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P2 orthorhombic Na0.7[Mn1-xLix]O2+y as cathode materials for Na-ion batteries Mi-Sook Kwon, Shin Gwon Lim, Yuwon Park, Sang-Min Lee, Kyung Yoon Chung, Tae Joo Shin, and Kyu Tae Lee ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b00058 • Publication Date (Web): 10 Apr 2017 Downloaded from http://pubs.acs.org on April 15, 2017

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ACS Applied Materials & Interfaces

P2 Orthorhombic Na0.7[Mn1-xLix]O2+y as Cathode Materials for Na-Ion Batteries

Mi-Sook Kwon,† Shin Gwon Lim,† Yuwon Park,† Sang-Min Lee,‡ Kyung Yoon Chung,§ Tae Joo Shin,⊥ Kyu Tae Lee*,†

†

School of Chemical and Biological Engineering, Institute of Chemical Processes, Seoul National

University, 1, Gwanak-ro, Gwanak-gu, Seoul 08826, South Korea ‡

Battery Research Center, Korea Electrotechnology Research Institute, 12 Bulmosan-ro 10 bean-gil,

Changwon 51543, South Korea §

Center for Energy Convergence Research, Korea Institute of Science and Technology, Hwarang-ro 14-

gil 5, Seongbuk-gu, Seoul 02792, South Korea ⊥

UNIST Central Research Facilities & School of Natural Science, Ulsan National Institute of Science

and Technology, 50 UNIST-gil, Eonyang-eup, Ulju-gun, Ulsan 44919, South Korea

1 ACS Paragon Plus Environment


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ABSTRACT

P2-type manganese-based oxide materials have received attentions as promising cathode materials for sodium ion batteries because of their low cost and high capacity, but their reaction and failure mechanisms are not yet fully understood. In this study, the reaction and failure mechanisms of βNa0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07, 0.25), α-Na0.7MnO2+y, and β-Na0.7MnO2+z are compared to clarify the dominant factors influencing their electrochemical performances. Using a quenching process with various amounts of a Li dopant, the Mn oxidation state in β-Na0.7[Mn1-xLix]O2+y is carefully controlled without the inclusion of impurities. Through various in situ and ex situ analyses including Xray diffraction, X-ray absorption near edge structure spectroscopy, and inductively coupled plasma mass spectrometry, we clarify the dependence of i) reaction mechanisms on disordered Li distribution in the Mn layer, ii) reversible capacities on the initial Mn oxidation state, iii) redox potentials on the JahnTeller distortion, iv) capacity fading on phase transitions during charging and discharging, and v) electrochemical performance on Li dopant vs. Mn vacancy. Finally, we demonstrate that the optimized β-Na0.7[Mn1-xLix]O2+y (x = 0.07) exhibits excellent electrochemical performance including a high reversible capacity of approximately 183 mA h g-1 and stable cycle performance over 120 cycles.

KEYWORDS. sodium ion battery, cathode, layered manganese oxide, mechanism, orthorhombic structure


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INTRODUCTION Li-ion batteries have been widely used as power sources for various applications, including electric vehicles and cell phones. However, the price of lithium reserves has been steeply increasing in the last few years because of the high demand for Li-ion batteries, leading to concerns of a lithium shortage. Consequently, Na-ion batteries have been considered as promising alternatives to Li-ion batteries because of the presence of abundant and low-cost sodium precursors.1-5 Various sodium intercalation materials have been introduced since the early 1980s, including layered transition metal oxide-,6-45 phosphate-,46-51 pyrophosphate-,52-55 fluorophosphates-,56-57 fluorosulfate-58-59 and sulfate-60-61 based materials. Among these, Mn-based layered transition metal oxides15-45 have been recently studied intensively as cathode materials because of their attractive low-cost, high reversible capacity, and lack of toxicity. NaxMnO2+γ (0 < x ≤ 1) forms 2D layered and 3D tunnel structures depending on the synthetic conditions and chemical compositions. Most 2D-structured NaxMnO2+γ are further classified into P2 and O3 types based on the Na+ coordination site and the number of MO2 (M = transition metal) slabs as suggested by Delmas et al.62 According to the phase diagram of NaxMnO2+γ (x ≤ 1) proposed by Parant et al.,63 O3-type NaxMnO2+γ, P2-type NaxMnO2+γ, and 3D NaxMnO2+γ can be obtained when the sodium composition x in NaxMnO2+γ is x = 1, 0.6 ≤ x ≤ 0.7, or x ≤ 0.44, respectively.2, 63 Na0.7MnO2+γ is the most Na-rich phase in P2-type NaxMnO2+γ, and two phases exist at x = 0.7, depending on temperature. One is α-Na0.7MnO2+z (0.05 ≤ z ≤ 0.25) with a hexagonal structure (space group: P63/mmc) and the other is β-Na0.7MnO2+y (y < 0.05) with an orthorhombic structure (space group: Cmcm). α-Na0.7MnO2+z and β-Na0.7MnO2+y are known to be stable at low temperature (< 600°C) and high temperature (> 600°C), respectively.21, 63 Therefore, β-Na0.7MnO2+y is obtained only through a quenching process after heating at high temperature (> 600°C). The average oxidation state of Mn in β-Na0.7MnO2+y is ca. +3.3, resulting in an orthorhombic structure due to the Jahn-Teller distortion of Mn3+. However, αNa0.7MnO2+z (0.05 ≤ z ≤ 0.25) has a hexagonal structure without Jahn-Teller distortion because of the 3 ACS Paragon Plus Environment


