Nanowires embedded porous TiO2@C nanocomposite anodes for enhanced stable lithium and sodium ion battery performance

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Nanowires embedded porous TiO2@C nanocomposite anodes for enhanced stable lithium and sodium ion battery performance Yu Wanga, Na Lia, Chuanxin Houa, Biao Hea, Jiajia Lia, Feng Danga,∗, Jun Wanga,∗∗, Yuqi Fanb,∗∗∗ a Key Laboratory for Liquid-Solid Structural Evolution and Processing of Materials (Ministry of Education), Shandong University, 17923 Jingshi Road, Jinan, 250061, China b Institute of Environment and Ecology, Shandong Normal University, Jinan, 250014, China

A R T I C LE I N FO

A B S T R A C T

Keywords: Nanocomposites Anode material Lithium ion batteries Sodium ion batteries

A porous carbon nanocomposite with embedded TiO2 nanowires (NWs) was synthesized using a two-step synthetic method in which carbon matrix was obtained by carbonizing a vacuum dried gel. This unique structure in which TiO2 nanowires uniformly distributed in and tightly bonded to the carbon matrix shortened the electron transport path and reduced the transmission resistance. Nanoporous structure ensured continuous transfer of Li+/Na+ and supplied a large specific surface area of 280.82 m2 g−1 to provide more active sites. Different from other existing works on TiO2@C anode materials with TiO2 loading higher than 60 wt%, the obtained very small amount of TiO2 (~12 wt%) improved the electrochemical and long-cycle performance of carbon substrate with TiO2 NWs embedded significantly, due to uniformly distributed TiO2 NWs throughout the carbon matrix. These TiO2@C composite anodes could deliver a specific capacity of 286 mA h g−1 at 0.3 C, 197 mA h g−1 at 0.15 C for lithium and sodium ion batteries, respectively. It maintained remarkably stable reversible capacities of 128 and 125 mA h g−1 for lithium and sodium ion batteries at 3 C during 2500 cycles, respectively. Smaller fluctuations and smoother curves demonstrated that sodium ion storage was more stable than lithium ion storage for the TiO2@C composite anode. In addition, the capacitive contributions of TiO2@C in both systems are quantified by kinetics analysis.

1. Introduction

accommodation of one Li+/Na+ per TiO2 [12]. The performance of TiO2 can be greatly influenced through controlling its phase and nanostructure. Moreover, one could tailor the quantum confinement effects, changes in the transport behavior of electrons and holes, and shifts in the electronic band structure [13–15]. Among their myriad of polymorphs, the TiO2–B (Bronze) shows a distinctive channel structure for ion migration, leading to fast charge/discharge capability of lithium/sodium batteries [16]. Due to unique sites and energetics of ion absorption and diffusion in structure, fast charge performance and long term cyclability can be achieved, and the pseudocapacitive process contributes to the majority of charge storage [17,18]. Large specific surface area and more active sites in crystal structure are of great importance to achieve high performance LIBs and SIBs. TiO2-based materials with different nanocrystalline forms have been widely used to reduce the diffusion length of ions and increase the contact area between electrolyte and electrode, accordingly improving both storage capacity and rate performance [19–21]. For instance, one-

Rechargeable batteries have penetrated into almost every aspect that required energy storage in our real-world [1–3]. Compared with lithium ion batteries (LIBs), sodium ion batteries (SIBs) have been beneficial from more abundant sodium to achieve low-cost batteries. The larger radius and heavier sodium atoms result in migration issues during the charge/discharge processes. In pursuit of high-performing electrodes, TiO2 has been regarded as an appealing anode alternative to carbon- or silicon-based anodes in LIBs or SIBs, not only due to its good stability, environment-friendly, intrinsic safety and cheap price but also because of the negligible volume change during the ion insertion/extraction [4,5]. Since the pioneering work by Fujishima and Honda to study the phenomenon of photocatalytic water splitting on a TiO2 electrode [6], the application of TiO2 has been expanded from photocatalysis [7] and photovoltaics to LIBs [8,9] and SIBs [10,11]. TiO2 has a comparable theoretical capacity of 335 mA h g−1 depending on the

∗

Corresponding author. Corresponding author. ∗∗∗ Corresponding author. E-mail addresses: [email protected] (F. Dang), [email protected] (J. Wang), [email protected] (Y. Fan). ∗∗

https://doi.org/10.1016/j.ceramint.2019.12.161 Received 25 November 2019; Received in revised form 16 December 2019; Accepted 19 December 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.

