High-performance lithium storage based on the

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Jul 26, 2017 - Jawayria Mujtaba a, 1, Hongyu Sun a, b, *, 1, Yanyan Zhao a, Guolei Xiang c, Shengming Xu d,. Jing Zhu a, ** a National Center for Electron ...
Journal of Power Sources 363 (2017) 110e116

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High-performance lithium storage based on the synergy of atomicthickness nanosheets of TiO2(B) and ultrafine Co3O4 nanoparticles Jawayria Mujtaba a, 1, Hongyu Sun a, b, *, 1, Yanyan Zhao a, Guolei Xiang c, Shengming Xu d, Jing Zhu a, ** a

National Center for Electron Microscopy in Beijing, School of Materials Science and Engineering, The State Key Laboratory of New Ceramics and Fine Processing, Key Laboratory of Advanced Materials (MOE), Tsinghua University, Beijing 100084, China Department of Micro- and Nanotechnology, Technical University of Denmark, Kongens Lyngby 2800, Denmark c Melville Laboratory for Polymer Synthesis, Department of Chemistry, University of Cambridge, Lensfield Road, Cambridge CB2 1EW, United Kingdom d Institute of Nuclear and New Energy Technology, Tsinghua University, Beijing 100084, China b

h i g h l i g h t s  Ultrafine Co3O4 nanoparticles @ TiO2(B) ultrathin nanosheets are synthesized.  The sample exhibits good lithium storage properties.  The excellent properties are related to the synergistic effect of the composites.

a r t i c l e i n f o

a b s t r a c t

Article history: Received 25 May 2017 Received in revised form 21 July 2017 Accepted 22 July 2017 Available online 26 July 2017

Lithium ion batteries (LIBs) are critical constituents of modern day vehicular and telecommunication technologies. Transition metal oxides and their composites have been extensively studied as potential electrode materials for LIBs. However, inefficient lithiation, poor electrical conductivity, and drastic volume change during cycling result in low reversible capacity and rapid capacity fading, and thus hinder the practical applications of those electrodes. In this work, we report a facile synthesis of a novel hierarchical composites, which consist of ultrafine Co3O4 nanoparticles uniformly dispersed on TiO2(B) nanosheets with atomic thickness (Co3O4 NPs@TiO2(B) NSs). When tested as anode material for LIBs, the Co3O4 NPs@TiO2(B) NSs sample with optimized composition shows a reversible capacity of ~677.3 mAhg1 after 80 cycles at a current density of 100 mAg1. A capacity of 386.2 mAhg1 is still achieved at 1000 mAg1. The synergistic effect of ultrafine Co3O4 nanoparticles and atomic-thickness TiO2(B) nanosheets is responsible for the enhanced electrochemical performance. © 2017 Elsevier B.V. All rights reserved.

Keywords: Lithium ion batteries Anode materials Co3O4 nanoparticles TiO2(B) nanosheets Synergistic effect

1. Introduction Rechargeable Lithium ion batteries (LIBs) have enticed significant attention as a favorable and sustainable green options for

* Corresponding author. Department of Micro- and Nanotechnology, Technical University of Denmark, Kongens Lyngby 2800, Denmark. ** Corresponding author. National Center for Electron Microscopy in Beijing, School of Materials Science and Engineering, The State Key Laboratory of New Ceramics and Fine Processing, Key Laboratory of Advanced Materials (MOE), Tsinghua University, Beijing 100084, China. E-mail addresses: [email protected] (H. Sun), [email protected] (J. Zhu). 1 These authors contributed equally to this work. http://dx.doi.org/10.1016/j.jpowsour.2017.07.076 0378-7753/© 2017 Elsevier B.V. All rights reserved.

