MnO@C nanorods derived from metal-organic frameworks as anode for superiorly stable and long-life sodium-ion batteries

Journal of Alloys and Compounds 710 (2017) 575e580

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MnO@C nanorods derived from metal-organic frameworks as anode for superiorly stable and long-life sodium-ion batteries Xiaojie Zhang a, Guang Zhu b, Dong Yan a, Ting Lu a, Likun Pan a, * a b

Shanghai Key Laboratory of Magnetic Resonance, School of Physics and Materials Science, East China Normal University, Shanghai, 200062, China Anhui Key Laboratory of Spin Electron and Nanomaterials, Suzhou University, Anhui, Suzhou, 234000, China

a r t i c l e i n f o

a b s t r a c t

Article history: Received 7 December 2016 Received in revised form 25 March 2017 Accepted 27 March 2017 Available online 29 March 2017

Porous MnO@C nanorods were synthesized simply by annealing Mn-based metal-organic frameworks precursor. The morphology, structure and electrochemical performance of MnO@C hybrid were characterized by scanning electron microscopy, nitrogen adsorption/desorption isotherms, galvanostatic charge/discharge tests, cyclic voltammetry and electrochemical impendence spectroscopy. When used as anode material for sodium-ion batteries, the MnO@C hybrid exhibits a high reversible specific capacity of 260 mAh g
1 after 100 cycles at a current density of 50 mA g
1. When the current density is increased to 2 A g
1, the MnO@C delivers a superior long-life cycling performance with a capacity of 140 mAh g
1 at very high current density of 2 A g
1. The excellent electrochemical performance of MnO@C can be attributed to its unique porous structure with MnO nanoparticles embedded in carbon matrix, which can apparently increase the electrical conductivity and buffer the volume change during the charge/ discharge process. © 2017 Elsevier B.V. All rights reserved.

Keywords: MnO nanorods Carbon coating Metal-organic frameworks Sodium-ion batteries Long-life cycling

1. Introduction Sodium ion batteries (SIBs), as a potential alternative to lithium ion batteries (LIBs), have attracted increasing attention in recent years for future large-scale energy-storage application because of low cost and high abundance of sodium, and its similar chemical and physical properties with lithium [1e5]. However, graphite, as commercial anode material in LIBs, delivers an unsatisfied capacity when used in SIBs owing to the large size of sodium ion [6]. Hence, it is urgent to explore suitable host materials for sodium ion storage in order to improve the electrochemical performance of SIBs [7]. Up to date, many attempts have been made to explore favorable anode materials for SIBs, including various carbonaceous materials [8], alloy materials (Sn [9], Sb [10]), metal sulfides (MoS2 [11], SnS2 [12], FeS [13]), phosphides [14] and transition metal oxides (CuO [15], TiO2 [16], Co3O4 [17], MnOx [18]). Among them, MnO is regarded as a superior choice owing to its significant advantages such as high theoretical capacity due to conversion reaction mechanism, natural abundance, low cost and environmental benignity [19,20]. Nevertheless, the application of MnO as anode material is hindered by the

* Corresponding author. E-mail address: [email protected] (L. Pan). http://dx.doi.org/10.1016/j.jallcom.2017.03.314 0925-8388/© 2017 Elsevier B.V. All rights reserved.

obstacles of intrinsically low electrical conductivity and large volume and structure changes during the charge/discharge process, which will lead to dramatic capacity fading, poor cycling performance and rate capability. Currently, great efforts have been made to fabricate MnO-based electrode materials with unique structures to address these drawbacks. One highly effective way is to coat the metal oxides with conductive carbon materials, which can effectively accommodate the volume change and improve the electrical conductivity of conversion-type electrode materials [21e23]. Another common approach is to prepare nano/submicro-scale materials to shorten the diffusion distance of sodium ions and electrons and provide more contact area between electrolyte and electrode [24]. The third strategy is to construct porous structure which can relieve the volume change and increase the active reaction sites, resulting in the enhanced electroactivity and cycling performance of the electrode materials [8,25]. Therefore, the combination of carbon coating and porous nanostructure should be an effective strategy to develop high-performance hybrid composites for SIBs. Porous materials derived from metal-organic frameworks (MOFs) precursors, such as porous carbons [26], metal oxides [27], metal oxide/carbon composites have been successfully used in LIBs with excellent performance [26,28e30]. The successful applications of MOFs derivatives in LIBs have aroused extensive interest in

