Cation-Assisted Formation of Porous TiO2–x Nanoboxes with High

Jan 12, 2018 - Recently, regulating the active grain boundaries (GBs) in electrocatalysts has become an efficient approach to increase the catalytical...
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Cation-Assisted Formation of Porous TiO2-x Nanoboxes with High Grain Boundaries Density as Efficient Electrocatalysts for Lithium-Oxygen Batteries Guoxue Liu, Wei Li, Ran Bi, Christian Atangana Etogo, Xin-Yao Yu, and Lei Zhang ACS Catal., Just Accepted Manuscript • DOI: 10.1021/acscatal.7b04182 • Publication Date (Web): 12 Jan 2018 Downloaded from http://pubs.acs.org on January 13, 2018

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Cation-Assisted Formation of Porous TiO2-x Nanoboxes with High Grain Boundaries Density as Efficient Electrocatalysts for LithiumOxygen Batteries Guoxue Liu,† Wei Li,† Ran Bi,† Christian Atangana Etogo,† Xin-Yao Yu,*, ‡ and Lei Zhang*, † † Key Lab of Heat Transfer Enhancement and Energy Conservation of Ministry of Education, School of Chemistry and Chemical Engineering, South China University of Technology, Guangzhou 510640, China ‡ School of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, China ABSTRACT: Despite its ultrahigh energy density, the practical use of rechargeable lithium-oxygen (Li-O2) batteries is still hindered by the poor performance of electrocatalysts. Recently, porous and hollow nanostructured catalysts have been found to be promising electrocatalysts toward ORR/OER to improve the performance of Li-O2 batteries. However, the catalytic activities of the catalysts are still insufficient and the nanopores are easily clogged by the insoluble discharge products. Herein, highly porous and hollow TiO2-x nanoboxes (TiO2-x NBs) with large surface area, increased oxygen vacancies, and high grain boundaries density are developed as an electrocatalyst for Li-O2 batteries via a cation-assisted synthetic approach. The Li-O2 batteries with TiO2-x NBs as the electrocatalyst exhibit a high specific capacity of 8569 mA h g−1, excellent rate capability, and long cycle life up to 200 cycles when the capacity is limited to 1000 mA h g−1. Moreover, the cation-assisted approach for fine tuning the nanostructure would be an innovative strategy to design efficient electrocatalysts for high-performance Li-O2 batteries.

KEYWORDS：titanium dioxide; electrocatalysts; hollow structure; grain boundaries; lithium-oxygen batteries 1. Introduction With the rapid development of portable electronics, long-range electric vehicles, and large-scale smart grids, electrochemical energy storage systems with high energy density and longlasting stability are urgently required.1 Recently, rechargeable lithium-oxygen (Li-O2) battery has been widely regarded as a competitive candidate for next-generation energy storage system in virtue of its ultrahigh energy density of 3505 Wh kg−1.25 However, the inherent sluggish kinetics of the electrocatalysts toward the reversible reaction of 2Li+O2 → Li2O2, including oxygen reduction reaction (ORR) and oxygen evolution reaction (OER), results in poor performance of Li-O2 batteries such as high overpotentials, low rate capability, and poor cycling stability.6-9 In addition, the electrocatalysts cannot offer sufficient space for accommodating the solid Li2O2 discharge product, which leads to low practical capacity.10-13 Therefore, designing an ideal electrocatalyst with high catalytic activities for both ORR and OER, and appropriate micro/nanostructures for Li2O2 deposition is highly expected to address the above problems. During the past decade, designing and synthesis of porous and hollow nanostructured catalysts,14-20 such as porous La0.75Sr0.25Mn3 nanotubes,14 Pt-coated hollow graphene nanocages,15 and 1D RuO2/Mn2O3 hollow architectures,16 have made enormous advances in the field of Li-O2 batteries, caused by their inherent advantages including large surface area to expose more active sites, short diffusion distances for fast and uniform transportation of mass/charge, and extra void space to store more Li2O2 products. However, the catalytic activities of these electrocatalysts for ORR/OER are still insufficient and the nanopores of which are easily clogged by the

