Yolk@Shell or Concave Cubic NiO–Co3O4@C Nanocomposites

Jul 31, 2017 - (29-32) Because MOFs have a large fraction of metal ion and organic ligands, the use of MOFs as promising precursors or sacrificial tem...
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Yolk@Shell or Concave Cubic NiO−Co3O4@C Nanocomposites Derived from Metal−Organic Frameworks for Advanced Lithium-Ion Battery Anodes Gang Huang,†,‡ Dongming Yin,†,‡ Feifei Zhang,†,‡ Qian Li,†,‡ and Limin Wang*,†,§ †

State Key Laboratory of Rare Earth Resource Utilization, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, Changchun 130022, P. R. China ‡ University of Chinese Academy of Sciences, Beijing 100049, P. R. China § Changzhou Institute of Energy Storage Materials and Devices, Changzhou 213000, P. R. China S Supporting Information *

ABSTRACT: Novel hybrid metal oxides with advanced architectures are extensively pursued to achieve synergetic properties with respect to improved lithium-ion storage properties. Here, rationally designed yolk@shell or concave NiO−Co3O4@C (YNCC or CNCC) nanocubes have been fabricated by the simple and versatile thermolysisinduced transformation of metal−organic frameworks (MOFs), aimed at simultaneously addressing the capacity fade and conductivity deficiency of metal oxides. The as-prepared nanocomposites with plentiful hierarchical pores integrate the distinct functionalities of the ternary components: NiO and Co3O4 as the major active materials can guarantee high capacity, while carbon can improve the conductivity and accommodate volume changes. Benefitting from the intrinsic material and architecture features, the YNCC and CNCC nanocomposites deliver excellent electrochemical performances with high reversible specific capacity, superior cycling stability (803 and 870 mAh g−1 at 100 mA g−1 after 100 cycles), and good rate capability (339 and 398 mAh g−1 at 2 A g−1) as anode materials for lithium-ion batteries.

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and alloying/dealloying processes (Si and SnO2).17,18 Among them, Co3O4 and NiO are advocated as promising anode candidates and inspired an explosion of research.10−12,19,20 Unfortunately, at least two issues still remain that greatly hinder the real-word implementation of Co3O4 and NiO electrodes: one is the substantial volume change during the repetitive lithiation/delithiation processes, which results in powdering and crushing of the electrode and loss of electrical connection with the current collector, leading to poor cycling stability; the other is the intrinsically low conductivity of TMOs, giving rise to inferior rate capability. To surpass the limitations mentioned above, various hybrid electrode materials with advanced architectures have been proposed because a synergetic effect can be achieved by integrating the functions of each component and hence making full use of the composite materials, which is not possible for single-component materials. On the basis of this concept, a range of composites have been built by different oxides for LIBs, such as Co3O4−MnO2, Fe3O4−TiO2, Fe2O3− MnO2, and so on.21−23 Another conventional solution to these dilemmas is coating/supporting TMOs with a better conductive material, which is generally achieved in several ways, including

INTRODUCTION There has been an increasing and urgent demand to seek clean and renewable energy sources because of the rapid consumption of nonrenewable fossil fuel and progressively more severe environmental pollution. Lithium-ion batteries (LIBs), major devices for energy conversion and storage, have attracted tremendous attention in the past 2 decades and are extensively explored for applications in portable electronic devices, electrical vehicles, and smart grids storing intermittent energy sources.1−6 However, the low capacity and poor safety of graphitic materials cause the existing commercial battery systems to approach the achievable performance limit and fail to fulfill the ever-growing requirements for the latter applications. Therefore, searching for alternative anode materials that are capable of delivering higher energy/power density and longer cycling life and possessing more secure charge−discharge potentials versus lithium has become an immediate need.7 Attractively, transition-metal oxides (TMOs) stand out from various anode materials because of their wide availability, low cost, and high theoretical capacities.8−14 Moreover, TMOs provide more choices to realize large reversible capacities by adopting different lithiation mechanisms, such as conversion reactions (Co3O4 and NiO),10−12 intercalation/extraction mechanisms (TiO2 and Li4Ti5O12),15,16 © 2017 American Chemical Society

