Effect of Defects on Decay of Voltage and Capacity for Li[Li0.15Ni0

Apr 26, 2016 - The electrochemical test indicated that the decay of discharge capacity and working ... Advanced Functional Materials 2017 27 (7), 1604...
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Effect of Defects on Decay of Voltage and Capacity for Li[Li0.15Ni0.2Mn0.6]O2 Cathode Material Wuwei Yan, Yongning Liu,* Shengwu Guo, and Tao Jiang State Key Laboratory for Mechanical Behavior of Materials, School of Material Science and Engineering, Xi’an Jiaotong University, Xi’an 710049, PR China S Supporting Information *

ABSTRACT: Lithium-rich manganese metal layered oxides are very promising cathode materials for high-energy-density lithium-ion batteries, but improvement in voltage decay and capacity fade is a great challenge, which is mainly related to the structural instability or reconstruction of material’s surface. Defects, such as part lattice distortions, local cation disordering and atomic ununiformity, often aggravate the further structural changes upon cycling. In this paper, we found that PEG contributed to form better layered structure, well crystallinity, uniform composition and polyhedral nanoparticles for Li[Li0.15Ni0.2Mn0.6]O2 (LNMO). On the basis of the comparative trial, a mechanism of electronegativity difference is proposed to elucidate cation nonuniform distribution. Higher electronegativity of Ni (1.91) than Mn (1.55) show a stronger ability of attraction between Ni and O atoms, and then led to Ni atoms show stronger diffusion driving force toward particle surface to contact the rich O atoms during sintering in air. However, PEG polymer can form a better barrier for more O atoms to attract Ni and Mn atoms on particle surface so that facilitated a uniform distribution. The electrochemical test indicated that the decay of discharge capacity and working voltage was mitigated, which was identified by the result of HRTEM analysis that the initial less defect structure obviously retarded the phase transformation from the layered to spinel after 50 cycles. Therefore, defects are crucial for understanding the voltage fade and capacity decay, and the improvement of performance also demonstrates that designing optimum compositions and ordering atomic arrangements will contribute to stabilize the surface structure and restrain inherent phase transitions. KEYWORDS: lithium-ion batteries, cathode, lithium-rich material, combustion method, PEG

1. INTRODUCTION The market growth of smart phones, laptop computers, energy storage systems, and hybrid electric vehicles has attracted an increasing demand for energy storage from the lithium-ion batteries. Lithium-rich manganese metal layered oxides are now arousing more attention, because they deliver a high reversible capacity greater than 250 mAh/g,1−8 which exceeds that of commercial cathode materials with a discharge capacity of 140−180 mAh/g, such as LiCoO 2 , LiFePO 4 and Li(Ni,Mn,Co)O2. Usually, this kind of materials that can be written as two-component notation of xLi2MnO3·(1−x)LiMO2 (M = Mn, Ni, Co, etc.) are considered to be a two-phase layered structure or a single solid solution structure.9,10 Although their structures have been debated, these lithiumrich materials show a plateau of lithium and oxygen (net Li2O loss) release above 4.45 V during the first charge,11−15 which results in large irreversible capacity.16 Additionally, poor cycle ability,17 low rate capability,18 and decaying voltage19,20 are also shortcomings of these materials, which must be overcome for their commercial applications. Of particular note is that voltage fade and capacity decay are obvious in rich-lithium materials during cycling, which decreased the energy output of batteries. © XXXX American Chemical Society

The major reasons for the fade of voltage and capacity are the dissolution of transition metal ions,13 the layered-spinel phase conversion, and the diffusion energy barriers for Li+ within different planes. To solve the problem in the dissolution of transition metal ions, researchers used a large number of surface coating technologies to improve the surface stability of lithium-rich materials, such as representative AlF3,21 LiMn2O4,22 LiNi0.5Mn1.5O4,23 and graphene oxide.24 Moreover, Oh et al.25 reported that lithium-rich material coated with graphene oxide exhibited the stability of discharge voltage and capacity during 100 cycles by a hydrazine treatment. Nevertheless, it is difficult to improve the voltage fade in long-term cycling stability, because the surface coating can not prevent the Li2MnO3 phase transformation from the layered structure to spinel.26 It is found that the layered-spinel phase conversion is correlated with transition metal atomic distribution. Recently, Gu et al.27 discovered that segregation of Ni ions was present Received: January 20, 2016 Accepted: April 26, 2016