insoluble Li2O2 products,21-23 Therefore, in order to improve the catalytic activities and optimize the mass/charge channels and pore structure in electrocatalysts, the increasing of the catalytic active sites is of great importance. Recently, regulating the active grain boundaries (GBs) in electrocatalysts has become an efficient approach to increase the catalytic active sites,24-31 For example, Wang and coworkers reported a Li (lithium) electrochemical tuning method, by which ZnO was converted into ultrasmall metallic zinc with a lot of GBs, achieving improved catalytic activity and high selectivity for carbon dioxide reduction.24 Cui’s group reported an efficient bi-functional ultrasmall nanoparticles electrocatalyst through obtained by a lithium-induced conversion reaction, which can create many grain boundaries and result in more exposed catalytic active sites, exhibit high catalytic activity for OER/HER and good stability.25 However, until now, it is rarely to apply this method to the synthesis of electrocatalysts for Li-O2 batteries, involved with a relatively complicated process and the solid electrolyte interface (SEI) usually generating on the surface of the catalysts, which may block the catalytic active sites. Among various electrocatalysts, titanium-based materials, such as TiO2,32-35 TiN,36 TiC,37 TiSi2,38 and Ti4O7,39 are promising electrocatalysts for Li-O2 batteries due to their good electrochemical stability, low cost, and environmental benignity. Nevertheless, the catalytic active sites and storage space for the Li2O2 products still need to be improved.36, 39 Herein, we develop a rational design and synthesis of highly porous and hollow TiO2-x nanoboxes (TiO2-x NBs), which are composed of interconnected ultra-small nanoparticles with large surface area, increased oxygen vacancies, and high GBs density as an

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electrocatalyst for Li-O2 batteries via a facile and efficient cation-assisted approach. The porous TiO2-x NBs can provide plenty of space for the discharge products deposition and a large number of small pores to enhance mass/charge transportation. In addition, the interconnected ultra-small nanoparticles with abundant oxygen vacancies and GBs will significantly offer numerous catalytic active sites toward ORR/OER. Therefore, the Li-O2 batteries based on porous TiO2-x NBs exhibit high specific capacity (8569 mA h g−1), superior rate capability, and long cycle life (200 cycles) with the capacity limited to 1000 mA h g−1.

Figure 1. Schematic illustration of the formation of porous TiO2-x nanoboxes. I) hydrothermal and annealing treatments to form pure perovskite phase, II) selective dissolution of barium and strontium ions to obtain porous TiO2-x nanoboxes.

2. Experimental Section 2.1. Chemicals. Iron(III) chloride hexahydrate [FeCl3•6H2O, 98%], sodium hydroxide [NaOH, 98%], titanium(IV) fluoride [TiF4], diethylene glycol [C4H10O3, 99.0%], and acetic acid [CH3CO2H, 99.7%] were purchased from Sigma-Aldrich. Oxalic acid [C2H2O4•2H2O, 98%], strontium hydroxide octahydrate [Sr(OH)2•8H2O, 99.5%], and barium hydroxide hydrate [Ba(OH)2•xH2O, 99.99%] were purchased from Aladdin. 2.2. Synthesis of TiO2 NBs. First, the Fe2O3 nanocubes were synthesized via a previously reported precipitation method.40 Then, 0.040 g of as-obtained Fe2O3 nanocubes were dispersed into a 45 mL mixed solution of acetic acid and diethylene glycol (1:8) by ultrasonication for 30 min, followed by addition of 0.050 g of TiF4. The suspension was transferred into a 50 mL Teflon-lined stainless steel autoclave and solvothermally treated at 180 oC for 8 h. The as-obtained products were collected by centrifugation, washed with deionized water and ethanol several times, respectively, and dried at 60 o C for 12 h. In order to improve the crystallinity of TiO2 layer, the as-obtained Fe2O3@TiO2 nanocubes were annealed at 500 o C in air for 2 h. Then, the annealed products were dispersed into 0.5 M oxalic acid treated in 70 oC an oven for 72 h to selectively remove the Fe2O3 nanocubes templates. Finally, the TiO2 nanoboxes were obtained by centrifugation and washed with deionized water and ethanol for several times, and dried at 60 oC for 12 h. 2.3. Synthesis of Ba0.5Sr0.5TiO3 NBs. The as-obtained TiO2 NBs were dispersed into a 40 mL solution of Ba(OH)2 and Sr(OH)2 with 0.020 M Ba2+ and Sr2+ ions and the suspension was transferred into a 50 mL Teflon-lined stainless steel autoclave and hydrothermally treated at 200 oC for 10 h. The asprepared products were collected by centrifugation and washed with a 0.1 M diluted HNO3 aqueous solution, deionized water and ethanol, respectively, and dried at 60 oC for 12

