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Surface-amorphized TiO2 nanoparticles anchored on graphene as anode materials for lithium-ion batteries
Journal of Power Sources 397 (2018) 162–169
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Surface-amorphized TiO2 nanoparticles anchored on graphene as anode materials for lithium-ion batteries
T
C. Li, M. Zhao, C.N. Sun, B. Jin, C.C. Yang∗, Q. Jiang∗∗ Key Laboratory of Automobile Materials (Jilin University), Ministry of Education, School of Materials Science and Engineering, Jilin University, Changchun, 130022, China
H I GH L IG H T S
G R A P H I C A L A B S T R A C T
novel strategy has been proposed to • Afabricate the SA-TiO /RGO hybrid. hybrid shows superior rate and • The cycling performances. mechanisms behind remarkable • The electrochemical properties are re2
vealed.
A R T I C LE I N FO
A B S T R A C T
Keywords: Surface amorphous structure TiO2 nanoparticles Graphene Hybrid materials Lithium-ion batteries
One demanding challenge in the development of high-power batteries is to achieve superior rate and cycling performances. Here, we report a two-step method to fabricate a hybrid material of surface-amorphized TiO2/ reduced graphene oxide (SA- TiO2/RGO), which exhibits superior electrochemical properties as an anode material in lithium-ion batteries, high rate property of 135.6 mAh g−1 at a current density of 10 A g−1 and cycling stability of 98 mAh g−1 after 2000 cycles at 5 A g−1. These excellent performances originate from the unique three-dimensional integrated structure of SA-TiO2/RGO, which provides: (i) a conductive substrate of RGO; (ii) higher conductivity of SA-TiO2 relative to crystallized-TiO2; (iii) ultra-small SA-TiO2 nanoparticles; and (iv) excellent contact between SA-TiO2 and RGO through C-O-Ti bonds.
1. Introduction The development of high-power lithium-ion batteries (LIBs) for applications in electric vehicles, power tools, military devices, etc., has attracted great attentions in recent years [1–3]. Numerous efforts have been devoted to develop electrode materials with excellent rate performance and cycling stability. Among various anode materials in LIBs, TiO2 has drawn intensive interest thanks to its high abundance, nontoxicity, high activity, and small volume expansion upon Li+ insertion/ extraction [4,5]. However, the electrochemical performance of TiO2 is hindered by low theoretical capacity, Li+ mobility and electronic conductivity, as well as severe agglomeration after long-term lithiation/
∗
Corresponding author. Corresponding author. E-mail addresses: [email protected] (C.C. Yang), [email protected] (Q. Jiang).
∗∗
https://doi.org/10.1016/j.jpowsour.2018.07.019 Received 19 April 2018; Received in revised form 20 June 2018; Accepted 4 July 2018 0378-7753/ © 2018 Elsevier B.V. All rights reserved.
delithiation cycles [6]. To solve the above problems, great efforts have been implemented to construct advanced TiO2 nanostructures and also their hybrid materials. The low Li+ mobility has been addressed by reducing the length of diffusion channels through mesoporous [7] or nanostructured materials [8,9]; or by enhancing ion transport capacity with less-dense structural phases [10,11]. The electronic conductivity has been improved by doping TiO2 with Nb [12]; or by coating TiO2 with conductive mediums [13,14], such as two-dimensional reduced graphene oxide (RGO) with merits of large specific surface area, and superior electrical conductivity, structural flexibility and assembly properties. Recently, it has been demonstrated that amorphous/disordered
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The SA-TiO2/RGO composite was obtained via a hydrothermal method. It should be noted that the supercritical water can exhibit strong reducing power during hydrothermal reaction. Therefore, the hydrothermal route is an efficient method for the deoxygenation reduction of GO to RGO [19–21]. In this work, graphene oxide (GO) was prepared through a modified Hummers' method. 0.2 g SA-TiO2 was added into 80 mL GO solution (0.57 mg/mL) and thus the mass ratio of TiO2/GO is about 4.39. After stirring for 2 h, the solution was transferred into a 100-mL Teflon-sealed autoclave and maintained at 150 °C for 5 h. The fabricated composite was recovered by centrifugation, rinsed with deionized water for several times, and fully dried in vacuum at 60 °C for 12 h, obtaining the SA-TiO2/RGO composite.
materials could form percolation pathways via the opening of active diffusion channels, which not only effectively enhances mass transport within these materials, but also potentially facilitates electron and ion diffusion [15,16]. For example, Guo et al. [17] synthesized a surfaceamorphized (SA) TiO2@graphene hybrid by an ultraviolet (UV)-assisted reduction technique, which reserves the merits of surface amorphous structure, high-conductivity matrices, and well-dispersed TiO2 nanoparticles. Thus, such a structure exhibits excellent electrochemical properties as an anode material in LIBs, a reversible capacity of 108 mAh g−1 after 1500 cycles at a high rate of 3.36 A g−1. However, the (UV)-assisted reduction technique cannot combine TiO2 and RGO tightly, and the electrochemical properties of the hybrid materials need further improvement. In this work, we fabricated a hybrid material by anchoring SA-TiO2 nanoparticles on RGO with a facial hydrothermal method. Compared with UV-assisted reduction technique, TiO2 nanoparticles and RGO could combine tightly under the hydrothermal condition and the asfabricated hybrid material shows superior rate performance (135.6 mAh g−1 at 10 A g−1) and cycling stability (98 mAh g−1 after 2000 cycles at 5 A g−1) in LIBs. The formation mechanism of SA-TiO2 and the structure-property relationship of the hybrid material have also been elaborately discussed.
2.2. Materials characterization The morphology and microstructure of the samples were characterized by field-emission scanning electron microscopy (FESEM, JSM6700 F, JEOL, 15 keV) and transmission electron microscopy (TEM, JEM-2100 F, JEOL, 200 keV). X-ray diffraction (XRD) was carried out on a D/max2500pc diffractometer using Cu-Kα radiation. X-ray photoelectron spectroscopy (XPS) analysis was performed on an ESCALAB 250 spectrometer (Thermo Fisher Scientific, UK) by using a monochromatic Al-Kα (1486.6 eV) source. Raman spectra were collected using a micro-Raman spectrometer (Renishaw) with a 532-nm laser. Thermogravimetric analysis (TGA) was performed in air at a heating rate of 10 °C min−1 over a temperature range of 50–800 °C using a SDT Q600 instrument.
