TiO2 nanorods confined in porous V2O5 nanobelts and interconnected carbon channels for sodium ion batteries

Applied Surface Science 473 (2019) 873–884

Contents lists available at ScienceDirect

Applied Surface Science journal homepage: www.elsevier.com/locate/apsusc

Full Length Article

TiO2 nanorods confined in porous V2O5 nanobelts and interconnected carbon channels for sodium ion batteries Fusheng Liua,b, Xinxin Suna,b, Yuting Liua,b, Xiuyan Songa,b, Jun Gaoc, Guohui Qina,b,

T

⁎

a

State Key Laboratory Base for Eco-Chemical Engineering in College of Chemical Engineering, College of Chemical Engineering, Qingdao University of Science and Technology, Qingdao 266042, Shandong, China b Shandong Synergetic Innovation Center of Ecological and Chemical Engineering, Qingdao 266042, China c College of Chemical and Environmental Engineering, Shandong University of Science and Technology, Qingdao 266590, China

A R T I C LE I N FO

A B S T R A C T

Keywords: TiO2 rods Pseudocapacity mechanism V2O5 nanobelts Volume mitigation Na-ion battery

The limited ionic diffusion and electron transport pathway confine the rate performance and stability of TiO2 based anode material for sodium ion batteries. As regards, a typical TiO2 nanorods with a heterojunction with V2O5 nanobelts encapsulated into continuous carbon frame work were designed to promote the transport of electrons and Na+ ions as well as mitigating the volume expansion during the sodiation/desodiation process. One-dimensional (1D) TiO2 hierarchical coupling with ultrathin V2O5 nanobelts confined in continuous carbon framework is designed by combining with hydro-thermal method and aqueous solution growth mechanism at room-temperature. Such structure advantages including three-dimensional (3D) building blocks, large surface area, and the optimum porous hybrid architecture afford rich accessible sites and multiple pathways for the transfer of charge carriers, which shorten the ion transport kinetics and facilitate the mass transfer as current expands to 200 times (from 0.1 Ag1 to 20 Ag−1) as well as the buffering volume expansion advantages during repeated cycles. Simultaneously, the composition synergistic effects including the oxygen vacancies due to the substitution of Ti by V introduce the enhancements in both electron conductivity and the preferred lower Na ion insertion/extraction energy barrier. In addition, these composite is benefited from superior capacity contribution from V2O5, the abundant active reaction sites because of the exposure of the large surfaces which are beneficial from the rich pores. Such three-dimensional TiO2-V2O5 composite can accelerate the electron transportation via the highly orientated interconnected structure as well as facilitating the ion diffusion because of the larger expanded interlayer. The approach supplies a promising material model for preferred orientation active planes and higher Na+ transport kinetics. Such composite with unique structural features presented remarkable highrate performance when tested as a anode material for sodium-ion batteries (366 mA h g−1 at ∼20 A g−1), and showed very stable sodium-storage performance (a capacity retention nearly to 100% at 5 A g−1). When employed as anode for sodium-ion hybrid capacitors (SIHCs), it delivered a maximum power density of 6.84 kW kg−1 (with 114.07 Wh kg−1 energy density) and a maximum energy density of 244.15 Wh kg−1 (with 152.59 W kg−1 power density). This work with highlights in composition synergistic effects and typical structure design supplies a promising approach to enhance the electrochemical property of sodium ion batteries, which also can be applied to different metal oxides for energy storage devices.

1. Introduction Sodium-ion batteries (SIBs) serving as one of the most alternative devices for Lithium ion batteries (LIBs), have been employed extensively in previous research, searching for the proper candidates for accommodating sodium ion still remains challenging due to the larger deionization radius and heavier weight of sodium ion, especially for

probing excellent anode material. The large volume expansion problem causes the poor cycling and rate performance. The fabrication of the combination of multi-dimension to hierarchical porous structure enables prosperous electron transport pathway, high electric conductivity, and effective stress relaxation, which could meet the demands of fabulous electrochemical properties. The structure superior as well as the composition synergistic effect take advantage of short Na+-diffusion

⁎ Corresponding author at: State Key Laboratory Base for Eco-Chemical Engineering in College of Chemical Engineering, College of Chemical Engineering, Qingdao University of Science and Technology, Qingdao 266042, Shandong, China. E-mail address: [email protected] (G. Qin).

https://doi.org/10.1016/j.apsusc.2018.12.026 Received 8 October 2018; Received in revised form 2 December 2018; Accepted 3 December 2018 Available online 22 December 2018 0169-4332/ © 2018 Elsevier B.V. All rights reserved.

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

derived method combining electro-polymerization method, which is free of harsh reaction conditions. The final materials were investigated as anode material for SIBs, which demonstrated high performance energy storage capability. It is noted that such TiO2 nanorods trapped in the walls of V2O5 enclosed by carbon carriers can well address several important challenging points related to SIBs: (a) such nanostructure can mitigate a larger space for allowing the large volumetric expansion of TiO2-V2O5-C during sodiation and preserving their structural integrity (Fig. 1). (b) The hierarchical nanostructure supplies multi-layer pathway for electron transportation and shortens the ion diffusion pathway. (c) The oxygen vacancies formed by the introduce of V increase the electrochemical conductivity and facilitate the Na ion diffusion due to the lowered energy barrier during (dis) charge process. Motivated by this intriguing technique, such TiO2-V2O5-C architectures deliver a maximum reversible specific capacity of 401.3 mA h g−1 even after 200 repetitive cycles at 5 A g−1 and a large reversible capacity of 366.5 mA h g−1 at a large current density of 10 A g−1.