h. Finally, the Ba0.5Sr0.5TiO3 NBs were obtained after being annealed at 650 oC in air for 2 h. 2.4. Synthesis of porous TiO2-x NBs. The Ba0.5Sr0.5TiO3 NBs dispersed into 0.5 M oxalic acid were treated in an oven at 70 oC for 24 h to selectively dissolve the Ba2+ and Sr2+ ions. The obtained porous TiO2-x NBs were collected by centrifugation and then washed with deionized water and ethanol, respectively, and finally dried in a vacuum oven at 60 oC for 12 h. 2.5. Materials Characterization. The morphology and microstructure of all the samples were characterized by fieldemission scanning electron microscope (FESEM, Hitachi SU8220), transmission electron microscope (TEM, JEOL JEM-2100F). The crystal structures of all the samples were characterized by X-ray diffraction (XRD, PANalytical X’pert Powder) with Cu-Kα radiation (λ = 1.54 Å). The nitrogen adsorption/desorption isotherms of all samples were measured at 77 K by a Micromeritics ASAP 2460 instrument. The X-ray photoelectron spectroscopy (XPS) analysis of all the materials was taken by a XPS instrument (Thermo Scientific Escalab 250Xi). 2.6. Electrochemical Measurements. The electrochemical properties of Li-O2 batteries were measured at room temperature by a 2032-type coin cell assembled in a glove box (MBRAUN MB-Unilab Pro SP 1800/780) in an Ar atmosphere (with O2 < 0.1 ppm and H2O < 0.1 ppm). The Li-O2 batteries consist of a cathode (oxygen electrode), a separator (glass fiber membrane, Whatman/F), an anode (Li-metal foil), and an electrolyte solution (1 M LiTFSI/TEGDME). The homogenous slurry of oxygen electrode was first prepared by mixing the porous TiO2-x NBs or the TiO2 NBs, polyvinylidene fluoride (PVDF) and Ketjen Black (KB) with a weight ratio of 30 : 60 : 10 with the addition of appropriate amount of N-methy1-2-pyrrolidone (NMP). Then the slurry was coated on Ni foam or carbon paper before being transferred into a vacuum oven at 100 oC for 12 h. The loading mass of the electrocatalysts is around 0.5 mg cm−2 in the electrodes. After resting for 5 to 10 h, the galvanostatical discharge and charge tests of the Li-O2 batteries were performed on a Neware Battery Testing System at room temperature within a voltage window of 2.0 − 4.5 V (vs. Li/Li+). Cyclic voltammetry tests were carried out at a rate of 0.1 mV s−1 within the potential range of 2.0 to 4.5 V (vs. Li/Li+) by an electrochemical workstation (Gamry 1000E).

Figure 2. FESEM images of a,b) TiO2 nanoboxes and c,d) porous TiO2-x nanoboxes.