2. Experimental section 2.1. Synthesis of SA-TiO2 and SA-TiO2/RGO The SA-TiO2/RGO hybrid material was synthesized by using a twostep method. The corresponding procedure is schematically illustrated in Fig. 1. In the first step, SA-TiO2 is obtained from hydrolysis of titanium glycolate balls, which were synthesized via the reaction of tetrabutyl titanate (TBOT) with ethylene glycol (EG) [17,18]. The combination of SA-TiO2 and GO was achieved in the second step by using a hydrothermal method, which resulted in further crystallization of SATiO2 and reduction of GO. The experimental details are given below. 4 mL tetrabutyl titanate (TBOT) and 20 mL ethylene glycol (EG) were mixed and stirred at room temperature for 12 h. Then, the above solution was added into 200 mL acetone and kept for 2 h. A white precipitate of titanium glycolate sphere was obtained by centrifugation, which was washed several times with deionized water and ethanol. After that, the white precipitate was added into the mixture of 40 mL isopropanol and 20 mL H2O. Here, isopropanol is beneficial for controlling the hydrolysis rate and preventing the precipitation of TiO2. Under the condition of oil bath, the above mixture was stirred for 10 h at 65 °C. The as-obtained sample was washed with deionized water and ethanol for several time, and dried in an oven at 60 °C for 12 h, obtaining SA-TiO2. The SA-TiO2 was annealed in air at 600 °C for 1 h to obtain crystallized-TiO2 (C-TiO2).
2.3. Electrode preparation and electrochemical measurement The working electrodes were fabricated by mixing 80% active materials, 10% conductive agent (Super P) and 10% binder polyvinylidene difluoride (PVDF) using N-methyl-2-pyrrolidone (NMP) as a solvent. The obtained slurry was pasted on Cu foil and dried in a vacuum oven at 110 °C for 12 h. To perform the electrochemical measurements for Liion battery, coin-type cells (CR2016) were assembled in an argon-filled glove box ([O2] < 1 ppm, [H2O] < 1 ppm) using as-synthesized materials as electrode active materials, lithium foil as a counter/reference electrode and Celgard 2500 membrane as the separator. The electrolyte was 1 M LiPF6 dissolved in the mixture of ethylene carbonate, dimethyl carbonate and ethyl methyl carbonate with a volume ratio of 1:1:1. The galvanostatic discharge/charge curves were measured using an LAND CT2001A battery testing system in a voltage range of 1.0–3.0 V. Cyclic voltammetry (CV) measurements were conducted using an IVIUM electrochemical analyzer with a potential scan rate of 0.2 mV s−1 between 1.0 and 3.0 V. The electrochemical impedance spectroscopy (EIS) measurements were carried out by applying an amplitude of 10 mV over the frequency range of 100 kHz to 10 mHz. 3. Results and discussion 3.1. Characterization of hybrids Fig. 2a compares XRD patterns of SA-TiO2 and SA-TiO2/RGO. The pattern of SA-TiO2 shows typical amorphous feature with a broad peak at 25°, implying ultra-small and low-crystallized TiO2 nanoparticles. After combining with GO under hydrothermal conditions, the fraction of the crystalline phase increases [22] as demonstrated by the XRD pattern of SA-TiO2/RGO in the figure, where the diffraction peaks at 25.3°, 37.8°, 47.9°, 53.9°, 54.9°, 62.7° and 74.0° correspond well to the (101), (004), (200), (105), (211), (204) and (215) planes of TiO2, which look similar to those of crystalline TiO2 (C-TiO2, see Fig. S1) with an anatase structure (space group: I41/amd (141), JCPDS No. 21-1272). Moreover, no obvious peak of RGO can be observed in Fig. 2a for SATiO2/RGO, which indicates that the anchored TiO2 nanoparticles prevent RGO from restacking [23]. The existence of RGO was confirmed by
Fig. 1. Schematic illustration of the preparation procedure of SA-TiO2/RGO. 163
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Fig. 2. (a) XRD patterns of SA-TiO2/RGO and SA-TiO2. (b) Raman spectrum of SA-TiO2/RGO. (c) Raman spectra of GO, RGO and SA-TiO2/RGO. (d) N2 adsorption and desorption isotherms (inset: pore size distribution) of SA-TiO2/RGO.
condensation takes place, leading to a sharp adsorption volume increase. N2 desorption, as the opposite of the adsorption, happens after saturation is reached. For mesopores, the evaporation usually takes place at a pressure lower than that of capillary condensation, resulting in a hysteresis as shown in Fig. 2d [30]. Moreover, the SA-TiO2/RGO hybrid exhibits a high Brunauer-Emmett-Teller (BET) specific surface area of 170.4 m2 g−1. The pore size distribution (the inset of the figure) calculated by the Barret-Joyner-Halenda (BJH) method shows a peak at around 7.2 nm. These small mesopores in the hybrid material could provide pathways for electrolyte/Li+ transport and accommodate large volume expansion/contraction of active materials during cycling [31]. Fig. 3a shows the survey spectrum of X-ray photoelectron spectroscopy (XPS) measurement results for SA-TiO2/RGO, where O, Ti and C peaks are observed. Fig. 3b plots high-resolution XPS spectra of O 1s, where four peaks located at 529.2, 530.7, 531.4 and 532.5 eV are attributed to Ti-O bond, C-O-Ti bond, adsorbed water and oxygen vacancy, respectively [32–34]. This confirms the strong interaction between SA-TiO2 and RGO through the formation of the C-O-Ti covalent bond, which prevents their self-agglomeration. As shown in Fig. 3c, the binding energies of Ti3+ 2p1/2, Ti4+ 2p1/2, Ti3+ 2p3/2 and Ti4+ 2p3/2 are 457.9, 458.9, 463.6 and 464.6 eV, respectively, indicating the presence of valence states of Ti3+ and Ti4+ in SA-TiO2/RGO [35]. The integral area ratio of Ti3+ to Ti4+ is 11.7:1, which implies that the surface of SA-TiO2/RGO contains a large number of oxygen vacancies. For C 1s spectra in Fig. 3d, the peak at 284.5 eV was assignable to the sp2 carbon species while the peaks at higher binding energies (286, 287.3 and 288.3 eV) were ascribed to oxygenated carbon species, such as hydroxyl, carboxyl and epoxide species [20,36]. Note that the characteristic peaks of oxygenated carbon are not strong, indicating that GO was reduced to RGO [20,36]. In order to confirm this point, we
the Raman spectrum of SA-TiO2/RGO in Fig. 2b, where two typical peaks at 1351 and 1585 cm−1 are assigned to the D band and G band, respectively, of carbon. Other characteristic peaks at 148, 399, 518 and 639 cm−1 correspond to the Eg(1), B1g(1), A1g + B1g(2) and Eg(2) modes of anatase, respectively [24]. Fig. 2c compares Raman spectra of GO, RGO and SA-TiO2/RGO. The calculated ID/IG value of RGO is 0.93, which is larger than that of GO (= 0.87), indicating the increase of topological defects and disorders in RGO [19]. For SA-TiO2/RGO, the ID/IG value further increases to 0.96, which implies that the surface oxygen defect of SA-TiO2 was filled by O atoms of GO [25], increasing the reduction degree of RGO. Furthermore, the hydrothermal condition is beneficial for the formation of covalent C-O-Ti bonds, which not only combine SA-TiO2 and RGO closely, but also accelerate the electron transport between them [26]. The content of SA-TiO2 in the hybrid material is determined by TGA in air (see Fig. S2). It has been demonstrated that RGO begins to decompose at 200-250 °C [27]. As shown in Fig. S2, there is a platform at about 230 °C, which should be assigned to the start of the decomposition of RGO. Moreover, it was reported that the water contained in the pores of the material could remain up to 250 °C [28]. Based on the above considerations, the small mass loss (6.5 wt%) for SA-TiO2/RGO below 230 °C can be attributed to the evaporation of absorbed solvent. The gradual weight loss (10.7 wt %) beginning at 230 °C is assigned to the decomposition of RGO. Thus, the content of SA-TiO2 in the hybrid is approximately 82.8%. The N2 absorption/desorption isotherms in Fig. 2d show porous characteristic of SA-TiO2/RGO, which display typical type-IV behavior with hysteresis loops in the relative pressure range of 0.45–1 [29]. This implies that a large number of mesopores exist in the samples. For type-IV isotherm, a monolayer of adsorbed N2 is formed at low relative pressures while at high relative pressures a multilayer adsorption occurs until 164
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Fig. 3. (a) XPS survey scan of SA-TiO2/RGO. (b) O 1s spectrum of SA-TiO2/RGO. (c) Ti 2p spectrum of SA-TiO2/RGO. (d) C 1s spectrum of SA-TiO2/RGO.
