Enhanced Electrochemical Performance of LiNi0.8Co0.1Mn0.1O2

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Enhanced Electrochemical Performance of LiNi0.8Co0.1Mn0.1O2 Cathode for Lithium-Ion Batteries by Precursor Pre-oxidation Congcong Zhang, Mengmeng Liu, Guangjie Pan, Siyang Liu, Da Liu, Chunguang Chen, Junming Su, Tao Huang, and Aishui Yu ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.8b00994 • Publication Date (Web): 01 Aug 2018 Downloaded from http://pubs.acs.org on August 7, 2018

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ACS Applied Energy Materials

Enhanced Electrochemical Performance of LiNi0.8Co0.1Mn0.1O2 Cathode for Lithium-Ion Batteries by Precursor Pre-oxidation Congcong Zhang,† Mengmeng Liu,† Guangjie Pan,‡ Siyang Liu,§ Da Liu,§ Chunguang Chen,§ Junming Su,† Tao Huang,† Aishui Yu*,†,§ †

Laboratory of Advanced Materials, Shanghai Key Laboratory of Molecular Catalysis and

Innovative Materials, Collaborative Innovation Center of Chemistry for Energy Materials, Institute of New Energy, Fudan University, Shanghai 200438, China ‡

Central Academe, Shanghai Electric Group Co., Ltd, Shanghai 200070, China

§

Department of Chemistry, Fudan University, Shanghai 200438, China

*

Corresponding author: Phone: +86-21-31249125; Fax: +86-21-31249125; E-mail: [email protected]

Abstract: Nickel-rich layered oxide LiNi0.8Co0.1Mn0.1O2 suffers from severe structural instability, causing inferior electrochemical performance. To solve this problem, a Na2S2O8 pre-oxidation method is employed to modify the surface structure of precursor Ni0.8Co0.1Mn0.1(OH)2. Transmission electron microscopy images show that the lattice orientations of precursor are well ordered, and the resulted product LiNi0.8Co0.1Mn0.1O2 with this precursor exhibits well-defined layered structure without cation-mixing layer on the surface. X-ray photoelectron spectroscopy and Rietveld refinement results indicate that the contents of Ni2+, Co2+, and Li+/Ni2+ disordering ratio are significantly reduced at the same time. ICP-AES and titration results suggest that the average oxidation state of Ni is enhanced after Na2S2O8 pre-oxidation. Further electrochemical kinetic analysis using electrochemical 1 ACS Paragon Plus Environment

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impedance spectroscopy and potentiostatic intermittent titration technique reveals that the LiNi0.8Co0.1Mn0.1O2 sample after precursor pre-oxidation owns the fast charge transfer and Li+ diffusion process. It also performs excellent cycling stability and rate capability. Remarkably, the sample with an optimum oxidation time of 30 minutes (S-NCM-30min) delivers a high discharge capacity of 203.5 mAh g-1 and retain 99.0% capacity after 100 cycles in the voltage range of 3.0-4.3 V. The superior electrochemical performance is attributed to the well-ordered surface structure with Na2S2O8 pre-oxidation, which can suppress the anisotropic shrinkage/expansion and meanwhile stabilize the original layered structure of LiNi0.8Co0.1Mn0.1O2 material during repeated charge-discharge cycling. Keywords:

LiNi0.8Co0.1Mn0.1O2;

precursor;

Na2S2O8

pre-oxidation;

cycling

performance; lattice defect; anisotropic shrinkage/expansion Introduction Rechargeable lithium-ion batteries (LIBs) have become a promising power source for portable electronic devices due to their high energy density, high voltage, long cycle life, and light weight.1-3 The rapid development of electric vehicles has stimulated the demand for LIBs with higher gravimetric and volumetric energy densities. Recently, Ni-rich layered oxide cathodes (LiNi1-x-yCoxMnyO2, 1-x-y ≥ 0.6) have been considered to be the competitive candidates to satisfy such requirements due to their higher discharge capacity (200~220 mAh g-1), higher energy density (>800 Wh kg-1) and lower cost.4 However, some problems still must be overcome for large-scale application, such as: (1) Li+/Ni2+ cation mixing due to the large energy barrier in the 2 ACS Paragon Plus Environment

