Biphase-Interface Enhanced Sodium Storage and Accelerated Charge

Research Article Cite This: ACS Appl. Mater. Interfaces 2017, 9, 43648−43656

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Biphase-Interface Enhanced Sodium Storage and Accelerated Charge Transfer: Flower-Like Anatase/Bronze TiO2/C as an Advanced Anode Material for Na-Ion Batteries Chenxiao Chu,† Jing Yang,† Qianqian Zhang,† Nana Wang,† Feier Niu,† Xuena Xu,† Jian Yang,*,† Weiliu Fan,*,† and Yitai Qian†,‡ †

Key Laboratory for Colloid and Interface Chemistry of State Education Ministry, School of Chemistry and Chemical Engineering, Shandong University, Jinan 250100, P. R. China ‡ Hefei National Laboratory for Physical Science at Microscale, Department of Chemistry, University of Science and Technology of China, Hefei 230026, P. R. China S Supporting Information *

ABSTRACT: Flower-like assembly of ultrathin nanosheets composed of anatase and bronze TiO2 embedded in carbon is successfully synthesized by a simple solvothermal reaction, followed with a high-temperature annealing. As an anode material in sodium-ion batteries, this composite exhibits outstanding electrochemical performances. It delivers a reversible capacity of 120 mA h g−1 over 6000 cycles at 10 C. Even at 100 C, there is still a capacity of 104 mA h g−1. Besides carbon matrix and hierarchical structure, abundant interfaces between anatase and bronze greatly enhance the performance by offering additional sites for reversible Na+ storage and improving the charge-transfer kinetics. The interface enhancements are confirmed by discharge/charge profiles, rate performances, electrochemical impedance spectra, and first-principle calculations. These results offer a new pathway to upgrade the performances of anode materials in sodium-ion batteries. KEYWORDS: TiO2, nanostructures, biphase interface, sodium-ion batteries, DFT calculations

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rutile TiO2.22 Dawson et al. thought that Na intercalation was preferred in bronze (TiO2(B)) as compared to the cases in anatase and rutile, on the basis of density functional theory (DFT).23 However, the experimental results on TiO2(B) are not so encouraging.24−26 For example, Cao and Gao examined the electrochemical performances of TiO2(B) nanotubes as an anode material for NIBs.24 These nanotubes only presented a capacity of ∼33 mA h g−1 at 400 mA g−1. Although the rate performance could be improved to ∼65 mA h g−1 at 300 mA g−1,25 it is still much lower than that of anatase.22 These results might be related to the poor crystallinity of TiO2(B) in these works. Passerini et al. achieved noodles-like TiO2(B) with an enhanced crystallinity, resulting in a capacity of 90 mA h g−1 at 335 mA g−1 (1 C).26 Even at 10 C, there was still a capacity of 50 mA h g−1. However, increasing the crystallinity of TiO2(B) by high-temperature annealing is always accompanied with a phase transition to TiO2(A).27 So, what would happen to the electrochemical performances, if TiO2(B) is partially converted into TiO2(A)? Moreover, biphase interfaces likely promote the electrode performances in NIBs, because the similar conclusion

INTRODUCTION Sodium-ion batteries (NIBs), as next-generation secondary batteries, have attracted extensive interest, due to a high abundance of sodium and a similar cell setup to lithium-ion batteries (LIBs).1,2 More important, NIBs share the similar “rocking-chair” principle and electrode materials as LIBs, which gives them advantages over other energy storage devices. However, the high electrode potential of Na+/Na and the large ionic radius of Na+ make high-performance electrode materials a great challenge.3 TiO2 as one of the promising anode materials in LIBs, owns the advantages of cost effectiveness, ease in processing, excellent stability, and stable intercalation chemistry.4−12 All these features also benefit its application in NIBs as an anode material. However, similar to the case of LIBs, the low capacity and low electron conductivity of TiO2 greatly affect its performances in NIBs. Thus, many efforts are devoted to address this issue, such as carbon hybridization,13−15 morphology/size control,16−19 and heteroatom doping,20,21 and so on. Compared to the above strategies, there are fewer reports about how to use the multiphase feature of TiO2 for highperformance anodes in NIBs. Su and Wang reported that anatase TiO2 (TiO2(A)) has a lower Na-diffusion barrier and more active sites for Na+ insertion than amorphous TiO2 and © 2017 American Chemical Society

Received: September 4, 2017 Accepted: November 7, 2017 Published: November 7, 2017 43648

DOI: 10.1021/acsami.7b13382 ACS Appl. Mater. Interfaces 2017, 9, 43648−43656

Research Article

ACS Applied Materials & Interfaces

Figure 1. (a) XRD pattern, (b, c) SEM, (d, e) TEM, and (f) HRTEM images of the product (TiO2(A)/TiO2(B)/C) obtained by a solvothermal reaction and a calcination at 500 °C in Ar.