plane of TiO2 [17]. The size of TiO2 nanoparticles (≈7 nm) in SA-TiO2/ RGO increases slightly compared with those in SA-TiO2 (< 7 nm), leading to a decrease of the amorphous layer thickness. This is because SA-TiO2 could further crystallize onto existing anatase particles during the hydrothermal process [39].
also measured the C 1s spectrum of GO (see Fig. S3). Obviously, the characteristic peaks of oxygenated carbon in SA-TiO2/RGO are much weaker than those in GO, confirming that most of the epoxide, hydroxyl and carboxyl functional groups were successfully removed. To further illustrate the reduction degree of GO, the peak area ratios of oxygencontaining bonds (such as -C-O, -C=O and O-C=O) to C=C/C-C bonds are calculated and the corresponding results are shown in Table S1. It is evident that the amounts of oxygen-containing groups in SA-TiO2/RGO decrease significantly, which further validates the reduction of GO in the SA-TiO2/RGO hybrid. The morphologies of as-fabricated samples are characterized by field-emission scanning electron microscopy (FESEM) and transmission electron microscopy (TEM). Fig. 4a and Fig. S4 show FESEM images of SA-TiO2 and C-TiO2 spheres, respectively. It is clear that SA-TiO2 has better dispersibility than C-TiO2. Fig. 4b presents FESEM images of SATiO2/RGO, where TiO2 nanospheres are anchored on the surface of RGO. Fig. 4c shows TEM images of SA-TiO2. The amorphous layer with a thickness of about 2.2 nm on the surface of SA-TiO2 can be confirmed by selected area electron diffraction (SAED). The TEM image of SATiO2/RGO (see Fig. 4d) shows that the SA-TiO2 is an aggregation of individual fine nanosized crystallites. The channels formed between particles by the hydrolysis of titanium glycol play a capillary role, which is more conducive to the penetration of electrolyte [18,37]. The average surface layer thickness of TiO2 in SA-TiO2/RGO is about 1.2 nm (see the inset of the figure). The TiO2/TiO2−x core/shell forms an n-i heterojunction at the amorphous state/crystalline state interface, which induces a built-in electric field across the interface, reducing the charge transport resistance significantly [38]. Fig. 4e and f show high-resolution TEM (HRTEM) images of SA-TiO2 and SA-TiO2/RGO, respectively, in which the interplanar spacing of 0.35 nm corresponds to the (101)
3.2. Electrochemical performances in LIBs The Li-storage properties of C-TiO2, SA-TiO2 and SA-TiO2/RGO were evaluated by assembling them into coin-type half cells for electrochemical tests. Fig. 5a presents the cyclic voltammetry (CV) curves of the SA-TiO2/RGO hybrid electrode for initial five cycles. For the first cathodic scan, the peaks at 1.71 V and 1.5 V are ascribed to the transitions of TiO2 → Li0.55TiO2 and Li0.55TiO2 → LiTiO2, respectively [40,41]. It has been demonstrated that the potentials of Li are different for extraction from amorphous and crystalline TiO2 [42]. The anodic peak at 1.98 V is associated with Li extraction from amorphous anatase (amorphous TiO2 on the surface), which is consistent with the reported result (1.9–2.1 V) in other literature [43]. The other anodic peak at 2.2 V is assigned to the deinsertion of Li from the crystalized TiO2 (CTiO2 in the inner) [44]. In order to confirm this point, we also measured the CV curve of C-TiO2 (see Fig. S5), where the peak is located at 2.2 V. Moreover, as shown in Fig. 5a and Fig. S5, a slight peak shift from 2.2 V to 2.1 V is observed during the first few cycles, which is caused by irreversible reactions [45]. Moreover, the CV curves almost coincide from the third cycle, suggesting excellent cycling stability of the SATiO2/RGO electrode. Fig. 5b shows the galvanostatic charge/discharge curves of SA-TiO2/RGO for the 1st, 2nd, 50th and 100th cycles at a current density of 0.2 A g−1. Note that the active materials are SA-TiO2/ RGO hybrid and all capacities reported in this work are calculated 165
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Fig. 4. (a) SEM image of SA-TiO2. (b) SEM image of SA-TiO2/RGO. (c) TEM image of SA-TiO2. (d) TEM image of SA-TiO2/RGO. (e) HRTEM image of SA-TiO2. (f) HRTEM image of SA-TiO2/RGO.
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Fig. 5. (a) CV curves of SA-TiO2/RGO electrode at a scan rate of 0.2 mV s−1. (b) Galvanostatic discharge/charge curves of SA-TiO2/RGO electrode at 0.2 A g−1. (c) Cycling performances of SA-TiO2, C-TiO2 and SA-TiO2/RGO electrodes at 0.2 A g−1. (d) Rate properties of SA-TiO2, C-TiO2 and SA-TiO2/RGO at different current densities. (e) Cycling performance and coulombic efficiency of SA-TiO2/RGO electrode at 5 A g−1.
