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Insights into the effects of zinc doping on structural phase transition of P2-type sodium nickel manganese oxide cathodes for high-energy sodium ion batteries Xuehang Wu, Gui-Liang Xu, Guiming Zhong, Zhengliang Gong, Matthew J McDonald, Shiyao Zheng, Riqiang Fu, Zonghai Chen, Khalil Amine, and Yong Yang ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b06701 • Publication Date (Web): 05 Aug 2016 Downloaded from http://pubs.acs.org on August 7, 2016

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ACS Applied Materials & Interfaces

Insights into the effects of zinc doping on structural phase transition of P2-type sodium nickel manganese oxide cathodes for high-energy sodium ion batteries †,∥

Xuehang Wu,

§

Gui-Liang Xu,

†

†

Guiming Zhong,

†

Zhengliang Gong, §

‡

Matthew J. §

McDonald, Shiyao Zheng, Riqiang Fu,┴ Zonghai Chen, Khalil Amine, and Yong †,‡

Yang*, †

State Key Laboratory Physical Chemistry Solid Surfaces, Department of Chemistry, Xiamen University, Xiamen, Fujian 361005, China ‡

§

School of Energy Research, Xiamen University, Xiamen 361005, China

Chemical Sciences and Engineering Division, Argonne National Laboratory, 9700 South Cass Avenue,

Argonne, Illinois 60439, USA ∥

Collaborative Innovation Center of Renewable Energy Materials, Guangxi University, Nanning, Guangxi 530004, China ┴

National High Magnetic Field Laboratory, 1800 E. Paul Dirac Drive, Tallahassee, Florida

32310, United States

Abstract: P2-type sodium nickel manganese oxide-based cathode materials with higher energy densities are prime candidates for applications in rechargeable sodium ion batteries. A systematic study combining in-situ high energy X-ray diffraction (HEXRD), ex-situ X-ray absorption fine spectroscopy (XAFS), transmission electron microscopy (TEM), and solid-state nuclear magnetic resonance (SS-NMR) techniques was carried out to gain a deep insight into the structural evolution of P2-Na0.66Ni0.33-xZnxMn0.67O2 (x = 0, 0.07) during cycling. In-situ HEXRD and ex-situ TEM measurements indicate that an irreversible phase transition occurs upon sodium insertion-extraction of Na0.66Ni0.33Mn0.67O2. Zinc doping of this system results in a high structural reversibility. XAFS measurements indicate that both materials are almost completely dependent on the Ni4+/Ni3+/Ni2+ redox couple to provide charge/discharge capacity. SS-NMR measurements indicate that both reversible and

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irreversible migration of transition metal ions into the sodium layer occurs in the material at the fully charged state. The irreversible migration of transition metal ions triggers a structural distortion, leading to the observed capacity and voltage fading. Our results allow a new understanding of the importance of improving the stability of transition metal layers. KEYWORDS: Sodium ion battery; Cathode material; Zn doping; Sodium nickel manganese oxide; Structural transition

1. Introduction Recently, a rapid increase in demand for renewable energy storage promotes further development of rechargeable batteries. Although lithium ion batteries (LIBs) have achieved great success thus far in many applications, the tension between their increasing market demand and the scarcity of lithium resources may hinder their use in large-scale energy storage applications.1,2 With sodium resources being substantially more abundant than lithium, sodium ion batteries (SIBs) have been investigated as a low-cost alternative to LIBs. However, designing a suitable cathode material with competitive electrochemical performance is the main challenge.3,4 In terms of the sequence of oxygen stacking and Na occupation environment, layered NaxMO2 materials are generally classified as either the P2-type or O3-type, using the notation proposed by Delmas et al.5 According to experimental results from the literature, P2-type NaxMO2 materials show higher sodium intercalation/deintercalation activity than that of corresponding O3-type compounds, which is considered to be attributed to the presence of larger trigonal prismatic sites in P2-type structures for sodium ions to occupy.6-8 Among these materials, P2-Na0.66Ni0.33Mn0.67O2 is one of the best candidates for use as a cathode in SIBs, due to its high Na+ ionic and electronic conductivity, high


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energy density, low cost, and low toxicity. P2-Na0.66Ni0.33Mn0.67O2 operates in the vicinity of 3.7 V vs. Na+/Na, which is at the same level as cathode materials for LIBs and is located in the stable voltage range of common Na-based electrolytes.9-12 Severe capacity and voltage fading during cycling has been the biggest barrier thus far to the development of P2-type Na0.66Ni0.33Mn0.67O2 cathode materials. Some researchers have already found that the electrochemical performance of P2-Na0.66Ni0.33Mn0.67O2 can be improved by doping with different metallic cations (Li+, Zn2+, Co3+, Fe3+, Ti4+, etc).13-19 Among them, zinc doping on the nickel site (P2-Na0.66Ni0.33–xZnxMn0.67O2) can make the charge-discharge curves smoother and greatly improve the capacity and voltage retention of this system. However, a lack of understanding still exists regarding how Zn2+ influences the structural reversibility of the system during the charge-discharge process. Considering the parallel case of LIBs, it is speculated that there are several possible causes responsible for the capacity and voltage fading, including (1) irreversible structural transition due to Jahn-Teller distortion, accompanied by the migration of transition metal ions,20-22 (2) irreversible loss of oxygen from the lattice, leading to the formation of oxygen vacancies,23-25 (3) electrolyte decomposition in the higher part of the working voltage range,26-28 and (4) the surface dissolution of manganese.29,30 To date, the primary degradation mechanism has not been fully understood, due to a lack of explicit experimental data on structural variations. In this paper, we studied the degradation mechanism of P2-Na0.66Ni0.33Mn0.67O2 and the positive effects of Zn2+ doping on the electrochemical performance of the P2-type Na-Ni-Mn-O system. In-situ high energy X-ray diffraction (HEXRD), ex-situ X-ray absorption fine structure spectroscopy (XAFS), transmission electron microscopy (TEM), and solid-state nuclear magnetic resonance (SS-NMR) were used

to

make

a

comparative

analysis

of

the

bulk

and local


structural

changes

of


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P2-Na0.66Ni0.33–xMn0.67O2 (x = 0, 0.07) during cycling. The cause of the accelerated degradation of capacity and voltage during high-voltage operation remains poorly understood. As this is the first direct structural observation of the effects of Zn2+ doping, the information gained here may lead to an in-depth knowledge of the system’s capacity and voltage fading mechanisms.