distance, large electrode/electrolyte contacting area, as well as abundant porosity, the high capacitance advantages, efficient electron and ion transportation along the longitudinal orientation [1,2]. The 3D mesoporous nanobelts facilitate a large surface area, huge pore volume, favorable mesoporous channels and well-crystalline thin mesopore walls, which are capable of supplying large accessible voids for sodium ion insertion and extraction as well as buffering volume deformability. Anatase TiO2 Titanium dioxide (TiO2) is a potential anode material due to its exceptional stability, environmentally benign and nontoxicity, low cost, and abundance. The preferable TiO2 based ordered 3D nanoarrays are able to accommodate Na-ion into their structures, serving as a promising anode material with adequately high reversible capacity and good cycleability [3,4]. The combined the nanoarchitectural design with conductive composite engineering as well as the size engineering has been proved to be effective for improving the fatal defects of TiO2 as anode material of SIBs, the poor conductivity and the limited ion diffusion capability and the low capacity [5,6]. In previous work, 1 D nanosized anatase TiO2 can achieve higher capacity utilization and rate capability when evaluated as anode material of SIBs [7]. However, the 1D TiO2 mesoporous, multi-channel rod-like nanomaterials tend to agglomerate during (de) sodiation process with the formed inferior conductive network causing the loss of particle connectivity. As regards, the mentioned hybrid model from 1D to 3D nanostructure by incorporation of composition contribution manifests the alleviation of aggregation tendency during the charge and discharge process and enhancing the rate capability of SIBs. The combination of 1D and 3D nanosize engineering can bring the positive effect to prevent the large irreversible capacity loss, and poor cycling stability and rate capability and low initial coulombic efficiency due to their huge volume change induced electrode pulverization during insertion and extraction and therefore cause cracking, fracture, and particle dissociation problem (Fig. 1) [8–10]. Interwaving to 1D crystals into 3D nanostructures with open framework for Na-ion intercalation and deintercalation exhibited fast sodium-ion conductivity and enhanced accomodation to the larger sodium ion and its higher ionization potential [11,12]. V2O5 with larger d-spacing as anodes/cathodes in Na-ion batteries presented an outstanding capability. Therefore, the hierarchical heterojunction of 1D TiO2 with 3D V2O5 nanobelts can be expected to be a promising anode material due to hybrid structure advantage and composition synergistic effects, in which V2O5 can afford pseudocapacitive capacity to such TiO2-V2O5 as well as the benefits including the creation of oxygen vacancies and robust structure stability to the assembly 3 D orientation structure [13,14]. In order to further fill with the conductivity fracture of such TiO2V2O5 composite and keep the structure integrity, such 3D nanostructure gets binding into the continuous conductive carbon carriers, in which their electrochemical kinetics was accelerated and the power density was improved, and adjustable carbon matrix can enable fast electron transfer and play as a buffer mediation to accommodate the strain with volume change [15–17]. While the continuous carbon framework facilitate strain relaxation and electron transport with suppressed mechanical fractures partly improving capacity and stability of SIBs. Such 3 D mesoporous nanostructure forms anti-aggregated robust architecture, maintaining the structure stability during cycling. 1D nanoscale subunits show extra structural advantages such as high cross-linking density, which predominates high open porosity and large accessible surface area, supplying huge electrolyte/electrode contacting area and ultrafast electron transport pathway. Simultaneously, the well-defined 3D V2O5 nanobelts and large mesoporosity as well as thin uniform mesopore walls are capable of supplying large accessible voids for sodium ion intercalation and extraction and as well as alleviating volume expansion. Therefore, such 3D mesoporous TiO2-V2O5-C nanostructure is expected to present superior rate performance and high stability. Herein, 3D hierarchical TiO2-V2O5-C composed by 1D meporous TiO2 intertwining of 3D V2O5 ordered nanobelts encapsulated into continuous carbon framework were successfully fabricated using sol-gel

2. Experimental section 2.1. Preparation of TiO2-V2O5-C nanostructure TiO2 nanorods were firstly prepared. In a typical synthesis, 30 ml of ethanol solution containing 6 ml of titanium (IV) isopropoxide under vigorous stirring at 10–80 °C was aged for 3 h. The precipitate was collected by centrifuge and repeatedly washed with ethanol and deionized water, and then air-dried at 80 °C. The prepared paste (1.6 g) and amphiphilic triblock copolymer Pluronic F127 (100 mg) were added into 100 ml water at a temperature of 100 °C and kept stirred until getting a viscous gel. And then the gel was dispersed in mixed solvents of ethanol and glycerol (3:1), the ammonium metavanadate (NH4VO3, 0.0690 g) powder was added into the above solution. In the following procedure, the mixture was heated at 120 °C in oil bath and refluxed for 6 h, and the assembly process of monomicelles occurred surrounding glycerol in spatially confined 1D direction with any solid interfaces free. The mesoporous heterojunction of TiO2-V2O5 was synthesized based on the removal of the amphiphilic block. Afterwards, the mixture was dried at 120 °C for 8 h, and then annealed at 350 °C with a detailed heating rate of 2 °C min−1 and placed in air ambient for 4 h. Subsequently, the carbon layer was coated on the surface of TiO2V2O5 composite via the electro-polymerization approach [18]. To be specific, the electrochemical-polymerization process was performed at a constant potential of 1 V for 600 s in a three-electrode cell at room temperature. The synthesized TiO2-V2O5 nanostructure was dispersed into ethanol and the solution with a titanium substrate serving as the working electrode. The PANI monomer (0.1 mol L−1) in sulfuric acid (0.1 mol L−1) was introduced as the electrolyte; and Hg/HgSO4 and Pt mesh were designated as the reference and counter electrodes, respectively. The products were washed with distilled water and acetone several times, before being calcined at 700 °C for 5 h in a tubular furnace under a N2 ambient to obtain the continuous carbon shell on the surface of TiO2-V2O5 composite, restricting the excess growth of TiO2V2O5 (Scheme 1). 2.2. Characterizations The X-ray diffraction (XRD) patterns of TiO2-V2O5-C samples were finished via a PANalytical X-pert diffractometer (PANalytical, Netherlands) with a Cu Ka radiation operated at 40 kV and 30 mA. Raman spectrum was tested by a Renishaw in ViaRaman microscope at room temperature with the 532 nm line of an Ar ion laser as an excitation source. X-ray photoelectron spectroscopy (XPS, VGMicro Tech) was used to record the functional group. Morphology and structure of the samples were tested via field-emission scanning electron microscopy (FE-SEM, S4800, Thermo Fisher) and high-resolution transmission electron microscopy (HR-TEM, Tecnai G2F20, Philips). 874

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

Fig. 1. (a) XRD patterns of the TiO2 samples; (b) Raman spectra of the TiO2 composites; (c) TGA curves of such TiO2-V2O5-C composites under Oxygen flow; (d) XPS survey spectrum of TiO2-V2O5-C at 700 °C in NH3 atmosphere for 1 h.

an argon-filled glove box. A Celgard 2400 microporous polypropylene membrane was used as the separator and sodium foil served as the counter electrode. The thickness of the prepared electrode film is 5 × 10−3 mm and the loading of the active material is 1.1 mg cm−2. The nonaqueous electrolyte in sodium ion battery was 1 M NaClO4 dissolved in a mixture (1:1 in wt%) of ethylenecarbonate (EC)- dimethylcarbonate (DMC). The coin cells were firstly activated at a current density of 50 mA g−1 for first cycle, and then cycled under different current densities within the voltage range of 0.01–2.2 V via a LAND-CT2001A battery test system (Jinnuo Wuhan Corp., China). And then the cells were cycled under different current densities within the voltage range of 0.01–2.2 V. Electrochemical impedance spectroscopy (EIS) measurements were carried out on an electrochemical workstation (CHI 660 D, CHI Company) at a frequency range of 0.1 Hz to 100 kHz.

3. Results and discussion

Scheme 1. A schematic illustration of the procedure for the preparation of three dimensional TiO2-V2O5-C nanostructure and their application as anode materials for Na batteries.

3.1. Structure and morphology Fig. 1a shows the X-ray diffraction pattern of as-synthesized TiO2 based samples. For all TiO2 materials, no diffraction peaks of C or V2O5 could be detected due to their low content or amorphous state, and all peaks belonged to the face-centered TiO2 phase (PDF Card No.121273). All the samples showed a characteristic signal of anatase phase, and the peaks at 2θ values of 25.8°, 38.3°,48.5.°, 53.9° and 55.5° were attributed to the (1 0 1), (0 0 4), (2 0 0), (1 0 5),(2 1 1), (2 0 4), and (1 1 6) crystal planes of anatase. The cell constants obtained from the Rietveld refinement are a = 3.79529 Å, b = 10.3045 Å, c = 4.6859 Å, a

The evaluation of electrochemical performance was performed via coin-type SIB cells (2025) assembled in an argon-filled glove box. As regards to the electrode material preparation, a mixture of active material, carbon black, and polyvinylidene fluoride (PVDF) binder with a weight ratio of 80:10:10 was added into Nmethylpyrrolidone (NMP) solution, in the following process, the mixtured slurry was then uniformly pasted into a Cu foil current collector. And then was dried at 120 °C for 24 h under vacuum before being assembled into coin cells in 875

