MagnÃ©li-Phase Ti4O7 Nanosphere Electrocatalyst Support for Carbon

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Magnéli-Phase Ti4O7 Nanosphere Electrocatalyst Support for Carbon-Free Oxygen Electrodes in Lithium–Oxygen Batteries Seun Lee, Gwang-Hee Lee, Jae-Chan Kim, and Dong-Wan Kim ACS Catal., Just Accepted Manuscript • DOI: 10.1021/acscatal.7b03741 • Publication Date (Web): 14 Feb 2018 Downloaded from http://pubs.acs.org on February 15, 2018

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Magnéli-Phase Ti4O7 Nanosphere Electrocatalyst Support for Carbon-Free Oxygen Electrodes in Lithium–Oxygen Batteries Seun Lee‡, Gwang-Hee Lee‡, Jae-Chan Kim, and Dong-Wan Kim* School of Civil, Environmental and Architectural Engineering, Korea University, Seoul 136-713, South Korea

ABSTRACT: Lithium–oxygen batteries have been considerably researched due to their potential for high energy density compared to some rechargeable batteries. However, it is known that the stability of a carbon-based oxygen electrode is insufficient owing to the promotion of carbonate formation, which results in capacity fading and large overpotential in lithium–oxygen batteries. To improve the chemical stability in organic-based electrolytes, alternative electrocatalyst support materials are required. The Ti–O crystal system appears to provide a good compromise between electrochemical performance and cost, and is thus an interesting material for further investigation. Here, we investigate a carbon-free electrode with the goal of identifying routes for its successful optimization. To replace carbon materials as an electrocatalyst support, Magnéli Ti4O7 nanospheres were synthesized from anatase TiO2 nanospheres via a controlled thermochemical reduction. The Magnéli Ti4O7 nanospheres demonstrated effective overpotential

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characteristics (1.53 V) compared to the anatase TiO2 nanospheres (1.91 V) during chargedischarge cycling at a current rate of 100 mA g–1. Additionally, RuO2@Magnéli-Ti4O7 nanospheres were prepared as a bi-functional catalyst-containing oxygen electrode for lithium– oxygen batteries, providing a remarkably reduced overpotential (0.9 V).

KEYWORDS: Magnéli phase, carbon-free, Ti4O7, RuO2, Li-O2 batteries


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INTRODUCTION Lithium–oxygen (Li–O2) batteries are one of the most anticipated energy storage technologies, with high theoretical gravimetric energy density (3505 W h kg–1) due to the lightweight Li metal and O2 gas.1,2 However, there are various chemical stability issues that must be overcome in order to achieve high performance.2,3 One of these challenges is the selection of highly conductive and chemically stable oxygen electrode materials to effectively perform oxygen reduction and oxygenation reactions (ORR/OER). Various carbon materials such as super P carbon black, ketjen black, and graphene are commonly used as an oxygen electrode for Li–O2 batteries.4-6 However, when such a carbon material is used, it accelerates the decomposition of organic-based electrolytes during the charge-discharge process of Li–O2 batteries7,8 and form unstable carbon species such as Lithium carbonate as an insulating by-product during the discharge process.

Li2 O2 + C +

1 O2 → Li2 CO3 (∆G = –542.4 kJ mol–1 ) 2

2Li2 O2 + C→Li2 O + Li2 CO3 (∆G = –533.6 kJ mol–1 )

The Lithium carbonate by-product does not completely decompose and accumulate on the electrode surface. Consequently, a rapid increase in charge voltage (> 4.2 V) and deterioration of cycle stability occur.9 To solve this problem, appropriate selection of carbon-free oxygen electrodes has become a major issue in Li–O2 batteries. For carbon-free oxygen electrodes, TiO2 is an attractive candidate for an electrocatalyst support and is one of the most studied materials owing to its stability, environmental friendliness, and low cost.10,11 In particular, TiO2 is chemically and


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electrochemically stable in a corrosive environment compared to carbon, thus avoiding corrosion in the operating environment of Li–O2 cells.12 When these characteristics are used as an electrocatalyst support in an oxygen electrode, side-reactions are suppressed in the reaction with organic-based electrolytes as compared to the reactions of carbon-based materials.13,14 Ti4O7 is known as a Magnéli phase, which has excellent electrical conductivity due to its structural characteristics, oxygen vacancies, and Ti3+ ions in TiO2–x and has been used in various fields such as water treatment and fuel cells.15,16 In particular, it has been proposed as a substitute for carbon-based materials in Li–O2 batteries. It has been reported that Ti4O7 maintains its state during charge-discharge processes wherein it is exposed to oxygen-rich environment. However, the specific capacity of Ti4O7 is relatively lower than that of the other non-carbon materials.17 Herein, we report a facile method for obtaining Magnéli phase TinO2n–1 via controlled thermochemical reduction under various H2 flow rates. To prevent grain growth during thermochemical reduction, a silica coating shell was formed onto the Magnéli-Ti4O7 (M-Ti4O7). In addition, we synthesized RuO2@M-Ti4O7 via nano-deposition of a highly active catalyst, RuO2 on the as-obtained Ti4O7.18,19 M-Ti4O7 supports the activation of catalytic activity of RuO2 nanoparticles during the discharge-charge process. As a result, it was confirmed that Ti4O7 could be used as a substitute for carbon-based materials, which is an oxygen electrode material for Li–O2 batteries.