Please cite this article as: Yu Wang, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.12.161

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Scheme 1. (a) Steps for preparing TiO2@C anode materials by gel carbonized process; (b) Photo images of samples at different steps in the synthesis of TiO2@C material.

charge/discharge, thereby enduring the electrode. Furthermore, the porous structure supplies a larger specific surface area to increase the activity and the rate performance. Therefore, it's promising to develop TiO2@C material with a nanoporous structure that combines the above characteristics. In this work, a porous carbon nanocomposite with lowly loaded TiO2 NWs was developed and demonstrated through a two-step methodology, i.e., carbonizing and etching a gel containing TiO2 NWs and SiO2 spheres. TiO2 NWs were synthesized using a hydrothermal method as described in our previous work [27]. Different from the reported TiO2-based TiO2@C composites as anode materials for LIBs, in which the content of TiO2 is usually higher than 60 wt% [35–37], only 12 wt% TiO2–B NWs were loaded into the porous carbon matrix. The resulting TiO2@C composites with a porous structure offered the following characteristics: (1) TiO2 NWs with length of ca. 15 μm and diameters of ca. 30 nm were assembled by Ti–O octahedral and thus featured TiO2–B structure (Fig. S1, supporting information). This construction increased the specific surface area and offered a reduced pathway for ion insertion/extraction. (2) TiO2 NWs tightly were bonded to the carbon matrix and formed a network structure, while the amount of TiO2 NWs was reduced greatly, and still reflected the electrochemical properties of TiO2. Using carbon as a matrix greatly enhanced the conductivity and reduced the impedance. (3) TiO2 NWs and SiO2 spheres were added and well mixed in the agar solution before heated to form a gel. Through etching the SiO2 spheres, a porous framework was obtained with pores of ca. 500 in size. The specific surface area of resulting TiO2@C reached up to 280 m2 g−1. The pore size distribution of the particles concentrated in 4 nm. Nanoporous structure and mesopores provided a large specific surface area which brought about superior lithium/sodium storage performance. As a result, the nanoporous TiO2@C composite anode presented a high capacity, i.e., 286 mA h g−1 at 0.3 C for

dimensional (1D) TiO2 NWs have a large specific surface area and allow maintaining firm electronic contacts with the conductive agents during the charge/discharge [22], yielding higher and faster electron-hole separation rate [23]. TiO2 nanostructures can be obtained through a sol-gel method, liquid deposition, and hydrothermal method [24–26]. Additional driving force is needed to form 1D TiO2 NWs because TiO2 has poor anisotropic properties in the crystal structure [27]. Currently, there exist several hurdles for realizing TiO2 as highperforming battery materials: i) its low electronic conductivity and ion diffusion ability; ii) the high porosity of aggregated TiO2 NWs leading to loss of connection between TiO2 NWs during cycling [8]; and iii) more expensive than most commercial carbon materials. One of the solutions is to compound it with the carbon matrix to improve the conductivity as well as capacity and stability. For instance, in LIBs, Yang et al. synthesized uniform TiO2@carbon composite nanofibers which exhibited a reversible capacity of 206 mA h g−1 to 100 cycles at 30 mA g−1 [28]. Zhao et al. synthesized carbon decorated TiO2 delivering a reversible capacity of 130 mA h g−1 at a specific current of 300 mA g−1 [29]. For SIBs, Kim et al. synthesized anatase TiO2 nanorods with carbon coating to deliver an initial capacity of 193 mA h g−1 at 10 mA g−1 [30]. Lee et al. prepared carbon/TiO2 composite and achieved 68 mA h g−1 at 440 mA g−1 [31]. Moreover, Huang et al. prepared a graphene-coated TiO2 material with a preferable rate capability and long cycle life [32]. However, the poor rate performance and relatively short cycle life still limited its application. Materials based on transformation reactions and alloying reactions suffer from a large volume change (> 200%) during charge/discharge, being a major failure mechanism of batteries [33]. Therefore, using porous nanostructures has been demonstrated as a good way to reduce the volume changes [34]. The internal pore structure is capable of accommodating the structural stress caused by the ion insertion/extraction during the 2

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Scheme 2. Schematic illustration of microscopic changes of TiO2@C during synthesis.

