Probing the Structural Transition Kinetics and Charge Compensation

Research Article Cite This: ACS Appl. Mater. Interfaces 2019, 11, 24122−24131

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Probing the Structural Transition Kinetics and Charge Compensation of the P2-Na0.78Al0.05Ni0.33Mn0.60O2 Cathode for Sodium Ion Batteries Yuansheng Shi,†,§ Shuai Li,† Ang Gao,† Jieyun Zheng,‡ Qinghua Zhang,*,‡ Xia Lu,*,†,§ Lin Gu,‡ and Dapeng Cao*,†

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†

State Key Laboratory of Organic−Inorganic Composites, Beijing Advanced Innovation Center for Soft Matter Science and Engineering, Beijing University of Chemical Technology, Beijing 100029, P. R. China ‡ Institute of Physics, Chinese Academy of Sciences, Beijing 100190, P. R. China § School of Materials, Sun Yat-sen University, Guangzhou 510275, P. R. China S Supporting Information *

ABSTRACT: Although the layered P2-type Na0.67Ni0.33Mn0.67O2 materials show high discharge voltage and specific capacity, they suffer from severe structural instabilities and surface reaction upon Na exchange for sodium-ion batteries (SIBs). Therefore, it is quite necessary to reveal the underlying structural evolution mechanism and diffusion kinetics to improve the structural/electrochemical stability for application in SIBs. Here, we synthesize a P2-type Na0.78Al0.05Ni0.33Mn0.60O2 material by a small dose of Al replacing the Mn, aiming at enhancing the structural stability without sacrificing the average discharge voltage and theoretical capacity. The etching X-ray photoelectron spectroscopy and energydispersive X-ray mapping/line scan results indicate that the Al doping induces dual effects of the Al2O3 surface coating and the bulk lattice doping, which efficiently suppress the accumulation of structural irreversible changes from P2 to O2, the volume changes, and surface reactions at high voltage. Obvious improvements are further found on the diffusion kinetics of Na ions as well as the decrease of overall voltage polarization. Interestingly, the dual effects of Al doping lead to the significant increase of capacity retention after 50 cycles and improvement of rate capability compared with the undoped counterpart between 2.0 and 4.5 V. Hence, this work sheds new light on stabilizing the P2-Na−Ni−Mn−O materials, which provides a rewarding avenue to develop better SIBs. KEYWORDS: Na0.78Al0.05Ni0.33Mn0.60O2, Al doping, in situ XRD, scanning transmission electron microscopy, Na ion batteries

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INTRODUCTION

occupation of Na ions among the MO6 slabs and therefore exhibit superior electrochemical performance from both the experimental results and the theoretical simulations4−6 with regard to the O3-type counterparts. The Ni-doped NaxMnO2 (P2-Na2/3Ni1/3Mn2/3O2) were first investigated by Dahn7 in 2001, where all of the Na ions in the layered P2Na2/3Ni1/3Mn2/3O2 can be intercalated and deintercalated reversibly based on the Ni2+/Ni4+ redox couples. It can deliver a capacity as high as 160 mAh g−1 with an average discharge voltage of ca. 3.6 V versus Na+/Na.8 Moreover, the P2Na2/3Ni1/3Mn2/3O2 cathode exhibits enhanced air stability9 and cost-effectiveness with respect to the other types of P2 and O3 materials.10 However, this material suffers from fast capacity fading and discharge voltage decay, which probably result from the P2−O2 phase transition at the voltage of ca. 4.3

Compared with the popular lithium ion batteries, sodium-ion batteries (SIBs), with the similar “rocking chair” mechanism,1 are found to be one of the good alternatives due to their low cost and wide applications in large-scale energy storage systems.2 However, developing a suitable cathode material is one of the most important steps to achieve the commercialization of SIBs. Among the available candidates, sodium layeredtype oxides (NaxTMO2, 0 < x < 1 and TM: transition metals) have attracted increasingly more attention due to their high theoretical capacity, moderate operation voltage, good structural stability, and easy availability, which show promising prospects to promote the availability of SIBs in the long run. The layered materials of NaxMO2 were categorized initially into two families by Delmas et al.,3 namely, the O3 and P2 types, according to the stacking sequence of alkali ions along the [001] direction. Previous investigations indicate that the P2-type cathode materials have a more stable Na+ ion diffusion trajectory and a larger interlayer spacing due to the prismatic © 2019 American Chemical Society

Received: April 13, 2019 Accepted: June 12, 2019 Published: June 12, 2019 24122

DOI: 10.1021/acsami.9b06233 ACS Appl. Mater. Interfaces 2019, 11, 24122−24131

Research Article

ACS Applied Materials & Interfaces

Figure 1. Phase and morphology of the as-prepared samples. The XRD Rietveld refinements of (a) Na0.66Ni0.33Mn0.67O2 and (b) Na0.78Al0.05Ni0.33Mn0.60O2 materials, where experimental XRD (red circles), simulated data (black solid line), Bragg reflection peaks (green solid ticks), and the difference curve (blue line) are shown. (c) TEM-SAED, (d) etching XPS profiles, and (e) elemental mapping results of the Na0.78Al0.05Ni0.33Mn0.60O2 sample.

