Ultrathin Porous NiCo2O4 Nanosheets for Lithium–Oxygen Batteries

May 22, 2019 - The adsorption energies of O2 were calculated by using the Vienna ab ... The GGA + U calculations were performed on the NiCo2O4 system ...
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Cite This: ACS Appl. Energy Mater. 2019, 2, 4215−4223

Ultrathin Porous NiCo2O4 Nanosheets for Lithium−Oxygen Batteries: An Excellent Performance Deriving from an Enhanced Solution Mechanism Xin Guo, Jinqiang Zhang, Yufei Zhao, Bing Sun,* Hao Liu,* and Guoxiu Wang* Centre for Clean Energy Technology, School of Mathematics and Physics, Faculty of Science, University of Technology Sydney, Broadway, Sydney, NSW 2007, Australia Downloaded via NOTTINGHAM TRENT UNIV on August 10, 2019 at 06:04:45 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

S Supporting Information *

ABSTRACT: Lithium−oxygen batteries are of interest for long-range electric vehicles owing to their high theoretical energy density. However, the poor cycling performance and low round-trip efficiency deriving from the sluggish oxygen reduction reaction (ORR) and oxygen evolution reaction (OER) kinetics severely impede their practical application. Ingenious design of cathode catalysts is imperative to overcome these challenges. Here, we report ultrathin porous NiCo2O4 nanosheets with abundant oxygen vaccines as an efficient cathode catalyst toward both OER and ORR for Li−O2 batteries. From combined theoretical calculation with experimental results, a unique enhanced solution mechanism is proposed in the ether-based electrolyte system. Benefiting from the porous 2D architecture of the cathode and the hierarchical toroidal products, the Li−O2 batteries using NiCo2O4 cathodes deliver a high discharge capacity of 16 400 mAh g−1 at 200 mA g−1 and an excellent cycling performance up to 150 cycles with a restricted capacity of 1000 mAh g−1. KEYWORDS: 2D materials, oxygen vacancies, OER/ORR catalyst, Li−O2 batteries, Li2O2 growth mechanism

1. INTRODUCTION The crisis of fossil fuel shortage and climate change has spurred the rapid development of clean energy technologies. Lithium−oxygen batteries, with the highest theoretical capacity among all rechargeable batteries, have shown great potential to fulfill the increasing demands from electric vehicles (EVs).1−3 However, suffering from the poor cycling performance and low round-trip efficiency, lithium−oxygen batteries are still far away from the practical application. Lithium peroxide (Li2O2), the discharge product of the oxygen reduction reaction (ORR) in most Li−O2 batteries, is a wide-bandgap insulator hindering charge transfer from the cathode to the Li2O2−electrolyte interface and leading to electrical passivation of the electrodes, which limits the overall discharge capacity to much lower than the theoretically possible.4,5 In addition, the produced Li2O2 with a sluggish oxygen evolution reaction (OER) kinetics requires a high overpotential to be decomposed reversibly; meanwhile, Li2O2 is also water-sensitive and readily induces parasitic chemistries, which decreases the efficiency and rechargeability of the batteries.6,7 Understanding and controlling of the Li2O2 formation process are essential to overcome these challenges.8−11 Currently, there are two well-recognized ORR mechanisms: the surface growth route and the solution growth route.12 In the surface electrochemical mechanism, Li2O2 grows directly on cathode surfaces as thin conformal films, leading to low discharge capacities. In the solution electrochemical process, © 2019 American Chemical Society

large Li 2 O 2 toroids grow in electrolyte solution via disproportionation of the soluble intermediate LiO2, rendering significantly enhanced discharge capacities.13,14 Therefore, Li− O2 batteries based on the solution mechanism are more favorable than those depending on the surface pathways.15,16 However, the formed toroidal aggregates via solution pathway are normally too large to be decomposed readily, reflecting by high charge overpotentials and fast decaying rates. Tremendous efforts have been made to address these issues, including cathode architectures design, selection of efficient OER catalysts, and exploration of electrolyte additives.17−26 Among these approaches, rational design for catalytic cathode has proven to be crucial to battery performances. In particular, transition metal oxides are attractive for practical Li−O2 batteries due to their low costs and good catalytic activities for both ORR and OER.27−34 Herein, we fabricated an ultrathin porous NiCo 2 O 4 nanosheets with oxygen vacancies as a high efficient cathode catalyst for Li−O2 batteries. In the meantime, a unique enhanced solution mechanism is established to account for the excellent electrochemical performance. Specifically, due to the strong adsorption strength between O2 and NiCo2O4 nanosheets, small toroidal arrays preferentially formed on the Received: March 2, 2019 Accepted: May 21, 2019 Published: May 22, 2019 4215