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oxidation process of Ni2+ to Ni3+;5 (2) irreversible structural transformation from the original layered phase to the disordered spinel-like or NiO-like rock-salt phases;6-7 (3) thermal instability and safety issues caused by the reaction between strong oxidizing Ni4+ and electrolyte;8-9 and (4) poor storage property.10 Many researchers have used surface modifications, such as surface coating11-15 and core-shell structures,16-19 to resolve the above problems, focusing mainly on the modification of final products. To some extent, precursors are vital to the performance of final products, and precursors with a perfect crystal lattice will produce high-quality products. However, almost all the primary particles of precursors have different crystallographic orientations and slip planes, and anisotropic lattice volume expansion or contraction between grains can result in microcracks.20 Intergranular fracture allows the penetrating of electrolyte through the cracks and develops new reaction sites at the electrode/electrolyte interface, producing a resistive cathode-electrolyte interface (CEI) layer on the surface of grains.21 The CEI layers inhibit the transport of electron and lithium ions through the grain-electrolyte interface, inhibiting the electrical performance.22 Therefore, it is impotant to stabilize the structure of final products by eliminating crystal defects of the precursors. Tang et al.

have

reported

a

surface-oxidation

method

in

which

the

precursor

Ni0.815Co0.15Al0.035(OH)2 is oxidized into Ni0.815Co0.15Al0.035OOH by Na2S2O8.23 The pre-oxidation of the hydroxide precursor can eliminate crystal defects and diminish cation mixing, producing a complete and ordered layered structure on the surface of LiNi0.815Co0.15Al0.035O2. 3 ACS Paragon Plus Environment


In this study, based on the fact that Ni(OH)2 can be oxidized by Na2S2O8 into β-NiOOH, we modified the surface structure of the precursor Ni0.8Co0.1Mn0.1(OH)2 with Na2S2O8 to increase the average valance state of Nickel and Cobalt ions. The final product, LiNi0.8Co0.1Mn0.1O2 produced from pre-oxidized precursor, was synthesized by a solid-state reaction method, and the relationship between oxidation time and its impact on electrochemical performance was discussed in detail. Electrochemical impedance spectroscopy (EIS), potentiostatic intermittent titration technique (PITT), scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray photoelectron spectroscopy (XPS), and X-ray diffraction (XRD) were adopted to gain deep insight into the fundamental mechanism of the enhanced performance. Experimental Section Material Synthesis. LiNi0.8Co0.1Mn0.1O2 (NCM) was synthesized by a solid-state reaction method. Precursor Ni0.8Co0.1Mn0.1(OH)2 (NCMOH, Guizhou Zoomwe Zhengyuan Advanced Materials Co., LTD., China) and LiOH·H2O with a molar ratio of 1: 1.05 were thoroughly mixed by a Fritsch planetary micro-miller (Pulverisette 7, Germany), and then preheated at 480ºC for 6 h in a muffle furnace. The obtained materials were mixed again and further calcinated at 800ºC for 18 h under a flowing oxygen atmosphere. To obtain pre-oxidized precursor, moderate Na2S2O8 and NaOH corresponding to a molar ratio of 2:1 were dissolved together in 100 mL deionized water with continuously magnetic stirring. After the mixed solution becoming transparent, 10 g 4 ACS Paragon Plus Environment

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NCMOH was added in the solution and stirred by magnetic stirring at 50ºC for 15, 30, and 60 min, respectively. The brown color of NCMOH was turned to black immediately in the beginning of the reaction, as shown in Figure S1e and f. The resultant powder was centrifuged and washed with deionized water, and then dried at 80ºC for 14 h in a vacuum oven. For convenience, the pre-oxidized precursor was abbreviated as S-NCMOH-15min, S-NCMOH-30min, and S-NCMOH-60min, respectively. Through the same prepare procedure as NCM materials, the pre-oxidized precursor was also reacted with LiOH·H2O to obtain final products (denoted as S-NCM-15min, S-NCM-30min, and S-NCM-60min hereafter). Material Characterization. The crystalline structure and morphology of electrodes were characterized by X-ray diffraction (XRD) on a Bruker D8 Advance X-ray diffractometer equipped with Cu Kα radiation (λ = 1.5406 Å), high resolution transmission electron microscopy (HRTEM, JEOL JEM-2100F), and field-emission scanning electron microscopy (FE-SEM, Hitachi S-4800). Rietveld refinement analysis was carried out using the TOPAS software package (Total Pattern Analysis Solution Software, version 4.2) in the 2θ range of 10 º-80º with a step size of 1.2º min-1. The surface chemical compositions of the samples were measured by X-ray photoelectron spectroscopy (XPS, ESCA PHI500C) with Al Kα radiation (hν =1486.6 eV), and all the spectra were calibrated with the C 1s peak at 284.8 eV. Inductively coupled plasma-atomic emission spectrometry (ICP-AES, 107 Thermo E. IRISDuo) was used to measure the content of Ni and Co. The oxidation state of Ni was determined by the chemical titration method. About 0.05 g samples were dissolved in 5 ACS Paragon Plus Environment


