Synthesis of Ge/C composites as anodes using glucose as a reductant and carbon source for lithium-ion batteries

Abstract

Ge-based materials as anodes in lithium ion batteries (LIBs) having a large theoretical reversible capacity are needed to overcome the unstable structural and electrochemical properties and pulverization of the electrodes for high-performance LIBs. Here, we synthesized Ge/C composites as anodes for use in LIBs via heating a mixture of GeO₂ powder and glucose as both a reductant and carbon source at 900 °C under a nitrogen atmosphere. The data from X-ray diffraction (XRD), Raman spectroscopy, and transmission electron microscopy (TEM) shows that the as-prepared samples consist of crystalline Ge particles and an amorphous carbon phase. Compared to pure Ge, the Ge/C samples exhibit discharge capacities of ∼627.1 mA h g⁻¹, improved cyclability, and excellent rate properties at a current of 3200 mA g⁻¹.

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1. Introduction

Recently, important efforts have been made to develop secondary lithium-ion batteries (LIBs) for various energy storage applications.^1–4 For high-performance LIBs, electroactive materials used as an alternative to carbon as a typical anode should have high specific capacity, large energy density, and improved cycling performance.^3–6

In particular, to further satisfy the demand, many intensive efforts have been made to develop high capacity materials. Group IV materials have attracted significant interest as anodes for LIBs due to their high theoretical capacity.^7–17 Especially, silicon (Si) and germanium (Ge) exhibit large theoretical capacities of 4200 and 1623 mA h g⁻¹ (4.4 Li⁺ per silicon or germanium atom), respectively. Although Ge has a relatively lower gravimetric capacity than Si, Ge shows the following advantages over silicon: (1) the electrical conductivity of Ge is four orders of magnitude higher than that of Si due to its smaller energy bandgap. (2) The diffusion coefficient of Li-ion in Ge is approximately two orders of magnitude greater than that in Si. (3) The volume expansion during Li insertion into Ge is smaller than that of Si.^{2,3,5,18–23} Despite these advantages, Ge exhibited a rapid capacity loss, accompanied by extremely large irreversible capacity, with nearly 0 mA h g⁻¹ after several cycles. The most significant reason for the severe loss in capacity is the large volume expansion of about 370% occurring during Li alloying/de-alloying processes, resulting in the pulverization and exfoliation from the current collector, leading to poor cyclability and rapidly declining capacity.^23–25 The following studies have been carried out to avoid the rapid capacity loss caused by the pulverization of Ge: (1) changing the morphology (e.g., nanostructures of various dimensionalities, porous structures). (2) Combining Ge with an inactive/active matrix. (3) Making Ge-based composites with carbonaceous materials.^{22–24,26–35} In particular, the utilization of carbon in composite structures has received a great deal of attention due to the light weight and relatively low cost of carbon.²⁰ In the composites consisting of Ge as an electroactive material and carbon as a matrix, carbon plays a structural buffering role in minimizing the mechanical stress induced by the volume change of Ge.^22,29,30 In the previous studies, composite electrodes having carbon as a buffer layer showed both improved specific capacity and cyclability in LIBs.³³ However, since the synthesis process for Ge embedded in a carbon matrix is rather complicated, the access of dispersed active materials embedded in carbon matrices still needs to be addressed.³⁶

In this work, we synthesized composites containing crystalline Ge and amorphous carbon (denoted as Ge/C) for use as an anode in LIBs via heating a mixture of GeO₂ powder and D-glucose with different weight ratios under nitrogen atmosphere at 900 °C. To understand the relation between specific capacity and cyclic properties for Ge/C composites, the portion of Ge nanoparticles (NPs) in the carbon structure as a matrix was adjusted using the different weight ratios of GeO₂ powder to glucose in the precursors. These as-prepared Ge/C composites showed well-dispersed Ge NPs in the carbon matrix, thus exhibiting excellent LIB performance during the lithiation/delithiation process.

2. Experimental

2.1 Synthesis of Ge/C composites

The precursor for Ge/C composites was prepared by mixing GeO₂ powder (99.999%, Alfa Aesar) and D-glucose (98%, anhydrous, SAMCHUN). In order to manipulate the ratio of carbon as a matrix and Ge as an active material in the electrodes, the weight ratios of GeO₂ to D-glucose in the synthesis were varied at 1 : 2, 1 : 3, and 1 : 5, denoted as Ge/C-2, Ge/C-3, and Ge/C-5, respectively. Carbonization of D-glucose and the reduction of GeO₂ powder were carried out at 900 °C under a N₂ atmosphere for 3 h. For comparison, 0.2 g of GeO₂ powder in the absence of D-glucose was heated at 600 °C under an N₂/H₂ atmosphere to fabricate pure Ge without carbon (denoted as Ge-only).