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ACS Catalysis 3. Results and Discussion Fe2O3 nanocubes, developed by our previous study, are employed as removable templates to obtain uniform TiO2 nanoboxes (denoted as TiO2 NBs).40-41 The Fe2O3 nanocubes are coated with a thin layer of TiO2 to form the Fe2O3@TiO2 coreshell nanocubes and subsequently annealed at 500 oC for two hours to increase the crystallinity of TiO2 layer. The TiO2 NBs are then obtained by removing the Fe2O3 core in oxalic acid solution. Next, the TiO2 NBs are converted to porous TiO2-x NBs via a newly developed cation-assisted approach illustrated in Figure 1. The above TiO2 NBs are dispersed into a mixed solution of Ba(OH)2 and Sr(OH)2 with 0.020 M Ba2+ and Sr2+ ions by ultrasonication and as-obtained Ba0.5Sr0.5TiO3 nanoboxes (Ba0.5Sr0.5TiO3 NBs) with perovskite phase are prepared via sequential hydrothermal and annealing processes (Step I). Finally, porous TiO2-x NBs are obtained by selectively dissolving Ba2+ and Sr2+ ions in Ba0.5Sr0.5TiO3 using oxalic acid solution (Step II). The Fe2O3 nanocubes template are prepared through a facile co-precipitation method. Field-emission scanning microscopy (FESEM) images show that the Fe2O3 nanocubes are highly uniform with an average diameter of around 450 nm (Figure S1a, b). Then, a TiO2 shell layer is deposited on the core Fe2O3 nanocubes via a modified solvothermal method. After following annealing treatment, it can be seen that the core-shell Fe2O3@TiO2 nanocubes are relatively uniform with a rough surface consisting of small nanoparticles (Figure S1c). When the core-shell Fe2O3@TiO2 nanocubes are dissolved into an oxalic acid solution, the Fe2O3 templates can be selectively etched. After removing the Fe2O3 templates, the uniform nanobox structure with a thin TiO2 shell is still maintained without obvious structural collapse (Figure 2a, b and Figure S2). After subsequent hydrothermal treatment in an alkaline solution containing barium ions and strontium ions and following annealing treatment, Ba0.5Sr0.5TiO3 NBs are synthesized. As shown in the low-magnification FESEM image (Figure S3a), the uniform Ba0.5Sr0.5TiO3 NBs entirely retain the hollow structures and no pulverization after the intercalation of barium ions and strontium ions into the crystal lattice

of TiO2 is observed. Interestingly, the particles size of Ba0.5Sr0.5TiO3 NBs become larger compared to those of TiO2 NBs and the shells still keep a dense surface (Figure S3b). Finally, the porous TiO2-x NBs are obtained via a facile selective dissolution of barium ions and strontium ions in oxalic acid solution. During this process, a lot of mesopores uniformly dispersed on the TiO2 shells are generated and the hollow structure is perfectly retained (Figure 2c). Fig. 2d clearly shows that the shells of the porous TiO2-x NBs are thicker than those of the TiO2 NBs, demonstrating that the porous TiO2-x NBs possess more pores (Figure 2d). Transmission electron microscope (TEM) is further used to examine the microstructure of the as-synthesized products. In agreement with the above FESEM observations, the internal space of TiO2 NBs can be easily observed by the distinct contrast between the TiO2 shell and the inner cavity (Figure 3a). The shell thickness of these nanoboxes is measured to be about 20 nm (Figure 3b). A magnified TEM image reveals that the shell of TiO2 NBs consists of nanoparticles with a size of about 10-20 nm (inset of Figure 3c and Figure S4a). The high-resolution TEM (HRTEM) image of the TiO2 NBs clearly shows the lattice fringes with an inter-planar spacing of 3.2 Å, in good accordance with the d-spacing of (100) planes of rutile TiO2 (Figure S4b).40, 42 The high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image further demonstrates the hollow structure of the TiO2 NBs with a mesoporous shell (Figure S5a). The corresponding elemental mapping images of one single nanobox reveals that the Ti and O elements are uniformly distributed on the surface of TiO2 NBs (Figure S5b-d). In addition, the microstructure of Ba0.5Sr0.5TiO3 NBs is further investigated by TEM (Figure S6). The hollow structure can be obviously seen by the clear contrast between dark Ba0.5Sr0.5TiO3 shell and light hollow cavity (Figure S6a, b). In contrast to TiO2 NBs, the shell of Ba0.5Sr0.5TiO3 NBs is thicker and denser (Figure S6c, d). The shell thickness is measured to be about 100 nm and the size of the nanoparticles on the Ba0.5Sr0.5TiO3 shell is around 40-50 nm (Figure S6c, d).Compared with Ba0.5Sr0.5TiO3 NBs, the NBs also porous TiO2-x