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high oxidation state of Mn; Mn oxidation is known to occur during a slow cooling process after heating at high temperature (> 600°C), leading to the formation of a Mn vacancy for charge compensation. P2-type NaxMnO2+γ has a structural advantage compared to O3-type NaxMnO2+γ. Na+ ions in P2type NaxMnO2+γ are distributed over the rectangular face-sharing trigonal prismatic sites in Na layers where Na+ ions diffuse through the wide rectangular face between adjacent trigonal prismatic sites, leading to high ionic diffusivity. However, for the diffusion of Na+ ions in O3-type NaxMnO2+γ, Na+ ions move through a relatively small tetrahedral site between adjacent octahedral sites with high activation energy barriers for cation hopping between adjacent sites.2 In addition, NaxMnO2+γ is known not to show a phase transition from a layered to a spinel structure during cycling, because the formation of spinel NaMn2O4 is thermodynamically unfavorable owing to unstable Na with a large radius in the tetrahedral site of spinel NaMn2O4.64 The layered-to-spinel transformation is one of the failure mechanisms for O3-type LiMO2 (M = Ni1-x-yCoxMny) because O3-type LiMO2 and spinel LiM2O4 have the same oxygen stackings and spinel LiM2O4 is thermodynamically more stable than O3-type Li0.5MO2.65 Moreover, in contrast to LiMn2O4 for Li-ion batteries, Mn dissolution is not a challenging issue for NaxMnO2+γ even though NaPF6-based electrolytes are used, because NaPF6 is more stable than LiPF6,66 resulting in lower formation of HF during cycling (Figure S1). This is another attractive aspect of NaxMnO2+γ. Based on these merits, P2-type NaxMnO2+γ-based materials have been considered promising cathodes for Na ion batteries. For example, NaxMnO2+γ,22-25 NaxMn1-yMyO2+γ (M = Fe,38-40 Co,25-28 Ni,43-45 and Cu41-42), and Mg-,29-32 or Li-doped NaxMn1-yMyO2+γ33-37 have exhibited promising electrochemical performance including a high reversible capacity (> 150 mA h g-1) and improved capacity retention. The reaction mechanism of P2-type NaxMnO2+γ is known to depend on Na+/vacancy ordering in Na layers.10, 67 The substitution of Mn with dopants causes Na+/vacancy disordering, leading to a onephase reaction during charging and discharging. Additionally, the capacity fading of P2-type NaxMnO2+γ was suggested to be attributable to the formation of the OP4 phase after full charging29 and 4 ACS Paragon Plus Environment

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Jahn-Teller distortion due to Mn3+.23 However, the reaction and failure mechanisms of P2-type NaxMnO2+γ are still not fully understood. Therefore, in this study, we clarify the role of dopants, vacancies, and the Mn oxidation state on the reaction and failure mechanisms of P2-type NaxMnO2+γ through the substitution of Mn with Li. Using various amounts of a Li dopant, the Mn oxidation state in P2-type NaxMnO2+γ was successfully controlled without the inclusion of impurities. Various compositions of β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07, 0.25) were prepared, and their reaction mechanism and electrochemical performance were compared to those of β-Na0.7MnO2+y and αNa0.7MnO2+z. Through the comparison of various in situ and ex situ analyses including X-ray diffraction (XRD), X-ray absorption near edge structure (XANES) spectroscopy and inductively coupled plasma mass spectrometry (ICP-MS), we clarify the dependence of i) reaction mechanisms on disordered Li distribution in the Mn layer, ii) reversible capacities on the initial Mn oxidation state, iii) redox potentials on the Jahn-Teller distortion, iv) capacity fading on phase transitions during charging and discharging, and v) electrochemical performance on Li dopant vs. Mn vacancy. Finally, we demonstrate that the optimized β-Na0.7[Mn1-xLix]O2+y (x = 0.07) exhibits excellent electrochemical performance including a high reversible capacity of approximately 183 mA h g-1 and stable cycle performance over 120 cycles.