2. Experimental Section Synthesis. The powder samples of Na0.66Ni0.33–xMn0.67O2 with x = 0, 0.07 were synthesized by a solid-state reaction method as described in our previous study.14 Stoichiometric amounts of Na2CO3 (≥ 99.8%, Sinopharm Chemical Reagent Co. Ltd.), NiO (≥ 98.0%, Sinopharm Chemical Reagent Co. Ltd.), ZnO (≥ 99.0%, Sinopharm Chemical Reagent Co. Ltd.), and MnO2 (99.9%, Alfa Aeasar) were mixed via ball milling for 3 h with acetone, followed by drying at 80 oC for 2 h to remove the dispersant. The mixture was compressed into pellets under a pressure of 10 MPa (23 mm diameter, 5 mm thickness). The pellets were calcined at 900 oC for 15 h in air and then naturally cooled down to room temperature in the furnace. The resulting samples were stored in an argon-filled glove box (M. Braun, O2 and H2O < 0.5 ppm) to avoid contact with moisture. Electrochemical Tests. Cathodes (except in-situ HEXRD) were prepared by mixing 80 wt. % active material, 10 wt. % acetylene black, and 10 wt. % poly(-vinylidene fluoride) binder (PVDF) in N-methyl-2-pyrrolidene (NMP) solvent by ball milling. The slurry was spread on aluminum current collectors and dried at 80 oC for 12 h in a vacuum oven. Metallic sodium discs were used as the negative electrode. The electrolyte was 1 M NaClO4 dissolved in PC with 2 vol. % FEC as an electrolyte additive. A piece of glass fiber filter was employed as a separator. The CR2025 coin-type cells were assembled in an argon-filled glove box. The cells were charged and discharged


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galvanostatically on a LAND CT-2001A (Wuhan, China) battery test system. Galvanostatic intermittent titration technique (GITT) was employed to obtain the equilibrium potentials at a pulse of 10 mA g–1 for 30 min, followed by a 2 h relaxation process between each pulse. Material Characterization. XRD patterns of powder samples were collected with a Rigaku Ultima IV powder X-ray diffractometer with Cu Kα radiation (λ = 1.5406 Å). Rietveld refinement was performed using the General Structure Analysis System (GSAS) program. Chemical compositions of the samples were determined using inductively coupled plasma-atomic emission spectrometry (ICP-AES, IRIS Intrepid II XSP, Thermo Electron). In-situ HEXRD measurements were carried out at Sector 11-ID-C of the Advanced Photon Source (APS) of Argonne National Laboratory, with the wavelength of X-ray used pre-set to 0.11798 Å. For the electrode fabrication, the active material was mixed with Super-P and PVDF in a weight ratio of 8/1/1. The slurry was coated onto carbon paper (Toray carbon paper 120) for the in-situ cells to avoid Al peaks interference. The charge-discharge current density of the cell was 15 mA g–1 during the test process, based on the mass of active material. The data collection time for each XRD scan was 28-30 min. Ex-situ XAFS data at the Ni and Mn K-edges was collected at the BL14W1 beamline of Shanghai synchrotron radiation facility (SSRF) in transmission mode at room temperature. A Si (111) double-crystal monochromator, cooled by liquid nitrogen, was employed to select a single X-ray wavelength. The monochromator was calibrated to reject higher harmonics of the selected wavelength (harmonic content < 10–4). All the spectra were recorded in a range of energies from 200 eV below to 1000 eV above the Ni and Mn absorption edge energies (E0). Energy calibrations were carried out using the first derivative of the spectra of the Ni and Mn metal foils, which were simultaneously collected during each measurement. XANES and EXAFS data were analyzed using the ATHENA



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software package. Before the XAFS experiments, coin cells were charged/discharged to the selected voltage states, and then disassembled in an Ar-filled glove box. The cathodes were taken out and cleaned with DMC to remove any residual electrolyte. Ex-situ 23Na MAS NMR spectra were acquired on a Bruker AVANCE III 400 MHz spectrometer, at a 23

Na Larmor frequency of 105 MHz. All experiments were carried out using a double resonance 1.3

mm MAS probe, spinning at frequencies of up to 60 kHz with a Hahn-echo pulse sequence (90o pulse – τ – 180o pulse – τ). The 90o pulse length for 23Na was 1.2 µs and the recycle delay was 20 ms. All the 23

Na shifts were referenced to 1 M NaCl (aq) (0 ppm). Measurements were conducted on cycled

positive electrodes after disassembly of the cell and removal of the electrolyte by rinsing with dimethyl carbonate (DMC) in an Ar-filled glovebox. The electrode materials were scraped carefully from the Al current collector and put into the sealed rotor before being characterized by SS-NMR. Transmission electron microscopy (TEM) characterization was performed by a JEM-2100 microscope (JEOL Ltd. Japan) at 200 kV. Measurements were conducted on cycled positive electrodes after disassembly of the cell and removal of the electrolyte by rinsing with dimethyl carbonate (DMC) in an Ar-filled glovebox. The electrode materials were scraped carefully from the Al current collector. The samples were prepared by dispersing the sample powders in ethanol under ultrasonication. The suspension was deposited on a copper grid with carbon film for TEM observation.