LIBs and 197 mA h g−1 at 0.15 C for SIBs, superior rate capability and extraordinary cycle stability almost without any loss up to 2, 500 cycles at 3 C in both systems.

carbon black), and Polyvinylidene Fluoride (PVDF) binder with a weight ratio of 8: 1: 1. Next, the Cu mesh was kept at 120 °C in vacuum for 8 h. Metallic lithium or sodium foil was used as the counter electrode. The electrolyte used for LIBs was different from that used for SIBs, and was a mixture of ethylene carbonate and dimethyl carbonate (1:1 (v/v)) and 1 M NaClO4 in a mixture of ethylene carbonate and dimethyl carbonate (1:1 (v/v)) with 5 wt% of fluoroethylene carbonate added, respectively. The mass loading on the electrode was around 1.4 mg cm−2. Cyclic voltammetry (CV) measurement was tested by a PARSTAT 2273 workstation. Electrochemical impedance spectroscopy (EIS) experiments were obtained on Autolab PGSTAT302 N with a frequency range of 100 kHz to 0.1 Hz. The charge/discharge tests were conducted at different current rates, where 1 C = 335 mA g−1 for LIBs and SIBs. The range of test voltage for LIBs and SIBs was from 0.01 to 3.0 V and 0.01–2.5 V, respectively. The electrochemical cell was galvanostatically discharged and charged using multichannel battery testers (LAND, CT2001A).

2. Experimental section 2.1. Synthesis of agar gel matrix containing TiO2 NWs and SiO2 spheres To obtain the agar gel matrix with uniform distribution of TiO2/ SiO2, a two-step gel carbonized method was applied, as shown in Scheme 1a and 1b. In this work, 0.48 g commercial SiO2 spheres (diameters of 400–600 nm) and 0.2 g prepared TiO2 NWs were added to 20 mL distilled water. After 1 h of ultrasonic dispersion, 1.2 g agar was dispersed under mechanical stirring. Then the solution was heated to 180 °C to form a uniform gel matrix with TiO2/SiO2 distributed. The matrix was cooled down to room temperature and then treated by vacuum freeze-drying method. In this step, as the water sublimes, the gel mixture would self-shrink and the silica spheres are ought to be squeezed together. In this process, TiO2 NWs are still distributed in the gel matrix and formed a network structure.

3. Results and discussion

2.2. Synthesis of TiO2@C composites

TiO2–B NWs embedded TiO2@C composites were prepared by a gel carbonized method (Scheme 1). TiO2 NWs and spherical SiO2 were dispersed in the agar gel matrix by ultrasonic dispersion during heattreatment. After the freeze-drying process and carbonized at 900 °C in N2 atmosphere, the sample was corroded by NaOH solution. Then the TiO2@C composites were obtained. This method greatly reduced the amount of TiO2 (12 wt%) and produced the sample on a large-scale. As illustrated in the ideal microscopic model (Scheme 2), SiO2 spheres are squeezed tightly by the shrinkage of the gel mixture during the freezedrying process. TiO2 NWs are disorderly distributed in the carbon matrix to form a three-dimensional network structure. After carbonization and corrosion treatment, the nanoscale pore structure with a disordered contribution of TiO2 was obtained. Fig. 1 displays the SEM images (Fig. 1a), TEM graphics (Fig. 1b and c), HRTEM image (Fig. 1d) and EDS analysis (Fig. 1e, f and 1g) of the TiO2@C material. TEM image of as-prepared pure TiO2 NWs shows that the measured diameters of the NWs were around 30 nm and the length could reach 10–15 μm (Fig. S2). After the corrosion, spherical SiO2 was corroded and spherical pores were in situ formed (Fig. 1a). No reaction between TiO2 NWs and NaOH was observed and TiO2 NWs still remained in the carbon matrix. HRTEM revealed the presence of (020) planes of TiO2–B. Eventually, a dense network composed of TiO2 NWs with diameters of ca. 30 nm was distributed in the carbon matrix and tightly bonded to the carbon matrix, in which nanopores with diameters of ca. 500 nm were formed after the corrosion of spherical SiO2 by NaOH solution (Fig. S3 and Fig. 1b). This nanoporous structure could create a large specific surface area and numerous active sites. A dense network consisting of TiO2 NWs in the carbon matrix could supply a fast transfer channel for electron which is favorable to the rate