Received: May 22, 2017 Published: July 31, 2017 9794

DOI: 10.1021/acs.inorgchem.7b01296 Inorg. Chem. 2017, 56, 9794−9801

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Inorganic Chemistry improvement of the conductivity, alleviation of the volume change, prevention of the agglomeration and isolation of electrode materials, and modification of the material surface with a stable solid electrolyte interface (SEI) layer.24−28 Although these procedures are more or less effective, each designed strategy alone brings limited improvement of the lithium storage properties of TMOs. Metal−organic frameworks (MOFs) with distinguishable properties, like diverse skeleton structures, well-defined porosities, and large surface areas together with tunable functionality, have opened new perspective applications in catalysis, drug delivery, gas adsorption and separation, and energy storage.29−32 Because MOFs have a large fraction of metal ion and organic ligands, the use of MOFs as promising precursors or sacrificial templates to synthesize porous metal oxides, primarily via thermolysis, is burgeoning.33−38 In thermolysis, the metal ion can be utilized as the metal source, while the organic component decomposes into gas molecules, such as H2O, CO2, and others, easily leading to the generation of abundant pores, generally maintaining the geometric morphology of the template/precursor. For example, when suitable MOFs are chosen, Co3O4−NiCo2O4 double-shelled nanocages, ZnO@ZnO quantum dots/C core@shell nanorod arrays on carbon cloth, and porous spinel ZnxCo3−xO4 hollow polyhedra have been fabricated and exhibit some benefits in energy-related applications.34−36 Despite the progress achieved to date, it is significantly challengeable to build up an integrated hierarchical porous smart NiO−Co3O4@C nanocomposite by exploring a simple approach, where the structure characteristics and electroactivities of each component can be fully utilized. Herein, a facile and versatile MOF-derived synthetic protocol for the controllable construction of yolk@shell or concave NiO−Co3O4@C (hereafter denoted as YNCC or CNCC) nanocubes is proposed, with the help of a coprecipitation process plus a polydopamine coating, chemical etching, and a subsequently short thermolysis treatment. The compositional and structural superiorities as well as the synergetic effect from different components cause YNCC and CNCC nanocomposites to exhibit enhanced electrochemical performances as anode materials for LIBs.

Figure 1. Schematic illustration of the procedure used to fabricate YNCC and CNCC.

RESULTS AND DISCUSSION Figure 1 schematically illustrates the synthetic process of YNCC and CNCC. To begin with, Ni3[Co(CN)6]2·3H2O (NiCo) Prussian blue analogue nanocubes are formed by the coprecipitation reaction between Ni2+ and [Co(CN)6]3− with sodium citrate as the surfactant. Then, dopamine, containing functional amine and catechol groups, self-polymerizes and spontaneously deposits a conformal layer of polydopamine on the surface of NiCo by dispersing NiCo in the dopamine/Tris buffer solution with stirring, forming core@shell NiCo@ polydopamine (NiCo@PDA). Further, NiCo@PDA is treated with a hot HCl solution to disassociate the defective parts of NiCo (H-NiCo@PDA).39 Finally, the thermolysis process is conducted at a moderate temperature to convert NiCo@PDA and H-NiCo@PDA to YNCC and CNCC, respectively. The phase purities and crystalline structures of NiCo, NiCo@PDA, and H-NiCo@PDA are examined by X-ray diffraction (XRD) measurement (Figure 2). It can be clearly seen that the positions and relative intensities of all of the diffraction peaks for the obtained products well match the standard data of Ni3[Co(CN)6]2·12H2O (JCPDS 89-3738). No additional peaks from impurities are detected, implying that the

polydopamine coating and chemical etching do not affect the crystalline structure of NiCo. The absence of diffraction peaks from polydopamine indicates its amorphous nature. The morphology and structure of NiCo is elucidated by scanning electron microscopy (SEM; Figure S1, Supporting Information). It reveals the existence of large-scale and uniform NiCo nanocubes. The transmission electron microscopy (TEM) images of NiCo further confirm the cubic morphology with solid inner structure and smooth surfaces (Figure S2, Supporting Information). After coating with polydopamine, the morphology of NiCo is almost unchanged except for an increase in size (Figure 3a,b), suggesting the conformal coating of polydopamine. The TEM images in Figure 4a,b reveal that the polydopamine layer smoothly, uniformly, and continuously covered the NiCo surface, confirming the core@shell structure of NiCo@PDA. Similar to the morphology of NiCo@PDA, HNiCo@PDA is a well-maintained nanocube possessing rough surfaces at the corners and edges after hot HCl etching, as shown in Figure 3c,d. Further insight into the microstructure and morphology of H-NiCo@PDA is obtained by TEM (Figure 4c,d). There is still a uniform polydopamine layer depositing on NiCo, indicating that the HCl etching cut no ice

Figure 2. XRD patterns of the as-prepared NiCo, NiCo@PDA, and HNiCo@PDA.