V versus Na+/Na and the large volume change upon overdesodiation. Consequently, many strategies were proposed to overcome these drawbacks of P2-Na2/3Ni1/3Mn2/3O 2 electrode materials. The first choice is surface modification, an effective way to suppress the unavoidable side reactions and particle pulverization.11,12 However, the lattice mismatch leads to the exfoliation of the coating layer after long-term cycling. In addition, no evidence is found that the surface coating works for the voltage decay upon cycling. The second one is elemental doping, that is, using low valence metal ions (Li+,13,14 Mg2+,15,16 Zn2+,17 Ca2+,18 Cu2+,19 etc.), and multimetal ion doping20−22 to increase the valence state of manganese ions and to reduce the contents of nickel as well. Of special interest is the fact that the decrease of nickel to some extent does not deteriorate the cycling reversibility as attested by Konarov et al.,23 but it will decrease the overall discharge voltage. Therefore, if the above two strategies are synergized together, better electrochemical performance would be expected for the P2-Na2/3Ni1/3Mn2/3O2 electrode. To date, the underlying mechanism of Al doping into a P2type-Na−Ni−Mn−O material is still unclear, although Aldoped materials with improved electrochemical performance have been reported in Na−Ni−Mn−O24,25 and Na−Co−Mn− O systems.26 Previous reports disclosed that there are several known factors responsible for battery capacity and voltage fading, including electrolyte decomposition,27 the Jahn−Teller distortion,28,29 transition-metal migration, intergranular cracking,30 dissolution of divalent manganese,12 instability of lattice oxygen,31−34 and structural rearrangement.7 Nonetheless, it is still difficult to elucidate the effect of Al doping on the structural/electrochemical stabilities of the P2Na2/3Ni1/3Mn2/3O2 cathode for SIBs. In this work, the layered Na0.78Al0.05Ni0.33Mn0.60O2 materials are prepared by a facile sol−gel method as the cathode material for SIBs. The dual effects of Al doping are identified by etching X-ray photoelectron spectroscopy (XPS) and elemental mapping. To analyze the capacity fading mechanism of Aldoped materials, the detailed structural evolution is monitored using in situ X-ray diffraction (XRD), ex situ spherical aberration-corrected scanning transmission electron micros-

copy (STEM) techniques, and electron energy loss spectroscopy (EELS). The results indicate that Al doping induces dual effects of Al2O3 surface coating and bulk lattice doping, which can significantly improve the stability of the P2-type Na−Ni− Mn−O layered cathodes because the dual effects efficiently suppress unavoidable side reactions, that is, the accumulation of structural irreversible changes from the P2 to O2 phase and the volume changes at high voltage after Al doping.

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EXPERIMENTAL SECTION

Preparation of P2-Type Al-Substituted Materials. All chemicals were sourced from Aladdin without further purifications. The Na0.78Al0.05Ni0.33Mn0.60O2 and Na0.67Ni0.33Mn0.67O2 samples were prepared by a facile sol−gel-based method. Sodium nitrate (NaNO3, Aladdin), aluminum nitrate (Al(NO3)3·9H2O, Aladdin), nickel nitrate(Ni(NO3)2·6H2O, Aladdin), and manganese nitrate with 50 wt % solution in H2O (Mn(NO3)2, Aladdin) were dissolved in 20.0 mL of deionized water stepwise according to the stoichiometric ratios. About 4−5 wt % excess of sodium nitrate was added to make up for the Na loss during synthesis. Then, the well-mixed solution was added dropwise into the citric acid solution. The resultant mixture was continuously stirred and evaporated to obtain the gel and then dried in air for 12 h. After that, the powder was preheated in a muffle furnace at 400 °C for 6 h. Then, it was ground in an agate mortar for 1 h and heated in a muffle furnace at 950 °C for 15 h in air and slowly cooled down to room temperature to obtain the targeted samples. Characterization. The X-ray diffraction (XRD) measurements were carried out on a Bruker D8 phaser diffractometer with a Cu Kα line (λ = 1.5418 Å) with the continuous scanning mode in the 2θ range of 10−120°. The XRD pattern was refined using FullProf Suite software. Chemical compositions of the samples were determined by an inductively coupled plasma/atomic emission spectrometer (ICP/ AES). The Al 2p electronic state analyses were carried out by X-ray photoelectron spectroscopy (XPS, Thermo Fisher ESCALAB 250 Xray equipped with a twin-anode Mg Kα X-ray source). The binding energy values were all calibrated using the C 1s peak at 284.8 eV. Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were carried out with JEOL JSM-6360 L V (Japan) and JEM-2100F (operating at 200 kV), respectively. The ultraviolet (UV)−visible (vis) light absorption spectrum was tested on an Agilent Cary 5000 UV spectrophotometer to get the exact optic band gap of the as-prepared samples. The in situ XRD diffraction was measured using the Bruker D8 phaser diffractometer with a 24123