of region C and thus higher capacity of SA-TiO2/RGO than the theoretical capacity. Moreover, the RGO electrode has low specific capacity and electrochemical reactivity in the voltage range of 1.0–3.0 V [51]. Therefore, the capacity of the SA-TiO2/RGO electrode is mainly contributed from TiO2. Fig. 5c compares the cycling stability of SA-TiO2/ RGO, SA-TiO2 and C-TiO2 electrodes at a current density of 0.2 A g−1. After 100 cycles, the SA-TiO2/RGO electrode shows a discharge capacity of 251.7 mAh g−1, which is much larger than those of SA-TiO2 (109 mAh g−1) and C-TiO2 (only 37.1 mAh g−1). Such an excellent electrochemical performance of SA-TiO2/RGO may be caused by its low internal resistance. The EIS measurement results (see Fig. S6) shows that the charge transfer resistance Rct is only 58.5 Ω for SA-TiO2/RGO, which is much lower than those of SA-TiO2 (93 Ω) and C-TiO2 (202 Ω). Fig. 5d exhibits rate performances of the C-TiO2, SA-TiO2 and SATiO2/RGO electrodes at different current densities from 0.2 to 10 A g−1. The SA-TiO2/RGO electrode delivers highly reversible average capacities of 319.3, 267.8, 235.1, 202, 159.2 and 135.6 mAh g−1 at 0.2, 0.5,
according to the total mass of SA-TiO2/RGO. The initial discharge capacity is 433.2 mAh g−1. With increasing cycle number, the discharge capacities are 363.6, 264 and 251.7 mAh g−1 for the 2nd, 50th and 100th cycles, which remain higher than the theoretical capacity (168 mAh g−1) of anatase [1]. Based on previous reports, the discharge curves for the SA-TiO2/RGO electrode can be divided into three different voltage regions, which are marked as A, B and C in Fig. 5b (here the discharge curve for the first cycle is taken as an example) [13,46,47]. In region A, a monotonic voltage drop to ∼1.75 V is observed, which is assigned to a homogeneous Li insertion into the bulk until the solid-solution limit of Li in TiO2 is reached. In region B, there is a biphase plateau at ∼1.75 V, indicating the coexistence of Li-rich and Li-poor phases. In region C, reversible storage of Li occurs at particle interfaces. The porous SA-TiO2/RGO hybrid and ultra-small SATiO2 nanoparticles with high specific surface area could readily induce the capacitive behavior during the electrochemical reaction, leading to an extra interfacial storage capacity [48–50]. This results in an increase 167
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density by Nb doping, which leads to a 3 orders of magnitude higher electronic conductivity of Ti0.94Nb0.06O2 compared with TiO2 [12]. Fig. 5e shows the long-term cycle stability of SA-TiO2/RGO electrode. The electrode was galvanostatically discharged and charged at 0.2 A g−1 for the first five cycles and then at 5 A g−1 for 2000 cycles. After 2000 cycles, the hybrid electrode still remains a capacity of 98 mAh g−1, exhibiting excellent cycling stability. Fig. 6 illustrates the mechanism behind superior electrochemical performances of SA-TiO2/RGO in LIBs. The excellent electrochemical performances of SA-TiO2/RGO originate from (i) the surface amorphous layer of TiO2, which provides plenty of oxygen vacancies for electrons/ ions transport, and alleviates the stress caused by volume changes during Li+ insertion/extraction [52]; (ii) high-conductivity RGO, which provides a conductive network, and also contributes a portion of capacity; (iii) the ultra-small SA-TiO2 nanoparticles with large surface/ volume ratio, which offers a large electrochemically active area and short ion diffusion distance; and (iv) the C-O-Ti bonds between SA-TiO2 and RGO, which reduce the inner resistance significantly.
Table 1 Comparisons of rate performances of some representative TiO2-based anodes materials for LIBs. Ti-based anodes
Rate capability
Ref.
Hollow TiO2/graphitic carbon Crystalline TiO2 hollow spheres
91 mAh g−1 at 10 A g−1 49.1 mAh g−1 at 5.19 A g−1 129 mAh g−1 at 1.68 A g−1 109 mAh g−1 at 3.74 A g−1 50 mAh g−1 at 0.168 A g−1 103 mAh g−1 at 8.4 A g−1 85 mAh g−1 at 3.36 A g−1
[53] [54]
70 mAh g−1 at 10 A g−1 120 mAh g−1 at 16.75 A g−1 135.6 mAh g−1 at 10 A g−1 121.9 mAh g−1 at 16.75 A g−1
[58] [12]
3-shelled TiO2 hollow microspheres TiO2-rGO nanocomposite hollow spheres Oxygen-deficient anatase TiO2@C Surface-amorphized TiO2@graphene Graphene supported ultralong TiO2 nanofiber Graphitic carbon-TiO2 Rutlie Ti0.94Nb0.06O2 SA-TiO2/RGO
[7] [55] [56] [17] [57]
This work
4. Conclusions In summary, a unique hybrid material of SA-TiO2/RGO has been fabricated through a two-step method, which combines SA-TiO2 and RGO tightly through the formed C-O-Ti covalent bonds during the hydrothermal reaction. As a result, such a hybrid material as an anode in Li-ion batteries exhibits superior electrochemical performances, high specific capacities (251.7 mAh g−1 after 100 cycles at 0.2 A g−1), excellent rate performance (135.6 mAh g−1 at 10 A g−1) and cycling stability (98 mAh g−1 after 2000 cycles at 5 A g−1). The developed hybrid material and its fabrication method reported in this work provide new insights into the fundamental understanding of high-performance surface amorphized metal oxides toward application in energy storage devices. Acknowledgements This project is financially supported by National Natural Science Foundation of China (Nos. 51671092 and 51631004), ChangBai Mountain Scholars Program, Natural Science Foundation of Jilin Province (No. 20160101315JC), Program for JLU Science and Technology Innovative Research Team (No. 2017TD-09) and the Fundamental Research Funds for the Central Universities. Appendix A. Supplementary data
Fig. 6. Schematic illustration of the mechanism behind superior electrochemical performances of SA-TiO2/RGO in LIBs.
Supplementary data related to this article can be found at http://dx. doi.org/10.1016/j.jpowsour.2018.07.019.
−1
1, 2, 5 and 10 A g , respectively, showing the highest capacity at each current density. When the current density comes back to 0.2 A g−1, the SA-TiO2/RGO electrode resumes a capacity of 261.9 mAh g−1 rapidly and then retains stable for further cycling. It is noteworthy that the capacity of SA-TiO2/RGO (135.6 mAh g−1) is 15.4 times that of C-TiO2 (8.8 mAh g−1) at 10 A g−1, displaying superior rate performance compared with other TiO2-based anode materials (see Table 1). It should be noted that a recent report indicates that the rutile Ti0.94Nb0.06O2 shows excellent rate property in Li-ion battery, for example, a high capacity of 120 mAh g−1 at a current density of 16.75 A g−1 [12]. For a comparison, we have measured the capacity of SA-TiO2/RGO at 16.75 A g−1, which is 121.9 mAh g−1, very similar to that of rutile Ti0.94Nb0.06O2 (120 mAh g−1). The superior rate properties of SA-TiO2/RGO can be attributed to: (i) the built-in electric field across the amorphous state/crystalline state interface of SA-TiO2, reducing the charge transfer resistance; and (ii) the excellent contact between SA-TiO2 and RGO through C-O-Ti bonds, decreasing the inner resistance. For rutile Ti0.94Nb0.06O2, its excellent rate performance originates from the significant increase of the conduction electron
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Journal of Power Sources journal homepage: www.elsevier.com/locate/jpowsour
Surface-amorphized TiO2 nanoparticles anchored on graphene as anode materials for lithium-ion batteries
T
C. Li, M. Zhao, C.N. Sun, B. Jin, C.C. Yang∗, Q. Jiang∗∗ Key Laboratory of Automobile Materials (Jilin University), Ministry of Education, School of Materials Science and Engineering, Jilin University, Changchun, 130022, China
H I GH L IG H T S
G R A P H I C A L A B S T R A C T
novel strategy has been proposed to • Afabricate the SA-TiO /RGO hybrid. hybrid shows superior rate and • The cycling performances. mechanisms behind remarkable • The electrochemical properties are re2
vealed.