EXPERIMENTAL SECTION Synthesis. α-Na0.7MnO2+z, β-Na0.7MnO2+y, and β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07, 0.25) were synthesized through a solid state reaction. Na2CO3, Mn3O4, and Li2CO3 precursors were ball-milled at 300 rpm for 1 hour. The mixtures were heated at 900°C in air for 10 hours. α-Na0.7MnO2+z was obtained by slow cooling, and β-Na0.7MnO2+y and β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07, 0.25) were obtained through a quenching process. All samples were stored in an Ar-filled glove box to avoid moisture.



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Material characterization. X-ray diffraction (XRD) patterns of powders and electrodes were obtained using a Bruker D2 PHASER with Cu Kα radiation (λ = 1.5418 Å) operated in the 2θ range of 10−80 °. Lattice parameters were obtained from full pattern matching refinements using the Topas program. For ex situ XRD analysis, airtight sample holders equipped with a Be window were used. The atomic composition of the samples was determined by a Perkin Elmer Elan 6100 inductively coupled plasma-mass spectrometer (ICP-MS). X-ray absorption spectroscopy was performed at the PLS-II 8C beamline of Pohang Accelerator Laboratory in Korea. X-ray absorption near edge structure (XANES) spectra of the Mn K-edge were obtained using a Si(111) double-crystal monochromator in transmission mode at an electron energy of 3 GeV and a current of 370 mA during top-up mode operation. Electrochemical characterization. Electrodes were prepared with active materials, carbon black (super P) and polyvinylidene fluoride (PVdF) in a 7:2:1 weight ratio. Electrochemical performance was evaluated using 2032 coin cells with a Na metal anode, a glass fiber separator, and 1 M NaPF6 in ethylene carbonate (EC) and diethyl carbonate (DEC) (1:1 v/v) electrolyte with 2 wt% fluoroethylene carbonate (FEC) additive. Galvanostatic cycling was performed using a WBCS3000 battery cycler (Wonatech, S. Korea) in a voltage range of 2-3.8 V at a 0.4C rate after the pre-cycle at a 0.1C rate (14 mA g-1 for α-Na0.7MnO2+z, 24 mA g-1 for β-Na0.7MnO2+y, 19 mA g-1 for β-Na0.7[Mn0.93Li0.07]O2+y). Galvanostatic experiments for the dQ/dV plots were carried out with 1 M NaClO4 in EC and PC (1:1 v/v) at a current density of 12 mA g-1. All cells were disassembled in the Ar-filled glove box for ex situ analyses, and electrodes were rinsed with dimethyl carbonate (DMC). In situ synchrotron X-ray Diffraction. For the in situ XRD experiments, 2032 coin cells with a 3.5 mm hole sealed by Kapton film with epoxy resin were used. NaClO4 (1 M) in EC and PC (1:1 v/v) electrolyte and a Na metal anode with a 3 mm hole for the penetration of X-rays were used. In situ galvanostatic experiments were performed at 0.1C rate in a voltage range of 2-3.8 V. In situ XRD patterns of β-Na0.7MnO2+y were obtained at the PLS-II 6D UNIST-PAL beamline of Pohang Accelerator Laboratory (PAL) using a 2D CCD detector (MX225-HS, Rayonix LLC, USA) every 1 min with 30 sec 6 ACS Paragon Plus Environment

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exposure time. In situ XRD patterns of α-Na0.7MnO2+z and β-Na0.7[Mn0.93Li0.07]O2+y were acquired at beamline 5A of PAL using a Mar 345 image-plate (marXperts GmbH, Germany) and were obtained every 3 min with a 30 sec exposure time. The used X-ray wavelengths of the 6D and 5A beamlines were 0.6119 and 0.6926 Å, respectively. For comparison, the resulting XRD patterns were converted into the Cu Kα wavelength (λ = 1.5418 Å).

RESULTS AND DISCCUSION To clarify the reaction and failure mechanisms of P2 β-Na0.7[Mn1-xMx]O2+y, we synthesized a solid solution series of β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07, 0.25) using a quenching process, and the structural and electrochemical characteristics of the compounds were compared with those of αNa0.7MnO2+z and β-Na0.7MnO2+y. The ionic radius of Li+ is so small that Li+ is more stable in octahedral Mn sites than prismatic Na sites. Hence, Li selectively replaces Mn, forming β-Na0.7[Mn1-xLix]O2+y.33-36 Figure 1 shows the XRD patterns of β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07, 0.25), α-Na0.7MnO2+z, and β-Na0.7MnO2+y. The XRD peaks assigned by triangles in Figure 1 are the characteristic peaks corresponding to the (02l) planes of an orthorhombic structure (β-Na0.7MnO2+y), revealing that βNa0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07) has an orthorhombic structure. Rietveld refinement of the XRD patterns in the space group Cmcm was carried out for β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07) (Figure 2). The lattice parameters, agreement factors, detailed structural information, and crystal structures are summarized in Table 1, Table S1-3, and Figure S2-4. As we substitute Mn3+ by Li+ to obtain β-Na0.7[Mn1-xLix]O2+y, two Mn3+ are oxidized into Mn4+ per Li+ for charge compensation. This leads to an increase in the average oxidation state of Mn in β-Na0.7[Mn1-xLix]O2+y. As a result, the JahnTeller effect is alleviated, thereby forming a less distorted orthorhombic structure with more Li in βNa0.7[Mn1-xLix]O2+y. This behavior is supported by a gradual shift in the characteristic (02l) peaks of the orthorhombic structure with increasing Li substitution (Figure 1). As x in β-Na0.7[Mn1-xLix]O2+y



