Core/shell TiO₂–MnO₂/MnO₂ heterostructure anodes for high-performance lithium-ion batteries

Abstract

Core/shell TiO₂–MnO₂/MnO₂ heterostructures were synthesized by combining an electrospinning technique with a hydrothermal reaction. To create the starting materials, porous TiO₂–carbon nanofibers were first prepared using a simple electrospinning technique followed by calcination. The porous structure in TiO₂–carbon nanofibers caused by the partial decomposition of polystyrene is beneficial to the diffusion of KMnO₄ from the outer surface into inner fibers to completely react with carbon and produce MnO₂ nanosheets. Some MnO₂ nanosheets in the TiO₂ core connect with other MnO₂ nanosheets surrounding the TiO₂ core to form core/shell TiO₂–MnO₂/MnO₂, which can enhance the stability of the structure. The large surface area of the resulting materials offers a sufficient electrode–electrolyte interface to promote the charge-transfer reactions, which yields a better rate capability. The porous structure of TiO₂–MnO₂/MnO₂ nanofibers not only facilitates Li-ion access, but also accommodates large volumetric expansion during the charging–discharging processes, resulting in an excellent cycle performance. As an anode, this material delivered a high reversible capacity of 891 mA h g⁻¹ at the first cycle and maintained the capacity of 888 mA h g⁻¹ after 50 cycles at the current density of 0.1 A g⁻¹; it also showed a remarkable rate capability of 2 A g⁻¹ while retaining a capacity of 185 mA h g⁻¹ after 500 cycles. Given their enhanced electrochemical performance, core/shell TiO₂–MnO₂/MnO₂ heterostructure nanofibers are promising anode candidates for lithium-ion batteries.

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Introduction

Nanostructured heterostructures play an important role in electrochemical energy storage devices such as lithium-ion batteries (LIBs) and supercapacitors, as they can be given diverse properties by designing their composition and morphology, and reassembling the basic primary building units.^1–3 Thus, heterostructures not only inherit the fascinating properties of the original building blocks, but they are also superior to their individual building units. For instance, a high capacity anode for LIBs has been fabricated by coating carbon nanohorns (CNHs) with nanoflaky MnO₂, combining the advantages of the good electrical conductivity and large surface area of CNHs, and the short path length of nanoflaky MnO₂.⁴ Additionally, 3D NiO–graphene nanosheet composites show enhanced electrochemical performance in LIBs due to the synergetic effect of NiO, graphene, and their distinct structure.⁵

Titanium dioxide (TiO₂) has been widely investigated as an anode material for LIBs due to its safety, low cost, and environmental friendliness. Moreover, its structural characteristics endow TiO₂ with a long cycle life.^6–8 However, low theoretical capacities (335 mA h g⁻¹) and low electrical conductivity have greatly hampered its practical application. Various strategies have been developed to address these problems, such as designing one-dimensional (1D) nanostructured TiO₂ can improve electronic conductivity due to the shorter electron transport along the 1D geometry;^9–12 coating TiO₂ with carbonaceous materials or a noble metal to increase the conductivity;^13–16 and incorporating metal oxides to enhance its capacity. Among these metal oxides, manganese dioxide (MnO₂) has been recognized as one of the intensively investigated metal oxides due to its high theoretical capacity of 1230 mA h g⁻¹, relatively low electrochemical motivation force, and natural abundance.^17–20

However, the implementation of MnO₂ for LIBs application has been greatly hampered by the large volume change that occurs during the lithiation/delithiation processes, which results in pulverization that breaks the electrical connection from the current collector, thus leading to the poor cyclability.