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

unit cell volume of 286.95 Å3, and average particle diameter of TiO2 (1 0 1) is 3.1100 Å, which is consistent with the ICDD card of No. 211273. V2O5 in these TiO2 based structure was mainly distributed near surface, however, the degree of segregation was too small to be tested by XRD investigation, owing to the relatively low concentration of the component (V/Ti = 0.03). However, V is confirmed in pure pattern (Fig. S1) and testified by XPS and EDS file (Fig. 4). To confirm the presence of carbon framework, Raman spectra of all TiO2 samples were investigated, in which there are two prominent peaks of D and G bands, typical characteristics of the chemically derived carbon (Fig. 1b). The Raman spectrum of pristine and characteristic peaks of carbonaceous materials at ∼1592 cm−1 (G-band, deriving from the in-plane stretch vibration of carbon atoms) and 1356 cm−1 (D-band, ascribed to the breathing-vibration of carbon atoms), confirm the existence of carbon in the prepared materials. According to Tuinstra-Koenig (TK) equation, the intensity ratio of D band to G band (ID/IG) is usually used to estimate the disordering degree of carbon [19,20]. The relative high ID/IG value (1.12) for our TiO2-V2O5-C suggests a more disordered phase compared with that of TiO2-V2O5 (0.81), which may be caused by the defects in carbon framework. It is noted that a downshift of the G band in the Raman spectra of TiO2-V2O5-C was observed caused by the inhomogeneous CeC and CeO (C]O) bond distances [21–23]. The Raman intensity and wavenumber of both the D and G bands change after carbon incorporated with TiO2/V2O5, indicating that the porous carbon framework has a good contact with TiO2-V2O5, which is consistent with the previous results. As depicted in TG curve (Fig. 1c), a weight loss of around 4.1% was observed from room temperature to 1000 °C in air with a corresponding optimized temperature of 652 °C for thermal treatment of the precursor to ensure its complete decomposition, which is attached with the loss of water and the combustion of carbon. To identify the surface composition and chemical states of the anode material, X-ray photoelectron spectroscopy (XPS) spectra was investigated (Figs. 1d and 2). Four elements including C, O, and Ti with a weak signal of V were detected. The binding energy of V mainly located at 517.5 eV, 521.32 eV and 531.6 eV are attributed to V2p3, V2p1 and occupied sites by obvious O signal. EPR file illustrates the surface Oxygen vacancies elevation of TiO2-V2O5-C after V2O5 and C incorporation (Fig. S2). The peak located at g = 2.02 is derived from O− signal mainly formed via oxygen vacancies. The gradually intensified signal of O− species from pure TiO2 to TiO2-V2O5-C implied the presence of surface defect sites which are good for the increment of contact area between active material and electrolyte. Such surface defects can superiorly improve the Na+ storage performance. Fig. 2a shows the fitted high-resolution spectra of the C1s, which display three peaks. The peak at around 284.5 eV can be assigned to CeC in the carbon. The peaks at 285.9 eV can be attributed to CeO, while the peak 288.2 eV is assigned to C]O bonds, respectively [24]. The intensities of CeO and C]O peak are weaker than that of CeC implies that most oxygen-containing functional groups are removed after solvothermal treatment, which are well consistent with the XRD and Raman results. The main peak at 284.5 eV is derived from to sp2hybridized carbon, while the weak peak at 284.8 eV is derived from sp3hybridized carbon. Two weak peaks at 286.1 and 288.5 eV can be assigned to oxidized carbons. The two symmetric peaks at 530.5, 532.8 and 534.1 eV in the spectrum of O1s peak can be attributed to VeO, the vacancies formed due to the V occupy and the TieO bond, respectively (Fig. 2b). The peaks (Fig. 2c) at about 459.2 and 464.79 eV are in accord with the Ti 2p3/2 and Ti2p1/2 spin–orbit. These results imply that TiO2 has not been reduced during the solvothermal process. Obviously, both carbon coating and V doping can improve the diffusion of Na ions. Introducing oxygen vacancies by Vanadium and coating carbon layer is validated to considerably enhance the surface pseudocapacitive process, assisting to achieve high capacities of TiO2 based composites at high currents [25].

It t is noted that the rich oxygen vacancies are beneficial to the increments of the electron conductivity and the diffusion ability of sodium ion batteries. The mesoporous texture of the TiO2-V2O5-C nanorods can be confirmed by the N2 adsorption-desorption isotherm. As shown in Fig. S3, the isotherm profile of the sample is ascribed to a type IV curve with a well-defined hysteresis loop at the relative pressure of 0.8–1.0 (Fig. S3a), which showed the existence of a large number of mesopores in the complicated TiO2 sample with a narrow size distribution of 5–20 nm (Fig. S3b), which complies with the result of the Barrett-Joyner-Halenda (BJH) pore-size distribution pattern (inset in Fig. 4d). The surface area of the TiO2-V2O5-C has been found to be 120.6 m2 g−1 and pore volume is 0.41 cm3 g−1, while that of TiO2-V2O5 is 87.6 m2 g−1 and 0.24 cm3 g−1, and that of pure TiO2 is 41.29 m2 g−1 and 0.14 cm3 g−1. Such mesoporous nature of the TiO2-V2O5-C composite and suitable pore distribution are conducive to accelerating mass diffusion of the electrolyte due to the abatement of diffusion resistance and facilitate for the electrolyte to penetrate efficiently into the pores and diffuse much easier to active sites due to less resistance, besides, such stable nanostructure is beneficial to buffer huge volume change during the Na+ insertion/extraction processes. The detailed particle morphology is studied via TEM image. As shown in Fig. 3a, pure TiO2 nanoparticles are consisted of densely packed different size of rod-like structure. The TiO2 nanocrystals exhibit clear lattice fringes, indicating single crystallinity of TiO2. Fig. 3b reveals the HTEM images of pure TiO2 reveal clear lattice fringes with d-spacing of around 0.178, which correspond to the (1 0 1) lattice planes of TiO2, indicating a high degree of crystallinity. The SAED pattern (Fig. 3c) also confirms the well corresponding to (0 0 4) spacing of the anatase TiO2 phase, which corroborates the pure phase of TiO2. Fig. 3d and e shows prepresentative TEM images of the TiO2-V2O5-C nanostructure, further corroborate the porous structure with nanorods pointing out from the surface, certifying its close incorporation with V2O5 nanobelt enclosed by continuous carbon framework. The hierarchical nanorods bundles with a longitudinal length of 20 nm and a transverse length of approximately 100 nm were assembled into welldefined nanorods in a uniform and ordered arrangement. Such small TiO2 nanorods bundle with other nanobelts from neighboring urchin V2O5 belts and continuous interconnected network carbon network. As shown in Figs. S4, S5, such nanobelts were comprised by numerous of porous interweaving nanobelts subunits with a thin wall thickness feature with a high aspect ratio and exhibit a mesoporous structure. Such an arrangement can ensure a high uniformity of the material and provide enough channels for the diffusion of the electrolyte, which may help to enhance the utilization of the active area. These TiO2-V2O5 heterostructure clearly possesses well-defined ordered nanostructure after incorporated with continuous carbon framework. The corresponding fast-Fourier transform (FFT) pattern (Fig. 3f) of the lattice-fringe pattern demonstrates well-resolved individual reflections, which indicates that the composite nanorods are single crystalline. It is clearly shown that TiO2-V2O5-C single crystals are highly crystalline with dominant highly active {1 0 1} facets, which is a typical structure for anatase based on the Wulff constructions. This preferred orientation is especially meaningful, because {1 0 1} facets possess superior Na storage capacity, as described in the introduction section. Some parallel channels (1 0 1), (0 0 4) were formed among samples, which have the possibility to increase the surface area of the materials and to shorten the transmission path for the ion and electron. Combined with the mesoporous nature, the preferred nanostucture can provide a short electron transfer path, sodium-ion transport length and an efficient electron transport channel when the materials are used as anode material for NIBs. The bright-field STEM images of TiO2-V2O5-C and its corresponding EDS and EDX results (Fig. 4) can further confirm the uniformly interwave of TiO2 nanorods into V2O5 nanobelts and carbon framework by such an solvothermal–calculation method, which fits well with the analysis results of XPS, XRD, Raman 876

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

Fig. 2. XPS spectrum of TiO2-V2O5-C at 700 °C inNH3 atmosphere for 1 h: (a) C1s narrow scan (b) O1s narrow scan (c) Ti2p narrow scan (d) V2p narrow scan.

