CuO Nanorod Arrays Formed Directly on Cu Foil from MOFs as Superior Binder-Free Anode Material for Lithium-Ion Batteries Dongming Yin,†,‡ Gang Huang,*,§ Zhaolin Na,†,‡ Xuxu Wang,†,‡ Qian Li,†,‡ and Limin Wang*,† †
State Key Laboratory of Rare Earth Resource Utilization, Changchun Institute of Applied Chemistry, CAS, Changchun 130022, China ‡ University of Chinese Academy of Sciences, Beijing 100049, China § WPI Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan S Supporting Information *
ABSTRACT: In this study, an in situ growth method is developed for the partial conversion of current collector into active materials for lithium ion batteries (LIBs). Through thermal treatment of a metal− organic framework (MOF) precursor, of which the metal ion is provided by a Cu foil current collector, porous CuO nanorod arrays (NRAs) can be directly formed on Cu foil. Importantly, this strategy can avoid the poor contact problem between the current collector and electrode material as well as circumvent the addition of insulating material (binder) and inhomogeneous distribution of conductive carbon material and active material on the current collector. When evaluated as binder-free electrodes for LIBs, porous CuO NRAs deliver a high specific capacity (1341 mA h g−1 at 100 mA g−1) and enhanced rate capability and cycling ability (671 mA h g−1 at 100 mA g−1 after 150 cycles).
N
pulverization during the discharge−charge processes still exist in the binder-free electrodes. In addition, techniques for the preparation of binder-free electrodes, such as electrodeposition and vapor deposition, suffer from a lot of shortcomings, including an intricate synthesis process, demand of special equipment, and a fancy price for large-scale production.19,20 Therefore, exploration of a binder-free electrode with suitable construction, facile preparation methods, and inexpensive cost needs urgent attention and implementation. Recently, one-dimensional (1D) nanomaterials (needles,21 wires,22−24 rods,25 tubes26) have attracted wide attention and shown some incomparable advantages in LIB materials. First, one-dimensional nanomaterials can shorten the Li+ and electron diffusion distance in the radial direction and provide more effective Li+ and electron transport channels in the axial direction, which are beneficial to obtain higher rate performance. Second, 1D nanomaterials with a stable architecture and large specific surface area effectively accommodate the stress exerted during lithiation−delithiation processes, resulting in better cycling stability.27−29 However, the rational design and
owadays, energy and environmental crises have aroused violent competition worldwide in developing efficient and effective power systems. In this context, lithium-ion batteries (LIBs) as clean and convenient highenergy rechargeable batteries are becoming more and more remarkable for the ever-enlarging markets of communication facilities, portable electronic products, and electric vehicles.1−3 Tremendous efforts have been motivated to develop novel electrode materials and construction of LIBs with higher power density, superior energy density, and longer cycle stability.4−6 Among the available anode materials, CuO attracts extensive attention because of its high theoretical specific capacity (about 670 mA h g−1), abundant natural resource, environmental friendliness, and moderate cost.7−11 Nonetheless, the huge volume expansion/contraction during discharge−charge cycles results in electrode pulverization and electric contact loss between the current collector and active material, leading to a large irreversible capacity loss and deteriorative cycling stability. Moreover, the polymer binder usually used for powder electrode materials will not only increase the fabricating cost of LIBs but also hinder further improvement of capacity and rate capability of the cells. To solve this issue, abundant efforts have been devoted to develop binder-free electrodes.12−18 Unfortunately, the morphological collapse and material © 2017 American Chemical Society
Received: March 10, 2017 Accepted: June 7, 2017 Published: June 7, 2017 1564
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Figure 1. Schematic illustration of the process used to synthesize CuO NRAs on copper foil.
Figure 2. XRD patterns of the as-prepared (a) Cu-TCNQ NRAs with different reaction times and (b) CuO NRAs.