A R T I C LE I N FO
A B S T R A C T
Keywords: Surface amorphous structure TiO2 nanoparticles Graphene Hybrid materials Lithium-ion batteries
One demanding challenge in the development of high-power batteries is to achieve superior rate and cycling performances. Here, we report a two-step method to fabricate a hybrid material of surface-amorphized TiO2/ reduced graphene oxide (SA- TiO2/RGO), which exhibits superior electrochemical properties as an anode material in lithium-ion batteries, high rate property of 135.6 mAh g−1 at a current density of 10 A g−1 and cycling stability of 98 mAh g−1 after 2000 cycles at 5 A g−1. These excellent performances originate from the unique three-dimensional integrated structure of SA-TiO2/RGO, which provides: (i) a conductive substrate of RGO; (ii) higher conductivity of SA-TiO2 relative to crystallized-TiO2; (iii) ultra-small SA-TiO2 nanoparticles; and (iv) excellent contact between SA-TiO2 and RGO through C-O-Ti bonds.
1. Introduction The development of high-power lithium-ion batteries (LIBs) for applications in electric vehicles, power tools, military devices, etc., has attracted great attentions in recent years [1–3]. Numerous efforts have been devoted to develop electrode materials with excellent rate performance and cycling stability. Among various anode materials in LIBs, TiO2 has drawn intensive interest thanks to its high abundance, nontoxicity, high activity, and small volume expansion upon Li+ insertion/ extraction [4,5]. However, the electrochemical performance of TiO2 is hindered by low theoretical capacity, Li+ mobility and electronic conductivity, as well as severe agglomeration after long-term lithiation/
∗
Corresponding author. Corresponding author. E-mail addresses: [email protected] (C.C. Yang), [email protected] (Q. Jiang).
∗∗
https://doi.org/10.1016/j.jpowsour.2018.07.019 Received 19 April 2018; Received in revised form 20 June 2018; Accepted 4 July 2018 0378-7753/ © 2018 Elsevier B.V. All rights reserved.
delithiation cycles [6]. To solve the above problems, great efforts have been implemented to construct advanced TiO2 nanostructures and also their hybrid materials. The low Li+ mobility has been addressed by reducing the length of diffusion channels through mesoporous [7] or nanostructured materials [8,9]; or by enhancing ion transport capacity with less-dense structural phases [10,11]. The electronic conductivity has been improved by doping TiO2 with Nb [12]; or by coating TiO2 with conductive mediums [13,14], such as two-dimensional reduced graphene oxide (RGO) with merits of large specific surface area, and superior electrical conductivity, structural flexibility and assembly properties. Recently, it has been demonstrated that amorphous/disordered
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The SA-TiO2/RGO composite was obtained via a hydrothermal method. It should be noted that the supercritical water can exhibit strong reducing power during hydrothermal reaction. Therefore, the hydrothermal route is an efficient method for the deoxygenation reduction of GO to RGO [19–21]. In this work, graphene oxide (GO) was prepared through a modified Hummers' method. 0.2 g SA-TiO2 was added into 80 mL GO solution (0.57 mg/mL) and thus the mass ratio of TiO2/GO is about 4.39. After stirring for 2 h, the solution was transferred into a 100-mL Teflon-sealed autoclave and maintained at 150 °C for 5 h. The fabricated composite was recovered by centrifugation, rinsed with deionized water for several times, and fully dried in vacuum at 60 °C for 12 h, obtaining the SA-TiO2/RGO composite.
materials could form percolation pathways via the opening of active diffusion channels, which not only effectively enhances mass transport within these materials, but also potentially facilitates electron and ion diffusion [15,16]. For example, Guo et al. [17] synthesized a surfaceamorphized (SA) TiO2@graphene hybrid by an ultraviolet (UV)-assisted reduction technique, which reserves the merits of surface amorphous structure, high-conductivity matrices, and well-dispersed TiO2 nanoparticles. Thus, such a structure exhibits excellent electrochemical properties as an anode material in LIBs, a reversible capacity of 108 mAh g−1 after 1500 cycles at a high rate of 3.36 A g−1. However, the (UV)-assisted reduction technique cannot combine TiO2 and RGO tightly, and the electrochemical properties of the hybrid materials need further improvement. In this work, we fabricated a hybrid material by anchoring SA-TiO2 nanoparticles on RGO with a facial hydrothermal method. Compared with UV-assisted reduction technique, TiO2 nanoparticles and RGO could combine tightly under the hydrothermal condition and the asfabricated hybrid material shows superior rate performance (135.6 mAh g−1 at 10 A g−1) and cycling stability (98 mAh g−1 after 2000 cycles at 5 A g−1) in LIBs. The formation mechanism of SA-TiO2 and the structure-property relationship of the hybrid material have also been elaborately discussed.
2.2. Materials characterization The morphology and microstructure of the samples were characterized by field-emission scanning electron microscopy (FESEM, JSM6700 F, JEOL, 15 keV) and transmission electron microscopy (TEM, JEM-2100 F, JEOL, 200 keV). X-ray diffraction (XRD) was carried out on a D/max2500pc diffractometer using Cu-Kα radiation. X-ray photoelectron spectroscopy (XPS) analysis was performed on an ESCALAB 250 spectrometer (Thermo Fisher Scientific, UK) by using a monochromatic Al-Kα (1486.6 eV) source. Raman spectra were collected using a micro-Raman spectrometer (Renishaw) with a 532-nm laser. Thermogravimetric analysis (TGA) was performed in air at a heating rate of 10 °C min−1 over a temperature range of 50–800 °C using a SDT Q600 instrument.
