Heterometallic Metal-Organic Frameworks approach to enhancing lithium storage for their derivatives as anodes materials

Inorganica Chimica Acta 494 (2019) 1–7

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Research paper

Heterometallic Metal-Organic Frameworks approach to enhancing lithium storage for their derivatives as anodes materials ⁎

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Bo-Wen Hua, , Ya-Jie Zhua, Lei Dua, Tian-Sheng Mua, Wen-Qi Zhua, Ge-Ping Yina, , Peng Chenb, , Quan-Wen Lic a MIIT Key Laboratory of Critical Materials Technology for New Energy Conversion and Storage, School of Chemistry and Chemical Engineering, Harbin Institute of Technology, Harbin 150001, China b Key Laboratory of Functional Inorganic Material Chemistry (MOE), Heilongjiang University, Harbin 150080, China c College of Chemistry, Nankai University, Tianjin 300071, China

A R T I C LE I N FO

A B S T R A C T

Keywords: Heterometallic approach 2D Metal-Organic Frameworks (MOFs) MOFs-derived bimetal oxides Lithium-ion batteries

Two novel isomorphous 2D NiII/CoII-NiII Metal-Organic Frameworks (MOFs) with tetrazole-1-acetate ligands have been synthesized and structurally characterized. And then they were calcined in air to obtain Ni/Co3O4doped NiO electrodes. Compared to the homometallic MOF-derived oxide, the heterometallic MOF-derived bimetal oxides presented higher coulombic efficiency, better cycle performance, larger specific capacity as lithiumion batteries anodes.

1. Introduction In the past decades, urgent requirements for rechargeable lithiumion batteries (LIBs) have triggered the communities to explore more efficient electrode materials. Currently, the application of LIBs has been significantly extended from conventional portable electronic devices to electric vehicles and smart grids, which requires enhanced specific capacity [1–4]. In this regard, the conventional graphite as anode materials for LIBs cannot meet the requirements due to its low theoretical capacity (∼370 mAhg−1) and poor rate performance [5–7]. To develop advanced alternative anode electrode materials, tremendous efforts have been made [8,9]. Among the proposed materials, transition metal oxides (TMOs) have been intensively investigated as the promising anode electrode because of high theoretical capacity (e.g. ∼890 mAhg−1 for Co3O4) [10,11]. However, the pristine Co3O4 electrode was found to be the subject to low rate performance and severe capacity fading during the lithium intercalation/deintercalation processes, which is attributed to the low natural conductivity and pore collapse during charging/discharging [12]. To address these challenges, binary cobalt-based metal oxides such as ZnCo2O4 [13], CoMn2O4 [14] and CoFe2O4 [15] were emerging. The incorporation of secondary metal in Co oxides has been proved to remarkably improve the rate performance of LIBs through fully utilizing the complementary and synergistic effects of two metals during charging/discharging process [16]. Binary cobalt-based metal oxides are hypothesized to storage

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more Li+ through both the conventional intercalation/deintercalation reactions and the alloying-dealloying process [17]. Furthermore, the improvement of bimetal oxides electrochemical conductivity and stability has a great advantage in the rate performance and stability of LIBs [18,19]. Therefore, the bimetallic oxides are promising for replacing graphite as anode materials. Metal-Organic Frameworks (MOFs), which also called porous coordination polymers (CPs), featured with structural controllability, high surface area and good stability have been used not only as electrode material of LIBs [20–26], but also as precursors for TMOs as anode materials [27–30]. Li’s group have proved that CPs as precursors derived the desired structure and morphology in the nanostructure metaloxides, through the template-engaged-quasi-topotactic transformation process [31], thus to well control the synthesis of bimetallic TMOs with rationally designed porosity, MOFs as precursors have many potentialities. While as self-sacrificial templates or precursors, MOFs are facilely prepared by the self-assembly of bimetal with bridging organic ligands, which include the required metal elements. Therefor the selection of ligands and metal ions for the synthesis of MOFs will affect the properties of MOFs-derived TMOs. Tetrazole-1-acetate as ligand provided aromaticity and abundant coordination sites to construct many porous distinguished multi-functional MOFs [32–38]. The pore structure of MOFs-derived metal oxides gives them certain advantages in electrochemical activity sites of LIBs, while the heterometallic MOFderived bimetal oxides for LIBs are still quite rare. And because of the

Corresponding authors. E-mail addresses: [email protected] (B.-W. Hu), [email protected] (G.-P. Yin), [email protected] (P. Chen).

https://doi.org/10.1016/j.ica.2019.04.049 Received 21 February 2019; Received in revised form 25 April 2019; Accepted 25 April 2019 Available online 26 April 2019 0020-1693/ © 2019 Elsevier B.V. All rights reserved.