Fig. 3. (a) TEM image, (b) HRTEM image, (c) elected area electron diffraction (SAED) pattern for pristine TiO2, (d) and (e) TEM image at different magnifications, (f) FFT pattern for TiO2-V2O5-C. 877

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

Fig. 4. (a) HAADF-STEM image and corresponding EDS mapping images of an individual TiO2-V2O5-C indicating the homogeneous distribution of carbon, vanadium, titanium, and oxygen, and (b) EDX HAADF image and corresponding EDX mapping.

TiO2 (Fig. S6a) are only 85% and 76%, respectively. Compared with the whole capacity contribution, TiO2, V2O5 and C contribute 53%, 23% and 24%, respectively, to the whole capacity. Hence, C and V2O5 play equally in enhancing the capacity to such composite while C contributes greater in improving the stability, which is further confirmed in the following investigations as regards to different contents of C and V2O5. The relative superior stability for TiO2-V2O5-C is derived from the characteristics of hybrid structure, the double insurance from flexible V2O5 and carbon layer keep the mesoporous TiO2 -V2O5-C from collapse during the sodium insertion/extraction and the enhanced conductivity supported by the carbon-base anode. V2O5 offers the specific capacities to such TiO2 based electrode material and leads to the creation of rich oxygen vacancies which are beneficial for the improvements of electron conductivity and ion diffusion. It is noted that the electrochemical performance of TiO2-V2O5-C is related to the content of V2O5 and C in the composite (Fig. 6a). V2O5 supplies pseudocapacitive capacity to such TiO2 based composite. As displayed in Fig. S7, pure V2O5 indeed contributes certain capacity to TiO2-V2O5-C in the voltage window of 0.01–2.2 V. The influence of different contents of V2O5 to the electrochemical performance of TiO2V2O5-C was studied. The charge-discharge capacity of such composite elevated from the initial 390.2 and 389.6 mA hg−1 (coulomb efficiency 99.8%) to 475.8 and 475.1 mA hg−1 (coulomb efficiency 99.8%) as the stoichiometric of V:Ti increased from 0.015:1: to 0.03:1. However, the charge and discharge capacity of such anode material decreased to 446.3 and 435.1 mA g−1 (coulomb efficiency 94%) when further increases such ration to 0.06:1. As shown in Fig. 6b, TiO2-V2O5-C (Ti:V = 1:0.015) achieved the stable discharge at second cycle, while TiO2-V2O5-C (Ti:V = 1:0.03) took 3 repeated cycle to get reversible capacity, and TiO2-V2O5-C (Ti:V = 1:0.06) didn’t achieve a stable discharge capacity until 8th repeated cycle. In addition, TiO2-V2O5-C (Ti:V = 1:0.06) sample showed a polarization tendency at a current as low as 0.5 A g−1, which explained the irreversible capacity caused by higher vanadium content. V2O5 nanobelts offers a cushion effect to cope with the partial stress of the volume change for TiO2-V2O5-C anode material and maintains the sufficient contact with electrolyte, what’s important, it supply a specific capacity to such composite material. When the content of V2O5 is too low, resulting in the low initial discharge capacity and inferior rate performance. On the other hand, when the content of V2O5 is high, some irreversible reactions can occur. Also excess V2O5 causes low electronic conductivity due to the thick coating

profiles. 3.2. Electrochemical performance The typical structure design encouraged us to investigate the detailed electrochemical performance of TiO2-V2O5-C serving as anode material in SIBs. The electrochemical behavior of the as-prepared TiO2 electrodes was evaluated using a CR2025-type coin cell. Fig. 5(a) shows the representative cyclic voltammetry (CV) curves of the TiO2-V2O5-C electrode in a range of 0.01–2.2 V at a scan rate of 1 mV s−1. The cathodic and anodic peaks can be ascribed to the Na ions insertion and extraction processes at the cathodes, respectively. As shown in Fig. 5a, the current beginning at 0.8 V has related to irreversible formation of the solid electrolyte interface (SEI) in the first cathodic scan, exclusion to the reversible sodiation of the material. In the following cycles, the stable oxidation and reduction peaks with better repeatability can be clearly observed [26]. The CV curves with distinct redox peaks are consistent with those reported in the literature for TiO2 materials [27]. The stable redox peaks upon cycling indicate the stability of the hybrid structure upon charge and discharge. Fig. 5b presents the typical charge/discharge profiles of TiO2/V2O5/ C for SIBs. The initial discharge/charge capacities for TiO2/V2O5/C were 555.9 and 472.8 mA h g−1 at 0.1 A g−1 with a Coulombic efficiency of 85% mainly ascribed to the decomposition of electrolyte and the irreversible formation of the solid/electrolyte interphase on the surface of the electrode. The capacity delivered by TiO2-V2O5-C are higher than those for TiO2-V2O5 (428.5 and 360.3 mA h g−1), and is far beyond of TiO2 (408.5 and 252.3 mA h g−1). They displayed the similar voltage plateau the CV measurement. The prominent performance of the TiO2-V2O5-C are attributed to their structural and compositional advantages. The building of TiO2-V2O5-C with TiO2 embedded in V2O5 nanobelts and carbon framework enables a short fast Na+ diffusion distance as well as effective contact between the active material and the electrolyte which is beneficial for the rapid charge-transfer reaction. The channel architectures constituted by the mesoporous and interior voids supply good structural integrity because of sufficient space for mitigating excellent large volume variations upon repeated cycling. Fig. 5c presents the 1st, 2nd, 10th, 50th, and 80th and 100th charge–discharge voltage profiles of the TiO2 anode between 0.01 and 2.2 V at a current density of 0.1 A g−1. 93% of capacity for TiO2-V2O5-C is retained after 100 cycles, while that for TiO2-V2O5 (Fig. S6b) and 878

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

Fig. 5. (a) Cyclic voltammogram curves with several cycles between 0.01 V and 2.2 V (vs. Na+/Na) at a scan rate of 1 mV s−1; (b) charge/discharge curves of TiO2 samples at 0.1 Ag−1; (c) the charge and discharge curves of TiO2-V2O5-C; (d) the cycle performances of the TiO2 samples at various current rates.

