Multi-Functional Surface Engineering for Li-Excess Layered Cathode

Jan 22, 2016 - Moreover, a high discharge capacity of 78 mAh g–1 could be obtained at 25 C. The exothermic temperature of the fully charged electrod...
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Multi-Functional Surface Engineering for Li-Excess Layered Cathode Material Targeting Excellent Electrochemical and Thermal Safety Properties Xiaofei Bian,† Qiang Fu,† Qiang Pang,† Yu Gao,† Yingjin Wei,*,† Bo Zou,‡ Fei Du,*,† and Gang Chen†,‡ †

Key Laboratory of Physics and Technology for Advanced Batteries (Ministry of Education), College of Physics, Jilin University, Changchun 130012, People’s Republic of China ‡ State Key Laboratory of Superhard Materials, Jilin University, Changchun 130012, People’s Republic of China S Supporting Information *

ABSTRACT: The Li(Li0.18Ni0.15Co0.15Mn0.52)O2 cathode material is modified by a Li4M5O12-like heterostructure and a BiOF surface layer. The interfacial heterostructure triggers the layered-to-Li4M5O12 transformation of the material which is different from the layered-to-LiMn2O4 transformation of the pristine Li(Li0.18Ni0.15Co0.15Mn0.52)O2. This Li4M5O12-like transformation helps the material to keep high working voltage, long cycle life and excellent rate capability. Mass spectrometry, in situ X-ray diffraction and transmission electron microscope show that the Li4M5O12-like phase prohibits oxygen release from the material bulk at elevated temperatures. In addition, the BiOF coating layer protects the material from harmful side reactions with the electrolyte. These advantages significantly improve the electrochemical performance of Li(Li0.18Ni0.15Co0.15Mn0.52)O2. The material shows a discharge capacity of 292 mAh g−1 at 0.2 C with capacity retention of 92% after 100 cycles. Moreover, a high discharge capacity of 78 mAh g−1 could be obtained at 25 C. The exothermic temperature of the fully charged electrode is elevated from 203 to 261 °C with 50% reduction of the total thermal release, highlighting excellent thermal safety of the material. KEYWORDS: lithium ion battery, Li-excess layered oxide, surface modification, electrochemical properties, thermal safety

1. INTRODUCTION Demands for lithium ion batteries have expanded from portable electronics to electric vehicles and stationary energy storage. The traditional LiCoO2 cathode cannot meet the requirements of these large-format lithium ion batteries because of its high cost, low energy density and poor safety at the overcharged state (>4.4 V). In the research of new cathode materials, Liexcess layered oxides, xLi2MnO3−(1 − x)LiMO2 or Li[LixM1−x]O2 (hereafter denoted as LLMO, M = Co, Ni, Mn), have attracted great attention since their first introduction in 2000s.1−4 LLMO is a highly economical cathode material because of its rich component in Mn and only a little fraction of Co. More remarkably, LLMOs exhibit rather high specific energy of about 1000 Wh kg−1 owing to their large discharge capacity of over 250 mAh g−1 coupled with an average discharge voltage of about 3.5 V. Due to these advantages, LLMO has been regarded as a promising cathode material for the next generation lithium ion batteries. The large discharge capacity of LLMO is originated from the initial oxygen loss from the material lattice during the first charge. This process is accompanied by a layered-to-LiMn2O4 phase transformation at the surface region of the material which results in a large initial irreversible capacity.5,6 The concurrent migration of varies ions (including O2−, Li+, and Mn+) during © 2016 American Chemical Society

this process retard the diffusion of Li ions. For instance, the lithium diffusion coefficients of Li1.13Ni0.3Mn0.57O2 decreased from 10−14 to 10−17 cm2 s−1 during the oxygen loss process. Moreover, the activation energy for charge transfer reactions also rapidly increased to a much larger value.7 These sluggish electrochemical kinetic properties are responsible for the poor rate capability of LLMOs. Once the LiMn2O4 spinel is formed after the first cycle, it would gradually grow inward the material with charge−discharge cycling. The material bulk then transforms to a nano composite where small LiMn2 O4 crystallites are embedded in a layered LiMO2 framework.8−10 As a result, the average discharge voltage of the electrode gradually decreases because of the lower discharge voltage of Li1+xMn2O4. In addition, the capacity fading of the electrode is accelerated due to dissolution of Mn and the Jahn−Teller effect of Mn3+ ions. As is known, the high reversible capacity of LLMOs can be only achieved by charging the materials above 4.6 V. This voltage is too high for most LiPF6-based electrolytes. Thus, the charge−discharge cycling of LLMOs is always accompanied by Received: November 19, 2015 Accepted: January 22, 2016 Published: January 22, 2016 3308

DOI: 10.1021/acsami.5b11199 ACS Appl. Mater. Interfaces 2016, 8, 3308−3318

Research Article

ACS Applied Materials & Interfaces

thermal runaway of LLMO. The thermal safety temperature is significantly elevated with 50% reduction of the total thermal release.