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Figure 2. Electrochemical performance of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode. (a) Charge/discharge curves of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode between 2.0 and 4.5 V at C/10; (b) CV curves of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode at 0.1 mV s−1. (c) Rate performance and (d) cycling performance of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode, where the electrochemical performance of the Na0.66Ni0.33Mn0.67O2 electrode is shown for comparison as well. homemade in situ cell. The current density was set to be 10 mA g−1 during cycling. Structural evolution at the atomic scale and electron energy loss spectroscopy (EELS) were carried out by JEM ARM200F (JEOL, Tokyo, Japan), equipped with double CEOS (Heidelberg, Germany) probe aberration correctors. High-angle annular dark-field STEM (HAADF STEM) and low-angle bright-field STEM (LABF STEM) were performed at 200 kV. The acquired raw data was processed with Mutivariate Statistical Analysis v4.3 and ABSF filters. Electrochemical Measurements. The working electrode was fabricated by casting a mixture of the active material, the conductive agent (acetylene black: AB), and poly(vinylidene difluoride) (PVDF) in a weight ratio of 8:1:1 (Na0.78Al0.05Ni0.33Mn0.60O2/AB/PVDF) onto an aluminum foil. The electrode was dried at 80 °C in a vacuum oven and then punched into circular pieces with a diameter of 12 mm. Before being transferred to the Ar-filled glovebox (MIKROUNA, O2 and H2O < 0.1 ppm), the electrode disks need to be dried at 100 °C for 10 h in a vacuum oven. The cycling tests of the working electrodes were performed in CR2032 coin cells. The sodium metal and the glass microfiber (Whatman, GF/D) were used as a reference electrode and a separator, respectively. The electrolyte was 1 M NaPF6 in a solution of propylene carbonate with 5% fluoroethylene carbonate. Galvanostatic cycling tests were conducted by the Land battery testing system (LAND CT-2001A Instrument, Wuhan, China) after resting for 10 h. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) were carried out using an electrochemical workstation (CHI 660D). The CV was measured in a voltage range of 2.0−4.5 V at a scan rate of 0.1 mV s−1. The alternating current impedance measurements were carried out in a frequency range from 0.1 MHz to 10 mHz with an amplitude voltage of 10 mV. The galvanostatic intermittent titration technique (GITT) was employed to get the apparent chemical diffusion coefficients of the sodium ion with a current density of 25 μA for 1.0 h, followed by a 5 h relaxation to reach probably electrochemical equilibrium states. All of these electrochemical measurements were carried out at room temperature.

Na0.78Al0.05Ni0.33Mn0.60O2 agree well with the hexagonal P2 structure (the space group of P63/mmc).35 The refinement results also indicate that the P2 phase remains unchanged after Al doping. Owing to the similar X-ray scattering factors of Ni and Mn, it is difficult to identify the Ni2+/Mn4+ arrangements in the (NiMn)O6 slabs inside the (002) plane. Compared with the high-Ni O3-type layered cathode materials,36 the P2Na0.78Al0.05Ni0.33Mn0.60O2 sample also encounters the air instability issue, where the splitting of (002), (004) peaks and a new peak at ∼38° are witnessed after being exposed in air for 1 month (Figure S1). The phase identification is further characterized by selected area electron diffraction (SAED), as shown in Figure 1c, where the bright diffraction dots correspond to the P2 phase exactly. Moreover, the ordered array of dark SAED spots around bright ones indicates the presence of superstructures in the as-prepared materials. The reasons for the superstructure formation are still under investigation. The chemical components of the samples are determined by means of ICP/AES. The results (Tables S3) suggest that the elemental ratio of the as-prepared samples matches well with the nominal ones. The SEM images (Figure S2) show that the particle size is about 2−5 μm, and the incorporation of Al makes the particle surface smoother. Figure 1d shows the etching Al 2p XPS peak at different depths of the Al-doped particle. The peak position at ∼74.3 eV is ascribed to the binding energy of aluminum atoms in an oxygen environment, that is, the aluminum oxide,37 whereas the other one at ∼73.3 eV is in good agreement with the profile of LiCo0.85Al0.15O2.38 With the increase of depth, the peak of the Al 2p shifts from the surficial 74.3 eV to the bulk 73.3 eV, suggesting the formation of Al2O3 at the surface and, simultaneously, the successful incorporation of Al into the bulk of the P2-Na−Ni−Mn−O samples. Then, the elemental mapping using STEM as shown in Figure 1e displays an inhomogeneous distribution of Al in the sample surface, which indicates the addition of a surface Al coating

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RESULTS AND DISCUSSION Figure 1a,b and Tables S1 and S2 demonstrate the refined powder XRD of the as-prepared samples, where the XRD patterns of P2-Na0.66Ni0.33Mn0.67O2 and 24124


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Al doping mainly works in the near-surface region of the asprepared electrode. We have no direct evidence of its influence on the Na+ ion/vacancy ordering behavior, especially in bulk materials as indicated by the dark SAED patterns in Figure 1c. However, the synergetic contribution of Al surface coating and bulk lattice doping does improve the electrochemical performance of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode with respect to the undoped one. The Coulombic efficiency of the formation cycle of Na0.66Ni0.33Mn0.67O2 is calculated to be 86%,11,17 whereas it is improved to 93.3% after Al doping (see Figure S8 in the Supporting Information (SI)). An obvious decrease of the plateau capacity at 4.3 V is also observed in the Na0.78Al0.05Ni0.33Mn0.60O2 electrode with respect to the undoped one, which implies higher structural stability, especially at high operation voltage after Al doping. This should be correlated with the suppression of the P2−O2 phase transition16 at high desodiation levels. Consequently, the structural irreversibility25 at the high-voltage region should be responsible for the capacity loss in the Na0.78Al0.05Ni0.33Mn0.60O2 electrode. Although the P2-type Na0.78Al0.05Ni0.33Mn0.60O2 electrode shows improved electrochemical performance as shown in Figure 2, obvious capacity fading is clearly seen upon cycling. Hence, in situ XRD was performed to investigate the underlying contribution of Al doping on structural evolution upon cycling. Figure 3a presents the contour plots of the in situ

