Rational fabrication of hybrid structure of SnO_x sandwiched between TiO₂ and carbon based on the complementary merits of SnO_x, TiO₂ and carbon, and its improved lithium storage properties

Abstract

In spite of high-profile theoretical capacity, the practical application of SnO₂ or Sn anode materials for lithium-ion batteries is severely impeded by poor electric conductivity and structural instability. Herein, a hybrid structure of SnO₂/Sn sandwiched between TiO₂ and carbon with rich porosity, good electric conductivity and stable structure, denoted as TiO₂@SnO_x@C, is fabricated based on the complementary merits of SnO₂, TiO₂ and carbon anode materials. The TiO₂@SnO_x@C exhibits a good electrochemical performance when used as anode material for lithium-ion battery, delivering a capacity of 629 mA h g⁻¹ at 200 mA g⁻¹ after 300 cycles. Moreover, a reversible capacity of 490.3 mA h g⁻¹ is obtained at 1000 mA g⁻¹ even after 1000 cycles and is much higher than theoretical capacity of graphite (372 mA h g⁻¹). The effectively complementary and synergic effect among structural stability of TiO₂, high theoretical capacity of SnO_x, and good conductive and flexible ability of carbon should be responsible for the superior electrochemical performance of TiO₂@SnO_x@C.

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1 Introduction

Extensive efforts have been made to explore alternative anode materials to replace graphite anode for meeting the pressing demands of next-generation lithium-ion batteries (LIBs) for high rate capability, safer power and higher capacity.^1–3 As one of the promising candidates, tin dioxide (SnO₂) or metallic tin (Sn) has attracted great research interest due to the high theoretical capacity (782 mA h g⁻¹ for SnO₂, 992 mA h g⁻¹ for Sn) and safe working potential.^4–6 However, the biggest obstacle for employing SnO₂ or Sn as applicable anode materials for LIBs is that it is suffering from huge volume variation during Li insertion/extraction cycle, which ultimately causes pulverization of the electrode, very rapid capacity decay and inferior cycling ability.⁷ Another promising anode material, titanium dioxide (TiO₂), has also been widely studied due to the superior structural stability during long-term cycling and higher working potential.⁸ Unfortunately, the practical application of TiO₂ as anode for LIBS is also impeded by low electronic conductivity and low practical capacity (ca.170 mA h g⁻¹) in particular.⁹

Considering SnO₂, TiO₂ and carbon anode materials have complementary merits, a lot of composites consisted of the three substances have been reported, such as SnO₂@TiO₂/carbon cloth,¹⁰ graphene-based TiO₂/SnO₂,¹¹ mesoporous TiO₂–Sn@C,¹² graphene–TiO₂–SnO₂,¹³ TiO₂/SnO₂/carbon hybrid nanofibers,¹⁴ and hierarchical TiO₂–SnO₂–graphene,¹⁵ all them exhibited improved electrochemical performance compared with respective component, which due to the synthetic strategy of components: (1) the conductive carbon substrate can serve as a current collector to maintain the high electron transfer pathways within the electrode and partly buffer the volume variation of SnO₂; (2) the structural stability of TiO₂ can serve a stable support to maintain the structural integrity of electrode and prevent the SnO₂ from aggregation; (3) the high theoretical capacity of SnO₂ can contribute to the desirable reversible capacity of composites.

Herein, inspired by the studies mentioned above, an interesting architecture of SnO₂/Sn sandwiched between TiO₂ and outermost carbon coating layer with mesoporous and hollow structure, denoted as TiO₂@SnO_x@C, has been prepared by a facile method. The sandwich structure supported by TiO₂ can effectively suppress the pulverization and aggregation of SnO₂/Sn, and strengthen the stability of electrode structure. The hollow and mesoporous can provide free space for buffering the huge volume change of SnO₂/Sn. The outermost carbon coating layer can enhance the electronic conductivity of whole electrode, facilitate electron and ion transport throughout the electrode, and further strengthen the integrity of whole electrode structure. As a result, the as-prepared TiO₂@SnO_x@C composite exhibited excellent cycling and rate performances, delivering a capacity of 629 mA h g⁻¹ at 200 mA g⁻¹ after 300 cycles, and a reversible capacity of 490.3 mA g⁻¹ at 1000 mA g⁻¹ even after 1000 cycles.

