A new strategy for developing superior electrode materials for advanced batteries: using a positive cycling trend to compensate the negative one to achieve ultralong cycling stability

Abstract

In this communication, in order to develop superior electrode materials for advanced energy storage devices, a new strategy is proposed and then verified by the (Si@MnO)@C/RGO anode material for lithium ion batteries. The core idea of this strategy is the use of a positive cycling trend (gradually increasing Li-storage capacities of the MnO-based constituent during cycling) to compensate the negative one (gradually decreasing capacities of the Si anode) to achieve ultralong cycling stability. As demonstrated in both half and full cells, the as-prepared (Si@MnO)@C/RGO nanocomposite exhibits superior Li-storage properties in terms of ultralong cycling stability (no obvious increase or decrease of capacity when cycled at 3 A g⁻¹ after 1500 cycles) and excellent high-rate capabilities (delivering a capacity of ca. 540 mA h g⁻¹ at a high current density of 8 A g⁻¹) as well as a good full-cell performance. In addition, the structure of the electrodes is stable after 200 cycles. Such a strategy provides a new idea to develop superior electrode materials for next-generation energy storage devices with ultralong cycling stabilities.

---

Conceptual insights

This work introduces a completely new concept for developing superior electrode materials for advanced batteries. The core idea of this concept is to use a positive cycling trend (the gradually increased capacities during cycling) to compensate the negative one, which can contribute towards the achievement of ultralong cycling stabilities. In this communication, this concept is verified for the first time using a nanoscale (Si@MnO)@C/RGO composite as the anode material for a lithium ion battery. Upon cycling, for the as-prepared (Si@MnO)@C/RGO, the MnO nanoparticles can deliver a positive cycling trend, due to the gradual activation of new a Mn²⁺/Mn⁴⁺ redox couple, to compensate the gradually decreased capacities of the Si nanospheres. Electrochemical tests demonstrate that the (Si@MnO)@C/RGO composite exhibits not only superior Li-storage properties in terms of ultralong cycling stability without an obvious decrease or increase in capacity (e.g., the capacities can stabilize at around 800–820 mA h g⁻¹ within 1500 cycles at 3 A g⁻¹) and excellent high-rate capabilities (for example, it can achieve a Li-storage capacity of 470 mA h g⁻¹ at 10 A g⁻¹) in the half cells, but also good full-cell performances when coupled respectively with commercial LiFePO₄ and LiNi_0.6Co_0.2Mn_0.2O₂ cathodes. Not only has the present work successfully prepared one outstanding anode nanocomposite for lithium ion batteries, but also more importantly it provides a completely new strategy for developing superior electrode materials for advanced energy storage devices.

Introduction

The increasing demand for lithium ion batteries (LIBs) with a stable long-cycling lifespan, as well as high energy and power densities, is being driven by the rapid development of portable electronic devices, electric vehicles (EVs), plug-in hybrid EVs, large-scale energy storage, and so on.^1–3 In order to achieve such advanced LIBs, a variety of electrode materials have been studied in the last two decades.^4–8 For the selection of potential candidate electrode materials, high theoretical Li-storage capacity should be of great importance. However, most of the high-capacity materials usually exhibit gradually decreased Li-storage capacities (defined as a “negative cycling trend”) during cycling due to structural decay processes. For example, the specific capacities delivered by bulk Si will degrade rapidly to almost zero within 5 charge/discharge cycles,^9,10 although it has the highest theoretical Li-storage capacity of 4200 mA h g⁻¹ among all anode materials for LIBs. In order to improve the electrochemical properties of a particular electrode material with a high theoretical capacity but with a negative cycling trend, a currently common research strategy is to inhibit the decay of the electrochemical properties by improving the electrode structures through means including building specific nanostructures,^11–18 constructing highly conductive networks to optimize the electron transfer,^19–25 reserving suitable voids to accommodate volumetric expansion/contraction in the successive Li-insertion/extraction,^26–29 and so on.

In addition to the vast majority of negative cycling trends, some electrode materials were interestingly reported to show gradually increased Li-storage capacities (conversely defined as a “positive cycling trend”) due to some specific reasons; e.g. carbon nanotube arrays grown on a stainless steel substrate exhibited 250% capacity increment after 1200 cycles due to the change of structure and surface defects.³⁰ Unfortunately, such increased capacities could not be effectively used in the common full cells as the amount of cathode materials is fixed. Hence, it is still a difficult challenge, as well as a very significant and interesting one, to make such a positive cycling trend useful in storing Li in full cell applications.