complex side reactions which result in gaseous compounds and a solid electrolyte interface (SEI) film.11−13 This brings serious safety concerns for the batteries because the side reactions would become violent at elevated temperatures. In addition, lattice oxygen tends to release from the material at high temperatures leading to thermal runaway of the battery cell. Different studies have shown that the onset thermal release temperature for LLMOs cathode is as low as 200 °C and the amount of released heat is even more than 1000 J g−1.14−16 Many efforts have been done to improve the electrochemical properties of LLMOs such as by cation doping,17−20 surface coating12,14,21−27 and preparing nano materials.28−30 In spite of these efforts, some shortcomings of LLMOs such as voltage decay, capacity fading and poor rate capability have not been effectively resolved. Alternately, recent proceedings of surface treatment have shown applausive improvements in the electrochemical properties of LLMOs. For example, it has shown that a NiO cubic phase could form on 0.4Li2MnO3· 0.6LiNi1/3Co1/3Mn1/3O2 by hydrazine treatment.25 This NiO layer can act as a surface stabilizer which improved the cycle stability of the material. But too thick of such a surface layer would depress the rate capability due to the poor conductivity of NiO. Song et al. modified the surface of Li(Li0.2Mn0.54Ni0.13Co0.13)O2 by carbon black which resulted in partial phase transformation from Li2MnO3 to LiMn2O4.31 Similar phase transformation was also obtained by AlF3 coating as reported by Sun et al.27 The surface LiMn2O4 spinel had high electronic/ionic conductivity so that it could improve the rate capability of the material. However, the material bulk still transformed to a LiMn2O4-like phase after several tens of cycles, resulting in continuous capacity fading and voltage decay. One can see from the above analysis that the discharge capacities, rate capability, and cycle stability of LLMOs are closely related with their ultimate crystal structures. In addition, it should be noted that the thermal safety of delithiated LLMOs initiates at the electrode surface and propagate into the material bulk. Therefore, in order to fundamentally improve the electrochemical performance of LLMOs, the following two issues should be well addressed: (1) how to control the phase transformations of LLMOs and (2) how to optimize the surface properties of LLMOs. Taking account of these concerns, herein we report a new surface engineering method for the LLMO cathode materials. The successfully prepared material is represented as LLMO@Li4M5O12@BiOF of which LLMO is a high-capacity cathode material with formula of Li(Li0.18Ni0.15Co0.15Mn0.52)O2. The LLMO cathode is in conjunction with a Li4M5O12-like interfacial heterostructure and a BiOF surface coating layer. Using this heterostructure, we successfully overcome the layered-to-LiMn2O4 transformation of LLMO. This heterostructure triggers LLMO to slowly transform to a Li4M5O12-like spinel during cycling which helps the material to keep high working voltage, long cycle life and excellent rate capability. In addition, the side reactions of the electrode and the electrolyte are effectively suppressed. The synergetic effects of the Li4M5O12-like spinel heterostructure and the BiOF coating layer lead to excellent capacity retention, low voltage decay and high rate capability. Moreover, we carried out a systematic work to study the thermal safety properties of LLMO. On the basis of combinative analysis, we show that the thermal runaway of LLMO is related with the oxygen release and simultaneous phase transformations of the material. Our new surface modification method inhibits the

2. EXPERIMENTAL SECTION 2.1. Materials Preparation. For preparation of Li(Li0.18Ni0.15Co0.15Mn0.52)O2, stoichiometric Li2CO3, Co(CH3COO)2, Ni(CH3COO)2, and Mn(CH3COO)2 were dissolved in deionized water with vigorous stir. Then citric acid was added into the solution with a molar ratio of citric acid: Li = 2:1. Ammonia−water was used to keep the pH value of the solution at about 7.3. The viscous solution was constantly stirred at 50 °C for 15 h and then dried at 120 °C in a vacuum oven. The obtained gel was pretreated at 450 °C for 5 h. Then, the precursor was pressed into pellet and heat treated at 900 °C for 12 h. After the annealing process, the pellet was quenched in liquid N2 and grounded to a Li(Li0.18Ni0.15Co0.15Mn0.52)O2 powder by hand milling. For preparation of LLMO@Li4M5O12@BiOF, 400 mg Li(Li0.18Ni0.15Co0.15Mn0.52)O2 was mixed with 4 mg Bi2O3, and then fluoridized in saturated NH4F solution for 3 h. The fluoridated powder was collected by centrifugation and washed with deionized water. Then the material was annealed at 450 °C for 3 h to obtain the LLMO@Li4M5O12@BiOF product. 2.2. Materials Characterizations. X-ray diffraction (XRD) was performed on a Rigaku AXS D8 diffractometer with Cu Kα radiation. The morphology of the materials was observed by a JSM-6700F field emission scanning electron microscope (FESEM). High-resolution transmission electron microscope (HRTEM) was studied on a FEI Tecnai G2 F20 S-TWIN coupled with a Bruker AXS energy dispersive spectroscopy (EDS). Raman scattering was performed on a Thermo Scientific FT-Raman using Nd-line laser. X-ray photoelectron spectroscopy (XPS) was collected on an ESCALAB spectrometer using Mg Kα light source. Fourier transform infrared spectroscopy (FTIR) was measured on a Thermo Scientific Nicolet 6700 spectrometer using the KBr method. Differential scanning calorimetry (DSC) was carried out on a TA-Q2000 thermal analyzer with a heating rate of 10 °C min−1. Mass spectrometry (MS) was performed on a NETZSCH STA499F3-QMS403D/Bruker V70 spectrometer. All the thermal analysis experiments were performed for the collected electrode composite after the first charge without removing the residual electrolyte. 2.3. Electrochemical Experiments. Electrochemical experiments were performed on 2032-type coin cells using metallic Li as the anode. The cathode was composed of 75 wt % active material, 15 wt % carbon black conductive additive and 10 wt % poly vinylidenefluoride (PVDF) binder which were pasted on an Al current collector. The cathode was 8 × 8 mm in size and contained about 2.2 mg active material. The cathode and anode were separated by a Celgard 2320 membrane. The electrolyte was a 1 mol·L−1 lithium hexafluoro-phosphate (LiPF6) solution dissolved in ethylene carbonate (EC), dimethyl carbonate (DMC) and ethylmethyl carbonate (EMC) (EC/DMC/EMC = 1:1:8 by v/v ratio). Galvanostatic charge−discharge was performed on a Land-2100 automatic battery tester. Electrochemical impedance spectroscopy (EIS) and galvanostatic intermittent titration technique (GITT) was studied on a Bio-Logic VSP multichannel potentiostaticgalvanostatic testing system. The EIS data were obtained by applying an ac voltage of 5 mV in the frequency range from 1 MHz to 1 mHz. For the GITT measurement, the battery was discharged with a current flux of 32 mA g−1 for 0.5 h, followed by an open circuit stand for 4 h to reach the quasi-equilibrium state.