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Figure 1. (a) Powder XRD patterns and (b) enlarged XRD patterns between 30~50 ° for βNa0.7MnO2+y, β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07, 0.25), and α-Na0.7MnO2+z. (Red line:β-Na0.7MnO2+y,

sky-blue

Na0.7[Mn0.96Li0.04]O2+y,

line:

orange

β-Na0.7[Mn0.98Li0.02]O2+y,

line:

yellow-green

β-Na0.7[Mn0.93Li0.07]O2+y,

green

line: line:

ββ-

Na0.7[Mn0.75Li0.25]O2+y, blue line: α-Na0.7MnO2+z)


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Figure 2. X-ray powder diffraction patterns (dotted curve), the Rietveld-refined profiles (red line), and

difference curves (blue line) for (a) β-Na0.7[Mn0.98Li0.02]O2, (b) β-

Na0.7[Mn0.96Li0.04]O2, and (c) β-Na0.7[Mn0.93Li0.07]O2.



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Table 1. Lattice constants for orthorhombic β-Na0.7[Mn1-xLix]O2 (x = 0.02, 0.04, 0.07).

a (Å)

b (Å)

c (Å)

β-Na0.7[Mn0.98Li0.02]O2

2.84648(22)

5.25402(62)

11.2973(13)

β-Na0.7[Mn0.96Li0.04]O2

2.84915(22)

5.24960(76)

11.2715(12)

β-Na0.7[Mn0.93Li0.07]O2

2.8848963

5.0996316

11.27833(29)

Figure 3. (a) Normalized Mn K-edge XANES spectra and (b) the corresponding second derivative spectra. (Region A: pre-edge, Region B: main-edge, Region C: main peak)


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increases, the XRD peak positions corresponding to the (02l) planes gradually shift toward those of (11l) planes with higher 2θ angles in the orthorhombic structure, finally merging into those of (10l) planes in the hexagonal structure (α-phase). This solid solution behavior of β-Na0.7[Mn1-xLix]O2+y indicates that Mn is gradually oxidized with increasing Li substitution, leading to a gradual phase transition from an orthorhombic to a hexagonal structure. Finally, the hexagonal structure (α-phase) without Jahn-Teller distortion was obtained, as x in Na0.7[Mn1-xLix]O2+y reached the value of 0.25. However, a Li2MnO3 impurity was also observed in the hexagonal α-Na0.7[Mn1-γLiγ]O2+z, when x in Na0.7[Mn1-xLix]O2+y is 0.25. The higher average oxidation state of Mn in β-Na0.7[Mn1-xLix]O2+y than in β-Na0.7MnO2+y was supported by XANES analysis. The average oxidation state of Mn in β-Na0.7[Mn1xLix]O2+y

was between those of Mn in α-Na0.7MnO2+z and β-Na0.7MnO2+y, as shown in Figure 3.

Figure 4 shows the voltage profiles of β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07), αNa0.7MnO2+z, and β-Na0.7MnO2+y. As x in β-Na0.7[Mn1-xLix]O2+y increases, we observe three characteristic tendencies, viz. i) a change in the voltage profiles from stepwise plateaus to sloping profiles, ii) a decrease in reversible capacities, and iii) a decrease in redox potentials. The complex stepwise plateaus were observed in the voltage profiles of β-Na0.7MnO2+y, suggesting that βNa0.7MnO2+y undergoes a complex reaction mechanism including various one-phase and two-phase reactions. The complex reaction mechanism of β-Na0.7MnO2+y is attributed to the formation of intermediate phases with various Na+/vacancy arrangements during sodiation and desodiation.10, 67 In contrast to β-Na0.7MnO2+y, α-Na0.7MnO2+z shows sloping voltage profiles, indicating that α-Na0.7MnO2+z proceeds through a one-phase reaction during sodiation and desodiation. This suggests that the Na+/vacancy distribution of α-Na0.7MnO2+z is disordered during sodiation and desodiation. Considering the voltage profiles of β-Na0.7[Mn1-xLix]O2+y, it is remarkable that their voltage profiles become similar to that of α-Na0.7MnO2+z as x in β-Na0.7[Mn1-xLix]O2+y increases. This indicates that fewer phase transitions occur during sodiation and desodiation as x in β-Na0.7[Mn1-xLix]O2+y increases. The Delmas group recently reported that Na+/vacancy ordering in P2-NaxCoO2 is affected by the electrostatic 11 ACS Paragon Plus Environment