DOI: 10.1021/acsaem.9b00450 ACS Appl. Energy Mater. 2019, 2, 4215−4223

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Figure 1. (a) Low-magnification and (b) high-magnification SEM images of NiCo2O4 material. (c) TEM image, (d) HRTEM image, and (e) corresponding SAED pattern of the NiCo2O4 sheets. (f) STEM image and EDS mapping analysis of the (g) Ni, (h) Co, and (i) O elements. (j) AFM image of NiCo2O4 nanosheets and the corresponding amplitude profile for the selected red line. The Brunauer−Emmett−Teller (BET) method was employed to compute the specific surface area. X-ray photoelectron spectroscopy (XPS) measurements were conducted on an ESCALAB250Xi (Thermo Scientic, UK) equipped with monochromated Al Kα (energy 1486.68 eV). 2.3. Electrochemical Measurements. The cathodes were prepared by the following procedure. First, a uniform slurry containing 50 wt % of as-prepared catalysts, 40 wt % of carbon black, and 10 wt % of poly(tetrafluoroethylene) in isopropanol was pasted on glass fiber disks. After drying at 110 °C in a vacuum oven for 12 h, the electrodes were quickly moved into an Ar-filled glovebox (Mbraun, H2O ≤ 1 ppm and O2 ≤ 1 ppm) for further usage. The typical loading of the electrodes is about 1 mg cm−2. Swagelok-type cells with an oxygen hole on the cathode side were used to evaluate the electrochemical performance. The Li−O2 cells were assembled in an glovebox using lithium foil as the anode, a glass microfiber as separator (Whatman), and 1 M LiTFSI in diethylene glycol dimethyl ether (DEGDME) or1 M LiClO4 in dimethyl sulfoxide (DMSO) as electrolytes. The as-prepared glass fiber electrodes with active materials were pressed on a stainless steel mesh to ensure the good conductivity of the cells. The cell was gastight except for the stainless steel mesh window that exposed the cathode to the dry oxygen atmosphere (1 atm). Galvanostatic discharge/charge was performed on a Neware battery testing system. The current densities and specific capacities were calculated based on the total mass of active material in the electrodes. 2.4. Computational Method. The adsorption energies of O2 were calculated by using the Vienna ab initio simulation package (VASP), employing the density functional theory (DFT) and the projected augmented wave (PAW) method.35,36 The Perdew− Burke−Ernzerhof (PBE) functional was used to describe the exchange and correlation effect.37 The cutoff energy was set to be 450 eV for all the geometry optimizations. The Monkhorst−Pack grids were set to be 4 × 3 × 1, 7 × 6 × 1, 4 × 5 × 1, and 4 × 4 × 1 for the calculations on NiCo2O4(220), Li2O2(100), Li2O2(101), and single-layer

NiCo2O4 surface at a low discharge depth. The formed toroids with sub-100 nm size have an intimate contact with the cathode material and thus can be decomposed easily during charging, which lowers the overpotential and improves the cycling performance of the batteries. Large toroidal products were subsequently grown on the small Li2O2 arrays through the solution pathway along with the increasing discharge depth, leading to an extremely high discharge capacity as well.