3. Results and Discussion XRD refinements (Figure 1, Table S1 and Table S2, Supporting Information) show that P2-Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 can be well indexed in the hexagonal system with the P63/mmc space group. The element compositions of Na0.66Ni0.33-xZnxMn0.67O2 (x = 0, 0.07)


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were also confirmed by means of ICP-AES analysis. The actual molar ratios of Na, Ni, Zn, and Mn were determined to be 0.663:0.328:0:0.670 for Na0.66Ni0.33Mn0.67O2 and 0.641:0.254:0.070:0.670 for Na0.66Ni0.26Zn0.07Mn0.67O2, which are very close to the designed values. Figure 2 shows the charge-discharge curves of P2-Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 in the first 10 cycles at 12 mA g−1. It can be observed that Na0.66Ni0.33Mn0.67O2 delivers a discharge capacity of 139 mAh g−1 in the first cycle (Figure 2a). There are three distinct voltage plateaus in the charge-discharge curves. However, the discharge voltage and capacity decay rapidly as cycling proceeds. The voltage and capacity losses of Na0.66Ni0.33Mn0.67O2 after 10 cycles are 8.8% and 30.9%, respectively. In contrast, it can be seen that the initial discharge capacity of Na0.66Ni0.26Zn0.07Mn0.67O2 is 127 mAh g−1 (Figure 2b). The voltage and capacity losses of Na0.66Ni0.26Zn0.07Mn0.67O2 after 10 cycles are 1.5% and 6.9%, respectively, which have a significant reduction compared to the undoped material. Furthermore, the charge-discharge curves become smoother after Zn doping, also suggesting that Zn doping has a significant impact on electrochemical cycling. The GITT curves (Figure S1a, b, Supporting Information) show kinetic processes of sodium ion extraction/insertion in the P2-Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 electrode materials. It can be seen that the

voltage polarization is larger in the voltage range above 3.8 V and below 3 V for Na0.66Ni0.33Mn0.67O2. After zinc doping, the voltage polarization becomes smaller in the voltage range of 3.8 − 4.1 V and the voltage plateaus becomes steeper. Obviously, it is not just due to the decrease of nickel content. These results reflect that zinc doping has a positive effect on the material structure. The diffusion coefficients (DNa ) can be determined from the galvanostatic +

intermittent titration pulses (Figure S1c, Supporting Information).31,32 In

order

to

reveal

the

structural

phase

transitions

of


Na0.66Ni0.33Mn0.67O2

and


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Na0.66Ni0.26Zn0.07Mn0.67O2 during the charge-discharge process, a synchrotron in-situ HEXRD technique was employed. Figure 3a,b show the in-situ HEXRD patterns of Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 along with the charge-discharge curve during the 1st cycle, respectively. Before cycling, the major diffraction peaks of these two materials can be attributed to the P2-type structure. In addition, there are some peaks which do not belong to the P2-type phase (hexagonal lattice, P63/mmc) that appear in the patterns. The peaks at 1.37o, 2.00o, and 2.87o can be attributed to signals of carbon paper used for current collector (Figure S2, Supporting Information for the HEXRD pattern of carbon paper). The (002) and (004) peaks of P2-type Na0.66Ni0.33Mn0.67O2 gradually shift to lower angles during the charge process, suggesting that the unit cell parameter c gradually increases with the extraction of sodium ions. This change is due to an increase in the repulsive force between adjacent transition metal layers after a decrease in sodium content, which leads to an expansion of the interlayer spacing (equal to c/2). When the voltage is higher than 4.05 V, the positions of the diffraction peaks such as (004), (102) and (104) show no obvious change, while the intensities of these peaks gradually weaken and the peak widths increase. This phenomenon shows that the long-range ordering of the P2 structure is clearly interrupted. At charge completion, the peak intensities of the P2-type phase significantly decrease compared to their original values, while some new peaks appear which can be assigned to a high voltage “Z” phase with a short-range disordered structure.33,34 Although the intensity of the (002) peak of the P2 phase is significantly lower in the 4.2 V plateau region and gradually shifts to that of the “Z” phase, its position does not show any obvious trend toward an increasing angle, indicating that the change in layer spacing is not large during the “Z” phase formation process. This means that the glide degree of the transition metal layer is small, and the repulsive interaction between vertically aligned


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oxygen atoms is weakened due to a certain local structural change. During the discharge process, the diffraction peaks of the P2-type phase do not recover to their original states, demonstrating that an irreversible phase transition has occurred. It can be inferred from the high voltage data that the distortion effect begins to have a significant impact on the structural reversibility of the Na0.66Ni0.33Mn0.67O2 material. When discharged to 3.3 V, the peak intensity of the P′2-type phase (orthorhombic, space group Cmcm) significantly increases.33 This result shows that the structure of P2-type Na0.66Ni0.33Mn0.67O2 undergoes a greater degree of distortion during cycling than that of P2-NaxMn1/2Fe1/2O2, which also exhibits capacity and voltage fading but to a lesser extent. During the 2nd cycle, the peaks of the P2-type phase do not reappear, demonstrating that the P2-P′2 phase transition is highly irreversible for the Na0.66Ni0.33Mn0.67O2 system (Figure S3, Supporting Information). Therefore, the capacity and voltage of Na0.66Ni0.33Mn0.67O2 are degraded until the complete P2-P′2 phase transition has been carried out. After doping with Zn2+, the structural variation of Na0.66Ni0.26Zn0.07Mn0.67O2 during the cycling process exhibits significant differences. Similar to the case of Na0.66Ni0.33Mn0.67O2, when charged to 4.05 V, the interlayer spacing increases with sodium extraction, indicated by changes in the (002) and (004) peaks. At the same time, the (100) peak shifts to a higher angle, suggesting that the transition metal layers shrink along the a direction with an increase in the valence state of the transition metal. However, when the voltage is higher than 4.05 V, the (002) peak gradually shifts to a higher angle, and the material structure is transformed from the P2-type (P63/mmc) to the O2-type (P63mc).35 During the discharge process, the diffraction peaks of Na0.66Ni0.26Zn0.07Mn0.67O2 can reappear, a phenomenon that does not exist with Na0.66Ni0.33Mn0.67O2. This result confirms that zinc doping can significantly help improve the reversibility of the P2-O2 phase transition, which can be realized via a glide of the