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similar coordination modes and atomic radius for CoII ions and NiII ions with the same valence state, it will be easy to build the isomorphous heterometallic MOFs. The combination of Co and Ni oxides is the efficient way to improve the cycling stability and the volumetric energy density of the materials [29,39]. Herein, two new isomorphous 2D MOFs, {[CoxNi2-x(tza)2(Htza)(μ3O)(H2O)]·H2O}n (x = 0 for 1, x = 1 for 2, Htza = tetrazole-1-acetic acid), were synthesized by solvothermal method. The electrochemical performances of two MOF-derived LIBs anode materials have been investigated in detail. Moreover, heterometallic MOF-derived hollow binary metal oxides presented significantly enhanced coulombic efficiency and better rate capacity compared with homometallic MOF-derived oxide.

Chemical composition of mixed-metal MOFs and Co3O4-NiO (Ni, Co) and was analysed by inductively coupled plasma atomic emission spectroscopy (ICP-AES). The morphologies and microstructure were characterized using scanning electron microscopy (SEM, Zeiss Supra55) with an energy-dispersive X-ray (EDX) attachment, transmission electron microscopy (TEM) and high-resolution TEM (HRTEM) on a TECNAT G2 TF20 electron microscope. The Brunauer-Emmett-Teller (BET) were using N2 adsorption/desorption isotherms measured on a nitrogen adsorption equipment (Micromeritics 3H-2000PS1) to calculated specific surface area and pore size distribution. 2.2. Synthesis of 1 and 2 A mixture of Ni(OAc)2·4H2O (0.70 mmol), Htza (1.40 mmol), and NaOH (1 mmol) dissolved in 10 mL EtOH and deionized H2O (1/1, v/v) was sealed in 25 mL teflon-lined stainless vessel. It was heated at 125 °C for 48 h under autogenous pressure and cooled down to room temperature at 2.1 °C h−1. The green block crystals of 1 were harvested in ca. 24% based on NiII ions. Anal. Calcd for C9H14N12Ni2O9: C, 19.59; H, 2.56; N, 30.47. Found: C, 19.53; H, 2.62; N, 30.48. FT-IR (KBr pellet, cm−1): 3469, 3137, 3004, 1636, 1442, 1387, 1104, 859, 805, 704. 2 was obtained using a similar method as for 1 but with a mixture of Ni(OAc)2·4H2O (0.35 mmol) and CoCl2·6H2O (0.35 mmol) instead of Ni (OAc)2·4H2O (0.70 mmol). Dark brown block crystals of 2 were harvested in ca. ∼25% yield based on NiII ions. Anal. Calcd for C9H14N12NiCoO9: C, 19.59; H, 2.46; N, 30.45. Found: C, 19.02; H, 2.81; N, 29.56. FT-IR (KBr pellet, cm−1):3469, 3137, 3004, 1636, 1442, 1387, 1104, 859, 805, 704. ICP-AES for the ratio of Ni:Co in 2 = 0.52: 0.48.

2. Experimental 2.1. Materials and physical measurements All reagents used for the synthesis were purchased commercially and used as received. Single-crystal X-ray diffraction measurements for 1 were carried out on a SuperNova diffractometer at 173(10) K with Mo-Kα radiation (λ = 0.71073 Å) in ω scan mode. The diffraction profiles were integrated using the program CrystalClear [40]. The structure was solved by direct methods using the SHELXS program of the SHELXTL package and refined with SHELXTL [41]. The final refinement was performed by full-matrix least-squares methods with anisotropic thermal parameters for non-hydrogen atoms on F2. Crystallographic data (excluding structure factors) for 1 have also been deposited on the Cambridge Crystallographic Data Centre as supplementary publication (no. 1855935). Copies of the data can be obtained free of charge on application to CCDC, 12 Union Road, Cambridge CB21EZ, UK (fax: (+44) 1223-336-033; e-mail: [email protected]. uk). Crystallographic data for 1 is shown in Table 1. The X-ray powder diffraction (XRPD) patterns were recorded on a Rigaku D/Max-2500 diffractometer at 40 kV and 100 mA using a Cu-target tube and a graphite monochromator. The XRPD pattern were simulated based on the single-crystal data of 1 with the Mercury (Hg) program, which is available free of charge via the Internet at http://www.iucr.org. Elemental analyses were performed on a PerkinElmer 240C analyzer. The FT-IR spectra were recorded from KBr pellets in the range 4000–450 cm−1 on a Nicolet iS5 FT-IR spectrometer. Thermal analyses of 1 and 2 were performed in the temperature range of 30–800 °C on a Rigaku standard TG-DTA analyzer in N2 with an increasing temperature rate of 10 °C/min−1. An empty Al2O3 crucible was used as reference.

