Undoubtedly, this poor kinetic response is avoided by the increase of annealing temperature, albeit the decrease in the overall capacity values.

4. Conclusions Sn–C composites electrochemically active for Li-ion batteries can be obtained by a conventional thermal treatment (pyrolysis and carbonisation) of a mechanical mixture of a vacuum residue and tin dioxide. The electrochemical profiles revealed the contribution of both components to the overall capacity. The conversion reaction of SnS was not as significant as in the case of other metal carbon composites previously studied. The presence of tin oxide modifies the pyrolysis process, which is reflected in the optical texture of the coke. During the pyrolysis and the further carbonisation, the SnO2 is reduced to Sn(II), as SnS (SnO is not observed at all), and metallic tin. The reduction is not straightforward, but the Sn(IV) is reduced to metallic tin or SnS; at the same time, the Sn can be reoxidised to SnS, while SnS cannot be transformed in Sn without the intervention of oxygen or hydrogen due to thermodynamic concerns. During all the thermal processes, the metallic tin formed is liquid, infiltrating the carbonaceous matrix as if this were a sponge. The metallic tin is eventually exuded and forms drops on the surface of the carbonaceous ‘‘sponges’’. This exudation is higher at higher temperatures. Nevertheless, a high amount of tin is maintained inside the carbonaceous matrix, forming a real composite as the X-ray and electron microscopy images show. The contribution of both main components, the tin and the carbonaceous phases, can be observed in the electrochemical behaviour of the composite material. The inclusion of tin involves a proportional increase in the capacity of the composite material. The samples carbonised at the highest temperatures show the lowest irreversible capacity in the first cycle, but it is still high. After the first cycle, the capacity of the material seems to stay stable near 500 mA h/g, mainly in the case of the sample VR800, probably due to the stabilizing and muffing effect of the carbonaceous matrix against the volumetric changes during the formation and deformation of the Li–Sn alloys. A better behaviour is obtained if the sample is treated by a process of carbon deposition. 119 Sn Mo¨ssbauer and 7Li MAS NMR spectroscopies have demonstrated to be useful techniques to reveal, respectively, the evolution of tin and lithium phases along the electrochemical reaction. For the selected sample, carbonised at 700 C, a number of Li–Sn intermetallic compounds were produced on discharging but only partially

removed at the end of the charge sweep. Although this fact may negatively affect to the electrochemical performance, a higher temperature of carbonisation could contribute to minimize its impact. Moreover, a close inspection of the charge process makes evident the low efficiency of charge retrieving over 1 V, while electrode degradation at high voltages must not be ruled out. In fact, better performances at long term cycling were achieved when the cell charging was limited to 1 V. Also, the deposition of a thin layer of carbon material by CVD significantly improved the capacity retention for a large number of cycles. Acknowledgements The authors are grateful to CICYT for financial support (Contract MAT2005-00374 and contract MAT2001-1694) and M.C. Mohedano for her technical support. We thank to SCAI (UCO Central Service for Research Support in the HRTEM and 7Li NMR characterization). A. Concheso is indebted to MCYT for his predoctoral grant. References [1] Flandrois S, Simon B. Carbon materials for lithium-ion rechargeable batteries. Carbon 1999;37:165–80. [2] Aurbach D, Zinigrad E, Cohen Y, Teller H. A short review of failure mechanisms of lithium metal and lithiated graphite anodes in liquid electrolyte solutions. Solid State Ionics 2002;148(3–4):405–16. [3] Aurbach D, Teller H, Koltypin M, Levi E. On the behavior of different types of graphite anodes. J Power Sources 2003;119–121:2–7. [4] Wakihara M. Recent developments in lithium ion batteries. Mater Sci Eng 2001;R33:109–34. [5] Alca´ntara R, Ferna´ndez Madrigal FJ, Lavela P, Tirado JL, Jime´nez Mateos JM, Stoyanova R, et al. 13C, 1H, 6Li magic-angle spinning nuclear magnetic resonance, electron paramagnetic resonance, and fourier transform infrared study of intercalation electrodes based in ultrasoft carbons obtained below 3100 K. Chem Mater 1999;11(1): 52–60. [6] Alca´ntara R, Ferna´ndez Madrigal FJ, Lavela P, Tirado JL, Jime´nez Mateos JM, Go´mez de Salazar C, et al. Characterisation of mesocarbon microbeads (MCMB) as active electrode material in lithium and sodium cells. Carbon 2000;38(7):1031–41. [7] Lavela P, Tirado JL. Baterı´as Avanzadas. Co´rdoba (Spain), Servicio de Publicaciones de la Universidad de Co´rdoba, 1999. pp. 219–229. [8] Balan L, Schneider R, Ghanbaja J, Willmann P, Billaud D. Electrochemical lithiation of new graphite-nanosized tin particle materials obtained by SnCl2 reduction in organic medium. Electrochim Acta 2006;51:3385–90. [9] Noh M, Kwon Y, Lee H, Cho J, Kim Y, Kim MG. Amorphous carbon-coated tin anode material for lithium secondary battery. Chem Mater 2005;17:1926–9. [10] Xie J, Varadan VK. Synthesis and characterization of high surface area tin oxide/functionalized carbon nanotubes composite as anode materials. Mater Chem Phys 2005;91:274–80. [11] Wang Y, Zeng HC, Lee JY. Highly reversible lithium storage in Porous SnO2 nanotubes with coaxially grown carbon nanotube overlayers. Adv Mater 2006;18:645–9. [12] Grigoriants I, Soffer A, Salitra G, Aurbach D. Nanoparticles of tin confined in microporous carbon matrices as anode materials for Li batteries. J Power Sources 2005;146:185–9. [13] Patel P, Kim IS, Maranchi J, Kumta P. Pyrolysis of an akyltin/ polymer mixture to form a tin/carbon composite for use as an anode in lithium-ion batteries. J Power Sources 2004;135:273–80.

