Synthesis, characterization and application of carbon nanocages as anode materials for high-performance lithium-ion batteries

Abstract

Carbon nanocages (CNCs) with diameters of about 200∼500 nm have been synthesized by a simple method. The electrochemical properties of the CNCs as anode materials were evaluated by discharge/charge measurement and electrochemical impedance spectroscopy. Results showed that the CNCs displayed excellent cycling performance and good rate capability with no noticeable capacity fading up to 50 cycles at current densities of 100, 300 and 500 mA g⁻¹. Their electrochemical properties were significantly improved after annealing treatment of the CNCs at 600 °C. For example, the CNCs(2)-annealed exhibited much better electrochemical performance with a high reversible capacity of 520 mAh g⁻¹ after 50 cycles at current density of 100 mA g⁻¹. A discharge capacity of 380 mAh g⁻¹ can be obtained after 50 cycles at high current density of 500 mA g⁻¹. The results of the Raman, thermal gravimetric and electrochemical impedance analysis indicated that the graphitization degree, electronic conductivity and charge-transfer rate of the CNCs have been improved after annealing treatment.

---

Introduction

Lithium-ion batteries have received considerable attention in applications, ranging from portable electronics to electric vehicles, due to their superior energy density over other rechargeable batteries.¹ In commercial lithium-ion batteries, graphite, one carbonaceous material with the high graphitization, is almost the only widely used anode material due to its stable specific capacity, small irreversible capacity, and good cycling performance.² Although carbonaceous materials have been widely used in commercial lithium-ion batteries, poor rate performance and relatively low storage capacity restrict their applications in wider areas.³ Furthermore, conventional carbon materials require usually high temperature treatment to form a graphitic layered structure.⁴ However, the rapid development of electronic devices and electric vehicles demands high storage capacity and good rate capability for anode materials.

In the last few decades, carbon nanomaterials such as carbon nanotubes, carbon nanofibers, graphene and composites based on these nanostructures have been widely investigated because nanoscale carbon materials have a relatively large number of lithium ion insertion sites on their surface, and charge-transfer resistances at the interface between the electrolyte and electrode materials are expected to be small.^5–13 These materials exhibit improved storage capacity but most anodes of pure carbon nanomaterials still suffer from the capacity decay in the discharge/charge cycles.

At present, few reports focus on the 3D hollow carbon nanomaterials such as hollow carbon nanospheres and hollow carbon nanocages (CNCs) in lithium-ion batteries. 3D carbon nanomaterials can be readily used for improving the rate capability of the lithium-ion batteries because the solid-state diffusion length is short and their relatively large surface area can also benefit the charge-transfer rate. In particular, 3D hollow carbon nanomaterials provide several advantages for applications in lithium-ion batteries: electrode materials with hollow structures can accommodate lager volume changes during the discharge/charge process, and hollow structures can provide extra space for Li ion storage.^14,15 Tien et al.¹⁶ synthesized uniform carbon spheres and the as-prepared anode exhibited a stable capacity above 330 mAh g⁻¹ after 130 cycles. Lee et al.¹⁷ synthesized a 3D porous carbon material by a sol–gel process. The porous carbon electrode showed a superior rate capability compared to spherical carbon and bulk carbon. Wang et al.¹¹ reported that the CNCs/graphene have been prepared by catalytic decomposition of xylene on MgO supported Co and Mo catalyst in supercritical CO₂ at a pressure of 10.34 MPa and temperature ranging from 650 to 750 °C. The reversible capacity of the CNCs could achieve 574 mAh g⁻¹ at a current density of 100 mA g⁻¹ after 60 discharge/charge cycles. However, fabricating 3D hollow carbon nanomaterials through a simple method to take advantages and restrain shortcomings is a challenge work for developing high-performance lithium-ion batteries.