Herein, we propose for the first time that the positive cycling trend can be employed to compensate the negatively charged ones of other materials working in the same voltage window via an easily scalable preparation processes, which provides a new approach towards developing superior electrode materials for advanced LIBs, or the possibility of expanding this approach for other energy storage devices. This concept was demonstrated by integrating MnO nanoparticles (NPs) and commercially available Si nanospheres (NSs) into one nanohybrid. Upon cycling, MnO NPs can deliver the gradually increased capacities originating from the gradual activation of one new Mn²⁺/Mn⁴⁺ redox couple,^31,32,33 to compensate the gradually decreased capacities of the Si NSs due to the pulverization and degradation of the electrode.^9,10 A proper combination of these two constituents in conjunction with the reduced graphene oxide (RGO) based network could deliver an ultrastable Li-storage property without any obvious capacity decrease or increase (e.g., the capacities were stabilized at about 800–820 mA h g⁻¹ within 1500 cycles at 3 A g⁻¹) but with an outstanding high-rate performance (for example, it can achieve a Li-storage capacity of about 470 mA h g⁻¹ at a very high current density of 10 A g⁻¹). The practical applications were demonstrated in two full-cell systems with commercial LiFePO₄ and LiNi_0.6Co_0.2Mn_0.2O₂ as the cathode materials, respectively. It should be noted that in addition to the MnO anode, several other electrode materials including anode (such as carbon nanotubes,³⁰ FeO_x@carbon nanowires,³⁴ MnO@ZnMn₂O₄/N–C nanorods,³⁵ and ZnCo₂O₄–ZnO–C composites³⁶) and cathode (such as a polyaryltriazine-derived composite³⁷) materials exhibit the similar “positive cycling trend”, implying that our proposed strategy is versatile for developing advanced electrode materials.

Results and discussion

In the designed nanohybrid, as illustrated in Fig. 1, many MnO NPs adhere to the surfaces of relatively large Si NSs, forming the primary Si@MnO subunit, which is further encapsulated with a thin carbon layer to obtain the (Si@MnO)@C structure. Due to the introduction of well-dispersed RGO nanosheets during the preparation processes, these (Si@MnO)@C building blocks are wrapped by the three-dimensional (3D) RGO conductive network to finally achieve the designed nanohybrid (hereafter abbreviated as (Si@MnO)@C/RGO according to its structural characteristics). Note that the proportions of Si NSs, MnO NPs and RGO-dominated carbonaceous materials in the prepared (Si@MnO)@C/RGO nanohybrids can be easily adjusted via controlling the usage amounts of the corresponding precursors. For the (Si@MnO)@C/RGO nanohybrid with optimized cycling properties discussed here, the mass ratio of Si, MnO and carbon is about 42 : 39 : 19, which is calculated from the results of inductively coupled plasma atomic emission spectroscopy (ICP-AES) and thermogravimetric analysis (TGA, Fig. S1, ESI†).

Fig. 1 Schematic of the designed (Si@MnO)@C/RGO nanohybrid.

The structure and morphology of the (Si@MnO)@C/RGO nanohybrid were firstly characterized using transmission electron microscopy (TEM), high-resolution TEM (HRTEM) and scanning electron microscopy (SEM). Fig. 2a and b show its typical TEM and HRTEM images, clearly disclosing that the MnO NPs of 10–20 nm adhere homogenously on the surface of ∼100 nm Si NSs. Furthermore, the Si@MnO subunits are coated by the partially graphitized carbon layers with a thickness of 3–5 nm. The SEM images (Fig. 2c and its inset) show that the (Si@MnO)@C building blocks have been embedded in the RGO-based conductive networks, forming secondary particles of tens of micrometers. The Barrett–Joyner–Halenda (BJH) fitting from the N₂ adsorption–desorption isotherm curves discloses the presence of mesopores in the (Si@MnO)@C/RGO nanohybrid with a wide size distribution centered at 10–40 nm (Fig. S2, ESI†), which may be originated from the voids in the 3D RGO networks and will be very beneficial in accommodating the successive volume changes of the active electrode materials. In addition, X-ray diffraction (XRD, Fig. 2d) and Raman spectroscopy (Fig. 2e) were further employed to check the phases of the (Si@MnO)@C/RGO nanohybrid. All diffraction peaks in the XRD pattern can be well indexed to Si and MnO. All the Raman peaks also originate from the Si and carbonaceous materials,³⁸ confirming that the as-prepared (Si@MnO)@C/RGO is composed of Si, MnO and carbon without any detectable impurities. Furthermore, the XRD peaks of MnO are broader than those of Si, indicating the smaller grain size of the MnO NPs in comparison to the Si NSs. Based on the Debye–Scherrer equation,³⁹ the grain sizes of the Si NSs and MnO NPs are about 50–100 nm and 10.5–12.5 nm, respectively. Both are consistent with the (HR)TEM observations.