2. EXPERIMENTAL SECTION 2.1. Preparation of the Cathode Materials. All chemicals used in the experiment are from Sigma-Aldrich, and the quality is analytical grade. First, 2.1 g of cobalt(II) nitrate hexahydrate was dissolved in 25 mL of methanol. Then benzyl alcohol (3.5 mL) and urea (0.2 g) were added to the mixture under stirring. Second, the well-mixed solution was transferred into a Teflon-lined stainless steel autoclave, which was subsequently kept at 180 °C for 18 h. The resulting precipitate was collected, washed with methanol and deionized water several times, and dried at 80 °C for 12 h. Finally, the dried precursor was calcined in air with a moderate heating rate of 5 °C min−1 to 450 °C for 2 h. 2.2. Characterization. Field-emission scanning electron microscopy (FE-SEM, Zeiss Supra 55VP) and scanning transmission electron microscopy (STEM, JEOL JEM-ARM200F) were employed to observe the morphologies of samples. The energy dispersive X-ray spectroscopy (EDX) results were also acquired by the STEM. The topography of material was measured by atomic force microscopy (AFM) using a Bruker Dimension 3100 instrument. X-ray diffraction was performed on a Bruker D8 Discovery X-ray diffractometer using Cu Kα radiation (λ = 0.154 06 nm). Before conducting of the ex- itu XRD characterization for the postcycled electrodes, the cycled electrode was washed with corresponding solvents and then dried and sealed with “Parafilm” to exclude the influence of air and moisture on the electrodes. Nitrogen-sorption measurements were carried out at 77 K with a Micromeritics 3Flex surface characterization analyzer. 4216

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Figure 2. (a) XRD pattern of the as-synthesized NiCo2O4 porous sheets. (b) Nitrogen adsorption/desorption isotherms and pore size distribution curve (inset) of NiCo2O4 materials. (c) XPS survey and (d) O 1s spectra of NiCo2O4 nanosheets. graphene, respectively. The GGA + U calculations were performed on the NiCo2O4 system using the model proposed by Dudarev et al.38 with the Ueff (Ueff = Coulomb U − exchange J) values of 6.0 and 3.3 eV for Ni and Co, respectively. To prevent vertical interactions between the slabs, more than 16 Å vacuum layer was applied in the zdirection of the slab models. The adsorption energy of O2 was defined as

diffraction (SAED) pattern is a characteristic spot pattern, suggesting that the monolayer NiCo2O4 sheet is single crystalline. The diffraction spots can be well indexed along the [112] zone axis of spinel NiCo2O4 (Figure 1e). STEM image and corresponding energy dispersive spectrum (EDS) mapping results confirm the components of the sheet, indicating the homogeneous distribution of Ni, Co, and O elements (Figure 1f−i). Atomic force microscopy (AFM measurement) was performed to determine the thickness and topography of the NiCo2O4 nanosheets. As shown in the Figure 1j, we can see the thickness of the as-prepared holey NiCo2O4 sheet is around 1.1 nm and the diameter of holes varies from several nanometers to 100 nm, which is consistent with the SEM and TEM observations (Figure S3). Figure 2a shows the XRD pattern of the as-prepared NiCo2O4 porous nanosheets, which can be well indexed as spinel NiCo2O4 phase (PDF no. 20-0781). The intensities of the peaks are relatively low since the porous sheets lack longrange order in the third dimension. Nitrogen adsorption and desorption measurements were used to identify the porous structure of the product. As shown in Figure 2b, the BET surface area of NiCo2O4 materials deduced from the sorption curves is 156.6 m2 g−1, which is double more than that of the corresponding precursor (Figure S4, 66.1 m2 g−1). The pore sizes calculated by the Barret−Joyner−Halenda (BJH) method are mainly in the range of 2.2−14 nm, and the pore volume is 0.57 cm−3 g−1. The porous architecture of NiCo2O4 with high surface area is able to facilitate electrolyte impregnation at the cathode side and provide plentiful active sites to catalyze the ORR and OER processes during battery cycling. Meanwhile, the high pore volume can afford ample space to accommodate the discharge products. X-ray photoelectron spectroscopy

ΔEads = Eads − Eslab − EO2 where Eads is the electronic energy of the slab with an adsorbed O2, Eslab is the electronic energy of the clean surface, and EO2 is the electronic energy of gaseous oxygen molecule. Under this definition, a more negative value indicates a stronger binding system.

