Green synthesis of 3D SnO₂/graphene aerogels and their application in lithium-ion batteries

Abstract

Three-dimensional (3D) tin oxide/graphene aerogels (SnO₂/GAs) were constructed by a simple, facile and environmentally friendly process. The small-sized SnO₂ nanoparticles (6 nm) are encapsulated within graphene-based aerogels with interconnected 3D networks for the SnO₂/GAs nanocomposite. When used as an anode material in lithium ion batteries, it delivers a high reversible capacity that is close to the theoretical capacities of SnO₂ and graphene after 50 cycles. TEM observations of the samples before and after 50 cycles illustrate that the structures of the graphene network and SnO₂ NPs are preserved, which explains well the good cyclic stability of the electrode. The excellent electrochemical performance of the nanocomposites can be explained by their unique 3D porous architecture and the combination of the advantages of both SnO₂ and graphene in Li ion storage and transport.

---

Introduction

As one of the most popular electrochemical energy storage devices, rechargeable lithium-ion batteries (LIBs) have been found to be a great success as power sources in portable electronics for their excellent characteristics such as high energy density, light weight, environmental benignity and no memory effect.^1–4 Graphite is the widely used material for commercial LIBs, but its limited specific capacity (372 mA h g⁻¹ in theory) cannot satisfy the increasing demand for batteries with higher energy density.^5,6 Up to now, various advanced anode materials have been studied to replace graphite. Among them, SnO₂ has been proposed as one of the most promising anode candidates for next-generation lithium ion batteries, owing to its high theoretical capacity (782 mA h g⁻¹) and low discharge potential (<1.5 V).^7–9 However, the practical use of SnO₂ is greatly hampered by the poor cycling performance arising from its low electronic conductivity, severe aggregation, and huge volume changes upon extended Li insertion/extraction.^10,11 To solve these problems, various nanostructured SnO₂ and SnO₂ composites have been proposed with the aim of stabilizing the active electrode material by minimize the volume change during cycling.^12–26 Up to now, reducing the size of SnO₂ to the nanometer range and incorporating electrochemically active/inactive matrix have been widely applied for improving the cycling stability of SnO₂-based anodes.^12–26 The fabrication of SnO₂/carbon (e.g., graphene, amorphous carbon, carbon nanotubes, etc.) nanocomposites has been extensively investigated, because the carbon may both accommodate the volume variation and provide an electrical conduction pathway.^16–26

According to recent studies, graphene is regarded as a superior carbon matrix to create SnO₂/graphene hybrids due to its excellent electrical conductivity, high carrier mobility and large surface area.²⁷ Several SnO₂/graphene hybrids with various assemblies such as directly decorated SnO₂/graphene, SnO₂/C-graphene, and sandwich-like graphene-supported hybrids have been developed.^19–26 Although improved electrochemical performances have been achieved in these hybrid anodes, the performance is still limited by the following obstacles. (1) The graphene sheets tend to irreversible aggregation or restacking due to the strong van der Waals interactions among individual graphene sheets, resulting in a seriously reduced active surface area and carrier mobility.²⁸ The conventional direct high-temperature calcination could easily result in severe agglomeration of the graphene sheets; (2) in some SnO₂/graphene hybrids, the exposed active SnO₂ nanostructures, which are simply decorated on the surface of graphene sheets, are hard to avoid volume expansion and aggregation during the cycle processes due to the non-intimate contact between graphene nanosheets and active materials;²⁹ (3) the shape, size and dispersion uniformity of SnO₂ nanoparticles which significantly influence the electrochemical performance of SnO₂/graphene composites are still difficult to control. Therefore, it is urgent to develop new SnO₂/graphene hybrids with unique structures to overcome these obstacles for higher electrochemical performances.