precursors, formed directly on Cu foil by a facile dipping method without an additional metal source, is successfully converted into porous CuO NRAs through morphologymaintained thermal treatment. As an anode material for LIBs, this designed binder-free electrode avoids the use of an insulating binder and inhomogeneous distribution of conductive carbon material on the current collector; meanwhile, it can achieve intimate adhesion between the material and the current collector. As a result, this binder-free electrode could not only improve the specific energy capacity when taking the whole electrode into consideration but also exhibit enhanced rate capability. The developed strategy is inexpensive and can be easily scaled up for commercial production of advance electrode materials. The manufacturing processes for highly ordered CuO NRAs on Cu foil are schematically described in Figure 1. First, Cu foil is treated with HCl to remove adsorbed dust and surface impurities. Subsequently, the preptreated Cu foil is dipped into 7,7,8,8-tetracyanoquinodimethane (TCNQ) acetonitrile solution for the formation of Cu-TCNQ. After that, CuO NRAs are obtained by thermal annealing of Cu-TCNQ at 250 °C under Ar and air atmosphere. The optical images of Cu-TCNQ NRAs and CuO NRAs are shown in Figure S1a (Supporting Information). It can be seen that the color of the copper foil
facile fabrication of 1D nanomaterials with robust construction and remarkable lithium storage properties still remain a big challenge. As a newly developing kind of porous materials, metal− organic frameworks (MOFs) with distinguishable advantages of large surface area, high porosity, tunable pore size, and functional flexibility have been generally applied in catalysis,30,31 gas adsorption,32 health monitoring,33 and electrochemistry.34−36 In recent years, there is gaining momentum to use MOFs as a precursor/template to design porous metal oxides with different morphologies for application in energy storage. For instance, hollow Cu2O/CuO cages,37 porous CuO hollow octahedra,38 and porous carbon-coated CuCo2O4 concave polyhedra39 have been successfully synthesized with appropriate MOF precursors/templates and demonstrated some benefits for boosting the electrochemical performances of LIBs. Though great progess has been made, the rational design of 1D porous nanostructures through MOF precursors/templates with preeminent construction and enhanced lithium storage performances still remains a huge challenge. Herein, we solve these challenges via the construction of copper-substrate-supported porous CuO nanorod arrays (NRAs), which are in situ derived from MOFs and used as binder-free electrodes for high-performance LIBs. The MOF 1565
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ACS Energy Letters turns from luminous yellow to light green after the formation of Cu-TCNQ NRAs, which is finally transformed into black when heat treated. The thermal behavior of Cu-TCNQ NRAs is researched by thermogravimetric (TG) and derivative thermogravimetric (DTG) analyses (Figure S2a, Supporting Information). It is found that the first weight loss corresponds to the loss of H2O. Further increasing the temperature from 200 to 300 °C, there is weight loss, probably attributed to the partial oxidation of the organic linker into CO2, NO2, and H2O. The sharp weight loss between 300 and 400 °C may be due to the completely oxidative decomposition of the organic linker. Because the thermal expansion coefficients of CuO NRAs and copper foil are different, when the temperature reaches 300 °C, the CuO NRAs on the surface of the Cu foil will peel off. For this reason, we choose 250 °C as the thermal treatment temperature. The thermal behavior of the obtained CuO NRAs is further investigated in Figure S2b (Supporting Information). As can be seen from the TG curve, there is still some weight loss after thermal treatment, corresponding to the oxidative decomposition of residual C and N. With a further increase of the temperature, the integral weight is slightly increased by the oxidation of Cu, which comes from the Cu foil when scraping CuO NRAs from Cu foil to perform TG analysis. The crystal structure and phase composition of Cu-TCNQ with different reaction times are characterized by X-ray diffraction (XRD). As can be seen from Figure 2a, the multiple peaks of Cu-TCNQ NRAs at 6 h are essentially identical to the Cu-TCNQ phase I reported in early literature.30,40 With the extension of reaction time, some products may be oxidized, resulting in the deviation of the XRD spectrum. The XRD pattern of CuO NRAs in Figure 2b shows a typical crystalline structure of CuO (JCPDS No. 48-1548). Because CuO NRAs is in situ growth on the surface of Cu foil, it is necessary to scrape off the CuO NRAs to conduct the XRD test. However, the copper is more or less introduced during the process of scratching, which will induce the inevitable occurrence of the crystalline Cu structure (JCPDS No. 04-0836) in the XRD pattern. There are no peaks of impurities that can be detected, confirming the successful transformation of Cu-TCNQ to CuO by the chosen temperature and designed method. The morphologies of Cu-TCNQ and CuO NRAs are investigated by a scanning electron microscope (SEM) and transmission electron microscopy (TEM) in Figures 3 and 4. It can be seen that the number of nanorods increases gradually with the extension of reaction time (Figure 3a,b). After 24 h, large-scale, aligned and dense Cu-TCNQ NRAs uniformly directly form on Cu foil (Figure 3c). The magnified crosssectional image (inset of Figure 3c) shows that the Cu-TCNQ NRAs have a straight axis and the average length is about 2−3 μm with a tapering outline (Figure 4a). When the concentration of TCNQ in acetonitrile solution is increased to 0.4 mg mL−1, many black lines are formed on the surface of the copper foil (Figure S1d, Supporting Information). This is further checked by SEM characterization in Figure S3 (Supporting Information); the increase of the TCNQ concentration leads to the growth of bowknot-like particles, highlighting the significant role of the concentration of TCNQ. As revealed by Figure 3d, the CuO NRAs well inherit the rodlike structure of MOF precursors and are composed of nanosized building blocks that are connected with each other forming a porous architecture, which can be further confirmed by TEM examination shown in Figure 4b. It is noticed that the
Figure 3. SEM images of Cu-TCNQ in a 0.2 mg mL−1 TCNQ acetonitrile solution at different times: (a) 2 min, (b) 1 h, and (c) 24 h. (d) Cu-TCNQ thermal treatment at 250 °C.