The resulting block was carbonized at 900 °C for 2 h in nitrogen atmosphere and then corroded by 1 M NaOH solution at 80 °C. After washed several times by distilled water and collected by centrifugation, the TiO2@C composites were finally obtained. 2.3. Material characterizations The X-ray diffraction (XRD) patterns were recorded on a D/Max 2550 X-ray diffraction using Cu Kα radiation (1.5418 Å) operating at 200 mA and 50 kV. The nanostructure of the resulting materials was observed by scanning electron microscopy (SEM, Keyence VE-9800) with an energy dispersive spectrometer (EDS). Transmission electron microscope (TEM) images were observed on a JEM-2100F TEM at 200 kV. Raman spectra were tested by a confocal microscopic Raman spectrometer (Horiba Jobin Yvon, LabRAMHR) with 532-nm laser excitation. The N2 adsorption/desorption isotherm was carried out to investigate the specific surface area and porosity from a Micrometritics ASAP 2020 analyzer at 77 K. Thermogravimetric analysis was determined using a thermal TGA/SDTA instrument in air condition. The compositions and elements states of the surface were examined by Xray photoelectron spectroscopy (XPS, ESCALAB 250). 2.4. Electrochemical characterization CR2032 coin-type cells were assembled in a glove box filled with argon. The working electrode was pressed uniformly on a Cu mesh, which was composed of the prepared materials, super P (conductive 3

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Fig. 1. (a) SEM images of the synthesized TiO2@C composites after the corrosion; (b and c) TEM and (d) HRTEM images of the synthesized TiO2@C composites; EDSelemental mapping images of TiO2@C composites at Fig. 1a: (e) C, (f) Ti and (g) O.

correspond to lattice oxygen of Ti–O bonds, OH- groups and absorbed water on the surface of the sample [38,42]. To study the electrochemical performance, Fig. 3a shows the CV curves of TiO2@C between cutoff voltages of 0.01 V–3 V vs. Li/Li+ at 0.2 mV s−1 with button-type cells. In the first discharge curve, the reduction peak of the electrode located at 1.65 V represented an irreversible insertion of Li ions [43]. A large irreversible peak below 1.5 V in the first cathodic scan was observed. Stated from the second cycle, TiO2@C anodes showed only couples of reversible redox peaks located at 1.0 V and 1.25 V, which were not apparent due to the higher percentage of carbon [39]. Therefore, it can be concluded that the irreversible peak originated from the formation of SEI layer during the first cycle and became stable from the initial cycles [37,39]. At the first three cycles, cathodic and anodic peaks around 0.2 and 0.4 V corresponded to the reduction and oxidation of graphitized carbon [44]. Fig. 3b displays the charge-discharge curves and cycling performance of TiO2@C at 0.3 C. The electrodes exhibited a long sloping curve without plateaus, owing to the presence of amorphous carbon [45,46]. In the first and second discharge processes, TiO2@C delivers high specific capacities of 569.5 and 325.8 mA h g−1 (based on the weight of TiO2@C) respectively. The significant reduction of capacity was mainly owing to the formed SEI film and the irreversible trapping of Li+ in TiO2@C during the first cycle [47]. Starting from the third cycle, the capacity was maintained at 270 mA h g−1 and the capacity loss became very low, indicating good reversibility of electrochemical reaction. The TiO2@C exhibited specific discharge capacities of 271, 221, 182, 130, 90, 51 and 32 mA h g−1 at 0.3, 0.5, 1.5, 3, 5, 15 and 30 C, respectively (Fig. 3c). Compared with pure TiO2 NWs, the rate performance has been significantly improved more than 120% in specific capacity. The specific capacity and rate performance were superior to reported TiO2/carbon composite used as anodes for LIBs [39,43]. Fig. 3d presents extraordinary cycling stability of TiO2@C anode at 3 C. After 2500 cycles, TiO2@C remained specific capacities of 128 mA h g−1, i.e., 98.5% remaining, indicating superior stable anode of TiO2@C material. Fig. 4 displays the EIS data of TiO2@C electrodes in LIB before and after the CV tests. The resulting EIS curves consist of a semicircle in the high frequency region and a straight line in the low frequency region, indicating that the electrochemical process is controlled by both charge transfer and Li+ diffusion [5]. In the equivalent circuit model in the insets of Fig. 4a and b, the intercept of high-frequency semicircle on the Z′ axis represents the resistance of electrolyte (Re), high-frequency