Three-dimensional (3D) macroscopic frameworks of graphene sheets could effectively prevent restacking of graphene, provide large specific surface area, porous structure, and fast electron conductivity due to the continuous graphene backbone, which might be used as ideal support for the nanoparticles and thus promising for higher electrochemical performances.^30–33 Moreover, in such geometric architectures, nanoparticles could confined within graphene layers which could enhance their interface interaction and mitigate the volume expansion and aggregation of nanoparticles, promoting the Li⁺ storage property.^33–39 For example, Chen et al. pioneered the capture of prepared Fe₃O₄ into 3D graphene network which delivered a high capacity of 730 mA h g⁻¹ even at a current density of 1600 mA h g⁻¹.³³ Besides, 3D Fe₂O₃/GAs,³⁴ TiO₂/GAs,³⁵ CoO/GAs,³⁶ and Co₃O₄/GAs³⁷ also demonstrated the effectiveness of macroscopic 3D architectures to the final lithium storage properties. These findings motivated our idea of synthesize 3D SnO₂/GAs as anode materials in lithium ion batteries. During our research process, Wang et al. reported their work on 3D N-doped SnO₂/GAs.³⁸ They successfully constructed SnO₂/GAs via a solvothermal process (high temperature) using DMF as reducing agent, and they found that the as prepared material exhibits superior rate capability. It is well known that, exploring facile and environmental friendly synthetic methods for function materials is a significant issue for practical industrial applications. It thus inspired us to find a simple and environmental friendly synthetic process for 3D SnO₂/GAs. From the literature, 3D GAs can be successfully constructed using nontoxic sodium ascorbate as reducing agent.³⁹ We believe that it is possible to find a promising eco-friendly approach to synthesize SnO₂/GAs.

Herein, we report a simple, facile and environmental friendly method for the preparation of 3D SnO₂/graphene aerogels (SnO₂/GAs). The 3D SnO₂/GAs are synthesized by a two step procedure, for which SnO₂ NPs were firstly prepared and then directly decorated on and between graphene sheets under mild conditions (95 °C) with nontoxic sodium erythorbate as reducing agent. This allows us to control the size and shape of SnO₂ NPs which would influence the electrochemical performance. The mild synthesis condition could further prevent agglomeration of the graphene sheets. Therefore, such prepared 3D SnO₂/GAs could overcome the disadvantages of SnO₂/graphene hybrids, and expect to exhibit better electrochemical performances for LIBs. Moreover, this synthetic method for SnO₂/GAs could extend to other 3D graphene hybrids.

Experimental

Materials preparation

Synthesis of SnO₂ nanoparticles (SnO₂ NPs). Typically, lysteine (Lys, C₆H₁₄N₂O₂, >99% purity, 10 mmol) was dissolved in 20 mL of deionized water to form solution A. SnCl₄·5H₂O (Analytic reagent, AR, 2 mmol) was dissolved in 20 mL of deionized water to form solution B. And then solution A was added into solution B with stirring for 12 h at room temperature. After that, the mixture was sealed into a 50 mL Teflon-lined autoclave, heated and maintained at 180 °C for 10 h. After cooled down to room temperature naturally, the products were collected and washed with deionized water and then absolute alcohol, and then dried at 60 °C for 4 h.

Synthesis of 3D graphene aerogels (GAs). Graphene oxide was synthesized from natural graphite by a modified Hummer's method.⁴⁰ 60 mg of GO was dissolved in 60 mL of distilled water with ultrasonic treatment for 4 h to form a suspension. Then, 120 mg sodium erythorbate was added into the suspension under magnetic stirring. The mixture was heated at 95 °C for 3 h under atmosphere pressure without stirring to form graphene hydrogels. The reduced graphite oxide hydrogels were cleaned by immersing in water for many times. The pure graphene hydrogels obtained were freeze-dried to prepare 3D graphene aerogels (GAs).

Synthesis of SnO₂/GAs. Composites were obtained by following the same procedure of GAs. In brief, 120 mg SnO₂ were dispersed to 60 mL 1 mg mL⁻¹ GO suspension by ultrasonication. 120 mg sodium erythorbate was added into the suspension under magnetic stirring, and SnO₂/graphene hydrogels were synthesized at 95 °C for 3 h under atmosphere pressure without stirring. The SnO₂/graphene hydrogels were immersed in water to remove impurities. Then the hydrogels were freeze-dried to get SnO₂/GAs.