Figure 4. (a) SEM image of Cu-TCNQ in a 0.2 mg mL−1 TCNQ acetonitrile solution for 24 h and structural characterizations of CuO NRAs: (b) TEM image, (c) HRTEM image, (d) the corresponding SAED pattern, (e) HAADF-STEM image, and (f− i) element mapping.
nanorods are about 2 μm in length, consistent with the SEM results. As shown in the high-resolution TEM (HRTEM) image in Figure 4c, the distances between adjacent lattice fringes are measured to be 0.231 and 0.158 nm, corresponding to the (200) and (202) planes of CuO, respectively. The SAED image of CuO NRAs (Figure 4d) demonstrates a set of well-defined spots, revealing its polycrystalline nature. As shown in Figure 4e, elemental mapping has been conducted to examine the chemical components and their corresponding spatial distribution in CuO NRAs. The Cu, O, C, and N are homogeneously distributed across the selected nanorod (Figure 4f−i). The 1566
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Figure 5. Tapping mode (a) (inset: height analysis), (b,c) AFM analysis of CuO NRAs in situ grown on Cu foil, and (d,e) XPS spectra for Cu 2p and O 1s in CuO NRAs.
Figure 6. (a) CV curves of CuO NRAs at a scan rate of 0.1 mV s−1 in the voltage range of 0.01−3.0 V vs Li/Li+. (b) Galvanostatic discharge− charge voltage profiles of CuO NRAs for the 1st, 2nd, 10th, 50th, and 100th cycles at a current rate of 100 mA g−1 between 0.01 and 3.0 V. (c) Discharge−charge capacities of CuO NRAs and S-CuO and corresponding Coulombic efficiency of CuO NRAs at a current rate of 100 mA g−1. (d) Rate capabilities of CuO NRAs and S-CuO.
functionality for NRAs construction with CuO NRAs as sacrificial templates. In order to further identify the surface texture of CuO NRAs, atomic force microscope (AFM) analysis is employed. Figure 5a−c exhibits the tapping-mode AFM images of CuO NRAs. The average topographic height of CuO NRAs is approximately 2.12 μm, which is in line with the results from SEM and TEM images. The chemical composition and valency of Cu-TCNQ
above-mentioned results indicate the successful in situ transformation of MOFs into CuO NRAs. This facile designed method for the preparation of aligned NRAs on Cu foil needs no mold directing agent or noble metal intermediates, making it a wise choice that has great promise for large-scale production. What’s more, the acid-removable instinct of CuO NRAs means that this method can be readily expanded to prepare other 1567
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manifesting remarkable capacity retention of the CuO NRA electrode. The morphology of the CuO NRA electrode after 100 discharge−charge cycles at 100 mA g−1 is displayed in Figure S6a (Supporting Information). It confirms that the CuO NRA’s architecture is still basically maintained after cycling, revealing the robustness of the designed porous structure, which may be responsible for the superior cycling stability. The electrochemical impedance spectroscopy (EIS) of CuO NRAs is analyzed in Figure S6b,c (Supporting Information). The Nyquist plot of CuO NRAs illustrates a depressed semicircle in the high-frequency region followed by a long line in the lowfrequency area, which are relative to the charge transfer resistance and the diffusion resistance of Li+ diffusion in the electrode. The bulk resistance (Rs) of the electrolyte, electrodes, and separator is about 2.06 Ω, manifesting preferable conduction of the electrode because of the strong adhesion between CuO NRAs and Cu foil. The charge transfer impedance (Rct) is about 18.12 Ω, which is smaller than that in a previous report.41 As a comparison, the lithium storage performance of S-CuO with PVDF as a binder and carbon black as a conductive agent has also been researched. An initial discharge capacity of 1033 mA h g−1 is achieved on the basis of the mass of active material, and a capacity of only 468 mA h g−1 is maintained after 150 discharge−charge cycles, which is obviously lower than that of in situ growth of CuO NRAs. The cyclic