energy storage [1], and become a vital part of the realm of portable equipment and vehicle electrification as energy storing candidate [2]. Nevertheless, the specific energy of the commercialized LIBs is insufficient due to the low capacity of electrode materials. Developing advanced electrodes is one of the key elements that relates to the battery performance improvement. In recent years, great efforts have been paid to design high performance anodes, especially transition metal composites (typically including oxides and sulfides) with different morphologies and structures. For example, iron oxides, cobalt oxides, tin oxide, nickel sulfide, molybdenum sulfide with one dimensional (nanowires, nanorods, nanotubes), two dimensional (nanosheets, nanoflakes), three dimensional hierarchical architectures, hybrids and complex nanostructures have been investigated as exhilarating anodes for LIBs [3e10]. Among

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them, nanostructured transition metal oxides, which are based on the conversion reaction mechanism, are of certain relevance for their diverse choices [11], such as high specific capacity (ranging from 500 to 1000 mAhg1, higher than the commercial used graphite with a capacity of 372 mAhg1), large specific surface area and a shorter path length for solid-state ion diffusion in comparison with their bulk counterparts [12,13]. However, there are two main issues generally associated with metal oxide anodes: (i) poor conductivity limits the high-rate performance [14], and (ii) severity in the volume variation during repetitive lithiation/delithiation processes results in the deterioration of reversible capacity and rapid capacity fading [15,16]. Cobalt oxide (Co3O4) and titanium dioxide (TiO2) possess theoretical specific capacities of 890 and 335 mAhg1, respectively, and have been extensively investigated as anode materials due to the advantageous of plentifulness and environmental amiability [1,2,6]. Both materials, however, show the drawbacks when used as anodes. For example, the poor cyclability and rate capability in Co3O4 and the low specific capacity in TiO2 anodes interfere with their using as energy storage materials in everyday applications [3,5,7,9]. To incorporate these problems, we synthesized a synergistic material by combining ultrafine Co3O4 nanoparticles with atomicthickness TiO2(B) nanosheets. The synthesized composite was employed as an anode material for LIBs and its electrochemical properties were studied. Although many nanocomposites based on Co3O4 and/or TiO2 structures, such as Co3O4/CNTs [4,10] and Co3O4/ graphene [8,11], have been reported for lithium storage [12e14], to the best of our knowledge, synthesis of ultrafine Co3O4 nanoparticles decorated on atomic-thickness nanosheets of TiO2(B) (Co3O4 NPs@TiO2(B) NSs) anodes have not been investigated. 2. Experimental 2.1. Materials synthesis Synthesis of Co3O4 ultrafine nanoparticles (sample 1, S1). All the chemicals were used directly without further purification. In a typical synthesis, Cobalt(II) nitrate hexahydrate (Co(NO3)2$6H2O, 0.363 g) was dissolved in ethanol (25 mL), then a mixture of oleylamine (2.5 mL) and ethanol (12.5 mL) was added into the solution under magnetic stirring for 1 h. The green mixed solution was transferred into an autoclave (50 mL) with a Teflon liner at 180  C for 12 h. The product was filtered and washed with cyclohexane and ethanol several times, and dried under vacuum at 80  C for 5 h, and finally calcined in air at 350  C for 4 h. Synthesis of TiO2(B) nanosheets with atomic thickness (sample 4, S4). TiO2(B) nanosheets were prepared as we reported previously [17]. In brief, titanium(III) chloride aqueous solution (TiCl3, 1.25 mL) and deionized water (1.25 mL) were mixed with ethylene glycol (37.5 mL) in a 50 mL Teflon-lined autoclave. After stirred with a glass rod, the mixture was sealed and then kept at 150  C for 4 h in an oven. The obtained products were separated and washed with water and ethanol several times. Powder samples were dried at 60  C in an oven, and then calcined at 250  C for 4 h in a vacuum oven. Synthesis of Co3O4 NPs@TiO2(B) NSs nanocomposite (samples 2, 3; S2, S3). In a typical procedure, different amount of Co(NO3)2$6H2O (0.363 g for S2; 0.073 g for S3) was dissolved in ethanol (25 mL), then a mixture of oleylamine (2.5 mL), ethanol (12.5 mL) and TiO2(B) powder (sample S4, 25 mg) was added into the solution under magnetic stirring for 1 h. The green mixed solution was transferred into an autoclave (50 mL) with a Teflon liner at 180  C for 12 h. The product was filtered and washed with cyclohexane and ethanol several times, and dried under vacuum at 80  C for 5 h, and finally calcined in air at 350  C for 4 h.

