Origin of Structural Evolution in Capacity Degradation for Overcharged

Jul 5, 2017 - In this work, the quasi-spherical Ni-rich layered LiNi0.6Co0.2Mn0.2O2 (QS-NMC622) material was successfully synthesized through the carb...
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Origin of Structural Evolution in Capacity Degradation for Overcharged NMC622 via Operando Coupled Investigation Qi Wang,†,§ Chong-Heng Shen,†,§ Shou-Yu Shen,† Yue-Feng Xu,† Chen-Guang Shi,† Ling Huang,*,† Jun-Tao Li,‡ and Shi-Gang Sun*,†,‡ †

Department of Chemistry, College of Chemistry and Chemical Engineering, State Key Laboratory of Physical Chemistry of Solid Surfaces, Xiamen University, Xiamen, 361005, China ‡ College of Energy, Xiamen University, Xiamen, 361005, China S Supporting Information *

ABSTRACT: The nickel-rich layered oxide materials have been selected as promising cathode materials for the next generation lithium ion batteries because of their large capacity and comparably high operating voltage. However, at high voltage (beyond 4.30 V vs Li/Li+), the members of this family are all suffering from a rapid capacity decay, which was commonly concerned with crystal lattice distortion and related cation disordering. In this work, the quasi-spherical Ni-rich layered LiNi0.6Co0.2Mn0.2O2 (QS-NMC622) material was successfully synthesized through the carbonate co-precipitation method. A coupled measurement, which is a combination of potentiostatic intermittent titration technique (PITT) and in situ X-ray diffraction (XRD), was deployed to simultaneously capture the structural changes and lithium ion diffusion coefficient of QS-NMC622 material during the first cycle. With help of in situ XRD patterns and high-resolution transmission electron microscope (HR-TEM) images, a defective spinel framework of Fd3̅m space group was detected along with a rapid decreasing lattice-parameter c and lattice distortion at deep delithiated state, which causes poor kinetics related to lithium ion mobility. The new-born framework seems to transform and remain as full spinel structure in the parent phase to the end of charge/discharge with high voltage, which could deteriorate both the surface and body structure stability during the subsequent cycles. This established coupled in situ measurement could be applied to simultaneously investigate the structure transformation and kinetics of cathode materials during charge/discharge. KEYWORDS: Ni-rich layered cathode material, in situ XRD, PITT, Li-ion diffusion, spinel framework

1. INTRODUCTION In the past decades, lithium-ion batteries (LIBs) have become the primary energy sources for portable electronic devices on account of their high capacity, long cycle life, and high safety factor.1−4 The layered LiCoO2 (LCO) has been widely used as commercial cathode materials of small LIBs for consumer electronics because of its high ionic conductivity and stable working voltage. However, the deliverance of no more than 150 mAh·g−1 at discharge for LCO material fails to meet demand for LIBs of high voltage transportation and stationary energy storage. To improve performance of the LCO, lots of efforts have been conducted.5,6 Recently, layered oxide LiNixCoyMn(1−x−y)O2 (LNCMO) materials have shown many attractive advantages such as high reversible capacity, relatively low cost, and low toxicity. Among the chemical components, the nickel ions play a role of providing the electron for redox reaction so as to achieve high capacity. Cobalt ions can maintain layered structure and increase electronic conductivity, resulting in an enhanced lithium storage performance at a high current density. Moreover, manganese ions can provide a stable skeleton to improve the stability and safety of the batteries.7,8 Therefore, © 2017 American Chemical Society

the electrochemical properties of LNCMO materials are greatly affected by the proportion of the metal ions (Ni, Co, and Mn). Nowadays, layered oxide LiNi1/3Co1/3Mn1/3O2 material has been used in commercial LIBs, but its capacity is still far from satisfaction of the practical needs.9 It has been reported that the Ni-rich layered oxide materials (LiNixCoyMn(1−x−y)O2 (x ≥ 0.5)) can deliver high specific capacity due to the high content of Ni.1,10 However, its poor cycle performance and structural instability at higher voltage strongly limited their commercialization. Considerable research work have been carried out to synthesize the Ni-rich layered (NRL) oxide materials with highly stabilized structures, such as hydroxide/carbonate coprecipitation,3,5,11−14 solid-state reaction,15 sol−gel,16,17 and combustion18 method. Among them, the co-precipitation route has been extensively used as the preparation of the Ni-rich layered metal oxide materials. Received: May 5, 2017 Accepted: July 5, 2017 Published: July 5, 2017 24731