performance and stability of LIBs/SIBs. EDS analysis displays the distribution of C, Ti and O atoms in the TiO2@C material, implying even distribution of TiO2 NWs in the carbon matrix. EDX comparative elemental analysis before and after corrosion (Fig. S4) suggests that SiO2 has been completely removed, leaving the composite of carbon matrix and TiO2 NWs. The fact that carbon does not protect the SiO2 spheres may be due to the fact that the NaOH solution can enter the carbon matrix from the mesopores on the carbon substrate. To give further insight into the nanoporous structures and pore-size distribution of TiO2@C, XRD and porosity analysis were performed. The relatively high diffraction peak of the (020) face of TiO2–B (Fig. S5a) indicates the growth of TiO2 NWs in that direction. Before and after treatment at 900 °C, the almost same diffraction peak position of the TiO2 NWs in the composite indicates the maintained linear structure of the NWs and amorphous structure of carbon. Fig. S5b shows the N2 adsorption-desorption isotherm loop and pore size distribution curve of TiO2@C. The specific surface area was up to 280.8 m2 g−1 which is twice or more times of reported TiO2/C composite (102.71 m2 g−1 for microsphere C–TiO2 [38], 123.31 m2 g−1 for TiO2/Carbon composite nanofibers [39] and 5.86 m2 g−1 for titania/carbon composite microspheres [11]. The size distribution of the composite was concentrated in 4 nm (inset of Fig. S5b), proving the existence of mesopores. The Raman spectra in Fig. S5c with two prominent peaks at about 1583 and 1362 cm−1, corresponding to the D and G bands of carbon, further confirm the presence of amorphous carbon in TiO2@C. To determine the content of TiO2 in the composite, the detected weight loss of about 10% from TGA under air under 200 °C (Fig. S5d) was caused by the evaporation of the adsorbed water. The weight loss from 400 to 550 °C was due to the burn-off of carbon, and TiO2 NWs remained stable. Thus, the content of TiO2 NWs in the composite was about 12 wt%. XPS has been employed to confirm the presence of Ti, O and C elements (Fig. 2a). The peaks of Na were also found, probably because the Na+ was not completely removed after the corrosion. In the high resolution XPS spectrum of Ti 2p (Fig. 2b), the Ti4+ 2p3/2 and Ti4+ 2p1/ 2 peaks are located at binding energies of 458.5 and 464.8 eV, respectively, in agreement with the value of Ti4+ in the TiO2 lattice [40]. Two additional peaks at 456.3 and 461.8 eV can be assigned to Ti3+ 2p3/2 and Ti3+ 2p1/2, respectively [11]. The high-resolution C 1s (Fig. 2c) peak can be fitted to three peaks. The peak located at 284.2 eV was assigned to C]C bonds and the peaks at 285.2 and 288.9 eV were assigned to C–O and COOR respectively caused by oxygen adsorbed on the surface of carbon substrate and not completely detached carboxyl group [41]. Three peaks in O 1s (Fig. 2d) at 529, 531 and 532 eV 4

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Fig. 2. (a) XPS survey spectrum of TiO2@C; High-resolution (b) Ti 2p, (c) C 1s, and (d) O 1s XPS spectra of TiO2@C.