In this study, we report a simple strategy to synthesize high-yield CNCs and Fe₃O₄/CNCs as anode materials for high-performance lithium-ion batteries. The as-obtained CNCs have diameters of 200∼500 nm. The CNCs and Fe₃O₄/CNCs display excellent cycling performance and good rate capacity. It is found that their electrochemical properties can be significantly improved after the annealing treatment at 600 °C for 10 h. To the best of our knowledge, there are few research reports about the CNCs for lithium-ion batteries.

Experimental

Preparation of CNCs

In a typical experimental procedure, 6.0 g citric acid and 2.5 g nickel oxalate were added into a stainless-steel autoclave of 20 ml capacity. The autoclave was sealed and put into an electronic furnace at room temperature. The temperature of the furnace was increased from room temperature to 550 °C in 55 min and maintained at 550 °C for 12 h, and then the autoclave was allowed to cool to room temperature naturally. The dark precipitates in the autoclave were collected and washed with absolute ethanol, dilute hydrochloric acid, and distilled water for several times. After that, these products were dried in a vacuum oven at 60 °C for 5 h for further characterization. This sample is denoted as CNCs(1).

Ethanol (15 ml) and ferrous oxalate (1.5 g) were added into a stainless-steel autoclave of 20 ml capacity. Other reaction parameters were similar with above mentioned processes. The obtained samples in the autoclave were washed with absolute ethanol and distilled water for several times. After that, the Fe₃O₄/CNCs composite materials were obtained. These composites were washed with hydrochloric acid for several times, and then CNCs also could be obtained. These samples were denoted as CNCs(2) and Fe₃O₄/CNCs(2). The detailed experiment process and growth mechanism of this sample can be found in our previous report.¹⁸

The as-obtained samples (CNCs(1), CNCs(2), and Fe₃O₄/CNCs(2)) were annealed at 600 °C in a tubular furnace under following argon atmosphere for 10 h. After this step, these annealed samples are denoted as CNCs(1)-annealed, CNCs(2)-annealed, and Fe₃O₄/CNCs(2)-annealed, respectively.

Characterization

The X-ray diffraction (XRD) analysis was performed on a Bruker D8 advanced X-ray diffractometer equipped with graphite-monochromatized Cu-Kα radiation (λ = 1.5418 Å). The Raman spectrum was recorded at ambient temperature on a LABRAM-HR confocal laser MicroRaman spectrometer with an argon-ion laser at an excitation wavelength of 514.5 nm. Inductively coupled plasma (ICP) analysis was performed using an ICP Perkin-Elmer Optima 7300DV. The morphologies of the samples were observed through scanning electron microscopy (SEM) and transmission electron microscope (TEM) measurements, which were carried out on a JEOL JSM-7600F field emission instrument and a Hitachi H-7000 TEM, respectively. High-resolution transmission electron microscope (HRTEM) images were carried out on a JEOL 2100 transmission electron microscope with an accelerating voltage of 200 KV. Thermal gravimetric analysis (TGA) was carried out on a Mettler Toledo TGA/SDTA851 thermal analyzer apparatus under air atmosphere.

Electrochemical measurements

The electrochemical discharge/charge tests of the samples were performed on a Land battery test system (CT2001A). The working electrodes were prepared by pasting a mixed slurry that consisted of 80 wt% active materials (CNCs or CNCs/Fe₃O₄), 10 wt% poly(vinylidene fluoride) (PVDF), and 10 wt% carbon black onto a copper foil. The fabricated working electrodes dried in a vacuum oven at 120 °C for 24 h. A Celgard 2300 microporous polypropylene membrane was used as separator. The electrolyte consisted of a solution of 1 M LiPF₆ in an ethylene carbonate/dimethyl carbonate/diethyl carbonate (EC/DMC/DEC, 1 : 1 : 1 w/w). Lithium foils were used as counter electrodes. The batteries were assembled in an argon-filled glove box and cycled at different current densities between voltage limits of 0.01 and 2.5 V. The electrochemical impedance spectroscopy (EIS) was measured with a Princeton Applied Research. The EIS were carried out by applying an alternating current voltage of 10 mV in the frequency from 10 mHz to 100 kHz at open-circuit voltage before the discharge/charge test.