3. RESULTS AND DISCUSSION 3.1. Structural and Physical Properties. Shown in Figure 1 are the X-ray diffraction patterns of the pristine LLMO and LLMO@Li4M5O12@BiOF samples. The pristine LLMO displays a typical diffraction pattern of Li-excess layered oxides. Most of the diffraction peaks are attributed to the rhombohedral LiM′O2 (M′ = Li0.18Ni0.15Co0.15Mn0.52) phase with R3̅m symmetry. The small peaks between 20 and 25° are 3309

DOI: 10.1021/acsami.5b11199 ACS Appl. Mater. Interfaces 2016, 8, 3308−3318

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Figure 1. XRD patterns of the LLMO and LLMO@Li4M5O12@BiOF samples.

due to the LiMn6 ordering of monoclinic Li2MnO3. The lattice parameters of the material are calculated based on the R3̅m structure, which are a = 2.8531 Å and c = 14.2320 Å. LLMO@ Li4M5O12@BiOF shows a similar XRD pattern as that of LLMO. But, a small shoulder appears at the right side of the (003) peak and the left side of the (101) peak, respectively. These features suggest formation of a spinel phase in LLMO as reported in literatures,31,32 but are not conclusive enough to enable unambiguous identification of the structure changes. The BiOF phase is not detected by XRD because it is only about 1.0 wt % of the total sample weight assuming that the Bi2O3 agent was completely consumed in the surface modification process. But the formation of BiOF could be confirmed by XPS as will be shown below. FESEM shows that the particle morphology of LLMO@ Li4M5O12@BiOF is generally similar as that of LLMO (Supporting Information, Figure S1). However, the particles are severely etched by the acidic NH4F solution leaving clear steps on the particle surface. EDS analysis shows that all the Ni, Co, Mn, Bi, and F elements are uniformly distributed in the material (Supporting Information, Figure S2). The HRTEM images in Figure 2 present detailed microstructures of the materials. The pristine LLMO shows well-ordered lattice fringes throughout the particle (Figure 2a). The distance of lattice fringes is 4.8 Å, which fits well with that of the (003) planes of layered LLMO. The layered structure is further confirmed by selected area electron diffraction (SAED) as shown in the inset of Figure 2a. The microstructure of LLMO@Li4M5O12@BiOF is different from that of LLMO. At the surface region, a uniform amorphous layer with thickness about 3 nm is assigned to the BiOF coating layer (Figure 2b). The SAED pattern of the material is composed of two integrated structures which are indexed to the (003), (104), (107) diffractions of layered LLMO and the (220), (240), (420) diffractions of spinel Li4Mn5O12. Note that the zone axis used for SAED analysis are [010]layered and [001]spinel, respectively. The Li4M5O12 phase is below the BiOF layer (Figure 2c), heterostructured with the layered LLMO (Figure 2d). Moreover, in the Raman spectrum of LLMO@Li4M5O12@ BiOF (Supporting Information, Figure S3), the shoulder peak at 656 cm−1 confirms that this spinel phase is Li4M5O12 rather than LiMn2O4. For the later spinel material, the corresponding shoulder peak should be located at 623 cm−1.24,33 Figure 3 shows the XPS patterns of the materials. The Ni 2p1/2, Co 2p1/2 and Mn 2p1/2 energies of LLMO are observed at 854.6, 780.8, and 642.4 eV, corresponding to those of Ni2+,

Figure 2. HRTEM images of (a) LLMO and (b−d) LLMO@ Li4M5O12@BiOF.