METHODS Preparation of single M-Ti4O7 phase via thermochemical reduction. Commercial TiO2 (99.7%, Sigma-Aldrich) powder was loaded into an alumina boat and placed in a horizontal tube furnace. Before the thermochemical reduction process, residual oxygen and water were removed


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for 30 min under 100 sccm of N2 gas. Thereafter, the N2 gas was replaced with H2 gas at various flow rates of 15, 25, 100, and 200 sccm. The tube furnace was then heated to 1050 °C at a rate of 8 °C min–1 and maintained for 2 h in a H2 gas flow. Finally, the furnace was cooled to room temperature at a rate of 5 °C min–1.

Synthesis of anatase TiO2 nanospheres. Initially, amorphous TiO2 nanospheres (amorphousTiO2 NS) were obtained via the sol–gel process.20 Hexadecylamine (0.994 g; CH3(CH2)15NH2, 90%, Sigma-Aldrich) was dissolved in 100 mL of ethanol, followed by the addition of 0.8 mL of 0.1 M KCl (99%, Sigma-Aldrich) solution. Then, 2.0 mL of titanium(IV) isopropoxide (Ti[OCH(CH3)2]4, 97%, Sigma–Aldrich) was added under stirring at room temperature for 1 h. The obtained suspension was washed with ethanol several times by centrifugation and subsequently dried at 70 °C in air. The obtained amorphous-TiO2 NS were crystallized to anatase-TiO2 nanospheres (A-TiO2 NS) in air at 450 °C for 4 h.

Synthesis of M-Ti4O7 nanospheres. To prevent grain growth, we applied a silica coating before synthesizing M-Ti4O7 nanospheres (M-Ti4O7 NS). A-TiO2 NS (0.5 g) were dispersed in 150 mL of a mixed solution of ethanol and distilled water, followed by the addition of 3.6 mL of ammonia solution (28 ~ 30 %, Samchun Chemical). Then, 6 mL of tetraethyl orthosilicate (Si(OC2H5)4, 99%, Sigma–Aldrich) was injected into the suspension. The mixture solution was stirred for 4 h and then washed repeatedly with distilled water. The resultant particles were dried at 70 °C in air. The obtained particles, which is silica-coated A-TiO2 NS, were thermally reduced for 2 h at 1050 °C under H2 (100 sccm). To remove silica coating, the silica-coated A-TiO2 NS were dispersed in 50 mL of distilled water with 5 mL of 2.5 M NaOH (98%, Samchun Chemical)


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solution, and the suspension was heated at 90 °C for 8 h. The final product was washed several times with deionized water and dried at 70 °C in air overnight.

Synthesis of RuO2@Magnéli-Ti4O7 nanospheres (RuO2@M-Ti4O7 NS). In 15 mL of distilled water, 0.15 g of the prepared M-Ti4O7 NS, 0.06 g of ruthenium chloride hydrate (RuCl3·xH2O, 99%,

Kojima

chemical),

and

1

g

of

cetyl

trimethyl

ammonium

bromide

(CH3(CH2)15N(Br)(CH3)3, 95%, Sigma–Aldrich) were dispersed. The solution was sealed in a 25-mL Teflon-lined stainless steel autoclave and heated to 150 °C for 10 h. The autoclave was cooled to room temperature, and the obtained particles were washed and filtered with a nylon membrane (Durapore, 0.22 mm, Millipore) several times using distilled water, and then dried at 70 °C overnight.

Material characterization. The morphological evolution of the samples was investigated using a transmission electron microscope (TEM; Tecnai G2 F30 S–Twin, FEI, USA) and a field emission scanning electron microscope (FESEM; S–4300, Hitachi, Japan). To examine the chemical composition, elemental mapping data were obtained using a scanning transmission electron microscope-energy dispersive X-ray spectroscope (STEM-EDS; JEM–2100F, JEOL). The chemical states of the elements were analyzed by X-ray photoelectron spectroscopy (XPS; PHI X–tool, ULVAC–PHI, Inc., Al 1486.6 eV Monochromatic at 23.1 W). The chemical reactions of the electrolyte during the discharge-charge process were analyzed by Fourier transform infrared spectroscopy (FTIR, LabRam ARAMIS IR2, HORIBA).