performance of CuO NRAs and S-CuO electrodes in the first 400 cycles at a high current density of 500 mA g−1 is shown in Figure S7 (Supporting Information), confirming the excellent stability of CuO NRAs. The CuO NRAs and SCuO electrodes deliver specific capacities of 521 and 394 mA h g−1, respectively, after 400 cycles. The prominent electrochemical performance of CuO NRAs in LIBs is further ascertained at different rates. It can be seen in Figure 6d that the CuO NRAs exhibit decent capacity retention despite sustaining rapid change of current density. The average discharge capacities are 730, 554, 434, and 367 mA h g−1 at current densities of 100, 200, 500, and 1000 mA g−1, respectively. Even at a cruelly high current density of 2000 mA g−1, the binder-free CuO NRA electrode could still receive a discharge capacity of 300 mA h g−1. After experiencing the rapid change of current densities, a discharge capacity of 685 mA h g−1 can still be recovered by reducing the density back to 100 mA g−1 and this value can be kept without marked attenuation after 100 cycles. However, for S-CuO, the average discharge capacities are 535, 459, 389, 308, and 218 mA h g−1 at rates of 100, 200, 500, 1000, and 2000 mA g−1, respectively, which are much lower than the corresponding capacities of the CuO NRAs. The better cycling and rate capability of CuO NRAs can be attributed to facile Li+ insertion−extraction and efficient electron transport in the designed porous binder-free electrode. The electrochemical performances of the as-prepared CuO NRAs for LIBs have been compared with those of some previously reported CuO anode materials and are listed in Table S1 (Supporting Information).11,46 The results show that the binder-free CuO NRAs in situ derived from MOFs exhibit superior or comparable electrochemical performances. On the basis of the above results, the inspiring lithium storage performances can be attributed to the following reasons. First, part of the copper foil current collector is converted into porous CuO NRAs, which can avoid the poor contact between the current collector and electrode material. Second, the binder-free electrode avoids the addition of an insulating binder
NRAs are checked by X-ray photoelectron spectroscopy (XPS) shown in Figure S4 (Supporting Information). Similar to the early literature reports,30,40 the spectra of Cu-TCNQ NRAs exhibited typical binding energies for Cu(I) 2p1/2 and 2p3/2 with no evidence of shakeup satellites attributed to Cu(II). As shown in Figure S5 (Supporting Information), the XPS data illustrate the existence of Cu, O, C, and N elements in the CuO NRAs. The Cu 2p spectrum in Figure 5d consists of two spin− orbit doublets and two shakeup satellites. The fitting energy bands at 933.6 and 953.7 eV are attributed to Cu 2p3/2 and Cu 2p1/2 characteristic peaks of CuO NRAs. Meanwhile, the satellite peaks at 940.4, 943.5, and 962.1 eV are three shakeup peaks of Cu at high binding energy sides of Cu 2p3/2 and Cu 2p1/2 edges, respectively. To give insight into the potential application, the electrochemical performances of CuO NRAs as binder-free anode materials for LIBs are evaluated. Figure 6a depicts the representative CV curves of CuO NRAs in the voltage window of 0.01−3.0 V at a scan rate of 0.1 mV s−1. In the first cathodic scan, a weak broad peak at 1.75 V, an intense peak at 1.19 V, and two moderate peaks at 0.85 and 0.62 V can be obviously observed, corresponding to a multistep electrochemical reaction. The detailed electrochemical description is as follows: (i) creation of Cu 1−x II Cu x I O 1−x/2 , (ii) conversion of Cu1−xIICuxIO1−x/2 into the Cu2O phase, and (iii) further lithiation of Cu2O into Cu and the formation of amorphous Li2O, which follows the growth of solid electrolyte interface (SEI) films.42−45 The anodic peaks at 2.42 and 2.67 V are associated with the formation of Cu2O and Li accompanied by the partial oxidation of Cu2O into CuO.41,42 In the subsequent cycles, the reductive and oxidative peaks mainly stay at the same potentials, suggesting a high reversibility of lithium storage. On the basis of the above analysis, the electrochemical conversion reactions of CuO NRAs can be expressed as follows: 2CuO + 2x Li+ + 2x e− → 2Cu1 − x IICux IO1 − x /2 + x Li 2O (1)