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has been well documented in LIBs.28−38 To date, there are only a few works referring to the mixed phase of TiO2(A) and TiO2(B) as an anode material for NIBs.39−41 The inspiring result was reported by Hu and Huang, who developed a microwave-assisted process for TiO2(A)/TiO2(B) nanosheets attached to reduced graphene oxide (rGO). This composite exhibited a high capacity (∼120 mA h g−1) after 4000 cycles at 0.5 A g−1.41 Even at 12 A g−1, the specific capacity was still around 90 mA h g−1. This work discussed the contribution of rGO/TiO2(B) interface to charge-transfer kinetics in detail. But the influence of TiO2(A)/TiO2(B) interface on the performances was missed. Herein, flower-like assembly of ultrathin nanosheets made by carbon-decorated TiO2(A)/TiO2(B), is synthesized by a modified method. This hierarchical structure exhibits excellent performances as an anode material of NIBs. It could deliver a capacity of 120 mA h g−1 over 6000 cycles at 10 C (1 C = 335 mA g−1). Even at 100 C, there is still a capacity of 104 mA h g−1. All these data are one of the best results for TiO2 anode in NIBs. More importantly, the detailed checks on galvanostatic discharge/charge profiles and electrochemical impedance spectra (EIS), demonstrate that the interface between TiO2(A) and TiO2(B) greatly enhances the performances in terms of both diffusion kinetics and storage capacity, which is further supported by first-principle calculations. This biphase interface, together with carbon materials, further improves the performances, leading to superb cycling stability and rate capability.

EXPERIMENTAL SECTION

Materials Synthesis. TiO2/C nanosheets were synthesized by a simple solvothermal method, followed by an annealing at a high temperature. Typically, 1 mL of TiCl3 (Aladdin, 15.0−20.0% TiCl3 in 30% HCl) was mixed with 30 mL of ethylene glycol (Sinopharm, 99.0%) and 3 mL of deionized water in a Teflon-lined autoclave. After being stirred for 5 min and kept at 150 °C for 4 h, the solution was cooled to room temperature. The product, denoted as TiO2(B), was obtained by centrifugation, washed with deionized water/ethanol several times, and dried at 60 °C overnight. The following calcination at different temperatures or in different atmospheres greatly affect the resultant products. As TiO2(B) was calcined at 500 °C for 4 h in air, it was partially converted to TiO2(A), resulting in TiO2(A)/TiO2(B). Elevating the temperature to 600 °C in air could promote this phase transition and convert TiO2(B) to TiO2(A) completely. However, if the atmosphere was switched to argon at 500 °C for 4 h, then the absorbed organic molecules would decompose to carbon, besides the phase transition. The as-obtained product was denoted as TiO2(A)/ TiO2(B)/C. Material Characterization. X-ray powder diffraction (XRD) patterns were measured on a rotation anode X-ray diffractometer (Bruker D8 Adv, Germany), using graphite monochromatized CuKa line as a radiation. SEM images were recorded on a field-emission scanning electron microscope (SUPRA 55, Germany), where the powders were attached to a conductive tape for less charge accumulation. TEM and HRTEM images were taken from a transmission electron microscope (JEM 1011, Japan), and a highresolution transmission electron microscope (JEOL 2100F, Japan), respectively. The powders were dispersed in absolute ethanol by ultrasound. Then, the suspension was dropped on a clean copper grid 43649