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their potential applications in SIBs [15]. Wang et al. prepared nitrogen-doped carbon-coated Co3O4 nanoparticles from MOFs ZIF67 and found that they exhibited a high reversible capacity of 506 mAh g
1 at 100 mA g
1 as anode for SIBs [31]. Despite the above progress to date, the exploration of porous materials, especially metal oxide/carbon hybrids is far from enough due to their potential ability for high performance anode materials for SIBs. Herein, a facile and scalable strategy is developed to fabricate MnO nanoparticles embedded in porous carbon matrix (MnO@C) by annealing Mn-based MOFs. Combining the advantages of ultrafine MnO nanocrystals with an elastic carbon matrix, and high porosity, the as-synthesized MnO@C hybrid structure exhibits excellent cycling stability and long-life cycling performance for SIBs. 2. Experimental 2.1. Preparation of MnO@C nanorods All agents were of analytical grade and used without further purification. In a typical synthesis, Mn(CH3COO)2$4H2O (2 mmol) and PVP (K-30, 1.0 g) were dissolved in a CH3CH2OH and H2O (100 ml, v/v 1:1) system under stirring to form a transparent solution A. Trimesic acid (BTC, 4 mmol) was also dissolved in a CH3CH2OH and H2O (100 ml, v/v 1:1) system under stirring. And then BTC solution was added into solution A drop by drop. After stirring for 10 min, the resulting solution was incubated at room temperature without interruption for 24 h. The white precipitates (Mn-BTC MOFs) were collected by centrifugation, washed several times with CH3CH2OH and H2O, and finally dried in an oven at 60  C for overnight. Finally, the obtained Mn-BTC MOFs were calcined at 700  C for 2 h under N2 atmosphere with a temperature ramp rate of 2  C min
1 to yield the target product (MnO@C sample). For comparison, pure MnO was prepared by calcining the precursor Mn-BTC MOFs at 450  C under air atmosphere for 2 h with a

temperature ramp rate of 2  C min
1, and then the obtained product was further calcined in an atmosphere of flowing 5% H2/N2 at the same temperature for 10 h. 2.2. Characterizations Field emission scanning electron microscopy (FESEM) images were obtained using a Hitachi S-4800 microscope. Transmission electron microscopy (TEM) images and selected area electron diffraction (SAED) were taken on a JEOL-2010 instrument. X-ray diffraction (XRD, Holland Panalytical PROPW3040/60) with Cu Ka radiation (V ¼ 30 kV, I ¼ 25 mA, l ¼ 0.15418 nm) was used for analyzing the structure of the samples. Raman spectra were collected using a confocal Raman microscope (DXR, Thermo-Fisher Scientific) with a 532 nm argon-ion laser. N2 adsorption/desorption isotherms were measured at 77 K using an ASAP 2020 Accelerated Surface Area and Porosimetry System (Micrometities, Norcross, GA) and used to calculate the specific surface area and pore size distribution. The samples were degassed under vacuum at 200  C overnight before analysis. 2.3. Electrochemical measurements Electrochemical properties of the as-prepared samples were evaluated as anode materials of SIBs in CR2032-type coin cells. The electrodes were fabricated by grinding porous MnO@C or MnO as active materials (70 wt%), Super P as conducting agent (20 wt%) and carboxyl methyl cellulose dissolved in deionized water as binder (10 wt%) to form a homogenously slurry. Then the slurry was doctor-bladed onto a copper foil (used as current collector) by an automatic thick film coater, and then dried under vacuum at 120  C for 24 h to remove excess water completely. The cells were fabricated in a glove box (MB-10-compact, MBRAUN) filled with Ar atmosphere where the contents of oxygen and water were less than 0.5 ppm. Pure metal sodium foil was used as both reference and

Fig. 1. (a) FESEM image, and (b, c, d) TEM images at different magnifications of MnO@C hybrid sample.