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Figure 3. SEM images of (a and b) NiCo@PDA and (c and d) HNiCo@PDA. Figure 5. XRD patterns of YNCC and CNCC.

in concave nanocubes. Inductively coupled plasma atomic emission spectroscopy (ICP-AES) is further used to determine the compositions of YNCC and CNCC, which show atomic ratios Ni/Co of 1.51 and 1.32, respectively. Figure 6 shows the SEM images of YNCC and CNCC. As presented in Figure 6a,b, YNCC basically copies the cubic

Figure 4. TEM images of (a and b) NiCo@PDA and (c and d) HNiCo@PDA.

with polydopamine. The size distribution histogram in Figure S3 (Supporting Information) indicates that the as-synthesized NiCo, NiCo@PDA, and H-NiCo@PDA have average sizes of 202, 268, and 232 nm, respectively (Figure S3 and Table S1, Supporting Information). The size decrease of H-NiCo@PDA compared with that of NiCo@PDA is probably due to acid etching of the defective parts of NiCo. Thanks to the abundant functional amine and catechol groups of dopamine, the polydopamine could still firmly coat the core part as before without peeling off. A thermolysis process has been carried out to obtain YNCC and CNCC, and the production rates are 42% and 39% (Figure S4, Supporting Information), respectively. The phases and specific structures of the calcined products are investigated by XRD, and the corresponding results are illustrated in Figure 5. The identified diffraction peaks can be well indexed to the coexistence of NiO and Co3O4 with cubic structures, implying the successful formation of NiO and Co3O4 composites. Because of the relatively low thermolysis temperature, the carbon is amorphous and the diffraction peaks from carbon are absent. A considerable broadness in the XRD pattern of CNCC may arise from the nanoparticles, which act as building blocks

Figure 6. SEM images of (a and b) YNCC and (c and d) CNCC.

morphology of NiCo@PDA without noticeable thermally induced alterations but a decrease in size of ∼196 nm (Figure S3 and Table S1, Supporting Information). From the holes on the surfaces, we can to some extent see the hollow structure of YNCC. The magnified SEM image in Figure 6b reveals the thin and transparent nature of the carbon shell with a solid core located within, forming a yolk@shell structure. For CNCC, an interesting feature is that the six plane surfaces clearly turn concave, probably because of the etching of NiCo and the structural contraction caused by thermal decomposition of the organic linker (Figure 6c,d). The concave nanocubes with a relatively rough surface and a size of ∼160 nm are composed of nanometer-sized building blocks, as indicated by closer inspection in Figure 6d. The specific morphologies and microstructures of YNCC and CNCC are further detailedly elucidated by carrying out TEM and energy-dispersive spectroscopy (EDS) characterizations. The TEM images in Figure 7a,b clearly give the yolk@ 9796

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Figure 8. HAADF-STEM images and elemental mapping of (a) YNCC and (b) CNCC.