Fig. 2 (a) TEM, (b) HRTEM and (c) SEM images, (d) XRD pattern and (e) Raman spectrum of the prepared (Si@MnO)@C/RGO nanohybrid. Insets of (b) and (c): The partially enlarged HRTEM and SEM images of (b) and (c), respectively.

The Li storage properties were firstly investigated in half cells and compared with two controls, MnO@C/RGO and Si@RGO. The phases of the two control samples were characterized as shown in Fig. S3 and S4 in the ESI.†Fig. 3a compares the cycling stability of all three samples cycled at 0.5 A g⁻¹. For Si@RGO derived from the commercial Si NSs, the specific capacities decrease rapidly to only about 420 mA h g⁻¹ from the initial capacity of 1345 mA h g⁻¹ after 150 cycles at 0.5 A g⁻¹. In contrast, for MnO@C/RGO, it shows gradually increased specific capacities from 665 to 1410 mA h g⁻¹ within 150 cycles, which should originate from the gradual activation of Mn²⁺/Mn⁴⁺ redox^31–33 and the strong interphase interaction between the MnO NPs and RGO nanosheets.⁴⁰ When such two constituents are integrated into one composite, the as-formed (Si@MnO)@C/RGO exhibits superior Li-storage stability without an obvious decrease or increase in specific capacity under the same test conditions. Fig. 3b shows the corresponding galvanostatic curves of the three samples at the 1st, 30th, 60th, 90th, 120th and 150th cycles. The curve profiles for (Si@MnO)@C/RGO overlap clearly, while those for the other two controls show obvious shifts. This confirms the outstanding cycling stability of (Si@MnO)@C/RGO, which can be also demonstrated at other current densities (e.g., 0.2 A g⁻¹, 2 A g⁻¹ and 5 A g⁻¹), as shown in Fig. S6 in the ESI.† In addition, (Si@MnO)@C/RGO exhibits a higher reversible voltage plateau (or mean voltage) in comparison to the Si@RGO control. A higher voltage plateau of anode materials can inhibit the Li-dendrite formation during the successive Li-insertion/extraction processes. Fig. S7 in the ESI† further shows the cyclic voltammetry (CV) curves of (Si@MnO)@C/RGO for the initial 5 cycles between 0.005 and 3 V vs. Li⁺/Li at a scan rate of 0.1 mV s⁻¹. In the profiles, the peak couples at 0.17 V/0.34 V and 0.38 V/0.53 V correspond to the alloying/dealloying reactions of the Si NSs, while the peak couple at 0.82 V/1.35 V should be attributed to the redox of Mn⁰/Mn²⁺.³³ From the inset of Fig. S7 in the ESI,† the oxidation process from the Mn²⁺ to Mn⁴⁺ species can also be clearly observed at about 1.89 and 2.1 V.³³ Furthermore, the as-prepared (Si@MnO)@C/RGO has an ultralong cycling life time. As shown in Fig. 3c, the specific capacities delivered by (Si@MnO)@C/RGO at 3 A g⁻¹ are at about 800–820 mA h g⁻¹ with Coulombic efficiencies of close to 100% for 1500 cycles.

Fig. 3 Comparison of (a) cycling performance and (b) the corresponding galvanostatic curves between Si@RGO, MnO@C/RGO and (Si@MnO)@C/RGO cycled at 0.5 A g⁻¹ after three cycles at 0.1 A g⁻¹. (c) The ultralong-term cycling preformance of (Si@MnO)@C/RGO at 3 A g⁻¹ after three cycles at 0.1 A g⁻¹.

In addition to the outstanding cycling stability, the as-prepared (Si@MnO)@C/RGO nanohybrid also exhibits excellent rate capabilities (Fig. S8, ESI†). For example, when the current densities increase from 0.1 A g⁻¹ to 8 A g⁻¹, a high capacity retention of 46.9% can be achieved. At a very high current density of 10 A g⁻¹, it can still deliver a specific capacity of ca. 470 mA h g⁻¹. In comparison, MnO@C/RGO delivers much lower specific capacities at the same testing current densities in the range from 0.1 to 10 A g⁻¹, implying that the incorporation of Si NSs into (Si@MnO)@C/RGO can increase the Li-storage capacities of the final composite. Additionally, (Si@MnO)@C/RGO also exhibits much higher capacity retentions in comparison to the Si@RGO control (e.g., 50.2% vs. 25.1% at 6 A g⁻¹).