In this study, we successfully synthesized novel core/shell TiO₂–MnO₂/MnO₂ heterostructures. MnO₂ nanosheets were produced by the in situ chemical oxidation–reduction reaction between KMnO₄ and carbon in TiO₂–carbon nanofibers. In addition, MnO₂ materials can exist in TiO₂ matrix, and connect with the outside MnO₂ nanosheets surrounding TiO₂ core to form a core/shell TiO₂–MnO₂/MnO₂ hybrid, which can enhance the stability of hybrid during charging–discharging process. The core/shell heterostructures incorporated the electrochemical function of the individual components such as the exceptional stability of TiO₂ and the high capacity of MnO₂. As a result, the prepared materials manifested excellent electrochemical performance as anodes, delivering a high reversible capacity of 891 mA h g⁻¹ at the first cycle and maintaining the capacity of 888 mA h g⁻¹ after 50 cycles at 0.1 A g⁻¹; they also showed a good rate capability of 2 A g⁻¹ while retaining a capacity of 185 mA h g⁻¹ after 500 cycles.

Experimental

Materials

The polystyrene (PS, M_w ∼ 280 000), tetraisopropyl titanate (Ti(OiPr)₄) and dimethylformamide (DMF) were purchased from Aldrich. The potassium permanganate (KMnO₄) was bought from the Guang Zhou Dong Hong Chemical Plant (Guangzhou, China).

Preparation of the core/shell TiO₂–MnO₂/MnO₂ heterostructures

The core/shell TiO₂–MnO₂/MnO₂ heterostructures were prepared with the electrospinning and hydrothermal methods. In the electrospinning process, 5 mL of Ti(OiPr)₄ and 4 mL of acetic acid were mixed in 13 wt% PS solution (DMF as solvent). The prepared mixture was loaded into a 20 mL syringe. A high voltage of 17 kV was applied to the syringe needle tip. The flow rate of the fluid was set to 0.04 mm min⁻¹. The distance between the needle and collector was 20 cm (KATO Tech Co., Ltd). The as-electrospun Ti(OiPr)₄–PS nanofibers were subjected to calcination at 450 °C in Ar for 3 h to obtain porous TiO₂–carbon nanofibers. Specifically, 0.05 g of the obtained porous TiO₂–carbon nanofibers was put into 50 mL of 0.03 M KMnO₄ solution with strong stirring, which was then transferred into a 60 mL Teflon-lined autoclave and subsequently heated at 160 °C for 5 h. The resulting solid sample was washed by deionized water and ethanol and then dried at 80 °C in a vacuum.

Characterization

The crystal structure was characterized by powder X-ray diffraction (XRD, Rigaku). Scanning electron microscopy (SEM; JEOL 6300F, 5.0 kV) and transmission electron microscopy (TEM; JEOL 2100F, 220 kV) equipped with an energy dispersive X-ray spectrometer (EDS, Oxford) and an Electron Energy Loss Spectroscopy (EELS) spectrometer (Gatan, Enfina) were used to determine the morphology, chemical composition and element distribution of the obtained materials. X-ray photoelectron spectroscopy was carried out on a Perkin-Elmer model PHI 5600 system. The analysis of the surface area was performed by N₂ adsorption and desorption isotherms at 77 K based on the Brunauer–Emmett–Teller method (BET, Micromeritics ASAP2020) and the pore-size distribution was calculated using the density functional theory (DFT) method.

Electrochemical characterization

The working electrode for the electrochemical measurements was produced by mixing 80 wt% active materials, 10 wt% carbon black, 10 wt% polyvintlidene fluoride (PVDF) binder, and some N-methylpyrrolidone (NMP) into a slurry. The slurry was pasted uniformly onto a copper foil current collector and was dehydrated in a vacuum oven at 120 °C for 12 h. The cell consisted of a working electrode, metallic lithium as a counter/reference electrode, Celgard 2400 film as a separator film, and an electrolyte of 1 M LiPF₆ in ethylene carbonate (EC)–diethyl carbonate (DEC) (1 : 1 in volume). The charge–discharge tests were performed using a LAND 2001 CT battery testing system at different current densities with a voltage window of 0–3 V. Cyclic voltammetric (CV) was carried out on a CHI 660C electrochemical workstation.