layer (it is probably an amorphous Al2O3 layer on the basis of surface Al XPS results in Figure 1d) on the electrode material. The inhomogeneous Al distribution results in an accumulation of Al in the surface region (Figures 1e, S3, and S4) in addition to the uniformly distributed Na, Mn, and Ni. In Figure S3, the accumulation of Al can be clearly seen in the particle surface and beyond the surface region; the concentration of Al decreases obviously with the bulk. The above STEM observations indicate the dual effects of Al doping, that is, surface coating and bulk doping (for illustration, see Figure S6). The Al2O3-coated surfaces could probably reduce the side reactions at the electrode−electrolyte interfaces and thus protect the Na0.78Al0.05Ni0.33Mn0.60O2 electrode against manganese dissolution in the electrolyte and the charge-transfer process,39,40 which will be beneficial for long cycling performance. On the other hand, the doped bulk Al ions may help postpone the phase transition, hence stabilizing the structure to facilitate Na+ ion transport, as shown later. Figure 2 shows the electrochemical performance of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode within a voltage range of 2.0−4.5 V versus Na/Na+ at a current density of 12 mA g−1. As shown in Figure 2a, the discharge capacity is 131.9 mAh g−1 with an average discharge potential of 3.41 V in the formation cycle. Figure 2b displays the first 10 CV curves of Na0.78Al0.05Ni0.33Mn0.60O2 at a scan rate of 0.1 mV s−1. It can be observed that a series of cathodic peaks correspond to the visible plateaus in the charge profile, which is a typical electrochemical feature of sodium cathode materials.41 Figure 2c shows the rate performances of Na0.78Al0.05Ni0.33Mn0.60O2 and Na0.66Ni0.33Mn0.67O2 electrodes. As expected, the Alsubstituted sample manifests remarkable better rate capability than the pristine one. Na0.78Al0.05Ni0.33Mn0.60O2 exhibits discharge capacities of 123.9, 115.7, 106, 91.3, 69.4, and 41.2 mAh g−1 at current rates of C/10, C/5, C/2, C, 2C, and 5C (600 mA g−1), respectively, within a voltage range of 2.0−4.5 V. These electrochemical results of the pristine sample in our work are comparable to those of the previously reported Na0.66Ni0.33Mn0.67O2 electrodes for SIBs,7,15 whereas there are only slight differences with the Co-doped Na0.66Ni0.33Mn0.67O2 electrode42 and high rate capability discrepancies in Mg15 or Ti43 substitution for Ni electrodes. Figure 2d further demonstrates the long-term cycling performance of the assynthesized samples at a current density of 0.1 C. As previously reported, the P2-type Na−Ni−Mn−O materials with high nickel contents often suffer from severe capacity fading in longterm cycling.23 However, it improves greatly after a small dose of Al doping, as shown in Figure 2d. The Al-doped sample presents a discharge capacity of 110 mAh g−1 with the capacity retention of 83.9%, whereas it is 52.5% for the pristine sample after 50 cycles. In addition, the discharge voltage decay is also suppressed significantly after Al doping (Figure S7). Throughout the entire 50 cycles, the Coulombic efficiency of the Al-doped material is greater than that of the undoped one. Moreover, for surface coating, the surface Al2O3 thin layer can be employed to decrease the surface catalytic effect of the Na0.67Ni0.33Mn0.67O2 electrode,44 and hence to suppress the side reactions. Furthermore, it also strengthens the surface structure by reducing the surface reconstruction and relaxation,38 which will help weaken the phase transition process upon electrochemical cycles. For Al bulk doping, as a matter of fact, our XPS etching profile only finds the appearance of Al at a nanometer-sized depth of the nearsurface region of the Na0.67Ni0.33Mn0.67O2 electrode, that is, the

Figure 3. In situ XRD of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode. The c o n to u r pl o t s o f t h e i n s i tu X R D pa t t er n s o f ( a ) Na0.78Al0.05Ni0.33Mn0.60O2 during the formation cycle, where the XRD peaks of the Al collector are denoted with “★”. (b) Schematics of the lattice planes of P2-type transition-metal oxides. (c) Interlayer d-spacing evolution during the formation cycle.

XRD patterns of Na 0.78 Al0.05Ni0.33Mn0.60 O2 during the formation cycle. As shown in Figure 3a, both of the (002) and (004) peaks shift significantly to lower angles during the first charge process and return back at the end of discharge, implying the gradual elongation and shrinkage of the interlayer d-spacing (equal to c/2, Figure 3b) of the (002) plane upon desodiation and sodiation. This kind of interlayer distance changes, like the “breathing effect”, is predominately ascribed to the variations of the electrostatic repulsion of the metastable 24125


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Figure 4. Atomic-scale observation of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode upon cycling. The STEM HAADF and ABF images of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode at the [010] zone axis, with (a, b) pristine, (c, d) 0.25 mol, and (e, f) 0.57 mol of Na extraction and (g, h) 0.5 mol of Na reinsertion. (i) O K edge EELS spectra of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode after being discharged to 2.0 V. (j) Measured dspacing of the TMO6 slabs (or the (002) plane) from the STEM images at different Na contents, each of which is determined by averaging the measured d-values of different particles in one specific charging/discharging state (Ch: charge; Dh: discharge).