3. RESULTS AND DISCUSSION The morphology and structure of as-prepared NiCo2O4 were characterized by scanning electron microscopy (SEM) and transmission electron microscopy (TEM) (Figure 1), showing a typical two-dimensional lamellar morphology. Numerous mesopores can be observed from the SEM images and lowmagnification TEM image (Figure 1a−c), which were generated by the structural shrinkage and release of gas from the layer precursor during the postcalcination step (Figure S1). Notably, a moderate thermal treatment condition is imperative to the preparation of the porous NiCo2O4 sheets. Both high heating rates and long sintering time were detrimental to the formation of desired porous lamellar structure and resulted in the collapse of the 2D structure (Figure S2). The high resolution TEM lattice image of the formed NiCo2O4 nanosheet (Figure 1d) displays a well-defined crystalline structure, in which the lattice spacing is 0.286 nm with an angle of 60°, corresponding to the (2̅02) and (02̅2) facets of cubic spinel NiCo2O4.39,40 The selected area electron 4217

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Figure 3. Electrochemical performance of the Li−O2 batteries with the DEGDME-based electrolyte. (a) Full range discharge tests of NiCo2O4 and CB cathodes at 200 mA g−1. (b) First discharge/charge curves of NiCo2O4 and CB cathodes with a restricted capacity of 1000 mAh g−1. (c) Selected discharge/charge curves of NiCo2O4 cathodes at 400 mA g−1. (d) Cycling performance of the NiCo2O4 and CB electrodes with a curtailing capacity of 1000 mAh g−1.

electrocatalytic performance of NiCo2O4 nanosheets for both OER and ORR could be attributed to the rich oxygen vacancies and the unique porous 2D structure, which expose more electrochemically active sites. The electrochemical performance of the NiCo2O4 cathodes was subsequently examined in nonaqueous Li−O2 batteries. Figure 3a exhibits the discharge curves of Li−O2 batteries using NiCo2O4 and carbon black (CB) cathode in the DEGDME-based electrolyte. Both electrodes show similar discharge plateaus at 2.75 V; however, the NiCo2O4 cathode delivers a capacity of 16 400 mAh g−1 at 200 mA g−1, which is much higher than that of the CB cathode (6562 mAh g−1). Besides this, it can be seen that the charge overpotential of NiCo2O4 electrode is also much lower than that of the CB electrodes at a fixed capacity of 1000 mAh g−1 (Figure 3b), indicating the good OER catalytic activity of NiCo2O4 nanosheets in the Li−O2 batteries. Figure 3c exhibits discharge−charge profiles of the NiCo2O4 cathode in the selected cycles, which are almost overlapped within the initial 60 cycles, followed by a gradual increase of the potential gap up to 150 cycles. Consequently, the NiCo2O4 electrode achieves a 6 times longer life span than the CB electrode under the same test conditions (Figure 3d). On the contrary, the NiCo2O4 material obtained at 450 °C for 6 h, which has fewer oxygen vacancies (Table S1) and collapsed nanosheets structures, only survives 43 cycles with larger polarizations under the same test conditions (Figure S8). To exclude the impact of CB and further verify the good catalytic activities of the porous NiCo2O4 nanosheets, we assembled the Li−oxygen cells using pure NiCo2O4 as the cathode material. As shown in Figure S9, the pure NiCo2O4 electrode also presents much higher round-trip efficiencies than that of CB electrodes.

(XPS) measurements were further performed to investigate chemical composition and surface state of the NiCo2O4 material. The full survey (Figure 2c) confirms the existence of Ni, Co, and O elements, and the high resolution of O 1s spectrum (Figure 2d) demonstrates the creation of superficial oxygen vacancies on the porous NiCo2O4 nanosheets. The O 1s peak deconvolution yields three peaks at 529.5, 531.2, and 532.8 eV, corresponding to the M−O bonds (M = Ni, Co), oxygen vacancies, and adsorbed hydroxyl species on the surface, respectively.41,42 The Co and Ni 2p spectra fitting results reveal that the NiCo 2 O 4 material contains a considerable proportion of Ni2+ and Co2+ (Figure S5), which can be ascribed to the short-term heat treatment for the preparation of the NiCo2O4. The low valence of Ni and Co can contribute to electron transport, and the oxygen vacancies could significantly improve the catalytic activity of NiCo2O4 material.43 The catalytic activity and durability of the NiCo2O4 material for OER were first evaluated in a three-electrode system as a reference. As shown in the Figure S6a, the NiCo2O4 material exhibits a small onset potential of 1.54 V vs RHE and a much smaller Tafel slope of 60.56 mV dec−1 compared with the previously reported NiCo2O4 nanowire arrays (90 mV dec−1).44 The small Rct value of 61.2 Ω at 600 mV in the Nyquist plot indicates a relative fast kinetics of electrocatalysis reaction (Figure S6b). Chronopotentiometry tests presented no obvious change of the potential with time, indicating the good OER stability of the porous NiCo2O4 nanosheets. (Figure S6c). The ORR catalytic activity was assessed by rotating disk electrode measurements in 1 M KOH solution, and a high limiting current density of 3.45 mA cm−2 was achieved at a potential of 0.1 V vs RHE (Figure S7). The high 4218