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transition metal layer. Figure 4a shows the variation in unit cell parameters of Na0.66Ni0.33Mn0.67O2 during the 1st cycle. In the charge process, c increases by 2.02% in the P2 region, while a decreases by 1.27%. In the P2 + Z and Z regions, c remains unchanged, and a decreases by 14% with the generation of the Z phase. In the discharge process, b and c decrease by 0.2% and 13% respectively, while a increases by 0.59%. Figure 4b shows the variation in unit cell parameters of Na0.66Ni0.26Zn0.07Mn0.67O2 during the 1st cycle. In the P2 region, c increases by 1.68%. After the P2-O2 phase transition, c rapidly decreases by 5.48%, and the total variation in c from the initial state to the fully charged state is about 3.80%. During the whole charge process, a shows a decreasing trend with a variation value of 1.76%. Because of the increase in c and the decrease in a in the P2 region, the variation of the unit cell volume V is not large. However, upon the occurrence of the P2-O2 phase transition, the rapid decrease in c leads to a rapid decrease in V. The total variation in V from the initial state to the fully charged state is 7.25%. The change trend of cell parameters V, c, and a in the discharge process is opposite to that of the charge process, and the two have high symmetry, which indicates that the Na0.66Ni0.26Zn0.07Mn0.67O2 structure has better reversibility in the charge-discharge process. In order to analyze the active redox couple and the changes in bond length in the local structure, X-ray absorption fine structure (XAFS) spectra were collected under ex-situ conditions. Figure 5a-d show the normalized XANES spectra of Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 electrodes at the Ni and Mn K-edges, collected at different charged/discharged states. The presence of pre-edge peaks in the XANES spectra is indicative of structural distortion in the NiO6 and MnO6 octahedra. The spectra of reference materials, such as NiO, Mn2O3, and MnO2, are also provided for the valence state analysis.


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The Ni K-edges of the pristine electrodes for the two materials correspond well with the NiO reference sample, suggesting that the valence state of the nickel ions is +2. As charging proceeds, the Ni K-edge absorption shifts toward the higher energy side, suggesting that nickel ions are involved in the electrochemical reaction in the two materials. When the voltage is higher than 3.9 V, there is only a slight shift in the Ni K-edge observed for the two materials. It is believed that this phenomenon is mainly due to the overlap between the eg energy band of Ni3+/4+ and the 2p band of O2–, leading to electron

delocalization

between

oxygen

and

nickel.

When

the

Na0.66Ni0.33Mn0.67O2

and

Na0.66Ni0.26Zn0.07Mn0.67O2 electrodes are discharged to 2.2 V, the Ni K-edges shift toward the lower energy side and return almost to their original position, indicating a reversible reduction of Ni4+/3+ to Ni2+. The Mn K-edges of the pristine electrodes for the two materials match well with the MnO2 reference sample, suggesting that the valence of manganese is +4. The position of the Mn K-edge remains unchanged during the entire cycling process, suggesting that the manganese ions in the two materials are not involved in any electrochemical reaction and instead play a major role in stabilizing the structure. The k3-weighted Fourier-transform (FT) magnitudes of the EXAFS spectra at the Ni and Mn K-edges for Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 at different charged/discharged states are shown in Figure 6a-d. In general, the first peak in the range of 1 – 2 Å can be attributed to the single scattering paths M→O by the nearest oxygen surrounding the absorbing Ni or Mn atoms, while the second peak in the range of 2 – 3 Å can be attributed to the single scattering paths M→M by the transition metals within the same a-b plane neighboring the absorbing Ni or Mn atoms.36 Significant change occurs in the first coordination layer of nickel during the charge-discharge process. It can be observed that the lengths of Ni-O bonds first decrease from around 1.59 Å to 1.46 Å for



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Na0.66Ni0.33Mn0.67O2 and from around 1.59 Å to 1.46 Å for Na0.66Ni0.26Zn0.07Mn0.67O2, due to the increase in oxidation states. The lengths of Ni-O bonds increase gradually as the discharge begins. When the material is discharged to 3.1 V, the lengths of Ni-O bonds in Na0.66Ni0.33Mn0.67O2 reach around 1.71 Å, which is significantly larger than that of the pristine state. By contrast, the lengths of Ni-O bonds in Na0.66Ni0.26Zn0.07Mn0.67O2 reach around 1.65 Å, suggesting that zinc doping helps relieve the distortion of Ni-O octahedra during cycling. This result also confirms that the irreversible structural transition mainly occurs during the discharge process. Although the phase transition occurs after a large amount of sodium extraction, the lengths of the Mn-O bonds exhibit no obvious change and remain at around 1.48 Å for both materials. This phenomenon is consistent with the results of the XANES spectra. In addition, the change in the intensity and position of the Ni-M and Mn-M peaks in the EXAFS spectra, especially those measured at 4.25 V, is caused by the distortion of adjacent Ni-O octahedra. Figure 7 shows TEM images and corresponding selected area diffraction (SAED) patterns of pristine Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2. Two particles with similar diameter (ca. 700 nm) were selected for high-resolution TEM (HRTEM) observation. The SAED patterns of the two samples show clear diffraction spots without any obvious impurity phase, and can be well indexed to the hexagonal crystal structure along the [0001] zone axis. The HRTEM images show well-defined lattice fringes with spacings of 2.497 Å and 2.504 Å for Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 respectively, both corresponding to (1010) crystal planes. Figure 8 shows HRTEM, fast Fourier transform (FFT) and inverse fast Fourier transform (IFFT) images of Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 after 15 cycles at 15 mA g–1. It can be seen that the lattice fringes of the Na0.66Ni0.33Mn0.67O2 electrode exhibit a locally disordered structure with a lot of stacking faults. There