Results and discussion

Fig. 1a shows a typical XRD pattern of the raw sample without hydrochloric acid treatment. The reflection peaks in this pattern can be indexed to face-centered cubic nickel (JCPDS card no. 04-0850). It is demonstrated that nickel were formed in the reaction system. The ICP result shows that the CNCs(1) without acid treatment contain Ni (53.5 wt%) element. Fig. 1b shows the XRD pattern of the sample after being washed with diluted hydrochloric acid. All of the diffraction peaks in this pattern can be indexed to the (002) and (101) diffraction planes of hexagonal graphite (JCPDS card no. 41-1487). Trace amount of Ni (0.42 wt%) remained in the final CNCs(1) according to the ICP test. Two broad diffraction peaks can be clearly observed at 26° and 45°, indicating that low crystallinity or disordered graphite was produced.

Fig. 1 (a) XRD patterns of the CNCs(1) before treated with hydrochloric acid and (b) after being washed with hydrochloric acid.

The morphology of the as-prepared samples is further characterized by SEM and TEM. The SEM and TEM images in Fig. 2a and b show hollow polygonal morphology of the CNCs(1) which indicate that large quantity of the CNCs(1) were obtained by this simple method. Most of the CNCs(1) have diameters ranging from 200 to 500 nm and the wall thickness is about 70 nm. The wall of a carbon nanocage is composed of amorphous or disordered structure through HRTEM observation (see Fig. 2c). The disordered character is also detected by the corresponding XRD pattern. Fig. 2d and e show the TEM images of the samples with different treatment process. It can be clearly seen that the solid nickel particles are encapsulated in the graphite layers, which clearly evidenced the growth mechanism of the CNCs(1). In the initial stage of the reaction, the decomposition of the nickel oxalate occurs at relatively low temperature, which leads to the formation of small polygonal nickel nanoparticles. With the increasing reaction temperature, citric acid decomposes into carbon atoms. Then, the above mentioned as-formed nickel nanoparticles serve as nucleation center and the carbon atoms coat around the nickel nanoparticles. Hollow CNCs(1) could be obtained after the Ni particles being washed with hydrochloric acid solution. Furthermore, it is found that the CNCs(1) could also be obtained when using other carbon source such as ethanol, glucose, and so on. A schematic image of the formation process for CNCs is shown in Fig. 2f.

Fig. 2 (a, b) SEM and TEM images of the CNCs(1); (c) HRTEM image of the wall structure of a carbon nanocage; TEM images of the CNCs(1) with different treatment processes: (d) before treated with acid; (e) washed with dilute hydrochloric acid for ∼2 h. (f) A schematic illustration of the formation process of the CNCs: (A) polygonal nickel nanoparticles were formed at first; (B) carbon atoms coated around these nickel nanoparticles; (C) graphitic layers were formed subsequently; (D) CNCs obtained after acid treatment process.

The SEM image (Fig. 3a) reveals that a large quantity of quadrangular CNCs(2) are also obtained by this approach. The length of the side of CNCs(2) ranges from 200 to 350 nm and the wall thickness is about 10∼15 nm (Fig. 3e). Fig. 3b shows a typical image of the Fe₃O₄/CNCs(2) composites. Fig. 3c and d reveal clearly structural characteristic of the Fe₃O₄/CNCs(2) composites, which display a general morphology and distinct core/shell structure characteristics. It is clearly seen that Fe₃O₄/CNCs(2) composites are composed of outer graphite layers and inner Fe₃O₄ nanoparticles. Trace amount of Fe (0.35 wt%) was also detected in the final CNCs(2) according to the ICP test after being washed with acid.

Fig. 3 (a) SEM image of the CNCs(2); (b) TEM image of the Fe₃O₄/CNCs(2) composites; (c, d) HRTEM image of the area marked with a rectangle in (b); (e) HRTEM image of the wall structure of a carbon nanocage.

