Probing the Structural Transition Kinetics and Charge Compensation

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Probing the Structural Transition Kinetics and Charge Compensation of 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 ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b06233 • Publication Date (Web): 12 Jun 2019 Downloaded from http://pubs.acs.org on June 12, 2019

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ACS Applied Materials & Interfaces

Probing the Structural Transition Kinetics and Charge Compensation of P2Na0.78Al0.05Ni0.33Mn0.60O2 Cathode for Sodium Ion Batteries Yuansheng Shia, c, Shuai Lia, Ang Gaoa, Jieyun Zhengb, Qinghua Zhangb,*, Xia Lua, c, *, Lin Gub, and Dapeng Caoa, * a. 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. b. Institute of Physics, Chinese Academy of Sciences, Beijing 100190, P. R. China. c. School of Materials, Sun Yat-sen University, Guangzhou 510275, P. R. China. E-mails: [email protected]; [email protected]; [email protected] 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 small dose of Al replacing the Mn, aiming at enhancing the structural stability without sacrificing the average discharge voltage and theoretical capacity. The etching XPS and EDX mapping/line scan results indicate that the Al doping induces the 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 - 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; transmission electron microscopy, Na ion batteries

1 ACS Paragon Plus Environment

In-situ

XRD,

Scanning

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Introduction Compared with the popular lithium ion batteries (LIBs), the sodium ions batteries (SIBs) with the similar “rocking chair” mechanism,1 are found to be one of the good alternatives due to its low cost and wide applications in large-scale energy storage systems.2 However, developing suitable cathode material is one of the most important steps to achieve the commercialization of SIBs. Among the available candidates, the sodium layered-type oxides (NaxTMO2, 0 < x < 1 and TM: transition metals) have attracted increasingly more attention due to their high theoretical capacity, the moderate operation voltage, the good structural stability and the easy availability, which show the promising prospects to promote the availability of SIBs in a 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 [001] direction. Previous investigations indicate that the P2-type cathode materials have more stable Na+ ion diffusion trajectory and the larger interlayer spacing due to the prismatic occupation of Na ions among the MO6 slabs, and therefore exhibit the superior electrochemical performance from both the experimental results and the theoretical simulations,4-6 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 the Na ions in the layered P2-Na2/3Ni1/3Mn2/3O2 can be intercalated and de-intercalated reversibly based on the Ni2+/Ni4+ redox couples. It can deliver a capacity as high as 160 mA h g-1 with an average discharge voltage of ca. 2 ACS Paragon Plus Environment

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3.6 V vs. Na+/Na.8 Moreover, the P2-Na2/3Ni1/3Mn2/3O2 cathode exhibits the enhanced air stability9 and the cost–effectiveness with respect to the other types of P2- and O3materials.10 However, this material suffers the fast capacity fading and discharge voltage decay, which is probably resulted from the P2 – O2 phase transition at the voltage of ca. 4.3 V vs. Na+/Na and the large volume change upon the overdesodiation. In consequence, many strategies were proposed to overcome these drawbacks of P2- Na2/3Ni1/3Mn2/3O2 electrode materials. The first choice is the 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 the elemental doping, i.e. using the low valence metal ions (Li+,13-14 Mg2+,15-16 Zn2+,17 Ca2+,18 Cu2+,19 etc.) and multimetal ions doping20-22 to increase the valence state of manganese ions and to reduce the contents of nickel as well. Of special interest is 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 synergize together, better electrochemical performance would be expected for P2- Na2/3Ni1/3Mn2/3O2 electrode. To date, the underlying mechanism of Al doping into P2 type-Na-Ni-Mn-O material is still unclear, although the Al-doped materials with improved electrochemical performance have been reported in Na-Ni-Mn-O24-25 and Na-Co-MnO systems.26 Previous reports disclosed that there are several known factors being 3 ACS Paragon Plus Environment