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repulsion between Co and Na+ in Naf and Nae sites, where Naf and Nae sites indicate the two prismatic sites sharing faces and edges, respectively.10 However, Li dopants are considered to be randomly distributed over the Mn layer in β-Na0.7[Mn1-xLix]O2+y, because no superstructure peaks due to the ordered Li/Mn distribution over Mn layer were observed in the XRD patterns (Figure 1). Consequently, the disordered Li/Mn distribution caused inhomogeneous electrostatic repulsion between Mn and Na+, leading to Na+/vacancy disordering. These reaction mechanisms are further supported by in situ XRD analyses, which will be discussed in the following section. The second tendency that is observed in the voltage profiles (Figure 4) is a change in redox potentials due to Jahn-Teller distortion. β-Na0.7MnO2+y and α-Na0.7MnO2+z have orthorhombic and hexagonal structures with and without Jahn-Teller distortion, respectively. As shown in Figure 4a, characteristic plateau or plateau-like sloping profiles near approximately 2 V vs. Na/Na+ were observed in the voltage profiles of β-Na0.7MnO2+y, α-Na0.7MnO2+z, and β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07). Plateaus in the voltage profiles shown in Figure 4a correspond to peaks in the dQ/dV plots shown in Figure 4b. α-Na0.7MnO2+z and β-Na0.7MnO2+y have dQ/dV peaks at approximately 2.17 and 2.36 V vs. Na/Na+, respectively. This suggests that the different redox potentials of α-Na0.7MnO2+z and β-Na0.7MnO2+y originate from their different crystal structures, where the redox potential of the orthorhombic structure is higher than that of the hexagonal structure. This is supported by the gradual peak shift in the dQ/dV plots for β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07) from 2.32 to 2.19 V vs. Na/Na+ with increasing x in β-Na0.7[Mn1-xLix]O2+y. As x in β-Na0.7[Mn1-xLix]O2+y increases, lessdistorted orthorhombic structures were formed because of the alleviated Jahn-Teller effect, resulting in a decrease in redox potential. The distortion degree due to Jahn-Teller effect can be roughly estimated using the lattice parameter ratio b/a.24 Figure 5 shows the linear relationship between distortion degree and redox potential, suggesting that the redox potential of β-Na0.7MnO2+y is controllable by modifying the crystal structure through Mn substitution. The distortion degrees of β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07) shown in Figure 5 were calculated using the lattice parameters in Table 1. Also, the 12 ACS Paragon Plus Environment

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distortion degrees of α-Na0.7MnO2+z and β-Na0.7MnO2+y were obtained from those of β-Na2/3MnO2 and α-Na2/3MnO2 in ref. 24.24 The third tendency that is observed in the voltage profiles (Figure 4) is a decrease in the reversible capacities with an increase in the initial Mn oxidation state. β-Na0.7MnO2+y (232 mA h g-1) and α-Na0.7MnO2+z (152 mA h g-1) delivered the highest and lowest reversible capacities. As x in βNa0.7[Mn1-xLix]O2+y increased, we observed a decrease in the reversible capacities of β-Na0.7[Mn1xLix]O2+y,

despite their similar theoretical specific capacities (260 mA h g-1 for β-Na0.7MnO2, 257 mA h

g-1 for β-Na0.7[Mn0.98Li0.02]O2, 255 mA h g-1 for β-Na0.7[Mn0.96Li0.04]O2, and 250 mA h g-1 for βNa0.7[Mn0.93Li0.07]O2). This is attributed to an increase number of immobile Na+ ions in β-Na0.7[Mn1xLix]O2+y,

depending on the initial Mn oxidation state. For example, the oxidation states of 0.7 Mn and

0.3 Mn in β-Na0.7MnO2+y are +3 and +4, respectively (assuming that y is negligible). Therefore, all Na+ ions (0.7 Na+) in β-Na0.7MnO2+y can be deintercalated during charging (desodiation), leading to the formation of MnIVO2, assuming y = 0. Consequently, 1.0 Na+ can be re-intercalated during the next discharging (sodiation), leading to the formation of NaMnIIIO2, because all the prismatic sites for Na+ in β-Na0.7MnO2+y empty after desodiation. For α-Na0.7MnO2+z, however, z is known to be approximately 0.05 - 0.25, which is not negligible. Therefore, the oxidation states of (0.7-2z) Mn and (0.3+2z) Mn in α-Na0.7□0.3+0.5z[Mn□0.5z]O2+z are +3 and +4, respectively (□ indicates a vacancy derived from a slow cooling process). Since only (0.7-2z) Mn3+ exists in the initial state, only (0.7-2z) Na+ among 0.7 Na+ in α-Na0.7□0.3+0.5z[Mn□0.5z]O2+z 1.5z[Mn