2. Experimental section 2.1. Synthesis of SA-TiO2 and SA-TiO2/RGO The SA-TiO2/RGO hybrid material was synthesized by using a twostep method. The corresponding procedure is schematically illustrated in Fig. 1. In the first step, SA-TiO2 is obtained from hydrolysis of titanium glycolate balls, which were synthesized via the reaction of tetrabutyl titanate (TBOT) with ethylene glycol (EG) [17,18]. The combination of SA-TiO2 and GO was achieved in the second step by using a hydrothermal method, which resulted in further crystallization of SATiO2 and reduction of GO. The experimental details are given below. 4 mL tetrabutyl titanate (TBOT) and 20 mL ethylene glycol (EG) were mixed and stirred at room temperature for 12 h. Then, the above solution was added into 200 mL acetone and kept for 2 h. A white precipitate of titanium glycolate sphere was obtained by centrifugation, which was washed several times with deionized water and ethanol. After that, the white precipitate was added into the mixture of 40 mL isopropanol and 20 mL H2O. Here, isopropanol is beneficial for controlling the hydrolysis rate and preventing the precipitation of TiO2. Under the condition of oil bath, the above mixture was stirred for 10 h at 65 °C. The as-obtained sample was washed with deionized water and ethanol for several time, and dried in an oven at 60 °C for 12 h, obtaining SA-TiO2. The SA-TiO2 was annealed in air at 600 °C for 1 h to obtain crystallized-TiO2 (C-TiO2).
2.3. Electrode preparation and electrochemical measurement The working electrodes were fabricated by mixing 80% active materials, 10% conductive agent (Super P) and 10% binder polyvinylidene difluoride (PVDF) using N-methyl-2-pyrrolidone (NMP) as a solvent. The obtained slurry was pasted on Cu foil and dried in a vacuum oven at 110 °C for 12 h. To perform the electrochemical measurements for Liion battery, coin-type cells (CR2016) were assembled in an argon-filled glove box ([O2] < 1 ppm, [H2O] < 1 ppm) using as-synthesized materials as electrode active materials, lithium foil as a counter/reference electrode and Celgard 2500 membrane as the separator. The electrolyte was 1 M LiPF6 dissolved in the mixture of ethylene carbonate, dimethyl carbonate and ethyl methyl carbonate with a volume ratio of 1:1:1. The galvanostatic discharge/charge curves were measured using an LAND CT2001A battery testing system in a voltage range of 1.0–3.0 V. Cyclic voltammetry (CV) measurements were conducted using an IVIUM electrochemical analyzer with a potential scan rate of 0.2 mV s−1 between 1.0 and 3.0 V. The electrochemical impedance spectroscopy (EIS) measurements were carried out by applying an amplitude of 10 mV over the frequency range of 100 kHz to 10 mHz. 3. Results and discussion 3.1. Characterization of hybrids Fig. 2a compares XRD patterns of SA-TiO2 and SA-TiO2/RGO. The pattern of SA-TiO2 shows typical amorphous feature with a broad peak at 25°, implying ultra-small and low-crystallized TiO2 nanoparticles. After combining with GO under hydrothermal conditions, the fraction of the crystalline phase increases [22] as demonstrated by the XRD pattern of SA-TiO2/RGO in the figure, where the diffraction peaks at 25.3°, 37.8°, 47.9°, 53.9°, 54.9°, 62.7° and 74.0° correspond well to the (101), (004), (200), (105), (211), (204) and (215) planes of TiO2, which look similar to those of crystalline TiO2 (C-TiO2, see Fig. S1) with an anatase structure (space group: I41/amd (141), JCPDS No. 21-1272). Moreover, no obvious peak of RGO can be observed in Fig. 2a for SATiO2/RGO, which indicates that the anchored TiO2 nanoparticles prevent RGO from restacking [23]. The existence of RGO was confirmed by
Fig. 1. Schematic illustration of the preparation procedure of SA-TiO2/RGO. 163
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Fig. 2. (a) XRD patterns of SA-TiO2/RGO and SA-TiO2. (b) Raman spectrum of SA-TiO2/RGO. (c) Raman spectra of GO, RGO and SA-TiO2/RGO. (d) N2 adsorption and desorption isotherms (inset: pore size distribution) of SA-TiO2/RGO.
condensation takes place, leading to a sharp adsorption volume increase. N2 desorption, as the opposite of the adsorption, happens after saturation is reached. For mesopores, the evaporation usually takes place at a pressure lower than that of capillary condensation, resulting in a hysteresis as shown in Fig. 2d [30]. Moreover, the SA-TiO2/RGO hybrid exhibits a high Brunauer-Emmett-Teller (BET) specific surface area of 170.4 m2 g−1. The pore size distribution (the inset of the figure) calculated by the Barret-Joyner-Halenda (BJH) method shows a peak at around 7.2 nm. These small mesopores in the hybrid material could provide pathways for electrolyte/Li+ transport and accommodate large volume expansion/contraction of active materials during cycling [31]. Fig. 3a shows the survey spectrum of X-ray photoelectron spectroscopy (XPS) measurement results for SA-TiO2/RGO, where O, Ti and C peaks are observed. Fig. 3b plots high-resolution XPS spectra of O 1s, where four peaks located at 529.2, 530.7, 531.4 and 532.5 eV are attributed to Ti-O bond, C-O-Ti bond, adsorbed water and oxygen vacancy, respectively [32–34]. This confirms the strong interaction between SA-TiO2 and RGO through the formation of the C-O-Ti covalent bond, which prevents their self-agglomeration. As shown in Fig. 3c, the binding energies of Ti3+ 2p1/2, Ti4+ 2p1/2, Ti3+ 2p3/2 and Ti4+ 2p3/2 are 457.9, 458.9, 463.6 and 464.6 eV, respectively, indicating the presence of valence states of Ti3+ and Ti4+ in SA-TiO2/RGO [35]. The integral area ratio of Ti3+ to Ti4+ is 11.7:1, which implies that the surface of SA-TiO2/RGO contains a large number of oxygen vacancies. For C 1s spectra in Fig. 3d, the peak at 284.5 eV was assignable to the sp2 carbon species while the peaks at higher binding energies (286, 287.3 and 288.3 eV) were ascribed to oxygenated carbon species, such as hydroxyl, carboxyl and epoxide species [20,36]. Note that the characteristic peaks of oxygenated carbon are not strong, indicating that GO was reduced to RGO [20,36]. In order to confirm this point, we
the Raman spectrum of SA-TiO2/RGO in Fig. 2b, where two typical peaks at 1351 and 1585 cm−1 are assigned to the D band and G band, respectively, of carbon. Other characteristic peaks at 148, 399, 518 and 639 cm−1 correspond to the Eg(1), B1g(1), A1g + B1g(2) and Eg(2) modes of anatase, respectively [24]. Fig. 2c compares Raman spectra of GO, RGO and SA-TiO2/RGO. The calculated ID/IG value of RGO is 0.93, which is larger than that of GO (= 0.87), indicating the increase of topological defects and disorders in RGO [19]. For SA-TiO2/RGO, the ID/IG value further increases to 0.96, which implies that the surface oxygen defect of SA-TiO2 was filled by O atoms of GO [25], increasing the reduction degree of RGO. Furthermore, the hydrothermal condition is beneficial for the formation of covalent C-O-Ti bonds, which not only combine SA-TiO2 and RGO closely, but also accelerate the electron transport between them [26]. The content of SA-TiO2 in the hybrid material is determined by TGA in air (see Fig. S2). It has been demonstrated that RGO begins to decompose at 200-250 °C [27]. As shown in Fig. S2, there is a platform at about 230 °C, which should be assigned to the start of the decomposition of RGO. Moreover, it was reported that the water contained in the pores of the material could remain up to 250 °C [28]. Based on the above considerations, the small mass loss (6.5 wt%) for SA-TiO2/RGO below 230 °C can be attributed to the evaporation of absorbed solvent. The gradual weight loss (10.7 wt %) beginning at 230 °C is assigned to the decomposition of RGO. Thus, the content of SA-TiO2 in the hybrid is approximately 82.8%. The N2 absorption/desorption isotherms in Fig. 2d show porous characteristic of SA-TiO2/RGO, which display typical type-IV behavior with hysteresis loops in the relative pressure range of 0.45–1 [29]. This implies that a large number of mesopores exist in the samples. For type-IV isotherm, a monolayer of adsorbed N2 is formed at low relative pressures while at high relative pressures a multilayer adsorption occurs until 164
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Fig. 3. (a) XPS survey scan of SA-TiO2/RGO. (b) O 1s spectrum of SA-TiO2/RGO. (c) Ti 2p spectrum of SA-TiO2/RGO. (d) C 1s spectrum of SA-TiO2/RGO.