2 Experimental section

2.1 Preparation of sample

2.1.1 The preparation of TiO₂ microsphere precursor. Briefly, 2.5 mL of tetrabutyl titanate (TBOT) was slowly dropped into 100 mL ethanol containing 0.4 mL of 0.1 M KCl aqueous solution. After stirred for 10 min, this white suspension was aged in a static condition for 24 h in a closed container at room temperature. Then, the white precipitate was collected, washed with ethanol and deionized water to gain monodispersed amorphous TiO₂ microsphere precursor.

2.1.2 The preparation of TiO₂@SnO₂ hollow microspheres. Typically, 0.1 g of TiO₂ precursor was dispersed in 60 mL of deionized water by ultrasound for 1 h. Then 0.2 g of SnCl₂·2H₂O was added into the suspension under stirring. After stirred for 30 min, 0.044 g of NH₄F was added into the yellow mixed suspension and then stirred another 30 min. After that, this yellow mixed suspension was transferred into a Teflon-lined stainless steel autoclave, and then placed in an oven at 180 °C for 2 h. After cooled down to room temperature, the TiO₂@SnO₂ hollow microspheres were collected by centrifugation, washed with deionized water and ethanol thoroughly, and dried in an oven at 60 °C overnight. For comparison, the TiO₂ hollow microspheres (TiO₂ HS) was prepared at the same conditions with TiO₂@SnO₂ but absence of SnCl₂·2H₂O. The control sample of SnO₂ hollow spheres was also prepared by according to our previous work, namely using amorphous SiO₂ microspheres to replace TiO₂ precursor and hydrothermal treatment for 48 h.¹⁶

2.1.3 The preparation of TiO₂@SnO_x@C. According to a previous literature with modifications,¹⁷ 0.5 g of TiO₂@SnO₂ was dispersed in a mixture of deionized water (70 mL) and ethanol (30 mL) by stirring, followed by the addition of 2.3 g of hexadecyltrimethyl ammonium bromide (CTAB), 0.35 g of resorcinol, and 0.1 mL of ammonium aqueous solution, and then stirring 1 h at room temperature. Then, 0.5 mL of formaldehyde solution was added to the suspension with continuous stirring for 8 h at 35 °C. After that, the products (TiO₂@SnO₂@RF) were collected by centrifugation, washed with deionized water and ethanol thoroughly, and dried in an oven at 60 °C overnight. Finally, the TiO₂@SnO_x@C was obtained by TiO₂@SnO₂@RF calcinations under Ar at 350 °C for 1.5 h, followed by a further annealing treatment at 600 °C for 3 h with a heating rate of 1 °C min⁻¹.

2.2 Materials characterization

The morphology and microstructure of the products were investigated by using field emission scanning electron microscopy (SEM, JEOL JSM-7401F) and transmission electron microscopy (TEM, JEOL JEM-2010) with an energy dispersive X-ray spectrometer (EDX) and a selected area electron diffraction pattern (SAED). The crystal structure and composition were characterized by X-ray diffraction measurement (XRD, Rigaku, D/max-Rbusing Cu K radiation) and X-ray photoelectron spectroscopy (XPS, AXIS ULTRA DLD instrument, using aluminum Kα X-ray radiation). The mesoporous structure and surface area were studied by nitrogen (N₂) adsorption/desorption isotherms (Micromeritics ASAP 2010 instrument). ICP analysis was conducted by an iCAP6300-type inductively-coupled plasma spectrometer.

2.3 Electrochemical characterization

Electrochemical measurements were conducted using 2016-type coin cells assembled in argon-filled glove box (German, M. Braun Co., [O₂] < 1 ppm, [H₂O] < 1 ppm). The working electrodes were composed of active material (TiO₂@SnO_x@C), conductive material (acetylene black, AB), and binder (sodium carboxymethyl cellulose, CMC) in a weight ratio of 7 : 2 : 1 and pasted on a Cu foil, followed by dried in a vacuum oven at 110 °C for overnight. The mass of active material loaded on electrode was ca. 1.3 mg and the diameter of electrode wafer was 14 mm. The pure lithium foil and glass fiber (GF/A) from Whatman were used as the counter electrode and separator, respectively. The electrolyte consisted of a solution of 1 M LiPF₆ in ethylene carbonate and dimethyl carbonate (EC + DMC) (1 : 1 in volume). The galvanostatic discharge/charge cycles were performed on a CT2001a cell test instrument (LAND Electronic Co.) over a voltage range of 0.01 to 3.0 V at room temperature (25 °C). Cyclic voltammetry (CV) was implemented on a CHI660D electrochemical workstation at a scan rate of 0.5 mV s⁻¹ between 0.0 and 3.0 V. Electrochemical impedance spectrum (EIS) measurements were also conducted using a CHI660D electrochemical workstation in the frequency range from 100 kHz to 0.01 Hz with an ac perturbation of 5 mV s⁻¹. The specific capacities reported and current densities used here were based on the total mass of TiO₂@SnO_x@C.