2.3. Synthesis of Ni-NiO and Co3O4-NiO Ni-NiO and Co3O4-NiO nanoparticles were prepared by calcining 1 and 2, respectively, at 400 °C for 10 min under air in a tube furnace with a heating rate of 1 °C min−1. ICP-AES results shows that the ratio of Co:Ni in Co3O4-NiO is nearly 0.51:0.49, which is agree with the SEMEDX result of Co/Ni for Co3O4-NiO (1.05:1, Fig. S3). 2.4. Electrochemical measurements Cyclic voltammetry measurements were executed using the CHI 660D electrochemical workstation using the voltage range from 0.01 to 3.0 V at a scanning rate of 0.2 mV s−1. Rate performances of the model cells were evaluated by galvanostatic charging and discharging that were carried out at different current densities in the range of 50–5000 mAg−1. All tests were performed at room temperature. The asobtained active materials (Ni-NiO and Co3O4-NiO) were mixed with Super P carbon and polyvinylidene fluoride with a weight ratio of 7:2:1 in N-methyl-2-pyridinone solvent to form a homogeneous slurry, which was then coated on a copper foil as a current collector and then dried at 120 °C under vacuum conditions for 10 h to form the electrodes. To measure the electrochemical performance, the CR2025-type coin cells were assembled in an Ar-filled glovebox by using lithium metal as the counter and reference electrodes, a porous polypropylene membrane (Celgard 2500) as the separator, and a 1 M solution of LiPF6 dissolved in ethylene carbonate/dimethyl carbonate/ethyl methyl carbonate (1/ 1/1, v/v/v) as the electrolyte. The cells were tested on the NEWAREBTS-5 V/5 mA instrument (Neware Co., Ltd., Shenzhen, China) with galvanostatic charge and discharge between 0.01 and 3 V (vs. Li+/Li).

Table 1 Crystal data and structure refinement parameters for 1. Compound

1

Empirical formula Formula Weight Crystal system Space group a/Å b/Å c/Å α (°) β (°) γ (°) V/Å3 Z D/g cm3 μ/mm−1 Independent reflections GOF on F2 R1,a wR2b [I > = 2σ(I)]

C9H14N12Ni2O9 551.74 triclinic Pī 8.2878(8) 9.2942(13) 12.5671(12) 101.871(9) 100.003(8) 98.351(9) 916.33(18) 2 2.000 2.135 4161 [Rint = 0.0782] 1.033 R1 = 0.0661, wR2 = 0.1821

a b

3. Results and discussion 3.1. Structure characterization

R1 = Σ||Fo|−|Fc||/Σ|Fo|. wR2 = [Σw(Fo2 − Fc2)2/Σw(Fo2)2]1/2.

Compound 1 crystallizes in the triclinic space group Pī with a 2D network. And the asymmetric unit of 1 contains two 2

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Fig. 1. Coordination modes of ligands in 1.