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Electrochemical performance of Li–O2 cells. The electrochemical performance of A-TiO2 NS, M-Ti4O7 NS, and RuO2@M-Ti4O7 NS were evaluated using Swagelok-type cells. The electrode was prepared by mixing each of the prepared samples (90 wt%) with carboxymethyl cellulose (10 wt%, Aldrich, Average Mw ~700,000) and dip-coating the obtained slurry on a nickel foam. The areal loading weights of all the electrodes were adjusted to be above 0.32 mg cm-2 based on the area of circular electrodes (area = 0.785 cm2). The Li–O2 cells were assembled in an Ar-filled glove box. The cells consisted of a lithium foil as an anode, a glass microfiber filter (Celgard 2400, Wellcos, Korea) as a separator, and 1 M LiNO3 (Alfa Aesar, anhydrous, ≥99.999%) in N,N-dimethylacetamide (DMAc; Alfa Aesar, anhydrous, ≥99.8%) as the electrolyte. All measurements were conducted in 1.5 atm dry oxygen to avoid any negative effect of humidity and CO2. The assembled cells were tested with an automatic battery cycler (WBCS 3000, WonAtech, Korea) in a voltage window of 2.0–4.5 V. Electrochemical impedance spectroscopy (EIS) was performed with an electrochemical workstation (Ivium-n-Stat electrochemical analyzer, Ivium Technologies B. V., Netherlands). The impedance response was collected by applying AC voltages of 10 mV while maintaining a constant DC voltage in the frequency range of 0.01 Hz to 100 kHz. All the above measurements were conducted at room temperature.

RESULTS AND DISCUSSION Effect of different gas flow rates on thermochemical reduction reaction. M-Ti4O7 NS were synthesized by thermochemical reduction reaction of A-TiO2 nanoparticles by annealing at various H2 flow rates. The Magnéli TinO2n–1 phase is based on the rutile TiO2 phase. Because the


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rutile/anatase phase boundary appears at 600 °C,21 the anatase phase transforms into the Magnéli phase through the intermediate rutile phase.

nTiO2 + H2 Tin O2n–1 (in H2 gas flow) Thermochemical reduction

The Magnéli TinO2n–1 is formed as two-dimensional chains of edge-sharing TiO6 octahedra, with every nth layer having oxygen deficiency.22 As shown in Figure 1, the two-dimensional chains of octahedra lead to shear planes in the crystal structure to accommodate the loss in stoichiometry. We investigated the intermediate rutile phase at 800 and 850 oC with a H2 flow of 100 sccm. At 800 oC annealing, the XRD pattern of the sample indicates mainly anatase phase containing very small amounts of rutile phase in which the (110) peak is less developed. At 850 oC annealing, the XRD pattern of the sample indicates the mixture of anatase and rutile TiO2, and Magnéli Ti9O17. Consequently, the anatase phase transforms into the Magnéli phase through the intermediate rutile phase (Figure S1 in the Supporting Information). Figure 1 presents the XRD patterns of the samples produced in the H2 flow rate range of 15 to 150 sccm, respectively, at the reaction temperature of 1050 °C for 2 h. The XRD pattern of the sample produced at a H2 flow rate of 15 sccm indicates the formation of various Magnéli phases such as Ti8O15, Ti7O13, and Ti6O11. At a higher H2 flow rate of 25 sccm, a single Ti6O11 phase was identified. In a more reducing atmosphere, a single Ti4O7 phase was identified for a H2 flow rate of 150 sccm. However, phase change to secondary Ti3O5 phase was detected when the H2 flow rate reached 200 sccm. The extent of reduction of the titania increased with an increase in H2 flow rate. The XRD results show that the optimal reduction atmosphere for the synthesis of Ti4O7 phase is the H2 flow rate of 100 sccm.


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To investigate the morphology of the Ti4O7 particles synthesized through thermochemical reduction, a typical FESEM image was collected (Figure S2 in the Supporting Information). As observed, commercial TiO2 nanoparticles had a small particle size of below 25 nm and highly aggregated morphology (Figure S1a in the Supporting Information). However, after thermochemical reduction at 1050 °C, the particles were sintered and grew to a few micrometers (Figure S1b in the Supporting Information).

Synthesis of monodisperse M-Ti4O7 NS. In order to solve issues such as coarsening and sintering of Ti4O7 particles, we synthesized monodisperse A-TiO2 NS and prevented their coarsening by applying a silica coating. Figure 2 schematically illustrates the fabrication of monodisperse M-Ti4O7 NS. The synthesis process involves the following four steps: (i) synthesis of A-TiO2 NS; (ii) silica coating of A-TiO2 NS; (iii) phase transformation from the silica-coated A-TiO2 NS to silica-coated M-Ti4O7 NS by thermochemical reduction under H2 gas flow; and (iv) removal of the silica shell using NaOH solution. In the initial step, A-TiO2 NS were synthesized by a combined sol–gel and heat treatment process. Upon hydrolysis of titanium isopropoxide, the Ti species and their oligomers participate in hydrogen-bonding interactions with the amino groups of hexadecylamine to form inorganic– organic micelles.20 The inorganic–organic micelles contain a hydrophobic long-chain alkyl group that induces self-organization to reduce interfacial free energy. During further hydrolysis and condensation, as the titanium oligomer associated with the inorganic–organic micelles further polymerizes, the condensed phase becomes denser and forms a gel, and finally precipitates in the solvent. The precipitated TiO2 precursors form a sphere to minimize the surface free energy as in