There are other structural merits in addition to the compensation effects for the enhancement of electrochemical properties in the as-prepared (Si@MnO)@C/RGO. Firstly, the complete carbon coating and large number of voids in the 3D RGO networks have the ability to accommodate the volume change of the (Si@MnO)@C subunits during the successive Li-insertion/extraction. Fig. 4 shows the TEM images of the (Si@MnO)@C/RGO material after 200 cycles at 3 A g⁻¹. As disclosed, the Si@MnO particles are still coated by the carbon layers and wrapped in the RGO-based networks, although the Si NSs and MnO NPs have already merged together due to the successive lithiation/delithiation processes. Secondly, the flexible and highly conductive 3D scaffold composed of the RGO nanosheets and carbon coating can effectively enhance the charge transport. Such a 3D scaffold is also robust enough to maintain the structural integrity of the material during the successive Li-insertion/extraction processes.^16,29,31 This is also verified by the TEM image (Fig. 4a) of the (Si@MnO)@C/RGO material after 200 cycles at 3 A g⁻¹, in which the cycled secondary particles are still wrapped in the RGO networks, as well as the elemental mappings of the (Si@MnO)@C/RGO electrode after the same cycling tests (Fig. S9, ESI†), in which carbon as well as other elements are uniformly distributed. In addition, the RGO nanosheets may also contribute to the positive cycling trend³² and promote the Mn²⁺/Mn⁴⁺ redox reaction. Thirdly, the utilized Si material is a commercially available product, which indicates a scalable material preparation process and the possibility of practical applications as anode materials for lithium full cells.¹⁵

Fig. 4 (a) TEM and (b) HRTEM images of the (Si@MnO)@C/RGO electrodes after 200 cycles at 3 A g⁻¹.

To demonstrate the potential practical applications, we further evaluated the full-cell performance of the as-prepared (Si@MnO)@C/RGO via coupling with commercial LiFePO₄ and LiNi_0.6Co_0.2Mn_0.2O₂ cathode materials, respectively. Fig. S10 in the ESI† shows the galvanostatic curve of the LiFePO₄ material in the half cells cycled at 34 mA g⁻¹, which shows the typical charge/discharge profiles for the LiFePO₄ cathode material as reported previously.³⁹ After it was fabricated into (Si@MnO)@C/RGO//LiFePO₄ full cells, Fig. 5a shows the CV patterns of the initial three cycles recorded at 1 mV s⁻¹ between 1 and 3.6 V. It is disclosed that all curves are well overlapping, implying the excellent reversible charge/discharge processes of the as-fabricated (Si@MnO)@C/RGO//LiFePO₄ full cells. As shown in Fig. 5c, such full cells can deliver stable specific capacities without an obvious decay after 150 cycles at 5C (charging/discharging the battery for 1/5 hour), except for a slight decrease of about 7% during the initial 10 cycles. The capacity decay in the initial cycles suggests that further optimization is very essential for the full cell assembly. In addition to the good cycling performance, the as-fabricated (Si@MnO)@C/RGO//LiFePO₄ full cells also exhibit a superior rate performance as shown in Fig. 5b. At a high rate of 10C, the cells show a high capacity retention of ca. 74% in comparison to the capacity at 0.2C. Note that it is the coin-type full cells with a capacity of 3–4 mA h, which can be further increased via fabricating larger cells in future practices. The inset of Fig. 5c further shows that one flexible pouch full cell fabricated using (Si@MnO)@C/RGO and LiFePO₄ can power a light-emitting diode (LED) bulb. In addition, the battery performances of the (Si@MnO)@C/RGO//LiNi_0.6Co_0.2Mn_0.2O₂ full cells are illustrated in Fig. S11 in the ESI.† Such full cells exhibit a charging/discharging ability at 1.0–4.0 V with an excellent rate performance (e.g., a high capacity retention of ca. 69% when tested at 10C in comparison to 0.2C), which demonstrates the successful fabrication and preliminary evaluation of the (Si@MnO)@C/RGO//LiNi_0.6Co_0.2Mn_0.2O₂ full cells.