crystallographically independent hexacoordinated Ni(II) ions, two tza ligands, one Htza ligand, one μ3-O ion, one coordinated water molecule, and one free water molecule. There are three coordination modes of ligands: 1) bridging three metal ions with μ3-κO: κO’: κN; 2) bridging three metal ions with μ3-κN: κN’: κO, which is the new coordination mode for Htza ligand, compared with the reported CPs with the same ligand (Table S1); and 3) bridging two metal ions with syn-syn COO– group (Fig. 1). Ni1 is coordinated by two μ3-O atoms [Ni1-O7 = 2.040(4) Å, Ni1O7A = 2.068(4) Å], one water molecule [Ni1-O8 = 2.104(4) Å], two O atoms from different carboxylate groups [Ni1-O2 = 2.045(5) Å, Ni1O5 = 2.077(4) Å], and one N atom on the tetrazole ring from the third tza ligand [Ni1-N9 = 2.078(6) Å]. Ni2 is hexacoordinated by one μ3-O atom [Ni2-O7 = 2.034(4) Å], three carboxylate oxygen atoms from three different tza ligands with the lengths of Ni-Ocarboxylate in the range of 2.036(5) − 2.082(4) Å, and two N atoms from another two ligands [Ni2-N8 = 2.086(5) Å, Ni2-N13 = 2.120(5) Å] (Fig. 2a). Ni1 and Ni2 are connected by μ3-O, κN,N-mode ligand, syn-syn COO– group and synanti COO– group to construct the tetranuclear [Ni4] entity (Fig. 2b). The [Ni4] entities are considered as connecting nodes and connected by eight different ligands to form 2D networks (Fig. 3). TG curves of compounds 1 and 2 shows gradual weight losses of 4.1 wt% for 1 and 3.6 wt% for 2 before 150 °C, corresponding to the loss of one water molecule (∼3.3 wt%) from the dissociating solvent. And the frameworks of 1 and 2 were stable until 210 °C (Fig. S2). With the temperature increasing, one tza ligand and one coordinated water molecule were gradually released from 215 °C and 350 °C and the decomposition was starting after that. Hence, annealing the as-prepared compounds 1 and 2 at 400 °C in an air flow for 10 min is sufficient to ensure complete obtain Ni-NiO and Co3O4-NiO, respectively. The phase purity of compounds 1 and 2 were confirmed by X-ray powder diffraction (XRPD). The XRPD diagrams of compounds crystallizing in the same phase were displayed in Fig. S1. The composition of MOFs-derived materials were also confirmed by the XRPD. The diffraction peaks of NiII MOFs-derivative, Ni-NiO, were well fitted by standard cubic Ni (JCPDS 04-0850) and NiO (JCPDS 47-1049) (Fig. 4a). And the ones of heterometallic MOFs-derivative, Co3O4-NiO, were exactly fitted by NiO (JCPDS 47-1049) and Co3O4 (JCPDS 43-1003)

Fig. 3. 2D network of 1.

(Fig. 4b). The PXRD of Co3O4-NiO did not show well defined diffraction peaks. It may due to the form of amorphous phase carbon by the effect originated from the carbonization of ligands [42].

3.2. Morphology characterization To investigate the hollow structure and morphology of MOFs and their derived composition metal oxides, SEM, TEM and representative high-resolution TEM (HRTEM) images were shown as follows. SEM images showed that heterometallic MOFs-derivatives Co3O4-NiO have more uniform pore distribution, bigger pore size and higher pore density than NiII MOFs-derivative Ni-NiO (Fig. 5). Such porous structure of heterometallic MOFs-derivative Co3O4-NiO might lead to enhanced specific capacity, rate performance and lifetime because of tremendous sites for Li intercalation/deintercalation, short lithium-ion diffusion pathway and accommodation for volume expansion [43–45]. The TEM images were consistent with the results from SEM images (Fig. 6). In particular, Co3O4-NiO exhibited larger pores than Ni-NiO electrodes. The HRTEM images of the two MOFs-derived materials are shown in Fig. 6c and 6f, respectively. It should be noted that Ni-NiO and Co3O4-NiO electrodes retained the high crystallinity and had distinct

Fig. 2. (a) View of coordination environment of nickel atoms in 1 (H atoms omitted for clarity, symmetry mode A: −x + 1, −y + 2,−z). (b) The linkage and coordination mode in 1. NiII ions are shown as light blue polyhedron, C as grey, O as red, and N as blue. 3

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Fig. 4. XRPD patterns of the (a) Ni-NiO and (b) Co3O4-NiO.

lattice fringes. The measured lattice fringes of ∼0.074 nm, 0.120 nm and 0.176 nm can be assigned to the cubic NiO (2 2 0), cubic NiO (1 1 1) and Ni (2 0 0) facets, respectively. In the meantime, as shown in Fig. 6f, the spacing of 0.120 nm is in good agreement with the (1 1 1) plane of NiO, and additional space lattices of 0.202 nm and 0.244 nm are observed to assign to the (4 0 0) and (3 1 1) planes of cubic Co3O4, which also confirmed the composition of Co3O4-NiO electrode. In order to better understand hollow microstructure of MOFs-derivative, N2 adsorption–desorption isotherm at 77.3 K were measured (Fig. S4). The N2 uptake of Ni-NiO and Co3O4-NiO at P/P0≈1 were 148.86 and 176.77 mLg−1. The BET specific surface area of Ni-NiO and Co3O4-NiO were calculated to be 22.65 and 25.05 m2g−1 with a pore volume of 0.23 and 0.27 mLg−1, respectively. In the insert of Figs. S4a and S4b, the pore sizes concentrate in 3.82 and 4.22 nm, respectively, which are much larger than Li+ radius (0.076 nm). The porosity of anode materials can influence the site density in Li-ion intercalation/ deintercalation [46].