responsible for the battery capacity and voltage fading, including the electrolyte decomposition,27 the Jahn–Teller distortion28-29 and the transition metal migration, the intergranular cracking30, the dissolution of divalent manganese,12 the instability of lattice oxygen,31-34 and the structural rearrangement.7 Nonetheless, it is still difficult to elucidate the effect of Al doping on the structural/electrochemical stabilities of 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 cathode material for SIBs. The dual effects of Al doping are identified by etching XPS and elemental mapping. To analyze the capacity fading mechanism of Al-doped materials, the detailed structural evolution is monitored using the in-situ X-ray diffraction (XRD), the ex-situ spherical aberration-corrected scanning transmission electron microscopy techniques (STEM) and electron energy loss spectroscopy (EELS). The results indicate that the Al doping induces the 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, i.e. the accumulation of structural irreversible changes from P2 to O2 phase and the volume changes at high voltage after Al doping.

Experimental 2.1 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, the 4 ACS Paragon Plus Environment

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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 DI 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 ℃ for 6 hours. Then it is ground in an agate mortar for one hour and heated in a muffle furnace at 950 ℃ for 15 h in air and slowly cooled down to room temperature to obtain the targeted samples. 2.2 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 an 2θ range of 10° to 120°. The XRD pattern was refined using the 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 an X-ray photoelectron spectroscopy (XPS, ThermoFisher ESCALAB 250 X-ray 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. The scanning electron microscopy (SEM) and transmission electron microscopy (TEM) were carried out with a JEOL JSM-6360 L V (Japan) and JEM-2100F (operating at 200 kV), respectively. The ultravioletvisible (UV) light absorption spectrum was tested on an Agilent Cary 5000 UV 5 ACS Paragon Plus Environment


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spectrophotometer to get the exact optic band gap of the as-prepared samples. The insitu XRD diffraction were measured at the Bruker D8 phaser diffractometer with the home-made in-situ cell. The current density was set to be 10 mA g-1 during cycling. Structural evolution at atomic scale and electron energy loss spectroscopy (EELS) were carried out by a JEM ARM200F (JEOL, Tokyo, Japan), equipped with a double CEOS (Heidelberg, Germany) probe aberration correctors. The high-angle annular dark-field (HAADF STEM) and low-angle bright-field STEM (LABF STEM) were performed at 200 kV. The acquired raw data was processed with the Mutivariate Statistical Analysis V4.3 and ABSF filters. 2.3 Electrochemical measurements. The working electrode was fabricated by casting the mixture of the active material, 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 ℃ 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 < 0.1ppm), the electrode discs need to be dried at 100 ℃ for 10 h in the vacuum oven. The cycling tests of the working electrodes were performed in the CR2032 coin cells. The sodium metal and the glass microfiber (Whatman, GF/D) were used as a reference electrode and separator, respectively. The electrolyte was 1 M NaPF6 in a solution of propylene carbonate (PC) with 5% fluoroethylene carbonate (FEC). Galvanostatic cycling tests were conducted by the Land battery testing system (LAND CT-2001A Instrument, Wuhan, China) after resting for 10 h. The cyclic 6 ACS Paragon Plus Environment

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voltammetry (CV) and the 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 AC 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 sodium ion with a current density of 25 μA for 1.0 hours then followed by a 5 hours relaxation to reach the probably electrochemical equilibrium states. All these electrochemical measurements were carried out at room temperature.

Results and Discussion Figure 1a, b and Tables S1, S2 demonstrate the refined powder XRD of the asprepared samples, where the XRD patterns of P2-Na0.66Ni0.33Mn0.67O2 and 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. Compare with the high-Ni O3-type layered cathode materials,36 the P2-Na0.78Al0.05Ni0.33Mn0.60O2 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 one month. (Figure S1) The phase identification is further characterized by the selected area electron diffraction (SEAD) as shown in Figure 1c, where the bright diffraction dots correspond to the P2 phase exactly. Moreover, the 7 ACS Paragon Plus Environment


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 investigations. 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.

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. Experimental (red circles) and calculated (black solid line), Bragg reflection peaks (green solid ticks) and the difference curve (blue line) are shown, respectively. (c) TEM-SEAD, (d) Etching XPS profiles, (e) elemental mapping results of Na0.78Al0.05Ni0.33Mn0.60O2 sample.