Raman spectra have been used to investigate the graphitization degree of the carbon materials.¹⁹ The Raman spectra of the CNCs(1) (Fig. 4a and b) show that there are two strong peaks centered at ∼1350 (D-band) and ∼1580 cm⁻¹ (G-band). The G-band at 1580 cm⁻¹ corresponds to the Raman-allowed optical mode E_2g, which is closely related to the vibration in all sp²-band carbon atoms in a two-dimensional hexagonal lattice. The D-band at 1350 cm⁻¹ could be assigned to the vibrations of carbon atoms with dangling bands in planar terminations of disordered graphite. Therefore, the intensity ratio of D-band and G-band (I_D/I_G) is usually used to determine the disorder degree in graphite layers. The higher ratio of I_D/I_G indicates more defects in graphite layers.^20,21 It is worth noting that the I_D/I_G of the CNCs(1)-annealed shows a noticeable decrease compared to that of the CNCs(1). This result reveals that the graphitization degree of the CNCs(1) were improved after annealing treatment. The improved graphitization degree and quality of the CNCs is very important to the performance of the CNCs for lithium-ion batteries.

Fig. 4 Raman spectra of the (a) CNCs(1) and (b) CNCs(1)-annealed.

The graphitization degree and quality of the CNCs is evaluated by the TGA measurements.^22,23Fig. 5a shows the TGA curves of the CNCs(1) and CNCs(1)-annealed. It is found that the CNCs(1)-annealed are more stable under high-temperature air oxidation than those of CNCs(1). The combustion temperature of the CNCs is remarkably increased after an annealing treatment. Compared with the CNCs(1) (∼500 °C), the combustion temperature of CNCs(1)-annealed is increased ∼100 °C. The similar phenomenon and tendency can also be found in Fig. 5b. It is known that the defect sites lead to a decrease in the thermal stability of the carbon materials at elevated temperature. High annealed temperature can facilitate the removal of functional groups and heal the defects in graphite layers.²⁴ This further proves that the graphitization degree and structure stability can be significantly improved after annealing treatment. The high graphitization degree and low defects play important roles in improving the cycling performance of the carbon anode materials in lithium-ion batteries.

Fig. 5 TGA curves of the CNCs: (a) CNCs(1) and CNCs(1)-annealed; (b) CNCs(2) and CNCs(2)-annealed.

The electrochemical properties of the CNCs for lithium-ion battery applications are evaluated using galvanostatic discharge/charge cycles. Discharge/charge cycles are carried out in the potential window of 0.01–2.5 V versusLi. The voltage versus capacity and capacity versus cycle number curves for the CNCs and Fe₃O₄/CNCs at different current densities of 100, 300 and 500 mA g⁻¹ are shown in Fig. 6. The first discharge capacities are 734 mAh g⁻¹ for CNCs(1), 875 mAh g⁻¹ for CNCs(1)-annealed, 808 mAh g⁻¹ for CNCs(2), 1050 mAh g⁻¹ for CNCs(2)-annealed, 1442 mAh g⁻¹ for Fe₃O₄/CNCs(2) and 1580 mAh g⁻¹ for Fe₃O₄/CNCs(2)-annealed at current density of 100 mA g⁻¹. It can be observed that the cycling performance and rate capability of the samples have been significantly improved after annealing treatment at 600 °C for 10 h. Because the graphitic carbon have a high electronic conductivity and stabilized structure compared with that of disordered carbon.²⁵ The capacity fades rapidly with high irreversible capacity in the first cycle. The first capacity fading mainly attribute to the decomposition of the electrolyte and the formation of the solid electrolyte interphase (SEI) layer at the surface of the CNCs.^26,27 The irreversible capacity of carbon materials at the initial cycles is related to the reaction of lithium with active sites and functional groups (such as –C=O or –OH etc.) on the electrode materials.^28,29 Furthermore, a relatively volume change during the cycles can result in structural instability and capacity fading.¹ During the subsequent cycles, the CNCs(1) and CNCs(1)-annealed present stable performance even at high current density of 500 mA g⁻¹. It can be seen that the reversible capacities of the CNCs(1) and CNCs(1)-annealed are about 335 mAh g⁻¹ and 380 mAh g⁻¹ at current density of 100 mA g⁻¹ after 50 cycles.