The large irreversible capacities for the first cycle arise from the formation of SEI film and the irreversible trapping of Na+ in TiO2@C. From the second cycle, the capacity of TiO2@C overlapped well, indicating good reversibility. A capacity of 190.3 mA h g−1 was retained after 50 cycles, corresponding to a capacity loss of 3.7% (against the second cycle). Fig. 5c depicts the rate performance of TiO2@C and pure TiO2 NWs anode in SIBs. Specific discharge capacities of 195, 176, 142, 110, 95, 75, 60, 42 and 24 mA h g−1 were measured at 0.15, 0.3, 0.5, 1.5, 3, 5, 10, 15 and 30 C, respectively. Compared with the rate data of pure TiO2 NWs anode, it shows an increase of more than 120% in specific capacity. The fluctuation of capacity was extremely small and remained around 125 mA h g−1 within 2500 cycles at 3 C (Fig. 5d). Fig. 6 displays the EIS data of TiO2@C electrodes in SIB before and after the CV test. In the corresponding equivalent circuit in the inset of Fig. 6a and b, Re is the internal impedance, Rct is the charge transfer resistance, and Rw is the Warburg impedance related to Na+ diffusion between TiO2@C electrode and electrolyte. Q represents the corresponding capacitance [53]. Different from in LIB, semicircle diameter in the high-frequency region for TiO2@C electrodes before the CV test is much higher than that after the loop, which could be due to the noncomplete wetting of the active substance by electrolyte and not activated electrode material before the cycle. After the initial several cycles, the larger sodium ions irreversibly expanded the microstructure of the material for subsequent insertion. The EIS data in Fig. 6b corresponded to the TiO2@C electrode cycled at 100, 500 and 2,000 times in long charge and discharge cycling at 3 C, respectively. With proceeding the cycle, the charge transfer resistances increased, resulting in the growth of Rct from 59.39 to 234 Ω, then to 284.3 Ω. The increase was caused by the continuously densified SEI layer, and the destruction of surface structure of the electrode at the cycling, which obstructed the electron transmission.

region of the semicircle represents the SEI layer resistance (Rs) and capacitance (C1), while middle-frequency semicircle is associated with the charge transfer resistance (Rct) and the double-layer capacitance (C2) [41]. The slope line in low frequency region is related to the Warburg impedance (Zw) of Li ion diffusion [48]. Semicircle diameter in the high-frequency region for TiO2@C electrodes before and after the CV test was 57.51 Ω and 77.63 Ω respectively, which suggested that the charge transfer resistance of the TiO2@C electrodes was small and increased slightly during the formation of the SEI film. Fig. 4b shows the EIS data for the TiO2@C electrode cycled 100, 500 and 2000 times at 3 C, respectively. The semicircle diameters of the electrode after 100 and 500 cycles remain unchanged. After 2, 000 cycles, the semicircle diameter was doubled, which is possibly due to the surface passivation, morphological change and continuously densified SEI layer of the TiO2@C that blocked the charge transport [49]. Compared to pure TiO2, the Rct of the TiO2@C electrode is simulated to be 57.51 Ω (Fig. S6), which is almost one quarter of pure TiO2 NWs, i.e., 214.3 Ω. It is mainly ascribable to the carbon substrates with good conductivity closely combined with poorly conductive TiO2, which accelerates the speed of electron transport. In case of SIBs, Fig. 5a shows the CV curves of the TiO2@C in the potential window of 0.1–2.5 V vs. Na+/Na. The irreversible broad peak located at 0.75 V was due to the generating of SEI film [50]. There were no obvious peaks but a long sloping curve above 0.3 V, owing to the presence of TiO2 NWs and amorphous carbon. The second and third cycles of TiO2@C overlapped well, indicating good reversibility. Fig. 5b shows the first charge-discharge profiles and cycle performance of the TiO2@C electrode between cutoff voltages of 0.01–2.5 V at 0.15 C. In the first and second discharge processes, TiO2@C delivers high specific capacities of 528.3 and 197.5 mA h g−1 respectively, which is comparable to the TiO2/carbon based anode materials reported [11,51,52]. 5

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Fig. 3. In LIBs: (a) cyclic voltammetry curves of the TiO2@C electrode for the first 3 cycles; (b) charge-discharge profiles of the TiO2@C electrode at different cycles at a current density of 0.3 C; (c) rate capabilities of the TiO2@C electrode and pure TiO2 NWs at current densities of 0.3 C–30 C; (d) cycling performance and coulombic efficiency of the TiO2@C electrode at a current of 3 C.