O2− ions between the TMO6 slabs.7 In the meantime, the shifts of the (100) peak to the higher angles correspond to the shrinkage along the a or b axes, meaning that the TM ions are getting closer with each other. This probably results from the radius reduction of Ni ions during oxidation, which could also refer to the in situ and ex situ X-ray absorption spectroscopy studies on Na0.66Ni0.33Mn0.67O2 materials.8,45,46 Then, similar shifts can be found in either the (012) or (103) peak to show the ionic gliding, mainly the TM ions in the respective intralayer. However, this shift trend disappears for the (104) peak, which seems to be pinned at the same diffraction angle in the whole formation cycle although it vanishes almost at high desodiation contents. After that, there is an eye-catching “turnover” on the interlayer and intralayer distance changes of TMO6 slabs upon desodiation and sodiation due to the shift of the (106) peak to the lower angles again. This phenomenon reflects clearly the physical picture of TM-ion vibration along with the Na+ ion migration. After all, the in situ XRD results imply that there is a severe competition between the Coulombic repulsion and the electrostatic attraction of the TM ions upon the elongation of the c axis, as shown in Figure 3b. First, they get increasingly closer with the obvious red shifts of (100), (101), (102), and (103) peaks to higher angles. Simultaneously, they repel each other due to the tiny blue shift of the (106) peak, which displays vividly the vibration of TM ions around its originally thermodynamic site along with desodiation. The underlying reason for this process is the dynamic charge-transfer process of TM ions upon desodiation, where the loss of electrons from the TM ions could alter the electrostatic interactions, and vice versa in the discharging (Figure 3a). In fact, the observations from the in situ XRD patterns here agree well with the changes of the Ni−O bonds

upon desodiation from the EXAFS spectra of Na0.66Ni0.33Mn0.67O2.8,17 The (002) and (004) peak shifts between 3.4 and 4.0 V during very early desodiation are highlighted in the blue curve region of the electrochemical profile in Figure 3a. This region corresponds to the desodiation contents at ca. 0.09 ≤ x ≤ 0.33, in line with the multiple redox couples from the CV curves in Figure 2b. In Figure 3a, a series of new phases are seized by the in situ XRD to confirm the multiphasic reactions upon desodiation from the intermediate Na0.69Al0.05Ni0.33Mn0.60O2 to the Na0.45Al0.05Ni0.33Mn0.60O2 state, which has not been reported ever before. With the decrease of sodium contents, the changes accompany with (1) the Na+ redistribution in the (002) plane, (2) the Coulombic repulsion between the Naf+TMO6 slab, (3) the chemical valences of cations, and (4) the regulation on the crystal lattice field.41 Consequently, even tiny changes in sodium content here are enough to induce new structural rearrangements during electrochemical cycling. When it is charged to 4.3 V, the in situ XRD patterns present a new peak at ∼18.7°, corresponding to a voltage plateau at the charge/discharge curves, as an indication of the typical phase transformation reaction region in Figures 2a and 3a. Then, the (012), (104), and (106) peaks almost disappear upon sodiation from 4.1 to 3.5 V (Figure 3a), which is similar in Zn-doped17 and Cu-substituted19 samples, whereas these peaks survived in the reported Na0.66Ni0.33Mn0.67O2.17,47 At present, this is ascribed to the particle size and morphology changes and/or to the underlying phase transition from the initial P2 phase to other phases, say, the O2 phase. This is partially consistent with the remarkable decrease in the peak intensity and increase in the full width at half-maximum of the XRD pattern (Figure 3a), showing particle pulverization. Moreover, 24126