Co3+, and Mn4+, respectively.34 For the LLMO@Li4M5O12@ BiOF sample, all these transitional metals are observed at the same positions as those of the pristine LLMO indicating that the oxidation states of Ni, Co, Mn are not changed by surface modification. When the LLMO/Bi2O3 composite was dispersed in saturated NH4F solution, a part of Li ions at the particle surface would be replaced by H+ via the ion exchange reaction. Note that the H ions are produced from the hydrolysis reaction of NH4F. In addition, the Li/H exchange would only take place in the Li layer because the energy of extracting one Li from the Li layer is much smaller than that from the transitional metal layer.35 When the H+-exchanged powder was heat treated, the protons and oxygens were extracted from the material (2H+ + O2− = H2O) leaving vacancies in the crystal lattice. In order to keep structure stability, the Li ions nearby the vacancies would migrate from the octahedral site to the tetrahedral site. Meanwhile, the ions in the transitional metal layer migrate to the octahedral vacancies, forming a Li4M5O12 spinel ordering at the interfacial region of the LLMO particles. In this Li4M5O12 spinel phase, the 8a tetrahedral site is occupied by lithium and the 16d octahedral site is occupied by both lithium and transition metals. XPS shows that the atomic ratio of Ni/Co/ Mn is 1.18:1:3.45, indicating that the major transitional metal in the Li4M5O12 spinel is Mn. But it is difficult to determine the exact content of Li. As we proposed above, the formation of the surface spinel involves the Li/H ion exchange, followed by migration of the Li and transition metal cations. Therefore, it is expected that the surface structure should be Li depleted or disordered, resulting in a Li4M5O12-like spinel rather than accurate Li4M5O12. At the surface region, the Bi2O3 oxide reacts with the F− anions under heat treatment forming the BiOF surface layer. Appearance of the Bi 1s (159.4 eV) and F 1s (685.0 eV) XPS confirm the formation of BiOF. 3.2. Galvanostatic Charge−Discharge Cycling. Figure 4a shows the initial charge−discharge profiles of the materials at the 0.2 C rate (1C = 200 mA g−1). The slope region of the first charge profile is attributed to Li+ extraction from the lithium layer and the voltage plateau is due to the oxygen loss process. For LLMO@Li4M5O12@BiOF, the specific capacities 3310

DOI: 10.1021/acsami.5b11199 ACS Appl. Mater. Interfaces 2016, 8, 3308−3318

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Figure 3. (a) Ni 2p, (b) Co 2p, (c) Mn 2p, (d) B 1s, and (e) F 1s XPS patterns of the LLMO and LLMO@Li4M5O12@BiOF samples.

Combining with its excellent cycle stability and slower voltage decay, a large specific energy of 852 Wh kg−1 is obtained for LLMO@Li4M5O12@BiOF after 100 cycles. In comparison, the pristine LLMO only shows a specific energy of 453 Wh kg−1 after 100 cycles (Supporting Information, Figure S5). Figure 4d shows the rate dependent cycling performance of the samples. At each cycle rate, the surface-modified material always shows larger discharge capacity than that of LLMO. Remarkably, LLMO@Li4M5O12@BiOF could still deliver 78 mAh g−1 at 25 C while the pristine LLMO has almost no available capacity at 20 C. Figure 5 shows the representative discharge curves and their dQ/dV profiles of the samples. LLMO shows an initial IR drop from 4.8 to 4.7 V which then gradually decreases to 4.3 V after 100 cycles. This indicates increased internal resistance of the battery cell. In comparison, the IR drop of LLMO@Li4M5O12@ BiOF is almost stabilized at 4.6 V in 100 cycles. From the dQ/ dV profiles, it is seen that the intensities of P1, P2, and P3 peaks of LLMO decrease with cycling. In addition, the P3 peak gradually shifts to 2.7 V in 50 cycles and then stabilizes at this voltage (marked by P5) in subsequent cycles. These changes are attributed to the continuous layered-to-LiMn2O4 transformation of LLMO. This phase transformation results in the voltage decay of LLMO. In addition, partial capacity fading of LLMO could be attributed to this transformation due to dissolution of Mn and the Jahn−Teller effect of LiMn2O4. For the LLMO@ Li4M5O12@BiOF material, the P5 peak disappears after 50

of these two fractions are all smaller than those of pristine LLMO because a part of LLMO was transformed to the Li4M5O12-like spinel during the surface modification process. In spite of this, LLMO@Li4M5O12@BiOF shows a much larger initial discharge capacity than that of LLMO (i.e., 248 vs 292 mAh g−1). Meanwhile, the columbic efficiency increases from 75 to 92%. The dQ/dV profiles help to understand the electrochemical reactions involved in the first discharge (Supporting Information, Figure S4). The P1, P2, and P3 peaks are due to Li+ interaction into layered LLMO accompanied by reduction of Co 4+ , Ni 4+ , and Mn 4+ , respectively.14 LLMO@Li4M5O12@BiOF shows two additional peaks at lower voltages as labeled by P4 and P5. These peaks are due to Li+ intercalation into the Li4M5O1 and LiMn2O4 spinels, respectively.34 This indicates that the large initial discharge capacity of LLMO@Li4M5O12@BiOF is partly attributed to its Li4M5O12-like interfacial structure. In addition, the surface modification facilitates the structure transformation of LLMO which results in some LiMn2O4-like phase after the first cycle. The cycle performances of the materials at 0.2 C rate are displayed in Figure 4b. LLMO@Li4M5O12@BiOF shows a discharge capacity of 269 mAh g−1 after 100 cycles which is much larger than that of LLMO. The corresponding capacity retention is 65 and 92%, respectively. Along with cycling, the average discharge voltage of LLMO decreases from 3.66 to 2.82 V as shown in Figure 4c. In contrast, a higher average discharge voltage of 3.16 V is obtained for LLMO@Li4M5O12@BiOF. 3311

DOI: 10.1021/acsami.5b11199 ACS Appl. Mater. Interfaces 2016, 8, 3308−3318

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Figure 4. Charge−discharge performances of LLMO and LLMO@Li4M5O12@BiOF: (a) initial charge−discharge profiles at 0.2 C; (b) cycling performance; (c) average discharge voltage; and (d) rate-dependent cycling performance.