Electrochemical measurements

Electrochemical experiments were carried out using 2032-type coin cells. A metallic Li foil was used as the anode electrode. The cathode electrode was composed of 70 wt% active material, 15 wt% carbon black conductive additive, and 15 wt% poly(vinylidene fluoride) binder dissolved in N-methyl-2-pyrrolidone. The cathode slurry was coated on Cu foil. Each electrode was 8 × 8 mm² in size with a SnO₂/GAs mass loading of ∼3 mg cm⁻². The cathode and anode electrodes were separated by the Celgard 2320 membrane. The electrolyte used was 1 mol L⁻¹ lithium hexafluorophosphate dissolved in ethylene carbonate (EC) and dimethyl carbonate (DMC) [EC : DMC = 1 : 1 (w/w) ratio]. Galvonostatic charge–discharge cycling was performed on a Land-2100 battery tester (Wuhan, China). EIS was performed on a Bio-Logic VSP multichannel potentiostatic–galvanostatic system (France). The impedance data were recorded by applying an alternating-current voltage of 5 mV in the frequency range from 1 MHz to 5 mHz.

Results and discussion

The 3D SnO₂/GAs were fabricated by a simple, fast and environmental friendly method described above. SnO₂ nanoparticles (see Fig. 1a), with an average diameter of 6 nm, uniformly disperse on the surface of GO when they were mixed together (see Fig. 1b). Noteworthy, even after a long time of sonication, SnO₂ NPs still stably anchored on GO, demonstrating the strong interaction between them. High-resolution TEM (Fig. 1c and d) revealed that well crystalline SnO₂ NPs were encapsulated by graphene sheets. Fig. 2a and b show the SEM images of SnO₂/GAs. The images revealed that an interconnected, porous 3D graphene framework with continuous macropores is formed in the composite. Graphene sheets decorated with SnO₂ NPs acted as a building block and self-assembled into a 3D interconnected network, driven by π–π stacking interactions due to the decrease of oxygenated groups on graphene. The pore size of the network is distributed in the range of submicrometer to several micrometers and the walls consist of several layers of stacked graphene sheets. Apart from the decoration of SnO₂ NPs on both sides of the graphene sheets, a large part of the NPs are encapsulated within the graphene layers, suggesting efficient assembly between the SnO₂ NPs and the graphene sheets. Such a geometric confinement of NPs within graphene layers has been reported in some other metal oxides, which is supposed to enhance their interface contact and to suppress the dissolution and agglomeration of NPs, thereby promoting the electrochemical activity and stability of the hybrids.^41,42

Fig. 1 (a) TEM image of SnO₂ NPs; (b) TEM image of SnO₂/GAs; (c and d) HRTEM images of SnO₂/GAs.

Fig. 2 SEM images of SnO₂/GAs.

Fig. 3a shows XRD patterns of as-prepared SnO₂/GAs nanocomposite and graphene aerogels. All diffracted peaks from SnO₂ and can be well indexed by a tetragonal phase with space group P42/mnm (JCPDS Card no: 8 8-0287). The broad diffracted peaks of the SnO₂/GAs composite should be caused by the nano-sized effect of the formed SnO₂ nanoparticles. A weak and broad diffracted peak in the XRD pattern of GAs can also be seen in Fig. 3a. The presence of this peak indicates that graphene sheets stacked into multilayers and aggregated to some extend. In contrast, no visible diffracted peak from stacked graphene sheets can be identified in the pattern of SnO₂/GAs, indicating that SnO₂ NPs were efficiently deposited on the graphene surface, suppressing the stacking of graphene layers.

Fig. 3 (a) XRD patterns of SnO₂/GAs and GAs; (b) Raman spectra of GAs and SnO₂/GAs; (c) XPS C1s spectra of SnO₂/GAs and GAs; (d) XPs Sn 3d spectra of SnO₂/GAs and SnO₂.

Raman spectroscopy is a convenient and nonvolatile technique for the characterization of graphitic materials. Fig. 3b shows the recorded Raman spectra of GAs and SnO₂/GAs. The D band and G band in both spectra are related to the defects/disorder in the graphitic layers and stretching vibration of sp² C in a graphene layer, respectively.⁴³ The intensity ratio I_D/I_G of SnO₂/GAs (∼1.26) is higher than that of the GAs (∼0.92). The result suggests a decrease of the sp² C upon reduction of the exfoliated GO.^43,44