can

be

deintercalated

during

desodiation,

forming

α-Na2z□1-

IV

□0.5z]O2+z after full desodiation. Thereby, not all of the Mn4+ can be reduced to Mn3+ during the

next sodiation, because some Na+ sites were still occupied even after full desodiation. Specifically, only (1-1.5z) Mn4+ can be reduced, forming α-Na1+0.5z[Mn(III)1-1.5zMn(IV)1.5z□0.5z]O2+z, because only (1-1.5z) Na+ sites are empty after full desodiation. In other words, (2z) Na+ in α-Na0.7MnO2+z are immobile, occupying Na+ sites during cycling. This results in the smaller reversible capacity of α-Na0.7MnO2+z than in β-Na0.7MnO2+y (y ≈ 0) in which 1.0 Na+ is reversibly intercalated during cycling. In the same 13 ACS Paragon Plus Environment


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Figure 4. (a) Voltage profiles and (b) dQ/dV plots for β-Na0.7MnO2+y, β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07), and α-Na0.7MnO2+z.


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Figure 5. Relationship between distortion degree and redox potential for β-Na0.7MnO2+y, βNa0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07), and α-Na0.7MnO2+z.



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manner, as the amount of Li dopant increases, the number of immobile Na+ ions in β-Na0.7[Mn1xLix]O2+y

increases. This is because, as more Li replaces Mn, more Mn is oxidized into +4 without

changing the composition of Na+ in the initial state, forming β-Na0.7[Mn(III)0.7-3xMn(IV)0.3+2xLix]O2+y (assuming y ≈ 0) where (3x) Na+ are immobile. To obtain evidence of the Li dopant effect causing fewer phase transitions during sodiation and desodiation (shown in the voltage profiles), we performed in situ XRD analyses for β-Na0.7[Mn1xLix]O2+y

(x = 0.07), α-Na0.7MnO2+z, and β-Na0.7MnO2+y (Figure 6). β-Na0.7MnO2+y shows complex

phase transitions during charging and discharging through various one-phase and two-phase reactions, which is consistent with the complex stepwise voltage profiles of β-Na0.7MnO2+y. In contrast to βNa0.7MnO2+y, α-Na0.7MnO2+z exhibits a one-phase reaction over the voltage range from 2.0 to 3.8 V vs. Na/Na+, where XRD peaks are gradually shifted without any evolution of new peaks during sodiation and desodiation. This is also consistent with the sloping voltage profiles of α-Na0.7MnO2+z. Moreover, a one-phase reaction mechanism was also observed for β-Na0.7[Mn1-xLix]O2+y (x = 0.07) during sodiation and desodiation, which is also consistent with the sloping voltage profiles of β-Na0.7[Mn1-xLix]O2+y (x = 0.07). Figure 7 presents the detailed change in the c-axis lattice parameter for β-Na0.7[Mn1-xLix]O2+y (x = 0.07), α-Na0.7MnO2+z and β-Na0.7MnO2+y during sodiation and desodiation. We observed expansions of ca. 1.0% and 0.14% during desodiation in the c-axis lattice parameter of α-Na0.7MnO2+z and βNa0.7[Mn1-xLix]O2+y (x = 0.07), respectively. The expansion along the c axis is ascribed to the increased electrostatic repulsion between oxygen layers due to the Na deficiency. However, it is remarkable that the change in the c-axis lattice parameter of β-Na0.7[Mn1-xLix]O2+y (x = 0.07) was much smaller than that of α-Na0.7MnO2+z, despite the fact that β-Na0.7[Mn1-xLix]O2+y (x = 0.07) delivered more reversible capacity (ca. 180 mA h g-1) than α-Na0.7MnO2+z (ca. 150 mA h g-1). This suggests that the Li dopant in β-Na0.7[Mn1-xLix]O2+y is more effective in alleviating the expansion along the c-axis during desodiation than the Mn vacancy in α-Na0.7MnO2+z.


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Figure

6.

In

situ

synchrotron

XRD

patterns

for

(a)

β-Na0.7MnO2+y,

(b)

β-

Na0.7[Mn0.93Li0.07]O2+y, and (c) α-Na0.7MnO2+z.



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Figure 7. c lattice parameters and the corresponding voltage profiles of (a) β-Na0.7MnO2+y, (b) β-Na0.7[Mn0.93Li0.07]O2+y, and (c) α-Na0.7MnO2+z.