2.2 Structural analysis

To analyze the crystal structure of the as-prepared samples, XRD analysis was performed using a Bruker X-ray diffractometer (D2 PHASER, Bruker AXS) with a Cu K_α source of λ = 0.15418 nm operated at 30 kV and 10 mA. Raman spectroscopy were performed using a high resolution Mirco-Raman spectrometer with an Ar ion laser with λ = 532.8 nm (LabRam Aramis, Horriba Jovin Yvon). Thermogravimetric analysis (TGA) were conducted using a thermal analyzer (SDTA851, Mettler Toledo) in the range of 25 to 800 °C under an air flow of 60 cm³ min⁻¹. The morphology and structure of the samples were confirmed using field emission-TEM operating at 300 kV (FE-TEM, Tecnai G2 F30 system).

2.3 Electrochemical measurement

The electrochemical measurement was conducted using a two-electrode coin cell system, which lithium foil was used as both the counter and reference electrodes (CR2032 type, Hohsen Corporation). To prepare the electrodes of the electrochemical measurement, 80 wt% active material, 10 wt% Ketjen black, and 10 wt% polyvinylidene difluoride (PVDF, Alfa Aesar) were homogeneously mixed in 1-methyl-2-pyrrolidinone (99%, Aldrich). The well-mixed slurry was pasted onto a copper foil as a current collector and dried in 100 °C vacuum oven for 12 h. For all electrodes, the average loading amount of active materials was 0.5–0.6 mg cm⁻². The electrodes of the cells were separated by a porous polypropylene membrane (Celgard 2400). The electrolyte solution contained 1.1 M LiPF₆ in a mixture of ethylene carbonate and dimethyl carbonate (Techno Semichem Co., Ltd). The cyclic voltammograms (CVs) were obtained using an AUTOLAB system with a scan rate of 0.1 mV s⁻¹. The charge–discharge cycling properties were measured in the range of 0.0 to 2.0 V vs. Li/Li⁺ using an instrument (PEBC 50.2 PnE Solution Co. Ltd). Electrochemical impedance spectra were obtained in the frequency range of 100 kHz to 0.01 Hz at 1.0 V using an AUTOLAB system.

3. Result and discussion

As shown in Fig. 1, Ge-only and Ge/C composites exhibited XRD peaks consisting of (111), (220), (311), (400), and (331), located at 27.28°, 45.31°, 53.68°, 66.02°, and 72.81°, respectively. In particular, the Ge state showed a diamond cubic structure (JCPDS #04-0545) in the metallic phase without any oxide structures. This results from the complete reduction of GeO₂ to Ge in the presence of D-glucose under an N₂ atmosphere. In contrast, in the Ge/C composites, the XRD peak at ∼25° in the (002) corresponding to carbon did not clearly appear, exhibiting the formation of an amorphous carbon phase. The amorphous carbon layer in the Ge/C can be beneficial to enhancing an electronic conductivity between particles and effectively decreasing transfer resistance of electrons, thus ensuring high initial specific capacity and rate capability.³⁷

Fig. 1 XRD patterns of the as-prepared Ge-only and Ge/C composites compared to the XRD reference data (JCPDS no. 04-0545).

Raman spectra of Ge-only, Ge/C-2, Ge/C-3, and Ge/C-5 were obtained, as shown in Fig. 2. The characteristic peaks corresponding to crystalline Ge and amorphous carbon were observed. A single sharp peak at 300 cm⁻¹, related to crystalline Ge, was detected in the absence of the GeO₂ phonon mode related to oxide phase.^32,38,39 Furthermore, compared to the presence of pure Ge peak in Ge-only, the Raman spectra of Ge/C-2, Ge/C-3, and Ge/C-5 contained the two characteristic peaks consisting of the G and D bands at ∼1598 and ∼1344 cm⁻¹, respectively (Fig. 2). In the Raman spectra, the G-band is associated with the E_2g vibration mode of sp² carbon crystal structures, whereas the D-band represents the structural defects and partially disordered structures on the graphitic plane.⁴⁰ The relative intensity ratio (I_D/I_G) values of Ge/C-2, Ge/C-3, and Ge/C-5 are 1.062, 1.045, and 1.059, respectively, indicative of amorphous carbon structure.^41,42 From XRD and Raman analysis, it is concluded that the as-prepared composites consist of crystalline Ge and amorphous carbon.