XPS was used to characterize the chemical information of the samples. Fig. 3c shows the high-resolution spectra of C1s in GAs and SnO₂/GAs samples. The C1s XPS spectrum of the GAs sample can be deconvoluted into four peaks with binding energies of 284.4 eV for C–C, 286.1 eV for C–O (including C–OH, C–O–C), 288.4 eV for C=O, and 291.1 eV which is identified as a shake-up satellite due to π–π* transitions.^45,46 In comparison with GAs, the peak position of C–O of SnO₂/GAs sample downshifted to 285.8 eV. Such a shift suggested the decrease of C–O–C (epoxide groups) in C–O groups. The XPS spectra of Sn 3d from the SnO₂ and SnO₂/GAs samples are shown in Fig. 3d. The symmetric peaks centered at 487.0 and 495.5 eV are assigned to Sn 3d_5/2 and Sn 3d_3/2, respectively, which are characteristic for Sn⁴⁺ in the SnO₂ NPs.⁴⁶ We noted that the peaks of Sn 3d_5/2 and Sn 3d_3/2 in SnO₂/GAs upshift to 487.4 and 495.8 eV, which are higher than those of as-prepared SnO₂ NPs. Such differences in binding energy probably indicate the presence of charge transfer between SnO₂ NPs and graphene sheets.

The porous nature of SnO₂/GAs was demonstrated by BET measurement (Fig. 4). The specific surface area of SnO₂/GAs reached up to 539 m² g⁻¹. This result highlighted that the building up of 3D frameworks was an effective way to achieve a high surface area for hybrid materials. BJH calculations disclosed the pore volume was 0.265 cm³ g⁻¹ with an average pore diameter of 3.5 nm (the inset in Fig. 4). The high porosity can provide not only more surface reaction sites but also sufficient buffer space to alleviate the volume expansion of SnO₂ during Li insertion/extraction and is therefore in favor of the electrochemical properties.

Fig. 4 Nitrogen adsorption and desorption isotherms of SnO₂/GAs.

The electrochemical performance of SnO₂/GAs was evaluated by the assembly of lithium half-cells with metallic lithium as the counter electrode. Fig. 5a shows the typical first three galvanostatic discharge–charge curves at current density of 50 mA g⁻¹ in the potential range of 0.05–3.0 V vs. Li⁺/Li. The initial discharge and charge capacity of SnO₂/GAs electrode are 1996 and 1036 mA h g⁻¹, respectively, giving a coulombic efficiency of 52%. The relatively low initial coulombic efficiency for the first cycle was mainly due to the irreversible amorphous lithium oxide and the formation of a solid electrolyte interface (SEI) layer over electrode. For the second cycle, the discharge and charge capacities were 1090 and 1012 mA h g⁻¹, respectively, and the coulombic efficiency increased to 94%.

Fig. 5 (a) The first three galvanostatic discharge–charge curves at a current density of 50 mA g⁻¹; (b) the cycling performance of SnO₂/GAs, GAs and SnO₂ measured at a current density of 50 mA g⁻¹.

For comparison, the cycling performance of SnO₂/GAs, GAs and SnO₂ measured at a current density of 50 mA g⁻¹ are shown in Fig. 5b. For the bare SnO₂ NPs electrode, there was a rapid decrease of capacity due to the severe pulverization. After 50 cycles, the charge capacity was only 207 mA h g⁻¹, which was about 23% retention of the initial reversible capacity. The GAs electrode displays higher electrochemical performance, showing a capacity of 542 mA h g⁻¹ after 50 cycles and about 55% retention of the initial reversible capacity. This result is obviously higher than the graphene electrode reported in previous literatures. For instance, the charge capacity of graphene electrode was only about 300 mA h g⁻¹ after tens cycles, which is close to that of graphite. The lower cycling performance should be caused by the stacking and aggregation of graphene layers, which decreases the accommodation of lithium ions. In our case, higher cycling performance of GAs is due to its unique 3D network architectures which prevents the agglomeration of the graphene sheets, as well as provides abundant reactive sites and fast electron transport channels in a 3D pathway. In comparison to bare SnO₂ NPs and GAs, the charge capacity and cycling performance of SnO₂/GAs composite electrode is enhanced significantly. The SnO₂/GAs composite shows a reversible discharge capacity of 760 mA h g⁻¹ after 50 cycles, 73% of the initial capacity (1036 mA h g⁻¹), exhibiting a high stability in the charge–discharge process. This reversible discharge capacity is close to the theoretical capacities of SnO₂ (782 mA h g⁻¹)⁷ and graphene (744 mA h g⁻¹),⁴⁷ indicating the superior advantage for electrochemical performance when both materials are hybrid to form SnO₂/GAs composite. As mentioned in previous literatures,^10,11 the main reason for the poor cyclic life of SnO₂ electrodes is the huge volume expansion produced by the alloying reaction of Li and Sn, leading to the pulverization and subsequent electrical disconnection of the electrodes. In our case of using SnO₂/GAs composite, both the higher mechanical strength of our 3D GAs compared to the reported graphene materials and the strong interaction between SnO₂ NPs and graphene sheets effectively suppress the volume change and cracking of the composite electrode, resulting in an enhanced cycling performance. On the other hand, the unique 3D network architectures together with the good separation effect from SnO₂ NPs prevent agglomeration of the graphene and the stacking of graphene layers (evidenced by XRD results), which preserve the intrinsic capacity performance of graphene sheets. The combination of the advantages of both SnO₂ NPs and GAs thus achieves high capacity and excellent cycling performance in their composites.