50 mL of 0.02 M FeSO4 which is in 5 M H2SO4 solution and corresponds to the Fe2+ excess. The solution was titrated with 0.02 M KMnO4 to react with the excess Fe2+ ions. Electrochemical Measurements. The working electrodes were fabricated by mixing 80 wt% active materials, 10 wt% carbon conductive agents (Super P), and 10 wt% polyvinylidene fluoride (PVDF), which were all dissolved in N-methyl-pyrolline (NMP) and compressed onto aluminum foil. After being thoroughly dried, the films were cut into wafers (12 mm×12 mm) as cathode electrodes and dried overnight at 80ºC in a vacuum oven. CR2016 coin-type cells were assembled in an Ar-filled glovebox (Mikarouna, Superstar 1220/750/900), using 1 M LiPF6 in an ethylene carbonate (EC)-dimethyl carbonate (DMC)- ethylmethyl carbonate (EMC) mixture (1:1:1 volume ratio) as electrolyte, Li metal as an anode, and Celgard 2300 as a separator. Galvanostatic charge-discharge measurements were performed on a battery test system (Land CT2001A, Wuhan Jinnuo Electronic Co., Ltd.) at room temperature at various rates (1 C = 200 mA g-1). Electrochemical impedance spectroscopy (EIS) measurements were conducted on a Zahner IM6 electrochemical workstation (Zahner-Elektrik GmbH & Co. KG, Kronach, Germany) in the frequency range from 100 kHz to 10 mHz with an AC amplitude of 5 mV. The EIS results were simulated using ZView software. To measure the Li+ diffusion coefficient, the potentiostatic intermittent titration technique (PITT) was performed on an electrochemical workstation (CHI 660D) by applying a voltage step of 25 mV. The voltage window was set from 3.2 to 4.8 V. 6 ACS Paragon Plus Environment

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Ex-situ Structural Analysis. The electrodes after 100 cycles at 25ºC were taken out from coin cells and washed with dimethyl carbonate solvent (DMC) for several times and dried at room temperature in an argon-filled glove box. Then the electrodes were transformed to the ex-situ XRD cell with PE film as a window. The active materials were scrapped off from electrodes and dispersed in ethanol. Then a droplet of the ethanol was deposited on a copper microgrid for HRTEM measurement. Results and Discussion Figure 1a shows the XRD patterns of NCMOH and S-NCMOHs. Overall, the diffraction peaks of all precursors are indexed to the Ni(OH)2 phase, other than an extra diffraction peak (002) present in S-NCMOHs, corresponding to the NiOOH phase (JCDPS 06-0141).24 The intensity of (002) peak increases with the extension of the oxidation time of Na2S2O8. The diffraction peaks of S-NCMOHs slightly shift to a higher angle and broaden, which results from the generation of the NiOOH phase and its poor crystallinity.23 Figure 1b shows the XRD patterns of final products. The diffraction peaks of all products are indexed to a typical hexagonal α-NaFeO2 _

structure with the space group R3m, and the clear splitting of the (006)/(012) and (018)/(110) doublets suggests the formation of a well-ordered layered structure.25 The crystal structures of NCM and S-NCMs were further analyzed by Rietveld refinements (see Figure 1c-1f), and the results are presented in Table 1. The increase of oxidation time gives rise to the growth of the volume values. The c/a ratios, which are used to evaluate the layered structure of cathode materials, are almost identical and greater than 4.95.26 Furthermore, the ratios of Li+/Ni2+ cation mixing decrease 7 ACS Paragon Plus Environment


with the increasing the oxidation time. The expansion of lattice volume and lower cation mixing may can facilitate Li+ diffusion and improve electrochemical performance.

Figure 1. XRD patterns of (a) NCMOH and S-NCMOHs, and (b) NCM and S-NCMs; Rietveld refinement of (c) NCM, (d) S-NCM-15min, (e) S-NCM-30min, and (f) S-NCM-60min. 8 ACS Paragon Plus Environment

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Table 1. Rietveld analysis results of NCM and S-NCMs Samples

a (Å)

c (Å)

Volume (Å3)

c/a

Ni2+ in Li site (%)

NCM

2.8712

14.1990

101.3718

4.9453

3.606

S-NCM-15min

2.8721

14.2046

101.4732

4.9457

3.371

S-NCM-30min

2.8719

14.2105

101.5010

4.9481

3.289

S-NCM-60min

2.8720

14.2096

101.5054

4.9476

2.522

SEM images of NCMOH and S-NCMOHs are shown in Figure 2a-2d. All precursors show a micro-sized spherical structure and compose of numerous nanosheets. Their secondary particles are homogeneously distributed, with an average particle diameter of 10 µm (see Figure S1). Interestingly, S-NCMOHs reveal loose packing of nanosheets, while these nanosheets of NCMOH agglomerate together compactly. Loose stacking is beneficial for the permeation of melted LiOH, resulting in decreased lithium residues on the surface of final products and uniform distribution of primary particles. Some fragments arise in S-NCMOH-60min due to the over-oxidation, as shown in Figure S1d, may slacking the cycling performance. Figure 2e-2h show SEM images of NCM and NCMs. The sheet-like structure of precursor changed into numerous nano-sized primary particles of 200-500 nm. The primary particles of S-NCM-15min and S-NCM-30min samples become rounded with 9 ACS Paragon Plus Environment


smooth skin and uniform size, while S-NCM-60min shows a bigger and more uneven primary particle. Large primary particle will extend the length of Li+ diffusion pathways and results in poor electrochemical performance.4