Results and discussion

Synthesis and characterization of core/shell TiO₂–MnO₂/MnO₂ heterostructures

The novel core/shell TiO₂–MnO₂/MnO₂ heterostructures were synthesized by combining an electrospinning technique with a hydrothermal reaction, as shown in Fig. 1. The porous TiO₂–carbon nanofibers were first produced by calcinating the electrospun Ti(OiPr)₄–PS nanofibers. Subsequently, the oxidation–reduction reaction occurred between carbon and KMnO₄ during the hydrothermal process. The porous structure in TiO₂–carbon nanofiber is beneficial to the diffusion of KMnO₄ from outer surface into inner fibers to thoroughly contact with carbon, when the small nanocrystalline MnO₂ is formed on the TiO₂ nanofiber owing to the redox reaction of eqn (1). Then, MnO₂ nanocrystalline continuously breeds to form the final MnO₂ nanosheet according to eqn (2).^21–23

4KMnO₄ + 3C + H₂O → 4MnO₂ + K₂CO₃ + 2KHCO₃ (1)

4KMnO₄ + 2H₂O → MnO₂ + 4KOH + 3O₂ (2)

Fig. 1 Schematic illustration of the preparation of the core/shell TiO₂–MnO₂/MnO₂ heterostructure.

Fig. 2a shows the XRD patterns of the core/shell TiO₂–MnO₂/MnO₂ heterostructures, confirming the presence of anatase TiO₂ (JCPDS 21-1272) and birnessite-type MnO₂ (JCPDS 42-1317). The elemental compositions of the prepared materials were collected by XPS. The peaks of O 1s, Ti 2p, Mn 2p, Mn 3p3, and K 2p are shown in Fig. S1,† where a trace of K was introduced during the hydrothermal reaction,^24,25 which is consistent with the results of the EELS in Fig. S2† and the EDS shown in Fig. S3.† Additionally, the mole ratio between MnO₂ and TiO₂ is about 2.2 displayed in Table S1.† Fig. 2b exhibits the Mn 2p XPS spectrum. The binding energies of 642.1 eV and 654.0 eV, corresponding to the Mn 2p_3/2 and the Mn 2p_1/2, indicate the presence of tetravalent manganese. Furthermore, the difference in their binding energies was 11.9 eV, which is in excellent agreement with the reported binding energy of MnO₂.^26,27 The valence state was estimated by measuring the branching ratio in the EELS spectrum. A typical EELS spectrum of Mn-L edge for the as-obtained TiO₂–MnO₂/MnO₂ heterostructures is shown in Fig. 2c. The branching ratio (I(L₃)/I(L₂)) for TiO₂–MnO₂/MnO₂ heterostructures was about 1.5, which agrees with the value for MnO₂.²⁸ Therefore, the valence state of the as-prepared materials was +4, matching well with the results of the XRD and XPS. The N₂ adsorption–desorption isotherms and corresponding pore-size distribution, shown in Fig. 2d, were calculated by the DFT method. The type IV isotherms with a type H3 hysteresis loop shown in Fig. 2d, indicates the existence of relatively large size pores.²⁹ It is clear that the core/shell TiO₂–MnO₂/MnO₂ heterostructures possess mesopores and macropores as shown in the inset of Fig. 2d, which were mainly the result of the packing of ultrathin MnO₂ nanosheets. As expected, core/shell heterostructures provide a relatively high BET specific surface area of 127 m² g⁻¹ with a total pore volume of 0.62 m³ g⁻¹.

Fig. 2 (a) XRD patterns of TiO₂–MnO₂/MnO₂ heterostructures, (b) XPS narrow scan spectra of Mn 2p for TiO₂–MnO₂/MnO₂ heterostructures, (c) EELS spectra of Mn-L edge, and (d) N₂ adsorption–desorption isotherms of TiO₂–MnO₂/MnO₂ heterostructures and corresponding pore-size distribution.