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DOI: 10.1021/acsami.6b00763 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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surfactants play an important role for morphology and crystallinity during the process of the precursor synthesis due to their multiple physical and chemical properties. Although some works have been reported that defects of nonuniform atomic distribution played a very important effect on the decay of voltage and capacity, little work has been shown how to reduce the defects in material preparation by introducing surfactant. In this paper, a cobalt-free low Li/M ratio material, Li[Li0.15Ni0.2M0.6]O2 material was synthesized by a simple combustion method. During the synthesis, PEG as a surfactant and a dispersant,41 was used to tune the morphology and structure of particles. It is found that PEG had a great influence on crystallinity of lithium-rich layered oxide synthesized by the combustion, so that the electrochemical properties were improved much in voltage decay and capacity retention. These studies also revealed that microcosmic defects in lithium-rich materials obviously influenced the electrochemical behavior and emphasized the fundamental mechanism in atomic level uniformity should be focused on.

on surface regions in the lithium-rich material, which could slow down lithium ions diffusion, so decreased battery capacity and deteriorated cyclic behavior. Boulineau et al.28 and Yan et al.29 also demonstrated that a segregation of nickel and manganese in particle surface decreased the working voltage upon cycling of this material. Zheng et al.8 obtained a uniform Ni distribution in the lithium-rich material so that the layeredspinel phase conversion was mitigated and voltage fade was suppressed. On the other hand, the optimized compositions of lithium-rich materials can also facilitate to alleviate the layeredspinel phase transformation. For instance, Shunmugasundaram et al.16 found that lithium-rich materials with low Li-to-total metal ratios (Li/M) exhibited a smaller first-cycle irreversible capacity loss of 4.0% originating from the presence of metal site vacancies and no Li atoms in the transition metal (TM) layer. Sathiya M. et al. found that the substitution of larger and less electropositive cation Sn4+ for Ru4+ in Li2Ru1−ySnyO3 facilitated maintaining stable structure and decreased voltage fade.30,31 It seemed as if the inherent phenomenon of capacity decay and voltage fade over long-term charge and discharge was associated with an unstable electrode interface and undesirable phase transitions, nevertheless, to the best of our knowledge, the fundamental understanding and intensive study of atomicscale composition and local structure in primary materials related with the phenomenon was less revealed. Defects, such as part lattice distortions, local cation disordering, nanoscale phase sepraration, are thought to be crucial factor for the long-term stability in the pristine lithiumrich mateials. For example, Jarvis at al.32 revealed that planar defects in lithium-rich material were formed during the transition of the transition metal layer from a disordered R3̅ state to a lithium- ordered C2/m state, and the synthesis process was the greatest factor for the planar defect formation. Boulineau at al.33 revealed that disorder resulting from lithium/ manganese exchange in LiMnO3 decreased with increase of the synthesis temperature, but the synthesis of a material without stacking faults was not possible by changes of temperature. Fell at al.34 observed in TEM that generation of the oxygen vacancy and activation of LiMnO3 caused the formation of crystal defects, which may lead to the first irreversible capacity and poor rate capability. Bareño at al.35 suggested that the local atomic structure, like strain fields, steric hindrances, influenced Li-ion removal in lithium-rich materials and their electrochemical performance during cycling. Zheng at al.8 synthesized fewer defects with Ni segregation in the lithium-rich material by mixing PVP and EG into the precursor via hydrothermal assisted method, so that the stabile energy density was maintained. The particle geometry and crystallinity of the products are highly dependent on the addition of surfactant during preparation of cathode materials, such as ethylene glycol (EG),36 hexadecyltrimethylammonium bromide (CTAB),37 citric acid (CA), 38 and aminopropyl-trimethoxysilane (APTMS),39 which are introduced and serve as precursors. For example, Wei at al. 5 synthesized single crystalline Li(Li0.17Ni0.25Mn0.58)O2 hexagonal nanobricks with a high percentage of {010} facets by introducing oxalic acid and CA into precursor, which exhibited excellent high rate performance for a cathode of lithium-ion batteries. Jouybari et al.40 studied three surfactants, such as triethanolamine (TEA), CA, and oxalic acid during the synthesis of LiNi0.8Co0.2O2 nanoparticles, they found TEA resulted in the smallest size distribution in range of 50−80 nm and the highest crystallinity. In this sense,