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are extra spots appearing in the FFT image that do not belong to the hexagonal lattice, indicating the occurrence of a structural transition. According to the in-situ HEXRD results, this structural disordering can be attributed to continuous hexagonal-orthorhombic phase transition during cycling. A clearer view of the lattice can be gained via the IFFT images, where three different areas are marked: (area 1) a relatively ordered lattice area with atomic arrangements close to that of the hexagonal system and (areas 2, 3) a significantly distorted lattice area characteristic of the orthorhombic system. However, the Na0.66Ni0.26Zn0.07Mn0.67O2 electrode shows a much different result. It can be seen from the relevant IFFT images that no great change has taken place in the appearance of the lattice fringes. It can thus be concluded that zinc doping plays a role in inhibiting irreversible phase transition during cycling. Figure 9 shows

23

Na MAS NMR spectra from Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2

electrodes at different charged/discharged states. Two 23Na resonances between 1300-1600 ppm can be observed for both materials. The stronger signal can be ascribed to sodium (including Nae and Naf) in the pristine P2-type structure. The weaker signal at the shoulder of the main signal most probably likely corresponds to the sodium-poor phase with x < 0.66 of NaxNi0.33-xZnxMn0.67O2, caused by moisture exposure during electrode preparation. The signals at 0 ppm correspond to a small amount of sodium carbonate impurity. During the charge process, the two sodium peaks coalesce into one symmetric peak and move to lower chemical shifts. Such a peak shift implies that the interactions between sodium and transition metal ions become weaker, indicative of the expansion of the layer spacing resulting in the increase in distance between sodium and transition metal ions. However, when the voltage is higher than 3.75 V, a new peak starts to appear at around 230-240 ppm. It is speculated that this new peak is associated with the “Z” phase. With further increasing the voltage to 4.25 V, the coalesced peak at ~1000 ppm in the



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Na0.66Ni0.33Mn0.67O2 electrode suddenly jumps to a higher chemical shift and is broadened significantly. By contrast, the sodium signal in the Na0.66Ni0.26Zn0.07Mn0.67O2 electrode always moves to a lower chemical shift throughout the charge process. A new peak (267 ppm) also appears at 4.05 V but disappears in subsequent cycling. In addition, a broad peak can be observed in the range of 540 − 920 ppm at 4.25 V, indicating the formation of the O2-type phase. A similar wide spectral peak can be also observed in the

23

Na MAS NMR spectra of the O3-type phase.35 These results suggest that the two

materials experience different structural changes at the end of charging. More transition metal ions (mainly nickel) may migrate from the octahedral sites into tetrahedral sites in the sodium layers of Na0.66Ni0.33Mn0.67O2 compared to Na0.66Ni0.26Zn0.07Mn0.67O2, blocking the glide of transition metal layers. During the discharge process, the peak from the Na0.66Ni0.33Mn0.67O2 electrode at 234 ppm disappears when the material is discharged to below 3.9 V, indicating that the migration of transition metal ions has a certain degree of reversibility. It can be observed that while the sodium peaks of Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 both move to lower chemical shifts, obvious changes take place in the relative intensity of the peaks when the materials are discharged to 2.2 V. The relative peak intensity of the sodium-poor phase of Na0.66Ni0.33Mn0.67O2 increases more obviously than that of Na0.66Ni0.26Zn0.07Mn0.67O2. This phenomenon shows that it is more difficult for the structure of Na0.66Ni0.33Mn0.67O2 to return to its original state, probably due to the presence of transition metal ions in the sodium layer as well as structural distortion. Zinc doping can improve the stability of the two-dimensional structure of the P2-type material and suppress the formation of a sodium-poor phase. The

23

Na MAS NMR spectra of Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 electrodes during

different cycles are also shown in Figure S4, Supporting Information. It can be observed for both


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materials that the intensity of sodium peaks reduces to different extents after cycling, demonstrating the occurrence of structural distortion and the formation of P′2-type phase.37 Compared to Na0.66Ni0.33Mn0.67O2, the local environment of sodium in Na0.66Ni0.26Zn0.07Mn0.67O2 appears to be relatively more stable. The above data and results lead to some insights about the observed severe capacity and voltage degradation with cycling of Na0.66Ni0.33Mn0.67O2. The overall phase transformation characteristics are shown in Figure 10. It is believed that there is an absence of Ni2+/Na+ interlayer exchange, due to the radius mismatch between the two ions in the pristine material.38 However, the migration of transition metal cations into the tetrahedral sites in the sodium layer occurs in the highly desodiated state, accompanied by a distortion of oxygen lattices. This can be analyzed in terms of ligand-to-metal charge-transfer (LMCT).39 Electron delocalization can be promoted by the larger overlapping Ni 3d-O 2p orbitals, especially in the case of high valence Ni4+. The electron localization of Ni4+ weakens the bonding between nickel and oxygen, allowing the migration of nickel ions to the nearest tetrahedral site as well as the release of oxygen. The presence of the transition metal ions in the tetrahedral site can increase the occupancy energy and hamper sodium ions from entering neighbouring octahedral sites, resulting in a decrease in the accessibility of sodium ions in the layer spacing. It is evident from the SS-NMR data that the migration of transition metal ions is partially reversible. On the other hand, they can also migrate elsewhere, residing in other available vacancies. As a result, the layered structure is progressively destroyed with cycling. This structural transition is the primary factor contributing to the capacity and voltage degradation of this system. By contrast, the presence of inactive zinc ions with octet stability forming Zn-O bonds in the material can effectively relieve the effect of LMCT on the structural stability and significantly improve the stability of the layered structure during cycling.