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a typical colloid formation process. Thereafter, we carried out heat treatment to crystallize the obtained amorphous-TiO2 NS into A-TiO2 NS. After the sol–gel process and crystallization, we observed the morphology and crystal structure of the obtained samples by FESEM and XRD (Figure 3). In case of amorphous-TiO2 NS, monodisperse particles are observed in Figure 3a. Amorphous-TiO2 NS are spherical with diameters in the range from 400 to 500 nm and have smooth surfaces. After crystallization, ATiO2 NS retained their particle size, monodispersity, and smooth surfaces compared with the amorphous-TiO2 NS, as seen in the FESEM images (Figure 3b). Figure 3c shows the XRD pattern of the as-synthesized particles. After annealing under air atmosphere, the amorphous phase completely transformed into anatase phase (PDF card No. 21–1272), as evident from the absence of any impurity. In the second and third steps, the crystalline phase of A-TiO2 NS completely transformed into MTi4O7 NS through a thermochemical reduction in H2 atmosphere (H2 flow rate of 100 sccm). ATiO2 NS can be sintered and grown during thermochemical reduction (Figure S3 in the Supporting Information). We carried out silica-protected thermochemical reduction to prevent particle growth and necking of A-TiO2 NS. Figure S4 in the Supporting Information shows the morphologies of silica coated of A-TiO2 NS containing tiny silica spheres. The silica coated of A-TiO2 NS reveals smooth surface without strong agglomeration. After thermochemical reduction, the silica-coated M-Ti4O7 NS exhibited homogeneous core/shell structures (Figure 4a). Moreover, the silica-coated M-Ti4O7 NS were similar in shape to the A-TiO2 NS. The silica-coated M-Ti4O7 NS retained their monodispersity with some shrinkage (diameter of approximately 250 nm) compared with the original A-TiO2 NS. From the typical TEM image of silica-coated M-Ti4O7 NS (Figure 4b), it can be confirmed that silica


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formed a thin shell structure around the M-Ti4O7 NS. The good agreement of the XRD patterns indicates that Ti4O7 is the primary crystalline phase (PDF card No. 72–1723, Figure 4c). In the final step, the silica shell structures were removed by using NaOH solution. M-Ti4O7 NS are composed of nanoparticles (diameter of > 100 nm), as seen from the uneven surface in the FESEM image and the particles in the TEM image (Figure 4d and e). Both the FESEM and TEM images indicate that M-Ti4O7 NS maintained a homogeneous spherical morphology. The XRD pattern of M-Ti4O7 NS was compared with that of standard Ti4O7. The good agreement between the XRD patterns indicates a M-Ti4O7 crystalline phase (PDF card No. 72–1723, Figure 4c). After the removal of silica shell, it was confirmed that the A-TiO2 NS completely converted into Ti4O7 without the formation of a secondary phase. The M-Ti4O7 NS were further investigated in detail using high resolution TEM (HRTEM) and the corresponding fast Fourier transform (FFT) of the electron diffraction pattern (Figure 4f). The HRTEM image indicates that the interplanar distance for the M-Ti4O7 NS is 0.235 nm, which is consistent with the (22 2 ) plane of Ti4O7. The FFT of the electron diffraction pattern could be indexed to the Ti4O7 crystalline phase (PDF card No. 72–1723), which is consistent with the XRD pattern. For further study of surface chemical composition, we performed the XPS analyses of A-TiO2 NS and M-Ti4O7 NS (Figure S5 in the Supporting Information). The structures of Magnéli phase titanium oxides can be treated as a rutile TiO2 derivative (Figure 1). The resultant vacant sites lead to the overlapping of Ti 3d orbitals in the crystal lattices, which formed during reorientation of TiO6 octahedrons, as a consequence of partial reduction of Ti(IV) to Ti(III). The overlap of Ti 3d orbitals has been considered as the main reason for the electrical conduction of oxygen deficient Magnéli phases, thereby broadening the conduction band and allowing electronic conduction.23 We believe that these results could contribute toward the deposition ability of the


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ORR products of the M-Ti4O7 NS electrode due to the corresponding inherent electronic conductivity of the Magnéli phase and the uneven surface effect of the M-Ti4O7 NS. The XPS O 1s spectra of the M-Ti4O7 NS were decomposed into Gaussian functions, which show three binding energies at 529.5, 530.4 and 531.7 eV whereas that of the A-TiO2 NS show two binding energies at 529.9 and 531.5 eV (Figure S5a and b in the Supporting Information).24 The O 1s spectrum in M-Ti4O7 NS represents Ti2O3 at the binding energy of 530.4 eV. Additionally, the Ti 2p spectrum of the A-TiO2 NS show only Ti4+ ions with the binding energy of 458.6 eV whereas that of the M-Ti4O7 NS show two binding energies at 457.9 and 459.3 eV (Figure S5c and d in the Supporting Information).25 The M-Ti4O7 NS samples represent Ti3+ ions binding energy at 457.9 eV. To maintain electrical neutrality, Ti3+ ions are expected to be adjacent to the oxygen deficient sites. In crystallographic point of view, the XPS spectra of the M-Ti4O7 NS can be interpreted as an ordered combination of rutile TiO2 and the oxygen deficient Ti2O3 parts, since not only TiO2 but also Ti2O3 may coexist with the Magnéli phase. Therefore, we can expect the improvement of the electrical conductivity according to the Ti3+ defect-structure of the M-Ti4O7 NS, and it can show higher discharge capacity compared with the A-TiO2 NS. Anatase TiO2 is a chemically stable material, but has lower electrical conductivity (electrical resistivity: 1 × 102 ~ 107 Ω cm) than the Magnéli phase.26,27 The electrical conductivity of Ti4O7 (electrical resistivity: 2.5 × 10–4 Ω cm) is comparable with that of graphite (electrical resistivity: 3.5 × 10–3 Ω cm).27,28 When Magnéli Ti4O7 is applied as electrocatalyst support, it possesses attractive features with respect to chemical stability and electrical conductivity.