2Cu1 − x IICux IO1 − x /2 + 2(1 − x)Li+ + 2(1 − x)e− → Cu 2O + (1 − x)Li 2O Cu 2O + 2Li+ + 2e− → 2Cu + Li 2O
(2) (3)
Figure 6b depicts the representative discharge−charge curves of CuO NRAs at a current density of 100 mA g−1 between 0.01 and 3.0 V. It is noted that the platforms on the voltage profiles are consistent with the CV curves. The first discharge capacity of CuO NRAs is 1341 mA h g−1, which is much larger than the theoretical value of CuO (about 670 mA h g−1). This additional capacity can be attributed to the irreversible formation of SEI films, a common phenomenon for most transition metal oxides. The cycling stabilities of CuO NRAs and S-CuO (CuO NRAs scraped off of the Cu foil) are investigated at a current density of 100 mA g−1 between 0.01 and 3.0 V (Figure 6c). The discharge capacity of CuO NRAs reduces mildly during the whole cycling process, and it can still achieve a reversible capacity of 671 mA h g−1 even after 150 cycles. Importantly, although the initial Coulumbic efficiency of CuO NRAs is only about 56%, which is derived from the construction of SEI films, part of the undecomposable Li2O, and irreversible decomposition of electrolyte, it demonstrates a rapid augment and can remain stable at around 98%, except for the initial two cycles, 1568
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of electrode materials reacting through conversion reactions. Adv. Mater. 2010, 22, E170−E192. (6) Sun, Y. M.; Hu, X. L.; Luo, W.; Xia, F. F.; Huang, Y. H. Reconstruction of conformal nanoscale MnO on graphene as a highcapacity and long-life anode material for lithium ion batteries. Adv. Funct. Mater. 2013, 23, 2436−2444. (7) Ko, S.; Lee, J. I.; Yang, H. S.; Park, S.; Jeong, U. Mesoporous CuO particles threaded with CNTs for high-performance lithium-ion battery anodes. Adv. Mater. 2012, 24, 4451−4456. (8) Wang, J.; Liu, Y.; Wang, S.; Guo, X.; Liu, Y. Facile fabrication of pompon-like hierarchical CuO hollow microspheres for high-performance lithium-ion batteries. J. Mater. Chem. A 2014, 2, 1224−1229. (9) Yang, W.; Wang, J.; Ma, W.; Dong, C.; Cheng, G.; Zhang, Z. Free-standing CuO nanoflake arrays coated Cu foam for advanced lithium ion battery anodes. J. Power Sources 2016, 333, 88−98. (10) Yuan, S.; Huang, X. L.; Ma, D. L.; Wang, H. G.; Meng, F. Z.; Zhang, X. B. Engraving copper foil to give large-scale binder-free porous CuO arrays for a high-performance sodium-ion battery anode. Adv. Mater. 2014, 26, 2273−2279. (11) Liu, Y.; Qiao, Y.; Zhang, W. X.; Hu, P.; Chen, C. J.; Li, Z.; Yuan, L. X.; Hu, X. L.; Huang, Y. H. Facile fabrication of CuO nanosheets on Cu substrate as anode materials for electrochemical energy storage. J. Alloys Compd. 2014, 586, 208−215. (12) Wang, J.; Zhang, Q.; Li, X.; Xu, D.; Wang, Z.; Guo, H.; Zhang, K. Three-dimensional hierarchical Co3O4/CuO nanowire heterostructure arrays on nickel foam for high-performance lithium ion batteries. Nano Energy 2014, 6, 19−26. (13) Botas, C.; Carriazo, D.; Singh, G.; Rojo, T. Sn- and SnO2graphene flexible foams suitable as binder-free anodes for lithium ion batteries. J. Mater. Chem. A 2015, 3, 13402−13410. (14) Li, N.; Jin, S. X.; Liao, Q. Y.; Wang, C. X. ZnO anchored on vertically aligned graphene: binder-free anode materials for lithium-ion batteries. ACS Appl. Mater. Interfaces 2014, 6, 20590−20596. (15) Liu, Y.; Wang, W.; Ying, Y.; Wang, Y.; Peng, X. Binder-free layered Ti3C2/CNTs nanocomposite anodes with enhanced capacity and long-cycle life for lithium-ion batteries. Dalton T. 2015, 44, 7123− 7126. (16) Song, J.; Zhou, M.; Yi, R.; Xu, T.; Gordin, M. L.; Tang, D.; Yu, Z.; Regula, M.; Wang, D. Interpenetrated gel polymer binder for highperformance silicon anodes in lithium-ion batteries. Adv. Funct. Mater. 