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measurements were performed on pristine materials to explore the electronic structures of O and TM ions (Figure S10c). The L edge of TM ions corresponds to the dipole-allowed electron hopping from the metal 2p orbitals to the unoccupied metal 3d orbitals.51 The valence of the TM ion changes with the onset energy variation of the TM L edge or L3/L2 ratios. The EELS line profiles of the pristine materials suggest that the oxidation states of Mn and Ni ions are +4 and +2, respectively, from the surface to the bulk.52 Upon charging to Na0.53Al0.05Ni0.33Mn0.60O2 (Figure 4j, ca. 0.25 mol Na extraction), the planar distance increases somewhat, and there is a gradual increase in the planar dspacing in the desodiation process. In the meantime, obvious stacking faults are found in the surface region, as shown in Figure 4c,d (see raw pictures in Figure S11). These structural changes should be closely correlated with the 23Na MAS NMR spectra of Na0.66Ni0.33Mn0.67O2, where it demonstrates that a new peak appears at ca. 230 ppm to show the migration of Ni ions with a voltage higher than 3.7 V.17 On further charging to the Na0.21Al0.05Ni0.33Mn0.60O2 state (about 0.57 mol Na extraction), a new metastable layered phase was observed at the surface region for the desodiated particle, whereas the bulk region still remained in the P2 phase (Figure 4e,f; see Figure S12 for raw STEM images). This new metastable phase corresponds to the new peak at ∼18.7° of the in situ XRD (Figure 3a). The average d-spacing of the (002) plane is about 4.65 Å in the Na0.21Al0.05Ni0.33Mn0.60O2 phase (Figure 4e), a little smaller than the value of 4.74 Å calculated from the in situ XRD; however, the interplanar d-spacing of the bulk P2 phase is ca. 5.65 Å, a little larger than that of the pristine electrode, which agrees with the shift of (002) and (004) peaks to the lower angles in the in situ XRD. At the same time, the cation migration is also captured in this new phase, as shown in Figure 4f. At this almost fully charged state, the atomicresolution EELS results demonstrate that no obvious peak shifts of the Ni, Mn, and O ions were found, indicating that the chemical valences of these ions become stable after oxidization. The EELS results on the lattice oxygens are different from the one reported by Ma et al.,46 where the lattice oxygen may participate in the charge compensation at the highly desodiated NaxAl0.05Ni0.33Mn0.60O2 electrode. Here, the shift of oxygen K edge seems undetectable to indicate the electrochemical inactivity upon Na extraction, which is probably ascribed to the deactivated effect by the doped Al. A small dose of Al would possibly disturb the short-range arrangements of TM ions and change the surface activity as well, which probably suppresses the electrochemical activity of the lattice oxygen at high voltage. On the other hand, the local chemical environment of AlIIIO6 octahedra could be manipulated by the adjacent coordination, such as the O2−, TM, and Na+ ions, which, in turn, enhances the structural stability and improves the electrochemical performance (Figure 2). Figure 4g,h shows the STEM HAADF and ABF images of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode after being discharged to 2.0 V (ca. 0.5 mol of Na reinsertion). It can be seen that the electrode material basically restores back to the P2 phase. Of special interest is the fact that the cation migration was kept at the bulk region, implying that the migration of TM ions to the Na+ ion layer is possibly irreversible. Although the main XRD peak of the newly formed phase at the high voltage (probably the O2 phase) already disappears, it still survives when discharged to 2.0 V (Figure S13), which might be affected by the migration of TM ions to the Na sites. Figure 4i shows the

in previous studies, a new peak located at 20.2°, corresponding to the (002′) plane, has been reported when the Na0.66Ni0.33Mn0.67O2 electrode was charged to around 4.3 V,7,15,47 where the d-spacing of this new peak was calculated to be 4.38 Å (Figure 3c). However, here this new peak at ∼18.7° corresponds to its interslab distance of 4.74 Å in the desodiated Na0.78Al0.05Ni0.33Mn0.60O2 electrode. Based on these structural changes, the volume changes of the pristine and Al-doped samples are estimated to be 21.8 and 15.4%, respectively, which suggests that a small dose of Al doping efficiently reduces the volume change of the Na0.78Al0.05Ni0.33Mn0.60O2 cathode at high desodiation levels. This may result from the smaller electronegativity of Al and higher ionicity of the M−O bond,48 postponing the P2-to-O2 phase transformation at high voltage. Consequently, a better structural/electrochemical reversibility is shown in the Na0.78Al0.05Ni0.33Mn0.60O2 electrode for SIBs. In addition, when the discharging voltage is lower than 2.5 V, there are two new peaks at 28.3 and 57.27°, as shown in the shadow regions of Figure 3a, which could be the formation of Ni−Mn superstructures.49 This phenomenon should be correlated with the partial reduction of Mn4+ and probably the transpositions among the TM ions, which are evidenced by the red shift of the Mn K edge in the fully discharged sample,8,46 and the subsequent STEM observations. From a structural point of view, the desodiation structures of the layered NaxAl0.05Ni0.33Mn0.60O2 electrode will be driven timely to approach the thermodynamic configurations through relaxation and reconstruction. While the fast structure changes could be significantly modulated by the structural defects, the atomic stacking faults as well as the interior lattice field changes with charge compensation from either the TM ions or the crystal oxygen ions, thus further influencing the electrochemical performance of the Na0.78Al0.05Ni0.33Mn0.60O2 electrodes. To elucidate this intimate structure−activity relationship, here postmortem scanning transmission electron microscopy (STEM) and electron energy loss spectroscopy (EELS) were performed at the atomic scale upon Na+ ion extraction/ insertion from/into Na0.78Al0.05Ni0.33Mn0.60O2 electrodes. In this respect, four different desodiated samplesincluding (1) pristine, (2) charged to 3.7 V (ca. 0.25 mol Na extraction), (3) charged to 4.5 V (ca. 0.57 mol Na extraction), and (4) discharged to 2.0 V (ca. 0.5 mol Na reinsertion)were studied by STEM HAADF and ABF imaging. Figure 4a,b displays the STEM images of the layered Na0.78Al0.05Ni0.33Mn0.60O2 electrode at the [010] zone axis. From STEM images, the atomic arrangements of the pristine sample can be well assigned to the P2 phase with the d-spacing of ca. 5.57 Å (close to 5.55 Å from XRD, as shown in Table S2) separating the transition-metal layers along the [001] direction, which is in good agreement with the simulated STEM images of the P2 phase (Figure S9). As reported, no experimental evidence can be clearly found on the formation of Na−Ni antisite (transposition) in P2-type Na−Ni−Mn−O materials because of the large radius difference between Na+ and Ni ions.50 In Figure S10a,b, the cation migration to the Na site was clearly captured in the near-surface region, which leads to the loss of normal sodium storage sites, partially accounting for the capacity fading of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode. This cation migration results in a relative gliding of the adjacent two transition-metal layers, whereas no evidence is found to support the Na+ ion migration into the TM layers to introduce the Na-TM antisite. Moreover, EELS 24127