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4. Conclusions In this work, we have comprehensively investigated and compared the structural transformations of P2-type Na0.66Ni0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 cathode materials by using the HEXRD, XAFS, TEM, and SS-NMR techniques. In-situ HEXRD results show that the structural distortion of Na0.66Ni0.33Mn0.67O2 leads to a phase transition from the hexagonal P2-type structure to the orthorhombic P′2-type structure during cycling, resulting in a rapid fade of capacity and voltage. However, after Zn2+ doping, the reversibility of the P2-O2 phase transition during cycling significantly improves, causing Na0.66Ni0.26Zn0.07Mn0.67O2 to show better capacity and voltage retention. XAFS results show that the vast majority of the capacity is related to the Ni4+/Ni3+/Ni2+ redox couple for both materials. The valence state of the manganese ions remains unchanged, with the ions playing a major role in stabilizing the structure. During the charge-discharge process, Zn2+ doping reduces the distortion degree of the Ni-O octahedrons in the Na-Ni-Mn-O structure and improves the reversibility of the distortion. HRTEM results show that the lattice fringes of Na0.66Ni0.33Mn0.67O2 significantly change after 15 cycles, indicating that a phase transition occurs during cycling. By contrast, the symmetry and lattice fringes of Na0.66Ni0.26Zn0.07Mn0.67O2 do not change significantly, suggesting that its structure exhibits a good stability during the charge-discharge process. SS-NMR results show that compared with Na0.66Ni0.26Zn0.07Mn0.67O2, more sodium ions in Na0.66Ni0.33Mn0.67O2 can migrate from the octahedral sites in the transition metal layer to the tetrahedral sites in the sodium layer in the fully charged state. These results will assist in the further development of P2-type Ni-Mn cathode systems with higher energy densities.


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ASSOCIATED CONTENT Supporting Information Detailed information involving parameters of XRD refinement, GITT curves, In-situ HEXRD, and NMR spectra. AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]. Author Contributions X.W. carried out the materials preparation and electrochemical test. G.X., Z.C. and K.A. acquired the in-situ XRD data, X.W., S.Z. and Y.Y. analyzed the in-situ data. X.W. and Z.G. carried out the XAFS experiments. G.Z. acquired the NMR data, G.Z. and R.F. analyzed the NMR data. X.W. and G.Z. processed and analyzed the XAFS data. X.W. recorded and analyzed the TEM results. M.M. modified the paper. Y.Y. and X.W. proposed the research, and Y.Y. attained the main financial support for the research. Notes The authors declare no competing financial interest. ACKNOWLEDGMENTS The authors acknowledge financial support of their research from the National Natural Science Foundation of China (Grant No. 21233004, 21473148, and 21428303) and the National Basic Research Program of China (973 program, Grant No. 2011CB935903). We sincerely acknowledge Dr. W. Wen and other staff of the XAFS beamline of Shanghai Synchrotron Radiation Facility for their support. R. Fu is also indebted to the support for being a PCOSS fellow by the State Key Lab of Physical Chemistry of



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Solid Surfaces, Xiamen University, China.

REFERENCES (1) Goodenough, J. B. Electrochemical Energy Storage in a Sustainable Modern Society. Energy Environ. Sci. 2014, 7, 14-18. (2) Wang, Y. X.; Liu, B.; Li, Q. Y.; Cartmell, S.; Ferrara, S.; Deng, Z. Q.; Xiao, J. Lithium and Lithium Ion Batteries for Applications in Microelectronic Devices: A review. J. Power Sources 2015, 286, 330-345. (3) Pan, H. L.; Hu, Y. S.; Chen, L. Q. Room-Temperature Stationary Sodium-Ion Batteries for Large-Scale Electric Energy Storage. Energy Environ. Sci. 2013, 6, 2338-2360. (4) Fang, C.; Huang, Y. H.; Zhang, W. X.; Han, J. T.; Deng, Z.; Cao, Y. L.; Yang, H. X. Routes to High Energy Cathodes of Sodium-Ion Batteries. Adv. Energy Mater. 2015, 6, 1501727. (5) Delmas, C.; Fouassier, C.; Hagenmuller, P. Structural Classification and Properties of the Layered Oxides. Physica B & C 1980, 99, 81-85. (6) Yabuuchi, N.; Kajiyama, M.; Iwatate, J.; Nishikawa, H.; Hitomi, S.; Okuyama, R.; Usui, R.; Yamada, Y.; Komaba, S. P2-Type Nax[Fe1/2Mn1/2]O2 Made From Earth-Abundant Elements for Rechargeable Na Batteries. Nat. Mater., 2012, 11, 512-517. (7) Palomares, V.; Casas-Cabanas, M.; Castillo-Martínez, E.; Han, M. H.; Rojo, T. Update on Na-Based Battery Materials. A Growing Research Path. Energy Environ. Sci. 2013, 6, 2312-2337. (8) Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Research Development on Sodium-Ion Batteries. Chem. Rev. 2014, 114, 11636-11682. (9) Lee, D. H.; Xu, J.; Meng, Y. S. An Advanced Cathode for Na-Ion Batteries with High Rate and