2014, 24, 5904−5910. (17) Xie, W. H.; Li, S. Y.; Wang, S. Y.; Xue, S.; Liu, Z. J.; Jiang, X. Y.; He, D. H. N-doped amorphous carbon coated Fe3O4/SnO2 coaxial nanofibers as a binder-free self-supported electrode for lithium ion batteries. ACS Appl. Mater. Interfaces 2014, 6, 20334−20339. (18) Zuo, Z. C.; Kim, T. Y.; Kholmanov, I.; Li, H.; Chou, H.; Li, Y. Ultra-light hierarchical graphene electrode for binder-free supercapacitors and lithium-ion battery anodes. Small 2015, 11, 4922−4930. (19) Wang, R.; Xu, C.; Sun, J.; Liu, Y.; Gao, L.; Lin, C. Free-standing and binder-free lithium-ion electrodes based on robust layered assembly of graphene and Co3O4 nanosheets. Nanoscale 2013, 5, 6960−6967. (20) Zhang, L.; Gong, H. Partial conversion of current collectors into nickel copper oxide electrode materials for high-performance energy storage devices. ACS Appl. Mater. Interfaces 2015, 7, 15277−15284. (21) Chen, X.; Zhang, N.; Sun, K. Facile fabrication of CuO 1D pineneedle-like arrays for super-rate lithium storage. J. Mater. Chem. 2012, 22, 15080−15084. (22) Huang, J.; Zhu, Y.; Yang, X.; Chen, W.; Zhou, Y.; Li, C. Flexible 3D porous CuO nanowire arrays for enzymeless glucose sensing: in situ engineered versus ex situ piled. Nanoscale 2015, 7, 559−569. (23) Sahay, R.; Suresh Kumar, P.; Aravindan, V.; Sundaramurthy, J.; Chui Ling, W.; Mhaisalkar, S. G.; Ramakrishna, S.; Madhavi, S. High aspect ratio electrospun CuO nanofibers as anode material for lithiumion batteries with superior cycleability. J. Phys. Chem. C 2012, 116, 18087−18092. (24) Cao, F.; Xia, X. H.; Pan, G. X.; Chen, J.; Zhang, Y. J. Construction of carbon nanoflakes shell on CuO nanowires core as
and inhomogeneous distribution of active material on the current collector. Third, the unique 1D porous structure and the large surface area can provide ample space and relaxed texture, thus greatly increasing the active material and electrolyte contact area, shortening the lithium ion and electron diffusion distance, providing fast electron transport paths, and releasing the stress during insertion/extraction of lithium. Moreover, the residual C and N could also promote the electrochemical performance. In summary, we have demonstrated a facile and effective strategy for the establishment of copper-substrate-supported porous CuO NRAs in situ derived from MOFs. When used as an anode material for LIBs, the CuO NRAs exhibits high specific capacity, superior cycling stability, and outstanding rate capability. This novel design method provides an efficient technique for promoting the electrochemical performances of CuO and can be extended to produce other novel electrode materials. This effective and versatile strategy can be further extended to fabricate other functional metal oxides with tunable compositions and constructions by choosing suitable MOFs as the precursor/template.
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ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsenergylett.7b00215. Experimental section, materials characterization, electrochemical measurements, optical images of the materials, TG curves, additional SEM images, XPS spectra, electrochemical impedance spectroscopy, cyclic performance, and comparison table (PDF)
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AUTHOR INFORMATION
Corresponding Authors
*E-mail:
[email protected]. Tel: +86-431-85262447. Fax: +86-431-85262836 ( L.W.). *E-mail:
[email protected]. Tel: +81-22-217-5991. Fax: +81-22-217-5959 (G.H.). ORCID
Limin Wang: 0000-0001-9618-9239 Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This research is financially supported by the National Nature Science Foundation of China (Grant No. 21521092). REFERENCES
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DOI: 10.1021/acsenergylett.7b00215 ACS Energy Lett. 2017, 2, 1564−1570
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DOI: 10.1021/acsenergylett.7b00215 ACS Energy Lett. 2017, 2, 1564−1570