the weight fractions of carbon in these two composites are 10.6% and 11.5%, respectively. With the aid of the above analysis, we could be safely informed of the feasibility of chemical conversion of MOFs@PDA into NiO, Co3O4, and carbon nanocomposites by a low-temperature thermal treatment. The textures and porous structures of YNCC and CNCC are characterized by N2 adsorption−desorption measurements. As plotted in Figure S8a,c (Supporting Information), a remarkable hysteresis occurs in both samples, underscoring their mesoporous characteristics. Evidently, the shift of the hysteresis loop toward relatively higher pressure for YNCC, compared to that of CNCC, reveals a larger pore size of YNCC.40 The Brunauer−Emmett−Teller surface areas and total pore volumes of YNCC and CNCC are 71.2 m2 g−1 and 0.41 cm3 g−1 and 142.6 m2 g−1 and 0.53 cm3 g−1, respectively. As calculated by the Barrett−Joyner−Halenda method, the average pore size of YNCC is 19.7 nm, larger than that of 16.5 nm for CNCC (Figure S8c,d, Supporting Information). The existence of macropores in YNCC originates from the holes on the nanocube surface and the void space between the yolk and shell, consistent with the results unveiled from the SEM and TEM images. The high surface areas render both samples with large electrode−electrolyte contact areas, and the porous characteristic facilitates the penetration of electrolyte molecules and the transport of Li+ and electrons. In view of the unique compositional and structural features, the hierarchical porous YNCC and CNCC nanocomposites might serve as high-performance anode materials for LIBs. Hence, their electrochemical performances are evaluated by cyclic voltammetry (CV) and galvanostatic charge−discharge measurements. It should be noted that all specific capacities are calculated on the basis of the total mass of the nanocomposites. Moreover, in the same voltage window, the specific capacity of the low-content amorphous carbon is negligible. Parts a and c of Figure 9 give the first three CV curves of YNCC and CNCC, which are almost the superposition of the CV curves of NiO and Co3O4 (Figure S9, Supporting Information). As can be seen in Figure 9a, there are three cathodic peaks in the first discharge process. The low-intensity peak near 1.45 V refers to the partially reversible formation of a polymeric gel-like layer,9 and the broad peak around 0.75 V and the intense peak at ∼0.50 V correspond to the reduction of Co3O4 and NiO to metallic cobalt and nickel, respectively, accompanied by irreversible decomposition of the electrolyte to form a SEI layer, which leads to an irreversible capacity loss.11,12 These cathodic peaks are substituted by one broad peak at ∼1.00 V in subsequent cycles, owing to the lithiation-

Figure 7. TEM images of (a and b) YNCC and (c and d) CNCC.

shell structure of YNCC with yolk sizes of 20−90 nm. Furthermore, the brighter contrast of the shell demonstrates the ultrathin state of polydopamine-derived carbon. The void space between the yolk and shell could effectively accommodate the volume changes of NiO−Co3O4 during repetitive lithium insertion/extraction processes, and the carbon can act as a buffer layer to preserve the framework. The high-resolution TEM (HRTEM) image of YNCC (Figure S5a, Supporting Information) discloses that the thickness of the carbon coating is around 10 nm. A representative HRTEM image of YNCC is provided in Figure S5b (Supporting Information). The measured interplanar distances of adjacent lattice planes are 0.203 and 0.209 nm, corresponding to the (400) plane of Co3O4 and the (200) plane of NiO, respectively. From the TEM images of CNCC (Figure 7c,d), the concave nanocubes with nearly hollow and loosely porous structure can be seen. The nanosized pores may come from the interstitial gaps between the nanoparticle and the liberation of gas molecules during calcination. All of these unique structural features are beneficial to lithium storage. Under HRTEM observation (Figure S5d, Supporting Information), this shows that NiO and Co3O4 are finely dispersed with a size of ∼6 nm and a carbon layer of ∼1.5 nm. Furthermore, the (222) plane of Co3O4 and the (111) plane of NiO can be easily recognized. EDS element mapping is used to inspect the element distribution of YNCC and CNCC. The coexistence and uniform dispersion of nickel and cobalt in the core part of YNCC and the whole particle of CNCC indicates that the two types of TMOs form a homogeneous mixture (Figure 8). The carbon is evenly distributed across the whole selected YNCC and CNCC. Moreover, both EDS spectra in Figure S6 (Supporting Information) show strong signals of nickel and cobalt, and the atomic ratio Ni/Co is calculated to be 1.59 and 1.36, consistent with the ICP results. Because the samples are dispersed on a carbon-coated copper grid, the actual carbon content in the nanocomposites could not be exactly reflected by the results from the EDS spectra. So, thermogravimetric analysis (Figure S7, Supporting Information) is conducted, and 9797

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NiO.11,12 Apart from the first cycle, all peaks are nearly superimposable and exhibit good reproducibility, implying the excellent reversibility of the reduction and oxidation during the electrochemical processes. The CV curves of CNCC (Figure 9c) display shapes similar to those of YNCC. In light of the above analysis and the previously reported storage mechanisms of Co3O4 and NiO,11,12,19,20 the whole lithiation/delithiation process of the as-synthesized nanocomposites can be summarized as follows. Co3O4 + 8Li+ + 8e− ↔ 4Li 2O + 3Co

(1)

NiO + 2Li+ + 2e− ↔ Li 2O + Ni

(2)