discharge platform at 0.6/0.3 V TiO2@C has a charge-discharge profile similar to TiO2 NWs, yet the curve is smoother and the slope is smaller. The specific capacity of the TiO2@C electrode was up to 289 and 197 mA h g−1 for LIBs/SIBs. Based on the above findings, the longcycling performance of TiO2@C and pure carbon substrate in SIBs was compared (Fig. S9). The incorporation of a small amount of TiO2 NWs (12 wt%) plays an important role in improving the electrochemical performance of carbon materials. The catalytic performance of TiO2–B and the conductivity of carbon matrix were fully utilized. To study the kinetics of the TiO2@C electrode, CV tests were conducted at different rates from 0.1 to 1.5 mV s−1. Typically, the capacity consists of diffusion and capacitive contributions which originated from

From the comparison of the EIS data between TiO2@C electrode and pure TiO2 NWs electrode (Fig. S7), the Rct is 98.46 and 362.81 Ω respectively. The significant reduction of Rct was due to the combination of amorphous carbon and TiO2 NWs which greatly enhances the transmission efficiency of electrons, therefore led to improved electrochemical performances. We tested the stabilized charge-discharge profiles of TiO2@C, TiO2 NWs and pure carbon matrix in LIBs and SIBs at 0.3 C and 0.15 C respectively (Figs. S8a and S8b). For LIB, TiO2 NWs and pure carbon matrix have a similar specific capacity of 150 mA h g−1 without obvious charge/discharge platform. Concerning SIB, pure carbon matrix can only provide a specific capacity of less than 30 mA h g−1 while TiO2 NWs has a capacity of 138 mA h g−1 with a charge/

Fig. 4. (a) Electrochemical impedance spectra and the equivalent circuit of the TiO2@C electrode (Re: resistance of the electrolytes; Rs: SEI layer resistance; C1: SEI layer capacitance; Rct: charge transfer resistance; Rw: Warburg resistance; C2: constant phase element); (b) EIS data of the TiO2@C electrode cycled 100, 500 and 2000 times in long charge and discharge cycling at a current densities of 3 C. 6

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Fig. 5. (a) In SIBs: Cyclic voltammetry curves of the TiO2@C electrode for the first three cycles; (b) Charge-discharge profiles of the TiO2@C electrode at different cycles at a current density of 0.15 C; (c) Rate capabilities of the TiO2@C electrode at current densities of 0.15 C–30 C; (d) Cycling performance and coulombic efficiency of the TiO2@C electrode at a current density of 3 C.

electrode was controlled by surface capacitive contribution and diffusion-controlled redox reaction. The current response at a given potential can be divided into a surface capacitive contribution (k1v1/2) and diffusion-controlled contribution (k2v), based on Equation (3&4) [48,49]:

surface charge-transfer and lithiation/delithiation reaction [18,54,55]. According to the power law, the current response (i) and sweep speed (v) obey a relationship of Equations (1) and (2): i = avb

(1)

logi = blogv + loga

(2)

i = k1v1/2 + k2v

where a and b are constants. While b-value reaches 0.5, the electrochemical reaction is diffusion-controlled; As b-value reaches 1.0, the capacitive process dominates [56]. The b-values for cathodic and anodic sweep at different potentials were calculated based on the logνlogi plots (Figs. S10a and S10b). the b-values varied from 0.58 to 0.98 (Fig. 7b), indicating that the electrochemical process of the TiO2@C

1/2

i/v

= k1 + k2v

(3) 1/2

(4)

where k1 and k2 are constants. According to the trend of the capacitance contribution in Fig. 7c, the proportion of capacitance was enlarged as the scan rate increased. The results showed that the capacitive contribution maintained an advantage at higher sweep speeds during

Fig. 6. (a) Electrochemical impedance spectra and the equivalent circuit of the TiO2@C electrode (Re: resistance of the electrolytes; Rct: charge transfer resistance; Rw: Warburg resistance; Q: constant phase element). Before and after CV test; (b) EIS data of the TiO2@C electrode cycled 100, 500 and 2000 time in long charge and discharge cycling at a current density of 3 C. 7

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Fig. 7. Kinetic analysis of the electrochemical behaviors of the TiO2@C electrode in LIBs: (a) CV curves at different scan rates in the range of 0.01–3.0 V; (b) b-values plotted against different potentials. The ratio of capacitive and diffusion contribution at different scan rates; (c) Contribution ratio of the capacitive and diffusion controlled charge versus scan rate. (d) The capacitive and diffusion contribution at a scan rate of 0.5 mV s−1.