be underneath the Li4M5O12-like interfacial structure. When the material was discharged to 2.7 V again (2nd discharge), Li ions would intercalate into the 16c site of LiMn2O4 through the lithiated Li4M5O12. At the same time, a part of the Li(16d) ions of lithiated Li4M5O12 could migrate to the vacant 16d site of LiMn2O4 due to the similar coordinate environments of these two 16d sites. This part of LiMn2O4-like phase is therefore converted to a Li4M5O12-like spinel where the 8a tetrahedral site is occupied by lithium and the 16d octahedral site is occupied by both lithium and transition metals. This looks like that some LiMn2O4-like phase is swallowed by Li4M5O12, as shown in Figure 7. Following discharge, the Li(16d) vacancies of the Li4M5O12-like spinel could be filled again when the cell was discharged to 2.9 V. Then, the swallowing process was repeated when the cell was further discharged to 2.7 V. Therefore, the LiMn2O4 phase is gradually swallowed with repeated cycling while the Li4M5O12-like phase keeps growing inward the material bulk. It can be seen from the dQ/dV profiles that this swallowing process would complete after 50 cycles. Afterward, the material shows almost no capacity fading during long-term cycling. 3.3. Electrochemical Kinetic Properties. Figure 8 shows the Nyquist plots of the samples after the first discharge. LLMO shows two semicircles which are attributed to the SEI film and charge transfer process, respectively. But, only one semicircle is

cycles. In the meanwhile, the P4 peak slightly shifts to 2.9 V and keeps grow with cycling. This implies some interesting structure changes occurring in the electrode. To reveal these structure changes, we performed HRTEM analysis on the materials after 140 cycles. Fast Fourier transform (FFT) analysis shows that not only the surface region (Figure 6a) but also the material bulk (Figure 6b) of LLMO transforms to a LiMn2O4-like spinel after long-term cycling. In contrary, the surface of LLMO@Li4M5O12@BiOF transforms to a Li4M5O12like spinel (Figure 6c), and the material bulk is coexistence of LLMO and Li4M5O12 (Figure 6d). This indicates that the LLMO@Li4M5O12@BiOF cathode slowly transforms to a Li4M5O12-like spinel with charge−discharge cycling. The above layered-to-Li4M5O12 transformation of LLMO@ Li4M5O12@BiOF is different from the layered-to-LiMn2O4 transformation of normal Li-excess layered oxides. As a cathode material, Li4M5O12 has higher discharge voltage than that of LiMn2O4. In addition, the Mn4+ ions in Li4M5O12 have minimal Jahn−Teller effect and are free of disproportionation reaction in the electrolyte. Due to these advantages, the LLMO@ Li4M5O12@BiOF cathode exhibits much better cycle stability than that of LLMO. We have proposed a “swallowing effect” to explain the layered-to-Li4M5O12 transformation of LLMO@ [email protected] First, the LiMn2O4 spinel formed after the first cycle is originally from the LLMO phase, and thus, it must 3312

DOI: 10.1021/acsami.5b11199 ACS Appl. Mater. Interfaces 2016, 8, 3308−3318

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ACS Applied Materials & Interfaces

such as ROCH2Li (832 cm−1), CO32− (878 cm−1), P−O−C (1012 cm−1), C−O (1184 cm−1), CO (1273 cm−1), and −CH3 (1392 cm−1).36 (Supporting Information, Figure S6). The SEI resistance of LLMO increases from 164 to 609 Ω after 100 cycles, whereas the SEI resistance of LLMO@Li4M5O12@ BiOF is only 45 Ω after 100 cycles. Meanwhile, the charge transfer resistance of LLMO increases to 1384 Ω which is much larger than that of LLMO@Li4M5O12@BiOF. This indicates that the surface modification not only suppresses the formation of SEI but also improves the charge transfer reactions at the electrode/electrolyte interface. One can see from HRTEM that a thick SEI film forms on the surface of the cycled LLMO (Figure 6a). Accumulation of such a SEI film is a major reason for the large IR drops, poor cycle stability and low rate capability of the material. Contrarily, only a very thin film is observed for LLMO@Li4M5O12@BiOF (Figure 6c). Along with EIS analysis, XPS was used to study the surface chemistry of the cycled materials. The XPS experiments were performed on the fully discharged electrodes after 100 cycles. As shown in Figure 9, the stronger Ni2+ XPS (854.4 eV) for the cycled LLMO@Li4M5O12@BiOF indicates that the Ni cations in this material are more electrochemically active than those of the pristine LLMO. The Co 2p XPS of LLMO is almost disappeared after 100 cycles, and its Mn 2p XPS is much weaker than that of LLMO@Li4M5O12@BiOF. This indicates that the dissolution of Co and Mn are effectively suppressed by surface modification. The P 2p XPS can be fitted by two components. The first one at 135.7 eV is due to the P−F bond of PF5, and the other one at 134.2 eV is attributed to the P− O−F bond of LixPOFy.37 The decomposition of LiPF6 at high voltage involves the following reaction sequence:38

Figure 5. (a) Representative discharge curves and (b) corresponding dQ/dV profiles of LLMO and LLMO@Li4M5O12@BiOF at the 0.2 C rate.