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2.2. Materials characterization Crystallographic information for the samples was checked using a Bruker Model D8 Advanced powder X-ray diffractometer (XRD) with Cu Ka irradiation (l ¼ 1.5418 Å). The morphology and microstructure were examined using transmission electron microscopy (TEM; JEOL, JEM-2100, 200 kV; FEI, Tecnai G2 20, 200 kV), X-ray photoelectron spectroscopy (XPS, Escalab 250, Al Ka), and Raman microscopy (Renishaw, UK, 633 nm excitation). The BrunauerEmmett-Teller (BET) surface area was analyzed by nitrogen adsorption-desorption isotherm at 77 K in a Micromeritics ASAP 2010 system. The samples were degassed at 180  C before nitrogen adsorption measurements. The surface area was determined by a multipoint BET method. A desorption isotherm was used to determine the pore size distribution via the Barret-Joyner-Halender (BJH) method. The nitrogen adsorption volume at the relative pressure (P/P0) of 0.994 was used to determine the pore volume and average pore size. 2.3. Electrochemical measurements The electrodes were constructed by mixing the active materials, conductive carbon black and carboxymethyl cellulose in a weight ratio of 70:20:10. The mixture was prepared as slurry and spread onto copper foil, then dried under vacuum at 120  C for 6 h to remove the solvent before pressing. The electrodes were cut into disks (12 mm in diameter) and dried again at 100  C for 24 h in vacuum. The loading mass of the active materials on the copper foil was ca. ~2.4e3 mgcm2 for each electrode. The cells were assembled inside an Ar-filled glove box by using a lithium metal foil as the counter electrode and the reference electrode and microporous polypropylene as the separator. The electrolyte used was 1M Lithium hexafluorophosphate (LiPF6) dissolved in a mixture of ethylene carbonate (EC), propylene carbonate (PC), and diethyl carbonate (DEC) with a volume ratio of EC/PC/DEC ¼ 1:1:1. Assembled cells were allowed to soak overnight, and then electrochemical tests on a LAND battery testing unit were performed. Galvanostatic charge and discharge of the assembled cells were performed at a current density of 100 mAg1 between voltage limits of 0.01 and 3.0 V (vs. Li/Liþ) for 80 cycles. For the high rate tests, the discharge current gradually increased from 100 mAg1 to 200, 500, and 1000 mAg1, and then decreased to 100 mAg1, step by step. The specific capacity was calculated on the basis of the weight of prepared products contained in the anode. The cyclic voltammogram (CV) curves were recorded between 0.01 and 3.0 V (vs. Li/Liþ) at a scan rate of 0.5 mVs1 by using a CHI 660D electrochemical workstation (Chenhua Instrument, Shanghai). 3. Results and discussion The as prepared samples were firstly analyzed by XRD as shown in Fig. 1 to study the phase structures. All peaks in samples S1 and S4 are well matched with the spinel structured Co3O4 (JCPDS No. 74e1657) and metastable polymorph TiO2(B) (JCPDS No. 46e1237). Meanwhile, no other diffraction peaks from possible impurities can be detected, indicating the high phase purity of the samples. For Co3O4 NPs@TiO2(B) NSs nanocomposites, especially sample S2, only the peaks belonging to Co3O4 phase can be identified. The absence or depressing some peaks of TiO2(B) phase may due to its low weight content in the composite (see the discussion below) or the preventing of the layer stacking of TiO2(B) nanosheets by the Co3O4 nanoparticles. Similar results have also been reported in other two dimensional materials [18,19]. The detailed structural investigation is studied by TEM and HRTEM analysis. Fig. S1 shows the morphology and microstructure