Figure 2. SEM images of (a) NCMOH, (b) S-NCMOH-15min, (c) S-NCMOH-30min, (d) S-NCMOH-60min, (e) NCM, (f) S-NCM-15min, (g) S-NCM-30min, and (h) S-NCM-60min.

Figure 3 exhibits the surface structure of the precursors and final products by HRTEM. Many nanocrystals with different orientations appear in NCMOH, as 10 ACS Paragon Plus Environment

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pointed out by the yellow rectangle in Figure 3a, indicating that crystal defects and uneven phases may exist in the primary particles.23 After Na2S2O8 oxidation, these nanocrystals disappear, but a uniform and ordered phase forms with a thickness of 25 nm on the surface of S-NCMOH-30min (Figure 3b). Lattice fringes with a spacing of 0.47 nm manifest that this phase can be indexed to the (001) plane of β-NiOOH. In the oxidation process, Ni2+ is first oxidized into Ni3+, followed by the nucleation and crystal growth of a β-NiOOH phase.23 In the following solid-state reaction, it can be easily turned into pure layered NCM without surface defects. As shown in Figure 3d, _

S-NCM-30min exhibits a well-defined layered structure with a space group of R3m, which is confirmed by the fast Fourier transform (FFT) patterns (inset of Figure 3d). _

However, a NiO-like cation-mixing layer with a space group of Fm3m (region II in Figure 3c) and thickness of 1~2 nm is shown in the NCM sample, arising from the lattice defects of NCMOH.

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Figure 3. HRTEM images of (a) NCMOH, (b) S-NCMOH-30min, (c) NCM, and (d) S-NCM-30min. Insets show corresponding FFTs.

XPS was recorded to confirm the contents of nickel and cobalt on the surface of precursors and final products. According to the literature, Ni 2p3/2 and 2p1/2 peaks of NCMOH located at 855.0 and 872.6 eV, respectively, are consistent with Ni2+, as shown in Figure 4a.27 After of Na2S2O8 pre-oxidation, the new peaks of the Ni 2p3/2 at 856.2 eV and Ni 2p1/2 at 874.1 eV indicate that the valence state of Ni near the surface is a mixture of Ni2+ and Ni3+.28 The new peaks of the Co 2p3/2 at 780.0 eV and Co 2p1/2 at 795.2 eV also suggest the appearance of Co3+ in the S-NCMOH-30min surface (Figure 4b). The shift to higher binding energy of a Co2+ peak for S-NCMOH-30min may be ascribed to the slight change of the transition-metal (TM) layer.23 In order to balance the valence state, the Mn2+ content of S-NCMOH-30min increase (Figure S2). As shown in Figure 4c, the deconvoluted peaks of the Ni 2p3/2 at 855.0 and 856.3 eV for NCM are very close to those for S-NCM-30min (854.8 and 856.1 eV), demonstrating that both samples near the surface are a mixture of Ni2+ and Ni3+.11 The contents of Ni3+ in NCM and S-NCM-30min are estimated at 42.9% and 47.0%, respectively. The existence of ordered β-NiOOH phase on the surface of S-NCMOH-30min clearly increases the fraction of Ni3+ in S-NCM-30min. After Na2S2O8 pre-oxidation, all Co2+ in precursor changes to Co3+ in S-NCM-30min due to the relatively lower energy barrier in the oxidation process of Co2+ to Co3+ (Figure 4d). Combined by ICP-AES and titration method, we further calculated their average 12 ACS Paragon Plus Environment

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oxidation states. As shown in Table S1-S4, the oxidation states of Ni for NCMOHs increase from +2.071 to +2.472 after Na2S2O8 pre-oxidation, while the finial products enhance from +2.623 to +2.879. As the average oxidation states of TM increase, Li+/Ni2+ mixing on the surface of S-NCMs decreases significantly because less Co2+ and Ni2+ migrate into Li+ slabs, producing well-defined layered structure and promoting Li+ diffusion. The lower Li+/Ni2+ cation mixing, which is in accordance with Rietveld refinement, is expected to cause better electrochemical performance.