2. EXPERIMENTAL SECTION 2.1. Material Preparation. Li[Li0.15Ni0.2M0.6]O2 powders were prepared by a combustion method. A stoichiometric ratio of raw materials, including lithium acetate (Li(CH3COO)·2H2O) with 5 wt % excess, nickel nitrate (Ni(NO3)2·6H2O) and manganese acetate (Mn(CH3COO)2·4H2O), was dissolved in ethanol solvent. PEG400 was added into the solution until a transparent solution was formed during continuous magnetic stirring. The molar ratio of (Li+Mn +Ni):PEG400 was 1:1. After slowly evaporating at 80 °C for 12 h, a viscous gel was obtained and it was transferred into a corundum crucible, which was put into tube furnace, and then combusted at 400 °C for 60 min in air to eliminate organic substances. The combusted precursor was ground in a carnelian mortar, then calcinated at 900 °C for 15 h in air, and finally cooled to room temperature. The materials prepared from the above procedure in ethanol only, and ethanol with PEG 400 additive, were named as E-sample, and EP-sample, respectively. 2.2. Material Characterization. The powder morphology was observed by scanning electron microscope (SEM) JSM-6390. The compositions of synthesized samples were measured by inductively coupled plasma atomic emission spectroscopy (ICP-AES, Varian 715ES, USA). The specific surface area of the particle was measured by the BET method. Powder X-ray diffraction (XRD) was measured with instrument (UltimaIV-185, Rigaku) using Cu−Kα radiation (λ = 1.5406 Å) at 40 kV and 40 mA. XRD data was acquired in the 2θ range of 10−80° at an interval of 0.02°. Thermogravimetric (TG) and differential thermal analyses (DTA) were conducted using a Seiko Dual TG/DTA 320 Analyzer, and samples, which were obtained from dry gel, were heated to 1000 °C at a rate of 10 °C/min in air. The powder particles were suspended on a copper grid with lacey carbon for TEM microscope (JEM-2100F). 2.3. Electrochemical Test. Cathode electrode was fabricated by mixing 80 wt % Li[Li0.15Ni0.2M0.6]O2 material, 10 wt % acetylene black and 10 wt % PVDF. The mixture was dissolved in NMP solvent and stirred to form homogeneous slurry, which was then coated onto an Al foil current collector and dried at 120 °C for 4−8 h. CR2025-type coin cells were assembled with the cathode electrode, metallic lithium foil as counter electrode, Cellgard 2300 as separator and a solution of 1 M LiPF6 as electrolyte, which was dissolved in a mixture of ethylene carbonate (EC)/dimethyl carbonate (DMC)/diethyl carbonate (DEC) (1:1:1, Vol%) in an argon-filled glovebox (4.9),43 I(003)/I(104) (>1.2),44 the splitting degree of the (006)/(102) and (108)/(110) peaks.45 Table 1 lists the values of a, c, c/a and intensity ratios of (003)/(104) peak from two samples, both of them possess the high c/a ratios and the big I(003)/I(104) peak intensity ratios. By contrast, the double peaks of (006)/(012) and (018)/(110) in the EP-sample show more distinct splitting than those of E-sample, which reveals a higher crystallinity and a well-ordered layered structure. The characteristic superstructure of EP-sample at 20−25° is more obvious compared with that of E-sample, it may be associated with a higher degree of LiMn6 cation ordering in the transition metal layer,9 lower amount of stacking defects,10 and so on. Thus, the XRD patterns suggest that PEG shows the ability to improve intrinsic degree of crystallite in LNMO. Figure 3 provides the distribution of O, Mn, and Ni elements on the LNMO particle surfaces (E- and EP-samples). On the basis of a liner sweep of one particle on the insertion photos in Figure 3a, b, respectively, more sharp edge and corner morphology of particle for EP-sample indicate its crystallinity is higher than that E-sample. Moreover, clearly, the Ni/Mn intensity ratio in EP-sample does not vary from outside to center on the particle surface. By contrast, the Ni/Mn intensity ratio on both ends of particle surface in E-sample is different from that in central region. Similarly, some researches also found that Ni ions segregated on surface of particles,8,27,29,46 which is ascribed to the surface energies and the diffusion kinetics of transition-metal ions in the layered material.27 It is reported that ethylene glycol (EG) exhibited a strong chelating ability for some transition metal ions.47 Herein, the PEG molecule has a similar structure and a longer chain compared with EG molecule, so the oxygen atoms from PEG have the same ability to coordinate with metal cations. Moreover, Ni atoms possess higher electronegativity (XNi = 1.91) than Mn atoms (XMn = 1.55), and show a stronger ability to capture O atoms when sintering in air. Meanwhile, plenty of the absorbed O atoms surrounded the LNMO particles, which has a stronger attraction to Ni atoms than to Mn atoms, so Ni atoms diffused toward surface and resulted in the higher content at the surface than it at the center for E-sample as shown in Figure 4a. More Ni ions occupy the surface position for stronger action with oxygen atoms, less position was remained for Mn ions to occupy, so Mn content is lower at surface and higher in the central region. When PEG was introduced for EP-sample, limited O atoms were provide by PEG, which isolated the attraction between Ni atoms and O atoms from air, and prevented more oxygen from air diffusing into and stopped Ni