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Figure 4. SEM images of NiCo2O4 electrodes discharged to (a) 2000 mAh g−1 and (b) 4000 mAh g−1 at 200 mA g−1. (c) Calculated oxygen adsorption energy on NiCo2O4|Ov, NiCo2O4, Li2O2, and CB; insets are the corresponding adsorption configurations. (d) Schematic representation of the electrochemical double layer at the interface of NiCo2O4 cathode and electrolytes. Section ① shows the one electron product superoxide chemisorbed to the NiCo2O4 catalyst in the inner Helmholtz plane (IHP). Section ② shows the outer Helmholtz plane one-electron charge transfer (OHP). (e) Schematic illustration of the probable mechanism of Li2O2 growth in the lithium−oxygen battery system.

S10). In contrast, similar toroidal-shaped discharge products but with lager size were observed in the pure CB electrode at the same discharged status (Figure S11). The toroidal-shaped discharge products with bigger size were preferentially formed on Li2O2 arrays instead of the carbon black additive when the battery using the NiCo2O4 cathode was discharged deeper to 4,000 mAh g−1 (Figure 4b). To assist in analyzing the observed phenomenon, we need to fundamentally understand the Li2O2 growth routes in the nonaqueous system.45 First, oxygen is caught by the active sites on the cathode and formed adsorbed oxygen (denoted as O2*).

However, due to the poor conductivity of the pure NiCo2O4 materials, the cells can be only discharged and charged at a low current density of 50 mAh g−1. Therefore, it can be deduced that the superior electrochemical performance of the NiCo2O4 catalyst could be attributed to the contribution of oxygen vacancies to good catalytic activities toward ORR/OER kinetics and the unique porous 2D architecture of the material as well. To unravel the mechanism underpinning the good electrochemical performance of the NiCo2O4 cathode catalyst, ex situ SEM characterization was conducted to observe the morphologies of the discharged products. Figure 4 shows SEM images of the NiCo2O4 cathodes at selected discharge states in the DEGDME-based electrolyte. When batteries were discharged to 2000 mAh g−1, homogeneous small toroids were found to be vertically grown on the NiCo2O4 nanosheets, blocking the pores of the 2D cathode material (Figures 4a and

O2 + S = O2 *

(1)

Second, the adsorbed oxygen accepts one electron and coupled with Li+ to form LiO2*. O2 * + e− + Li+ = LiO2 * 4219

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Figure 5. (a) XRD patterns of (i) fully discharged NiCo2O4 cathode, (ii) charged NiCo2O4 cathode after a full discharge−charge cycle, and (iii) NiCo2O4 cathode after 60 cycles at 400 mA g−1 with a restricted capacity of 1000 mAh g−1. (b) SEM, (c) TEM, and (d) HRTEM images of the NiCo2O4 cathode material after 60 cycles at 400 mA g−1 with a restricted capacity of 1000 mAh g−1.