3.3. Electrochemical performance In order to evaluate the effect of NiII MOFs and heterometallic MOFs as precursors on lithium ion batteries, the lithium storage performance of the derived oxides obtained after calcination as anodes were explored. CV curves of Ni-NiO and Co3O4-NiO electrode operated at a scan rate of 0.2 mV s−1 (Fig. 7a and b). For the Ni-NiO electrode, it was observed that there appears a cathodic peak at 0.51 V and an anodic band at 2.20 V in the first cycle, which was attributed to the lithiation and delithiation of Ni-NiO electrode (NiO + 2Li+ + 2e-1 ↔ Li2O + Ni) [47]. For the Co3O4-NiO electrode (Fig. 7b), in the first cathodic sweep, two peaks located at potentials of 0.57 and 0.94 V were attributed to the Li+ insert in both NiO and Co3O4 (Co3O4 + 8Li+ + 8e−1 ↔ 4Li2O + 3Co) [47]. The main cathodic peak shifted to a higher potential at 1.08 V in the subsequent cathodic scans, due to the formation of solid electrolyte interface (SEI) in the first cycle [48,49]. The two anodic peaks at 1.4 and 2.21 V correspond to the reverse reaction,

Fig. 5. SEM images of (a) NiII MOFs, (b) Ni-NiO electrode, (c) NiII-CoII MOFs and (d) Co3O4-NiO electrode. 4

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Fig. 6. TEM images of (a) NiII MOFs, (b) Ni-NiO electrode, (d) NiII-CoII MOFs and (e) Co3O4-NiO electrode. HRTEM images of (c) Ni-NiO and (f) Co3O4-NiO electrodes.

during the initial cycles are significantly different from those after 50 cycles, which may be related to the phase transformation of oxides. Turning back to Fig. 7a, the phase transformation takes place after 2 cycles, likely leading to changes in Li-ion storage mechanism, which is under investigation. For Co3O4-NiO electrode (Fig. 7d), the 1st, 2nd, 50th and 85th charge/discharge profiles were measured between 0.01 and 3.0 V at a current density of 200 mA g−1, and two long voltage plateau at about 0.7 V and 1.0 V were observed, corresponding to Li+ insertion into Co3O4 and NiO in the first discharge process [53,54]. It showed a higher specific capacity of 474.9 mAhg−1 after 85 cycles than Ni-NiO electrode (332.7 mAhg−1), because of the synergy of heterometallic oxides [55]. The coulombic efficiency of Co3O4-NiO electrode is continuously improved, maintaining a high coulombic efficiency of 98.1% after 85 cycles (Fig. 7f). Fig. 7e and f showed the cycle performance of Ni-NiO and Co3O4NiO composite at a current density of 200 mAg−1. As shown in Fig. 7e, as the cycle numbers increased, the specific capacity of Ni-NiO gradually decreased. After 69 cycles the specific capacity began to rise slowly, and at 85 cycles the discharge capacity rose to 332.7 mAhg−1, and the corresponding capacity retention rate is 45.6%. For Co3O4-NiO composite, it showed a slightly higher specific capacity than that of Ni-NiO

including the Ni to NiO and Co to Co3O4. From the second cycle, the intensity of the cathodic and anodic peaks overlapped very well, indicating good reversibility of the Co3O4-NiO electrode. As NiO and Co3O4 have different electrochemical activities at the same charging/ discharging potential stage, it provide the stepwise lithium storage mechanism to present the good cycling performances by their synergetic effect. As shown in Fig. 7c, the galvanostatic charge/discharge curves of the Ni-NiO electrode were displayed between 0.01 and 3.0 V at a current density of 200 mAg−1 for the 1st, 2nd, 50th and 85th cycles. For the first discharge curve of the Ni-NiO electrode, there was a longer discharge platform around at 0.7 V. The first discharge and charge specific capacities of Ni-NiO were 729.6 and 534.4 mAhg−1, respectively, and the corresponding initial coulomb efficiency is 73.2% (Fig. 7e). This large irreversible capacity is corresponding to the irreversible electrochemical reaction of forming a SEI film on the surface of electrode materials [27,50], and electrolyte deterioration decrease the electric contact between Ni-NiO and also between current collector [51,52]. In the subsequent cycle, the coulombic efficiency of Ni-NiO was continued to increase, and the high coulombic efficiency of 97.6% was maintained after 85 cycles. In addition, the profiles of curves 5