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Figure 5. GITT and EIS tests of Na0.66Ni0.33Mn0.67O2 and Na0.78Al0.05Ni0.33Mn0.60O2 electrodes. (a, b) GITT measurements and the derived overpotentials of Na0.66Ni0.33Mn0.67O2 and Na0.78Al0.05Ni0.33Mn0.60O2 in the second cycle. The dotted lines represent the quasi-equilibrium potentials. (c) EIS with the fitting results of the Na0.66Ni0.33Mn0.67O2 and Na0.78Al0.05Ni0.33Mn0.60O2 electrodes at the open-circuit voltage resting for 10 h after being assembled in half-cells. The corresponding equivalent circuit is also shown in the inset. (d) Profiles of Zr vs ω−1/2 from 0.1 to 0.01 Hz. Solid dots are the experimental data, whereas the dotted line represents the fitting results to characterize the Na+ ion diffusion kinetics.

overpotential at the beginning of the first multiphasic reaction region during the charging process. As shown in Figure 3a, there is a complex phase transition process between 3.4 and 4.0 V, where the initial P2-type structure experiences remarkable changes/distortions and acceleration of TM ion migration to the Na layer (Figure 4c,d), contributing to the large overpotential. Compared with the Na0.67Ni0.33Mn0.67O2 electrode (Figure 5a), the electrochemical polarizations of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode are smaller in the whole cycling process, which implies the dual effects of Al doping on electrochemical stabilization (Figure 5b). Another important contribution is the improvement in electronic conductivity after Al doping. Figure S15 demonstrates the measured optical band gap, where that of the Al-doped electrode is apparently smaller than the undoped one, although both samples have good electronic conductivity (∼1.3 eV) comparable with other electrode materials.17 To this end, the transport of the Na ion is thought to be the rate-determining step in SIBs, consistent with the reported results.15 Upon sodium insertion, after going through the electrode− electrolyte interface, the Na+ ions first reach the outmost surface region and then diffuse into the bulk region of the electrode.57 The different diffusion rates of Na+ ions contribute dominantly to electrochemical polarization during this migration. As shown in Figure 5a,b, the maximum points of electrochemical overpotential appear always in the transition region from a slope voltage profile to a plateau upon charging and vice versa upon discharging at the high-voltage region. This means that the transition from the solid solution reaction to the phase separation reaction (even the multiphasic reaction or the P2−O2 transition) is the main factor contributing to the electrochemical overpotential in this Na−Ni−Mn−O material system. At the end of both charging and discharging, the diffusion pathway elongation of Na+ ions from the surface to the bulk region and the phase boundary movement observed from in situ XRD lead to the sluggish diffusion kinetics, which could produce large electrochemical polarizations.

spatially resolved O K edge EELS spectra of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode after being discharged to 2.0 V, where two peaks (a, b) can be identified clearly. The former peaks, labeled a, result from the electron hopping from the 1s core states to the oxygen 2p and TM 3d hybridized states.53 The latter peaks, labeled b, are caused by the electron hopping to the hybridized states of TM 4s and 4p with oxygen 2p orbitals.54 Both peaks of a and b are important indications of the local structural changes upon cycling.55 It is observed that the peaks (a, b) are split at the region when the TM ion migrates into the Na+ ion layer (see TM L edge EELS spectra in Figure S14), indicating the local distortion of TMO6 octahedra. This corresponds to the cation migration that could be accelerated by the appearance of Na vacancies and the variation of the TM−O bonds as mentioned above. The above structural evolution observations suggest that (1) the P2−O2 phase transition happens in the electrochemical cycle of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode; (2) the O2−P2 phase transition is not fully reversible, which is probably affected by the cation migration to the Na+ ion layer; and (3) the phase transition probably experiences a back-and-forth process and could be deteriorated by the accumulation of structural irreversible changes upon Na+ ion exchange. The above observation may explain the slow capacity fading of the Al-doped materials after 50 cycles. The effect of inactive elemental Al helps retard the electroactive P2 phase irreversibly transforming into the electroinactive O2 phase. The galvanostatic intermittent titration technique (GITT)56 and electrochemical impedance spectroscopy (EIS) were further performed to elucidate the Na+ diffusion kinetics of both Na0.66Ni0.33Mn0.67O2 and Na0.78Al0.05Ni0.33Mn0.60O2 electrodes (see the Supporting Information for experimental details). Due to the significant structural reformation and side reactions in the formation cycle, Figure 5a,b shows the GITT results and the derived overpotentials of both electrodes in the second cycle. It can be seen that both electrodes demonstrate similar maximum points of electrochemical 24128