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Figure 2. (a) CV curve, (b) Raman spectrum, (c) TG curve, and (d) N2 sorption curve of TiO2(A)/TiO2(B)/C. coated with an amorphous carbon film. After the evaporation of ethanol, the particles were deposited on the copper grid for observation. Thermal gravimetric analysis (TGA) was carried out on a thermal analyzer (Mettler Toledo TGA/SDTA851, Germany) in air at a rate of 10 °C min−1. Nitrogen sorption isotherms were acquired on a gas sorptometer (Micromeritics ASAP-2020HD88 U.S.A.). Raman spectrum was recorded by a Raman spectrometer (Nicolet NEXUS 670, U.S.A.) at room temperature. Electrochemical Measurements. Electrochemical performances of TiO2 as an anode material in NIBs were evaluated by CR2032 coin cells. The specific capacity of TiO2/C is calculated based on the total mass of TiO2 and carbon. In a typical protocol, TiO2 (or TiO2/C), acetylene black, and polyvinylidene fluoride (PVdF, DodoChem) were mixed in a weight ratio of 7:2:1 in N-methylpyrrolidinone (NMP, Aladdin, >99%). These materials were ground for 30 min, leading to a gray slurry. Then, this slurry was spread on a clean copper foil, and dried at 60 °C in vacuum for 12 h. The loaded foil was roll-pressed and punched into discs with an average diameter of 12 mm. The loading of TiO2 was controlled to be approximately ∼1.2 mg cm−2. The disc was assembled with homemade Na foil as the reference and counter electrode, and glass microfibers of Whatman GF/F as the separator in an argon-filled glovebox (Mikrouna, Super 1220/750/900). The electrolyte was 1 M NaClO4 in ethylene carbonate (EC)/diethyl carbonate (DEC) (1:1 vol., DodoChem) containing 2% fluoroethylene carbonate (FEC, Alladin, >98%). Cyclic voltammograms (CV) were obtained from electrochemical station (LK 2005A, China). Galvanostatic discharge/charge profiles were measured on battery cyclers (LAND CT-2001A, China) within a range of 0.01−2.5 V. Electrochemical impedance spectra were achieved from an electrochemical workstation (AUTOLAB PGSTAT302N, Switzerland) in a frequency range of 100 kHz to 0.01 Hz. In half cells of lithium-ion batteries, the working electrode was made by the same protocol as that of sodium-ion batteries. But a lithium foil was used as the counter and reference electrode, and a Celgard 2400 membrane as the separator. The electrolyte was 1 M LiPF6 in a mixed solvent of dimethyl carbonate (DMC) and ethylene carbonate (EC) (1:1 vol.) purchased from Dodochem. These components were assembled into CR2032 coin cells in an Ar-filled glovebox for CV tests. DFT Calculations. First-principle calculations were conducted using the Vienna Ab Initio Simulation Package (VASP)42,43 within the project augmented wave (PAW)44 approach. The exchange-correlation function based on local density approximation (LDA) was used,45 and the energy cutoff was set at 300 eV for all the calculations. The

Brillouin zone was sampled by a Γ-centered Monkhorst−Pack grid of 2 × 2 × 1 for the unit cell. The structure relaxation and energy calculation would be completed, until Hellmann−Feynman force on each atom was less than 0.05 eV Å−1 and energy converged within 10−5 eV. The resultant lattice parameters were a = b = 3.74 Å and c = 9.53 Å for TiO2(A), and a = 12.10 Å, b = 3.70 Å, and c = 6.49 Å for TiO2(B), close to the reported data. To model the interface of TiO2(A) and TiO2(B), 4 × 4 supercell of Ti-terminated (001) slab of TiO2(A) was constructed to match with a 4 × 2 supercell of Oterminated (100) slab of TiO2(B). A vacuum thickness of 10 Å was set along c axis to avoid the interaction between neighboring slabs. During the relaxation, the atoms at the leftmost layer and the rightmost layer were fixed, and the other was allowed to relax.