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Fig. 1. XRD patterns of pure Co3O4 (sample S1), TiO2(B) (sample S4), and Co3O4 NPs@TiO2(B) NSs nanocomposites (samples S2 and S3).

of pure Co3O4 nanoparticles (sample S1). The as-prepared particles are monodispersed with a uniform size of ca. 5.2 nm (Fig. S4). Selected area electron diffraction (SAED) pattern (Fig. S1d) demonstrates the polycrystalline nature. Fig. 2 shows typical TEM and HRTEM images of Co3O4 NPs@TiO2(B) NSs nanocomposite (sample S2). Compared to the pure TiO2(B) nanosheets (Fig. S3), the surface of the sample S2 is coarse and wrinkled (Fig. 2a and b). The higher magnification TEM images (Fig. 2d and e) indicate a large number of ultrafine Co3O4 nanoparticles are uniformly distributed on atomicthickness TiO2(B) nanosheets. The dark regions in the sample are due to the fold of TiO2(B) nanosheets. The clear lattice fringes with spacing values of 0.47 and 0.35 nm are in good agreement with the

Fig. 2. TEM, HRTEM images and SAED pattern of Co3O4 NPs@TiO2(B) NSs nanocomposite (sample S2).

(111) and (110) crystalline planes of Co3O4 and TiO2(B), respectively (Fig. 2f). SAED pattern (Fig. 2c) illustrates the polycrystalline nature of the sample. The obtained Co3O4 NPs@TiO2(B) NSs nanocomposite forms a three dimensional hierarchical architecture, which integrates zero dimensional ultrafine nanoparticles and two dimensional ultrathin nanosheets together. This special configuration facilitates the dispersion and prevents the agglomeration of each nanostructure, and thus plays a vital role in lithium storage. The structure of sample 3 (Fig. S2) shows a similar character with that of sample 2. The transparent nature of pure TiO2(B) nanosheets (Fig. S3) confirms the thickness is down to atomic scale. The composition of the samples was further validated by using energy dispersive X-ray spectroscopy (EDX) technique (Fig. S5, the Cu and C signal comes from the supported grid). For samples S2 and S3, the weight ratios of TiO2 to Co3O4 in the composites were estimated to be 0.05 and 3.25 by quantitative EDX calculations. Fig. S4 also shows the size distribution of the Co3O4 nanoparticles in samples S2 and S3. The average diameters are ~4.0 and ~2.8 nm for samples S2 and S3, respectively. Therefore, the addition of TiO2(B) nanosheets during the synthesis is favorable for the formation of ultrafine Co3O4 nanoparticles. To further understand the nature of the Co3O4 NPs@TiO2(B) NSs nanocomposites, Raman spectra were recorded (Fig. 3). The results for pure Co3O4 (sample S1) and TiO2(B) (sample S4) are also shown for comparison. For sample S1, there are five active Raman modes, which are approximately located at ~189, ~470, ~510, ~608 and ~677 cm1. All the observed modes are in great agreement with the values of pure Co3O4 (191, 470, 510, 608 and 675 cm1) [20,21]. In sample S4, eight active Raman modes approximately located at ~207, ~251, ~382, ~421, ~478, ~554, ~631 and ~831 cm1 can be well distinguished. The values are consistent with that of pure TiO2(B) phase [22]. The Raman spectra for samples S2 and S3 show a combined characteristic of both Co3O4 and TiO2 (B), demonstrating the co-existence of the both phases in the nanocomposites. The chemical composition and surface oxidation state of the samples were analyzed by XPS in the region of 0e1350 eV. The survey XPS spectra (Fig. S6) indicate according elements contained

Fig. 3. Raman spectra of ultrafine Co3O4 nanoparticles (sample S1), TiO2(B) nanosheets with atomic thickness (sample S4), and Co3O4 NPs@TiO2(B) NSs nanocomposite (samples S2 and S3).