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Fig. 7. Electrochemical performances of the Ni-NiO and Co3O4-NiO electrodes at 0.01–3 V. CVs of the (a) Ni-NiO electrode and (b) Co3O4-NiO electrode. Discharge/ charge curves at 200 mA g−1 for the (c) Ni-NiO electrode and (d) Co3O4-NiO electrode. Cycle performance and coulombic efficiency curves at 200 mA g−1 for the (e) Ni-NiO electrode and (f) Co3O4-NiO electrode. Rate performance at diffident current densities for the (g) Ni-NiO electrode and (h) Co3O4-NiO electrode.

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(Fig. 7f), the first discharge and charge specific capacities are 779.9 mAhg−1 and 539.6 mAhg−1, respectively, and the coulomb efficiency is 69.2%. The specific capacity of Co3O4-NiO gradually decreased in the subsequent cycle process. After 55 cycles, the discharge capacity decreases to 412.5 mAhg−1. Subsequently, the specific capacity began to slowly rise. At 85 cycles, the discharge capacity arrived to 474.9 mAhg−1, and the capacity retention rate is 60.9%, which verified the high specific capacity of the Co3O4-NiO composite for LIBs. Based on the above facts, compared to Ni-NiO electrode, Co3O4-NiO composite showed higher discharge specific capacity, because of heterometallic synergy effect providing larger pore structure and more electrochemical active sites, which showed benefit for LIBs [56,57]. For example, the larger pore structures are available to facilitate the electrode material to store lithium and improve its utilization rate; on the other hand, it can make the electrolyte infiltrate into the electrode material better and shorten the diffusion path of lithium ion and electron [58,59]. In short, the porous Co3O4-NiO exhibited better cycle performance and higher specific capacity than Ni-NiO composite mateiral. Furthermore, compared with the electrochemical properties of Co3O4 and NiO reported in the literature (Table S2), this Co3O4-NiO nanocomposite showed good specific capacity and high coulomb efficiency, because of their larger surface area and shorter lithium ion diffusion path, which inherited from the morphological features of NiIICoII MOFs precursors. As shown in Fig. 7g and h, the rate performance of Ni-NiO and Co3O4-NiO were further evaluated at various current densities of 50, 100, 200, 500, 1000, 2000, 3000 and 5000 mAg−1. The Ni-NiO delivers average reversible capacities of 536.7, 497.3, 406.1, 293.9, 200.9, 122.1, 73.7 and 57.5 mAhg−1, respectively. Meanwhile, the average reversible capacities of Co3O4-NiO composite are 592.4, 584.7, 538.0, 448.2, 358.7, 270.8, 207.8 and 167.3 mAhg−1 at various current densities, respectively. When the current density returned back to 50 mAg−1, Co3O4-NiO composite still maintained high specific capacity of 571.4 mAhg−1 (96.5% of the capacity recovered), while Ni-NiO electrode (442.07 mAhg−1) only had 82.4% of the capacity recovered. That indicated good structure stability of the nanocomposite, and also confirmed that the Co3O4-NiO has a great promise as a high rate anode material for LIBs.

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4. Conclusions In summary, two novel isomorphous MOFs have been constructed by tetranuclear [M4] clusters and organic linker tza with 2D network. The Ni-NiO and Co3O4-NiO electrodes inherited the structure of MOFs precursors were formed by calcination. The heterometallic MOFs-derived Co3O4-NiO electrode exhibited good specific capability, better cycle stability and perfect coulombic efficiency. This work provides an effective heterometallic approach to enhancing electrochemical performance for LIBs anodes. Conflicts of interest There are no conflicts to declare. Acknowledgements This work was supported by the National Natural Science Foundation of China (No. 21805064), the China Postdoctoral Science Foundation (2016 M591519), Heilongjiang Province Postdoctoral Science Foundation (LBH-Z16069) and Heilongjiang Province Science Foundation (LC2018005). And the International Clean Energy Talent Program of China Scholarship Council is acknowledged. Appendix A. Supplementary data Supplementary data to this article can be found online at https:// 7