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RESULTS AND DISCUSSION Figure 1a shows the XRD pattern of the product obtained by a solvothermal reaction then a calcination at 500 °C in Ar. The diffraction peaks are in good agreement with those from anatase TiO2 (TiO2(A), JCPDS Card No. 21-1272). In spite of this, the presence of a small amount of bronze TiO2 (TiO2(B)) in the product could not be excluded, due to the similar peak positions between TiO2(A) and TiO2(B) and the broad diffraction peaks. SEM and TEM images (Figure 1b−e) show that the product is composed of spherical nanoparticles assembled by ultrathin nanosheets (ca. 10 nm in thickness) (Figure S1 of the Supporting Information, SI). This hierarchical structure possesses many features that are highly desirable for anode materials in sodium-ion batteries, like large specific surface area, well-defined building units, narrow size distribution, and so on. The close check on crystal lattices discloses the existence of TiO2(B) in the structure (Figure 1f). Moreover, there is an obvious lattice distortion at the interface of TiO2(A)/TiO2(B). Such a lattice disorder may provide fast diffusion channels and active sites for Na+, resulting in an additional capacity. The formation process of this unique nanostructure was followed by TEM images (Figure S2). At the beginning of this solvothermal reaction (t < 1.5 h), few precipitates could be collected. Even if the reaction time prolonged to 1.9 h, only a small quantity of nanoparticles was produced (Figure S2a). 43650


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Figure 3. (a) CV curves of TiO2(A)/TiO2(B)/C at a scan rate of 0.1 mV s−1. (b) CV curves of TiO2(A)/TiO2(B)/C at various scan rates. (c) Plot of peak current (ip) vs scan rate (v). (d) Plot of capacity versus scan rate−1/2.

also observed, which could be assigned to graphitic carbon and disordered carbon, respectively. In order to identify the content of carbon, TG analysis was done in air for this structure. As presented in Figure 2c, there is a weight loss in the temperature range of 300−500 °C, indicating the content of carbon at ∼10 wt %. The small amount of carbon might come from thermal decomposition of organic molecules adsorbed on the surface. The uniform distribution of carbon is also supported by element mapping on the structure. As displayed in Figure S5, the signal of carbon uniformly distributes throughout the structure, confirming the existence of carbon in the structure again. Carbon material is important to electrochemical performances of this structure in NIBs, because both TiO2(A) and TiO2(B) are poorly conductive. On the basis of the above results, these spherical aggregates of ultrathin nanosheets could be described as TiO2(A) and TiO2(B) embedded in carbon, denoted as TiO2(A)/TiO2(B)/C. The specific surface area of this structure is characterized by N2 sorption isotherms (Figure 2d). There is a small hysteresis in the P/P0 range of 0.5−1.0, indicating the formation of mesoporous structure. The specific surface area calculated by Brunauer−Emmett−Teller (BET) equation, is approximately 120 m2 g−1. The pore volume and average pore size determined by Barrett−Joyner−Halenda (BJH) model, are 0.4 m3 g−1 and ∼23 nm, as supported by the inset of Figure 2d. The electrochemical performance of TiO2(A)/TiO2(B)/C is measured as an anode material in NIBs. Figure 3a shows the CV curves of this structure for the first five cycles. In the first cathodic scan, the broad peak ranging over 0.6−1.3 V greatly degrades in the following scans, indicating that the underlying reactions are partially reversible. These reactions include Naion insertion into crystal defects, the formation of solid electrolyte interphase (SEI) layer, electrolyte decomposition, and so on.20,40,41 The cathodic peak at 0.02 V probably comes from Na-ion intercalation into carbon materials, including conductive carbon in electrode, or carbon in TiO2(A)/ TiO2(B)/C.47,48 In the first anodic scan, Na-ion extraction from carbon materials and NaxTiO2 appear at 0.1 and 0.85 V as

Some of nanosheets already aggregated or stacked together to minimize surface energy. As the solvothermal reaction underwent (t = 2.7 h), more and more nanosheets were generated, resulting in a prototype of flower-like structures (Figure S2b). After 4 h, the nanosheet aggregates were well developed, as shown in Figure S2c, d. Since TiO2(B) is not a layer structure, the preferential growth to nanosheets driven by crystal anisotropy could be excluded. Another possible reason for this preferential growth, probably comes from the surface adsorption of hydrocarbon moieties related to ethylene glycol or its derivatives. To confirm this point, FTIR spectrum was measured. As shown in Figure S3a, the vibration modes belonging to ethylene glycol could be easily observed in the product after solvothermal reaction. Meanwhile, as the solvent was replaced by ethanol, DMF, or THF, nanosheets disappear from the product (Figure S3b−d), also indicating the role of EG in the formation of nanosheets. This conclusion was the same as that reported by Yang and Che.9 Mixed-phase TiO2(A)/TiO2(B) in the product is also supported by CV curve and Raman spectrum. Figure 2a shows the CV curve of the product, using a lithium foil as the reference and counter electrode. The characteristic peaks of TiO2(B) between 1.4 and 1.7 V are easily observed, accompanied by that of TiO2(A) at 1.7−1.9 V.9,10 The small polarization of the peaks from TiO2(B) (ΔE = 0.1 V) indicates the fast diffusion kinetics, which could be attributed to its openchannel frameworks.11,12 The content of TiO2(B) in the product is estimated by the method proposed by Plouet and Brohan,46 on the basis of the peak areas associated with TiO2(A) and TiO2(B). Using this method, the weight ratio of TiO2(A) and TiO2(B) in the product is about ∼53:47. The similar result is confirmed by Raman spectrum (Figure 2b), where it is fitted by four peaks. The peak at 152 cm−1 originates from TiO2(A), and the other peaks from TiO2(B).37,41f On the basis of the peak areas of TiO2(A) at 152 cm−1 and TiO2(B) at 123 cm−1, the ratio of TiO2(A) to TiO2(B) could be obtained as ∼46:54, close to that deduced from CV curves. Besides these peaks, two broad peaks at 1342 and 1592 cm−1 (Figure S4) are 43651