LiPF6 → LiF + PF5

(1)

PF5 + H 2O → POF3 + 2HF

(2)

POF3 + Li 2O → LiF + LixPOFy

(3)

The P 2p XPS of the cycled LLMO is dominated by the P− O−F group. This indicates that the electrolyte decomposition in this battery is vigorous, which goes to the last step of the above reaction sequence (i.e., step 3). In contrast, the strong P−F peak of LLMO@Li4M5O12@BiOF indicates that the electrolyte decomposition in this battery is limited at the first step (i.e., step 1). On this basis, we concluded that the BiOF coating layer effectively suppresses the decomposition of the electrolyte. Figure 10a shows the lithium diffusion coefficients of the materials during the first discharge which are determined by GITT (Supporting Information, Figure S7). It is seen that the diffusion coefficients are improved by 4 orders of magnitude after surface modification. The improved diffusion kinetics can be partly attributed to the interfacial Li4M5O12-like heterostructure whose three-dimensional (3D) framework is favorable for Li ion diffusion. The lithium diffusion coefficients of LLMO@Li4M5O12@BiOF are almost unchanged after 100 cycles (Figure 10b). This is reasonable considering that the material is still coexistence of LLMO and Li4M5O12 after longterm cycling. The pristine LLMO, however, is almost completely transformed to 3D LiMn2O4 after 100 cycles. As a result, the diffusion coefficients of the cycled LLMO increase to 10−14 ∼ 10−11 cm2 s−1. But, these values are still much lower than those of LLMO@Li4M5O12@BiOF.

Figure 6. HRTEM images of (a, b) LLMO and (c, d) LLMO@ Li4M5O12@BiOF after 140 cycles.

observed for LLMO@Li4M5O12@BiOF indicating that the SEI film of this material is very small. To study the chemical composition of the SEI film, we charge−discharged the LLMO and LLMO@Li4M5O12@BiOF electrodes for three cycles. FTIR analysis shows that the cycled materials have the same SEI composition which contains some typical functional groups 3313

DOI: 10.1021/acsami.5b11199 ACS Appl. Mater. Interfaces 2016, 8, 3308−3318

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Figure 7. Scheme of the layered-to-Li4M5O12 phase transformation of LLMO@Li4M5O12@BiOF during charge−discharge cycling.

Figure 8. Nyquist plots of LLMO and LLMO@Li4M5O12@BiOF (a) after the first cycle and (b) after 100 cycles.

3.4. Thermal Safety Properties. For studying the thermal safety of the materials, DSC analysis was performed for the fully charged electrodes after the first cycle (4.8 V). As shown in Figure 11, the delithiated LLMO shows a large exothermic peak at 203 °C and another small peak at 267 °C which result in total thermal release of 856 J g−1. In contrast, the delithiated LLMO@Li4M5O12@BiOF only shows a mild exothermic process. Not only does the onset temperature increase to 261 °C, but the total thermal release is also remarkably reduced to 421 J g−1. This indicates that the surface modification significantly improve the thermal safety of LLMO. To gain deep insight into the thermal behavior of the delithiated electrodes, we performed mass spectrometry and in situ XRD from 25 to 450 °C (Figures 12 and 13). One can see from the room-temperature XRD that the delithiated LLMO has an R3̅m layered structure without any Li2MnO3 superlattice. Moreover, the corresponding HRTEM image could also be characterized by the R3̅m layered structure (Figure 14a). Combining with the

DSC, MS, in situ XRD, and HRTEM analysis, the thermal behavior of the delithiated LLMO can be generally divided into three stages. Stage I. This stage occurs in the temperature range of 25− 230 °C. Only O2 evaporation is observed in this stage. The inverse of the (003)/(104) intensity ratio indicates significant cation mixing between the slab and the interslab space. Gu et al. shows that the Ni ions tend to segregate at the surface of Li[Li0.2Ni0.2Mn0.6]O2 during high-temperature synthesis driven by the thermodynamic force.39 Similarly, we suggest that the heat treatment of delithiated LLMO will result in phase segregation forming a Li/Ni-rich phase at the particle surface and a Co/Mn-rich phase in the material bulk. The (220) peak is characteristic of the Fd3m space group. The appearance of this peak indicates transformation from the LiNiO2-type layered structure to a LiNi2O4-type spinel phase (Fd3m). This transformation involves migration of 1/4 of the Ni4+ cations from the slab to the interslab space, as well as 3314

DOI: 10.1021/acsami.5b11199 ACS Appl. Mater. Interfaces 2016, 8, 3308−3318

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Figure 9. (a) Ni 2p, (b) Co 2p, (c) Mn 2p, and (d) P 2p XPS patterns of LLMO and LLMO@Li4M5O12@BiOF after 100 cycles.

Figure 10. Lithium diffusion coefficients of LLMO and LLMO@Li4M5O12@BiOF: (a) during the first discharge and (b) during the 100th discharge.

Figure 11. DSC curves of the fully charged LLMO and LLMO@ Li4M5O12@BiOF electrodes.

Figure 12. MS patterns of the fully charged LLMO and LLMO@ Li4M5O12@BiOF electrodes.

displacement of the Li+ ions from the octahedral site to the tetrahedral site. As the temperature increases, the decrease of the (003) peak indicates that the LiNi2O4-type phase tends to decompose to rock-salt LixNiyO at higher temperature. Both the LiNiO2-to-LiNi2O4 and LiNi2O4-to-LixNiyO transformations are associated with oxygen release from the material lattice.40 HRTEM shows that the initially well-crystallized layered material deforms to numerous nanosized regions after

heating the material at 220 °C (Figure 14b). The FFT pattern is ambiguous but can be indexed to either spinel LiNi2O4 or rock salt LixNiyO. Stage II. This stage corresponds to the temperature range of 230−340 °C. Several gases are released during this stage including H2O, CO2, CH3-group, and a small amount of O2. The LiNi2O4-type phase would continue to transform to 3315