further increases to 6%. As well known that, the addition of carbon at a relative low content can bring the enhancement of conductivity as well as the improvement of cycle stability, however, the relative high content doping of carbon causes detrimental effect to such composite. In this case, too excess carbon can block the well contact of active material with electrolyte, in which the abatement of Na+ diffusion ability can generate. Therefore, the optimized carbon doping content is 4.1% for such TiO2-V2O5-C.

of V2O5 with the interface of electrode and electrolyte. Therefore, TiO2V2O5-C (Ti/V = 1:0.03) displayed higher reversible capacity and longer cycling stability than other ratio. The influence of different contents of carbon to the electrochemical performance of such TiO2-V2O5-C was also investigated (Fig. S8). The discharge capacity of TiO2-V2O5-C increases from the initial 427 mA g−1 to 472 mA g−1 as the increments of carbon content from 0 to 4.1% at 0.1 A g−1, however, it drops to 449 mA g−1 when carbon

Fig. 6. (a) charge-discharge curves of TiO2-V2O5-C at 0.1 Ag−1 with different vanadium contention (b) the rate performance of TiO2-V2O5-C from 0.1 to 0.5 Ag−1. 879

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

8.5 mol cm−3, while the loading of the active material (m) is 1.1 mg cm−2, the thickness of the prepared electrode film (T) is 5 × 10−3 mm and the pore volume (Vg) is 0.41 cm3 g−1, C can be calculated to be 7.69 × 10−3 mol cm−3 based on Eq. (2)

Fig. 5d shows the rate and cycle performance of charge and discharge curves at various current densities. The excellent electrochemical capacities of 472, 459, 444, 433, 427, 401 and 366 mA h g−1 for TiO2-V2O5-C at various current densities of 0.1, 0.2, 0.5, 1, 2, 5 A g−1 and 10 A g−1, respectively. Even at a high current density of 10 A g−1, it can maintain a reversible capacity of 366 mA h g−1, obviously, when the current density back to 0.1 A g−1 after 140 cycles, the capacity recovered to 472 mA h g−1, demonstrating an excellent rate capability. Comparatively, TiO2-V2O5 showed a lower rate performance, just delivered 342, 329, 313, 288, 257, and 220 mA hg−1 at 0.1, 0.2, 0.5, 1, 2, and 5 A g−1. After 100 repeated cycles, there was an obvious degeneration at discharge capacity, but only 190 mA hg−1 was delivered for TiO2-V2O5. While pure TiO2 delivered the worst discharge capacity, it exhibited a high specific capacity of 140 mA h g−1 at 2 Ag−1 and the decline tendency was the most serious. As confirmed by the CV and galvanic charge and discharge files, urchin V2O5 and carbon wall have the large surface area and pore volume. The large surface area can increase the interface of the active material and electrolyte, thus facilitating the Na+ ion diffusion. The mesoporous structure is beneficial to ion diffusion and facilitates easy permeability of the electrolyte. Large-volume change upon charging and discharging can also be alleviated by the porosity. Thus, the best cycling and capacity-retention performances of TiO2-V2O5-C among all of the samples are accredited to its peculiar nanostructure with the most excellent rate performance. The porous TiO2-V2O5-C nanorods exhibit an excellent rate capability as compared to the TiO2-V2O5 and TiO2 samples, however, the differences in conductivity and ion diffusion between them should be taken into account. Electrochemical impedance spectroscopy (EIS) measurements were performed to investigate the conductivity of the cathodes composed of TiO2 based electrodes. As shown in Fig. 7a, the TiO2-V2O5-C electrode possesses a much lower resistance than that of the TiO2-V2O5 and TiO2 electrodes (ca. 49.1 vs. 80.6 and 140.6 Ω, respectively), implying that the incorporation of a mesoporous structure with V2O5 and carbon matrix can significantly enhance the electrochemical conductivity and ion transmission kinetics. The synergistic effect of the successful integration of the TiO2 structure and interwaved porous V2O5 and carbon matrix is beneficial to the structure stability during charge and discharge progresses. The sodium ion diffusion coefficient in the bulk can be calculated according to the following equation [28]:

D= R2T 2/2A2n4F 4C2σW2

C=

Co MVg (2)

T −1

and F is the Faraday constant (96,486 C mol ). The value of σw is the slope of the lines between Z′ and ω−1/2 (Fig. 7b). Clearly, the DNa for TiO2-V2O5-C is almost 6 times of that to TiO2 and 2 times higher than that for TiO2-V2O5. This is likely associated with the 3D conductive matrix, as discussed in the TEM part and the V2O5 and carbon network fabricate for ionic and electronic diffusion. For the sodium ion intercalation and dissociation reaction, the electronic and Na ion diffusion are equally important. Ideal electrode materials are both good electronic and ionic conductors. The electrons and Na ions must reach or leave the reaction point simultaneously. However, for the pure TiO2, it not only suffers from extremely low intrinsic electronic conductivity due to lacking of effective conductive medium on its surface. Thus the larger particle size for pure TiO2 may decrease the electronic conductivity in the bulks and result in the higher charge-transfer impedance Rct. Furthermore, its less favored mesoporous framework leads to lower diffusion coefficient DNa, too. It also should be noted that TiO2V2O5-C shows the largest Rct and the largest DNa, these are attributable to the fact that the TiO2-V2O5-C composite is more amenable to the diffusion of Na+ ions than the pure TiO2 and TiO2-V2O5 [26,29], which agrees well with the aforementioned discussions. In particular, the lower electron transfer resistance can effectively accelerate the transfer of electrons and ions in the hetero-structured samples, positively enhances the reactive activities of these samples, which is beneficial for extending the sunlight utilization. Therefore, the synergistic function of TiO2, V2O5 and carbon together with the particular mesoposous nanostructure introduce a promising candidate for energy store. It is noted that EIS file for pure TiO2 was electron transferred controlled progress while they were ionic controlled progress for TiO2-V2O5 and TiO2-V2O5C, respectively, further verified the great contribution of V2O5 nanobelts and carbon network which possess rich Na+ storage sites, which are beneficial to shorten the sodium ion diffusion distance and facilitate the ion diffusion kinetics. The large surface area, which facilitates the Na+ ion diffusion and electron transport, endowing TiO2-V2O5-C the extraordinary electrochemical performances. To disclose the more advantaged features of TiO2-V2O5-C during the cycle life, electrochemical impedance spectroscopy (EIS) measurements of the two electrodes (Fig. 8a) were conducted before and after cycling for different cycles. The slight variations of the electron transfer impedance and ion diffusion impedance confirm the structure stability of such composite. The EIS data was analyzed by

(1) 2 −1

where D is Na-ion diffusion coefficient (cm s ), R is gas constant (8.314 J mol−1 K−1), T is the absolute temperature (K), A is the electrode area (cm2), n is the number of electrons involved in the redox process (1 in our case), C is the ion concentration inserted into the electrode material, in this case, the concentration of electrolyte (Co) is

Fig. 7. (a) EIS of TiO2 based electrode, (b) profiles of real parts of the complex impedance versus ω−1/2. 880

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

Fig. 8. (a) is electrochemical impedance spectroscopy (EIS), (b) Long cycle performance of TiO2-V2O5-C at 5 Ag−1.

Fig. 9. (a) CV curves at different scan rates. (b) Log (i) versus log(v) at different voltages for TiO2-V2O5-C.