Fig. 5 Performances of the (Si@MnO)@C/RGO//LiFePO₄ full cells: (a) CV profiles at a scan rate of 1 mV s⁻¹, (b) galvanostatic curves at various rates from 0.2C to 10C, and (c) cycling performance and the corresponding coulombic efficiencies at a rate of 5C. The inset of (c) shows that one fabricated coin-type full cell can light up a bulb.

Conclusions

In summary, in order to develop superior electrode materials for advanced LIBs, a completely new strategy was firstly proposed and then demonstrated in a designed (Si@MnO)@C/RGO anode by using the positive cycling trend of one component to compensate the negative one of another, to achieve ultralong cycling stability. As demonstrated in both half and full cells, the as-prepared (Si@MnO)@C/RGO hybrid exhibits superior Li-storage properties in terms of ultralong cycling stability, excellent high-rate capabilities and good full-cell performance, suggesting the effectiveness of the strategy proposed here. In view of the ease of preparation, the present strategy could be easily extended to other anode as well as cathode composites for LIBs, and possibly to other energy storage devices including sodium ion batteries. For example, polyaryltriazine-based cathode materials can be employed as constituents for developing an ultrastable cathode composite according to the present strategy because of its remarkable positive cycling trend, above 2 times of capacity increment after 5000 cycles in 3–4.5 V vs. Li⁺/Li.³⁷ Therefore, not only has the present work successfully prepared a superior anode composite with ultralong cycling stability for LIBs, but more importantly, it also provides a completely new strategy for developing superior electrode materials for advanced energy storage devices.

Acknowledgements

This work was supported by the National Natural Science Foundation of China (51602048), the Science and Technology Program of Jilin Province (20140101087JC and 20150520027JH), and Singapore MOE AcRF Tier 1 grants RG2/13 and RG113/15.