Table 1 shows that the chemical compositions of two samples are close to the theoretical formula, which suggested that the LNMO materials were successfully prepared by the combustion method. Therefore, adding PEG into precursor can suppress the fast decompositions from nitrate and acetate, and keeps the precursor stable during the process of decomposition. The particles of E-sample in Figure 2a are aggregated and exhibit irregular features, but EP-sample in Figure 2b shows a relatively uniform shape and a primary polyhedron nanoparticle. Moreover, the naonoparticles of EP-sample show smooth surface and homogeneous dimension range between 200−300 nm in Figure 2c. The specific surface area of EPsample is 5.3 m2 g−1 by BET test, which is nearly 3 times that of E-sample (1.8 m2 g−1). The addition of PEG into the solution created a quite viscous environment, once the colloidal particles are formed, the long PEG (HO-(CH2CH2O)nH) chain structure can bond with metal ions by the chelation of HOgroup (Figure S1). As a result, −O− and −OH groups can connect with the transition metal and alkyl chain toward the ethanol. After evaporation of ethanol, PEG formed a protective layer of polymer film on the colloidal particle surface and surrounded it, and finally led to reduction of the particle agglomeration because of the effect of steric hindrance. To obtain good morphology and crystallinity, we also investigated the influence from different concentrations of (Li +Mn+Ni)/PEG (0:1, 0.5:1, 1:1 and 1.5:1), and different molecular weights of PEGs (PEG200, PEG400, PEG600 and PEG800) in precursor, respectively (Figures S3−S6). As a result, the proper ratio (PEG400/(Li+Mn+Ni) = 1:1) can very effectively coated the colloidal particle surface and surrounded it, so that exhibited smooth surface and uniform dimension distribution (Figure S3c). On the basis of c/a ratios, I(003)/I(104) peak intensity ratios, and the double peaks of (006)/(012) and (018)/(110), LNMO product from molecular weight of PEG400 and concentration of PEG/(Li+Mn+Ni) (1:1), indicated a highest crystallinity and a well-ordered layered structure (Figure S6). Figure 2d presents the XRD patterns of E- and EP-samples. All peaks can be indexed as the layered hexagonal α-NaFeO2 structure with R3̅m space group, except for these weak peaks in the range of 20−25°, which can be indexed to (020), (110), (11̅ 1) lattice planes of a Li2MnO3-type unit cell with lower symmetry C2/m space group.42 It has been reported that the electrochemical properties of layered oxide materials were affected by the degree of cation ordering structure, which can be indicated from XRD patterns with the lattice constant ratios D