Fig. 10. CV profiles of capacitive contribution at scan rates from 0.1 to 2 mV/s (a)–(e). (f) The specific capacities generated from battery contribution and capacitive contribution at different scan rates for TiO2-V2O5-C.

fitting to an equivalent electrical circuit (Fig. S9), such fitting patterns are consistent well with the experimental EIS data. The interface between TiO2, V2O5 and the carbon layer is good to the buffer of volume swelling force and facilitates to the rearrangements during the (de)sodiation, greatly prolonging the stability of such TiO2 based anode

material. The cycle performance of such anode material at a more tough higher current density in the long term was also invested, Fig. 8b makes a comparison about the cycling performance of TiO2, TiO2-V2O5 and TiO2-V2O5-C electrodes at 10 A g−1 in the voltage range of 0.01–2.2 V 881

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

Fig. 11. (a) Galvanostatic charge/discharge curves of TiO2-V2O5-C at current densities from 25 mA g−1 to 1 A g−1. (b) Corresponding Ragone plots of SIHCs (based on the total weight of active materials) and most energy storage devices.

serious polarization tendency occurred for both TiO2 and TiO2-V2O5, however, no obvious attenuation was observed for that of TiO2-V2O5-C because of the enhanced electron conductivity and the improvement of the ion diffusion ability as well as buffering ability to the volume swelling force endowed by the interconnected channel [30–34]. The synergistic effects of the components ensure the extraordinary electrochemical performance of such composite, which is consistent well with the above-discussed results. TiO2-V2O5-C showed a long cycling performance at high current densities, which paves a way to practical application. The crystal phase information and morphology variations were recorded by XRD and SEM, TEM investigations. As shown in Fig. S10, the crystallinity of TiO2-V2O5-C was the best compared with that of TiO2V2O5 and TiO2, indicating its structure integrity even after long tough repeated cycle, which match well with the aforementioned electrochemical performances. Consistently, the cracks and pulverization with some broken particles were observed for pure TiO2 (Fig. S11a), which there were hardly any collapses for TiO2-V2O5-C were observed (Figs. S11b and S12). These results corroborate that the dual TiO2-V2O5-C strategy could buffer the volume change of TiO2 during cycling, thus ensuring superior cycling stability. In order to further investigate the high rate performance of TiO2V2O5-C, the corresponding CV scans at different scan rates from 0.1 to 2 mV s−1 were recorded (Fig. 9). The correlation ship between current (mA) and scan rate (mV/s) was listed as follows:

Table 1 Specific capacities of TiO2-based nanostructures in the literature. Samples

Current density (A g−1)

Discharge capacity (m Ah g−1)

Ref. and published year

Microsphere C@TiO2 3D Ni-TiO2 core-shell nanoarrays TiO2@C-NPs TiO2(B) nanotubes Nanocrystalline TiO2(B) Nb-Doped Rutile TiO2 TiO2/carbon Hollow spheres G-TiO2-N TiO2 nanorods TiO2 Nanocubes Anatase TiO2 Anatase TiO2 nanorods Microsphere C-TiO2 TiO2-V2O5-C

10C 5

82.7 95

[32], 2014 [35], 2015

0.8 0.4 10C 16.75 2

127.6 33 50 120 60

[36], [37], [38], [39], [40],

2 10 1.68 4.059 33 2 20

462.8 82 84 100 53 82.7 366

[41], 2013 [42], 2014 [43], 2015 [44], 2014 [43], 2014 [32], 2014 this work

2015 2013 2015 2015 2015

up to 200 cycles. Well consistent with the above discussed, TiO2-V2O5-C exhibited better cycling stability and delivered a much higher and more stable discharge capacity (401 mA h g−1 in the 200th cycle) than TiO2V2O5 (184 mA h g−1) and TiO2 (53 mA h g−1) in the long cycle. The

Scheme 2. Schematic illustration of the 3D TiO2-V2O5-C hybrid electrode with large electrode–electrolyte contact area, short Na-ion diffusion distances and fast electron transport network. 882

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

i= av b

(3)

Furthermore, this dual buffering architecture can inhibit the side reactions that resulted from the direct contact between the active materials with the electrolyte. The presence of a carbon matrix and V2O5 mesoporous nanobelt cage effectively prevent the agglomeration of TiO2 nanorods and greatly enhance the conductivity of the electrode material. It is noted that the oxygen vacancies generated by the substitution of V to Ti creat a narrowed band gap of TiO2 thus lead to an improved intrinsic conductivity, which is favorable for the electron transfer and Na+ diffusion. The lowered energy barrier of sodiation due to the existence oxygen vacancies is benefit for a more favorable Na+ insertion into the anatase structure. Simultaneously, the specific capacity of V2O5 enhances the capability of such TiO2 composite. Consequently, the smart TiO2-V2O5-C structure endows the electrode extremely Na+ storage properties.

where the a and b are adjustable parameters. It is diffusion-controlled battery reaction process in the case of b = 0.5, while it belongs to a surface capacitive response in the case of b = 1. The values of b were calculated via the relationship of log(i) and log(v). The values of b at reduction peaks (0.001 V, 0.6 V, 1.3 V) and oxidation peak (2 V) can be calculated to 0.74246, 0.8898, 1.01964, 0.9538, respectively, which are close to 1, implying high surface capacitive contribution. Eq. (4) was adjusted for the convenience to obtain the surface capacitive contribution:

i= k1 v+ k2v1/2

(4)

where k1 and k2 are coefficient of fixed potential. As demonstrated in Eq. (3), k1v and k2v1/2 are ascribe to surface capacitive and diffusion contribution, respectively. The values of k1 can be obtained by rearranging Eq. (4) to

i/v1/2 = k1v1/2 + k2

4. Conclusions In summary, a strategy is developed to encapsulate the 1D TiO2 nanorods into aligned V2O5 nanoblets bundles, and simultaneously confine TiO2-V2O5 in interconnected carbon channels. As a proof of this strategy, such possesses excellent electrical/ionic conductivity, kinetics favorable structure, and synergistic compositional effects. The electrochemical performance proves that the as-prepared TiO2-V2O5-C possesses the synergistic structure and composition effect of V2O5, TiO2 and carbon and exhibits superior electrochemical capacitance and outstanding cycle stability. On one hand, the encapsulation of TiO2 nanorods into mesoporous V2O5 nanobelts and continuous carbon network immobilizes the TiO2 rods during the charge and discharge and buffers the huge volume expansion during the charge and discharge progress. The oxygen vacancies generated bring lower energy desodiation/sodiation energy and lead to improved electrochemical conductivity and attenuated ion diffusion resistance. On the other hand, V2O5 accelerates the pseudocapacitance for sodium storage, while carbon effectively increases the electrochemical conductivity of such composite and offers rich active sites for sodium storage. Consequently, such TiO2-V2O5-C electrode achieves high-rate response up to 20 A g−1 and stable capacity retention for 200 cycles with superior capacities (366 mA g−1) due to the extraordinary redox stability and the volume deformability. When assembled to SIHCs, such devices displayed high power density and energy density which is comparable to most supercapacitors. This work builds a hierarchical nanocomposite using 1D and 3D nanocomponents for the fabrication of 3D porous composite nanostructure for electrochemical applications.