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Electrochemical performance of M-Ti4O7 NS. We evaluated the electrochemical performance of oxygen electrodes containing A-TiO2 NS and M-Ti4O7 NS for Li–O2 batteries. As shown in Figure 5a, the initial galvanostatic discharge and charge tests of both the A-TiO2 NS electrode and M-Ti4O7 NS electrode were performed at a current rate of 100 mA g–1 when cycling between 2.0 and 4.5 V. The A-TiO2 NS electrode exhibited a discharge capacity of 523 mA h g–1, and the M-Ti4O7 NS electrode exhibited a discharge capacity of 3,125 mA h g–1, which is approximately 5 times the discharge capacity of A-TiO2 NS electrode. To further investigate the discharge and charge performance, cyclic voltammetry (CV) measurements were made for the A-TiO2 NS and M-Ti4O7 NS electrodes in the voltage window of 2.0–4.5 V at a scan rate of 0.1 mV S–1 (Figure 5b and c). The M-Ti4O7 NS shows strong reduction peaks during the first cathodic scan, indicating the good ORR activity of Li–O2 batteries. In the subsequent scanning cycles, the CV curves clearly show obvious cathodic peaks, suggesting the good reversibility and structural stability of the M-Ti4O7 NS. During the first anodic scan process, an obvious OER peak can be observed at 3.8 V for the M-Ti4O7 NS. During the subsequent anodic scan process the CV curves show stable anodic peaks with relatively lower peak intensities than that in the first cycle. However, no oxidation peak can be observed below 4.0 V for the A-TiO2 NS. Furthermore, the anodic peak intensity of the M-Ti4O7 is much higher than that of the A-TiO2 NS. For the anodic scan process in the subsequent cycles, the MTi4O7 NS still displays apparent oxidation peaks, demonstrating the good reversibility for OER. The CV and galvanostatic full discharge-charge curves confirm that the discharge and charge overpotential of ORR/OER in the Li–O2 battery system with the M-Ti4O7 NS is much lower than that with the A-TiO2 NS (Figure 5a–c). We also compared the discharge-charge cycling abilities of the M-Ti4O7 NS electrode and A-TiO2 NS electrode at a current rate of 100 mA g–1 and a fixed


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capacity regime of 500 mA h g–1, as shown in Figure 5d and e. The overall overpotentials of the M-Ti4O7 NS electrode are smaller than the overpotentials of the A-TiO2 NS electrode during 50 discharge-charge cycles. In the initial discharge curves, the final voltage value of the M-Ti4O7 NS electrode (2.52 V) is remarkably higher than that of the A-TiO2 NS electrode (2.04 V). We believe that these results could contribute toward the deposition ability of the ORR products of the M-Ti4O7 NS electrode due to the corresponding inherent electronic conductivity of the Magnéli phase and the uneven surface effect of M-Ti4O7 NS. Figure S6a-c compares the surface components by the Ti 2p XPS spectrum of the pristine, discharged, and charged M-Ti4O7 NS electrodes. These three XPS core level spectra were decomposed by a Gaussian function. The pristine M-Ti4O7 NS electrode shows two binding energies at 457.9 and 459.4 eV. Both discharged and charged M-Ti4O7 NS electrodes retain two peaks for Ti4+ and Ti3+, but tend to increase the concentration of Ti3+, which appears to result in the formation of surface defect. The negative shift in binding energy of about 1.5 eV indicates the reduction from Ti4+ to Ti3+. The M-Ti4O7 NS has become electrically conductive by introducing Ti3+ into the lattice after discharge and charge cycles. These results suggest that the M-Ti4O7 NS electrodes are stable from intermediate species attack and can support reversible Li2O2 formation/decomposition maintaining stable electrolyte interface during discharge and charge process.17 However, as shown the Ti 2p spectra of the M-Ti4O7 electrodes after 30th cycling (Figure S6d and e in the Supporting Information), the Ti 2p spectra peaks almost disappeared. These results indicate that irreversible ORR products were deposited on the MTi4O7 surface and could cause high overpotential. To demonstrate the enhanced cycling performance, we employed ex situ FESEM to observe the creation and extinction of ORR products such as Li2O2 and Lithium carbonate after 1st