diffusion during the charge/discharge process. Carbon matrix with its excellent electrical conductivity improved the electronic conductivity (Figs. S6 and S7) and the structural stability of the prepared electrodes, meanwhile, the unique carbon matrix structure can further release the volume changes of active electrodes. Moreover, the ultrafine nanowire morphology of TiO2 in carbon matrix can ensure a high rate capability due to the reduced distance of ion and electron transport. Furthermore, the high ratio of the capacitive contribution allows fast uptake and release of Li+/Na+ with little degradation of the electrode structure, which is responsible for the excellent rate capability and long-term cycling stability. In a word, the advantages of the TiO2@C anode in composition and structure greatly improved the reaction kinetics and stability during the electrochemical process, facilitating the superior rate and long-cycle performance.

charge/discharge process. Moreover, by calculating the value of k1 and k2 at different potentials, the proportion of capacitive contribution in the whole process can be obtained. For instance, at a scan rate of 0.5 mV s−1, the charge generated by the capacitive process accounts for 60.18% of the total charge (Fig. 7d), suggesting most charge storage in TiO2@C from the capacitive mechanisms. The identical test was performed on the logν-logi plots of SIB (Figs. S11a and S11b). The b-value located from 0.67 to 0.86 (Fig. 8b) indicated that the capacity of the SIBs was also composed of surface capacitive contribution and diffusion controlled contribution. With the increase of scan rate in Fig. 8c, the portion of capacitive contribution increased. A 76.25% ratio of the total charge generated by the capacitive process was determined at a scan rate of 0.5 mV s−1 (Fig. 8d). The higher capacitive contribution is owing to the short Li/Na ion diffusion length and increased electron transfer, being favorable for fast charge storage and long-term cyclability. The TiO2@C anode shows superior high rate cycle performance in both lithium-ion and sodium-ion battery systems, which is mainly due to its unique structure. Firstly, TiO2 NWs forming a network structure in the carbon matrix feature excellent chemical stability, structural stability, and extremely low volume effect during Li and Na ions insertion/ extraction [4,57–59]. Secondly, Nanoporous structure with diameters of ca. 500 nm and mesopores as supported by Fig. 1a and b and S5b not only alleviated mechanical stress to maintain the structural integrity of the electrode, but also supplied a large specific surface area of 280.8 m2 g−1 with sufficient active sites, which contributed to the large specific capacity and cycling performance. In addition, the large amount of nanopores of carbon matrix increased the contact area between electrolyte and electrode, which contributed to the facile diffusion of electrolyte through overall electrode and facilitated the ion

4. Conclusions In conclusion, a porous carbon nanocomposite with lowly loaded TiO2 NWs was successfully synthesized using a gel-carbonized method. This composite anode material could reduce the length of the diffusion path for lithium/sodium ions by embedding TiO2 NWs in the carbon matrix. Nanoporous structure with diameters of ca. 500 nm and mesopores were formed and provided a large specific surface area, which contributed to superior lithium/sodium storage and rate performance. Combined with the ultrafine nanowire structure of TiO2–B and enhanced the conductivity of electron, the obtained TiO2@C material delivered a reversible specific capacity of 286 mA h g−1 at 0.3 C in LIB and 197 mA h g−1 at 0.15 C in SIBs. In both cases, the stability remained close to 100% after 2500 cycles as a proof of superior durable TiO2@C anode materials. By calculation, a higher capacitive 8

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Fig. 8. Kinetic analysis of the electrochemical behaviors of the TiO2@C electrode in SIBs: (a) CV curves at different scan rates in the range of 0.01–2.5 V; (b) b-values plotted against different potentials. The ratio of capacitive and diffusion contribution at different scan rates; (c) Contribution ratio of the capacitive and diffusion controlled charge versus scan rate. (d) The capacitive and diffusion contribution at a scan rate of 0.5 mV s−1.

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Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements This work was supported by Qilu Young Scholar Program in Shandong University, Open Program in Tsinghua University State Key Laboratory of New Ceramic and Fine Processing (KF201814, KF201805) and Open Program in Guangxi Key Laboratory of Information Materials (171002-K).

Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.ceramint.2019.12.161. 9

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