plane of TiO2 [17]. The size of TiO2 nanoparticles (≈7 nm) in SA-TiO2/ RGO increases slightly compared with those in SA-TiO2 (< 7 nm), leading to a decrease of the amorphous layer thickness. This is because SA-TiO2 could further crystallize onto existing anatase particles during the hydrothermal process [39].
also measured the C 1s spectrum of GO (see Fig. S3). Obviously, the characteristic peaks of oxygenated carbon in SA-TiO2/RGO are much weaker than those in GO, confirming that most of the epoxide, hydroxyl and carboxyl functional groups were successfully removed. To further illustrate the reduction degree of GO, the peak area ratios of oxygencontaining bonds (such as -C-O, -C=O and O-C=O) to C=C/C-C bonds are calculated and the corresponding results are shown in Table S1. It is evident that the amounts of oxygen-containing groups in SA-TiO2/RGO decrease significantly, which further validates the reduction of GO in the SA-TiO2/RGO hybrid. The morphologies of as-fabricated samples are characterized by field-emission scanning electron microscopy (FESEM) and transmission electron microscopy (TEM). Fig. 4a and Fig. S4 show FESEM images of SA-TiO2 and C-TiO2 spheres, respectively. It is clear that SA-TiO2 has better dispersibility than C-TiO2. Fig. 4b presents FESEM images of SATiO2/RGO, where TiO2 nanospheres are anchored on the surface of RGO. Fig. 4c shows TEM images of SA-TiO2. The amorphous layer with a thickness of about 2.2 nm on the surface of SA-TiO2 can be confirmed by selected area electron diffraction (SAED). The TEM image of SATiO2/RGO (see Fig. 4d) shows that the SA-TiO2 is an aggregation of individual fine nanosized crystallites. The channels formed between particles by the hydrolysis of titanium glycol play a capillary role, which is more conducive to the penetration of electrolyte [18,37]. The average surface layer thickness of TiO2 in SA-TiO2/RGO is about 1.2 nm (see the inset of the figure). The TiO2/TiO2−x core/shell forms an n-i heterojunction at the amorphous state/crystalline state interface, which induces a built-in electric field across the interface, reducing the charge transport resistance significantly [38]. Fig. 4e and f show high-resolution TEM (HRTEM) images of SA-TiO2 and SA-TiO2/RGO, respectively, in which the interplanar spacing of 0.35 nm corresponds to the (101)
3.2. Electrochemical performances in LIBs The Li-storage properties of C-TiO2, SA-TiO2 and SA-TiO2/RGO were evaluated by assembling them into coin-type half cells for electrochemical tests. Fig. 5a presents the cyclic voltammetry (CV) curves of the SA-TiO2/RGO hybrid electrode for initial five cycles. For the first cathodic scan, the peaks at 1.71 V and 1.5 V are ascribed to the transitions of TiO2 → Li0.55TiO2 and Li0.55TiO2 → LiTiO2, respectively [40,41]. It has been demonstrated that the potentials of Li are different for extraction from amorphous and crystalline TiO2 [42]. The anodic peak at 1.98 V is associated with Li extraction from amorphous anatase (amorphous TiO2 on the surface), which is consistent with the reported result (1.9–2.1 V) in other literature [43]. The other anodic peak at 2.2 V is assigned to the deinsertion of Li from the crystalized TiO2 (CTiO2 in the inner) [44]. In order to confirm this point, we also measured the CV curve of C-TiO2 (see Fig. S5), where the peak is located at 2.2 V. Moreover, as shown in Fig. 5a and Fig. S5, a slight peak shift from 2.2 V to 2.1 V is observed during the first few cycles, which is caused by irreversible reactions [45]. Moreover, the CV curves almost coincide from the third cycle, suggesting excellent cycling stability of the SATiO2/RGO electrode. Fig. 5b shows the galvanostatic charge/discharge curves of SA-TiO2/RGO for the 1st, 2nd, 50th and 100th cycles at a current density of 0.2 A g−1. Note that the active materials are SA-TiO2/ RGO hybrid and all capacities reported in this work are calculated 165
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Fig. 4. (a) SEM image of SA-TiO2. (b) SEM image of SA-TiO2/RGO. (c) TEM image of SA-TiO2. (d) TEM image of SA-TiO2/RGO. (e) HRTEM image of SA-TiO2. (f) HRTEM image of SA-TiO2/RGO.
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Fig. 5. (a) CV curves of SA-TiO2/RGO electrode at a scan rate of 0.2 mV s−1. (b) Galvanostatic discharge/charge curves of SA-TiO2/RGO electrode at 0.2 A g−1. (c) Cycling performances of SA-TiO2, C-TiO2 and SA-TiO2/RGO electrodes at 0.2 A g−1. (d) Rate properties of SA-TiO2, C-TiO2 and SA-TiO2/RGO at different current densities. (e) Cycling performance and coulombic efficiency of SA-TiO2/RGO electrode at 5 A g−1.