DOI: 10.1021/acsami.6b00763 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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when heated in 400 °C for 1 h (Figure S2 and Table S1), which proved that a higher crystalline layered structure already appeared in 400 °C. It disclosed that primary product of low temperature decomposition influenced the crystallinity and a layered structure of the final product after sintered in high temperature. In addition, an independent study revealed a similar phenomenon that longer calcination time was easy to form Ni segregation in LNMO particle surface.27 Therefore, the electronegativity difference can explain the phenomenon of Ni segregation well and the effect of PEG further supported the explanation. More important, the original product with uniform atomic distribution played a vital role for the final crystallinity. TEM images and diffraction patterns of particles from E- and EP-sample are displayed in Figure 5. The six weak dots in the SAED pattern of E-sample in Figure 5a can be well-indexed with planes of the C/2m symmetry, namely, a layered [001] Li2MnO3 phase. In Figure 5c, the lattice fringe image is obtained by the Fast Fourier Transformation (FFT) for HRTEM image (Figure 5b), which obviously indicates the presence of lattice distortion or planar defect as shown in red circles. Not surprisingly, these defects were coincidently reflected in the weak intensity of XRD pattern in the range of 20−25° (Figure 2d) and inhomogeneous atomic distribution (Figure 3a) in E-sample. In Figure 5d, although the LNMO material was ground to powders for effective observation, a large number of polyhedral particles can be seen from the image, which is consistent with the corresponding SEM findings. In addition, Figure 5d also presents the SAED pattern along the [001] zone axis with C/2m space group, proving that these polyhedral particles also has a monoclinic Li2MnO3 layered structure. The FFT pattern in Figure 5f of the

Figure 4. Schematic illustration of formation of (a) Ni segregation without PEG and (b) Ni homogeneity with PEG on LNMO particle surface.

and Mn atoms migrating toward particle surface, so that it made sure the stoichiometric ratio of O, Ni, and Mn during calcination process as shown in Figure 4b. Therefore, more uniform composition and good crystallinity structure are formed in the EP-sample. Generally, higher temperature and longer time is an effective way to make materials uniform in composition and crystal structure. However, the composition and the crystal structure in nanosized particles did not show uniform for E-sample even though it was treated at 900 °C for 15 h because of the formation of lower crystallinity in 400 °C. By contrast, it was observed that the EP-sample heated in 400 °C for 1 h exhibits higher c/a and I(003)/I(104) than E-sample

Figure 5. (a) TEM image and SAED pattern of the E-sample, corresponding (b) HRTEM image and (c) FFT pattern. (d) TEM image and SAED pattern of the EP-sample, corresponding (e) HRTEM image and (f) FFT pattern. E

DOI: 10.1021/acsami.6b00763 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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Figure 6. First discharge curves of (a) E- and (b) EP-samples at different rates between 2.0−4.8 V. (c) Comparison of rate discharge capabilities. (d) Cycling performance at discharge rate of 0.3C rate. Discharge curves of (e) E- and (f) EP-samples for d.