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counter electrodes, and whatman glass fiber membrane was used as separator. The electrolyte used in the coin cells was composed of 1 mol l
1 NaClO4 dissolved in ethylene carbonate and propylene carbonate (1:1, w/w) with the addition of 5% fluoroethylene carbonate. Cycling tests were carried out using a LAND2001A battery test system at a current density of 50 mA g
1 in a voltage range of 0.005e3.0 V (vs. Naþ/Na) unless otherwise specified. The specific capacity of the anodes was calculated based on the total mass of the active materials. The cyclic voltammetry (CV) was recorded on an electrochemical workstation (AUTOLAB PGSTA302N) between 0.005 and 3.0 V at a scan rate of 0.2 mV s
1. The electrochemical impedance spectroscopy (EIS) measurement was conducted on the same electrochemical workstation after 100 charge/discharge cycles over the frequency ranging from 0.1 to 100 kHz with an ac amplitude voltage of 5 mV. 3. Results and discussion FESEM image shows that the Mn-BTC MOFs display a rod-like morphology with a diameter of ~200 nm and a length of 3e10 mm, and their surface are very smooth, as shown in Fig. S1(a). The samples can retain the pristine shape after calcining in N2 atmosphere (Fig. 1(a)). TEM images at low (b) and high (c) magnifications demonstrate that the surface of porous MnO@C becomes rough due to the release of CO2 and H2O. In more detail, from Fig. 1(c) it can be seen that the MnO nanoparticles are 3e5 nm in size and densely and uniformly distributed in elastic porous carbon matrix. The enlarged TEM image in Fig. 1(d) shows that MnO nanocrystals with an interplane distance of 0.21 nm, corresponding to the (311) lattice plane, are well distributed in the amorphous carbon matrix. Moreover, the SAED pattern (Fig. S2) indicates that the as-prepared MnO nanoparticles are polycrystalline. The XRD pattern of Mn-BTC is consistent with that reported previously (Fig. S3(a)) [32]. The main peaks of XRD pattern can be identified to the orthorhombic MnO phase (JCPDS No. 04-0326.), as shown in Fig. 2(a). In addition, a broad peak at about 24 belongs to the (002) plane of carbon, indicating the presence of amorphous carbon, which agrees with the results of TEM measurement. From the FESEM image of MnO in Fig. S1(b), it is found that the rod-like morphology of Mn-BTC MOFs is not maintained and a number of MnO particles are formed. As shown in Fig. S3(b), the XRD pattern of MnO shows that all the diffraction peaks correspond to the orthorhombic MnO phase [33]. Fig. 2(b) displays the Raman spectra of MnO@C and MnO. The peak at 640 cm
1 in the Raman spectra of both samples is assigned to Mn-O vibration. Moreover, two obvious peaks appear at around 1345 and 1590 cm
1 in Raman spectra of MnO@C, corresponding to the feature peaks of the disordered carbon (D-band) and the ordered graphitic carbon (G-band), respectively. The stronger G-band than D-band indicates a high graphitization degree of carbon in MnO@C, which is beneficial to the increase of the electrical conductivity [34,35]. The porosity and BET surface areas of MnO@C and MnO were further investigated by N2 adsorption and desorption isotherms, as shown in Fig. 3(a) and (b). The MnO@C shows a typical II isotherm along with obvious H2 hysteresis, which is the typical feature of mesoporous materials. Furthermore, a fraction of macropores (>50 nm) also exist in the MnO@C sample. However, the MnO exhibits a typical III isotherm, indicating the existence of a large part of macropores. The pore size distributions are illustrated in Fig. 3(b), and the mean pore sizes of MnO@C and MnO are 3.85 and 36.95 nm, respectively. The specific surface area of MnO@C is 175.23 m2 g
1, which is more than ten times larger than that of MnO (14.86 m2 g
1). Apparently, the mesoporous structure will be more beneficial to the alleviation of the volume changes and large specific surface area will make a good contact between active materials and electrolyte, resulting in

Fig. 2. (a) XRD pattern of MnO@C, (b) Raman spectra of MnO@C and MnO.

excellent cycling performance and fast transport of sodium ions [36]. CV curves of MnO@C electrode are scanned at a rate of 0.2 mV s
1 in a range of 0.005e3 V, as shown in Fig. 4(a). In the first discharge cycle, an irreversible reduction peak at around 0.25 V appears due to the decomposition of electrolyte and the formation of solid electrolyte interface (SEI) layer on the surface of electrode. The peak at 0.01 V corresponds to the complete reduction of Mn2þ to Mn0, which can be confirmed by ex-situ XRD pattern at fully discharged state [37]. From the second sweep, a sharp peak around 0.75 V in the discharge process and a smooth peak at about 0.4 V in the charge scan can be ascribed to the reversible redox conversion reaction from sodium and MnO to form metallic Mn and Na2O, which is similar to the results for LIBs [38,39]. It is worth noting that an oxidation peak at about 2.2 V shows that Mn2þ is reoxidized to a higher oxidation state (Mn3þ) [40]. Ex-situ XRD measurement was usually used to understand the reaction mechanism for the typical MnO@C hybrid electrode during the sodiation/desodiation process. As shown in Fig. 4(b), when the cell is discharged to 1.5 V, the diffraction peak of MnO disappears, and metal Mn phase at 38 is formed when the cell is discharged to 0.8 V. When the cell is fully discharged to 0.005 V, metallic Mn phases at 45 and 66 are formed due to the electrochemical reaction. Whereas in the corresponding charging process, the metal Mn phases gradually