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galvanostatic charge-discharge process of M-Ti4O7 NS electrode (Figure 6a). Figure 6b and c exhibit the typical FESEM images of M-Ti4O7 NS electrode after initial full discharge and charge, respectively. When the M-Ti4O7 NS electrode was fully discharged, it was covered with ORR products. As shown in Figure 6b, the ORR products were formed in sheet-like shape. Lowdonor solvent (DMAc: 26.6, dimethyl ether: 20) showed much weaker solvation capability; therefore, the intermediate LiO2* cannot be stably discharged from the electrode surface to the electrolyte. The surface growth mechanism due to LiO2* and its instability on the electrode surface is due to the rapid secondary reduction to Li2O2 generally occurring at the electrode surface and depositing a thin layer of ORR products. It can be seen that the solvent used in our experiment belongs to DMAc and is produced in sheet-like Li2O2 form.29,30 Compared with the typical toroidal Li2O2 formation, these sheet-like Li2O2 are uniformly distributed on the surface of the electrode structures and have the advantage of promoting charge transfer at the Li2O2/electrode interface.31,32 As shown in Figure 6b, after full charge, all the ORR products disappeared from the surface of the M-Ti4O7 NS electrode. The formation of Li2O2 was confirmed by XRD and TEM analysis of pristine, discharged, and charged M-Ti4O7 NS electrodes. As shown in Figure S7a and b in the Supporting Information, no diffraction peaks for amorphous Li2O2 were detected in XRD patterns of all samples. Also, the TEM images showed that the formation of sheet-like Li2O2 in the oxygen electrode could lead to their growth during the discharge process, as can be seen FESEM images (Figure S7c–e in the Supporting Information). The HRTEM image confirm no crystallinity appears, and corresponds to the the XRD pattern of the discharged M-Ti4O7 NS (Figure S7e in the Supporting Information). To provide chemical evidence for the creation and extinction of ORR products after the discharge and charge processes, we used XPS to examine the ORR products in detail. Figure 6d–


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f shows the Li 1s XPS spectra of the pristine, discharged, and charged electrodes. The Li 1s spectrum for the fully discharged M-Ti4O7 NS electrode separated into two peaks at binding energies of 54.5 and 55.3 eV corresponding to the Li2O2 and Lithium carbonate species, respectively (Figure 6e).33,34 The generation of Lithium carbonate species results from the decomposition of organic electrolyte on the electrode materials.35 As in the Li 1s XPS spectrum of the fully charged electrode (Figure 6f), the peak for Li2O2 disappeared, which confirmed the complete decomposition of Li2O2 on the M-Ti4O7 NS electrode. However, the binding energy peak at 55.3 eV indicated that a small amount of lithium carbonate species did not decompose completely. Superoxide radical anion (O2·-) and lithium superoxide (LiO2) attack the electrolyte and air-electrode during the discharge and charge process under a high overpotential, resulting in the formation of irreversible by-products such as lithium carboxylates and Li2CO3.36-39 This phenomenon is susceptible to nucleophilic attack by the reduced oxygen species (O2·- and LiO2) at the air-electrode, instead of reacting with Li+ cations to form the final reaction products (Li2O2). Lithium carbonate species can be formed by the decomposition of organic electrolytes during charge as well as during discharge.37 To identify the origin of the by-product deposit during the discharge and charge, we measured the nitrogen species bonding derived from 1 M LiNO3-DMAc electrolyte by using XPS and FTIR analysis to demonstrate that a side reaction occurred from the electrolyte (Figure S8 in the Supporting Information). In our experimental, the FTIR spectra of the M-Ti4O7 NS electrodes show the bands around 3200–3500 cm-1 after discharge and charge process, which can be assigned to the NH and CH stretching.40 The observed bands at 3400 cm-1 in the FTIR spectra assigned as the NH stretching. The origin of the NH stretching might be related to the decomposition of LiNO3 or DMAc. The similar behavior was observed in the C 1s, O 1s, and N 1s peaks of corresponding XPS spectra after discharge and


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charge. Therefore, it is believed that the by-products result from the decomposition of the electrolyte during discharge and charge.

RuO2@M-Ti4O7 NS for carbon-free oxygen electrodes. We synthesized RuO2@M-Ti4O7 NS containing the well-known RuO2 electrocatalyst and applied it to carbon-free oxygen electrodes to evaluate the properties of M-Ti4O7 NS as an electrocatalyst support. As shown in the schematic illustration in Figure 7a, we carried out a hydrothermal process for formation of RuO2@M-Ti4O7 NS. To confirm the morphology and crystal structure of RuO2@M-Ti4O7 NS, typical FESEM and TEM images were obtained, as shown in Figure 7b–d. The FESEM image in Figure 7b indicates that the M-Ti4O7 NS maintained their morphology during the hydrothermal process and were decorated with RuO2 nanoparticles on their surfaces. Additionally, the TEM image in Figure 7c shows that the RuO2@M-Ti4O7 NS retained their original spherical shape, and that RuO2 nanoparticles deposited on the surface of the M-Ti4O7 NS. The RuO2 nanoparticles are highly aggregated on the surface of the M-Ti4O7 NS. To further confirm the microstructure and chemical composition of RuO2@M-Ti4O7 NS, STEM-EDS and selected area electron diffraction (SAED) analyses were performed for the particles shown in Figure 7c. As can be seen from the STEM-EDS element mapping profiles (Figure 7d), the yellow points representing Ti-Kα were captured by the existing M-Ti4O7 NS, and the magenta points representing Ru-Lα were captured by the RuO2 nanoparticles. The SAED pattern of RuO2 (inset A) indicates its amorphous nature, and that for M-Ti4O7 NS (inset B) shows a random orientation of the scattered Ti4O7 crystalline phase (PDF card No. 72–1723). The diffraction patterns could be indexed to the (22 2 ) and (14 2 ) crystal planes, which is in good agreement with the XRD patterns (Figure S9 in the Supporting Information). The XRD pattern of RuO2@M-Ti4O7 NS