3 Results and discussion

Fig. 1a and b show the TEM images of TiO₂ precursor and TiO₂ HS, respectively, as well as the XRD patterns of both samples are showed in Fig. 2. It can be seen that the solid, surface smooth and amorphous TiO₂ precursor transformed to anatase TiO₂ HS with a hollow structure and rough surface after hydrothermal treatment under the presence of NH₄F. Then, the TiO₂@SnO₂ was prepared by amorphous TiO₂ precursor hydrothermal treatment under the simultaneous presence of NH₄F and SnCl₂, as shown in Fig. 1c. Obviously, the TiO₂@SnO₂ preserved the hollow structure of TiO₂ HS but the shell transformed from urchin-like to more rough and loose, indicating the SnO₂ was successfully introduced onto the shell of TiO₂ HS. The XRD patterns of TiO₂@SnO₂ is showed Fig. 2, in which the XRD profile of TiO₂@SnO₂ is comprised of two parts, one is consistent with anatase TiO₂ indicated by (1),¹⁸ and the other one is accord well with rutile SnO₂ marked by (2).¹⁹ The result of XRD patterns further verified the desirable TiO₂@SnO₂ was successfully obtained. Fig. 3a shows the SEM image of TiO₂@SnO_x@C, it can be clearly seen that the TiO₂@SnO_x@C preserved the hollow structure after carbon coating. The hollow structure of the TiO₂@SnO_x@C is further confirmed by the TEM characterization due to the clearly contrast between the centre and edge, as shown in Fig. 3b. Compared with TiO₂@SnO₂, the as-prepared TiO₂@SnO_x@C has a relative smooth surface, implying the carbon successfully coated on the surface of TiO₂@SnO₂. And the SAED patterns (inset of Fig. 3b) show a series of bright diffraction rings, indicating the polycrystalline nature of TiO₂@SnO_x@C, which is good agreement with the XRD result. The magnified TEM image of an outermost part of shell of TiO₂@SnO_x@C is showed in Fig. 3c, in which the outermost layer with a light region may be attributed to the carbon coating layer. For confirming this inference the HRTEM characterization was carried out on a part of the outermost layer of the shell, as shown in Fig. 3d. Distinctly, the lattice fringes corresponding to (110) lattice planes of SnO₂ with a spacing of 0.33 nm are observed and surrounded by amorphous substance, further indicating the TiO₂@SnO₂ is successfully coated by amorphous carbon. The XRD patterns of TiO₂@SnO_x@C is shown in Fig. 2, besides the peaks of TiO₂ and SnO₂ became stronger compared to TiO₂@SnO₂, the peaks for β-Sn is also observed which due to the reduction of SnO₂ by carbon in a high temperature of 600 °C. The EDX elemental mapping was also carried on TiO₂@SnO_x@C, as shown in Fig. 4. The EDX elemental mappings for C, O, Sn and Ti showed in Fig. 4b–e are match well the result shown in the SEM image (Fig. 4a), indicating SnO_x and C are introduced into TiO₂ matrix successfully. The Fig. 4b is almost covered by carbon element due to the additional carbon element derived from the conductive adhesive used as the sample stage in SEM measurements. In addition, the sandwich structure is also effectively confirmed by the TEM mapping images as shown in Fig. S1.†

Fig. 1 TEM images of TiO₂ precursor (a), TiO₂ HS (b), TiO₂@SnO₂ HS (c) and SnO₂ HS (d).

Fig. 2 XRD patterns of different samples.

Fig. 3 (a) SEM image of TiO₂@SnO_x@C; (b) TEM image of TiO₂@SnO_x@C, the inset showing the SAED patterns; (c) magnified TEM image of a part of (b) indicated by red rectangle; (d) HRTEM image of a part of (c).