Page 18 of 34

Page 19 of 34

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Excellent Structural Stability. Phys. Chem. Chem. Phys. 2013, 15, 3304-3312. (10) Manikandan, P.; Ramasubramonian, D.; Shaijumon, M. M. Layered P2-type Na0.5Ni0.25Mn0.75O2 as A High Performance Cathode Material for Sodium-Ion Batteries. Electrochim. Acta 2016, 206, 199-206. (11) Sun, X.; Ji, X. Y.; Xu, H. Y.; Zhang, C. Y.; Shao, Y.; Zang, Y.; Chen, C. H. Sodium Insertion Cathode Material Na0.67[Ni0.4Co0.2Mn0.4]O2 with Excellent Electrochemical Properties. Electrochim. Acta 2016, 208, 142-147. (12) Li, Z. Y.; Gao, R.; Zhang, J. C.; Zhang, X. L.; Hua, Z. B.; Liu, X. F. New Insights into Designing High-Rate Performance Cathode Materials for Sodium Ion Batteries by Enlarging the Slab-spacing of the Na-Ion Diffusion Layer. J. Mater. Chem. A 2016, 4, 3453-3461. (13) Xu, J.; Lee, D. H.; Clément, R. J.; Yu, X. Q.; Leskes, M.; Pell, A. J.; Pintacuda, G.; Yang, X. Q.; Grey, C. P.; Meng, Y. S. Identifying the Critical Role of Li Substitution in P2-Nax[LiyNizMn1-y-z]O2 (0 < x, y, z < 1) Intercalation Cathode Materials for High-Energy Na-Ion Batteries. Chem. Mater. 2014, 26, 1260-1269. (14) Yoshida, H.; Yabuuchi, N.; Kubota, K.; Ikeuchi, I.; Garsuch, A.; Schulz-Dobrick, M.; Komaba, S. P2-type Na2/3Ni1/3Mn2/3-xTixO2 as a New Positive Electrode for Higher Enegy Na-Ion Batteries. Chem. Commun. 2014, 50, 3677-3680. (15) Yuan, D. D.; He, W.; Pei, F.; Wu, F. Y.; Wu, Y.; Qian, J. F.; Cao, Y. L.; Ai, X. P.; Yang, H. X. Synthesis and Electrochemical Behaviors of Layered Na0.67[Mn0.65Co0.2Ni0.15]O2 Microflakes as a Stable Cathode Material for Sodium-Ion Batteries. J. Mater. Chem. A 2013, 1, 3895-3899. (16) Yuan, D. D.; Hu, X. H.; Qian, J. F.; Pei, F.; Wu, F. Y.; Mao, R. J.; Ai, X. P.; Yang, H. X.; Cao, Y. L. P2-type Na0.67Mn0.65Fe0.2Ni0.15O2 Cathode Material with High-Capacity for Sodium-ion Battery.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Electrochim. Acta 2014, 116, 300-305. (17) Hasa, I; Passerini, S.; Hassoun, J. A Rechargeable Sodium-Ion Battery Using A Nanostructured Sb–C Anode and P2-type Layered Na0.6Ni0.22Fe0.11Mn0.66O2 cathode. RSC Adv. 2015, 5, 48928-48934. (18) Li, Z. Y.; Gao, R.; Sun, L. M.; Hua, Z. B.; Liu, X. F. Designing An Advanced P2-Na0.67Mn0.65Ni0.2Co0.15O2 Layered Cathode Material for Na-ion Batteries. J. Mater. Chem. A, 2015, 3, 16272-16278. (19) Wu, X. H.; Guo, J. H.; Wang, D. W.; Zhong, G. M.; McDonald, M. J.; Yang, Y. P2-type Na0.66Ni0.33-xZnxMn0.67O2 as New High-Voltage Cathode Materials for Sodium-Ion Batteries. J. Power Sources 2015, 281, 18-26. (20) Mohanty, D.; Li, J. L.; Abraham, D. P.; Huq, A.; Payzant, E. A.; WoodIII, D. L.; Daniel, C. Unraveling the Voltage-Fade Mechanism in High-Energy-Density Lithium-Ion Batteries: Origin of the Tetrahedral Cations for Spinel Conversion. Chem. Mater. 2014, 26, 6272-6280. (21) Pan, J.; Deng, J. Q.; Yao, Q. R.; Zou, Y. J.; Wang, Z. M.; Zhou, H. Y.; Sun, L. X.; Rao, G. H. Novel LiNi0.5Mn1.5O4 Porous Microellipsoids as High-Performance Cathode Materials for Lithium Ion Batteries. J. Power Sources 2015, 288, 353-358. (22) Zhao, H. Y.; Liu, S. S.; Wang, Z. W.; Cai, Y.; Tan, M.; Liu, X. Q. Enhanced Elevated-Temperature Performance of LiAlxSi0.05Mg0.05Mn1.90-xO4 (0≤x≤0.08) Cathode Materials for High-Performance Lithium-Ion Batteries. Electrochim. Acta 2016, 199, 18-26. (23) Gu, M.; Genc, A.; Belharouak, I.; Wang, D. P.; Amine, K.; Thevuthasan, S.; Baer, D. R.; Zhang, J. G.; Browning, N. D.; Liu, J.; Wang, C. M. Nanoscale Phase Separation, Cation Ordering, and Surface Chemistry in Pristine Li1.2Ni0.2Mn0.6O2 for Li-Ion Batteries. Chem. Mater. 2013, 25, 2319-2326. (24) Mohanty, D.; Sefat, A. S.; Payzant, E. A., Li, J. L.; Wood III, D. L.; Daniel, C. Unconventional


Page 20 of 34

Page 21 of 34

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Irreversible Structural Changes in A High-Voltage Li–Mn-rich Oxide for Lithium-Ion Battery Cathodes. J. Power Sources 2015, 283, 413-428. (25) Qian, D.; Xu, B.; Chi, M. F.; Meng, Y. S. Uncovering the Roles of Oxygen Vacancies in Cation Migration in Lithium Excess Layered Oxides. Phys. Chem. Chem. Phys. 2014, 16, 14665-14668. (26) Huang, W; Xing, L; Zhang, R; Wang, X.; Li, W. A Novel Electrolyte Additive for Improving the Interfacial Stability of High Voltage Lithium Nickel Manganese Oxide Cathode. J. Power Sources 2015, 293, 71-77. (27) Choi, N. S.; Han, J. G.; Ha, S. Y.; Park, I.; Back, C. K. Recent Advances in the Electrolytes for Interfacial Stability of High-Voltage Cathodes in Lithium-Ion Batteries. RSC Adv. 2015, 5, 2732-2748. (28) Zhu, Y. M.; Rong, H. B.; Mai, S. W.; Luo, X. Y.; Li, X. P.; Li, W. S. Significantly Improved Cyclability of Lithium Manganese Oxide under Elevated Temperature by An Easily Oxidized Electrolyte Additive. J. Power Sources 2015, 299, 485-491. (29) Wang, L. F.; Fang, B. J.; Chen, J. S. Rotating Ring-Disc Electrode Measurements of Manganese Dissolution and Capacity Loss of Li1+xMn2-xO4 and Li1+xAlyMn2-x-yO4 Spinel Electrodes for Lithium-Ion Batteries. J. Power Sources 2005, 150, 1-10. (30) Put, B.; Vereecken, P. M.; Labyedh, N.; Sepulveda, A.; Huyghebaert, C.; Radu, I. P.; Stesmans, A. High Cycling Stability and Extreme Rate Performance in Nanoscaled LiMn2O4 Thin Films. ACS Appl. Mater. Interfaces 2015, 7, 22413-22420. (31) Kleiner, K.; Melke, J.; Merz, M.; Jakes, P.; Nagel, P.; Schuppler, S.; Liebau, V.; Ehrenberg, H. Unraveling the Degradation Process of LiNi0.8Co0.15Al0.05O2 Electrodes in Commercial Lithium Ion Batteries by Electronic Structure Investigations. ACS Appl. Mater. Interfaces 2015, 7, 19589-19600. (32) Li, Y.; Bai, W. Q.; Zhang, Y. D.; Niu, X. Q.; Wang, D. H.; Wang, X. L.; Gu, C. D.; Tu, J. P.