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, i.e. the aluminum oxide,37 while the other one at ∼73.3 eV is in good


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agreement with the profile of LiCo0.85Al0.15O2.38 With the increase of the 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 that an inhomogeneous distribution of Al in the sample surface, which indicates the addition surface Al coating layer (it is probably an amorphous Al2O3 this 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 as shown in Figure 1e, Figure S3, and Figure 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 to the bulk. The above STEM observations indicate the dual effects of the Al doping, i.e., the surface coating and the bulk doping (The illustration sees 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 the manganese dissolution in the electrolyte and the charge-transfer process,39-40 which will be beneficial for the long cycling performance. On the other hand, the doped bulk Al ions may help on postponing the phase transition, hence stabilizing the structure to facilitate the Na+ ions transport as shown later. Figure 2 shows the electrochemical performance of Na0.78Al0.05Ni0.33Mn0.60O2 electrode within a voltage range of 2.0 - 4.5 V vs. 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 9 ACS Paragon Plus Environment


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 performance of Na0.78Al0.05Ni0.33Mn0.60O2 and Na0.66Ni0.33Mn0.67O2 electrodes. As expected, the Al-substituted sample manifests remarkable better rate capability than the pristine one. The Na0.78Al0.05Ni0.33Mn0.60O2 exhibits the discharge capacities of 123.9, 115.7, 106, 91.3, 69.4, and 41.2 mAh g-1 at the current rates of C/10, C/5, C/2, C, 2C and 5C (600 mA g-1) within a voltage range of 2.0 and 4.5 V. These electrochemical results of pristine sample in our work are comparable with the previously reported Na0.66Ni0.33Mn0.67O2 electrodes for SIBs,7, 15 while there are only slight differences with the Co-doped Na0.66Ni0.33Mn0.67O2 electrode,42 and the high rate capability discrepancies in the Mg15 or Ti43 substitution for Ni electrodes. Figure 2d further demonstrates the long-term cycling performance of the as-synthesized samples at a current density of 0.1 C. As previously reported, the P2-type Na−Ni−Mn−O materials with the high nickel contents often suffer severe capacity fading in long cycling.23 However, it gets a great improvement 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%, while it is 52.5% for the pristine sample after 50 cycles. In addition, the discharge voltage decay also is suppressed significantly after Al doping. (Figure S7) Throughout the entire fifty cycles, the coulombic efficiency of Al-doped material is greater than the undoped one. 10 ACS Paragon Plus Environment

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Moreover, For the surface coating, the surface Al2O3 thin layer can be employed to decrease the surface catalytic effect of Na0.67Ni0.33Mn0.67O2 electrode,44 hence to suppress the side reactions; Furthermore, it also strengthens the surface structure by reducing the surface reconstruction and relaxation,38 which will help on weakening the phase transition process upon electrochemical cycles. For the Al bulk doing, as a matter of fact, our XPS etching profile only finds the appearance of Al in a nanometer-sized depth of the near surface region of Na0.67Ni0.33Mn0.67O2 electrodethat is-the Al doping mainly works in the near surface region of the as-prepared electrode. We have no direct evidence on its influence on the Na+ ion/vacancy ordering behavior, especially in the 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

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 while it is improved to 93.3% after Al doping (refers to Figure S8 in 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 the 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 the high desodiation levels. Consequently, the structural irreversibility25 at high voltage region should be responsible for the capacity loss in Na0.78Al0.05Ni0.33Mn0.60O2 electrode. 11 ACS Paragon Plus Environment


Figure 2. Electrochemical performance of Na0.78Al0.05Ni0.33Mn0.60O2 electrode. (a) Charge/ discharge curves of Na0.78Al0.05Ni0.33Mn0.60O2 electrode between 2.0 and 4.5V at C/10; (b) The CV curves of Na0.78Al0.05Ni0.33Mn0.60O2 electrode at 0.1 mV s-1. (c) Rate performance and (d) cycling performance of Na0.78Al0.05Ni0.33Mn0.60O2 electrode, where the electrochemical performance of Na0.66Ni0.33Mn0.67O2 electrode is shown for comparison as well.