Fig. 4 (b–e) EDX elemental mappings of the area of the SEM image (a) of TiO₂@SnO_x@C, for carbon, oxygen, tin and titanium, respectively.

The chemical composition and element states of TiO₂@SnO_x@C were also investigated by XPS measurement. The Fig. 5a shows the survey spectrum, there are four peaks located at around 284.8, 530.1, 487 and 459.1 eV, which are assigned to the C 1s, O 1s, Sn 3d and Ti 2p spectra, respectively. Fig. 5b shows the typical high-resolution XPS spectra of Sn 3d, besides respective two peaks located at 487 and 495.4 eV are corresponding to binding energies of Sn⁴⁺, another pair of peaks located at 484.9 and 493.7 eV, respectively, are also observed and corresponding to binding energies of Sn⁰, which is good agreement with the XRD patterns.²⁰ The Fig. 5c shows typical high-resolution XPS spectra of Ti 2p, binding energies for Ti 2p_3/2 at 458.8 eV and Ti 2p_1/2 at 464.6 eV which are assigned to Ti⁴⁺ in TiO₂.²¹ Fig. 5d shows the typical high-resolution XPS spectra of O 1s, it can be deconvoluted into three individual peaks located at 530.2, 532, and 533.3 eV, respectively, which can be corresponding to O²⁻ in TiO₂, O²⁻ in SnO₂, and oxycarbide or adsorbed water, respectively.^21,22 These results analyzed above demonstrated the TiO₂@SnO_x@C is successfully prepared.

Fig. 5 (a) Survey spectrum of TiO₂@SnO_x@C; (b–d) high-resolution spectra of Sn 3d, Ti 2p and O 1s, respectively.

The contents of respective TiO₂, SnO₂, Sn and C in TiO₂@SnO_x@C composite were also determined by ICP and TG analysis. The mass ratio of Sn/Ti is 44.43/53.83 determined by ICP. The weight increase between 500 and 750 °C can be attributed to the oxidation of β-Sn by O₂ during TG analysis, as shown Fig. 6. Thus, the contents of respective TiO₂, SnO₂, β-Sn and C in TiO₂@SnO_x@C composite are estimated to be 41.9, 15.1, 13.6, and 28.6% based on the following equations:

m(SnO₂) + m(TiO₂) = M₂ (1)

m(Sn_total)/m(Ti) = 0.83 (2)

m(β-Sn) ≈ (M₂ − M₁)/M(O₂) × 118 (3)

m(Sn_total) = m(β-Sn) + m(Sn of SnO₂) (4)

m(TiO₂) + m(SnO₂) + m(β-Sn) + m(C) = M₀ (5)

where the M(O₂) is molar mass of O₂, m(X) implies the mass of X, and M₀, M₁ and M₂ are the remaining mass during TGA test showed in Fig. 6.

Fig. 6 TG analysis of TiO₂@SnO_x@C.

Then, the N₂ adsorption/desorption test was conducted on TiO₂@SnO_x@C as shown in Fig. 7a. The N₂ adsorption/desorption isotherm can be identified as type IV, which is the characteristic isotherm of mesoporous materials. The pore sizes distribution is mainly distributed at ca. 3.9 and 11 nm, as exhibited in Fig. 7b. The BET surface area is measured to be 160 m² g⁻¹ based on the N₂ adsorption/desorption tests. It is believed that the mesoporous structure will bring TiO₂@SnO_x@C improved electrochemical performances due to the high porosity can increase the contact interface of electrode/electrolyte, partly buffer the huge volume change of SnO_x, accelerate the diffusion of lithium-ion and electron throughout the electrodes.

Fig. 7 (a) N₂ adsorption and desorption isotherms of TiO₂@SnO_x@C. (b) The distribution of pore sizes.