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Figure 4. (a) Discharge/charge profiles of TiO2(A)/TiO2(B)/C at 0.1 C. (1 C = 335 mA g−1). (b) Discharge/charge profiles of TiO2(A), TiO2(B), TiO2(A)/TiO2(B), and TiO2(A)/TiO2(B)/C at the first cycle. (c) Cycling performance, (d) rate performance, and (f) Nyquist plots of TiO2(A), TiO2(B), TiO2(A)/TiO2(B), and TiO2(A)/TiO2(B)/C. (e) Comparison of the rate capability of TiO2(A)/TiO2(B)/C with the reported data. (g) Long-term cycling performance of TiO2(A)/TiO2(B)/C at 10 C.

structure and abundant interfaces of TiO2(A)/TiO2(B),36 promoting the charge-transfer kinetics. Figure 4a shows the discharge/charge profiles of TiO2(A)/ TiO2(B)/C, where they are slope lines except the first discharge curve. This shape is in good agreement with capacitive feature of these sodiation/desodiation processes. The average voltage is ∼0.7 V, much lower than that in LIBs (∼1.5 V). This would improve the output voltage and power density of full cell, when it is paired with the same cathode. The first discharge/charge capacity is 947/353 mA h g−1, indicating an initial Coulombic efficiency at 37.3%. The low initial Coulombic efficiency could be attributed to irreversible reactions in the first discharge process, as listed in the discussion about the first cathodic scan of CV curves. These data could be greatly improved by using proper binder, new electrolyte or presodiation.20,41,49 In our case, using ether as the electrolyte, could improve the initial Coulombic efficiency to 54.3% (Figure S6). After several cycles, the Coulombic efficiency quickly increases to 99.0%, indicating the high reversibility. Figure 4b depicts the first discharge/

very broad peaks. From the second cycle, the scans basically repeat the curve, indicating the good stability. The shallow and obscure peaks in CV curves imply that the sodiation/ desodiation processes in this case are mainly surface-controlled. This conclusion is confirmed by CV curves at different scan rates. As shown in Figure 3b, these CV curves display a similar shape at different scan rates. The relationship of peak currents (ip) and various scan rates (v) could be described by the following equation, ip = avb. b = 0.5 implies an ideally diffusioncontrolled process, whereas b = 1.0 indicates a surfacecontrolled process, i.e., a capacitive behavior. In our case, the b-values of anodic and cathodic peaks could be quantified as 0.91 and 0.88 (Figure 3c), reflecting the predominance of the capacitive behavior in these processes. The plot of capacity versus v−1/2 suggests that the capacity slowly degrades, as the scan rate increases from 0.2 mV s−1 to 3 mV s−1 (Figure 3d). This result confirms the capacitive characteristics of these processes again. It could be correlated with the hierarchical 43652