Fig. 2 Raman spectra associated with Ge and carbon in the as-prepared samples.

TEM analysis of the as-prepared Ge-only and Ge/C composites was conducted, as shown in Fig. 3, to investigate the encapsulation of crystalline Ge particles in the carbon. The Ge NPs in Ge/C-2, Ge/C-3, and Ge/C-5 were homogeneously mixed in the carbon without any serious agglomeration between the Ge NPs. On the other hand, Ge-only exhibited a large particle size due to the over-growth during the reducing process with GeO₂ in the absence of D-glucose. Furthermore, the Ge NPs exhibited the {111} facets with a d-spacing of 3.26 Å of the Ge metallic phase corresponding to a diamond cubic structure, as shown in Fig. 3(b). Furthermore, to identify the elemental distribution of the Ge/C composite, as shown in Fig. 3(i), mapping images of Ge and carbon were obtained using high-angle annular dark-field scanning TEM-EDX spectroscopy. It can be observed that the as-prepared sample has Ge and carbon homogeneously distributed in the composite structure. According to TEM analysis, the well-mixed Ge NPs in the carbon matrix are expected to exhibit improved electrochemical properties during the lithiation/delithiation process due to the prevention of the pulverization of the Ge NPs by the carbon material as a buffer.

Fig. 3 FE-TEM and HR-TEM images of (a and b) Ge-only, (c and d) Ge/C-2, (e and f) Ge/C-3, and (g and h) Ge/C-5. (i) Elemental mapping images of Ge and carbon for the Ge/C.

The weight ratio of Ge to C in the composites was determined using TGA as indicated in Fig. 4. In the case of Ge/C composites, the 1st weight loss corresponds to the evaporation of water in the range of 25 °C to 110 °C. The weight loss around at 500 °C is mainly due to the oxidation of carbon (C + O₂ → CO₂) and the subsequent weight gain after 600 °C is due to the oxidation of Ge (Ge + O₂ → GeO₂).²⁰ The Ge contents in Ge/C-2, Ge/C-3, and Ge/C-5 were determined to be 80, 66, and 53 wt%, respectively. On the other hand, Ge-only showed no loss of weight around at 500 °C due to the oxidation of carbon and rapid weight gain at 600 °C due to the oxidation of Ge.

Fig. 4 TGA curves of Ge-only, Ge/C-2, Ge/C-3, and Ge/C-5.

Fig. 5 shows scheme of a facile synthesis of the composites containing Ge and carbon via heating a mixture of GeO₂ powder and D-glucose at 900 °C under a nitrogen atmosphere. In the present research, D-glucose is believed to serve both as a carbon source for the in situ formation of carbon matrix with Ge NPs and as a reducing agent for GeO₂ into Ge as the following reactions:

GeO₂ + D-glucose → GeO₂ + C + H₂O + CO₂ (1)

GeO₂ + C → Ge/C + CO₂ + CO (2)

Fig. 5 Schematic illustration of the formation of Ge/C composites by heating GeO₂ powder in the presence of D-glucose at 900 °C under an N₂ atmosphere.

D-Glucose is carbonized under 200 °C (reaction (1)) and then the reduction reaction of GeO₂ by carbon occurs with the release of gas phases (CO₂, CO) at above 750 °C (reaction (2)), thus forming a porous Ge/C composite.^43,44

Moreover, D-glucose as a capping agent can prevent the overgrowth or agglomeration of the Ge NPs, as shown in the TEM image. The Ge/C composites containing different weight ratios of Ge to C as anodes in LIBs can be synthesized by varying the weight ratio of GeO₂ to glucose.