DOI: 10.1021/acsami.7b06326 ACS Appl. Mater. Interfaces 2017, 9, 24731−24742

Research Article

ACS Applied Materials & Interfaces

2. EXPERIMENTAL SECTION

In order to understand the capacity fade of Ni-rich layered materials at a relatively higher voltage, it is vital to investigate the structural transformation of the cathode material during the charge/discharge processes. By means of in situ X-ray diffraction (XRD) experiment,19,20 Su et al. investigated the irreversibility of structural changes of LiNi0.5Co0.3Mn0.2O2 material in a higher working voltage region of 2.0−4.9 V.10 In situ XRD patterns show the formation of hexagonal phase H3 at around 4.7 V in the charge process. With aid of ex-TEM experiment, Kang et al. demonstrated that the LiNi0.5Co0.3Mn0.2O2 material was subjected to an irreversible phase transformation from rhombohedral phase to spinel and rock salt phase in high voltage region.21 Besides structural evolution, surface chemical instability is another well-acknowledged failure factor for NRL materials at high voltage. Contacting with electrolyte, NRL materials could have side reactions happening on the surface. Moreover, the byproducts of electrolyte decomposition can further contaminate the active materials.22 Anderson et al. took electrode materials out from cycled cells, and they found surface layer on cathode and anode. This formed film contained compounds such as LiF, LixPFy, and LixPFyOz mixture, different types of polycarbonates, and Li2CO3. These species can hinder lithium ion diffusion and jeopardize the electrochemical properties by insulating the transportation pathways from outside.23 Through symmetric cells, Chen et al. determined that impedance growth was primarily caused by charge transfer resistance increase on the electrode/electrolyte interface. Containing different inorganic and organic species of decomposing electrolyte, the solid electrolyte interface (SEI) film could suppress lithium ion mobility and diffusivity during electrochemical process.24 The lithium ion mobility is a determined factor to detect kinetics information at high voltage for NRL materials, which provide another perspective to capacity fade. To measure the detailed content of lithium ion mobility, tests based on lithium ion diffusion coefficient (DLi+) were carried out, such as galvanostatic intermittent titration technique (GITT) or potentiostatic intermittent titration technique (PITT) method. Whittingham et.al demonstrated that the DLi+ value of different LiNiyMnyCo1−2yO2 (0.50 ≥ y ≥ 0.33) materials was lower than the DLi+ value of LixCoO2.25 Nevertheless, the formation time of the new-born phase for NRL materials in high voltage region has not been confirmed and the relationships between structural changes and kinetics process have not been found out yet. Although much effort has been carried out to understand and improve the structure evolutions and electrochemical performance, the NRL oxide materials are still suffering from many remaining challenges as a result of structural instability and the sluggish Li+ diffusion before they could be used practically. To further understand the structure−performance relationships, the layered oxide QS-NMC622 material was first synthesized through the carbonate co-precipitation method. We then established a continuous coupled technique which combines PITT with in situ XRD measurement. This combined in situ technique could not only simultaneously examine structural changes and kinetics of the layered oxide such as QS-NMC622 in this work but help to understand and determine degradation mechanisms for other high-voltage cathode materials based on Li+ mobilization or structure-dependent activity during the first charge/discharge process.