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Ω). The similar case also happens to Rct of TiO2(A)/TiO2(B) (63.9 Ω), as compared to those of TiO2(A) (110 Ω) and TiO2(B) (107 Ω). These results imply that the interfaces within TiO2(A)/TiO2(B) effectively reduce the resistances in the course of Na-ion insertion/extraction, thereby benefiting the related kinetics. TiO2(A)/TiO2(B)/C is even better than TiO2(A)/TiO2(B), thereby resulting in the highest capacity retention. The check on Na ion diffusion coefficients of these materials also agrees with this conclusion. The diffusion coefficient of Na ions in TiO2(A)/TiO2(B)/C is 5.72 × 10−11 cm2 s−1, much faster than 1.26 × 10−11 cm2 s−1 for TiO2(A)/TiO2(B), 5.48 × 10−12 cm2 s−1 for TiO2(A), and 2.68 × 10−12 cm2 s−1 for TiO2(B). Figure 4g shows the cycling performance of TiO2(A)/TiO2(B)/C at a high current density of 10 C. The reversible capacity is remained at 120 mA h g−1 throughout the whole cycling (6000 cycles), indicating the excellent electrochemical stability. This stability is also confirmed by discharge/charge profiles (Figure S10), which displays constant electrode polarization upon this high-rate cycling. Moreover, the flower-like structure organized from ultrathin nanosheets could be easily observed in the product after 1200 cycles at 10 C (Figure S11). Encouraged by these results, the electrode with a high mass loading of TiO2(A)/ TiO2(B)/C at 2.5 mg cm−2 was also tested (Figure S12), where the specific capacity still remains 143 mA h g−1 after 300 cycles at 0.5 C. Although these data are lower than those obtained at a lower mass loading, the areal capacity increases from 0.24 mA h cm−2 at ∼1.2 mg cm−2 to 0.36 mA h cm−2 at ∼2.5 mg cm−2. The excellent performances of TiO2(A)/TiO2(B)/C could be ascribed to the following aspects. First, the mixed-phase interface between TiO2(A) and TiO2(B) not only offers extra active sites for redox reactions but also improves reaction kinetics by reducing the resistances. The similar effects of the interface on electrochemical performances were already documented for anode materials in LIBs.32−38 To be specific, Jamnik and Maier proposed a so-called “job-sharing” mechanism to explain excess Li storage at the interfaces, analogue to the case of Li storage at solid/solid interfaces of metal/oxide.33−35 In this mechanism, the synergistic cooperation of Li+-accepting phase (oxide) and electron-accepting phase (metal) makes the additional Li-storage at the interface reversible. Fan and Wu applied this theory to account for high capacity of TiO2(A)/TiO2(B) in LIBs, where the interface is between two oxides.36 Xi, Guo, and Yang further investigated the Li storage at the interface of TiO2(A)/TiO2(B) using firstprinciple calculations, and confirmed the charge separation at the interface.37 In order to clarify if the interfacial storage also works for Na+ ions in our case, the similar calculations based on DFT are carried out. At first, TiO2(A) and TiO2(B) bulks are relaxed individually. The resultant lattice parameters are a = b = 3.74 Å and c = 9.53 Å for TiO2(A), and a = 12.10 Å, b = 3.70 Å, and c = 6.49 Å for TiO2 (B), close to the experimental data.50,51 These results confirm the validity of our calculation. Then, the interface of TiO2(A)/TiO2(B) is constructed by Ti-termined (001) plane of the TiO2(A) slab and O-termined (100) plane of the TiO2(B) slab (Figure 5a, b), due to their similar orientation. The atoms at the leftmost and rightmost layer remain fixed, whereas the other atoms are allowed to relax. After the relaxation, the atom arrangement at the interface is carefully checked. At this interface, Ti atoms at the (001) plane of TiO2 (A) prefer to act as electron acceptors, thus accompanied by the transition from Ti4+ to Ti3+. Accordingly, O atoms at the (100) plane of TiO2(B) likely trap Na+ ions,