To reveal the Li-ion behavior of the as-prepared samples during the cycling process, the CVs were obtained in the range of 0.0 to 2.0 V vs. Li⁺/Li, as shown in Fig. 6. In the 1st CV curves, the electrodes exhibited irreversible reductions at 1–0.5 V due to the formation of a solid electrolyte interface layer. The deep cathode peak below 0.5 V represents a phase transition of Ge to Li_xGe_y alloys. The anodic peaks at 0.57 V represent a phase transition of Li_xGe_y to Ge due to lithium extraction from Ge. After the first cycle, the distinct cathodic and anodic current curves are well overlapped, suggesting that the structural integrity was maintained and the electrochemical reaction of Ge/C composites was stable. In the subsequent cycles, three reversible reduction peaks between 0.7 and 0 V are assigned to the lithium insertion into Ge to form Li_xGe_y alloys as follows: Ge → Li₉Ge₄ → Li₇Ge₂ → Li₁₅Ge₄ + Li₂₂Ge₅.^31,35

Fig. 6 CVs of (a) Ge-only, (b) Ge/C-2, (c) Ge/C-3, and (d) Ge/C-5 measured after the 1st, 2nd, and 3rd cycles in the voltage range of 2.0 to 0 V vs. Li/Li⁺ at a scan rate of 0.1 mV s⁻¹.

Fig. 7 shows the galvanostatic voltage profiles of the samples in the 1st, 2nd, 3rd, and 100th cycles at a current density of 200 mA g⁻¹. The 1st discharge capacities of Ge-only, Ge/C-2, Ge/C-3, and Ge/C-5 were 1467.6, 1922.5, 1798.4, and 1753.4 mA h g⁻¹, respectively, which were higher than the theoretical capacity (1623 mA h g⁻¹) of Ge as an anode. The initial coulombic efficiencies of Ge-only, Ge/C-2, Ge/C-3, and Ge/C-5 were 69.6, 31.8, 49.3, and 51.1%, respectively. The Ge/C composites exhibited the lower efficiency than that of Ge-only due to a large solid-electrolyte-interface (SEI) layer caused by the porous composite structure having an increased contact area between electrode and electrolyte.^45,46 As shown in the 1st voltage profiles of Ge/C composites (Fig. 7(b)–(d)), the plateaus apparently appeared at ∼0.7 V due to the dominant formation of the SEI layer, compared to Ge-only.

Fig. 7 Charge-discharge profiles of (a) Ge-only, (b) Ge/C-2, (c) Ge/C-3, (d) Ge/C-5 for the 1st, 2nd, 3rd, and 100th cycles between 2.0 and 0 V at a current density of 200 mA g⁻¹.

To clarify the low coulombic efficiencies of the composites with amorphous carbon compared to the Ge-only, the charge–discharge profiles and CVs of the cells were obtained using the electrodes fabricated with 80 wt% Ketjen black and 20 wt% polyvinylidene difluoride under the same condition as our paper, as shown in Fig. 8. The Ketjen black as an active material exhibited more severe irreversibility in the capacity caused by a large contact area between electrode and electrolyte due to its high specific surface area, as compared to Ge/C composites including amorphous carbon phase. Despite the low initial efficiency of Ge/C composites caused by amorphous carbon phase, however, after several cycles, Ge/C-3 and Ge/C-5 showed the coulombic efficiencies close to 100% whereas Ge/C-2 exhibited the lowest efficiency due to insufficient carbon content in the composite. Furthermore, Ge-only, Ge/C-2, Ge/C-3, and Ge/C-5 at the 100th cycle exhibited discharge capacities of 48.5, 412.5, 627.1, and 459.1 mA h g⁻¹, respectively, and the capacity retention ratios (100th/3rd cycle) of 4.9%, 43.9%, 65.1%, and 58.3%, respectively.

Fig. 8 (a) Charge-discharge profiles of the electrodes fabricated with Ketjen black for the 1st, 2nd, and 3rd cycles between 2.0 and 0 V at current density of 200 mA g⁻¹. (b) CVs of the electrodes fabricated with Ketjen black measured after the 1st, 2nd, and 3rd cycles in the voltage range of 2.0 to 0 V vs. Li/Li⁺ at a scan rate of 0.1 mV s⁻¹.

The cycling performance of the as-prepared anodes for 100th cycles at a current rate of 200 mA g⁻¹ is presented in Fig. 9. The Ge/C composites exhibited high discharge capacities and improved capacity retention ratios after 100th cycles. The rapid capacity fading of Ge-only is attributed to the pulverization of the electrode caused by the volumetric expansion during the cycling process. However, in particular, the Ge/C-3 having an optimal ratio of conductive carbon matrix exhibited considerably improved cyclability and coulombic efficiency.

Fig. 9 Specific discharge capacity and coulombic efficiency vs. cycle number of the samples at a current density of 200 mA g⁻¹ for 100 cycles.