2.1. Synthesis of LiNi0.6Co0.2Mn0.2O2 Material. All chemical reagents are of analytical purity and used without further purification. The precursor was prepared with the carbonate co-precipitation method. In the typical synthetic process, NiSO4·6H2O, CoSO4·7H2O, and MnSO4·H2O with stoichiometric mole ratio of Ni/Co/Mn = 6:2:2 were dissolved in deionized water to form 1.5 M metal sulfate solution (solution A). The 1.5 M precipitating agent (solution B) was composed of Na2CO3 and NH4HCO3 with mole ratio of 1:1, and 1.5 M alkaline buffer agent NH3·H2O was called solution C. First, the solutions A and B were pumped into a container at a fixed rate with mechanical stirring at 50 °C, respectively. Then, solution C was used to adjust the pH to 8.0. After stirring for 24 h, the carbonate precipitates were washed with distilled water, then filtered and dried under vacuum at 120 °C overnight. The as-prepared carbonate precursor and LiOH·H2O were mixed at a certain proportion in aqueous solution. The obtained brown mixture was collected by solvent evaporation at 60 °C. The LiNi0.6Co0.2Mn0.2O2 material was obtained after preheating the mixture at 450 °C for 5 h and then 750 °C for 12 h in air atmosphere. 2.2. Material Characterization. The phase structure of the asprepared powder sample was measured by X-ray diffraction (XRD) on the Rigaku Ultima IV powder diffractometer with Cu Kα radiation (λ = 1.548 Å), while in situ XRD test was conducted by another X-ray diffractometer (XRD2, Bruker D8 Discovers X-ray analytical systems with Cu Kα radiation). Scanning electron microscopy (SEM, Zeiss Sigma SEM) and transmission electron microscopy (TEM, JEOLJEM2100 microscope) were used to characterize the morphology and crystal structure of the products. 2.3. Electrochemical Measurements. The electrochemical properties of the LiNi0.6Co0.2Mn0.2O2 electrodes were measured by assembling them into coin cells (type CR2025) in an Ar-filled glovebox. The cathode was prepared by spreading the slurry onto an Al foil which is used as current collector and then dried under vacuum at 120 °C for 12 h. The slurry consisted of 80 wt % active material, 10 wt % acetylene black, and 10 wt % polyvinylidene fluoride (PVDF) dissolving in N-methylpyrrolidone (NMP). The loading active material mass of each Al foil was around 4 mg on a disk sliced electrode of 16 mm in diameter. Lithium foil was used as the anode electrode separated from the cathode by a piece of Celgard 2400 separator. The electrolyte was 1 M LiPF6 dissolved in a mixture solution of ethylene carbonate (EC) and dimethyl carbonate (DMC) (3:7, v/v). The charge/discharge tests were conducted at room temperature with the voltage range between 2.5 and 4.6 V using multichannel battery test system (LAND-V34, Land Electronic Co., Ltd., Wuhan). Cyclic voltammetry (CV) test was carried out on electrochemical workstation (CHI 660C) at a scan rate of 0.1 mV s−1 from 2.5 to 4.6 V. According to our definition, 1 C rate is by definition as equivalent to 160 mA g−1 in this work. 2.4. In Situ XRD during PITT Process. On the basis of our group’s previous works,19,26 several improvements about the method and device of the in situ XRD technique were achieved. In order to avoid the overlap for the diffraction peaks of Al foil and the LiNi0.6Co0.2Mn0.2O2 material, carbon substrates were used as the current collector to replace the former. The active material loaded on carbon substrates was around 10 mg. The device for in situ XRD studies was built based on 2025 coin cell which was punched and left a sizable circular hole in the center and then sealed with Kapton film. The device was prepared inside a glovebox with the same step of preparing ordinary coin cell. In the following test step, the device was fixed on a bracket and then adjusted by its coordinate. The XRD spetra were collected by two-dimension Bruker D8 Discovers X-ray diffractometer (XRD2) with operating voltage of 40 kV and current of 40 mA. The connected electrochemical workstation (CHI 660C) was used to gather data of PITT step at the same time. The voltage difference (ΔE) of each step was set to 25 mV with staying time of 1800 s at each process. Through this coupled measurement method, the XRD pattern and chronoamperometric curve (I−t curve) of each set voltages in the initial charge/discharge processes could be obtained. 24732

DOI: 10.1021/acsami.7b06326 ACS Appl. Mater. Interfaces 2017, 9, 24731−24742

Research Article

ACS Applied Materials & Interfaces

3. RESULT AND DISCUSSION 3.1. Structural and Morphological Characterization. The XRD pattern of the LiNi0.6Co0.2Mn0.2O2 (NMC622) powder is shown in Figure 1. It is clear that all diffraction peaks

Figure 2. (a) Low and (b) high magnification SEM images of precursor. (c) Low and (d) high magnification SEM images of the asprepared LiNi0.6Co0.2Mn0.2O2 particles.

Figure 1. XRD pattern of the as-prepared LiNi0.6Co0.2Mn0.2O2 powder material. “R” in the subscript of index refers to “rhombohedral”.

can be indexed to the crystal planes of a well-defined hexagonal α-NaFeO2 structure of R3̅m space group with no impurity. Normally for layered LNCMO materials, a less than 1.2 in the intensity ratio of (003) to (104) peaks [I(003)/I(104)] means that undesirable cation mixing would appear.27 In this regard, the I(003)/I(104) value of the as-prepared LiNi0.54Co0.23Mn0.23O2 material is 1.612 (>1.2), which implies no obvious cation mixing in the lattice. In addition, it can be observed that the (006)/(102) and (108)/(110) peaks are splitted obviously, indicating that the pristine sample is highly crystallized and well-layered. The lattice parameters of the structure can be obtained through the XRD Rietveld refinement by TOPAS software. The results indicate that the lattice parameters of the as-prepared layerd oxide materials, within the error range (Rp = 9.9%, Rwp = 18.0%, GOF = 3.26), are c = 14.213 Å, a = b = 2.870 Å, respectively. The ratio of c/a equals 4.953 which is higher than 4.899. It is another proof of the well layered structure of the as-prepared material.28 The SEM images of the precursor and LiNi0.6Co0.2Mn0.2O2 are shown in Figure 2. As shown in Figure 2, the precursor consists of uniform spherical particles with an estimated diameter of around 10 μm. It can be also observed that many nanorods are distributed on the surface of spherical aggregation. After two-step annealing, the quasi-spherical shape could be maintained, while the nanorods had transferred into small particles, as shown in Figure 2c and Figure 2d. Therefore, a lot of holes formed along the primary nanoparticles. The porous structure of the layered oxide material may be suitable for the infiltration of electrolyte and the transportation of lithium ions during the charge/discharge processes. Figure 3a and Figure 3b display TEM images of the asprepared QS-NMC622 material in different magnification. The as-prepared QS-NMC622 material presents hierarchical micro-/nanostructures, which consists of many primary nanoparticles with sizes of 100−200 nm. Figure 3c is a highresolution TEM image (HR-TEM) of the particle region pointed out by the red circle in Figure 3b. Figure 3d is a fast