Obviously, EP-sample exhibits the more excellent rate stability than E-sample at high-rate discharge. In particular, compared to E-sample, the discharge capacities of EP-sample show a gradually increasing tendency and reach stabilization during initial 5 cycles. Usually, the discharge capacity of the lithiumrich transition metal oxides mainly depended on the degree of the activated Li2MnO3 phase (Li2MnO3 → Li2O + MnO2) because more Mn ions participated in the electrochemical redox reactions during the subsequent discharge/charge cycles.48 Consequently, the initial increasing discharge capacities suggested that the high proportion of active surface areas containing Li2MnO3 phase was activated in EP-sample. To avoid the electrolyte decomposition at high voltage, cells from E- and EP-samples were measured during 200 cycles between 2.0−4.6 V under 0.3 C rate as shown in Figure 6d−f. Herein, the charging rate of 0.1 C was applied to prevent the polarization so that fully charging was achieved almost. Compared to E-sample, the cell of EP-sample delivered a higher reversible capacity of about 204−195 mA hg−1, and the

HRTEM image (Figure 5e) shows the clearly crystal lattice stripe image which is absence of defect and distortion, suggesting EP-sample possesses a high-quality crystal structure. In conclusion, this structural variation discloses that the synthesis condition affects the presence of nanoscale defects in the layered LNMO structure, moreover, the low defect concentration plays a significant role for the quick and frequent electron transport and lithium-ion insertion/extraction within this material. 3.2. Electrochemical Performances. The discharge performance of E- and EP-samples at high rates is presented in Figure 6a−c. The cells were initially operated for one cycle between 2.0−4.8 V at a 0.1 C rate in order for sufficient activation, and then discharged between 2.0−4.8 V at 0.3, 0.5, 1, 3, 5, and 7 C rates. The corresponding first discharge capacities of EP-sample in Figure 6a are 231, 221, 213, 198, 165, and 139 mAh g−1, respectively, which are superior to that of E-sample in Figure 6b. Figure 6c displays the stepped ratecycling performance of two samples during 35 cycles. F

DOI: 10.1021/acsami.6b00763 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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Figure 7. (a) HRTEM image of the E-sample, corresponding (b) FFT pattern and (c) SAED pattern. (d) HRTEM image of the EP-sample, corresponding (e) FFT pattern and (f) SAED pattern. Two samples were cycled for 50 cycles between 2.0 and 4.8 V.

fine capacity retention of 95.6% after 200 cycles as well as the Coulombic efficiency of nearly 100% was obtained, which suggested the absence of parasitic reactions or side reactions between electrode and electrolyte. The better high-rate cyclic stability and capacity retention of EP-sample can be associated with the following reasons, the well-crystallized layer structure, a nanosized polyhedral shape, and the primary nanoparticles with high specific surface area. The high layer structure can keep the phase stable and facilitate the fast Li+ diffusion inside the crystal structure so that exhibits superior electrochemical behavior. Furthermore, the nanosized polyhedral shape can shorten distance for Li+ ions diffusion and provide a large active material/electrolyte contact area, which contributes to the high rate of discharge and cycle ability. More importantly, the high specific surface area structure of these nanoparticles can offer the sufficient contact between electrolyte and inner particles, and provide an extensive Li+ diffusion pathway, which may be associated with high rate capability. Meanwhile, the primary nanoparticles can sustain the volume change and the structural strain during the repeated Li+ insertion/extraction. In Figure 6d−f, it proves that EP-sample displays less voltage fade than E-sample as shown in red arrows. The phenomenon of voltage fade occurs in many layered oxides during charge/ discharge process, and it is affected by many factors, including cycling voltage window, temperature and cycle life. Except the side reactions between the electrode and electrolyte at high voltages, which can cause a deterioration of electrode/ electrolyte interface and led to voltage fade, the main intrinsic reason on voltage fade of lithium-rich materials is related with layered-to-spinel phase transformation as reported in literatures.8,49 The formation of the irreversible spinel phase on the surface of the material are caused by the surface transition metal ions migrating into Li sites, combined with tetrahedral Li ions that migrated from the octahedral site after oxygen release at high voltage.50 Nevertheless, the study revealed that a uniform atomic distribution could greatly suppress the phase conversion