Undoubtedly, the electrochemical properties of TiO₂@SnO_x@C were investigated as anode materials for lithium-ion batteries. Fig. 8a shows the cyclic voltammograms curve (CV) of the initial cycle. Three cathodic peaks respectively located at 1.7, 1 and 0.3–0 V are observed and respectively corresponding to the lithium insertion into TiO₂, reduction of SnO₂ by lithium and formation of solid electrolyte interface (SEI) layer, and Sn alloying with lithium. The reduction of SnO₂ by lithium and formation of solid electrolyte interface (SEI) layer are the major reasons of larger initial irreversible capacity happened. Followed by two anodic peaks respectively located at 2.1 and 0.6 V are observed and assigned to the lithium extraction from TiO₂ (TiO₂ + xLi⁺ + xe⁻ ↔ Li_xTiO₂ (x ≤ 1)) and Sn de-alloying with lithium, respectively. It is worth noting that the pair redox peaks located at 1 (cathodic) and 1.26 (anodic) V can be assigned to the redox reaction between SnO₂ and lithium, which meaning the partly reversible reaction of SnO₂ and lithium, and hence improved lithium storage properties.^23,24 Thus, we believe that the observed partial reversibility of reaction between SnO₂ and lithium is due to the small SnO₂ nanoparticles sandwiched between TiO₂ and carbon coating layer (Fig. 3d). Fig. 8b shows the galvanostatic discharge/charge profiles for the 1st, 2nd, 20th, 50th, 100th, 200th and 300th cycles at a current density of 200 mA g⁻¹ between 0.01 and 3 V. A plateau potential and two slope potential appear in first discharge, and respectively corresponding to lithium insertion into TiO₂, reduction of SnO₂ by lithium, and Sn alloying with lithium, which is good agreement with the results of CVs. The first discharge and charge capacities are 1285.2 and 732.2 mA h g⁻¹, respectively, coupled with a coulombic efficiency of 57%, the initial irreversible capacity is ascribed to the reduction SnO₂ by lithium and the formation of SEI layer.²⁵ The discharge capacity at 2nd, 20th, 50th, 100th, 200th, and 300th is 743.9, 522.7, 496.4, 503.4, 569 and 629 mA h g⁻¹, respectively, displaying a slowly decreased trend until to 50th cycle and then fast increases. Fig. 8c and d show the long-term cycling performances of TiO₂@SnO_x@C at 200 and 1000 mA g⁻¹ between 0.01 and 3 V, respectively. Obviously, the TiO₂@SnO_x@C exhibits a superior cycling performance, a high capacity of 629 mA h g⁻¹ is remained even after 300 cycles at 200 mA g⁻¹, and the coulombic efficiencies are maintained at above 98% after 20 cycles, as shown in Fig. 8c. Moreover, a reversible capacity of 490.3 is obtained at 1000 mA g⁻¹ even after 1000 cycles and is much higher than theoretical capacity of graphite (372 mA h g⁻¹), as shown in Fig. 8d. The superior electrochemical performances make TiO₂@SnO_x@C a potential anode material for LIBs with long cycle life and high rate capability. The phenomenon of capacities first decrease and then quickly increase is also reported in other metal oxide materials and can be attributed to reversible formation of a polymeric gel-like layer via electrolyte decomposition, activation of active materials, and the electrocatalytic reversible conversion of some components of SEI films after long-term cycling.^26–28 In addition, the large specific surface area of 160 m² g⁻¹ can bring TiO₂@SnO_x@C extra surface lithium storage such as surface defects and voids. And the electrolyte can penetrate into active materials deeper after long cycling, which finally results in capacity increase. For demonstrating the superior cycling performance and high lithium storage of TiO₂@SnO_x@C were benefited from the effectively complementary and synergic effect among TiO₂, SnO_x and carbon coating layer, two control samples of TiO₂@SnO₂ HS and SnO₂ HS (Fig. 1d and 2)¹⁶ were prepared and tested as anode materials at the same conditions with TiO₂@SnO_x@C (at 200 mA g⁻¹ between 0.01 and 3 V), as shown in Fig. 8e. Comparing with TiO₂@SnO_x@C, both the two control samples show inferior cycling performances. The SnO₂ HS shows the worst cycling performance among the three samples, accounting for a discharge capacity drop from 1448.5 mA h g⁻¹ at the first cycle to 69.9 mA h g⁻¹ at the 200th cycle quickly. Although the TiO₂@SnO₂ HS shows a better cycling performance compared to SnO₂ HS but worse than TiO₂@SnO_x@C. The first discharge capacity of TiO₂@SnO₂ HS is 1342.2 mA h g⁻¹, but only 236.4 mA h g⁻¹ remained after 200 cycles. The morphologies of three samples were also studied by TEM measurement after different cycles, as shown Fig. 9. After 116 cycles, the morphology of TiO₂@SnO_x@C is still preserved well as shown in Fig. 9a, meaning the good structural stability during cycling process. The other two control samples, however, the original morphologies are lost seriously after 20 cycles.