Figure 3. (a) Low and (b) high magnification TEM image, (c) highresolution TEM image, (d) FFT graph, and (e) IFFT graph of LiNi0.6Co0.2Mn0.2O2 particle.

Fourier transform (FFT) graph of the chosen area marked by red dash rectangle, while Figure 3e is a corresponding inverse fast Fourier transform (IFFT) graph. From the FFT image, the hexagonal spots pattern is observed from [1̅2̅1]R zone axis of rhombohedral phase (“R” is rhombohedral phase for short). The hexagonal spots in Figure 3d could be indexed as (1̅11)R, (1̅01̅)R, and (01̅2̅)R crystal lattice plane, corresponding to the layered structure. The interfacial angle between (1̅11)R and (1̅01̅)R lattice plane is 62.4°. The distance between two lattices fringes in IFFT graph was measured to be 0.246 and 0.243 nm, which corresponds to the d-spacing of (101) plane from the rhombohedral phase.29 24733

DOI: 10.1021/acsami.7b06326 ACS Appl. Mater. Interfaces 2017, 9, 24731−24742

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Figure 4. High-resolution TEM image (a), FFT graph (b), IFFT graph (c), intensity plot from two different regions (d, e), and brief presentation of cationic order and disorder microstructural (f) LiNi0.6Co0.2Mn0.2O2 material.

cycle, an oxidation peak at 3.85 V and a homologous reduction response at 3.60 V can be obviously observed, corresponding to redox transition of Ni2+/Ni4+ with voltage interval of 0.25 V. However, a small peak appears at a relatively higher voltage caused by some side reactions on the particle surface. A slight slip in the oxidation process during the following two cycles can be observed, indicating a poor reversibility during Li+ insertion and extraction process in high voltage range. Figure 5b displays the charge/discharge curves of the asprepared layered oxide electrode during different cycles in the voltage range of 2.5−4.6 V at a rate of 0.2 C. The initial discharge capacity is 190.5 mAh g−1 with a Coulombic efficiency of 78.7%. The large irreversible capacity loss in the first cycle should be attributed to the decomposition of electrolyte and side reactions at high voltage.30 The capacity retention is 74.9% after 100 cycles and Coulombic efficiency is close to 100% for each cycle as shown in Figure 4c. The rate capability of the LiNi0.54Co0.23Mn0.23O2 electrode is shown in Figure 4d. The material cannot resist the large-current impact, as judged from the result of test at a high rate of 5 C. Overall, the as-prepared LiNi0.54Co0.23Mn0.23O2 material exhibited rapid capacity degradation at the cutoff voltage of 4.6 V. 3.2. In Situ XRD during PITT Process. The changes of (003)R, (018)R, and (110)R diffraction peaks position and intensity between charge and discharge of potential stepping are shown in Figure 6. The in situ 2D XRD view of the asprepared cathode material is shown in Figure 6a and Figure 6c, where the charge process corresponds to the pattern number

Figure 4 provides a direct demonstration of how the cations were arranged within the crystal lattice in one of the particles from pristine LiNi0.6Co0.2Mn0.2O2 material. Divided by the yellow dash line, the regions presented different crystal plane configuration in Figure 4a. Through FFT and IFFT graphs shown in Figure 4b and Figure 4c, the distance between lattice fringes can be determined and calculated to be presented in Figure 4d and Figure 4e. In general, the indexed bright spots in Figure 4b correspond to mainly (003)R, (100)R, and other diffractions viewed down from [010]R zone axis. Notably, (100)R and (103)R diffractions are both forbidden according to their structural factor, specifically in XRD patterns. The distinction between atomic sequences within the two yellow rectangle in Figure 4c can be seen from the relatively weak signal intensity peaks marked by the red arrows, which indicate the occupation of transition metal (TM) ions in sites of Li+ layers. This particular fact in the layered Ni-rich cathode materials is schematically sketched in Figure 4f: the TM ions, mostly nickel ions, migrate from TM layers to Li+ layers and take the place of Li+ ions due to their similar ionic radius (Ni2+ of 0.69 Å and Li+ of 0.76 Å), causing cationic disorder between the layers. The Rietveld refinement of XRD pattern gave a cation disorder of 3.14% for the pristine LiNi0.6Co0.2Mn0.2O2 material. Cyclic voltammetry experiments of the QS-NMC622 material between 2.5 and 4.6 V at a scan rate of 0.1 mV s−1 were carried out to investigate the redox potential of QSNMC622 electrode, as shown in Figure 5a. During the first 24734