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Fig. 3. (a) N2 adsorption-desorption isotherms and (b) pore size distributions of MnO@C and MnO.

disappear, and when the cell is fully charged state, the MnO can be regenerated. It can be seen that the XRD pattern (Fig. 4b) is almost recovered to that of pristine state after the cell is fully charged to 3 V [21]. In addition, the weak diffraction pattern of Mn2O3 is formed when the cell is charged to 3 V, indicating the enhanced kinetics of conversion reactions during cycling. The result is in agreement with the CV analysis. It is noteworthy that from the second cycle the CV curves almost overlap, indicating a good reversibility of the electrochemical reaction for SIBs. In order to further understand the reason for the different electrochemical performances between MnO@C and MnO electrodes, EIS spectra were conducted after 100 cycles, as shown in Fig. S4. The Nyquist plots can be fitted based on the equivalent circuit model, as shown in the inset of Fig. S4. Rf in the high frequency region represents the SEI layer resistance. The large semicircle in the medium frequency region is related to charge transfer resistance (Rct) and the straight line in the low frequency region corresponds to the Warburg impedance of sodium ion diffusion, similar to that in LIBs. The MnO@C exhibits a Rct value of 150 U, which is considerably smaller than that of MnO (383 U). The low Rct can favor the electron transfer and diffusion, which is beneficial to the battery performance. Fig. 5(a) shows the cycling performance with corresponding coulombic efficiency (CE) of MnO@C and MnO at a current density of 50 mA g
1. It can be seen that the initial discharge capacity of MnO (648.9 mAh g
1) is higher than that of MnO@C (560 mAh g
1). However, the MnO electrode exhibits poor cycling stability with

Fig. 4. (a) Initial four cycles CV curves of MnO@C, (b) XRD pattern of MnO@C electrode at different states.

sharp capacity loss and the capacity is only remained at about 20 mAh g
1 after 50 cycles due to the huge volume expansion of MnO during the sodiation/desodiation process. On the contrary, the MnO@C electrode shows excellent cycling stability. The first discharge and charge capacities of the MnO@C electrode are 560 and 233 mAh g
1, respectively, and the initial CE is 41.6%. The relative low initial CE can be attributed to the irreversible capacity loss, including the decomposition of electrolyte and the formation of SEI film on the surface of electrode. From the fifth cycle, the MnO@C electrode reveals excellent cycling stability and the CE reaches over 99%. The capacity of MnO@C electrode shows a gradual increase during cycling at the beginning 10 cycles, which is attributed to the activation of the porous structure and the increasing electrical conductivity of the MnO@C electrode during cycling, which can be confirmed by the EIS measurement at different cycles (Fig. S5(a)) [40,41]. In addition, the pores in the electrode should be helpful for more electrolyte to gradually access the internal porous structure and provide more active reaction sites of the electrode materials, which results in an increase in capacity during the charge/discharge processes. Fig. S5(b) shows the quasiequilibrium redox potential of the MnO@C electrode via the galvanostatic intermittent titration technique (GITT) at a current density of 100 mA g
1. The shape of the GITT curve is very similar to the charge/discharge profile, indicating that the electrode in the continuous charge/discharge process is close to the equilibrium due