Figure 4. XPS spectra of (a) Ni 2p for NCMOH and S-NCMOH-30min, (b) Co 2p for NCMOH and S-NCMOH-30min, (c) Ni 2p for NCM and S-NCM-30min, and (d) Co 2p NCM and S-NCM-30min. To study the impact of Na2S2O8 pre-oxidation on electrochemical performance, initial charge-discharge cycling were tested at 25ºC. As shown in Figure 5a, NCM sample delivers a discharge capacity of 202.8 mAh g-1, while the discharge capacities 14 ACS Paragon Plus Environment

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of S-NCM-15min, S-NCM-30min, and S-NCM-60min are 202.8, 203.5, and 202.5 mAh g-1, respectively. All samples clearly exhibit similar electrochemical behavior, despite the higher discharge capacity of S-NCM-30min. Similarly, the first coulombic efficiency increases with the augment of oxidation time at first, and then diminishes. The

coulombic

efficiency

of

NCM,

S-NCM-15min,

S-NCM-30min

and

S-NCM-60min are 86.9%, 86.3%, 88.4%, and 86.0%, respectively (Table S5). The obvious decrease of irreversible capacity loss for S-NCM-30min is ascribed to the longer time of Na2S2O8 oxidation compared to S-NCM-15min, which lowers the Li+/Ni2+ mixing and accelerates the kinetics of Li+ ion diffusion. Continuing to increase the oxidation time, however, the S-NCM-60min obtains a larger primary particle (see Figure 2h), which extends the Li+ diffusion distance and leads to the inferior first coulombic efficiency. To highlight the effect of Na2S2O8 pre-oxidation, a higher charge voltage of 4.8 V was used to evaluate its electrochemical performance. Since a higher charge voltage causes more Li+ to insert/extract from the crystal lattice, their first discharge capacities increase greatly (Figure 5b). Similar to the circumstance of 3.0-4.3 V, S-NCM-30min still shows a superior initial specific capacity (219.1 mAh g-1), while the NCM, S-NCM-15min, and S-NCM-60min have values of 216.3, 218.7 and 216.1 mAh g-1, respectively. This result proves that S-NCM-30min, with few cation-mixing layer and a lower activation energy barrier for Li+ migration, harvests the largest reversible capacity. In the next section, we will systematically investigate the kinetics of Li+ ion diffusion during the first charge process by EIS and PITT analysis. 15 ACS Paragon Plus Environment


Figure 5c and d present the Nyquist plots of NCM and S-NCM-30min during the first charge process. Based on the equivalent circuit in Figure 5g, the fitted values of Rsf and Rct are shown in Figure 5e (exact data is in Table S7). Rsf and Rct are assigned to the overall resistance of interparticle contact of the electrode and Li+ migration through the electrode surface film, and charge-transfer resistance respectively.7 At the beginning of charge, almost all Li+ diffusion pathways are closed due to an inactive film covering on the electrode surface,29 as indicated by the very large Rsf and Rct values. Starting from 3.8 V, the rapid decline in Rsf and Rct are observed, which is ascribed to the surface inactive film that has been activated by the current flow and provides more open channels for Li+ transfer.7 Although the Rsf of S-NCM-30min is somewhat larger compared to that of NCM, and Rct values are comparable for both samples in the voltage range of 3.8-4.4 V, both of these impedance parameters increase remarkably for NCM sample when being charged above 4.4 V. The increase of Rsf values of NCM during a higher potential region suggests an inactive and fast-growing CEI layer on the particle surface. This CEI layer is attributed to the intergranular fracture driven by mechanical stress between grains during lithium deintercalation, which provides new reactive sites for the electrode and electrolyte.30 Fortunately, Na2S2O8 pre-oxidation reduces the crystal defects on the surface of S-NCM-30min and helps establish a stable CEI layer beneficial for Li+ diffusion, and thus brings about relatively invariant Rsf values. Also, the values of Rct for NCM increase quickly, as indicated by the enlarged graph in the inset of Figure 5e. We attribute the rise of Rct for NCM at high potential to the surface phase transformation 16 ACS Paragon Plus Environment

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that provides an extra barrier to Li+ transfer. In contrast to NCM, S-NCM-30min exhibits a limited increasing trend in Rct values, implying its structural stability and lower activation energy for charge transfer even at high potential.31 These results correlate well with the improved reversible capacity and coulombic efficiency of S-NCM-30min during the first charge-discharge cycling (see Figure 5a and b). Based on the evolution of a one-dimensional finite-space diffusion model, PITT was deployed to identify the lithium diffusion coefficient (DLi+) during the first charge process. By a particular assumption regarding the diffusion process along with the solution of a partial differential equation to Fick’s second law with mathematical manipulation, DLi+ can be obtained as

‫ܦ‬௅௜ା = −

ୢ୪୬(ூ) ସ௅మ ୢ௧

∙

(1)