Figure 3. TEM images of a-c) TiO2 nanoboxes and d-f) porous TiO2-x nanoboxes. The inset pictures in panel c and d are corresponding HRTEM images of TiO2 nanoboxes and porous TiO2-x nanoboxes, respectively.

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ACS Catalysis duced by the transformation of Ti4+ to Ti4+/Ti3+.45 The Ti XPS spectrum of the porous TiO2-x NBs shows that the binding energies of Ti 2p are a little lower than those of the TiO2 NBs, suggesting that Ti ions of the porous TiO2-x NBs have lower oxidation state than those of the TiO2 NBs, leading to higher oxygen vacancies.32, 42 The O 1s spectrum (Figure 6d) of both porous TiO2-x NBs and TiO2 NBs can be spliced into three main peaks with binding energies located at 530.2 eV (lattice oxygen), 532.2 eV (oxygen vacancies), and 533.7 eV (adsorbed molecular water), respectively.45-46 The porous TiO2-x NBs possess higher ratio of oxygen vacancies/lattice oxygen than that of TiO2 NBs.

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perfectly maintain hollow structure with the same shell thickness (Figure 3d-e). Interestingly, the shell of these nanoboxes is highly porous and composed of ultra-small nanoparticles with a size of about 2-4 nm (Figure 3f and Figure S7). It is antici-pated that the porous TiO2-x NBs with high GBs density can provide numerous catalytic active sites toward ORR/OER.25 Furthermore, the lattice fringes in the HRTEM image of the porous TiO2-x nanoboxes are separated by 3.2 Å, matching well with d-spacing of (100) planes of rutile TiO2 (Figure S7b). The HAADF-STEM image further verifies the hollow and mesoporous structure of the porous TiO2-x NBs (Figure 4a). The corresponding elemental mapping images of the porous TiO2-x NBs are shown in Figure 4b-d, indicating the uniform distribution of Ti, O and a few Ba and Sr elements in the porous shells.

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Figure 4. a) HAADF-STEM and b-f) corresponding EDX elemental mapping images of b) mix, c) Ti, d) O, e) Ba, and f) Sr elements in porous TiO2-x nanoboxes.

The crystal phase change of TiO2 before and after barium ions and strontium ions intercalation/dissolution is investigated by X-ray diffraction (XRD) analysis. The XRD patterns of both porous TiO2-x NBs and TiO2 NBs show well-defined diffraction peaks (Figure 5a), which can be well indexed to rutile TiO2 (JCPDS card no. 89-4202),40 implying that the crystal phase of TiO2 is not influenced by barium and strontium ions intercalation/dissolution. The nitrogen adsorption-desorption isotherms of the porous TiO2-x NBs and TiO2 NBs are shown in Figure 5b. The porous TiO2-x NBs exhibit a type-IV curve with an apparent hysteresis loop, suggesting a large number of mesopores in the shell.42 The porous TiO2-x NBs have a high BET surface area of 60.6 m2 g−1, while the TiO2 NBs only possess a BET surface area of 26.3 m2 g−1, indicating that the porous TiO2-x NBs could offer more catalytic active sites enhancing both the ORR and OER, and more tri-phase regions for the electrochemical reaction between Li+ and O2.43-44 In addition, the pore size distribution (Figure S8) shows that the porous TiO2-x NBs are with hierarchical porous structure, which could enhance Li+ and O2 transfer. The surface chemistry environment of the porous TiO2-x NBs and TiO2 NBs are further characterized by X-ray photoemission spectroscopy (XPS). Typical Ti and O signals are all detected in the porous TiO2-x NBs and TiO2 NBs (Figure 6a-c). Compared with the TiO2 NBs, the porous TiO2-x NBs have a few Ba and Sr signals, which is in consistent with the results of EDX elemental mapping (Figure S12). The remaining Ba and Sr cations in the porous TiO2-x NBs might derive from charge imbalance in-