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in the samples, which are in accordance with the above EDX measurements. The high-resolution XPS spectra of Ti 2p, O 1s, and Co 2p regions for the samples are shown in Fig. 4. For the pure Co3O4 nanoparticles (sample S1), the core level spectrum of Co 2p shows two peaks assigned to Co 2p3/2 and Co 2p1/2 at binding energies of 778.7 and 795.1 eV, respectively (Fig. 4a). The satellite peaks in Co 2p spectrum suggest the formation of Co3O4 phase. While for the Ti 2p scan of pure TiO2(B) nanosheets (sample S4), two distinct binding energy peaks located ~464.8 and ~459.1 eV correspond to the electronic states of Ti 2p1/2 and Ti 2p3/2, respectively (Fig. 4d). Similar to sample 1, the high-resolution XPS spectra of Co 2p for samples S2 and S3 (Fig. 4b and c) can be best

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fitted by eight peaks with two pairs of spin-orbit doublets and their four shakeup satellites (denoted as “sat”). The doubles are characteristic of the peaks of Co2þ and Co3þ. In all the samples, the binding energy component observed at ~530.0 eV is attributed to O 1s, which means the formation of oxides with titanium and/or cobalt [23]. The O 1s peak can be deconvoluted into three components, with peaks at ~529, 530 and 532 eV, and assigned as lattice oxygen, OH group, and molecular water in the sample and/or adsorbed on the surface. Among them, the OH group is directly related with defects [24e26]. The relationship between the ratio of OH group to lattice oxygen for all samples is shown in Fig. S7. The results reveal that the oxygen vacancies in the samples can be adjusted by

Fig. 4. High-resolution XPS spectra of the Ti 2p, O 1s, and Co 2p regions of the samples: (a) ultrafine Co3O4 nanoparticles (sample S1), (b, c) Co3O4 NPs@TiO2(B) NSs nanocomposites (samples S2 and S3), and (d) TiO2(B) nanosheets with atomic thickness (sample S4).

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changing the ratio between Co3O4 nanoparticles and TiO2(B) nanosheets, which is important to optimize the lithium storage properties. The specific surface area and porous nature of the products were determined by using nitrogen adsorption-desorption isotherm. The N2 adsorption-desorption isotherms for all samples at 77 K are presented in Fig. 5a. Hysteresis loops for the products at relative pressure are in the range of 0.4e1.0 P/P0. Barret-Joyner-Halenda (BJH) pore size distribution for the corresponding samples are shown in the inset in Fig. 5a, the average pore diameter approximately being ~ 4 nm. A more detailed graphical representation of pore diameter and surface area is given in Fig. 5b. The specific surface area of sample S1, S2, S3 and S4, being ~99.5, ~86.42, ~226 and ~320 m2 g-1, respectively. The higher surface area of Co3O4 NPs@TiO2(B) NSs nanocomposites may be attributed to the mesoporous nature of the atomic-thickness nanosheets and the void spaces between the ultrafine nanoparticles. The porous nature of Co3O4 NPs@TiO2(B) NSs is of importance for lithium storage as they have the capability to provide extra active reaction sites and facilitate mass diffusion and ion transport. The obtained samples were directly used as active materials and mixed with conductive carbon black and carboxymethyl cellulose to construct CR 2032 coin-type half cells for electrochemical studies. Fig. 6a shows typical CV curves of S2 electrode for the first three cycles. A dominant cathodic peak at ~0.78 V can be observed in the first cycle and then disappears in the subsequent cycling process. This peak is related to the formation of solid electrolyte interface (SEI) layer on the electrode surface and Li2O due to the electrochemical reduction (lithiation) reaction of active materials with Li. Moreover, a small reduction peak at ~1.6 V is assigned to the intercalation of Liþ in TiO2(B) [27]. The corresponding oxidation peak at ~2.1 V is ascribed to the reaction of metallic cobalt into oxides and the deintercalation of Liþ from TiO2(B) lattice. In the following two cycles, the reduction peaks shift to higher potential direction, which may due to the irreversible structural change of the oxides during the cycling [28]. The discharge-charge profiles of sample S2 at a current density of 100 mAg1 in the voltage cutoff window of 0.01e3 V (vs. Li/Liþ) are shown in Fig. 6b. In the first discharge curve, the potential value quickly falls to a plateau of ~1.0 V. The extended plateau with a capacity of ~784 mAhg1 may likely be assigned to the conversion of Co3O4 to Co, and then gradually declines to the cutoff voltage (0.01 V). This decrease in discharge capacity could be associated with TiO2 coating effect [29]. It should be noted here that reversible capacity related to TiO2(B) at ~1.6 V [27] cannot be detected in the profile. We attribute it to the