Although the P2-type Na0.78Al0.05Ni0.33Mn0.60O2 electrode shows the improved electrochemical performance as shown in Figure 2, the obvious capacity fading is clearly seen upon cycling. Hence, the 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 XRD patterns of Na0.78Al0.05Ni0.33Mn0.60O2 during the formation cycle. As shown in Figure 3a, both of the (002) and (004) peaks shift significantly to the 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 plane (002) upon desodiation and


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sodiation. This kind of interlayer distance changes, like the “breathing effect”, are predominately ascribed to the variations of the electrostatic repulsion of the metastable O2- ions between the TMO6 slabs.7 In the meantime, the shifts of (100) peak to the higher angles correspond to the shrinkage along the a- or b- axes, meaning the TM ions getting closer with each other. This is probably resulted from the radius reduction of Ni ions during oxidation, which could also refer to the in-situ and ex-situ X-ray absorption spectroscopy (XAS) studies on Na0.66Ni0.33Mn0.67O2 materials.8, 45-46 Then similar shift can be found either in (012) or in (103) peaks to show the ionic gliding, mainly the TM ions in the respective intralayer. However, this shift trend disappears for (104) peak, which seems to be pinned at the same diffraction angle in the whole formation cycle although it vanishes almost at the high desodiation contents. After that, there is an eye-catching “turn-over” on the interlayer and intralayer distance changes of TMO6 slabs upon desodiation and sodiation due to the shift of (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 closer and closer as the obvious red shifts of (100), (101), (102) and (103) peaks to the higher angles. Simultaneously, they repel with each other due to the tiny blue shift of (106) peak, which displays vividly the vibration of TM ions around its originally thermodynamic site along with the desodiation. The underlying reason for this process is the dynamic charge transfer 13 ACS Paragon Plus Environment


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 of 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) peaks shift between 3.4 and 4.0 V at the very early desodiation were highlighted in the blue curve region of electrochemical profile (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 multi-phasic reactions upon the desodiation from the intermediate Na0.69Al0.05Ni0.33Mn0.60O2 to the Na0.45Al0.05Ni0.33Mn0.60O2 states, which has not been reported ever before. With the decrease of the sodium contents, the changes accompany with (1) the Na+ redistribution in (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 the new structural rearrangements during electrochemical cycling.


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Figure 3. In-situ XRD of Na0.78Al0.05Ni0.33Mn0.60O2 electrode. The contour plots of the in-situ XRD patterns of (a) Na0.78Al0.05Ni0.33Mn0.60O2 during the formation cycle, where the XRD peaks of Al collector are denoted with ‘★’. (b) Schematics of the lattice planes of P2 type transition metal oxides. (c) The interlayer d-spacing evolution during the formation cycle.

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 disappear almost upon sodiation from 4.1 to 3.5 V (Figure 3a), which is a similar case in Zn-doped 17 and Cu substituted19 samples, while these peaks survived in the reported Na0.66Ni0.33Mn0.67O2.17, 47 At present, the 15 ACS Paragon Plus Environment


reasonable explanations are ascribed to the particle size and morphology changes, and/or the underlying phase transition from the initial P2 phase to other phases, say, O2 phase. This is partially in consistence with the remarkable decrease on the peak intensity and increase on the full width at half maximum (FWHM) of the XRD pattern (Figure 3a) to show the particle pulverization. Moreover, in previous investigations, 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 dspacing 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 the 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 the high desodiation levels. This may be resulted 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 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 the 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 16 ACS Paragon Plus Environment

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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 the 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 the charge compensation from either the TM ions or the crystal oxygen ions, thus further influence the electrochemical performance of the Na0.78Al0.05Ni0.33Mn0.60O2 electrodes. In order to elucidate this intimate structure-activity relationship, here the postmortem scanning transmission electron microscopy (STEM) and electron energy loss spectroscopy

(EELS)

were

performed

at

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 of Na extraction), (3) charged to 4.5 V (ca. 0.57 mol of Na extraction), (4) discharged to 2.0 V (ca. 0.5 mol of Na reinsertion) were studied by STEM HAADF and ABF imaging.