గమ

where I, t, and L refer to step current, step time, and diffusion distance, which can be regarded as the cathode electrode thickness.32-33 By combining the slopes of lnI-t plots with the electrode thickness (L=130 µm, see Figure S3), DLi+ can be calculated through formula (1). Figure 5f shows that the DLi+ values in the initial charge process (˂ 3.8 V) are 1-2 orders of magnitude smaller than the subsequent charge process, which is in accordance with the large impedance parameters at the beginning of Li+ extraction (Figure 5e). From 3.8 to 4.1 V, the Li+ diffusion kinetics enhances quickly, and the maximum DLi+ of S-NCM-30min reaches 8.66×10-8 cm2 s-1 (exact data are shown in Table S7). When charging beyond 4.1 V, the enhancement of electrostatic repulsion between TM and lithium ions,33 the different phases (rhombohedral phase and spinel or rock-salt phases),34 and the byproducts of electrolyte decomposition35 all 17 ACS Paragon Plus Environment


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inhibit Li+ diffusion. Therefore, DLi+ values of both samples begin to decrease. Under a higher potential region (> 4.4 V), the lithium diffusion process continues to slow down, and the DLi+ value of NCM declines to 6.15×10-9 cm2 s-1 at 4.8 V. Among the two samples, DLi+ of S-NCM-30min is higher than for NCM during the entire charge process,

which

again

explains

its

enhanced

reversible

charge-discharge cycling.


capacity

at

first

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Figure 5. Initial charge-discharge curves for NCM and S-NCMs at 0.1 C at 25ºC (a) between 3.0 and 4.3 V, and (b) between 3.0 and 4.8 V; variation of Nyquist plots of (c) NCM and (d) S-NCM-30min during the first charge process measured at several potentials up to 4.8 V; (e) variations of fitted Rsf and Rct values for NCM and S-NCM-30min as a function of step potential, with enlarged view of high potential 19 ACS Paragon Plus Environment


region at inset; (f) lithium ion diffusion coefficients of NCM and S-NCM-30min during the first charging process; (g) equivalent circuit for the impedance spectra.

To evaluate the effect of oxidation time on the cycling performance of NCM, the electrodes were cycled at 0.5 C between 3.0 V and 4.3 V, as shown in Figure 6a. After Na2S2O8 pre-oxidation, all samples deliver higher discharge capacities and cycling retentions than NCM after 100 cycles; detailed data are shown in Table S5. NCM exhibits a continual decline of discharge capacity, dropping to 171.6 mAh g-1 after 100 cycles with 93.7% capacity retention. Such a rapid capacity fade of NCM is ascribed to structural instability due to the phase transformation and the increased impedance as a result of surface side reactions during repeated charge-discharge cycling.36 However, Na2S2O8 pre-oxidation efficiently stabilizes the surface structure of NCM and produces remarkable cycling performance. Accordingly, S-NCM-15min, S-NCM-30min, and S-NCM-60min show discharge capacities of 178.3, 183.1, and 176.2 mAh g-1 after 100 cycles, corresponding to the capacity retentions of 97.3%, 99.0% and 96.2%, respectively. It is apparent that S-NCM-30min shows the optimized cycling performance. The obvious capacity reduction for S-NCM-15min and S-NCM-60min compared to S-NCM-30min is ascribed to oxidation insufficiency and the cracked particles due to the over-oxidation. When being charged to a higher voltage of 4.8 V (Figure S4), despite the large initial discharge capacities, all samples suffer a drastic capacity loss, and their discharge capacities after 100 cycles are even less than between 3.0 V and 4.3 V. However, the improved effect on cycling stability 20 ACS Paragon Plus Environment

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under a higher charge voltage is more predominant; detailed data are shown in Table S6. To further explain the effect of pre-oxidation, dQ/dV profiles of NCM and S-NCM-30min under various cycles are displayed in Figure 6b and c. At the first cycle, NCM exhibits a sharp oxidation peak at about 3.74 V, and secondary peaks appear at about 3.80, 4.01 and 4.21 V during charging. According to previous literature, these peaks arise from multiphase transitions of hexagonal to monoclinic (H1 → M), monoclinic to hexagonal (M → H2), and hexagonal to hexagonal (H2 → H3).25, 37 The first oxidation peak of NCM shifts to a lower voltage at the 25th cycle, hinting to a large activation energy in the initial cycle. During extended cycling, the redox peaks of NCM change into a more polarized state and shift farther apart. Especially, the redox peaks at 4.2 V get more irreversible. This H2 → H3 phase transition near charge-end causes an abrupt anisotropic shrinkage (or expansion during discharge), and then destroys the bulk structure of cathode materials, thus aggravating capacity fading by allowing the internal exposed surfaces to contact with electrolyte.38 After Na2S2O8 pre-oxidation, the lattice orientations of precursor are controlled to take on an ordered configuration (see Figure 3b), so that the primary particles of S-NCM-30min contract/expand uniformly rather than in random directions, and the adverse impact of phase transition is greatly weakened. As shown in Figure 6c, the potential shifts are remarkably relived in S-NCM-30min. We further investigated the rate capability of NCM and S-NCM-30min at various current densities from 0.1 C to 5 C between 3.0 and 4.3 V, and then back to 0.1 C, sustaining each rate for five cycles. Apparently, Figure 6d shows that S-NCM-30min 21 ACS Paragon Plus Environment