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site. This indicates that a part of the Mn/Co ions in the material bulk migrate into the tetrahedral site during high temperature treatment. Meanwhile, excessive oxygen releases from the material accompanied by the reduction of Mn4+ and Co4+. The O2 gas generated in both Stage I and Stage II would trigger decomposition of the electrolyte solvent producing CO2 and H2O gases. In addition, self-decomposition of the electrolyte components could occur in this temperature range, which generates CO 2 and CH 3 -group gaseous compounds.41 Stage III. No obvious structure change happens in the temperature range of 340−450 °C. The observed CO2 gas could be attributed to combustion of the PVDF binder and the active carbon in the electrode composite. The thermal reactions involved in this stage or higher temperatures are less concern for battery safety issues. Obviously, Stage I is the most harmful for the thermal safety of LLMO not only because it occurs at low temperature but also because it contributes 62% of the total thermal release. All these harmful effects originate from the poor structure stability of the delithiated LLMO. In addition, the released oxygen resulted from the structure deformation triggers decomposition of the electrolyte solvents. DSC shows that these harmful effects could be effectively eliminated by our LLMO@ Li4M5O12@BiOF material. The onset thermal release temperature is increased to 261 °C by surface modification, and the total thermal release is reduced about 50% with respect to the pristine LLMO. Mass spectrometry does not detect any oxygen release. In addition, the in situ XRD patterns do not show any changes during the whole heat treatment process. This indicates that the structure of delithiated LLMO@Li4M5O12@ BiOF does not change at high temperature. HRTEM shows that the Li4M5O12 interfacial structure does not change up to 220 °C (Figure 14c,d). This stable interfacial structure prevents accumulation of Li/Ni-rich phase at the particle surface. In addition, possible oxygen release is prohibited due to the high kinetic barriers of oxygen migration in the close packed array. As a result, the crystal structure of the material is perfectly reserved. Moreover, the decomposition of electrolyte solvents is prohibited without participation of O2. The thermal behavior of the material therefore only involves the self-decomposition of the electrolyte, which results in excellent thermal safety.

Figure 13. XRD patterns the fully charged LLMO and LLMO@ Li4M5O12@BiOF electrode.

4. CONCLUSION In summary, we have reported a new surface engineering method for Li-excess layered cathode materials. The successfully prepared material is composed of a high-capacity Li(Li0.18Ni0.15Co0.15Mn0.52)O2 cathode in conjunction with a Li4M5O12-like heterostructure and a BiOF surface layer. The layered-to-LiMn2O4 transformation occurring in normal Liexcess layered oxides during charge−discharge cycling is replaced by a layered-to-Li4M5O12 transformation. This unique phase transformation is very important for the material to keep high working voltage, good cycle stability, and high rate capability. The material shows much improved electrochemical kinetic properties with respect to its pristine counterpart because the 3D framework of Li4M5O12 is much favorable for Li ion diffusion; and the BiOF coating layer could suppress the side reactions of the electrode and the electrolyte. The Li4M5O12-like interfacial structure prohibits oxygen release from the material, thus improving structure stability of the electrode at elevated temperatures. The above synergetic effects of the Li4M5O12-like surface spinel and the BiOF coating layer

Figure 14. HRTEM images of the fully charged (a, b) LLMO and (c, d) LLMO@Li4M5O12@BiOF.

LixNiyO in this high temperature range. In addition, the increase of the (220) peak and the progressive merging of the (018)/(110) doublets indicate formation of another spinel phase which is probably MnCo2O4 (PDF No. 23-1237). The Mn and Co ions in spinel MnCo2O4 are both in mixed valence states, in which Mn2+ and Co2+ are located at the tetrahedral stacking interstices, while Mn3+ and Co3+ occupy the octahedral 3316