(5)

As displayed in Fig. 10, capacitive contribution of TiO2-V2O5-C is proportional to the scan rate with the value increases from the initial 59.8% at a scan rate of 0.1 mV s−1 to 90.6% (2 mV s−1) (Fig. 10a–f), which can be derived from the fine tailoring of each composition and porous structure of TiO2-V2O5-C. It is derived from that the lower ion diffusion barrier indeed contributes greatly to the excellent rate performance [30]. Such porous TiO2-V2O5-C composite can be expected to display excellent electrochemical performance when evaluated as SIHCs due to the considerable contritution of pseudocapacitive capacity. Encouraged by such concern, the coin-type asymmetric supercapacitors were assembled with TiO2-V2O5-C as anodes and commercial active carbon (AC) serving as cathodes, while 1 M NaClO4 in EC-DEC (1:1 v/v) with 5 wt% FEC was introduced as electrolyte. The current were calculated according to the total weight of both electrodes active materials. The electrochemical test was performed in the scope voltage window of 1.0–4.5 V. The galvanostatic charge/discharge curves at different current densities were shown in Fig. 11 with a high Coulombic efficiency confirmed by the symmetric charge/discharge curves at different current densities. The specific capacitance for such TiO2-V2O5-C was about 44.2 F g−1 (corresponding to energy density of 7528.6 Wh kg−1). As shown in Table S1, the energy density and power density for such SIHCs exceed most current SIBs and can rival the most supercapacitors [31]. We also compare the electrochemical performance of TiO2-V2O5-C in this study with the reported TiO2-based nanostructures in literature, the results are listed in Table 1. The previous work that TiO2 based anode composites [32–39] showed super cycleability and rate capability for sodium ion batteries, however, our work showed obvious advantages in rate performance. The reference [40] reported the discharge capacity 140.4 mA hg−1 at repeated 100 cycles and 462.8 mA hg−1 at 200 cycles [41], 108 mA g−1 at 2C up to 1000 long cycles. Although we didn’t perform the long life cycle as long as 1000 cycles, the rate and cycle performances were obtained at various current densities from 0.1 to 20Ag−1, and it still delivered 366 mA hg−1 after repeated 20 cycles and it maintained a discharge capacity as high as 401.3 mA hg−1 after 200 cycles. Compared with other reported work [32,42–44], our work displayed some superiority in both rate and cycle performances, which are sufficient prove such composite material is a promising candidate for Na-ion battery. The extraordinary differences of sodium storage performance of the samples cover the structure advantages and the composition synergistic effects from V2O5 and carbon. The structural design advantages combine the merits of the 3D structure and the high specific capacity of the V2O5 anode. As shown in scheme 2, the 3D structure network serves as a buffer layer to alleviate structural degeneration and increases the effective contact area between electrode and electrolyte. Such mesoporous nanostructure ensures fast electron-transport pathways.

Acknowledgements The authors acknowledge the financial support of the doctor foundation of Shandong province (No. ZR2018BB033), the Postdoctor Science Foundation of China (No. 2018M640616) and the National Science Foundation of China (No. 21805152, 51673106). Appendix A. Supplementary material Supplementary data to this article can be found online at https:// doi.org/10.1016/j.apsusc.2018.12.026. References [1] Y. Liu, X.D. Yan, J.L. Lan, Y.H. Yu, X.P. Yang, Y.H. Lin, Phase-separation induced hollow/porous carbon nanofibers with in-situ generated ultrafine SnOx as anode materials for lithium-ion batteries, Mater. Chem. Front. 1 (2017) 1331. [2] P. Kumar, J. Sundaramurthy, S. Sundarrajan, V. Babu, G. Singh, S. Allakhverdiev, S. Ramakrishna, Hierarchical electrospun nanofibers for energy harvesting, production and environmental remediation, Energy Environ. Sci. 7 (2014) 3192. [3] Y. Tang, Y. Zhang, J. Deng, J. Wei, H.L. Tam, B.K. Chandran, Z. Dong, Z. Chen, X. Chen, Mechanical force-driven growth of elongated bending TiO2-based nanotubular materials for ultrafast rechargeable lithium ion batteries, Adv. Mater. 26 (2014) 6111.

883

Applied Surface Science 473 (2019) 873–884

F. Liu et al.

[4] S. Wang, D. Qu, Y. Jiang, W.S. Xiong, H.Q. Sang, R.X. He, Q. Tai, B. Chen, Y. Liu, X.Z. Zhao, Three-dimensional branched TiO2 architectures in controllable bloom for advanced lithium-ion batteries, ACS Appl. Mater. Interfaces 8 (2016) 20040. [5] H. Liu, S. Zhang, Y. Chen, J. Zhang, P. Guo, M. Liu, et al., Rational design of TiO2@ nitrogen-doped carbon coaxial nanotubes as anode for advanced lithium ion batteries, Appl. Surf. Sci. 458 (2018) 1018. [6] S.R. Hu, Q.T. Huang, Y. Lin, C. Wei, H.Q. Zhang, W.X. Zhang, Z.B. Guo, X.X. Bao, J.G. Shi, A.Y. Hao, Reduced graphene oxide-carbon dots composite as an enhanced material for electrochemical determination of dopamine, Electrochim. Acta 130 (2014) 805. [7] Z. Wang, M. Liu, G. Wei, P. Han, X. Zhao, J. Liu, J. Zhang, Hierarchical self-supported C@TiO2-MoS2 core-shell nanofiber mats as flexible anode for advanced lithium ion batteries, Appl. Surf. Sci. 423 (2017) 375–382. [8] Y. Luo, C. Chen, L. Chen, M. Zhang, T. Wang, 3D reticular pomegranate-like CoMn2O4/C for ultrahigh rate lithium-ion storage with re-oxidation of manganese, Electrochim. Acta 241 (2017) 244. [9] F. Zou, Y.M. Chen, K. Liu, Z. Yu, W. Liang, S.M. Bhaway, M. Gao, Y. Zhu, Metal organic frameworks derived hierarchical hollow NiO/Ni/graphene composites for lithium and sodium storage, ACS Nano 10 (2016) 377. [10] M. Mao, F. Yan, C. Cui, J. Ma, M. Zhang, T. Wang, C. Wang, Pipe-wire TiO2–Sn@ carbon nanofibers paper anodes for lithium and sodium ion batteries, Nano Lett. 17 (2017) 3830. [11] B. Lu, B. Ma, X. Deng, B. Wu, Z. Wu, J. Luo, X. Wang, G. Chen, Dual stabilized architecture of hollow Si@TiO2@C nanospheres as anode of high-performance Liion battery, Chem. Eng. J. 351 (2018) 269. [12] W. Tang, X. Gao, Y. Zhu, Y. Yue, Y. Shi, Y. Wu, K. Zhu, A hybrid of V2O5 nanowires and MWCNTs coated with polypyrrole as an anode material for aqueous rechargeable lithium batteries with excellent cycling performance, J. Mater. Chem. 22 (2012) 20143. [13] S. Tepavcevic, H. Xiong, V.R. Stamenkovic, X. Zuo, M. Balasubramanian, V.B. Prakapenka, T. Rajh, Nanostructured bilayered vanadium oxide electrodes for rechargeable sodium-ion batteries, ACS Nano 6 (2011) 530. [14] M. Cao, Y. Bu, X. Lv, X. Jiang, L. Wang, S. Dai, Y. Shen, Three-dimensional TiO2 nanowire@ NiMoO4 ultrathin nanosheet core-shell arrays for lithium ion batteries, Appl. Surf. Sci. 435 (2018) 641–648. [15] X. Zhu, Z. Meng, H. Ying, X. Xu, F. Xu, W. Han, A novel CoS2 /reduced graphene oxide/multiwall carbon nanotubes composite as cathode for high performance lithium ion battery, Chem. Phys. Lett. 684 (2017) 191. [16] F. Han, C. Zhang, B. Sun, W. Tang, J. Yang, X. Li, Dual-carbon phase-protective cobalt sulfide nanoparticles with cable-type and mesoporous nanostructure for enhanced cycling stability in sodium and lithium ion batteries, Carbon 118 (2017) 731. [17] Y. Zhang, Q. Zhou, J. Zhu, Q. Yan, S.X. Dou, W. Sun, Nanostructured metal chalcogenides for energy storage and electrocatalysis, Adv. Funct. Mater. 27 (2017) 1702317. [18] W. Qiu, J. Jiao, J. Xia, H. Zhong, L. Chen, A self-standing and flexible electrode of yolk-shell CoS2 spheres encapsulated with nitrogen-doped graphene for high-performance lithium-ion batteries, Chem.-A Eur. J. 21 (2015) 4359. [19] X. Huang, Z. Yin, S. Wu, X. Qi, Q. He, Q. Zhang, Q. Yan, F. Boey, H. Zhang, Graphene-based materials: synthesis, characterization, properties, and applications, Small 7 (2011) 1876. [20] D. Wei, Y. Liu, Controllable synthesis of graphene and its applications, Adv. Mater. 22 (2010) 3225. [21] D. Deng, X. Pan, L. Yu, Toward N-doped graphene via solvothermal synthesis, Adv. Mater. 23 (2011) 1188. [22] X. Li, H. Wang, J. Robinson, H. Sanchez, G. Diankov, H. Dai, Simultaneous nitrogen doping and reduction of graphene oxide, J. Am. Chem. Soc. 131 (2009) 15939–15944. [23] X. Wang, Q. Weng, X. Liu, X. Wang, D.M. Tang, W. Tian, C. Zhang, W. Yi, D. Liu,