displays excellent rate performance, especially at high C-rates (from 1 C to 5 C), and its discharge capacities are 176.5, 162.0, and 130.3 mAh g-1, respectively. However, the discharge capacity of NCM decreases dramatically and drops to 110.8 mAh g-1 at 5C. NCM and S-NCM-30min retain 96.1% and 97.4% of their initial capacities, respectively, when the current density returns to 0.1 C. This implies that S-NCM-30min has a good capability to capacity recover after rate testing. The improved capacity retention and rate performance confirm that an optimum oxidation time of 30 minutes enables accelerated Li+ diffusion and maintains structural stability during repeated cycling. To explain the mechanism of electrochemical degradation, EIS was conducted on NCM and S-NCM-30min after different cycles charged to 4.3 V. As shown in Figure 6e and f, both Nyquist plots consist of two semicircles and a sloped line, corresponding to surface film resistance (Rsf), charge transfer resistance (Rct), and Warburg impedance (Zw), respectively.11 The equivalent circuit shown in Figure 5g is used to fit EIS spectra, and Table S8 shows the corresponding fitting data. The Rsf values of S-NCM-30min gradually decrease and stabilize at 13.85 Ω during cycling, indicating that a stable CEI layer is established after Na2S2O8 pre-oxidation. However, despite the slightly lower Rsf value of NCM after the initial charge process, its Rsf increases from 16.31 to 18.92 Ω after 100 cycles, which is larger than that of S-NCM-30min. This implies that the CEI layer on the NCM surface is unstable and tends to thicken as cycling proceeds. In addition, NCM exhibits larger variation of Rct values than S-NCM-30min upon cycling, suggesting that NCM suffers from serious 22 ACS Paragon Plus Environment

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structural degradation. The high overall resistance of NCM causes serious electrode polarization and poor electrochemical performance, which are effectively suppressed in S-NCM-30min. Therefore, it is imperative to do postmortem analysis of cathode materials.

Figure 6. Electrochemical performance at 25ºC between 3.0 and 4.3 V: (a) cycling performance of NCM and S-NCMs at 0.5 C after first five cycles at 0.1 C; differential capacity curves of (b) NCM and (c) S-NCM-30min; (d) rate capability of NCM and S-NCM-30min; Nyquist plots of (e) NCM and (f) S-NCM-30min after different 23 ACS Paragon Plus Environment


cycles. Figure 7 shows the ex-situ XRD patterns of NCM and S-NCM-30min electrodes before and after the charge-discharge process (black lines are electrodes before cycling, red lines are after 100 cycles between 3.0 V and 4.3 V, and blue lines are after 100 cycles between 3.0 V and 4.8 V). An Al current collector centered at 65.099º was used as a standard to calibrate the diffraction patterns. As shown in Figure 7a, the intensity of diffraction peaks for NCM declined after repeated delithiation/lithiation processes in both voltage ranges, indicating that its crystal structure suffers from destruction. The (006) peak of NCM shifts to a smaller degree, moving close to the (101) peak after 100 cycles in the potential range of 3.0-4.3 V, while its (012) peak shifts in the reverse direction. This peak-shifting phenomenon becomes more serious when the NCM electrode is cycled under high charge voltage of 4.8 V, as indicated by the blue strip. On the other hand, for S-NCM-30min electrode (Figure 7b), the intensity of all diffraction peaks and the position of (006)/(012) peaks hardly change even in 3.0-4.8 V. These results indicate that Na2S2O8 pre-oxidation successfully modifies the surface structure of NCM, and consequently, its structural integrity is guaranteed.