Fig. 6 The first discharge/charge voltage curves and cycling performance at current densities of 100 mA g⁻¹, 300 mA g⁻¹ and 500 mA g⁻¹: (a) CNCs(1); (b) CNCs(1)-annealed; (c) CNCs(2); (d) CNCs(2)-annealed; (e) Fe₃O₄/CNCs(2); (f) Fe₃O₄/CNCs(2)-annealed.

Fig. 6c and d show the discharge/charge curves and cycling performance of the CNCs(2) and CNCs(2)-annealed. The CNCs(2)-annealed anode exhibits a good cycle stability and large reversible capacity. Although the discharge capacity dropped from 1050 mAh g⁻¹ to 480 mAh g⁻¹ after 10 cycles, the discharge capacity of the CNCs(2)-annealed anode gradually restored during the subsequent cycles, even after 50 cycles the anode remained a capacity of ∼520 mAh g⁻¹ at current density of 100 mA g⁻¹. More importantly, this sample exhibits a much better rate capability. After 50 cycles, the discharge capacity of this sample can be stably retained at about 380 mAh g⁻¹ at high current density of 500 mA g⁻¹.

Fig. 6e and f show the discharge/charge curves of the Fe₃O₄/CNCs(2) composite electrodes in the first cycle. In the first discharge cycle, both of them display a long voltage plateau at 0.5–0.75 V, and followed by a sloping curve down to the cut off voltage of 0.01 V, indicating a typical characteristic of voltage trends for the Fe₃O₄ electrode.^19,30 The initial discharge capacity of Fe₃O₄/CNCs(2) is about 1442 mAh g⁻¹ at a current density of 100 mA g⁻¹, and the capacity is about 380 mAh g⁻¹ after 50 cycles. However, at the large current density of 500 mA g⁻¹, the capacity is only about 250 mAh g⁻¹ after 50 cycles. The discharge capacity of the Fe₃O₄/CNCs(2) is poor as compared with that of other samples. After annealing treatment, the discharge capacity can reach 450 mAh g⁻¹ at a current density of 100 mA g⁻¹ and 320 mAh g⁻¹ at a current density of 500 mA g⁻¹ after 50 cycles. The improved cycling performance and rate capability can be attribute to the high electronic conductivity and stabilized structure. The stabilized graphitization layer can effective restrict the volume effect of the Fe₃O₄ nanoparticles.²⁵

The excellent cycling performance and high-rate capability of the CNCs are ascribed to the unique hollow structure, which can improve the electron transportation, electrolyte penetration and shortened Li ion diffusion pathway in the discharge/charge cycles.¹¹

In order to further compare the electrochemical performance of the typical pure carbon nanomaterials, some relevant information are summarized in Table 1. It is found that the first discharge capacities of the different CNCs have not been obviously improved when comparing with other pure carbon nanomaterials. But the as-prepared CNCs show both good cycling performance and enhanced high-rate performance. The reversible capacity after 50 cycles for CNCs(2)-annealed at current density of 100 mA g⁻¹ is 520 mAh g⁻¹, which is much larger than many previous reported pure carbon nanomaterials.