low weight ratio of TiO2 to Co3O4 in the composite (0.05). The sample S2 delivers first discharge and charge capacities of 1205.8 and 882.6 mAhg1, respectively, yielding an irreversible capacity loss of 26.8%. Such initial irreversible capacity loss originates from the formation of SEI layer due to the irreversible degradation of the electrolyte and other irreversible side effects, the trapping of lithium ions inside the active material, and the electrolyte decomposition catalyzed by Co, which is common for electrode materials based on conversion reaction mechanism [30e36]. The following two discharge/charge curves tend to be stable and exhibit similar electrochemical behavior. The comparison of cycling performance of the four electrodes at a current rate of 100 mAg1 up to 80 cycles is shown in Fig. 6c. It is obvious that the sample S2 shows the highest lithium storage capacity and the best cycling performance. The discharge specific capacity steadily drops from 1200 mAhg1 (the first cycle) to ~635 mAhg1 after around 20 cycles. Accordingly, the Coulombic efficiency increases from 73% in the first cycle to ~99.5% in the following cycles. The capacity changes in the first several cycles may due to the irreversible side reactions and microstructure adjustment of the electrodes. After 80 cycles, the reversible capacity of S2 cell is 677.3 mAhg1, which is higher than that of S1 (224.4 mAhg1), S3 (430 mAhg1), as well as S4 cells (179 mAhg1). The morphology and microstructure of the S2 cell after the cycling performance test were characterized by TEM observations as shown in Fig. S8. It find that the sample still maintains the initial morphology and structure even after 80-cycle repetitive charging-discharging test, demonstrating the good structural stability of the hierarchical architectures. The rate performance curves for all four samples are given in Fig. 6d, which were tested successively at current densities of 100, 200, 500, 1000 and 100 mAg1 (each for 10 cycles). It can be seen that the reversible capacity decreases regularly with an increased current density for all the samples. Samples S2 clearly shows the best rate capacities as compared to the other samples since S2 cell possesses the highest reversible capacities at various current rates. Specifically, the reversible capacity changes from 758.8 mAhg1 to 386.2 mAhg1 at the current densities of 100 mAg1 and 1000 mAg1, and the capacity returns to 816 mAhg1 when the current density decreases back to 100 mAg1. The above lithium storage measurements show that the Co3O4 NPs@TiO2(B) NSs nanocomposites, especially the sample S2, possess good lithium storage properties in terms of specific capacity, rate capability, and cycling stability due to the synergistic effects of ultrafine nanoparticles and atomic-thickness nanosheets. Firstly, the mesoporous nature of the Co3O4 NPs@TiO2(B) NSs

Fig. 5. (a) Nitrogen adsorptionedesorption isotherms and corresponding pore size distribution curve (inset) of the ultrafine Co3O4 nanoparticles (sample S1), TiO2(B) nanosheets with atomic thickness (sample S4), and Co3O4 NPs@TiO2(B) NSs nanocomposites (samples S2 and S3), (b) Surface area (purple) and pore diameter (Black) of the samples. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

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Fig. 6. Electrochemical properties of the samples: (a) CVs at a scan rate of 0.5 mV s1 between 0.01 and 3.0 V, (b) Galvanostatic charge/discharge voltage profiles for the 1st, 2nd, and 3rd cycles between 0.01 and 3.0 V versus Li/Liþ at a current density of 100 mA g1, (c) cycling performance at a constant current rate of 100 mA g1 between 0.01 and 3.0 V and (d) rate capability at various current rates between 100 and 1000 mA g1.