The SEM image in Fig. 3a shows that the obtained pristine TiO₂–carbon nanofibers with diameters ranging from 200 to 500 nm as a result of the calcination of electrospun Ti(OiPr)₄–PS nanofibers had no secondary nanostructures. As shown in Fig. 4a, highly porous structure caused by partial decomposition of PS are observed in TiO₂–carbon nanofibers.⁸ This porous structure is beneficial to the diffusion of KMnO₄ from outer surface into inner fibers to completely react with carbon and produce MnO₂ nanosheets. As shown in Fig. 3b, the fiber diameter became larger after the thermal reaction. Fig. 3c is a typical TEM-SEI image of an individual hybrid nanostructure. The TiO₂ nanofiber was decorated by many thin MnO₂ nanosheets with thicknesses of ∼4 nm, which accords with the TEM image in Fig. 4c and S4.† Fig. 3d is the magnification of Fig. 3c. It is obvious that the thin MnO₂ nanosheets grew on the TiO₂ nanofibers are similar to the structure of carambola, which had distinctive ridges running down its sides. The high-resolution TEM (HRTEM) image in Fig. 4b shows that the materials consist of carbon and TiO₂. The measured interplanar spacing of 0.35 nm corresponded to the d-spacing of (101) plane of anatase TiO₂ for TiO₂–carbon and TiO₂–MnO₂/MnO₂. Additionally, the inter-planar distance of ∼0.67 nm is in excellent agreement with the (001) plane of birnessite-type MnO₂ shown in Fig. 4d.^30,31 Furthermore, the ring-like pattern of the TiO₂–MnO₂/MnO₂ heterostructures in the selected area for electron diffraction (inset in Fig. 4d) further confirms the presence of anatase TiO₂ and birnessite MnO₂. The EDS line scanning analysis shown in Fig. 5 also demonstrates that pure core/shell TiO₂–MnO₂/MnO₂ heterostructures were prepared and MnO₂ nanosheets distributed uniformly on the TiO₂ matrix.

Fig. 3 (a) SEM image of porous TiO₂–carbon nanofibers. (b) SEM and (c and d) TEM-SEI images of core/shell TiO₂–MnO₂/MnO₂ heterostructures.

Fig. 4 (a and b) TEM and HRTEM images of TiO₂–carbon nanofiber and (c and d) TEM and HRTEM images of the core/shell TiO₂–MnO₂/MnO₂ heterostructure.

Fig. 5 (a) Dark-field TEM image and (b–d) corresponding scanning line of core/shell TiO₂–MnO₂/MnO₂ heterostructures.