The representative galvanostatic charge−discharge voltage profiles of YNCC and CNCC at a current density of 100 mA g−1 in the range of 0.01−3.0 V are illustrated in Figure 9b,d. A distinct potential plateau around 1.0 V is clearly identified for both samples in the first discharge process, and it shifts to 1.2 V and remains stable in the following cycles. The upshift of the discharge potential plateau means a more facile electrochemical reduction process after the first cycle. In the meantime, two poorly defined plateaus near 1.6 and 2.2 V are observed in the charge process, agreeing well with the CV curves. The initial discharge and charge capacities are 1196 and 834 mAh g−1 for YNCC and 1142 and 806 mAh g−1 for CNCC, corresponding to irreversible capacity losses of 30.0% and 29.4%, respectively, which could be ascribed to the formation of a SEI layer, the irreversible decomposition of the electrolyte, and the incomplete extraction of lithium from the active materials.11,12 Parts a and c of Figure 10 show the capacity and Coulombic efficiency versus cycle number profiles of YNCC and CNCC at

Figure 9. (a and c) Representative CV curves of YNCC and CNCC at a scan rate of 0.1 mV s−1 between 0.01 and 3 V versus Li/Li+. (b and d) Charge−discharge voltage profiles of YNCC and CNCC for the 1st, 2nd, and 10th cycles in the voltage range of 0.01−3 V at a current rate of 100 mA g−1.

induced morphological and structural changes as well as the electrochemical milling effect in the first cycle.41,42 In the anodic sweep, the peak at ∼1.58 V represents the reversible decomposition of a polymeric gel-like film, which could be attributed to the fact that the in situ generated nickel nanocrystal during the discharge process could act as an effective catalyst.43 The other anodic peak near 2.25 V is associated with the oxidation of cobalt and nickel to Co3O4 and

Figure 10. (a and c) Capacity and Coulombic efficiency versus cycle number profiles of YNCC and CNCC at a current rate of 100 mA g−1 in the voltage range of 0.01−3 V. (b and d) Rate capabilities of YNCC and CNCC. 9798

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Inorganic Chemistry a constant rate of 100 mA g−1. Both samples exhibit a capacity rise in the first 10 cycles, which may be correlated with the insufficient utilization of active materials caused by the inadequate wetting of the electrode and the formation of a polymeric gel-like film by electrolyte decomposition. After that, the capacity experiences a slight drop and stabilizes around 800 mAh g−1. High discharge capacities of 803 and 870 mAh g−1 for YNCC and CNCC can still be retained after 100 cycles. However, under identical test conditions, NiO−Co3O4 exhibits a sharp decline in the discharge capacity, and a capacity of only 376 mAh g−1 is left after 100 cycles (Figure S11a, Supporting Information), probably because of the structure collapse during the continuous volume changes. Although the initial Coulombic efficiencies of YNCC and CNCC are only 70.0% and 70.6%, they promptly increase to 95% within three cycles and stabilize at over 95% afterward, suggesting facile Li+ insertion−extraction and efficient electron transport in the nanocomposites. Furthermore, the morphology and microstructure evolution of the YNCC, CNCC, and NiO−Co3O4 composite electrodes during the charge−discharge processes have been checked. Figure S12 (Supporting Information) shows the SEM and TEM images of the three electrodes after 100 cycles at a current density of 100 mA g−1. This indicates that the yolk@shell structure of YNCC and the concave architecture of CNCC are basically maintained even over a long running time, while the hollow cubic structure of NiO−Co3O4 can hardly be seen, confirming the compositional and structural advantages of YNCC and CNCC. In addition to the inspiring cycling performance, YNCC and CNCC demonstrate splendid rate capabilities as well, as depicted in Figure 10b,d, in which the current density is stepwise increased from 100 to 5000 mA g−1 for every 10 cycles in the range of 0.01−3.0 V. As can be seen from the profiles, even suffering from rapid changes of the current density, both samples still manifest stable capacities at each rate. YNCC (CNCC) exhibits decent average discharge capacities of 865 (911), 727 (874), 578 (676), 448 (475), and 339 (398) mAh g−1 at 100, 200, 500, 1000, and 2000 mA g−1, respectively. At a rigorously high rate of 5000 mA g−1, discharge capacities as high as 281 and 315 mAh g−1 are delivered for the two nanocomposites. Furthermore, upon a decrease in the current density to 100 mA g−1, 744 and 745 mAh g−1 capacities are resumed and with no obvious fading up to 100 cycles, indicating the robustness of the nanocomposites. The superior electrochemical performance of CNCC over that of YNCC may be attributed to its unique configuration. The nanosized building blocks and concave surfaces can reduce the Li+ diffusion distance, make the active material more in order to take part in the electrochemical reaction, improve its ability to accommodate the volume change, and enhance its conductivity and reaction kinetics. The lithium storage performances of the CNCC and YNCC composite electrodes in terms of the cycling capacity and rate capability are superior or comparable to those of other representative anode materials (Table S3, Supporting Information), rendering their use as next-generation anode materials for future LIBs. To gain further insight into the superior electrochemical performances of YNCC and CNCC over the NiO−Co3O4 composite, electrochemical impedance spectroscopy (EIS) measurement has been conducted. The Nyquist plots (Figure S11b, Supporting Information) of the three samples display a depressed semicircle at high-to-medium frequency and a sloped line at low frequency. The diameter of the semicircle is