Multiple-current galvanostatical tests were carried out to investigate the high rate performance SnO₂/GAs and SnO₂ NPs. Fig. 6 shows the 10th discharge–charge curves at different current densities. It can be seen that the specific capacities of the SnO₂ NPs obviously decreased from 822 mA h g⁻¹ at 25 mA g⁻¹ to 19 mA h g⁻¹ at 1000 mA g⁻¹, while the specific capacities of the SnO₂/GAs show a significantly higher stability as current densities increase, leading to 929, 819, 730, 668, 582, 502, 424, and 361 mA h g⁻¹ at 25, 50, 100, 200, 400, 600, 800 and 1000 mA g⁻¹, respectively. All these values from the measurements with current densities below 1000 mA g⁻¹ are higher than the theoretical capacity of the commonly used graphite anode material (372 mA h g⁻¹). The improved rate capability of SnO₂/GAs should also be related to its unique 3D network structure. The 3D conductive network increases the electronic conductivity of the nanocomposite and facilitates the liquid electrolyte diffusion into the materials. Moreover, the nano-sized SnO₂ particles can shorten the path length for electronic and Li⁺ transport, which also favors the high rate capability.

Fig. 6 (a) The 10th discharge–charge curves of SnO₂ NPs and (b) of SnO₂/GAs at different current densities.

The structural stability of the samples after electrochemical measurements has been further examined by TEM. Fig. 7 shows TEM and HRTEM images of bare SnO₂ NPs and SnO₂/GAs samples after 50 cycles at a current density of 200 mA g⁻¹. For bare SnO₂ NPs electrode after the charge and discharge cycling, TEM observation shows that NPs tend to form agglomeration (Fig. 7a) and the HRTEM image (Fig. 7b) clearly shows the loss of crystallinity. For SnO₂/GNs electrode, it is clearly seen that the SnO₂ NPs are still highly dispersed after 50 cycles (Fig. 7c), due to the support of the stable graphene network. Furthermore, the HRTEM image (Fig. 7d) indicated that the SnO₂ NPs still preserve crystalline structure, which can be well indexed by the starting rutile phase. The results indicate that the charge/discharge reactions between SnO₂ and Li are reversible in the SnO₂/GAs nanocomposite. As shown in Fig. 7d, a uniform solid electrolyte interface (SEI) layer was formed and fully covered a SnO₂ nanoparticle after Li-ion battery measurement. Although the formation of the SEI layer might cause partial loss of the irreversible capacity, the SEI layer favors to maintain the robust electrode structure by preventing the aggregation and pulverization of SnO₂ NPs in the cycling process.

Fig. 7 TEM image (a) and HRTEM image (b) of pure SnO₂ NPs after 50 cycles at a current density of 200 mA g⁻¹; TEM image (c) and HRTEM image (d) of SnO₂/GAs after 50 cycles at a current density of 200 mA g⁻¹. The d-spacing of 0.33 nm corresponds to the (110) plane of a rutile phase.