J.L. Tirado et al. / Carbon 45 (2007) 1396–1409 [14] Jung YS, Lee KT, Ryu JH, Im D, Oh SM. Sn-carbon core–shell powder for anode in lithium secondary batteries. J Electrochem Soc 2005;152(7):A1452–7. [15] Billaud D, Balan L, Schneider R, Willmann P. The influence of the synthesis conditions of graphite/tin nanoparticle materials on their electrode electrochemical performance in Li-ion battery anodes. Carbon 2006;44:2508–15. [16] Concheso A, Santamarı´a R, Blanco C, Mene´ndez R, Jime´nez-Mateos JM, Alca´ntara R, et al. Composite electrode materials for lithium-ion batteries obtained by metal oxide addition to petroleum vacuum residua. Carbon 2005;43(5):923–36. [17] Goldstein JI, Newbury DE, Echlin P, Joy DC, Romig Jr AD, Lyman CE, et al. Scanning electron microscopy and X-ray microanalysis: a text for biologists, materials scientists, and geologists. Second edition. New York, USA: Plenum Press; 1992. [18] Friel JJ, Lyman CE. X-ray mapping in electron-beam instruments. Microsc Microanal 2006;12:2–25. [19] Robert F, Lippens PE, Olivier-Fourcade J, Jumas JC, Morcrette M. Mo¨ssbauer spectroscopy as an efficient tool to study Li insertion mechanisms in negative electrodes for Li ion batteries. J Power Sources 2005;146:492–5. [20] Weast RC, editor. Handbook of Chemistry and Physics. 62nd ed. Boca Raton, USA: CRC Press Inc.; 1982. p. D52–64. [21] Carreira Anto´n P. Produccio´n de materiales grafı´ticos mediante copiro´lisis de un residuo de petro´leo y compuestos de silicio o boro (Production of graphitic materials by co-pyrolysis of a petroleum residue and compounds of silicon and boron). University of Alicante, Spain. PhD Thesis, 2002. [22] Ellis PJ, Paul CA. Delayed coking fundamentals. In: Proc of Int Conf on Refining Processes, Topical Conf Preprints. AIChE Spring National Meetings, 1998. Atlanta, USA. [23] Tirado JL. Inorganic materials for the negative electrode of lithium ion batteries: state of the art and future prospects. Mater Sci Eng R 2003;40:103–36.

1409

[24] Wu YP, Wan CR, Jiang CY, Fang SB, Jiang YY. Mechanism of lithium storage in low temperature carbon. Carbon 1999;37: 1901–8. [25] Courtney IA, Dhan JR. Key factor controlling the reversibility of the reaction of lithium with SnO2 and Sn2BPO6 glass. J Electrochem Soc 1997;144(9):2943–8. [26] McDonald JR. Impedance spectroscopy, emphasizing solid materials and systems. USA: John Wiley and Sons; 1987. [27] Dunlap RA, Small DA, MacNeil DD, Obrovac MN, Dahn JR. A Mo¨ssbauer effect investigation of the LiSn system. J Alloys Comp 1999;289(1–2):135–42. [28] Idota Y, Kubota T, Matsufuji A, Maekawa Y, Miyasaka T. Tinbased amorphous oxide: a high-capacity lithium-ion-storage material. Science 1997;276(5317):1395–7. [29] Wang Y, Sakamoto J, Huang CK, Surampudi S, Greenbaum SG. 7 Lithium NMR investigation of electrochemical reaction of lithium with SnO. Solid State Ionics 1998;110(3-4):167–72. [30] Goward GR, Leroux F, Power WP, Ouvrard G, Dmowski W, Egami T, et al. Tin composite oxide glasses. Electrochem Solid-State Lett 1999;2:367–70. [31] Goward GR, Taylor NJ, Souza DCS, Nazar LF. Electrochemical and multinuclear solid-state NMR studies of tin composite oxide glasses as anodes for Li ion batteries. J Mater Chem 2000;10(5):1241–9. [32] Furuya K, Ogawa K, Mineo Y, Matsufuji A, Okuda J, Erata T. Solid 7 Li NMR and in situ XRD studies of the insertion reaction of lithium with tin oxide and tin-based amorphous composite oxide. J Phys Condens Matter 2001;13(14):3519–32. [33] Letellier M, Chevallier F, Be´guin F, Frackowiak E, Rouzaud JN. The first in situ 7Li NMR study of the reversible lithium insertion mechanism in disorganised carbons. J Phys Chem Solids 2004;65: 245. [34] Chouvin J, Branci C, Sarradin J, Olivier-Fourcade J, Jumas JC, Simon B, et al. Lithium intercalation in tin oxide. J Power Sources 1999;81–82:277–81.