reflective of the charge-transfer resistance between the electrode−electrolyte interface, and the sloped line corresponds to the Warburg impedance associated with the diffusion of Li+ within the bulk electrode. The smaller diameters of YNCC and CNCC indicate faster charge transfer and lower charge-transfer resistance than those of the NiO−Co3O4 composite, which confirms that the introduction of carbon could improve the electronic conductivity of the as-prepared electrode materials. It can be seen that the slope of CNCC is larger than those of YNCC and NiO−Co3O4, showing unambiguously the structural advantage of CNCC for rapid diffusion of Li+ and thus resulting in enhancement of the rate capability. The remarkable lithium storage capabilities of YNCC and CNCC are believed to originate from their unique compositional and structural features. First, both the porous structure and hollow architecture could enhance the electrode−electrolyte contact area, facilitate the transport of Li+ and penetration of electrolyte molecules into the electrode, and partially accommodate the stress induced by volume variation during repetitive cycles. Second, the carbon layer not only improves the conductivity of the composite and protects the active materials from disintegration and self-aggregation but also improve the mechanical robustness of the electrode materials. Third, the combination of unique properties from NiO and Co3O4 provides a synergetic effect and richer redox chemistry, contributing to enhanced electrochemical performances.44,45 All in all, the above factors favor the superior cycling stability and rate capability of YNCC and CNCC.

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CONCLUSIONS In summary, the yolk@shell or concave NiO−Co3O4@C (YNCC or CNCC) nanocubes have been successfully fabricated by a feasible thermally induced chemical conversion of MOF templates. By virtue of the rational combination of Co3O4, NiO, and carbon and the yolk@shell or concave hierarchical porous architectural characteristics, the two nanocomposite electrodes exhibit excellent lithium storage performances in terms of their high specific capacity, long cycling stability, and good rate capability. Furthermore, these MOFderived high-performance nanocomposites open up a new direction for the development of a wide range of advanced electrode materials potentially applicable in the field of energy storage devices.

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EXPERIMENTAL METHODS

Synthesis of Ni3[Co(CN)6]2 Nanocubes. All of the purchased chemicals and solvents were of analytical grade and were directly used without any further purification. The synthesis of Ni3[Co(CN)6]2 nanocubes was adopted in our previously reported methods with minor modification.33 Briefly, 1.426 g of NiCl2·6H2O and 2.647 g of sodium citrate dihydrate were dissolved in 200 mL of distilled water with magnetic stirring to form a homogeneous solution. Then 100 mL of an aqueous solution containing 1.330 g of K3[Co(CN)6] was slowly added with stirring for 10 min. After that, the above-prepared mixture was incubated at room temperature for 10 h. The pale-blue product was harvested by centrifugation and washed several times with distilled water and ethanol before drying at room temperature overnight. Synthesis of Core@Shell Ni3[Co(CN)6]2@polydopamine. A total of 80 mg of Ni3[Co(CN)6]2 nanocubes along with 80 mg of dopamine were added to 100 mL of a Tris buffer solution (100 mM, pH = 8.5), and a homogeneous suspension was formed by ultrasonic treatment for 30 min. The suspension was then subjected to continuous stirring for 3 h. Afterward, the gray precipitate was 9799