Then they can directly accept one electron from the cathode in the inner Helmholtz plane (IHP), leading to the growth of tiny toroidal arrays on the NiCo2O4 surface (section ①, Figure 4d). The formed arrays have close contact with cathode surface and can be easily decomposed during the charging process, resulting in the low charge overpotential when batteries are cycling at a restricted capacity of 1000 mAh g−1. As the discharge process goes deeper, the increasing amount of Li2O2 on the electrode surface weakens the adsorption strength of NiCo2O4 toward oxygen/superoxides, and dissolved oxygen in the electrolyte can directly trap an electron from the outer Helmholtz plane (OHP) (section ②, Figure 4d).46 The conventional solution reaction starts to take over the discharge process and form large toroidal particles, leading to the high discharge capacity (Figure 4e). Moreover, the large toroidalshaped Li2O2 becomes the primary discharge product when the battery is fully discharged to 2.0 V (Figure S12). The ex situ XRD characterizations demonstrate that the dominating discharge product in the Li−O2 batteries was Li2O2, as shown in Figure 5a. No other diffraction peaks belonging to side products were detected probably because of their low percentage in the electrodes. After the batteries were fully charged to 4.3 V, the characteristic peaks of Li2O2 disappeared, indicating that most of the discharge products were reversibly decomposed. The major diffraction peak of NiCo2O4 can be found even after 60 cycles with a restricted capacity of 1000 mAh g−1, demonstrating the good structural stability of the cathode catalyst. Ex situ XRD, SEM, and TEM characterizations exhibit that the porous nanosheet structure of NiCo2O4 cathodes were mostly reserved in spite of the

Finally, the Li2O2 is generated through the further reduction process or disproportionation of the soluble intermediate LiO2. LiO2 * + Li+ + e− = Li 2O2

or 2LiO2 = Li 2O2 + O2 (3)

The O2 adsorption is the first step for the ORR process, which determines the sites of occurrence of the subsequent steps. The observed phenomenon could be ascribed to the difference of oxygen adsorption abilities of the different components in the Li−O2 cells. Therefore, we evaluated the O2 adsorption energy of NiCo2O4 with oxygen vacancies (denoted as NiCo2O4|Ov), NiCo2O4, Li2O2, and carbon black by DFT calculation. As shown in Figure 4c, NiCo2O4|Ov has the strongest strong chemisorption capability with an oxygen adsorption energy (AE) of −2.68 eV, followed by the NiCo2O4 (AE = −1.34 eV), indicating the better ORR catalytic activities of NiCo2O4 with oxygen vacancies compared with the pure one. The catalytic activity of pure NiCo2O4 originates from the exposed Ni2+/Co3+ redox couple. A much weaker oxygen adsorption (AE = −0.08 eV) occurs on the Li2O2 surface. However, this value is still much lower than that of CB (AE = 0.88 eV), which could be the reason that large toroids prefer to grow on the formed Li2O2 arrays instead of the CB surface after the NiCo2O4 catalyst surface is fully covered by the formed tiny toroidal arrays. On the basis of the aforementioned results, we deduce plausible mechanisms for the NiCo2O4 cathode in the etherbased electrolyte. At the initial stage of the discharge process, the fully exposed porous NiCo2O4 nanosheets have strong adsorption energy to capture the O2 species on the surface. 4220

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Last, the selection of DEGDME-based electrolyte with weaker dissolution of lithium superoxide also contributes to better performance compared with the highly soluble and unstable DMSO-based electrolyte. As a result, the Li−O2 batteries can achieve a high capacity and good cycling performance simultaneously.