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The O2,44 OP429 and Z-phases45 are known to be observed after full desodiation of Mn-based P2-type layered oxide materials. However, these phases were not observed in our in situ XRD analyses of β-Na0.7[Mn1-xLix]O2+y (x = 0.07), α-Na0.7MnO2+z, and β-Na0.7MnO2+y in the voltage range from 2.0 to 3.8 V vs. Na/Na+ (Figure 6). In contrast to the in situ XRD analyses, the OP4 phase was observed in the ex situ XRD pattern of β-Na0.7MnO2+y at 3.8 V vs. Na/Na+, as shown in Figure 8. This is because the P2-type Na-poor phases obtained after desodiation are metastable. This is supported by the additional ex situ XRD analyses that observed the phase transition of the Na-poor phase during rest for 4 weeks after charging to 3.5 V vs. Na/Na+ (Figure 8). The OP4 phase was not initially observed immediately after charging to 3.5 V vs. Na/Na+, but observed after storage for 4 weeks. The peak ca. 17 ° corresponds to the OP4 phase. This reveals that not being able to observe the OP4 phase during charging and discharging (Figure 6) is attributed to the phase transition from P2 to OP4 being kinetically hindered, although the OP4 phase is more thermodynamically stable than P2 for Na-poor phases. Moreover, we compared the cycle performance of β-Na0.7[Mn1-xLix]O2+y (x = 0.07), αNa0.7MnO2+z, and β-Na0.7MnO2+y, as shown in Figure 9a. The corresponding voltage profiles are presented in Figure 10. While poor capacity retention was observed for β-Na0.7MnO2+y, β-Na0.7[Mn1xLix]O2+y

(x = 0.07) and α-Na0.7MnO2+z exhibited excellent cycle performance, including negligible

capacity fading over 120 cycles. This suggests that the stable cycle performances of β-Na0.7[Mn1xLix]O2+y

(x = 0.07) and α-Na0.7MnO2+z are attributed to their one-phase reaction mechanism (no phase

transition during sodiation and desodiation), while the poor cycle performance of β-Na0.7MnO2+y is due to the complex phase transitions during sodiation and desodiation. Galvanostatic cycling of βNa0.7MnO2+y was performed in the two independent voltage ranges of 2.0-2.9 and 2.4-3.8 V vs. Na/Na+ (Figure 9b). Over both voltage ranges, β-Na0.7MnO2+y delivered similar reversible capacities of approximately 112 and 135 mA h g-1, respectively, but poorer cycle performance was observed in the higher voltage range despite the fact that Jahn-Teller distortion occurs in the lower voltage range.



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Figure 8. Ex situ XRD patterns of β-Na0.7MnO2+y at (a) 3.8 V with no rest, (b) 3.5 V with no rest, and (c) 3.5 V vs. Na/Na+ after resting for 4 weeks.


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Figure 9. Cycle performance of (a) β-Na0.7MnO2+y, β-Na0.7[Mn0.93Li0.07]O2+y, and αNa0.7MnO2+z in the voltage range from 2.0 to 3.8 V vs. Na/Na+, and (b) cycle performance of β-Na0.7MnO2+y in the different voltage ranges.



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Figure 10. Voltage profiles of (a) β-Na0.7MnO2+y, (b) β-Na0.7[Mn0.93Li0.07]O2+y, and (c) αNa0.7MnO2+z at various cycle numbers.


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In addition to the Jahn-Teller distortion, the OP4 phase was also not observed after desodiation, as shown in Figure 6. Therefore, the poor cycle performance of β-Na0.7MnO2+y is not due to the JahnTeller distortion or the formation of the OP4 phase during charging. However, the binary phase transition is known to be one of the principal factors causing the capacity fading of layered cathode materials (i.e., Li[Ni1-x-yCoxMny]O2) for Li-ion batteries because interfacial strain can lead to the formation of cracks at grain boundaries within a particle, resulting in increasing polarization.68-69 Therefore, the poor cycle performance of β-Na0.7MnO2+y is considered more dependent on the binary reaction mechanism than the Jahn-Teller distortion or OP4 phase formation. We further performed ex situ ICP analysis to examine whether Li+ dopants are immobile during cycling. Figure 11a reveals that the amount of Li in β-Na0.7[Mn1-xLix]O2+y (x = 0.07) decreases as the cycle number increases. This indicates that Li+ dopants in β-Na0.7[Mn1-xLix]O2+y (x = 0.07) were deintercalated during charging (desodiation), leading to the formation of vacancies such as β-Na0.7φ[Mn1-xLix-δ□δ]O2+y.