from the layered to spinel. Such as, Zheng et al.8 proved that a uniform distribution of Ni ions on the surface of nanoparticles facilitated to alleviate voltage decay and maintain cyclic stability for batteries. Therefore, the result suggested that the EP-sample displayed less degradation of voltage fade and retained a more stable cyclic behavior than E-sample during cycling, it is mainly attributed to the relatively uniform distribution of transition metal ions as mentioned before. 3.3. Phase Transformation Analyses. Figure 7 shows microstructures, corresponding FFT images and SAED patterns of LNMO particles from E- and EP-samples after 50 charge/ discharge cycles. To distinguish components more clearly in this local distribution in Figure 7a, the inverse FFT was used to examine this distribution. Obviously, a great number of lattice distortions (A2) along the region (A1) were observed in Esamples as shown in Figure 7b. Based on SAED patterns from FFT imagines in Figure 7c, two different phases could be obtained, their corresponding indexing were identified as the original layered structure of the [001] Li2MnO3 phase and the distorted spinel-like [1̅12] LiMn2O4 phase. By contrast, Figure 7d and e show the well-organized crystal lattice stripe image (B1), and less distorted lattice fringe (B2). As can be seen from Figure 7f, the two diffraction patterns in B1 and B2 are indexed to a [001] Li2MnO3 and a distorted [110] LiMn2O4, respectively. The results from HRTEM image and SAED pattern revealed that the original layered structure partly transformed into a spinel structure in both of E- and EPsamples, but the EP-sample exhibited less layered-spinel phase transformation. On the basis of the above observation, an explanation was presented as follows. More layered to spinel phase transition occurred in E-sample, in which initial crystal structure is not perfect, and there are a lot of defects and distortions in this material. These crystal defects obviously promoted the phase transition by providing more nucleation sites for the formation of new phase embryos. Moreover, during the electrochemical activation, these embryos are easy to transform into a new G

DOI: 10.1021/acsami.6b00763 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

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phase. Therefore, controlling the low percentage of defects and optimizing the selective morphology feature play a crucial role for the lithium-ion transformation and electron transport in these materials, which may assist the improvement of electrochemical behavior for lithium-ion batteries with favorable energy density. Moreover, the electrochemical test result and phase analysis gives a more comprehensive structure−property relationships of lithium-rich materials, and selecting an appropriate surfactant is very significant to achieve nanoscale structural homogenities and inherent high crystallinities of the pristine materials themselves.

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.6b00763. Schematic illustration of PEG structure in ethanol by the coordination with transition metal ions, the detailed structure information on the samples after calcination at 400 °C, morphology, and crystallinity influenced by varying the concentration of PEG and the molecular weights of PEG (PDF)

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4. CONCLUSIONS In this study, the Li[Li0.15Ni0.2Mn0.6]O2 cathode material was synthesized in ethanol with PEG by a combustion method. The results showed that PEG not only facilitated the formation of primary polyhedral nanoparticles with a size of 200−300 nm, but also benefited to form well crystallized structure and uniform composition. The electronegativity difference is the reason for the nonuniform distribution of Ni and Mn elements on the LMNO particle surface. The higher electronegativity (1.91) of Ni shows stronger ability of capturing oxygen than Mn (1.55), so Ni atom has a higher diffusion driving force to migrate to particle surface. However, PEG makes a better barrier for more O atoms to attract Ni atoms in particles, which prevented more oxygen from air diffusing into and Ni atoms migrating toward particle surface. The LMNO material exhibits less capacity decay and voltage fade, better initial discharge capacity of 139 mAh g−1 at a 7 C rate and maintains capacity retention of 95.6% at discharge of 0.3 C rate after 200 cycles. The less defect layered structure is in favor of the fast and efficient transport for lithium ions and electron in the LNMO cathode material during lithiation/delithiation. Characterization of the initial defects explained the decay of capacity and voltage, and it make us acutely aware of the need for further reducing the inherent ununiformity of these materials by controlling synthesis with appropriate chemical and physical method.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +86 29 8266 4602. Fax: +86 29 8266 3453. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the China Postdoctoral Science Foundation (no. 2012M521760) and the Fundamental Research Funds for the Central Universities (xjj2014052). H

DOI: 10.1021/acsami.6b00763 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX

Research Article

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DOI: 10.1021/acsami.6b00763 ACS Appl. Mater. Interfaces XXXX, XXX, XXX−XXX