charge profiles of TiO2(A)/TiO2(B)/C at 0.1 C. To gain insights about the performance, TiO2(A), TiO2(B), and TiO2(A)/TiO2(B) were prepared by the similar protocol (Figure S7), except the last calcination was conducted in air. Then, their performances were measured under the same conditions as that for TiO2(A)/TiO2(B)/C. Compared to TiO2(A) and TiO2(B), TiO2(A)/TiO2(B) exhibits a larger charge capacity and a higher initial Coulombic efficiency (Table S1), confirming the benefit of the mixed phase on the extraction/insertion of Na ions. Because the results were obtained at a low current density on ultrathin sheet-like structure, the influences of reaction kinetics to a high capacity are almost negligible. In this case, it is believed that the high capacity of TiO2(A)/TiO2(B) likely comes from enhanced Na+ storage at the interface, which will be discussed later. The similar enhancement was also observed for the comparison between TiO2(A)/TiO2(B) and TiO2(A)/TiO2(B)/C, confirming the interface contribution again. Figure 4c presents the cycling performance of TiO2(A), TiO2(B), TiO2(A)/TiO2(B), and TiO2(A)/TiO2(B)/C at 0.5 C. It is noted that the capacity of TiO2(A) gradually increases during the first tens of cycles, which could be attributed to electrochemical activation.17−19 Since Na-ion diffusion in TiO2(A) is quite sluggish, the active materials far away below the surface might be not utilized, particularly at a high rate like 0.5 C. So, after several cycles, the trapped Na ions or structure reorganization helps Na-ion diffusion easily, achieving the capacity increase. If the current density is lowered to 0.1 C, then this capacity increase should attenuate, as supported by Figure 4d. After the first tens of cycles, all the capacities level off for the following cycles. The reversible capacity of TiO2(A)/TiO2(B)/C is kept at ∼200 mA h g−1 after 1000 cycles, much better than those of TiO2(A) (∼107.3 mA h g−1), TiO2(B) (∼68.4 mA h g−1), and TiO2(A)/ TiO2(B) (∼141.2 mA h g−1). The excellent electrochemical performance of TiO2(A)/ TiO2(B)/C is also validated in terms of rate capability (Figure 4d), where it delivers a capacity of 217 mA h g−1 for 1 C, 179 mA h g−1 for 5 C, 154 mA h g−1 for 10 C, and 149 mA h g−1 for 20 C. Even at 100 C, the reversible capacity is still 104 mA h g−1 (Figure S8). These results are not only superior to those of anatase and bronze, but also better than many of reported TiO2 (Figure 4e).15−20,41 For instance, petal-like rutile TiO2 nanostructures covered by a graphene-like carbon layer, only exhibited a small capacity of 59.8 mA h g−1 at 12.5 C.15 These data could be promoted to 102 mA h g−1 at 6000 mA g−1 (∼18 C), by using TiO2(A)/TiO2(B) nanosheets deposited on rGO.41 The similar result was also reported for N-doped yolklike anatase TiO2 wrapped by carbon, where the capacity at a rate of 20 C is 116 mA h g−1.17 The rate capability evaluated by capacity retention could reflect the effect of mixed-phased interface and carbon on reaction kinetics. As shown in Figure S9, TiO2(A)/TiO2(B)/C shows the highest capacity retention of ∼51.3%, which is higher than TiO2(A)/TiO2(B) (43.1%), TiO 2 (A) (35.8%), and TiO 2 (B) (25.5%). These data demonstrate that the mixed-phase interface promotes the reaction kinetics and the carbon matrix further intensifies this trend. The superior kinetics of TiO2(A)/TiO2(B)/C is also supported by EIS (Figure 4f). TiO 2(A)/TiO 2(B) and TiO2(A)/TiO2(B)/C presents a similar diameter for the depressed semicircle, closely related to surface-film resistance (Rf), and the charge-transfer resistance (Rct). As fitted by equivalent circuit (Table S2), Rf of TiO2(A)/TiO2(B) (14.3 Ω) is smaller than those of TiO2(A) (19.1 Ω) and TiO2 (B) (30.6 43653


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principle calculations,41 and found that the interface effectively reduce the diffusion barriers of Na ions, thereby facilitating the charge-transfer kinetics. This phenomenon was supported by the decreasing of Rf and Rct and the increasing of diffusion coefficient in TiO2(A)/TiO2(B) in our case. The similar conclusion has been documented for TiO2 in lithium-ion batteries.36,37 So, excess active sites, together with enhanced charge-transfer kinetics, make the mixed phase of TiO2 promising in NIBs. The ratio of TiO2(A)/TiO2(B) in the composite could be tuned by the annealing temperature (Figure 6a), which allows us to further understand the interface effect on electrochemical performances. As displayed in Figure 6b, all these products show the similar cycling stability. But the composite obtained by annealing at 500 °C, in which the ratio of TiO2(A)/TiO2(B) is roughly 1:1, shows the highest capacity. The deviation of this ratio from 1:1 would lower the interfacial storage capacity, but increase the bulk intercalation either in TiO2(A) or in TiO2(B). Thus, the balance between bulk intercalation and interfacial storage endows the composite annealed at 500 °C the best performances. Second, uniform distribution of carbon throughout the composite further promotes the charge-transfer kinetics. Moreover, the big size and spherical shape of this structure also facilitate the material process during the electrode fabrication. Third, large specific surface area caused by nanosheets could promote the infiltration of electrolyte into electrode. The large surface area also reduces the areal current density, benefiting the reaction kinetics. All the features make the hierarchical structure as anode materials promising, which has been demonstrated in many anode materials in lithium-ion batteries.36,52−54