of region C and thus higher capacity of SA-TiO2/RGO than the theoretical capacity. Moreover, the RGO electrode has low specific capacity and electrochemical reactivity in the voltage range of 1.0–3.0 V [51]. Therefore, the capacity of the SA-TiO2/RGO electrode is mainly contributed from TiO2. Fig. 5c compares the cycling stability of SA-TiO2/ RGO, SA-TiO2 and C-TiO2 electrodes at a current density of 0.2 A g−1. After 100 cycles, the SA-TiO2/RGO electrode shows a discharge capacity of 251.7 mAh g−1, which is much larger than those of SA-TiO2 (109 mAh g−1) and C-TiO2 (only 37.1 mAh g−1). Such an excellent electrochemical performance of SA-TiO2/RGO may be caused by its low internal resistance. The EIS measurement results (see Fig. S6) shows that the charge transfer resistance Rct is only 58.5 Ω for SA-TiO2/RGO, which is much lower than those of SA-TiO2 (93 Ω) and C-TiO2 (202 Ω). Fig. 5d exhibits rate performances of the C-TiO2, SA-TiO2 and SATiO2/RGO electrodes at different current densities from 0.2 to 10 A g−1. The SA-TiO2/RGO electrode delivers highly reversible average capacities of 319.3, 267.8, 235.1, 202, 159.2 and 135.6 mAh g−1 at 0.2, 0.5,
according to the total mass of SA-TiO2/RGO. The initial discharge capacity is 433.2 mAh g−1. With increasing cycle number, the discharge capacities are 363.6, 264 and 251.7 mAh g−1 for the 2nd, 50th and 100th cycles, which remain higher than the theoretical capacity (168 mAh g−1) of anatase [1]. Based on previous reports, the discharge curves for the SA-TiO2/RGO electrode can be divided into three different voltage regions, which are marked as A, B and C in Fig. 5b (here the discharge curve for the first cycle is taken as an example) [13,46,47]. In region A, a monotonic voltage drop to ∼1.75 V is observed, which is assigned to a homogeneous Li insertion into the bulk until the solid-solution limit of Li in TiO2 is reached. In region B, there is a biphase plateau at ∼1.75 V, indicating the coexistence of Li-rich and Li-poor phases. In region C, reversible storage of Li occurs at particle interfaces. The porous SA-TiO2/RGO hybrid and ultra-small SATiO2 nanoparticles with high specific surface area could readily induce the capacitive behavior during the electrochemical reaction, leading to an extra interfacial storage capacity [48–50]. This results in an increase 167
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density by Nb doping, which leads to a 3 orders of magnitude higher electronic conductivity of Ti0.94Nb0.06O2 compared with TiO2 [12]. Fig. 5e shows the long-term cycle stability of SA-TiO2/RGO electrode. The electrode was galvanostatically discharged and charged at 0.2 A g−1 for the first five cycles and then at 5 A g−1 for 2000 cycles. After 2000 cycles, the hybrid electrode still remains a capacity of 98 mAh g−1, exhibiting excellent cycling stability. Fig. 6 illustrates the mechanism behind superior electrochemical performances of SA-TiO2/RGO in LIBs. The excellent electrochemical performances of SA-TiO2/RGO originate from (i) the surface amorphous layer of TiO2, which provides plenty of oxygen vacancies for electrons/ ions transport, and alleviates the stress caused by volume changes during Li+ insertion/extraction [52]; (ii) high-conductivity RGO, which provides a conductive network, and also contributes a portion of capacity; (iii) the ultra-small SA-TiO2 nanoparticles with large surface/ volume ratio, which offers a large electrochemically active area and short ion diffusion distance; and (iv) the C-O-Ti bonds between SA-TiO2 and RGO, which reduce the inner resistance significantly.
Table 1 Comparisons of rate performances of some representative TiO2-based anodes materials for LIBs. Ti-based anodes
Rate capability
Ref.
Hollow TiO2/graphitic carbon Crystalline TiO2 hollow spheres
91 mAh g−1 at 10 A g−1 49.1 mAh g−1 at 5.19 A g−1 129 mAh g−1 at 1.68 A g−1 109 mAh g−1 at 3.74 A g−1 50 mAh g−1 at 0.168 A g−1 103 mAh g−1 at 8.4 A g−1 85 mAh g−1 at 3.36 A g−1
[53] [54]
70 mAh g−1 at 10 A g−1 120 mAh g−1 at 16.75 A g−1 135.6 mAh g−1 at 10 A g−1 121.9 mAh g−1 at 16.75 A g−1
[58] [12]
3-shelled TiO2 hollow microspheres TiO2-rGO nanocomposite hollow spheres Oxygen-deficient anatase TiO2@C Surface-amorphized TiO2@graphene Graphene supported ultralong TiO2 nanofiber Graphitic carbon-TiO2 Rutlie Ti0.94Nb0.06O2 SA-TiO2/RGO
[7] [55] [56] [17] [57]
This work
4. Conclusions In summary, a unique hybrid material of SA-TiO2/RGO has been fabricated through a two-step method, which combines SA-TiO2 and RGO tightly through the formed C-O-Ti covalent bonds during the hydrothermal reaction. As a result, such a hybrid material as an anode in Li-ion batteries exhibits superior electrochemical performances, high specific capacities (251.7 mAh g−1 after 100 cycles at 0.2 A g−1), excellent rate performance (135.6 mAh g−1 at 10 A g−1) and cycling stability (98 mAh g−1 after 2000 cycles at 5 A g−1). The developed hybrid material and its fabrication method reported in this work provide new insights into the fundamental understanding of high-performance surface amorphized metal oxides toward application in energy storage devices. Acknowledgements This project is financially supported by National Natural Science Foundation of China (Nos. 51671092 and 51631004), ChangBai Mountain Scholars Program, Natural Science Foundation of Jilin Province (No. 20160101315JC), Program for JLU Science and Technology Innovative Research Team (No. 2017TD-09) and the Fundamental Research Funds for the Central Universities. Appendix A. Supplementary data
Fig. 6. Schematic illustration of the mechanism behind superior electrochemical performances of SA-TiO2/RGO in LIBs.
Supplementary data related to this article can be found at http://dx. doi.org/10.1016/j.jpowsour.2018.07.019.
−1
1, 2, 5 and 10 A g , respectively, showing the highest capacity at each current density. When the current density comes back to 0.2 A g−1, the SA-TiO2/RGO electrode resumes a capacity of 261.9 mAh g−1 rapidly and then retains stable for further cycling. It is noteworthy that the capacity of SA-TiO2/RGO (135.6 mAh g−1) is 15.4 times that of C-TiO2 (8.8 mAh g−1) at 10 A g−1, displaying superior rate performance compared with other TiO2-based anode materials (see Table 1). It should be noted that a recent report indicates that the rutile Ti0.94Nb0.06O2 shows excellent rate property in Li-ion battery, for example, a high capacity of 120 mAh g−1 at a current density of 16.75 A g−1 [12]. For a comparison, we have measured the capacity of SA-TiO2/RGO at 16.75 A g−1, which is 121.9 mAh g−1, very similar to that of rutile Ti0.94Nb0.06O2 (120 mAh g−1). The superior rate properties of SA-TiO2/RGO can be attributed to: (i) the built-in electric field across the amorphous state/crystalline state interface of SA-TiO2, reducing the charge transfer resistance; and (ii) the excellent contact between SA-TiO2 and RGO through C-O-Ti bonds, decreasing the inner resistance. For rutile Ti0.94Nb0.06O2, its excellent rate performance originates from the significant increase of the conduction electron
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