DOI: 10.1021/acsami.7b06326 ACS Appl. Mater. Interfaces 2017, 9, 24731−24742

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ACS Applied Materials & Interfaces

Figure 5. (a) Cyclic voltammetry (CV) curves at a scan rate of 0.1 mV s−1. (b) Charge/discharge voltage profiles at 1st, 2nd, 10th, 50th, and 100th cycle at a rate of 0.2 C. (c) Cycling performance at the constant rate of 0.2 C. (d) Rate capacity at various rates between 0.1 and 5 C of LiNi0.6Co0.2Mn0.2O2 electrodes in the voltage range of 2.5−4.6 V.

Figure 6. (a) 2D view of selected in situ XRD patterns from 17.9 to 20.1°. (b) Interpreted 1D view of 17.7−20.2°. (c) 2D view of selected in situ XRD patterns from 63.6 to 67.1°. (d) Interpreted 1D view of 62.5−67.5° of as-prepared electrode during voltage stepping. “DS” and “FS” mean defect and full spinel, respectively.

from 1 to 31 and the discharge process corresponds to the pattern number from 32 to 76 in the Y-axis, respectively.

Different colors represent changes in peak intensity of the 2D view. 1D XRD patterns at selected voltage are displayed in 24735

DOI: 10.1021/acsami.7b06326 ACS Appl. Mater. Interfaces 2017, 9, 24731−24742

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ACS Applied Materials & Interfaces

Figure 7. Subregional comparison of the position and intensity of (003)R diffraction peak with respect to d-spacing during the whole voltage stepping (a), subregional XRD Rietveld refinement utilizing α-NaFeO2 structure, defective spinel, and combination of both from 4.45−4.73 Å (b−d) and 1.39−1.43 Å (e−g).

Figure 6b and Figure 6d to give a clear look at the shape of the particular diffraction peaks. Carbon substrates (peaks fixed at 26.0° and 54.5° in 2θ value) are chosen as reference substance for pattern comparison. Compared with the XRD pattern of the pristine powder, the diffraction patterns at the open circuit voltage (OCV) are nearly the same except the intensity of (003)R peak (low angle) is reduced by open battery shell edge, Kapton film, and electrolyte. From Figure 6b and Figure 6d, it could be observed that (003)R and (018)R peaks shifted toward lower 2θ angle during the initial charge from OCV to 4.05 V, while (110)R peak shifted toward higher 2θ angle slowly. The detail situations for other featured diffraction peaks are shown in Figure S1 in Supporting Information. Moreover, with the voltage increasing from 4.00 to 4.60 V, the position of (003)R and (018)R peaks abruptly shifted toward the opposite direction while the position of (110)R peak still moved to higher 2θ angle rapidly. From 4.50 to 4.60 V for charging, the split (018)R and (110)R peaks were merged into one single diffraction peak, indicating a degradation to the layered stacking. According to Figure 7a, the d-spacing distance has decreased about 0.2 Å from OCV to 4.60 V charged state, which indicates a serious lithium-poor configuration or even a microstructural transformation. On the basis of single peak search−match and refinement, simulated results of candidate diffractions from two different crystal structures to fit (003)R and (018)R/(110)R peaks at 4.60 V charged state are shown in Figure 7b−g. Figure 7b−d are respectively refinement pattern presentation of (003)R by using α-NaFeO2 structure (JCPDS No. 00-85-1968, Li0.89Ni1.01O2), defective spinel phase (JCPDS No. 00-82-0343, Li0.5CoO2), and compound structure consisting of both. So is Figure 7e−g for (018)R/(110)R. In terms of the three R-factors (Rp, Rwp, Rexp), at 4.60 V charged state, there is a higher probability that both defective spinel and defective O3 phases could coexist in Li1−xNi0.6Co0.2Mn0.2O2 (0.5 < x < 1) structure.31,32 Therefore, Rietveld refinement of full range XRD pattern with respect to 2θ value was run to help see the structural components of Li1−xNi0.6Co0.2Mn0.2O2 electrode at 4.60 V charged state. It turned out, as shown in Figure 8a−c, the combination of defective spinel and O3 phases not only makes the calculated XRD pattern more matched to the observed one but also