samples provide enough free space to accommodate the local volume change upon charge/discharge cycling, and the hierarchical architectures consisting zero dimensional ultrafine nanoparticles and two dimensional ultrathin nanosheets are beneficial for preventing the aggregation and keeping the integrity of the active materials. All those are beneficial for improving the cycling performance. Secondly, the hybrid Co3O4 NPs@TiO2(B) NSs structures can provide extra active sites for the storage of lithium ions, which increases the specific capacity. Thirdly, the ultra-small dimension of the TiO2(B) nanosheets and Co3O4 nanoparticles in the composites offers rapid conductive path in the electrode and reduce effective diffusion distance for lithium ions and electrons, respectively, which are essential to decreases the inner resistance of LIBs and stabilize the electronic and ionic conductivity, therefore leading to an excellent rate capability.

4. Conclusions In conclusion, we report the synthesis of novel Co3O4 NPs@TiO2(B) NSs nanocomposite, in which ultrafine Co3O4 nanoparticles are uniformly dispersed on TiO2(B) nanosheets with atomic thickness. When tested as anode material for LIBs, the Co3O4 NPs@TiO2(B) NSs sample shows better lithium storage properties compared to the pure Co3O4 nanoparticles and TiO2(B) nanosheets. The Co3O4 NPs@TiO2(B) NSs sample with optimized composition shows a reversible capacity of ~677.3 mAhg1 after 80 cycles at a current density of 100 mAg1. A capacity of 386.2 mAhg1 can still be achieved at the rate of 1000 mAg1. The synergistic effect of ultrafine Co3O4 nanoparticles and atomic-thickness TiO2(B) nanosheets in the composites is responsible for the lithium storage enhancement. The present facile route to synthesize

nanocomposites with ultra-small dimension can also be extended to other materials system for the diverse applications in energy fields. Acknowledgments This work was financially supported by the National 973 Project of China (2015CB654902) and Chinese National Natural Science Foundation (11374174, 51390471, 51401114 and 51571054). This work made use of the resources of the National Center for Electron Microscopy in Beijing and Tsinghua National Laboratory for Information Science and Technology. Appendix A. Supplementary data Supplementary data related to this article can be found at http:// dx.doi.org/10.1016/j.jpowsour.2017.07.076. References [1] X. Wu, S. Zhang, L. Wang, Z. Du, H. Fang, Y. Ling, Z. Huang, J. Mater. Chem. 22 (2012) 11151e11158. [2] J.-Y. Liao, V. Chabot, M. Gu, C. Wang, X. Xiao, Z. Chen, Nano Energy 9 (2014) 383e391. [3] B.L. Ellis, P. Knauth, T. Djenizian, Adv. Mater. 26 (2014) 3368e3397. [4] H. Fang, S. Zhang, W. Liu, Z. Du, X. Wu, Y. Xing, Electrochimica Acta 108 (2013) 651e659. [5] H. Huang, W. Zhu, X. Tao, Y. Xia, Z. Yu, J. Fang, Y. Gan, W. Zhang, ACS Appl. Mater. Interfaces 4 (2012) 5974e5980. [6] J.-Y. Liao, A. Manthiram, Adv. Energy Mater. 4 (2014), 1400403-n/a. [7] X. Su, Q. Wu, X. Zhan, J. Wu, S. Wei, Z. Guo, J. Mater. Sci. 47 (2012) 2519e2534. [8] H. Sun, Y. Liu, Y. Yu, M. Ahmad, D. Nan, J. Zhu, Electrochimica Acta 118 (2014) 1e9. [9] Z. Wang, X.W. Lou, Adv. Mater. 24 (2012) 4124e4129. [10] L. Zhuo, Y. Wu, J. Ming, L. Wang, Y. Yu, X. Zhang, F. Zhao, J. Mater. Chem. A 1

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