Figure 5. A proposed model about sodium storage at the interface of TiO2(A)/TiO2(B). Relaxed interface structure doped by Na along [100] direction of TiO2(A) and [001] direction of TiO2(B) (a), or along [010] directions of TiO2(A) and TiO2(B) (b). (c, d) The charge-density difference at the interface before and after Na doping. The blue region indicates the electron density has been depleted, while the red region indicates the density has been accumulated.

due to high affinity between them. Such a distribution leads to a charge separation at the interface, corresponding to electron transfer from Na-doped TiO2(B) to TiO2(A). To further illustrate this charge separation, Bader charge analysis is conducted for the interface with and without Na doping. Before that, the location of doped Na should be identified. Although many sites in TiO2(B) like C2 are taken into accounts, C1 sites at the interface are more stable, suggesting the preferential location of Na. As summarized in Table S3, the doped Na atom donates 0.77e to TiO2. But there is only an average increase of electrons about ∼0.03e for Ti atoms and ∼0.02e for O atoms on the side of TiO2(B). Associated with the positive charge from doped Na atom, there is still a net electron flow out of Na-doped TiO2(B). This result is also confirmed by electron increase on the side of TiO2(A), ∼ 0.03e for both Ti and O atoms. Such a charge separation at the interface could be directly visualized by charge density difference in Figure 5c, d. These results clearly prove the charge separation at the interface of TiO2(A)/TiO2(B), leading to extra sites for Na storage. Besides offering extra sites for Na storage, the interface also enhances the charge-transfer kinetics. Hu and Huang investigated the interface between TiO2 and rGO by first-

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CONCLUSIONS In summary, spherical assembly of ultrathin nanosheets made by TiO2(A) and TiO2(B) in carbon matrix were successfully prepared by a facile process. The mixed-phase interface, including crystal boundaries and lattice distortion, effectively increase the reversible capacity and promote the reaction kinetics, as supported by discharge/charge profiles, rate capability, EIS, and first-principle calculations. Moreover, this feature does not compromise the cycling stability, leading to excellent performances as an anode material in NIBs together with carbon matrix in the composite. At 100 C, the reversible capacity of this composite is preserved at 104 mA h g−1. The long-term cycling at 10 C for 6000 cycles shows excellent stability, giving a capacity of 120 mA h g−1. These results offer new opportunities to improve performances of anode materials in NIBs.

Figure 6. (a) CV curves and (b) cycling performances of the products calcined at 400 °C, 500 °C, 600 °C for 4 h in argon, denoted as T400, T500, and T600. (1 C = 335 mA g−1). 43654


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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b13382. SEM images, Raman spectra, element mapping and initial galvanostatic curves of TiO2(A)/ TiO2(B)/C; CV curves and SEM images of TiO2(A)/TiO2(B), TiO2(A), and TiO2(B); galvanostatic discharge/charge profiles of TiO2(A)/TiO2(B)/C at different rates and at selected cycles; long-term cycling performance of TiO2(A)/ TiO2(B)/C at 10 C and SEM image after 1200 cycles; and cycle and rate performance of TiO2(A)/TiO2(B)/C at a high mass loading (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (J.Y.). *E-mail: [email protected] (W.F.). ORCID

Jian Yang: 0000-0002-6401-276X Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the National Nature Science Foundation of China (Nos. 21471090 and 61527809), Key Research and Development Programs of Shandong Province (2017GGX40101), and Taishan Scholarship in Shandong Provinces (No. ts201511004). J.Y., Q.Z., and W.F. conducted the theoretical calculation and prepared the related text.

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Biphase-Interface Enhanced Sodium Storage and Accelerated Charge

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