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CONCLUSIONS In summary, the highly crystallized P2-type Na0.78Al0.05Ni0.33Mn0.60O2 cathode material was successfully synthesized by a sol−gel method. The Na0.78Al0.05Ni0.33Mn0.60O2 electrode exhibited a much higher capacity retention of 83.9%, as compared with 52.5% for the Na0.66Ni0.33Mn0.67O2 electrodes after 50 cycles in SIBs. Moreover, remarkable improvements were found in the suppression of voltage decay and rate capabilities upon cycling after Al doping. The structural analysis from etching XPS and energy-dispersive X-ray mapping indicates that Al plays dual roles in both surface coating and bulk doping. The influences on electrochemistry of Al doping were disclosed using both in situ XRD and ex situ STEM. The results show that the accumulation of the irreversible structural changes from the original P2 to the electrochemically less active O2 phase and the volume changes at high voltage were significantly suppressed. Of special interest was the cation migration upon cycling, which may intrigue the structural degradation to contribute to capacity fading. The STEM-EELS characterizations on the highly desodiated samples showed that the electrochemical activity of the lattice oxygen was probably deactivated by Al doping. Further studies exhibited that the Na + ion diffusion kinetic was greatly enhanced after introducing this inactive Al using the GIIT and EIS. The dual effects of Al doping promise good electrochemical performance of the Na0.78Al0.05Ni0.33Mn0.60O2 cathode material. These findings on the Na+ ion storage and transport mechanism deepen the comprehension on the electrode process dynamic and battery communities as well. This work provides new insights on the design and optimization on P2 Na−Ni−Mn−O materials and the development of high-energy density sodium ion batteries.

As shown in Figure 5b, the electrochemical overpotentials of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode change with the sodium insertion/extraction (Figure 5b). The underlying reason is that the Na+ ion diffusion kinetic changes with Na contents. From the GITT characterizations, the diffusion coefficient56 can be obtained by presuming sodium ion transport in the electrode obeying Fick’s second law. The calculated DNa+ is shown in Figure S16. It can be seen that throughout the whole charging and discharging processes, both electrodes show similar variations although the Na0.78Al0.05Ni0.33Mn0.60O2 electrode demonstrates an overall higher diffusion coefficient. Upon discharging, the derived diffusion coefficient increases to a value higher than 10−10 S cm−2 with the discharge capacity, whereas during desodiation, it decreases to about 10−14 S cm−2, approx. 4 orders of magnitude lower than the diffusion coefficient at the beginning of charge. Such an obvious conversion of the diffusion coefficient first originated from the available Na+ ion and the vacancy and also their interactions (distributions and rearrangements) during electrochemical cycling. Another important contribution comes from the phase transition that seems to hinder the diffusion of Na ions, consistent with the previous report.58 Of special note is that the charge capacity of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode is 110 mAh g−1 before the electrochemical polarization reaches the maximum point, much higher than the 95.7 mAh g−1 of the undoped one, as shown in Figure 5a,b. This, again, supports the above conclusion that doping Al in the Na0.78Al0.05Ni0.33Mn0.60O2 electrode retards the unfavorable structural transition and improves the electrochemical performance. The EIS measurements are performed on both of the uncycled electrodes in coin cells. As shown in Figure 5c, three major parts are found in the EIS plots, that is, two semicircles at high and middle frequencies and a slope line at low frequency (0.1 MHz to 0.01 Hz). The semicircles at high and middle frequencies represent the resistances from the cathode−electrolyte interphasic film (Ri) and the charge-transfer process (Rct),47 respectively. The slope line at low frequency corresponds to the Warburg impedance (W). By fitting to the measured EIS results, the respective resistances (Figure 5c) are derived as shown in Table S4. It can be seen that the main contribution after Al doping is the significant decrease of the charge-transfer resistance, where it reduces from 1702.1 Ω in the undoped electrode to 829.5 Ω in the Na0.78Al0.05Ni0.33Mn0.60O2 electrode. This should be partially responsible for the superior electrochemical performance, as shown in Figure 2d. Moreover, the diffusion coefficient DNa+ could be estimated by fitting to the Warburg impedance part in the low frequencies. The most important point is the calculation of the Warburg coefficient, which is obtained from the slope of Z′ versus ω−1/2, as shown in Figure 5d. As a result, the obtained Na+ ion diffusion coefficient is ca. 1.32 × 10−13 in the Al-doped electrode, almost 1 order of magnitude higher with respect to that in the Na0.66Ni0.33Mn0.67O2 electrode. The Na+ ion diffusion coefficient here is close to the one at the beginning of charging, whereas it is smaller than that during desodiation from the GITT method in Figure S16. This results from the increase of the Na vacancy to enhance the mobility of the residue Na+ ions as aforementioned. In short, the experimental GITT and EIS results demonstrate a clear consistence of the Na + ion diffusion kinetics in the Na0.78Al0.05Ni0.33Mn0.60O2 electrode with better electrochemical performance, as shown in Figure 2.

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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/acsami.9b06233. Calculation methods, refinement results, EIS, SEM, high-resolution TEM and mapping results, STEM and EELS results, UV−vis absorption, and GITT results on Na ion diffusion kinetics (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (Q.Z.). *E-mail: [email protected] (X.L.). *E-mail: [email protected] (D.C.). ORCID

Xia Lu: 0000-0003-3504-9069 Lin Gu: 0000-0002-7504-031X Dapeng Cao: 0000-0002-6981-7794 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by National Natural Science Foundation of China (Distinguished Young Scholars Program No. 21625601 and General Project No. 11704019), Outstanding Talent Fund from BUCT, and The Hundreds of Talents program of Sun Yat-sen University. 24129


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Probing the Structural Transition Kinetics and Charge Compensation

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