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Based on the above characterization data, the formation process of the porous TiO2-x NBs includes two steps explained in detail as follows. Firstly, in the phase-conversion process, the rutile TiO2 NBs are transformed into the perovskite Ba0.5Sr0.5TiO3 NBs. In this process, a thin Ba0.5Sr0.5TiO3 layer is initially generated on the surface of TiO2 nanoparticles through the reaction between TiO2 and dissolved Ba2+ and Sr2+ under alkaline solution. With the diffusing of Ba2+ and Sr2+ through the Ba0.5Sr0.5TiO3 layer, the inner TiO2 nanoparticles are further transformed into Ba0.5Sr0.5TiO3 until fully transformed into polycrystalline Ba0.5Sr0.5TiO3 nanoparticles.47 Although the atomic ration of Ba to Sr in Ba0.5Sr0.5TiO3 is not measured by EDX because of the overlap between characteristic X-ray of Ti-Kα (4.511 KeV) and Ba-Lα (4.466 KeV), the perovskite phase of Ba0.5Sr0.5TiO3 NBs is confirmed by XRD

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analysis (Figure S13). During the in-situ growth process, the hollow structure can be well retained and the shells of nanoboxes become denser (Figure S6). Secondly, in the selective dissolution process, the porous TiO2-x NBs are obtained via selective dissolution of Ba and Sr cations in Ba0.5Sr0.5TiO3 NBs with oxalic acid solution. It should be emphasized that the coordination between Ba2+/Sr2+ and oxalic acid is crucial to the successful synthesis of porous TiO2-x NBs. Compared with annealing under H2 treatment method and Li electrochemical tuning method, the present method is more facile and efficient to synthesize highly porous TiO2-x NBs with increased oxygen vacancies and high GBs density, which are expected to offer favorable porous structures and numerous catalytic active sites for achieving rechargeable Li-O2 batteries. The electrochemical performance of the porous TiO2-x NBs as the electrocatalyst for Li-O2 batteries is investigated in 1 M LiTFSI of TEGDME electrolyte. The fabrication details of the Li-O2 cells are shown in the Experimental Section. Cyclic voltammetry (CV) is conducted in the voltage range from 2.0 to 4.5 V (vs. Li/Li+) at a scan rate of 0.1 mV s−1 to investigate the electrocatalytic activity of the porous TiO2-x NBs. As shown in Figure 7a, the porous TiO2-x NBs/KB cathode exhibits higher ORR onset potential and larger ORR peak current density, compared with the TiO2 NBs/KB and pure KB cathodes, indicating that the porous TiO2-x NBs possess higher ORR electrocatalytic activity to enhance the formation of discharge products. In addition, the porous TiO2-x NBs/KB cathode exhibits lower OER onset potential and larger OER peak

current density than TiO2 NBs or KB, suggesting its higher electrocatalytic activity toward OER to effectively decompose the discharge products. Therefore, these results indicate that the porous TiO2-x NBs electrocatalyst possesses excellent bifunctional activities toward ORR/OER to improve the performance of the Li-O2 batteries. The Li-O2 batteries performance of the porous TiO2-x NBs electrocatalyst is further evaluated by galvanostatic discharge/charge analysis. Figure 7b displays the first discharge/charge curves of Li-O2 batteries with the porous TiO2-x NBs, TiO2 NBs, and KB. Notably, the Li-O2 batteries with the porous TiO2-x NBs exhibit a discharge capacity of 8569 mA h g−1 at a current density of 100 mA g−1, which is higher than that of TiO2 NBs (7818 mA h g−1) and that of KB (6804 mA h g−1) at the same current density. Furthermore, the discharge voltage plateau of Li-O2 batteries with the porous TiO2-x NBs is ~2.71 V, which is higher than that of TiO2 NBs or that of KB by about 10-20 mV. In addition, the charge voltage plateau of the Li-O2 batteries with the porous TiO2-x NBs is obviously lower than that of TiO2 NBs (about 90 mV) or that of KB (about 200 mV). The low overpotentials could enhance the energy efficiency of the Li-O2 batteries. These results further demonstrate that the porous TiO2-x NBs can offer numerous catalytic active sites toward ORR/OER to facilitate the formation and decomposition of Li2O2. The rate capabilities of the Li-O2 batteries with the porous TiO2-x NBs, TiO2 NBs and KB are also investigated at various current densities. Figure 7c and Figure S9 shows that the Li-O2 batteries with the porous TiO2-x NBs display discharge