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confirms a Ti4O7 crystalline phase without any by-products after the hydrothermal process. In addition, no diffraction peaks for amorphous RuO2 were detected. We compared the galvanostatic cycling performance of RuO2@M-Ti4O7 NS electrode and MTi4O7 NS electrode at a current rate of 100 mA g–1 and a fixed capacity regime of 500 mA h g–1 (Figure 8a and b). At the 20th discharge-charge cycle, the overpotential of RuO2@M-Ti4O7 NS electrode (0.9 V) was half of that of the M-Ti4O7 NS electrode (1.8 V). This result is believed to be due to the synergistic effect of the stability of M-Ti4O7 NS as an electrocatalyst support and the high ORR/OER activities of the RuO2 nanoparticles as an electrocatalyst. We investigated EIS analysis for a more secure approach (Figure S10 and Table S1 in the Supporting Information). In Li-O2 battery, the semicircle in the middle frequency region indicates the charge transfer resistance (Rct), which is associated with the charge transfer reaction at the electrode/electrolyte interface.41 Rct is related to the rate coefficient of the chemical reaction in the oxygen electrode. Decreased Rct represents efficient oxygen diffusion pathway and less agglomeration of the oxygen electrode. This means that the electrolyte and the electrode are activated due to intimate contact and effective lithium ions and oxygen diffusion. Therefore, the reduction of Rct in RuO2@M-Ti4O7 NS electrode provides an evidence of activation of electrodes, which leads to lower overpotential upon cycling. In addition, the CV curves of RuO2@M-Ti4O7 NS, M-Ti4O7 NS, and A-TiO2 NS electrodes obtained at a scan rate of 0.1 mV s−1 were compared, as shown in Figure 8c and Figure S11 in the Supporting Information. The RuO2@M-Ti4O7 NS electrode exhibits larger cathodic and anodic peaks than the M-Ti4O7 NS electrode and A-TiO2 NS electrode do. This suggests that the RuO2@M-Ti4O7 NS electrode shows higher ORR/OER activity than the other two electrodes do. However, during charging, the RuO2@M-Ti4O7 NS electrode appears to exhibit side reactions


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that decompose the organic solvent-based electrolyte or Lithium carbonate over 4.1 V. Importantly, the galvanostatic cycling performance of the RuO2@M-Ti4O7 NS electrode was obtained below 4.0 V in the fixed capacity region of 500 mA h g-1, avoiding side reactions, at which it exhibited a stable cycling performance. We carried out the full discharge-charge cycling test of the RuO2@M-Ti4O7 NS electrodes at current rates of 100 and 300 mA g-1. Large specific capacity of over 3500 mA h g-1 are achieved for 15 discharge and charge cycles at a current rate of 100 mA g-1, indicating an apparent ORR and OER catalytic activity of the RuO2@M-Ti4O7 NS electrodes (Figure S12a and b in the Supporting Information). Similar to the M-Ti4O7 NS electrodes, when the RuO2@M-Ti4O7 NS electrode is fully discharged, it is covered with sheet-like ORR products (Figure S13a in the Supporting Information). As shown in Figure S13b, all the ORR products disappeared from the surface of the RuO2@M-Ti4O7 NS electrodes after full charge. Figure S12c and d show the discharge and charge cycling performance of the RuO2@M-Ti4O7 NS electrodes at a current rate of 300 mA g-1, delivering over 500 mA h g-1 even after 70 discharge and charge cycles. Therefore, it is concluded that the RuO2@M-Ti4O7 NS have high capacity and stable cycle performance as carbon-free oxygen electrodes.

CONCLUSION We propose a carbon-free RuO2@M-Ti4O7 NS as a concept to replace the prevailing carbonbased oxygen electrodes to enhance the performance of Li–O2 batteries. For the low overpotential of Li–O2 batteries (< 1.0 V for discharge and charge voltage gap), the instability of carbon materials will pose a significant problem. Our approach is focused on other conductive


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and stable materials to support the ORR/OER activity of electrocatalysts. The outstanding stability of the M-Ti4O7 NS electrode compared to the A-TiO2 NS electrode is responsible for the low overpotential and high specific capacity during discharge-charge cycles, as verified by comparing the CV and galvanostatic cycle curves of M-Ti4O7 NS and A-TiO2 NS. The M-Ti4O7 NS support the activation of catalytic ability of the RuO2 nanoparticles during the discharge-charge processes. Li–O2 batteries containing the RuO2@M-Ti4O7 NS as oxygen electrodes exhibited enhanced performance with a low overpotential of 0.9 V. We believe that the carbon-free RuO2@MTi4O7 NS are promising for superior ORR/OER performance for Li–O2 batteries.