Notes and references

1. J. W. Choi and D. Aurbach, Nat. Rev. Mater., 2016, 1, 16013.
2. J. B. Goodenough, Energy Storage Mater., 2015, 1, 158–161.
3. M. Armand and J. M. Tarascon, Nature, 2008, 451, 652–657.
4. M. N. Obrovac and V. L. Chevrier, Chem. Rev., 2014, 114, 11444–11502.
5. M. S. Whittingham, Chem. Rev., 2014, 114, 11414–11443.
6. B. C. Melot and J. M. Tarascon, Acc. Chem. Res., 2013, 46, 1226–1238.
7. P. Xiong, J. Zhu, L. Zhang and X. Wang, Nanoscale Horiz., 2016, 1, 340–374.
8. R. Chen, T. Zhao, X. Zhang, L. Li and F. Wu, Nanoscale Horiz., 2016 10.1039/c6nh00016a.
9. H. Kim, M. Seo, M. H. Park and J. Cho, Angew. Chem., Int. Ed., 2010, 49, 2146–2149.
10. X. H. Liu, L. Zhong, S. Huang, S. X. Mao, T. Zhu and J. Y. Huang, ACS Nano, 2012, 6, 1522–1531.
11. K. Zhang, X. Han, Z. Hu, X. Zhang, Z. Tao and J. Chen, Chem. Soc. Rev., 2015, 44, 699–728.
12. Q. Zhang, E. Uchaker, S. L. Candelaria and G. Cao, Chem. Soc. Rev., 2013, 42, 3127–3171.
13. X. Su, Q. Wu, J. Li, X. Xiao, A. Lott, W. Lu, B. W. Sheldon and J. Wu, Adv. Energy Mater., 2014, 4, 1300882.
14. A. T. Tesfaye, R. Gonzalez, J. L. Coffer and T. Djenizian, ACS Appl. Mater. Interfaces, 2015, 7, 20495–20498.
15. N. Liu, H. Wu, M. T. McDowell, Y. Yao, C. Wang and Y. Cui, Nano Lett., 2012, 12, 3315–3321.
16. J.-Z. Guo, F. Wan, X.-L. Wu and J.-P. Zhang, J. Mol. Sci., 2016, 32, 265–279.
17. Y. Tang, X. Rui, Y. Zhang, T. M. Lim, Z. Dong, H. H. Hng, X. Chen, Q. Yan and Z. Chen, J. Mater. Chem. A, 2013, 1, 82–88.
18. F. Wang, X. Wang, Z. Chang, Y. Zhu, L. Fu, X. Liu and Y. Wu, Nanoscale Horiz., 2016, 1, 272–289.
19. S. Xin, Y.-G. Guo and L.-J. Wan, Acc. Chem. Res., 2012, 45, 1759–1769.
20. K. Feng, W. Ahn, G. Lui, H. W. Park, A. G. Kashkooli, G. Jiang, X. Wang, X. Xiao and Z. Chen, Nano Energy, 2016, 19, 187–197.
21. M. Zhou, X. Li, B. Wang, Y. Zhang, J. Ning, Z. Xiao, X. Zhang, Y. Chang and L. Zhi, Nano Lett., 2015, 15, 6222–6228.
22. Z. Zhang, Y. Wang, W. Ren, Q. Tan, Y. Chen, H. Li, Z. Zhong and F. Su, Angew. Chem., Int. Ed., 2014, 126, 5265–5269.
23. W. Shi, J. Zhu, X. Rui, X. Cao, C. Chen, H. Zhang, H. H. Hng and Q. Yan, ACS Appl. Mater. Interfaces, 2012, 4, 2999–3006.
24. C. Xu, S. Peng, C. Tan, H. Ang, H. Tan, H. Zhang and Q. Yan, J. Mater. Chem. A, 2014, 2, 5597–5601.
25. L. P. Tan, Z. Lu, H. T. Tan, J. Zhu, X. Rui, Q. Yan and H. H. Hng, J. Power Sources, 2012, 206, 253–258.
26. N. Liu, Z. Lu, J. Zhao, M. T. McDowell, H. W. Lee, W. Zhao and Y. Cui, Nat. Nanotechnol., 2014, 9, 187–192.
27. X. Li, C. Yan, J. Wang, A. Graff, S. L. Schweizer, A. Sprafke, O. G. Schmidt and R. B. Wehrspohn, Adv. Energy Mater., 2015, 5, 1401556.
28. Q. Xiao, M. Gu, H. Yang, B. Li, C. Zhang, Y. Liu, F. Liu, F. Dai, L. Yang, Z. Liu, X. Xiao, G. Liu, P. Zhao, S. Zhang, C. Wang, Y. Lu and M. Cai, Nat. Commun., 2015, 6, 8844.
29. X.-L. Wu, Y.-G. Guo and L.-J. Wan, Chem. – Asian J., 2013, 8, 1948–1958.
30. C. Masarapu, V. Subramanian, H. Zhu and B. Wei, Adv. Funct. Mater., 2009, 19, 1008–1014.
31. D.-H. Liu, H.-Y. Lü, X.-L. Wu, B.-H. Hou, F. Wan, S.-D. Bao, Q. Yan, H.-M. Xie and R.-S. Wang, J. Mater. Chem. A, 2015, 3, 19738–19746.
32. D.-H. Liu, W. Li, F. Wan, C.-Y. Fan, Y.-Y. Wang, L.-L. Zhang, H.-Y. Lü, Y.-M. Xing, X.-H. Zhang and X.-L. Wu, ChemElectroChem, 2016, 3, 1354–1359.
33. Y. M. Sun, X. Hu, W. Luo, F. F. Xia and Y. Huang, Adv. Funct. Mater., 2013, 23, 2436–2444.
34. X. Li, Y. Ma, G. Cao and Y. Qu, J. Mater. Chem. A, 2016, 4, 12487–12496.
35. M. Zhong, D. Yang, C. Xie, Z. Zhang, Z. Zhou and X. H. Bu, Small, 2016 DOI:10.1002/smll.201601959.
36. Z. Li and L. Yin, J. Mater. Chem. A, 2015, 3, 21569–21577.
37. Y. Su, Y. Liu, P. Liu, D. Wu, X. Zhuang, F. Zhang and X. Feng, Angew. Chem., Int. Ed., 2015, 54, 1812–1816.
38. R. P. Wang, G. W. Zhou, Y. L. Liu, S. H. Pan, H. Z. Zhang, D. P. Yu and Z. Zhang, Phys. Rev. B: Condens. Matter Mater. Phys., 2000, 61, 16827.
39. X. L. Wu, Y. G. Guo, J. Su, J. W. Xiong, Y. L. Zhang and L. J. Wan, Adv. Energy Mater., 2013, 3, 1155–1160.
40. S. Zhang, L. Zhu, H. Songn, X. Chen and J. Zhou, Nano Energy, 2014, 10, 172–180.

---

Footnotes

† Electronic supplementary information (ESI) available. See DOI: 10.1039/c6nh00150e

‡ Equal contribution.

---