existence of some splinters (Figure 5a−c). A more interesting finding is that the edge of NiCo2O4 sheets changes from the crystalline structure (Figure S3b) to amorphous phase (Figure 5d). The EDX mapping results confirm that the cycled nanosheets including the amorphous layer are also composed of Ni, Co, and O elements (Figure S11). We deduced that the long-term exposure to catalytic potential presumably led to the partial conversion of crystalline NiCo2O4 to amorphous phase near the active surface. A similar phenomenon was reported in Co-containing catalysts for water oxidation, and the structural disorder normally improves the catalytic activity because of the presence of more defects, vacancies, or enhanced oxidation state of metal ions.47,48 It is worth mentioning that the electrolytes have a significant impact on the electrochemical performance. We compared the performance of NiCo2O4 electrode in DMSO-based electrolyte (NiCo2O4−DMSO) with the above used DEGDME-based electrolyte (NiCo2O4−DEGDME). The overpotential of NiCo2O4−DMSO is lower than that of the CB electrode, demonstrating the universal positive effects of NiCo2O4 catalyst on the OER kinetics in both kinds of electrolytes (Figure S12). However, the overpotential is higher than that of NiCo2O4−DEGDME and the cycling performance was diminished, which only sustained about 40 cycles with a restricted capacity of 1000 mAh g−1 (Figure S13). Ex situ SEM characterization of postcycled electrodes was performed to clarify the reason for the worse performance using the DMSO electrolyte compared with the DEGDME-based electrolyte. Figure S14 shows that the toroidal-shaped particles formed in the DMSO-based electrolyte are inhomogeneous and the average size of the discharge products is much larger than the particle size of the NiCo2O4−DEGDME. The discharge products grew even bigger during deeper discharging and severely destroyed the cathode structure, which caused the poor electrochemical performance of NiCo2O4−DMSO. The DMSO has a strong solvation of LiO2* species, which might surpass the O2 adsorption energy of the NiCo2O4 catalyst, with the result that discharge products continuously grow in the solution apart from the NiCo2O4 nanosheets. Consequently, the toroidal aggregates with micrometre sizes are too large to decompose in the charging process, inducing a higher charge overpotential and poorer cycling performance of NiCo2O4− DMSO than that of NiCo2O4−DEGDME. Another reason probably leading to the worse performance might be the worse stability of DMSO, which is a high soluble solvent and vulnerable to superoxides.49,50 Therefore, it has been proved that the DEGDME-based electrolyte is better than the DMSObased electrolyte for NiCo2O4 cathode in the Li−O2 battery system. Overall, the excellent performance of NiCo2O4 cathodes can be ascribed to the porous 2D architecture of NiCo2O4 and the unique enhanced solution mechanism in the suitable DEGDME-based electrolyte. First, the ultrathin porous architecture facilitates the electrolyte impregnation and oxygen diffusion. Meanwhile, the high surface area of the NiCo2O4 nanosheets with oxygen vacancies enables ample active sites for catalyzing the ORR and OER reactions. Second, the strong adsorption energy of NiCo2O4 triggers the growth of small toroidal arrays on the cathode surface under a restricted capacity, which can be easily decomposed during charging. Third, the enhanced solution mechanism results in the formation of hierarchical Li2O2 toroids and delivers a high discharge capacity, which prevent the premature death of cells.

4. CONCLUSION In summary, ultrathin 2D porous NiCo2O4 materials with abundant oxygen vacancies were prepared as cathode catalysts for high performance Li−O2 batteries. The porous 2D architecture with high surface area is beneficial for electrolyte impregnation, and the strong absorption strength of NiCo2O4 to oxygen contributes to the enhanced solution mechanism in the ether-based electrolyte. Consequently, the Li−O2 batteries using porous NiCo2O4 nanosheets catalyst achieve a high discharge capacity, improved round-trip efficiency, and excellent cycling performance with a fixed capacity of 1000 mAh g−1. The findings in this work are expected to shed light on the mechanisms underpinning Li−O2 batteries and would be instructive for the design of advanced cathodes for Li−O2 batteries.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsaem.9b00450. SEM, HRTEM, and STEM images; adsorption/ desorption isotherms; pore size distribution; XPS spectra; OER catalytic performance; ORR polarization data; charge/discharge curves; calculated oxygen vacancies; cycling performance (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*B.S.: e-mail, [email protected]. *H.L.: e-mail, [email protected]. *G.W.: e-mail, [email protected]. ORCID

Jinqiang Zhang: 0000-0001-5476-0134 Hao Liu: 0000-0003-0266-9472 Guoxiu Wang: 0000-0003-4295-8578 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by the Australian Research Council (ARC) through ARC DECRA project (Grant DE180100036), Future Fellow projects (Grant FT180100705), and Discovery projects (Grants DP160104340, DP170100436, and DP180102297). The authors acknowledge TEM support and use of the JEOL JEM-ARM200F microscope within the University of Wollongong (UoW) Electron Microscopy Centre.

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REFERENCES

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DOI: 10.1021/acsaem.9b00450 ACS Appl. Energy Mater. 2019, 2, 4215−4223

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DOI: 10.1021/acsaem.9b00450 ACS Appl. Energy Mater. 2019, 2, 4215−4223