Figure 8. General XRD Rietveld refinement patterns and results utilizing α-NaFeO2 structure (a), defective spinel (b), and combination of both (c) at 4.60 V charged state. The amount of full spinel (d), defective spinel (e), and layered O3 (f) during the whole voltage stepping.

obtains a more rational refinement result (GOF = 1.21 is more approached to 1 than 1.63 or 2.57) compared to the other two options. The high-resolution image of one particle at 4.60 V can evidence the two phase coexistence as shown in Figure 9. Figure 9b and Figure 9c are FFT graphs with distinct diffraction 24736

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Figure S2. Back to Figure 6, during discharge, (003)R/DS(111)C and (018)R/(110)R/DS-(440)C diffraction peaks moved back to lower 2θ angle rapidly until the voltage had reduced below 4.00 V. Then they were stabilized in 2θ position to the end of the discharge. Compared to the original shape and position at OCV state, the peak at around 64.5 in 2θ value seems to belong to a parent layered phase with spinel structural trait below 3.20 V. Considered the lithiated process, the new phase could be full spinel constructed distinguished from defective spinel. Hence, by application of standard A1+xB2−xO4like spinel phase (JCPDS No. 00-89-8149, Li1.14Mn1.86O4) to discharge, along with standard layered O3 and defective spinel structures into the whole cycle, the quantitive amount of components through XRD Rietveld refinement can be achieved, and the results are shown in Figure 8d and Figure 8e. According to the quantitive analysis, the whole cycle could roughly be divided into four stages: parent layered phase in stage I; mixture of layered and defective spinel phases in stage II; mixture of layered, defective, and full spinel phases in stage III; mixture of layered and full spinel phases in stage IV. The shift of peak position reflects the variations of crystal structure during Li+ insertion and extraction according to the Bragg equation. A series of changes in structure parameter can also be obtained through the XRD Rietveld refinement. The results of lattice-parameter c and a, cation-disorder indicator Ni3b (Ni in Li+ layer), and layered-factor c/a are shown in Figure 10a−c, respectively. The overall results of lattice-

Figure 9. HR-TEM image (a), FFT graphs from two regions (b, c) corresponding IFFT graphs (d, e), intensity plots from different subregions (f−i), and brief schematic exhibition of different crystal lattice (j−l) of the particle at 4.60 V of the initial charge.

spot patterns obtained from two areas of the surface in Figure 9a. The two FFT graphs are indexed to be cubic and rhombohedral scenarios viewed along [011]C and [010]R zone axis, respectively. The corresponding IFFT graphs are shown in Figure 9d and Figure 9e. In Figure 9d, two related crystal facets arrangement can be found if carefully observed, which illustrates two crystal lattice structures. Figure 9f indicates the appearance of defect spinel with TM ions disordering in Li+ layers, while Figure 9g infers the presence of NiO-like framework with space group of Fm3m ̅ , which could mostly be caused by massive lithium ion extraction from the parent layered O3 structure. Figure 9j and Figure 9k are, respectively, the brief atomic occupation and ordering demonstration of these two crystal microstructures. In Figure 9e, rhombohedral and R3̅m featured (003)R and the forbidden (1̅00)R diffractions can be determined by the lattice fringe distances in Figure 9h and Figure 9i despite dislocations, which results from the d-spacing shrinkage of (003)R due to the formation of more Li+ site vacancies as shown in Figure 9l. TEM analysis of another particle at 4.60 V charged state also confirmed the sign of cubic crystal structure and is shown in

Figure 10. (a) Lattice-parameters c. (b) Lattice-parameters a. (c) Amount of Ni3b in Li+ layer. (d) c/a ratio through XRD Rietveld refinement during voltage stepping. 24737

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Figure 11. Brief schematic interpretation of atomic distribution and ordering in microstructural framework during four stages.

were oxidized and migrated back to TM layers due to unfitness of the ionic radius, which caused Ni3b to decrease. From 4.10 to 4.30 V, lattice-parameter c started to decrease due to shrinkage of c-axis, which could be caused by much lithium ions loss. During this voltage section, all nickel ion completed migrating from Li+ layers to TM layers, leaving Ni3b to 0 as shown in Figure 10c. However, part of the remaining lithium ions hopped into the tetrahedral sites face shared with the Mn octahedron to form the so-called “Li−Mn dumbbell” as shown in Figure 11, which is a pre-sign of defective spinel framework.33 The amount of defective spinel in the parent O3 phase increased to 5.6% at 4.30 V charged state as shown in Figure 8e. Started from 4.35 V, caused by the compression of Li+ layers after massive Li+ extraction, lattice-parameter c declined at a huge rate along with the c/a ratio, which decreased to 4.858 ( 0.1), the I(t) shown as eq 1 can be obtained.