Electrochemical performance

The investigation of the electrochemical performance of the core/shell TiO₂–MnO₂/MnO₂ heterostructures for lithium ions storage is illustrated in Fig. 6. The initial three charge–discharge cycle profiles for the prepared free TiO₂, MnO₂, and TiO₂–MnO₂/MnO₂ heterostructures with a current rate of 0.1 A g⁻¹ are shown in Fig. 6a–c. The initial charge capacities were 142 mA h g⁻¹ for TiO₂, 330 mA h g⁻¹ for MnO₂, and 891 mA h g⁻¹ for the TiO₂–MnO₂/MnO₂ electrode, respectively, indicating that the increased capacity for TiO₂–MnO₂/MnO₂ heterostructures was not only from mechanical amalgamation, but an inter-enhancement effect of these two oxides. The large capacity loss in the first cycle for these electrodes was caused by the formation of a solid electrolyte interface layer (SEI) on the electrode surface due to the electrolyte decomposition.^32–34 Fig. 6b and c clearly show a voltage plateau at ∼0.45 V in the first discharge, reflecting the conversion reaction of MnO₂ to Mn metal,²⁹ MnO₂ + 4Li⁺ + 4e⁻ ↔ Mn + 2Li₂O. The charge curve shows a sloping plateau at ∼1.2 V caused by the reversible nature of this electrochemical reaction. Fig. 6d shows the CV curve of the core/shell TiO₂–MnO₂/MnO₂ heterostructures for the first five cycles at a scan rate of 0.1 mV s⁻¹ in the voltage range from 0 to 3 V. The current peak at 0.3 V vs. Li⁺/Li in the first cathodic scan can be attributed to the complete reduction of Mn(IV) to Mn(0), the formation of Li₂O, and SEI film.^29,32,33 In the first anodic scan, the current peak located at 1.2 V was ascribed to the reversible oxidation of Mn(0) to Mn(IV). From the third scans, one peak appeared at 0.4 V with declined intensity, while the anodic peak remained unchanged, suggesting the formation of SEI film in the first cycle and good reversibility of the electrochemical reaction.^34,35 Rate capability is another important factor in the practical application of LIBs. The rate performance of the core/shell TiO₂–MnO₂/MnO₂ heterostructures was studied at different current densities as shown in Fig. 6e. The core/shell electrode delivered capacities of 347, 212, 157, and 122 mA h g⁻¹ at rates of 1, 2, 3, and 4 A g⁻¹, respectively. Fig. 6f shows the cycling performance of the core/shell TiO₂–MnO₂/MnO₂ at 0.1 and 2 A g⁻¹ in the voltage range of 0–3 V. It delivered a reversible capacity of 888 mA h g⁻¹ at 0.1 A g⁻¹ after 50 cycles, which is much higher than that of free TiO₂ (70 mA h g⁻¹) and MnO₂ (137 mA h g⁻¹), as shown in Fig. S5.† When the current density increased to 2 A g⁻¹, a capacity of 185 mA h g⁻¹ remained after 500 cycles. Furthermore, the Coulombic efficiencies at both current densities remained higher than 98% after the first few cycles. The low initial coulombic efficiency of the anode is a detrimental problem when assembling full cell, which can be enhanced by creating a passivation layer on the surface of the active material.¹ The capacity also increased gradually with cycling, which is very common in most metal oxide anode materials.^36–38 The reasons for the increase of capacity after cycling could be ascribed to the reversible growth of a polymer gel-like film caused by kinetically activated electrolyte degradation and the activation of the active materials.^39–43 To further demonstrate the enhanced performance of core/shell TiO₂–MnO₂/MnO₂ prepared in this work, Table S2† compares the capacities of various MnO₂-based material electrodes. The high electrochemical performance of the core/shell nanomaterials was mainly due to the following properties: (1) the unique 1D core/shell heterostructure provides a better structure stability; (2) the large surface area of the resulting materials offers sufficient electrode–electrolyte interface to promote the charge-transfer reactions; (3) the porous structure not only facilitates Li-ion access, but also accommodates large volumetric expansion during the charging–discharging processes; (4) the MnO₂ active materials have intrinsically high capacities.

Fig. 6 Electrochemical properties of the core/shell TiO₂–MnO₂/MnO₂ heterostructure anodes for Li-ion batteries. (a–c) The first three charge–discharge curves of (a) TiO₂, (b) MnO₂, and (c) TiO₂–MnO₂/MnO₂ at 0.1 A g⁻¹. (d) CV curve of the core/shell TiO₂–MnO₂/MnO₂ at a scan rate of 0.1 mV s⁻¹. (e) The rate performance at various current densities and (f) cycling performance of core/shell TiO₂–MnO₂/MnO₂ heterostructure anodes at 0.1 and 2 A g⁻¹.

Conclusion

In summary, a facile combination of an electrospinning technique and a hydrothermal method was developed to synthesize core/shell TiO₂–MnO₂/MnO₂ heterostructures. The novel core/shell TiO₂–MnO₂/MnO₂ heterostructures provided sufficient electrode–electrolyte interface and buffer to alleviate the large strain caused by the conversion reaction of MnO₂, and were also beneficial to the Li-ion and electron transportation. As a result, the TiO₂–MnO₂/MnO₂ heterostructure anode showed a high reversible capacity of 888 mA h g⁻¹ at 0.1 A g⁻¹, excellent rate performance, and a long cycling lifetime. The core/shell TiO₂–MnO₂/MnO₂ heterostructure nanofibers are promising anode materials for LIBs.