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(8) Gu, M.; Belharouak, I.; Zheng, J.; Wu, H.; Xiao, J.; Genc, A.; Amine, K.; Thevuthasan, S.; Baer, D. R.; Zhang, J.-G.; Browning, N. D.; Liu, J.; Wang, C. Formation of the Spinel Phase in the Layered Composite Cathode Used in Li-Ion Batteries. ACS Nano 2013, 7, 760−767. (9) Yu, S.-H.; Yoon, T.; Mun, J.; Park, S.; Kang, Y.-S.; Park, J.-H.; Oh, S. M.; Sung, Y.-E. Continuous Activation of Li2MnO3 Component upon Cycling in Li1.167Ni0.233Co0.100Mn0.467Mo0.033O2 Cathode Material for Lithium ion Batteries. J. Mater. Chem. A 2013, 1, 2833−2839. (10) Croy, J. R.; Kim, D.; Balasubramanian, M.; Gallagher, K.; Kang, S.-H.; Thackeray, M. M. Countering the Voltage Decay in High Capacity xLi2MnO3-(1-x)LiMO2 Electrodes (M = Mn, Ni, Co) for Li+Ion Batteries. J. Electrochem. Soc. 2012, 159, A781−790. (11) Yabuuchi, N.; Yoshii, K.; Myung, S. T.; Nakai, I.; Komaba, S. Detailed Studies of a High-capacity Electrode Material for Rechargeable Batteries, Li2MnO3-LiCo(1/3)Ni(1/3)Mn(1/3)O2. J. Am. Chem. Soc. 2011, 133, 4404−4419. (12) Liu, J.; Manthiram, A. Functional Surface Modifications of a High Capacity Layered Li[Li0.2Mn0.54Ni0.13Co0.13]O2 Cathode. J. Mater. Chem. 2010, 20, 3961−3967. (13) Martha, S. K.; Nanda, J.; Veith, G. M.; Dudney, N. J. Surface Studies of High Voltage Lithium-rich Composition: Li1.2Mn0.525Ni0.175Co0.1O2. J. Power Sources 2012, 216, 179−186. (14) Fu, Q.; Du, F.; Bian, X.; Wang, Y.; Yan, X.; Zhang, Y.; Zhu, K.; Chen, G.; Wang, C.; Wei, Y. Electrochemical Performance and Thermal Stability of Li1.18Co0.15Ni0.15Mn0.52O2 Surface Coated with the Ionic Conductor Li3VO4. J. Mater. Chem. A 2014, 2, 7555−7562. (15) Qiao, Q.-Q.; Qin, L.; Li, G.-R.; Wang, Y.-L.; Gao, X.-P. Snstabilized Li-rich Layered Li(Li0.17Ni0.25Mn0.58)O2 Oxide as a Cathode for Advanced Lithium-ion Batteries. J. Mater. Chem. A 2015, 3, 17627−17634. (16) Sun, Y. Y.; Li, F.; Qiao, Q. Q.; Cao, J. S.; Wang, Y. L.; Ye, S. H. Surface Modification of Li(Li0.17Ni0.2Co0.05Mn0.58)O2 with LiAlSiO4 Fast Ion Conductor as Cathode Material for Li-ion Batteries. Electrochim. Acta 2015, 176, 1464−1475. (17) Croguennec, L.; Bains, J.; Bréger, J.; Tessier, C.; Biensan, P.; Levasseur, S.; Delmas, C. Effect of Aluminum Substitution on the Structure, Electrochemical Performance and Thermal Stability of Li1+x(Ni0.40Mn0.40Co0.20−zAlz)1−xO2. J. Electrochem. Soc. 2011, 158, A664−A670. (18) Singh, G.; Thomas, R.; Kumar, A.; Katiyar, R. S. Electrochemical Behavior of Cr− Doped Composite Li2MnO3-LiMn0.5Ni0.5O2 Cathode Materials. J. Electrochem. Soc. 2012, 159, A410−A420. (19) Wang, Y.; Yang, Z.; Qian, Y.; Gu, L.; Zhou, H. New Insights into Improving Rate Performance of Lithium-Rich Cathode Material. Adv. Mater. 2015, 27, 3915−3920. (20) Zhang, H. Z.; Qiao, Q. Q.; Li, G. R.; Gao, X. P. PO43− Polyanion-doping for Stabilizing Li-rich Layered Oxides as Cathode Materials for Advanced Lithium-ion Batteries. J. Mater. Chem. A 2014, 2, 7454−7460. (21) Li, G. R.; Feng, X.; Ding, Y.; Ye, S. H.; Gao, X. P. AlF3-coated Li(Li0.17Ni0.25Mn0.58)O2 as Cathode Material for Li-ion Batteries. Electrochim. Acta 2012, 78, 308−315. (22) Kim, I. T.; Knight, J. C.; Celio, H.; Manthiram, A. Enhanced Electrochemical Performances of Li-rich Layered Oxides by Surface Modification with Reduced Graphene oxide/AlPO4 Hybrid Coating. J. Mater. Chem. A 2014, 2, 8696−8704. (23) Kang, S.-H.; Thackeray, M. M. Enhancing the Rate Capability of High Capacity xLi2MnO3·(1−x)LiMO2 (M = Mn, Ni, Co) Electrodes by Li−Ni−PO4 Treatment. Electrochem. Commun. 2009, 11, 748−751. (24) Li, B.; Li, C.; Cai, J.; Zhao, J. In situ nano-coating on Li1.2Mn0.52Ni0.13Co0.13O2 with a Layered@spinel@coating Layer Heterostructure for Lithium-ion Batteries. J. Mater. Chem. A 2015, 3, 21290−21297. (25) Oh, P.; Ko, M.; Myeong, S.; Kim, Y.; Cho, J. A Novel Surface Treatment Method and New Insight into Discharge Voltage Deterioration for High-Performance 0.4Li 2 MnO 3 -0.6LiNi1/3Co1/3Mn1/3O2 Cathode Materials. Adv. Energy Mater. 2014, 4, 1400631.

lead to excellent capacity retention, low voltage decay, and high rate capability. Especially, the thermal safety of the electrode is remarkably improved, resulting in a high onset temperature and minimal thermal release.

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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.5b11199. FESEM images of the LLMO and LLMO@Li4M5O12@ BiOF samples; elemental mappings of the LLMO@ Li4M5O12@BiOF sample ; Raman patterns of the LLMO and LLMO@Li4M5O12@BiOF samples; dQ/dV profiles of the LLMO and LLMO@Li4M5O12@BiOF samples during the first discharge; specific energies of the LLMO and LLMO@Li4M5O12@BiOF samples; FTIR patterns of the LLMO and LLMO@Li4M5O12@BiOF materials after three cycles; GITT analysis of the LLMO; and LLMO@Li4M5O12@LBO samples. (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. Tel/Fax: 86-431-85155126. *E-mail: [email protected]. Tel/Fax: 86-431-85155126. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by Ministry of Science and Technology of China (No. 2015CB251103), National Natural Science Foundation of China (No. 51472104, 21473075, 51272088, and 51572107), Defense Industrial Technology Development Program of China (No. B1420133045), Jilin Provincial Science and Technology Department (No. 20140101083JC, 20150204078GX).

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REFERENCES

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