component, making the amount of defective spinel phase surge to 22.2% at 4.60 V charged state. From 4.55 V down to 3.725 V for discharge, latticeparameter c bounced back from the lowest point and increased rapidly due to the c-axial pillaring by massive inserting lithium ions. Lattice-parameter a showed the same trend owed to the reduction of TM ions. During this voltage section, driven by the filling lithium ions, nickel ions used to staying in Li+ layers could migrate to TM layers, leaving Ni3b dropping and deconstruction of the defective spinel phase. From 3.70 to 3.60 V, except for continuously increasing lattice-parameter a, the amount of defective spinel decreased all the way down since much more lithium ions took back the positions and compromised the cubic framework. (3) Stage III. From 3.55 to 3.30 V, lattice-parameter c was kept stabilized because of the O3 structural integrity recovery. Lattice-parameter a increased back to original number. Notably, based on the local defective spinel framework, more tetrahedral lithium ions appeared neighboring around TM ions in Li+ layers as shown in Figure 11 to make themselves into single spinel unit. Therefore, a full spinel phase of Li1−x(Ni0.6Co0.2Mn0.2)2+xO4 formed during this stage associated with the increasing Ni3b, leading the amount of defective spinel to 0 at 3.30 V discharged state, which suggests that the latter is an intermedium during Li+ de-/intercalation. (4) Stage IV. From 3.25 to 2.50 V, the Ni3b kept increasing and gave ingredients for spinel unit to grow, which made lattice-parameter a increase to the end of the discharge because full spinel with Fd3̅m space group is more enlarged in a−b plane than layered O3 phase. At 2.50 V discharged state, the parent O3 structure has taken 96.5% and the new-born full spinel has taken 3.5%. As an impurity to the layered structure, the Li1−x(Ni0.6Co0.2Mn0.2)2+xO4 may not only sabotage the

⎛ π 2tD + ⎞ Li ⎟ I(t ) = I0 exp⎜ − 2 4L ⎠ ⎝

(1)

According to the ln I(t)−t plots, the lithium ion diffusion coefficient (DLi+) of the as-prepared material can be calculated from the slope of the line. A computational formula for lithium ion diffusion coefficient is shown as follows (eq 2): D Li+ = −

4L2 d ln(I ) π 2 dt

(2)

In eq 2, L refers to the diffusion distance which can be regarded as the thickness of cathode electrode material, I represents the step current, and t is the step time in test procedure.34 The average thickness of the cross section of electrode material was measured to be 13.25 μm from SEM characterization as shown in Figure S3. Figure 12a and Figure 12b show the chronoamperometry curves of charge and discharge processes, respectively. The XRD patterns were collected between 900 and 1500 s to ensure the current value is nearly constant. The corresponding ln I(t)− t curves of relevant voltages between 900 and 1800 s during charge and discharge processes are shown in Figure S4. Through eq 2, the DLi+ can be calculated. The values of log DLi+ at selected voltages are shown in Figure 8c,d. The variation range of DLi+ in initial charge and discharge processes is 10−9− 24739