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Inorganic Chemistry collected by centrifugation, rinsed thoroughly with distilled water three times, and then dried at 60 °C in a vacuum oven for 10 h. Synthesis of HCl-Treated Ni3[Co(CN)6]2@polydopamine. A total of 60 mg of Ni3[Co(CN)6]2@polydopamine was dispersed in 60 mL of a 2 M HCl solution under ultrasonication for 10 min. The suspension was then stirred for 1 h, transferred into a Teflon-lined stainless steel autoclave, and heated at 120 °C for 3 h. After cooling to room temperature naturally, the precipitate was collected by centrifugation and washed with distilled water and ethanol several times. Finally, HCl-treated Ni3[Co(CN)6]2@polydopamine was obtained after drying at room temperature for 12 h. Synthesis of YNCC and CNCC Nanocubes. The as-prepared Ni3[Co(CN)6]2@polydopamine or HCl-treated Ni3[Co(CN)6]2@ polydopamine were placed in a porcelain boat and transferred to a tube furnace. Then, they were calcined at 350 °C for 1 h with a ramp of 2 °C min−1 under the protection of argon gas. After that, the argon gas flow was switched off, and the furnace was still kept at this temperature in air for another 1 h. Last, the products were obtained after the furnace cooled to room temperature. NiO−Co3O4 was synthesized by the same procedure except using Ni3[Co(CN)6]2 nanocubes as the precursor. Characterization. The phase purities of all of the products were characterized by a Bruker D8 Focus powder X-ray diffractometer using Cu Kα radiation. The morphologies and structures of the samples were examined by a Hitachi S-4800 field-emission scanning electron microscope. Transmission electron microscopy (TEM), high-resolution transmission electron microscopy (HRTEM), high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM), and energy-dispersive X-ray spectroscopy (EDS) spectra and elemental mapping were recorded on a FEI Tecnai G2 STwin instrument with a field-emission gun operating at 200 kV. Elemental analysis was performed on a JY2000 Ultrace inductively coupled plasma atomic emission spectrometer equipped with a JY AS 421 autosampler and a 2400 g mm−1 holographic grating. Thermogravimetric analysis (TGA) was carried out on a STA 449 °C Jupiter (Netzsch) thermogravimetry analyzer with a heating rate of 10 °C min−1 under air. The Brunauer−Emmett−Teller specific surface areas of the samples were calculated from N2 adsorption−desorption measurement using a Micromeritics ASAP 2010 instrument at −196 °C. Electrochemical Measurements. A CR 2025 coin-type half-cell configuration was used to evaluate the electrochemical behaviors of the products. The working electrodes were prepared by spreading the Nmethyl-2-pyrrolidone slurry containing the active materials carbon black and poly(vinylidene fluoride) in a weight ratio of 70:20:10 on a piece of copper foil by an automatic film coater with a micrometer doctor blade and a vacuum pump. The slurry-coated copper foil was dried at 80 °C under vacuum for 12 h and punched into circular pieces. The load mass of the active material for the testing electrode was about 1.0−1.2 mg. The coin cells were assembled in a glovebox using lithium foil as the counter electrode, a Celgard 2400 membrane as the separator, and 1 M LiPF6 dissolved in ethyl carbonate and diethyl carbonate (1:1 in volume) as the electrolyte. The CV measurements were performed on a BioLogic VMP3 electrochemical workstation in the range of 0.01−3.0 V versus Li/Li+ at a scan rate of 0.1 mV s−1. The charge−discharge performances were conducted on a LAND CT2001A multichannel battery testing system with a cutoff voltage window of 0.01−3.0 V at room temperature. EIS of the electrode was measured with a BioLogic VMP3 electrochemical workstation over a frequency range of 0.01−100 kHz.

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size-distribution histograms, and electrochemical characterization (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Limin Wang: 0000-0001-9618-9239 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS

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REFERENCES

This work is supported by the Creative Research Groups of the National Natural Science Foundation of China (Grant 21521092).

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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/acs.inorgchem.7b01296. SEM and TEM images, EDS spectra, TGA curves, N2 adsorption−desorption isotherms, CV curves, particle9800

DOI: 10.1021/acs.inorgchem.7b01296 Inorg. Chem. 2017, 56, 9794−9801

Article

Inorganic Chemistry

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DOI: 10.1021/acs.inorgchem.7b01296 Inorg. Chem. 2017, 56, 9794−9801