Table 1 Cycling performance and capacity of pure carbon nanomaterials reported in previous works

Typical materials	Current density mA g^−1	First capacity mAh g^−1	Cycle number	Remaining capacity mAh g^−1	Ref.
Carbon nanotubes	57	1800	36	750	27
Carbon nanotubes	75	310	20	320	31
Carbon nanotubes	75	1000	30	200	32
Carbon nanotubes	25	1200	50	450	33
Single-wall carbon nanotubes	50	942	60	150	34
Ordered mesoporous carbon	50	3083	50	500	35
Ordered mesoporous carbon	100	3100	20	850	36
Hierarchical porous carbon	18.6	710	20	220	37
Hierarchical porous carbon	75	1500	40	500	38
Carbon nanosprings	50	750	100	420	3
Carbon nanofibers	100	430	30	220	5
Carbon nanofibers	100	680	25	260	7
Carbon spheres	50	950	135	400	16
Hollow carbon nanospheres	50	540	15	290	39
CNCs/graphene	100	1271	60	574	11
CNCs(1)	100	734	50	335	this work
CNCs(1)	500	586	50	238	this work
CNCs(1)-annealed	100	875	50	373	this work
CNCs(1)-annealed	500	507	50	280	this work
CNCs(2)	100	808	50	314	this work
CNCs(2)	500	427	50	242	this work
CNCs(2)-annealed	100	1050	50	520	this work
CNCs(2)-annealed	500	728	50	380	this work

---

Fig. 7 shows the Nyquist plots of the CNCs and CNCs-annealed. The depressed semicircles in the high frequency are related to the charge transfer process. The numerical value of the diameter of the semicircle on the Z_re axis is approximately equal to the charge transfer resistance (R_ct).⁴⁰ It can be seen that there is an obvious decrease in R_ct after the annealing treatment. This indicates that the electronic conductivity of the CNCs was improved after the annealing treatment.^41,42 The inclined lines in the low frequency are attributed to the Warburg impedance, which is associated with lithium ion diffusion in the electrode materials (CNCs). The lithium ion diffusion coefficient could be calculated using the following equation:⁴³

D = 0.5 (RT/AF²Cσ)² (1)

Herein, R is the gas constant, T is the absolute temperature, A is the surface area of the electrode, F is the Faraday constant, C is the concentration of lithium ions, and σ is the Warburg impedance coefficient which is associated with Z_re.

Z_re = R_e + R_ct + σω^−1/2 (2)

Herein, R_e is the electrolyte resistance, R_ct is the charge transfer resistance, ω is the angular frequency in the low frequency. The parameters of the equivalent circuit and diffusion coefficients of the CNCs are calculated and recorded in Table 2.

Fig. 7 Nyquist plots of the CNCs at open-circuit voltage.

Table 2 Electrochemical impedance parameters of the CNCs and CNCs-annealed cell samples

Samples	Re (Ω)	Rct (Ω)	σ (Ω cm^2 S^−0.5)	D (cm^2 S^−1)	i^o (mA cm^−2)
CNCs(1)	1.5	502.4	165.1	2.11E−15	5.11E−5
CNCs(1)-annealed	2.2	109.5	114.4	4.93E−15	2.34E−4
CNCs(2)	1.1	352.9	228.4	1.10E−15	7.28E−5
CNCs(2)-annealed	1.8	101.4	47.7	2.52E−14	2.53E−4

---

It is found that the lithium ion diffusion coefficient of the electrode increased after annealing treatment. The diffusion coefficient of the CNCs(2)-annealed shows the higher mobility for lithium ions diffusion than other electrode materials.³⁸ Furthermore, the exchange current density (i^o = RT/nFR_ct) of the CNCs(2)-annealed is higher than in other electrode materials. Therefore, the charge transfer reaction of the CNCs(2)-annealed is stronger than that of other electrode materials.^44,45 It is demonstrated that the R_ct obviously decreases and the lithium ion diffusion coefficient increases after annealing treatment, and therefore improves the discharge/charge performance and rate capability of the batteries.

The reason for the different capacities in lithium-ion batteries with the use of the different kinds of CNCs can also be explained by the Nyquist plots. It is obvious that the charge transfer resistance (R_ct) of the CNCs(2) was smaller than that of CNCs(1), indicating the better electronic conductivity (means the higher Li ion transfer speed across interfaces between the electrolyte and active electrode materials)^3,5 of the former sample. In addition, the wall thickness of the CNCs(1) was about 70 nm while that of the CNCs(2) was only about 10∼15 nm (see Fig. 2b and Fig. 3e). Thinner graphite layers are favorable for the electrolyte penetration and shortened Li ion diffusion paths in discharge/charge cycling.¹⁴ Therefore, the capacities and performances of the CNCs(2) were better than those of CNCs(1) as anode materials in lithium-ion batteries in this study.