Figure 4．Atomic-scale observation of Na0.78Al0.05Ni0.33Mn0.60O2 electrode upon cycling. The STEM HAADF and ABF images of Na0.78Al0.05Ni0.33Mn0.60O2 electrode at [010] zone axis, where (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) The O K-edge EELS spectra of the Na0.78Al0.05Ni0.33Mn0.60O2 electrode after being discharged to 2.0 V. (j) The measured d-spacing of the TMO6 slabs (or plane (002)) from the STEM images at different Na contents, each of which is determined by averaging the measured dvalues of different particles in one specific charging/discharging state (Ch: Charge; Dh: Discharge).

Figure 4a, b displays the STEM images of layered Na0.78Al0.05Ni0.33Mn0.60O2 electrode at [010] zone axis. From STEM images, the atomic arrangements of pristine sample can be well assigned to the P2 phase with the d-spacing of ca. 5.57 Å (close to the 5.55 Å from XRD as shown in Table S2) separating the transition metal layers along [001] direction, which is in good agreement with the simulated STEM images 18 ACS Paragon Plus Environment

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of P2 phase (Figure S9). As reported, no experimental evidences are 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, while no evidence is found to support the Na+ ion migration into the TM layers to introduce the Na-TM antisite. Moreover, the EELS measurements were performed on pristine material to explore the electronic structures of O and TM ions (Figure S10c). The Ledge 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 TM ion changes with the onset energy variation of TM L-edge or L3/L2 ratios. The EELS line profiles of the pristine materials suggest that the oxidation states of Mn, Ni ions are +4, +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 d-spacing in the desodiation process. In the meantime, the 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 19 ACS Paragon Plus Environment


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, while the bulk region was still kept in 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 (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; while 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 atomic-resolution 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 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 are 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 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-, 20 ACS Paragon Plus Environment

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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 that the cation migration was kept at the bulk region, implying that the migration of TM ions to 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 spatially resolved O K-edge EELS spectra of 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 as a, are resulted from the electron hopping from the 1s core states to the oxygen 2p and TM 3d hybridized states.53 The latter peaks, labeled as 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 to the locally structural changes upon cycling.55 It is observed that the peaks (a, b) are split at the region when the TM ion migration into the Na+ ion layer (see TM L-edge EELS spectra in Figure S14), which means 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 to O2 phase transition happens in the electrochemical cycle of the 21 ACS Paragon Plus Environment


Na0.78Al0.05Ni0.33Mn0.60O2 electrode; (2) the O2 to P2 phase transition is not fully reversible, which is probably affected by the cation migration to the Na+ ion layer; (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 to retard the electro-active P2-phase irreversibly transforming into the electro-inactive 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 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 at the second cycle. It can be seen that both electrodes demonstrate the similar maximum points of electrochemical overpotential at the beginning of the first multi-phasic 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 the remarkable changes/distortions and the acceleration of TM ion migration to the Na layer (Figure 4c, d) to contribute the large overpotential here. Compared to the Na0.67Ni0.33Mn0.67O2 electrode (Figure 5a), the electrochemical polarizations of Na0.78Al0.05Ni0.33Mn0.60O2 electrode are smaller in the whole cycling process, which implies the dual effects of Al doping on the electrochemical stabilization (Figure 5b). Another important 22 ACS Paragon Plus Environment

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contribution is the improvement on the electronic conductivity after Al doping. Figure S15 demonstrate 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 the other electrode materials.17 To this end, the transport of Na ion is thought to be the rate determining step in SIBs, in consistence with the reported results.15 Upon the sodium insertion, after going through the electrode-electrolyte interface, the Na+ ions firstly reach the outmost surface region, then diffuse into the bulk region of the electrode.57 The different diffusion rates of Na+ ions contribute dominantly to the 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 high voltage region. This means that the transition from the solid solution reaction to the phase separation reaction (even the multi-phasic reaction or the P2-O2 transition) is a main contribution to the electrochemical overpotential in this Na-Ni-Mn-O material system. At the ends 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 the large electrochemical polarizations. As

shown

in

Figure

5b,

the

electrochemical

overpotentials

of

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 23 ACS Paragon Plus Environment