Acknowledgements

The authors are grateful for the support received from the Research Grants Council of the Hong Kong Special Administration Region (grants: PolyU 5349/10E and PolyU 5312/12E) and the Hong Kong Polytechnic University (grants: G-YK47 and 1-BD08).

Notes and references

1. Y. M. Chen, X. Y. Li, K. Park, J. Song, J. H. Hong, L. M. Zhou, Y. W. Mai, H. T. Huang and J. B. Goodenough, J. Am. Chem. Soc., 2013, 135, 16280–16283.
2. Y. M. Chen, Z. G. Lu, L. M. Zhou, Y. W. Mai and H. T. Huang, Energy Environ. Sci., 2012, 5, 7898–7902.
3. Z. L. Xu, B. Zhang, Z. Q. Zhou, S. Abouali, M. A. Garakani, J. Q. Huang, J. Q. Huang and J. K. Kim, RSC Adv., 2014, 4, 22359–22366.
4. H. Lai, J. X. Li, Z. G. Chen and Z. G. Huang, ACS Appl. Mater. Interfaces, 2012, 4, 2325–2328.
5. L. Tao, J. Zai, K. Wang, Y. Wan, H. Zhang, C. Yu, Y. Xiao and X. Qian, RSC Adv., 2012, 2, 3410–3415.
6. J. F. Lei, K. Du, R. H. Wei, J. Ni, L. B. Li and W. S. Li, RSC Adv., 2013, 3, 13843–13850.
7. Z. Y. Weng, H. Guo, X. M. Liu, S. L. Wu, K. W. K. Yeung and P. K. Chu, RSC Adv., 2013, 3, 24758–24775.
8. X. Y. Li, Y. M. Chen, L. M. Zhou, Y. W. Mai and H. T. Huang, J. Mater. Chem. A, 2014, 2, 3875–3880.
9. S. W. Kim, T. H. Han, J. Kim, H. Gwon, H. S. Moon, S. W. Kang, S. O. Kim and K. Kang, ACS Nano, 2009, 3, 1085–1090.
10. F. X. Wu, Z. X. Wang, X. H. Li and H. J. Guo, J. Mater. Chem., 2011, 21, 12675–12681.
11. H. B. Wu, H. H. Hng and X. W. Lou, Adv. Mater., 2012, 24, 2567–2571.
12. J. M. Li, W. Wan, H. H. Zhou, J. J. Li and D. S. Xu, Chem. Commun., 2011, 47, 3439–3441.
13. H. Han, T. Song, E. Lee, A. Devadoss, Y. Jeon, J. Ha, Y. Chung, Y. Choi, Y. Jung and U. Paik, ACS Nano, 2012, 6, 8308–8315.
14. X. Zhang, P. S. Kumar, V. Aravindan, H. H. Liu, J. Sundaramurthy, S. G. Mhaisalkar, H. M. Duong, S. Ramakrishna and S. Madhavi, J. Phys. Chem. C, 2012, 116, 14780–14788.
15. H. W. Bai, Z. Y. Liu and D. D. Sun, J. Mater. Chem., 2012, 22, 24552–24557.
16. J. S. Chen, D. Y. Luan, C. M. Li, F. Y. C. Boey, S. Z. Qiao and X. W. Lou, Chem. Commun., 2010, 46, 8252–8254.
17. A. L. Reddy, M. M. Shaijumon, S. R. Gowda and P. M. Ajayan, Nano Lett., 2009, 9, 1002–1006.
18. Z. Xiao, Y. Xia, Z. H. Ren, Z. Y. Liu, G. Xu, C. Y. Chao, X. Li, G. Shen and G. R. Han, J. Mater. Chem., 2012, 22, 20566–20573.
19. Z. Q. Li, N. N. Liu, X. K. Wang, C. B. Wang, Y. X. Qi and L. W. Yin, J. Mater. Chem., 2012, 22, 16640–16648.
20. J. Z. Zhao, Z. L. Tao, J. Liang and J. Chen, Cryst. Growth Des., 2008, 8, 2799.
21. X. B. Jin, W. Z. Zhou, S. W. Zhang and G. Z. Chen, Small, 2007, 3, 1513–1517.