[24]

[25]

[26]

[27]

[28] [29]

[30]

[31]

[32]

[33]

[34] [35]

[36]

[37]

[38] [39]

[40] [41]

[42]

[43] [44]

884

Y. Bando, D. Golberg, Atomistic origins of high rate capability and capacity of Ndoped graphene for lithium storage, Nano Lett. 14 (2014) 1164–1171. W.W. Chen, T.T. Li, Q. Hu, C.P. Li, H. Guo, Hierarchical CoS2@C hollow microspheres constructed by nanosheets with superior lithium storage, J. Power Sources 286 (2015) 159. Q. Gan, H. He, K. Zhao, Z. He, S. Liu, S. Yang, Plasma-induced oxygen vacancies in urchin-like anatase titania coated by carbon for excellent sodium-ion battery anodes, ACS Appl. Mater. Interfaces 10 (2018) 7031. I.R. Kottegoda, N.H. Idris, L. Lu, J.Z. Wang, H.K. Liu, Synthesis and characterization of graphene–nickel oxide nanostructures for fast charge–discharge application, Electrochim. Acta 56 (2011) 5815–5822. H. Cha, H.M. Jeong, J.K. NKang, Nitrogen-doped open pore channeled graphene facilitating electrochemical performance of TiO2 nanoparticles as an anode material for sodium ion batteries, J. Mater. Chem. A 2 (2014) 5182. G. Qin, Q. Ma, C. Wang, A porous C/LiFePO4/multiwalled carbon nanotubes cathode material for Lithium ion batteries, Electrochim. Acta 115 (2014) 407–415. X. Zhao, C. Yan, X. Gu, L. Li, P. Dai, D. Li, H. Zhang, Ultrafine TiO2 nanoparticles confined in N-doped porous carbon networks as anodes of high-performance sodium-ion batteries, ChemElectroChem 4 (2017) 1516–1522. C. Zhao, C. Yu, B. Qiu, S. Zhou, M. Zhang, H. Huang, B. Wang, J. Zhao, X. Sun, J. Qiu, Ultrahigh rate and long-life sodium-ion batteries enabled by engineered surface and near-surface reactions, Adv. Mater. 30 (2018) 1702486. C. Zhao, C. Yu, S. Liu, J. Yang, X. Fan, He Huang, J. Qiu, 3D porous N-doped graphene frameworks made of interconnected nanocages for ultrahigh-rate and long-life Li-O2 batteries, Adv. Functional Mater. 25 (2015) 6913–6920. S.M. Oh, J.Y. Hwang, C.S. Yoon, J. Lu, K. Amine, I. Belharouak, Y.K. Sun, High electrochemical performances of microsphere C-TiO2 anode for sodium-ion battery, ACS Appl. Mater. Interfaces 6 (2014) 11295. Z. Bi, M.P. Paranthaman, P.A. Menchhofer, R.R. Dehoff, C.A. Bridges, M. Chi, S. Dai, Self-organized amorphous TiO2 nanotube arrays on porous Ti foam for rechargeable lithium and sodium ion batteries, J. Power Sources 222 (2013) 461. L. Wu, D. Buchholz, D. Bresser, L.G. Chagas, S. Passerini, J. Anatase, TiO2 nanoparticles for high power sodium-ion anodes, J. Power Sources 251 (2014) 379. J. Yang, H. Wang, P. Hu, J. Qi, L. Guo, L. Wang, A high-rate and ultralong-life sodium-ion battery based on NaTi2 (PO4)3 nanocubes with synergistic coating of carbon and rutile TiO2, Small 11 (2015) 3744. Y. Ge, H. Jiang, J. Zhu, Y. Lu, C. Chen, Y. Hu, X. Zhang, High cyclability of carboncoated TiO2 nanoparticles as anode for sodium-ion batteries, Electrochim. Acta 157 (2015) 142. J.P. Huang, D.D. Yuan, H.Z. Zhang, Y.L. Cao, G.R. Li, H.X. Yang, Electrochemical sodium storage of TiO2 (B) nanotubes for sodium ion batteries, RSC Adv. 3 (2013) 12593. L. Wu, D. Bresser, D. Buchholz, S. Passerini, Nanocrystalline TiO2 (B) as anode material for sodium-ion batteries, J. Electrochem. Soc. 162 (2015) A3052. H. Usui, S. Yoshioka, K. Wasada, M. Shimizu, H. Sakaguchi, Nb-doped rutile TiO2: a potential anode material for Na-ion battery, ACS ACS Appl. Mater. Interfaces 7 (2015) 6567. F. Yang, Z. Zhang, Y. Han, K. Du, Y. Lai, J. Li, TiO2/carbon hollow spheres as anode materials for advanced sodium ion batteries, Electrochim. Acta 178 (2015) 871. G. Qin, X. Zhang, C. Wang, Design of nitrogen doped graphene grafted TiO2 hollow nanostructures with enhanced sodium storage performance, J. Mater. Chem. A 2 (2014) 12449. K.T. Kim, G. Ali, K.Y. Chung, C.S. Yoon, H. Yashiro, Y.K. Sun, J. Lu, K. Amine, S.T. Myung, Anatase titania nanorods as an intercalation anode material for rechargeable sodium batteries, Nano Lett. 14 (2014) 416. X. Yang, C. Wang, Y. Yang, Y. Zhang, X. Jia, J. Chen, X. Ji, Anatase TiO2 nanocubes for fast and durable sodium ion battery anodes, J. Mater. Chem. A 3 (2015) 8800. L. Wu, D. Buchholz, D. Bresser, L. Gomes Chagas, S. Passerini, Anatase TiO2 nanoparticles for high power sodium-ion anodes, J. Power Sources 251 (2014) 379.