1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Synthesis and Electrochemical Performance of Lithium Vanadium Phosphate and Lithium Vanadium Oxide Composite Cathode Material for Lithium Ion Batteries. J. Power Sources 2015, 282, 100-108.

(33) Boisse, B. M.; Carlier, D.; Guignard, M.; Bourgeois, L.; Delmas, C. P2-NaxMn1/2Fe1/2O2 Phase Used as Positive Electrode in Na Batteries: Structural Changes Induced by the Electrochemical (De)intercalation Process. Inorg. Chem. 2014, 53, 11197-11205. (34) Talaie, E.; Duffort, V.; Smith, H. L.; Fultz, B.; Nazar, L. F. Structure of the High Voltage Phase of Layered P2-Na2/3-z[Mn1/2Fe1/2]O2 and the Positive Effect of Ni Substitution on Its Stability. Energy Environ. Sci. 2015, 8, 2512-2523. (35) Paulsen, J. M.; Thomas, C. L.; Dahn, J. R. O2 Structure Li2/3[Ni1/3Mn2/3]O2: A New Layered Cathode Material for Rechargeable Lithium Batteries. J. Electrochem. Soc. 2000, 147, 861-868. (36) Lyu, Y. C.; Zhao, N. J.; Hu, E. Y.; Xiao, R. J.; Yu, X. Q.; Gu, L.; Yang, X. Q.; Li, H. Probing Reversible Multielectron Transfer and Structure Evolution of Li1.2Cr0.4Mn0.4O2 Cathode Material for Li-Ion Batteries in a Voltage Range of 1.0–4.8 V. Chem. Mater. 2015, 27, 5238-5252. (37) Singh, G.; López del Amo, J. M.; Galceran, M.; Pérez-Villar, S.; Rojo, T. Structural Evolution during Sodium Deintercalation/Intercalation in Na2/3[Fe1/2Mn1/2]O2. J. Mater. Chem. A 2015, 3, 6954-6961. (38) Cabana, J.; Chernova, N. A.; Xiao, J.; Roppolo, M.; Aldi, K. A.; Whittingham, M. S.; Grey, C.

P. Study of the Transition Metal Ordering in Layered NaxNix/2Mn1-x/2O2 (2/3≤x≤1) and Consequences of Na/Li Exchange. Inorg. Chem. 2013, 52, 8540-8550. (39) Wang, X. F.; Liu, G. D.; Iwao, T.; Okubo, M.; Yamada, A. Role of Ligand-to-Metal Charge Transfer in O3-Type NaFeO2-NaNiO2 Solid Solution for Enhanced Electrochemical Properties. J. Phys. Chem. C 2014, 118, 2970-2976.


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Figure 1. Rietveld refinements of the XRD data for (a) Na0.66Ni 0.33Mn0.67O2 and (b) Na0.66Ni0.26Zn0.07Mn0.67O2 materials. Observed (black circles) and calculated (red solid line), Bragg reflection peaks (green solid ticks) and the difference curve (blue) are shown, respectively.


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Figure 2. Charge-discharge curves of (a) Na0.66Ni0.33Mn0.67O2 and (b) Na0.66Ni0.26Zn0.07Mn0.67O2 between 2.2 and 4.3 V at 12 mA g−1 in the first 10 cycles.



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Figure 3. In-situ HEXRD patterns of (a) Na0.66Ni0.33Mn0.67O2 and (b) Na0.66Ni0.26Zn0.07Mn0.67O2 during the 1st cycle. The XRD peaks marked with symbol (◆) are assigned to the carbon paper current collector.


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Figure 4. Lattice parameter evolutions of (a) Na0.66Ni 0.33Mn0.67O2 and (b) Na0.66Ni0.26Zn0.07Mn0.67O2 during the 1st cycle.



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Figure 5. Normalized XANES spectra of (a, c) Na0.66Ni0.33Mn0.67O2 and (b, d) Na0.66Ni0.26Zn0.07Mn0.67O2 electrodes at the Ni and Mn K-edges collected at different charged/discharged states.


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Figure 6. The k3-weighted Fourier-transform (FT) magnitudes of the EXAFS spectra at the Ni and Mn K-edges for (a,c) Na0.66Ni0.33Mn0.67O2 and (b,d) Na0.66Ni0.26Zn0.07Mn0.67O2 at different charged/discharged states.



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Figure 7. TEM images and corresponding selected area diffraction (SAED) patterns of (a,c,e) Na0.66Ni 0.33Mn0.67O2 and (b,d,f) Na0.66Ni0.26Zn0.07Mn0.67O2.


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Figure 8. HRTEM, fast Fourier transform (FFT) and inverse fast Fourier transform (IFFT) images of (a) Na0.66Ni0.33Mn0.67O2 and (b) Na 0.66Ni0.26Zn0.07Mn0.67O2 after 15 cycles.



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Figure 9. 23Na MAS NMR spectra of (a) Na0.66Ni0.33Mn0.67O 2 and (b) Na0.66Ni0.26Zn0.07Mn0.67O2 electrodes at different charged/discharged states. Asterisk (＊) indicates the spinning side bands. (c) Comparison of 23Na MAS NMR spectra of Na 0.66Ni 0.33Mn0.67O2 and Na0.66Ni0.26Zn0.07Mn0.67O2 at 4.25 V charged state.


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Figure 10. Schematic diagram showing the structural transformation of Na0.66Ni0.33Mn0.67O2 observed experimentally.



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