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Figure 7. Ex-situ XRD patterns of (a) NCM and (b) S-NCM-30min before and after 100 cycles at 25ºC under different charge cutoff voltage. 25 ACS Paragon Plus Environment


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To study the local structural variation, HRTEM, FFT and inverse fast Fourier transform (IFFT) analysis, which are complementary to XRD data, were conducted on the cycled electrodes. Figure 8a shows the remarkable structural change in NCM after 100 cycles between 3.0 V and 4.3 V. In bulk region b, the corresponding FFT (Figure _

8b2) indicates that the diffraction spots are attributed to the layered structure with R3 m space group. In transition region c, the (003)R diffraction spot of a layered phase _

_

vanishes, leaving the mixed R3m and Fm3m phases as determined by FFT (Figure 8c2). This mixed phases arise from the migration of the reduced Ni2+ ions to Li layers, which occupy the Li slabs.39 The structure change is exacerbated in the surface region d, and the original layered structure is extensively converted to rock-salt phase in the large-scale region from the surface. As shown in Figure 8d2, the diffraction spots of (003)R and (101)R of the layered structure disappear together, and the lattice fringes with a spacing of 0.24 nm suggest the formation of a (111)C crystal plane, which confirm the formation of NiO-like rock-salt phase. Such structural transformation causes the increased impedance and inferior cycling performance. An HRTEM image of the S-NCM-30min electrode is also displayed in Figure 9a. Different from the NCM (Figure 8a), all regions from the bulk to near-surface and even to the outermost _

surface, retain a well-ordered layered structure. The layered structure with R3m can be verified by FFTs in Figure 9b2-9c2. This result indicates that the structural stability of S-NCM-30min is much higher than that of NCM, and that is responsible for its better cycling performance.


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Figure 8. (a) HRTEM image of NCM after 100 cycles at 25ºC between 3.0 and 4.3 V; (b1, c1, and d1) IFFTs from region (b, c, and d) in (a); (b2, c2, and d2) FFTs from _

region (b, c, and d) in (a). The subscripts R and C represent R3m rhombohedral and _

Fm3m cubic phases, respectively. 27 ACS Paragon Plus Environment


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Figure 9. (a) HRTEM image of S-NCM-30min after 100 cycles at 25ºC between 3.0 and 4.3 V; (b1, c1, and d1) IFFTs from region (b, c, and d) in (a); (b2, c2, and d2) _

FFTs from region (b, c, and d) in (a). The subscript R represent R3m rhombohedral phase.


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In addition, we studied the structural variations of cycled electrodes under a higher charge voltage of 4.8 V, as shown in Figure S5. The surface structure of NCM performs numerous corrosion holes, which arise from the repeated lattice volume expansion/contraction and attack from destructive HF.40 Obviously, the structural variation of cycled NCM charged to 4.8 V is more serious than that of cycled NCM charged to 4.3 V. Charge-discharge in depth causes drastic structural instability, leading to inferior cycling retention, as shown in Figure S4. Conversely, we only observe the rock-salt phase in a very narrow region (1-7 nm) from the surface for S-NCM-30min, while the bulk maintains the original layered structure (Figure S5c). Conclusions In summary, a Na2S2O8 pre-oxidation strategy was employed to modify the surface structure

of

precursor

Ni0.8Co0.1Mn0.1(OH)2.

After

Na2S2O8

pre-oxidation,

LiNi0.8Co0.1Mn0.1O2 shows superior electrochemical performance. Remarkably, S-NCM-30min exhibits an improved rate capability with a discharge capacity of 130.3 mAh g-1 at 5 C, and remarkable cycling performance at 0.5 C, corresponding to 99.0% capacity retention after 100 cycles between 3.0 V and 4.3 V. When charged to a higher potential of 4.8 V, S-NCM-30min can even display a discharge capacity of 219.1 mAh g-1 and retain 93.8% capacity after 100 cycles, while the capacity retention of pristine NCM is only 84.5%. We believe the excellent cycling stability demonstrated in this work is the most preferable ever reported for LiNi0.8Co0.1Mn0.1O2 materials. This excellent electrochemical performance could be attributed to the well-ordered surface structure of precursor with Na2S2O8 pre-oxidation. On the one 29 ACS Paragon Plus Environment


hand, the ordered lattice orientations suppress the anisotropic shrinkage/expansion and stabilize the original layered structure of LiNi0.8Co0.1Mn0.1O2. On the other hand, the decreased content of Ni2+ inhibits the Li+/Ni2+ cation mixing and increases the kinetics of charge transfer and Li+ diffusion. The key findings of this work provide a new avenue for developing LIBs with a large discharge capacity, stable cycling performance, and high charge cutoff voltage required for the next-generation energy storage systems. Associated Content Supporting Information SEM images and photographs, XPS spectra of Mn 2p, cross section SEM image, plots of chronoamperometry and lnI, cycling performance, TEM images. Author Information Corresponding Author * Phone: +86-21- 31249125; Fax: +86-21- 31249125; E-mail: [email protected]. Notes The authors declare no competing financial interest. Acknowledgements The authors acknowledge funding supports from the 973 Program (No. 2013CB934103) and Science & Technology Commission of Shanghai Municipality (No. 08DZ2270500), China. References (1) Tarascon, M. A. a. J.-M. Building Better Batteries. Nature 2008, 451, 652-657. 30 ACS Paragon Plus Environment

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