To further evaluate the rate cycling performance of the samples as anodes, the discharge current rate was elevated from 100 to 3200 mA g⁻¹ (Fig. 10). The discharge capacities of Ge/C-2 were 2332.5, 815.0, 666.3, 585.0, 510.0, and 405.0 mA h g⁻¹ at current densities of 100, 200, 400, 800, 1600, and 3200 mA g⁻¹, respectively; the discharge capacities of Ge/C-3 were 2731.9, 776.6, 664.1, 603.1, 525.0 and 412.5 mA h g⁻¹ at current densities of 100, 200, 400, 800, 1600, and 3200 mA g⁻¹, respectively; the discharge capacities of Ge/C-5 were 2047.7, 658.0, 562.5, 502.3, 446.6, and 353.4 mA h g⁻¹ at current densities of 100, 200, 400, 800, 1600, and 3200 mA g⁻¹, respectively; the discharge capacities of Ge-only were 1665.4, 535.7, 229.4, 99.2, 37.9, and 22.4 mA h g⁻¹ at current densities of 100, 200, 400, 800, 1600, and 3200 mA g⁻¹, respectively. As the discharge rate increased, the Ge/C samples exhibited the increased specific discharge capacities more than Ge-only. It has been reported that pure Ge anodes for LIBs show a remarkable reduction of both capacity and rate performance due to their cracking and pulverization during cycling.^47,48 On the other hand, the Ge/C composites could maintain the electrochemical performance due to carbon matrix to prevent the volume change during the lithiation/delithiation process. Especially, Ge/C-3 having an optimal amount of amorphous carbon phase in the composite showed the highest capacity retention and improved rate cycling performance.

Fig. 10 Rate cycling performance of the samples with different current densities.

To further understand the improved LIB performance of all samples, Nyquist plots of the electrodes were obtained after the 1st and 100th cycles at 1.0 V, as shown in Fig. 11(a) and (b), respectively. The values of R_ct of Ge-only, Ge/C-2, Ge/C-3, and Ge/C-5 after the 1st cycles were 15.0, 8.2, 11.5, and 20.7 Ω, respectively. After the 100th cycles, the values of R_ct of Ge-only, Ge/C-2, Ge/C-3, and Ge/C-5 were 70.0, 34.0, 30.6, and 25.6 Ω, respectively. The values of R_ct of all composite electrodes after the 100th cycles were lower than that of Ge-only, as a result of preventing the pulverization of the Ge anode. On the other hand, a continuous formation of SEI layers caused by pulverization in Ge-only can be responsible for the high R_ct, compared to Ge/C composites. Plots of Z_re versus square root of frequency (ω^−1/2) for the electrodes were obtained in the low frequency range to evaluated the diffusion behaviour of Li ions within the electrodes (Fig. 11(c) and (d)).^49,50 Using the relationship between Z_re and ω^−1/2, the Li-ion diffusion coefficients (D_Li) for Ge-only, Ge/C-2, Ge/C-3, and Ge/C-5 are determined as 0.505 × 10⁻¹², 1.375 × 10⁻¹², 9.643 × 10⁻¹², 2.382 × 10⁻¹² cm² s⁻¹ after the 1st cycle, respectively, and 1.774 × 10⁻¹², 2.411 × 10⁻¹², 3.233 × 10⁻¹², and 2.419 × 10⁻¹² cm² s⁻¹ after the 100th cycle, respectively, indicating a rapid Li-ion diffusion process for the Ge/C composites due to the conductive carbon matrix. As a result, the improved properties of the Ge/C composites as anodes for lithium-ion reaction, i.e. high capacity retention, improved cycle life, and high rate performance, are attributed to the low transport resistance and high diffusion coefficient of the lithium ion in the composite electrode consisting of Ge particles as an active material and carbon as a matrix. In particular, Ge/C-3 having an optimal amount of carbon as a matrix exhibited the most improved LIB performance. This is attributed to the facilitation of Ge lithiation and prevention of electrode pulverization through amorphous carbon matrix.

Fig. 11 Nyquist plots obtained (a) after the 1st cycle and (b) after the 100th cycle at 1.0 V between 100 kHz and 10 mHz. Plots of Z_re vs. square root of frequency (ω^−1/2) for the samples (c) after the 1st cycle and (d) after the 100th cycle in the low frequency range.