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to the good electrical conductivity of MnO@C [42]. The MnO@C electrode shows excellent cycling stability and at 100th cycle the reversible capacity can be still retained at 262 mAh g
1. This high specific capacity should be ascribed to the following reasons: (1) The carbon matrix can effectively cushion the volume expansion and reduce the aggregation of nanoparticles during the sodiation/ desodiation processes, resulting in excellent cycling performance. (2) The carbon matrix serving as conductive paths can increase the electrical contact between MnO nanoparticles for fast electron transport and facilitate the easy access of electrolyte into the electrode surface, which enhances the electrochemical kinetics of the MnO electrode [43]. The SEM image of MnO@C after 100 galvanostatic charge/discharge cycles is shown in Fig. S6. It is clearly observed that the morphology of MnO@C keeps integrity after 100 cycles, indicating its structure stability during cycling [44]. However, the rods are covered with a thin SEI layer, resulting in their fuzzier edge. Fig. 5(b) shows the rate performance of MnO@C electrode at various current densities. The average capacities can reach 255, 222.5, 186.3, 155.3, 138.7 and 115.7 mAh g
1 at the current densities of 50, 100, 500, 1000, 2000 and 4000 mA g
1, corresponding to the time periods for a complete discharge in 316, 132, 43, 19, 8 and 2 min, respectively. A capacity of 265 mAh g
1 can be obtained again when the current density is recovered to 50 mA g
1. The MnO@C electrode exhibits excellent rate capability because carbon matrix can largely increase the electrical conductivity of MnO@C, which facilitates the diffusion of sodium ions and electrons. Fig. 6(a) presents the potential profiles at different current densities. The profiles retain similar shapes without polarization, indicating that the transport of sodium ions and electrons into and

Fig. 6. (a) Charge/discharge voltage plots of MnO@C electrode at various current densities, (b) long-life cycling performance of MnO@C electrode at 2 A g
1 for 5000 cycles.

out of the frameworks is sufficiently rapid to satisfy the fast charge/ discharge process. Furthermore, the MnO@C electrode was tested (cycled after 100 cycles at 50 mA g
1) at 2 A g
1 for 5000 cycles to study the long-life cycling performance at high current density (Fig. 6(b)). The specific charge capacity of MnO@C can still be maintained at 140 mAh g
1 after 5000 cycles, and the CE is remained at ~100%. The CE of slightly over 100% is mainly caused by a minor parasitic reaction (electrolyte oxidation), as reported in the literature [45,46]. The excellent cycling performance at ultra-fast charge/discharge process indicates the superior structure stability of MnO@C during the sodiation/desodiation.

4. Conclusions

Fig. 5. (a) Cycling stability along with CE of MnO@C and MnO electrodes, (b) rate performance of MnO@C electrode.

In summary, a novel porous MnO@C hybrid structure was successfully synthesized through the direct calcination of Mn-BTC MOFs precursor and investigated as anode materials for SIBs. The carbon matrix can act as a dispersive buffer to prevent the agglomeration of MnO nanoparticles and cushion the volume expansion during the charge/discharge process. In addition, the carbon matrix can increase the electrical conductivity of MnO, which results in an enhanced electrochemical performance of the MnO@C electrode. The MnO@C exhibits a maximum capacity of 260 mAh g
1 at a current density of 50 mA g
1 after 100 cycles. Even at a high current density of 2 A g
1, a capacity of 140 mAh g
1 can be remained after 5000 cycles. The results suggest that MnO@C