Conclusions

We have successfully prepared SnO₂/GAs nanocomposites via a simple, facile and environmental friendly process, for which ultra-small SnO₂ nanocrystals with a size of 6 nm are well dispersed in the 3D graphene networks. Such unique 3D network structure effectively prevents the agglomeration of graphene sheets and provides open channels for lithium-ion transport and storage. On the other hand, the 3D network exhibit excellent structural stability and suppress the pulverization of SnO₂ nanoparticles during Li ion insertion/extraction. As a result, the reversible capacity of SnO₂/GAs after 50 cycles still remains 760 mA h g⁻¹, which is close to the theoretical capacities of SnO₂ (782 mA h g⁻¹) and graphene (744 mA h g⁻¹). TEM observations on the samples before and after 50 cycles at a current density of 200 mA g⁻¹ illustrates that the structures of the graphene network and SnO₂ NPs are preserved, which well explains the good cyclic stability of the electrode. We believe that this simple, fast and environmental friendly synthetic route can be further extended to produce other 3D graphene composites, for various applications in lithium-ion batteries, catalysts, sensors, and so on.

Acknowledgements

This work was supported financially by the National Basic Research Program of China (2011CB808200), NSFC (51025206, 51320105007, 51032001, 11104105, 11374120), the Cheung Kong Scholars Programme of China and 973 Program (no. 2015CB251103), and Open Project of State Key Laboratory of Superhard Materials (Jilin University) (no. 201404, no. 201406).