Conclusions

CNCs with a diameter of about 200∼500 nm have been successfully synthesized by a simple pyrolysis method. Electrochemical measurements show that the CNCs exhibited excellent cycling performance and high-rate capability. The cycling performance and rate capability have been significantly improved after annealing treatment of the samples at 600 °C for 10 h. It is demonstrated that the annealing treatment process not only increase the electronic conductivity and lithium ion diffusion coefficient but also decrease the charge transfer resistance, and obviously improve the cycling performance and rate capability of the CNCs. These results reveal that the CNCs would become promising anode materials for lithium-ion batteries.

Acknowledgements

This work was supported by the 973 Project of China (No. 2011CB935901), the National Nature Science Found of China and Academy of Sciences large apparatus United Found (Grant Nos. 11179043, 20971079 and 20871075). and the Independent Innovation Foundations of Shandong University (Grant Nos. 2009TS017, 2009JC019), The authors thank Congming Yin for helpful discussions and also thank Liancheng Wang for help in the ICP test.

References

1. J. B. Brian, J. G. Matthew, D. C. Ganter, A. D. Roberta and P. R. Ryne, Energy Environ. Sci., 2009, 2, 638.
2. N. Du, H. Zhang, B. D. Chen, X. Y. Ma and W. F. Zhang, Mater. Res. Bull., 2009, 44, 211.
3. X. Wu, Q. Liu, Y. Guo and W. Song, Electrochem. Commun., 2009, 11, 1468.
4. L. Wang, Y. Yu, P. Chen and C. Chen, Scr. Mater., 2008, 58, 405.
5. G. Zou, D. Zhang, C. Dong, H. Li, K. Xiong, L. Fei and Y. Qian, Carbon, 2006, 44, 828.
6. A. Varzi, C. Taubert, M. Wohlfahrt, M. Kreis and W. Schutz, J. Power Sources, 2011, 196, 3303.
7. D. Deng and Y. Lee, Chem. Mater., 2007, 19, 4198.
8. Y. Wang, H. C. Zeng and J. Y. Lee, Adv. Mater., 2006, 18, 645.
9. S. Zheng, J. Hu, L. Zhong, W. Song, L. Jun and Y. Guo, Chem. Mater., 2008, 20, 3617.
10. C. Zhai, N. Du, H. Zhang, J. Yu, P. Wu, C. Xiao and D. Yang, Nanoscale, 2011, 3, 1798.
11. K. Wang, Z. Li, Y. Wang, H. Liu, J. Chen, J. Holmes and H. Zhou, J. Mater. Chem., 2010, 20, 9748.
12. Z. Wu, W. Ren, L. Wen, L. Gao, J. Zhao, Z. Chen, G. Zhou, F. Li and H. Cheng, ACS Nano, 2010, 4, 3187.
13. J. Zhu, T. Zhu, X. Zhou, Y. Zhang, X. W. Lou, X. Chen, H. Zhang, H. H. Hng and Q. Yan, Nanoscale, 2011, 3, 1084.
14. C. Liu, F. Li, L. Ma and H. Cheng, Adv. Mater., 2010, 22, E28.
15. Y. Lin, J. Duh and M. hung, J. Phys. Chem. C, 2010, 114, 13136.
16. B. Tien, M. Xu and J. Liu, Mater. Lett., 2010, 64, 1465.
17. K. T. Lee, J. C. Lytle, N. S. Ergang, S. M. Oh and A. Stein, Adv. Funct. Mater., 2005, 15, 547.