4. Conclusions

In summary, we prepared Ge/C composites as anodes for LIBs by heating a mixture of GeO₂ powder and D-glucose as a carbon source at 900 °C under a nitrogen atmosphere. Since the as-prepared composites consisted of crystalline Ge NPs and amorphous carbon phases and revealed homogeneous distribution of Ge NPs in a carbon matrix, the Ge/C samples exhibited relatively improved performance compared to Ge-only. Especially, the Ge/C-3 showed enhanced properties for LIBs due to an optimal amount of carbon among these composite anodes for the LIBs.

Acknowledgements

This research was supported by Basic Science Research Program through the National Research Foundation of Korea (NRF) funded by the Ministry of Education (NRF-2013R1A1A2012541).

Notes and references

1. J. M. Tarascon and M. Armand, Nature, 2001, 414, 359–367.
2. S. Fang, L. Shen, H. Zheng and X. Zhang, J. Mater. Chem. A, 2015, 3, 1498–1503.
3. A. Magasinski, P. Dixon, B. Hertzberg, A. Kvit, J. Ayala and G. Yushin, Nat. Mater., 2010, 9, 353–358.
4. R. Hu, H. Liu, M. Zeng, J. Liu and M. Zhu, Chin. Sci. Bull., 2012, 57, 4119–4130.
5. S. Liu, J. Feng, X. Bian, Y. Qian, J. Liu and H. Xu, Nano Energy, 2015, 13, 651–657.
6. L.-F. Cui, Y. Yang, C.-M. Hsu and Y. Cui, Nano Lett., 2009, 9, 3370–3374.
7. M. H. Kim, S. H. Ahn, J. W. Park and J. A. Ascencio, J. Korean Phys. Soc., 2006, 49, 1107–1110.
8. L. C. L. Yang, Q. S. Gao, L. Li, Y. Tang and Y. P. Wu, Electrochem. Commun., 2010, 12, 418–421.
9. L. Y. Lim, N. Liu, Y. Cui and M. F. Toney, Chem. Mater., 2014, 26, 3739–3746.
10. W.-M. Zhang, J.-S. Hu, Y.-G. Guo, S.-F. Zheng, L.-S. Zhong and W.-G. Song, et al., Adv. Mater., 2008, 20, 1160–1165.
11. S.-H. Woo, S. J. Choi, J.-H. Park, W.-S. Yoon, S. W. Hwang and D. Whang, J. Electrochem. Soc., 2013, 160, A112–A116.
12. M. Yoshio, H. Wang, K. Fukuda, T. Umeno, N. Dimov and Z. Ogumi, J. Electrochem. Soc., 2002, 149, A1598–A1603.
13. W.-S. Chang, C.-M. Park, J.-H. Kim, Y.-U. Kim, G. J. Jeong and H.-J. Sohn, Energy Environ. Sci., 2012, 5, 6895–6899.
14. Z. Li, Y. Wang, H. Sun, W. Wu, M. Liu and J. Zhou, et al., J. Mater. Chem. A, 2015, 3, 16057–16063.
15. M. Wu, J. Qiu, J. Liu, M. Tan, Z. Li and W. Wu, et al., RSC Adv., 2014, 4, 25189–25194.
16. B. Guo, J. Shua, Z. Wanga, H. Yang, L. Shi and Y. Liu, et al., Electrochem. Commun., 2008, 10, 1876–1878.
17. Z. Li, M. Wu, G. Wu, D. Liu, W. Wu and B. Jiang, et al., J. Mater. Chem. A, 2014, 2, 7471–7477.
18. D. Li, C. Feng, H. kun Liu and Z. Guo, J. Mater. Chem. A, 2015, 3, 978–981.
19. J. D. Ocon, J. K. Lee and J. Lee, Appl. Chem. Eng., 2014, 25, 1–13.
20. D. T. Ngo, R. S. Kalubarme, H. T. T. Le, J. G. Fisher, C.-N. Park and I.-D. Kim, et al., Adv. Funct. Mater., 2014, 24, 5291–5298.
21. M. H. Park, K. Kim, J. Kim and J. Cho, Adv. Mater., 2010, 22, 415–418.
22. J. Liu, K. Song, C. Zhu, C. Chen, J. Maier and Y. Yu, et al., ACS Nano, 2014, 8, 7051–7059.