Fig. 8 (a) Cyclic voltammogram of TiO₂@SnO_x@C at a scan rate of 0.5 mV; (b) galvanostatic discharge/charge profiles of different cycles for TiO₂@SnO_x@C at 200 mA g⁻¹; (c and d) cycling performances of TiO₂@SnO_x@C at 200 and 1000 mA g⁻¹, respectively; (e) cycling performances of different samples at 200 mA g⁻¹; (f) rate capability of TiO₂@SnO_x@C.

Fig. 9 TEM images of TiO₂@SnO_x@C after 116 cycles (a), TiO₂@SnO₂ HS after 20 cycles (b), and SnO₂ HS after 20 cycles (c).¹⁶

Moreover, electrochemical impedance spectroscopy (EIS) measurements were carried out on three samples to better demonstrate the synergic effect of components on the electronic conductivity of the TiO₂@SnO_x@C. Fig. 10 shows the Nyquist plots of the three samples after 5 cycles at 200 mA g⁻¹. All electrodes show Nyquist plots composing of a depressed semicircle at high frequency range and an angled line in the low frequency range. The diameter of the depressed semicircle is correlated with the electron transfer resistance on the electrode interface, and the angled line is related to a diffusion controlled process. Obviously, the TiO₂@SnO_x@C exhibits much smaller diameter of the high frequency semicircles than other two control samples, confirming the outermost carbon coating layer plays a key role in improving electric conductivity. Moreover, more rough and incompact surface of TiO₂@SnO₂ HS brings better electric conductivity than SnO₂ HS (compact shell), which is well reflected in Nyquist plots as shown in Fig. 10. Comparing a series of characterizations of three samples, it is confirmed that the improved cycling stability and lithium storage properties of the TiO₂@SnO_x@C are attributed to effectively complementary and synergic effect among TiO₂, SnO_x and carbon coating layer. The excellent rate property is another indispensable capability for electrodes as advanced LIBs. Fig. 8f displays the rate capability of TiO₂@SnO_x@C at different current densities from 400 to 3000 mA g⁻¹. As expected, the capacity decreases gradually as the current rate increases. The reversible capacity in fifth cycle is 569.7, 435.8, 378.7, 349.4, and 302.5 mA h g⁻¹ when cycles at 400, 1000, 1600, 2000, and 3000 mA g⁻¹, respectively. And it can be clearly seen that the TiO₂@SnO_x@C keep a stable cycle performance at each current density. When the current density finally backs to 400 mA g⁻¹ after rate test, a capacity of 449.1 mA h g⁻¹ can be restored at 60th cycle, confirming the good rate performance and stability of the TiO₂@SnO_x@C composite.

Fig. 10 The Nyquist plots of different samples.

4 Conclusions

In summary, a hybrid structure of TiO₂@SnO_x@C is successfully fabricated based on rational design. When used as anode material for lithium-ion battery, the TiO₂@SnO_x@C exhibits a good electrochemical performance, delivering a capacity of 629 mA h g⁻¹ at 200 mA g⁻¹ after 300 cycles, and a reversible capacity of 490.3 mA h g⁻¹ at 1000 mA g⁻¹ even after 1000 cycles. By comparing with control samples, it is suggested that the superior cycling performance and high lithium storage of TiO₂@SnO_x@C should be benefited from the effectively complementary and synergic effect among structural stability of TiO₂, high theoretical capacity of SnO_x, and good conductive and flexible ability of carbon due to four merits: (1) the high theoretical capacity of SnO_x can contribute to the desirable reversible capacity of TiO₂@SnO_x@C; (2) the conductive and flexible carbon coating can serve as a current collector to maintain the high electron transfer pathways within the electrode and partly buffer the volume variation of SnO_x; (3) the structural stability of TiO₂ can serve a stable support to maintain the structural integrity of electrode and prevent the SnO₂ from aggregation; (4) hollow, mesoporous, and sandwich structure supported by TiO₂ and carbon coating layer can effectively suppress the pulverization and aggregation of SnO_x, and further strengthen the stability of electrode structure during long-term cycling process.

Acknowledgements

The authors are grateful for financial support from the National Natural Science Foundation of China (grant nos 21103108 and 21173148) and materials characterization from Instrumental Analysis Center of Shanghai Jiao Tong University.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c5ra04629g

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