Notes and references

1. J. M. Tarascon and M. Armand, Nature, 2001, 414, 359–367.
2. C. Liu, F. Li, L. P. Ma and H. M. Cheng, Adv. Mater., 2010, 22, E28–E62.
3. J. B. Goodenough and Y. Kim, Chem. Mater., 2010, 22, 587–603.
4. E. J. Yoo, J. Kim, E. Hosono, H. S. Zhou, T. Kudo and I. Honma, Nano Lett., 2008, 8, 2277–2282.
5. G. Armstrong, A. R. Armstrong, P. G. Bruce, P. Reale and B. Scrosati, Adv. Mater., 2006, 18, 2597–2600.
6. N. A. Kaskhedikar and J. Maier, Adv. Mater., 2009, 21, 2664–2680.
7. Z. Y. Wang, L. Zhou and X. W. Lou, Adv. Mater., 2012, 24, 1903–1911.
8. C. Wang, G. H. Du, K. Stahl, H. X. Huang, Y. J. Zhong and J. Z. Jiang, J. Phys. Chem. C, 2012, 116, 4000–4011.
9. P. Wu, N. Du, H. Zhang, J. X. Yu and D. R. Yang, J. Phys. Chem. C, 2010, 114, 22535–22538.
10. J. O. Besenhard, J. Yang and M. J. Winter, Power Sources, 1997, 68, 87–90.
11. J. Fan, T. Wang, C. Z. Yu, B. Tu, Z. Y. Jiang and D. Y. Zhao, Adv. Mater., 2004, 16, 1432–1436.
12. Y. L. Zhang, Y. Liu and M. L. Liu, Chem. Mater., 2006, 18, 4643–4646.
13. C. C. Chang, S. J. Liu, J. J. Wu and C. H. Yang, J. Phys. Chem. C, 2007, 111, 16423–16427.
14. J. S. Chen and X. W. Lou, Small, 2013, 9, 1877–1893.
15. Z. Y. Wang, Z. C. Wang, S. Madhavi and X. W. Lou, Chem.–Eur. J., 2012, 18, 7561–7567.
16. J. Liu, W. Liu, K. F. Chen, S. M. Ji, Y. C. Zhou, Y. L. Wan, D. F. Xue, P. Hodgson and Y. C. Li, Chem.–Eur. J., 2013, 19, 9811–9816.
17. G. Chen, Z. Y. Wang and D. G. Xia, Chem. Mater., 2008, 20, 6951–6956.
18. L. Zhang, G. Q. Zhang, H. B. Wu, L. Yu and X. W. Lou, Adv. Mater., 2013, 25, 2589–2593.
19. X. F. Li, X. B. Meng, J. Liu, D. S. Geng, Y. Zhang, M. N. Banis, Y. L. Li, J. L. Yang, R. Y. Li, X. L. Sun, M. Cai and M. W. Verbrugge, Adv. Funct. Mater., 2012, 22, 1647–1654.
20. S. Yang, W. B. Yue, J. Zhu, Y. Ren and X. J. Yang, Adv. Funct. Mater., 2013, 23, 3570–3576.
21. J. F. Liang, W. Wei, D. Zhong, Q. L. Yang, L. D. Li and L. Guo, ACS Appl. Mater. Interfaces, 2012, 4, 454–459.
22. H. Kim, M. K. Yun and H. H. Park, Phys. Status Solidi A, 2011, 208, 1869–1872.
23. S. G. Ding, D. Y. Luan, F. Y. C. Boey, J. S. Chen and X. W. Lou, Chem. Commun., 2011, 47, 7155–7157.
24. Y. P. Tang, D. Q. Wu, S. Chen, F. Zhang, J. P. Jia and X. L. Feng, Energy Environ. Sci., 2013, 6, 2447–2451.
25. X. S. Zhou, L. J. Wan and Y. G. Guo, Adv. Mater., 2013, 25, 2152–2157.
26. F. Ye, B. T. Zhao, R. Ran and Z. P. Shao, Chem.–Eur. J., 2014, 20, 4055–4063.
27. A. K. Geim and K. S. Novoselov, Nat. Mater., 2007, 6, 183–191.
28. Y. Y. Li, Z. S. Li and P. K. Shen, Adv. Mater., 2013, 25, 2474–2480.
29. Y. Z. Su, S. Li, D. Q. Wu, F. Zhang, H. W. Liang, P. F. Gao, C. Cheng and X. L. Feng, ACS Nano, 2012, 6, 8349–8356.
30. Z. P. Chen, W. C. Ren, L. B. Gao, B. L. Liu, S. F. Pei and H. M. Cheng, Nat. Mater., 2011, 10, 424–428.
31. H. Wang, K. Sun, F. Tao, D. J. Stacchiola and Y. H. Hu, Angew. Chem., Int. Ed., 2013, 52, 9210–9214.
32. Y. M. He, W. J. Chen, X. D. Li, Z. X. Zhang, J. C. Fu, C. H. Zhao and E. Q. Xie, ACS Nano, 2013, 7, 174–182.
33. W. F. Chen, S. R. Li, C. H. Chen and L. F. Yan, Adv. Mater., 2011, 23, 5679–5683.
34. R. H. Wang, C. H. Xu, M. Du, J. Sun, L. Gao, P. Zhang, H. L. Yao and C. C. Lin, Small, 2014, 10, 2260–2269.
35. B. C. Qiu, M. Y. Xing and J. L. Zhang, J. Am. Chem. Soc., 2014, 136, 5852–5855.
36. M. Zhang, Y. Wang and M. Q. Jia, Electrochim. Acta, 2014, 129, 425–432.
37. H. Y. Sun, Y. G. Liu, Y. L. Yu, M. Ahmade, D. Nan and J. Zhua, Electrochim. Acta, 2014, 118, 1–9.
38. R. H. Wang, C. H. Xu, J. Sun, L. Gao and H. L. Yao, ACS Appl. Mater. Interfaces, 2014, 6, 3427–3436.
39. K. X. Sheng, Y. X. Xu, C. Li and G. Q. Shi, New Carbon Mater., 2011, 26, 9–15.
40. W. S. Hummers and R. E. Offeman, J. Am. Chem. Soc., 1958, 80, 1339.
41. Z. S. Wu, S. B. Yang, Y. Sun, K. Parvez, X. L. Feng and K. Müllen, J. Am. Chem. Soc., 2012, 134, 9082–9085.
42. D. Y. Chen, G. Ji, Y. Ma, J. Y. Lee and J. M. Lu, ACS Appl. Mater. Interfaces, 2011, 3, 3078–3083.
43. S. Stankovich, D. A. Dikin, R. D. Piner, K. A. Kohlhaas, A. Kleinhammes, Y. Jia, Y. Wu, S. T. Nguyen and R. S. Ruoff, Carbon, 2007, 45, 1558–1565.
44. J. L. Zhang, H. J. Yang, G. X. Shen, P. Cheng, J. Y. Zhang and S. W. Guo, Chem. Commun., 2010, 46, 1112–1114.
45. D. X. Yang, A. Velamakanni, G. Bozoklu, S. Park, M. Stoller, R. D. Piner, S. Stankovich, I. Jung, D. A. Field, J. C. A. Ventrice and R. S. Ruoff, Carbon, 2009, 47, 145–152.
46. Y. M. Li, X. J. Lv, J. Lu and J. H. Li, J. Phys. Chem. C, 2010, 114, 21770–21774.
47. G. X. Wang, B. Wang, X. L. Wang, J. Park, S. X. Dou, H. J. Ahn and K. Kim, J. Mater. Chem., 2009, 19, 8378–8384.

---