23. D.-J. Xue, S. Xin, Y. Yan, K.-C. Jiang, Y.-X. Yin and Y.-G. Guo, et al., J. Am. Chem. Soc., 2012, 134, 2512–2515.
24. H. Kim, Y. Son, C. Park, J. Cho and H. C. Choi, Angew. Chem., Int. Ed., 2013, 52, 5997–6001.
25. K. H. Seng, M. H. Park, Z. P. Guo, H. K. Liu and J. Cho, Nano Lett., 2013, 13, 1230–1236.
26. Y. Yu, C. Yan, L. Gu, X. Lang, K. Tang and L. Zhang, et al., Adv. Energy Mater., 2013, 3, 281–285.
27. J. Hwang, C. Jo, M. G. Kim, J. Chun, E. Lim and S. Kim, et al., ACS Nano, 2015, 9, 5299–5309.
28. J.-G. Ren, Q.-H. Wu, H. Tang, G. Hong, W. Zhang and S.-T. Lee, J. Mater. Chem. A, 2013, 1, 1821–1826.
29. W. Li, J. Zheng, T. Chen, T. Wang, X. Wang and X. Li, Chem. Commun., 2014, 50, 2052–2054.
30. Y. Xu, G. Yin, Y. Ma, P. Zuo and X. Cheng, J. Mater. Chem., 2010, 20, 3216–3220.
31. S. Yoon, C.-M. Park and H.-J. Sohn, Electrochem. Solid-State Lett., 2008, 11, A42–A45.
32. K. H. Seng, M. H. Park, Z. P. Guo, H. K. Liu and J. Cho, Angew. Chem., Int. Ed., 2012, 51, 5657–5661.
33. C. Wang, J. Ju, Y. Yang, Y. Tang, J. Lin and Z. Shi, et al., J. Mater. Chem. A, 2013, 1, 8897–8902.
34. X.-L. Wu, Y.-G. Guo and L.-J. Wan, Chem.–Asian J., 2013, 8, 1948–1958.
35. C. Zhong, J.-Z. Wang, X.-W. Gao, D. Wexler and H.-K. Liu, J. Mater. Chem. A, 2013, 1, 10798–10804.
36. G. Jo, I. Choi, H. Ahn and M. J. Park, Chem. Commun., 2012, 48, 3987–3989.
37. H. Tang, L. Zan, W. Mao and Z. Tang, J. Electroanal. Chem., 2015, 751, 57–64.
38. G. Cui, L. Gu, N. Kaskhedikar, P. a. van Aken and J. Maier, Electrochim. Acta, 2010, 55, 985–988.
39. G. Cui, L. Gu, L. Zhi, N. Kaskhedikar, P. a. Van Aken and K. Müllen, et al., Adv. Mater., 2008, 20, 3079–3083.
40. X. Xu, H. Tan, K. Xi, S. Ding, D. Yu and S. Cheng, et al., Carbon, 2015, 84, 491–499.
41. H. Lee, H. Kim, S.-G. Doo and J. Cho, J. Electrochem. Soc., 2007, 154, A343–A346.
42. C. Li, G. Fang, L. Yuan, N. Liu, L. Ai and Q. Xiang, et al., Nanotechnology, 2007, 18, 155702–155705.
43. M. Hurtta, I. Pitkänen and J. Knuutinen, Carbohydr. Res., 2004, 339, 2267–2273.
44. D. T. Ngo, H. T. T. Le, C. Kim, J.-Y. Lee, J. G. Fisher and I.-D. Kim, et al., Energy Environ. Sci., 2015, 8, 3577–3588.
45. K. Evanoff, A. Magasinski, J. Yang and G. Yushin, Adv. Energy Mater., 2011, 1, 495–498.
46. Z. Y. Zeng, J. P. Tu, X. H. Huang, X. L. Wang, X. B. Zhao and K. F. Li, Electrochem. Solid-State Lett., 2008, 11, A105–A107.
47. X. Li, J. Liang, Z. Hou, W. Zhang, Y. Wang and Y. Zhu, et al., J. Power Sources, 2015, 293, 868–875.
48. F.-W. Yuan, H.-J. Yang and H.-Y. Tuan, ACS Nano, 2012, 6, 9932–9942.
49. H. Liu, Q. Cao, L. J. Fu, C. Li, Y. P. Wu and H. Q. Wu, Electrochem. Commun., 2006, 8, 1553–1557.
50. Z. Zhang, J.-Z. Wang, S.-L. Chou, H.-K. Liu, K. Ozawa and H. Li, Electrochim. Acta, 2013, 108, 820–826.

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