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Figure 7. a) CV profiles of porous TiO2-x nanoboxes, TiO2 nanoboxes, and bare KB electrodes between 2.0 and 4.5 V at 0.1 mV s−1 in coin-type cells. b) Discharge/charge profiles and c) rate capabilities of porous TiO2-x nanoboxes, TiO2 nanoboxes, and bare KB electrodes. d-f) Discharge/charge profiles of porous TiO2-x nanoboxes, TiO2 nanoboxes, and bare KB electrodes, respectively at different cycles under a specific capacity limit of 1000 mA h g−1. g) The terminal voltages of discharge process on vs. cycle numbers of porous TiO2-x nanoboxes, TiO2 nanoboxes, and KB electrodes.

capacities of 8569, 8392, 8187 and 7834 mA h g−1 at current densities of 100, 200, 250 and 500 mA g−1, respectively. By contrast, the Li-O2 batteries with the TiO2 NBs exhibit a sharply decreased discharge capacity of 7818, 7251, 6877 and 6452 mAh g−1 under the same circumstance. And the Li-O2 batteries with the KB deliver a poorer discharge capacities of 6804, 6456, 5496 and 4770 mA h g−1. The Li-O2 batteries performances with porous TiO2-x NBs also exhibit higher charge capacity than those of TiO2 NBs, and bare KB electrodes. In addition, with the increasing of current density, the initial Coulombic efficiency (CE) of porous TiO2-x NBs, TiO2 NBs, and bare KB electrodes decreases and the CE of porous TiO2-x NBs are higher than that of TiO2 NBs or bare KB (Figure S10), suggesting that Li-O2 batteries with porous TiO2-x NBs exhibit better reversibility. As shown in Figure S11, the overpotentials of porous TiO2-x NBs are lower than those of TiO2 NBs or bare KB, especially at high current density, indicating that the porous TiO2-x NBs have more catalytic active sites toward ORR/OER. The cycling performance of the Li-O2 batteries with the porous TiO2-x NBs, TiO2 NBs, and KB are further tested with a specific capacity limit of 1000 mA h g−1 at a current density of 200 mA g−1. As shown in Figure 7d-f, the Li-O2batteries with the porous TiO2-x NBs exhibit lower discharge/charge overpotentials than those with TiO2 NBs or KB. Furthermore, the Li-O2 batteries with the porous TiO2-x NBs show good stable specific capacity (Figure 7d and g) and retain a stable terminal discharge above 2.0 V after 200 cycles, indicating excellent cycling stability. Figure 7g shows that the discharge voltage of the TiO2-x NBs electrode decrease obviously at about the 25th cycle, which could be attributed to formation of the by-products. The small pores on the surface of