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FIGURES

Figure 1. XRD patterns of commercial TiO2 and Magnéli phase TinO2n-1 prepared by thermochemical reduction process at different H2 flow rates. Crystal lattice illustrations of rutile TiO2 phase and Magnéli Ti4O7 phase.


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Figure 2. Schematic illustration of fabrication of M-Ti4O7 NS.


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Figure 3. FESEM images of (a) amorphous-TiO2 NS and (b) A-TiO2 NS. (c) XRD patterns of amorphous-TiO2 NS and A-TiO2 NS.


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Figure 4. (a) FESEM and (b) TEM images of silica-coated M-Ti4O7 NS. (c) XRD patterns of silicacoated M-Ti4O7 NS and M-Ti4O7 NS. (d) FESEM, (e) TEM, and (f) HRTEM images of M-Ti4O7 NS. The inset of (f) is FFT diffraction pattern of M-Ti4O7 NS corresponding to the yellow circle in (e).


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Figure 5. (a) Galvanostatic full discharge-charge curves for A-TiO2 NS electrode and M-Ti4O7 NS electrode. CV curves of (b) A-TiO2 NS and (c) M-Ti4O7 NS electrodes. Galvanostatic dischargecharge curves of (d) A-TiO2 NS electrode and (e) M-Ti4O7 NS electrode.


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Figure 6. (a) Galvanostatic full discharge-charge curves of M-Ti4O7 NS electrode at a current density of 100 mA g–1. FESEM images of (b) discharged and (c) charged M-Ti4O7 NS electrodes. (b) corresponds to the red circle and (c) corresponds to the blue circle in (a). Li 1s XPS spectra of (d) pristine, (e) discharged, and (f) charged M-Ti4O7 NS electrodes.


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Figure 7. (a) Schematic illustration of fabrication of RuO2@M-Ti4O7 NS. (b) FESEM and (c) TEM images of RuO2@M-Ti4O7 NS. (d) EDS element mapping profile and SAED patterns of RuO2@MTi4O7 NS. The SAED patterns correspond to zone ‘A’ (red square) and zone ‘B’ (blue square) in (c), respectively.


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Figure 8. Galvanostatic discharge-charge curves of (a) M-Ti4O7 NS electrode and (b) RuO2@MTi4O7 NS electrode. (c) Comparison of CV curves of RuO2@M-Ti4O7 NS electrode, M-Ti4O7 NS electrode, and A-TiO2 NS electrode.


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ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge on the ACS Publications website at DOI: Representative XRD patterns of commercial TiO2 by thermochemical reduction process at 800 and 850 °C; representative FESEM images of commercial TiO2 nanoparticles and Ti4O7 particles; representative FESEM images of sintered Ti4O7; representative FESEM images of silica coated A-TiO2 NS; representative XPS spectra of M-Ti4O7 NS and A-TiO2 NS; representative XPS spectra M-Ti4O7 NS electrode; representative XRD patterns and TEM images of M-Ti4O7 NS electrode; representative FTIR and XPS spectra of M-Ti4O7 NS electrode; representative XRD patterns of M-Ti4O7 NS and RuO2@M-Ti4O7 NS; nyquist plots of RuO2@M-Ti4O7 NS electrode; Fitting values of RuO2@M-Ti4O7 NS electrode; cyclic voltammetry measurements of RuO2@M-Ti4O7 NS, M-Ti4O7 NS and A-TiO2 NS electrode; galvanostatic full discharge-charge cycle of RuO2@M-Ti4O7 NS; representative FESEM images of RuO2@M-Ti4O7 NS electrodes.(PDF)

AUTHOR INFORMATION Corresponding Author * E-mail: [email protected] ORCID Dong-Wan Kim: 0000-0002-1635-6082 Notes


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The authors declare no competing financial interest. Author Contributions ‡S. L., G.H.L. contributed equally to this work. S. L. performed, analyzed the experiments. G.H.L. performed theoretical analysis and drafted the manuscript. J.C.K. carried out drawing illustrations for experimental concepts. D.W.K. conceived and designed the study, led the discussion of the results, and performed the final edits of the manuscript. All authors made critical contributions to the work, discussed the results and commented on the manuscript.

ACKNOWLEDGMENT This work was supported by the National Research Foundation of Korea (NRF) Grant funded by the Ministry of Science, ICT, and Future Planning [NRF-2016R1A2B2012728, NRF2016M3A7B4909318] and by the institutional research program of the Korea Institute of Science and Technology [2E26081-16-054]. This work was also supported by a Korea University Grant.


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s for Nonaqueous Lithium-Oxygen Batteries. Sci. Rep. 2017, 7, 9495.


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TOC

As a concept to replace the prevailing carbon-based oxygen electrodes, we propose a carbon-free RuO2@M-Ti4O7 nanospheres containing oxygen electrode for Li–O2 batteries.


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MagnÃ©li-Phase Ti4O7 Nanosphere Electrocatalyst Support for Carbon

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