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the Na contents. From the GITT characterizations, the diffusion coefficient56 can be obtained by presuming sodium ion transport in the electrode obeying the Fick’s second law. The calculated

D Na +

is showed in Figure S16. It can be seen that

throughout the whole charging and discharging processes, both electrodes show the similar variations although the Na0.78Al0.05Ni0.33Mn0.60O2 electrode demonstrates the overall higher diffusion coefficient. Upon discharging, the derived diffusion coefficient increases to the higher value than 10-10 S/cm2 with the discharge capacity, while during desodiation, it decreases to about 10-14 S/cm2, approx. four orders of magnitude lower than the diffusion coefficient at the beginning of charge. Such an obvious conversion of the diffusion coefficient is firstly 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 hindering the diffusion of Na ions, in consistence with the previous

report.58

Of

special

notice

is

that

the

charge

capacity

of

Na0.78Al0.05Ni0.33Mn0.60O2 electrode is 110 mAh g-1 before the electrochemical polarization reaching 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 semi-circles at high and middle frequencies and a slope line at the low frequency (0.1 MHz - 0.01 Hz). The 24 ACS Paragon Plus Environment

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semi-circles 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.

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 at the 2nd 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 (OCV) resting for 10 h after being assembled in half-cells. The corresponding equivalent circuit is also shown inset. (d) Profiles of the Zr vs. ω-1/2 from 0.1 to 0.01 Hz. Solid dots are the experimental data, while the dotted line represent the fitting results to characterize the Na+ ion diffusion kinetics.

It can be seen that the main contribution after Al doping is the significant decrease of the charge transfer resistance, where it reduces from the 1702.1 Ω in the 25 ACS Paragon Plus Environment


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undoped electrode to the 829.5 Ω of 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

D Na +

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’ vs. 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 one order of magnitude higher with respect to that in Na0.66Ni0.33Mn0.67O2 electrode. The Na+ ion diffusion coefficient here is close to the one at the beginning of charging, while it is smaller than that during desodiation from the GITT method in Figure S16. This is resulted 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 Na0.78Al0.05Ni0.33Mn0.60O2 electrode with the better electrochemical performance as shown in Figure 2.

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 to 52.5% for the Na0.66Ni0.33Mn0.67O2 electrodes after 50 cycles in SIBs. Moreover, remarkable improvements were found on the suppression of voltage decay and rate capabilities upon cycling after Al doping. The structural analysis from the etching XPS and EDX mapping indicates that the Al plays the dual 26 ACS Paragon Plus Environment

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roles in both of the surface coating and the bulk doping. The influences on electrochemistry of Al doping were disclosed using both the in-situ XRD and the exsitu STEM. The results show that the accumulation of the irreversibly 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 the capacity fading. The STEM-EELS characterizations on the highly desodiated samples shown that that the electrochemical activity of the lattice oxygen was probably deactivated by the Al doping. Further studies exhibited that the Na+ ion diffusion kinetic was greatly enhanced either after introducing this inactive Al using the GIIT and EIS. The dual effects of Al-doping promise the good electrochemical performance of Na0.78Al0.05Ni0.33Mn0.60O2 cathode material. These findings on Na+ ion storage and transport mechanism deepens the comprehension on the electrode process dynamic and battery communities as well. This work provided the new insight of the design and optimization on P2 Na-Ni-MnO materials and the development of high-energy density sodium ion batteries.

ASSOCIATED CONTENT Supporting Information: Calculation methods, Refinement results, EIS, SEM, HRTEM and Mapping results, STEM and EELS results, Uv-Vis. absorption and GITT results on the Na ion diffusion kinetics.

This material is available free of charge via the Internet at http://pubs.acs.org.



Acknowledgements This work is 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.

Conflict of Interest The authors declare no conflict of interest.

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