the porous TiO2-x NBs could be blocked by by-products at about the 25th cycle, which results in the decrease of discharge voltage. After 25th cycle, the discharge voltage keeps stable because the porous TiO2-x NBs have hierarchical porous structure which could provide enough storage space for discharge product and the catalytic active sites towards ORR/OER. In comparison, the Li-O2 batteries with TiO2 NBs and KB only exhibit 53 cycles and 9 cycles, respectively (Figure 7e-g).The good performance of the Li-O2 batteries with the porous TiO2-x NBs can be mainly ascribed to their unique and favorable structural features. First, the unique porous hollow structure can facilitate Li+/O2 transfer and offer large void space to support Li2O2 growth, which can improve the discharge capacity and rate capability. In addition, the porous TiO2-x NBs possessing high GBs density, large surface area, and numerous oxygen vacancies could provide abundant catalytic active sites to accelerate Li2O2 formation and decomposition, which is beneficial to improve the discharge capacity, rate capability, and cycle life.48-49 To further understand the enhanced performance of the Li-O2 batteries, the morphology of the cathode with the porous TiO2-x NBs electrocatalyst after discharge/charge processes are observed via FESEM. Figure S14a shows the discharge/charge profiles of the Li-O2 batteries with the porous TiO2-x NBs after two cycles. First, after discharged to 2000 mA h g−1 at current density of 200 mA g−1, the local surface of the porous TiO2-x NBs electrocatalyst is uniformly covered by rod-like products with a length of ~200 nm (Figure S14b). Then, after recharging, the rod-like products disappear on the surface of the porous TiO2-x NBs and the porous structure still can be retained (Figure S14c). In the following cycle, rod-like products generate and decompose on

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ACS Catalysis the surface of the porous TiO2-x NBs again (Figure S14d,e), indicating the porous TiO2-x NBs possess high catalytic activities toward ORR/OER and enough space to support discharge products.8, 14 These results further demonstrate the excellent rechargeability of Li-O2 batteries with the porous TiO2-x NBs as electrocatalysts. In addition, compared with the reported titanium-based electrocatalysts (Table S1), the porous TiO2-x NBs shows better electrochemical performance. 4. Conclusion In summary, we have developed a facile and effective cationassisted method for the synthesis of porous TiO2-x NBs as an electrocatalyst for Li-O2 batteries. The porous and hollow structure could not only promote mass (such as Li+ and O2) transfer into the cathodes but also supply appropriate space to maintain the formation and storage of the discharged Li2O2 products. The porous TiO2-x NBs with large surface area, increased oxygen vacancies, and high GBs density could effectively enhance the catalytic activities toward ORR/OER, which are beneficial for the growth and decomposition of Li2O2. When applied as an electrocatalyst for Li-O2 batteries, the Li-O2 batteries with porous TiO2-x NBs exhibit significantly improved discharge capacity, excellent rate capability, and superior cycling stability. The developed cation-assisted approach via phase conversion and cation selective dissolution for fine tuning the micro/nanostructure will provide a novel strategy to synthesize bi-functional electrocatalysts for highperformance Li-O2 batteries.

ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acscatal.XX. Detailed experimental procedures; the formation mechanism of porous TiO2-x NBs; FESEM images of Fe2O3 NCs and Fe2O3@TiO2 core-shell NCs (Figure S1); FESEM image of TiO2 NBs (Figure S2); FESEM images of Ba0.5Sr0.5TiO3 NBs (Figure S3); HRTEM images of TiO2 NBs (Figure S4); HAADF-STEM and corresponding EDX elemental mapping images of TiO2 NBs (Figure S5); TEM mages of Ba0.5Sr0.5TiO3 NBs (Figure S6); HRTEM images of porous TiO2-x NBs (Figure S7); pore size distribution of porous TiO2-x NBs and TiO2 NBs (Figure S8); discharge/charge profiles (Figure S9), the initial Coulombic efficiency (Figure S10) and the overpotentials (Figure S11) of porous TiO2-x NBs, TiO2 NBs, and bare KB electrodes at current densities; Ba 3d and Sr 3d high-resolution XPS spectra of porous TiO2x NBs (Figure S12); XRD pattern of Ba0.5Sr0.5TiO3 NBs (Figure S13); FESEM images of the cathode with porous TiO2-x NBs at different conditions (Figure S14).

AUTHOR INFORMATION Corresponding Author * E-mail: [email protected]; * E-mail: [email protected]

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT This work was financially supported by the National Key Research and Development Program of China (2016YFA0202604), the Natural Science Foundation of China (21606088, 51621001) and the “Thousand Talents Program”.

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