In addition, we observed that the amount of Li dopants decreased and increased

during charging and discharging, respectively, as shown in Figure 11b. This means that not only Na+ but also Li+ in β-Na0.7[Mn1-xLix]O2+y are reversibly intercalated and deintercalated as charge carriers during cycling.36-37 However, no change in voltage profiles was observed during cycling, as shown in Figure 10b. This suggests that Li dopants and the Mn vacancy in β-Na0.7[Mn1-xLix-δ□δ]O2+y play the same role in phase transitions during charging and discharging. This is consistent with αNa0.7[Mn□0.5z]O2+z proceeding through a one-phase reaction. Similar to the disordered Li distribution in β-Na0.7[Mn1-xLix]O2+y (x = 0.07), the random distribution of Mn vacancies in α-Na0.7MnO2+z also caused inhomogeneous electrostatic repulsion between Mn and Na+, leading to the Na+/vacancy disordering during sodiation and desodiation. As a result, a one-phase reaction is observed for α-Na0.7MnO2+z.



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Figure 11. The ratios of Na and Li per Mn in β-Na0.7[Mn0.93Li0.07]O2+y (a) after charging at various cycle numbers and (b) for charging and discharging during the 10th cycle.


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CONCLUSIONS Through various in situ and ex situ analyses including XRD, XANES and ICP-MS, we compared the reaction and failure mechanisms of β-Na0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07, 0.25), β-Na0.7MnO2+y, and α-Na0.7MnO2+z to clarify the dominant factors influencing their electrochemical performances. We further demonstrated that the optimized β-Na0.7[Mn1-xLix]O2+y (x = 0.07) exhibited excellent electrochemical performance including a high reversible capacity of approximately 183 mA h g-1 and stable cycle performance over 120 cycles. Using a quenching process with various amounts of a Li dopant, the Mn oxidation state in P2-type β-Na0.7[Mn1-xLix]O2+y could be controlled without the inclusion of impurities. As the level of lithium substitution in β-Na0.7[Mn1-xLix]O2+y increased, the average oxidation state of Mn in β-Na0.7[Mn1-xLix]O2+y increased, resulting in the formation of lessdistorted orthorhombic structures due to the attenuated Jahn-Teller effect. A decrease in the redox potential of β-Na0.7[Mn1-xLix]O2+y was observed with increasing Li, depending on the extent of the JahnTeller distortion. Li dopants in β-Na0.7[Mn1-xLix]O2+y were randomly distributed over the Mn layer, leading to an inhomogeneous electrostatic repulsion between Mn and Na+. This causes the Na+/vacancy disordering over the Na layer, resulting in a one-phase reaction of β-Na0.7[Mn1-xLix]O2+y (x = 0.07). Compared to β-Na0.7MnO2+y (with its complex phase transitions during sodiation and desodiation), the more stable cycle performances of β-Na0.7[Mn1-xLix]O2+y (x = 0.07) and α-Na0.7MnO2+z were attributed to their one-phase reaction mechanisms (undergoing no phase transition during sodiation and desodiation). The poor cycle performance of β-Na0.7MnO2+y was not due to Jahn-Teller distortion or the formation of an OP4 phase during charging. In particular, the OP4 phase was not observed during charging and discharging, because the transition from the metastable P2 Na-poor phases to the OP4 phase was kinetically hindered. The amount of Li+ dopants in β-Na0.7[Mn1-xLix]O2+y gradually decreased during cycling, indicating that Li+ ions deintercalated with the formation of vacancies. However, vacancies generated during cycling played the same role in phase transitions during charging and discharging. In addition, as x in β-Na0.7[Mn1-xLix]O2+y increased, the number of immobile Na+ ions in β25 ACS Paragon Plus Environment


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Na0.7[Mn1-xLix]O2+y increased leading to a decrease in reversible capacities. Finally, we believe that these reaction and failure mechanisms provide a new avenue for developing promising P2-type cathode materials.

ASSOCIATED CONTENT Supporting Information. The Supporting Information is available free of charge via the Internet at http://pubs.acs.org. Mn concentrations dissolved in electrolytes for various cathode materials and atomic parameters for βNa0.7[Mn1-xLix]O2+y (x = 0.02, 0.04, 0.07) refined from powder XRD data (PDF)

AUTHOR INFORMATION Corresponding Author * E-mail: [email protected] Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT This research was supported by the National Research Foundation of Korea (NRF) Grant (No. NRF2016R1A2B3015956), by Korea Electrotechnology Research Institute (KERI) Primary research 26 ACS Paragon Plus Environment

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program through the National Research Council of Science & Technology funded by the Ministry of Science, ICT and Future Planning (MSIP) (No. 14-12-N0101-69), by the Materials and Components Technology Development Program of MOTIE/KEIT, Republic of Korea [10050477, Development of separator with low thermal shrinkage and electrolyte with high ionic conductivity for Na-ion batteries], and by the KIST Institutional Program (Project No. 2E26330). Experiments at PLS-II 6D UNIST-PAL beamline were supported in part by MEST, POSTECH, and UNIST UCRF.

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