DOI: 10.1021/acsami.7b06326 ACS Appl. Mater. Interfaces 2017, 9, 24731−24742

Research Article

ACS Applied Materials & Interfaces 10−10 cm2 s−1, which is in the same order of magnitude as for LiNi1/3Co1/3Mn1/3O2 material (10−9−10−10 cm2 s−1).25,35 In the initial charge process (Figure 12c), the DLi+ decreases rapidly to minimum value (1.75 × 10−10 cm2 s−1) at 3.60 V and then quickly increases a lot. The reduction of DLi+ from OCV to 3.60 V may be affected by the following factors: (1) There are rare octahedral vacancies in the Li+ layers for Li+ diffusion, since the charge process is just at the beginning and only few lithium ions have been extracted from the octahedral site;36 (2) latticeparameter c decreased causing Li−O2 slab to contract to make migration space smaller for lithium ions. When the voltage increases from 3.60 to 4.00 V, a lot of lithium ions are extracted from the Li+ layers leaving behind several octahedral vacancies, and the lattice-parameter c increases and the Ni3b decreases. Therefore, the channels for Li+ diffusion become abundant and free from the disturbance of cationic disorder. Meanwhile, the octahedral vacancies of Li+ layer are dominated, so the value of DLi+ recovers partly. The value of DLi+ decreases a lot from 4.00 to 4.60 V, which can be attributed to three main factors: (1) the decrease of lattice-parameter c and the increase of crystal defects such as dislocation; (2) in this process, the proportion of Ni3+/Ni4+ increases, and therefore, the electrostatic repulsion between TM ions and lithium ions is strengthening which can inhibit the movement of Li+; (3) the massive formation of the defect spinel phase. Li+ diffuses slowly between two phases (rhombohedral phase and spinel phase); (4) the number of nickel ions migrating into vacant Li+ sites was increasing and caused hindering to Li+ diffusion; (5) the byproducts of electrolyte decomposition at high voltage (>4.40 V) have an insulating influence to the Li+ diffusivity and transportation at the surface and reduce the kinetic process. The values of DLi+ in discharge are shown in Figure 12d. Generally, intercalation is dominated and energetically favored for discharge rather than diffusion out for lithium ion. When the voltage reaches 4.55 V, the value of DLi+ suddenly grows larger than that of the end in charging state, which may be attributed to a lot of space in Li+ layer that is still empty at the beginning of the discharge process. Then the DLi+ value is reduced greatly from 4.55 to 3.60 V on account of the Li+ continuous insertion into lithium layer occupying octahedral vacancies, which could hinder the diffusion of Li+ and make available migration sites in short. The conversion from defective spinel structure to full spinel phase and increasing Ni3b caused the value of DLi+ to decrease from 3.55 to 2.50 V. However, the decreasing rate is much smaller than that of the previous voltage section, which might be attributed to the reduction of the electrostatic repulsions between TM ion and lithium ion resulting from the transformations of Ni4+/Ni3+ to Ni2+.

lattice-parameter a remains decreasing almost throughout the whole charge process. Above 4.40 V charged state, the degree of cationic disorder in Li+ layers became deeper and triggered the formation of Li0.5CoO2-like defective spinel framework, which can be evidenced by HR-TEM analysis and PITT results. The crystal defects and new-born structure can severely hinder the mobility of lithium ions on the surface at high charged voltages. The variation tendencies of lattice-parameters c and a in the discharge process are nearly inverse compared with that of the charge process. Notably, at below 3.60 V discharged state, TM/Li+ disorder started increasing, which cause a AB2O4like full spinel phase to rise based on the foundation of defective spinel associated with the existence of tetrahedral Li+. According to the XRD quantitative analysis, the amount of full spinel Li1−x(Ni0.6Co0.2Mn0.2)2+xO4 in the parent O3 phase is about 3.5%, which indicates that it will sustain as impurity during successive cycling. It can be rationally speculated that the growing full spinel would jeopardize the body structural stability and impact the mobility and migration of lithium ions if it accumulates itself during subsequent high voltage cycling. On the basis of the above-mentioned speculation, massive formation of defective spinel framework by charging to a deeper delithiated degree brings in the full spinel appearance in the body structure, which is the cause of capacity degradation of Ni-rich layered cathode materials in high voltage range. The coupled PITT and in situ XRD measurement system has played an important role in the determination of fading mechanism for the quasi-spherical LiNi0.6Co0.2Mn0.2O2 material, and it will be useful for exploring structural changes and kinetics processes of other high voltage cathode materials of LIBs.

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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.7b06326. XRD patterns, TEM analysis results, cross section of electrode slice, charge and discharge results, XRD Rietveld parameters, and parameter refinement results (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*L.H.: phone, +86-18959283629; fax, +86-592-2180181; email, [email protected]. *S.-G.S.: e-mail, [email protected]. ORCID

4. CONCLUSION In summary, the quasi-spherical Ni-rich layered oxide LiNi0.6Co0.2Mn0.2O2 material was successfully synthesized through the carbonate co-precipitation method. The XRD and TEM results illustrated that the as-prepared material is a well-defined hexagonal layered α-NaFeO2 structure with R3̅m space group. Moreover, lithium-ion storage performance with the initial discharge capacity of 190.5 mAh g−1 in the voltage range of 2.5−4.6 V at 0.2 C could be obtained and the discharge capacity retention was relatively lower for only 74.9% after 100 cycles. In situ XRD results indicated that the lattice-parameter c value increases at low voltage (below 4.00 V) and drastically reduces at higher voltage in the charge process, while the

Ling Huang: 0000-0003-1092-5974 Jun-Tao Li: 0000-0002-9650-6385 Author Contributions §

Q.W. and C.-H.S. contributed equally to this work

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was jointly supported by the NSFC (Grants 21673194, 21621091, and 21373008), the National Key Research and Development Program (2016YFB0100202), and NFFTBS (Grant J1310024). 24740

DOI: 10.1021/acsami.7b06326 ACS Appl. Mater. Interfaces 2017, 9, 24731−24742

Research Article

ACS Applied Materials & Interfaces

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