18. G. D. Li, H. X. Yu, L. Q. Xu, Q. Ma, C. Chao, Q. Hao and Y. T. Qian, Nanoscale, 2011, 3, 3251.
19. C. Klinke, R. Kurt, J. M. Bonard and K. Kern, J. Phys. Chem. B, 2002, 106, 11191.
20. K. Niwase, T. Homae, K. G. Nakamura and K. Kondo, Chem. Phys. Lett., 1997, 362, 47.
21. A. C. Ferrari and J. Robertson, Phys. Rev. B: Condens. Matter, 2000, 61, 14095.
22. B. David, A. Rodney, J. David, C. Bailin and M. Mark, Nano Lett., 2002, 2, 615.
23. C. M. Chen, M. Chen, F. C. Leu, S. Y. Hsu, S. C. Wang and S. C. Shi, Diamond Relat. Mater., 2004, 13, 1182.
24. Q. Liu, W. Ren, F. Li, H. Cong and H. M. Cheng, J. Phys. Chem. C, 2007, 111, 5006.
25. L. Wang, Y. Yu, P. C. Chen, D. W. Zhang and C. H. Chen, J. Power Sources, 2008, 183, 717.
26. W. Xing, P. Bai, Z. F. Li, R. J. Yu, Z. F. Yan, G. Q. Lu and L. M. Lu, Electrochim. Acta, 2006, 51, 4626.
27. D. T. Welna, L. Qu, B. E. Taylor, L. Dai and M. F. Durstock, J. Power Sources, 2011, 196, 1455.
28. L. Ji, Z. Lin, M. Alcoutlabi and X. Zhang, Energy Environ. Sci., 2011, 4, 2682.
29. Y. Matusumura, S. Wang and J. Mondori, J. Electrochem. Soc., 1995, 142, 2914.
30. Z. Cui, L. Jiang, W. Song and Y. Guo, Chem. Mater., 2009, 21, 1162.
31. R. Lv, L. Zou, X. Gui, F. Kang, Y. Zhu, H. Zhu, J. Wei, J. Gu, K. Wang and D. Wu, Chem. Commun., 2008, 44, 2046.
32. L. Zou, R. Lv, F. Kang, L. Gan and W. Shen, J. Power Sources, 2008, 184, 566.
33. X. Xia, J. N. Wang, H. Chang and Y. F. Zhang, Adv. Funct. Mater., 2007, 17, 3613.
34. J. X. Li, Y. Zhao and L. H. Guan, Electrochem. Commun., 2010, 12, 592.
35. H. Li, R. Liu, D. Zhao and Y. Xia, Carbon, 2007, 45, 2628.
36. H. Zhou, S. Zhu, M. Hibino, I. Honma and M. Ichihara, Adv. Mater., 2003, 15, 2107.
37. J. Yi, X. P. Li, S. J. Hu, W. S. Li, L. Zhou, M. Q. Xu, J. F. Lei and L. S. Hao, J. Power Sources, 2011, 196, 6670.
38. Y. Hu, P. Adelhelm, B. M. Smarsly, S. Hore, M. Antonietti and J. Maier, Adv. Funct. Mater., 2007, 17, 1873.
39. C. Wu, X. Zhu, L. Ye, C. OuYang, S. Hu, L. Lei and Y. Xie, Inorg. Chem., 2006, 45, 8543.
40. A. Y. Shenouda and H. K. Liu, J. Alloys Compd., 2009, 477, 498.
41. M. M. Rahman, S. Chou, C. Zhong, J. Wang, D. Wexler and H. Liu, Solid State Ionics, 2010, 180, 1646.
42. X. H. Huang, J. P. Tu, C. Q. Zhang, X. T. Chen, Y. F. Yuan and H. M. Wu, Electrochim. Acta, 2007, 52, 4177.
43. A. J. Bard and L. R. Faulkner, Electrochemical Methods, Second ed., Wiley, New York, 2001, p. 321.
44. B. Jin, E. M. Jin, K. Park and H. Gu, Electrochem. Commun., 2008, 10, 1537.
45. A. Y. Shenouda and H. K. Liu, J. Electrochem. Soc., 2010, 157, A1183.

---