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should be suitable for the next generation of anode materials for SIBs with superior stability and long-life cycling performance. Acknowledgments Financial support was provided by Basic Research Project of Shanghai Science and Technology Committee (No. 14JC1491000), Research Team of Anhui Provincial Education Department (No. 2016SCXPTTD) and Key Project of National Natural Science Foundation of China (No. 61231003). Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jallcom.2017.03.314. References [1] V. Palomares, P. Serras, I. Villaluenga, K.B. Hueso, J. Carretero-Gonzalez, T. Rojo, Na-ion batteries, recent advances and present challenges to become low cost energy storage systems, Energy Environ. Sci. 5 (2012) 5884e5901. [2] Y. Song, Y. Chen, J. Wu, Y. Fu, R. Zhou, S. Chen, L. Wang, Hollow metal organic frameworks-derived porous ZnO/C nanocages as anode materials for lithiumion batteries, J. Alloys Compd. 694 (2017) 1246e1253. [3] M. Xie, Y. Huang, M. Xu, R. Chen, X. Zhang, L. Li, F. Wu, Sodium titanium hexacyanoferrate as an environmentally friendly and low-cost cathode material for sodium-ion batteries, J. Power Sources 302 (2016) 7e12. [4] R. Chen, Y. Huang, M. Xie, Q. Zhang, X. Zhang, L. Li, F. Wu, Preparation of prussian blue submicron particles with a pore structure by two-step optimization for Na-Ion battery cathodes, ACS Appl. Mater. Interfaces 8 (2016) 16078e16086. [5] R. Chen, Y. Huang, M. Xie, Z. Wang, Y. Ye, L. Li, F. Wu, Chemical inhibition method to synthesize highly crystalline prussian blue analogs for sodium-ion battery cathodes, ACS Appl. Mater. Interfaces 8 (2016) 31669e31676. [6] T. Chen, L. Pan, T. Lu, C. Fu, D.H.C. Chua, Z. Sun, Fast synthesis of carbon microspheres via a microwave-assisted reaction for sodium ion batteries, J. Mater. Chem. A 2 (2014) 1263e1267. [7] Y. Lu, N. Zhang, Q. Zhao, J. Liang, J. Chen, Micro-nanostructured CuO/C spheres as high-performance anode materials for Na-ion batteries, Nanoscale 7 (2015) 2770e2776. [8] H. Liu, M. Jia, N. Sun, B. Cao, R. Chen, Q. Zhu, F. Wu, N. Qiao, B. Xu, Nitrogenrich mesoporous carbon as anode material for high-performance sodium-ion batteries, ACS Appl. Mater. Interfaces 7 (2015) 27124e27130. [9] X. Fan, J. Mao, Y. Zhu, C. Luo, L. Suo, T. Gao, F. Han, S.-C. Liou, C. Wang, Superior stable self-healing SnP3 anode for sodium-ion batteries, Adv. Energy Mater. (2015) 1500174. [10] H. Hou, M. Jing, Y. Yang, Y. Zhang, W. Song, X. Yang, J. Chen, Q. Chen, X. Ji, Antimony nanoparticles anchored on interconnected carbon nanofibers networks as advanced anode material for sodium-ion batteries, J. Power Sources 284 (2015) 227e235. [11] S.H. Woo, L. Yadgarov, R. Rosentsveig, Y. Park, D. Song, R. Tenne, S.Y. Hong, Fullerene-like re-doped MoS2 nanoparticles as an intercalation host with fast kinetics for sodium ion batteries, Isr. J. Chem. 55 (2015) 599e603. [12] J. Wang, C. Luo, J. Mao, Y. Zhu, X. Fan, T. Gao, A.C. Mignerey, C. Wang, Solidstate fabrication of SnS2/C nanospheres for high-performance sodium ion battery anode, ACS Appl. Mater. Interfaces 7 (2015) 11476e11481. [13] M. Walter, T. Zünd, M.V. Kovalenko, Pyrite (FeS2) nanocrystals as inexpensive high-performance lithium-ion cathode and sodium-ion anode materials, Nanoscale 7 (2015) 9158e9163. [14] Y.-L. Ding, P. Kopold, K. Hahn, P.A. van Aken, J. Maier, Y. Yu, A lamellar hybrid assembled from metal disulfide nanowall arrays anchored on a carbon layer: in situ hybridization and improved sodium storage, Adv. Mater. 28 (2016) 7774e7782. [15] X. Zhang, W. Qin, D. Li, D. Yan, B. Hu, Z. Sun, L. Pan, Metal-organic framework derived porous CuO/Cu2O composite hollow octahedrons as high performance anode materials for sodium ion batteries, Chem. Commun. 51 (2015) 16413e16416. [16] Z. Chen, D. Zhang, X. Wang, X. Jia, F. Wei, H. Li, Y. Lu, High-performance energy-storage architectures from carbon nanotubes and nanocrystal building blocks, Adv. Mater. 24 (2012) 2030e2036. [17] M.M. Rahman, I. Sultana, Z. Chen, M. Srikanth, L.H. Li, X.J. Dai, Y. Chen, Ex situ electrochemical sodiation/desodiation observation of Co3O4anchored carbon nanotubes: a high performance sodium-ion battery anode produced by pulsed plasma in a liquid, Nanoscale 7 (2015) 13088e13095. [18] Y.-T. Weng, T.-Y. Huang, C.-H. Lim, P.-S. Shao, S. Hy, C.-Y. Kuo, J.-H. Cheng, B.J. Hwang, J.-F. Lee, N.-L. Wu, An unexpected large capacity of ultrafine manganese oxide as a sodium-ion battery anode, Nanoscale 7 (2015) 20075e20081. [19] S. Wang, Y. Xing, C. Xiao, H. Xu, S. Zhang, A peapod-inspired MnO@C core-

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