22. W. Xiao, D. L. Wang and X. W. Lou, J. Phys. Chem. C, 2010, 114, 1694–1700.
23. Y. Chen, Y. Zhang, D. S. Geng, R. Y. Li, H. L. Hong, J. Z. Chen and X. L. Sun, Carbon, 2011, 49, 4434–4442.
24. W. M. Chen, L. Qie, Q. G. Shao, L. X. Yuan, W. X. Zhang and Y. H. Huang, ACS Appl. Mater. Interfaces, 2012, 4, 3047–3053.
25. X. D. Zhao, H. M. Fan, J. Luo, J. Ding, X. Y. Liu, B. S. Zou and Y. P. Feng, Adv. Funct. Mater., 2011, 21, 184–190.
26. L. Q. Mai, F. Dong, X. Xu, Y. Z. Luo, Q. Y. An, Y. L. Zhao, J. Pan and J. N. Yang, Nano Lett., 2013, 13, 740–745.
27. H. Xia, M. O. Lai and L. Lu, J. Mater. Chem., 2010, 20, 6896–6902.
28. T. Riedl, T. Gemming and K. Wetzig, Ultramicroscopy, 2006, 106, 284–291.
29. W. Xiao, J. S. Chen, Q. Lu and X. W. Lou, J. Phys. Chem. C, 2010, 114, 12048–12051.
30. H. T. Zhu, J. Luo, H. X. Yang, J. K. Liang, G. H. Rao, J. B. Li and Z. M. Du, J. Phys. Chem. C, 2008, 112, 17089–17094.
31. G. P. Wang, L. Zhang and J. J. Zhang, Chem. Soc. Rev., 2012, 41, 797–828.
32. B. Zhang, Y. S. Liu, Z. D. Huang, S. W. Oh, Y. Yu, Y. W. Mai and J. K. Kim, J. Mater. Chem., 2012, 22, 12133–12140.
33. Y. M. Chen, Z. G. Lu, L. M. Zhou, Y. W. Mai and H. T. Huang, Nanoscale, 2012, 4, 6800.
34. L. Wang, Y. H. Li, Z. D. Han, L. Chen, B. Qian, X. F. Jiang, J. Pinto and G. Yang, J. Mater. Chem. A, 2013, 1, 8385–8397.
35. J. Y. Liao, D. Higgins, G. Lui, V. Chabot, X. C. Xiao and Z. W. Chen, Nano Lett., 2013, 13, 5467–5473.
36. P. L. Taberna, S. Mitra, P. Poizot, P. Simon and J. M. Tarascon, Nat. Mater., 2006, 5, 567–573.
37. Y. Fu, X. Li, X. Sun, X. Wang, D. Liu and D. He, J. Mater. Chem., 2012, 22, 17429–17431.
38. X. H. Wang, Z. B. Yang, X. L. Sun, X. W. Li, D. S. Wang, P. Wang and D. Y. He, J. Mater. Chem., 2011, 21, 9988–9990.
39. S. Grugeon, S. Laruelle, L. Dupont and J.-M. Tarascon, Solid State Sci., 2003, 5, 895–904.
40. S. Laruelle, S. Grugeon, P. Poizot, M. Dolle, L. Dupont and J. M. Tarascon, J. Electrochem. Soc., 2002, 149, A627–A634.
41. J. S. So and C. H. Weng, J. Power Sources, 2005, 146, 482–486.
42. Y. Le, Z. Y. Wang, L. Zhang, H. B. Wu and X. W. Lou, J. Mater. Chem. A, 2013, 1, 122–127.
43. Y. M. Chen, X. Y. Li, X. Y. Zhou, H. M. Yao, H. T. Huang, Y. W. Mai and L. M. Zhou, Energy Environ. Sci., 2014, 7, 2689–2696.

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Footnote

† Electronic supplementary information (ESI) available: TEM, EDS, XPS, EELS, and electrochemical performance. See DOI: 10.1039/c4ra06981a

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