Mixed Electronic and Ionic Conductor-Coated Cathode Material for

Apr 29, 2016 - Electrochemistry Communications 2017 83, 106-109 ... Sujith Kalluri , Moonsu Yoon , Minki Jo , Suhyeon Park , Seungjun Myeong , Junhyeo...
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Mixed Electronic and Ionic Conductor-Coated Cathode Material for High-Voltage Lithium Ion Battery Jae-Hyun Shim,† Jung-Min Han,† Joon-Hyung Lee,*,‡ and Sanghun Lee*,§ †

Battery R&D Center, Samsung SDI, Suwon, Gyunggido 16677, Republic of Korea Automotive & ESS Business, Samsung SDI, Yongin, Gyunggido 17084, Republic of Korea § Department of BioNano Technology, Gachon University, Seongnam, Gyunggido 13120, Republic of Korea Downloaded via UNIV OF CALIFORNIA SANTA BARBARA on August 6, 2018 at 13:22:24 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

‡

ABSTRACT: A lithium ionic conductor, Li1.3Al0.3Ti1.7(PO4)3 (LATP), is introduced as a coating material on the surface of Mg-doped LiCoO2 to improve electrochemical performances for high-voltage (4.5 V) lithium ion batteries. Structure, morphology, elemental distribution, and electrical properties of the materials are thoroughly characterized by SEM, TEM, EELS, EDS, and C-AFM. The coating layer is electrically conductive with the aid of Mg ions which are used as a dopant for the active materials; therefore, this mixed electronic ionic conductor strongly enhances the electrochemical performances of initial capacity, cycling property, and rate capability. The LATP coating layer also demonstrates very promising applicability for 4.4 V prismatic full cells with graphite anode, which correspond to the 4.5 V half-cells with lithium anode. The 2900 mA h full cells show 85% of capacity retention after 500 cycles and more than 60% after 700 cycles. KEYWORDS: lithium ion battery, mixed electronic and ionic conductor, lithium cobalt oxide, high-voltage cell, full cell application

1. INTRODUCTION For the well-balanced electrochemical performances of energy density, power, rate capability, and cyclability, LiCoO2 has been the most commonly used cathode material for the lithium ion battery (LIB) since the first commercialization by SONY in 1991.1 However, the recent development of high-performance mobile electronics and transportations powered by electric energy, such as electric vehicles and hybrid electric vehicles, strongly requires improving performances of the LIB.2,3 In particular, increasing the energy density is one of the priority tasks for future advanced energy storage. With finding new materials possessing many lithium ions in a unit volume or mass, for example, lithium-rich oxides,4−6 raising cutoff voltages of currently known materials is a main strategy to obtain the high energy density. Because the layer-structured materials including LiCoO2 without any modification, in which transition metal ions are directly coordinated by oxygen, may suffer from easy structure collapse during cycling, particularly, at high voltages, the recent progress in the high-voltage LIBs has been made primarily in polyanion oxides or spinel materials.7 However, because the theoretical capacities of these materials are much lower than those of the layered materials, they are not practically employed in commercial batteries and LiCoO2 is still one of the most popular cathode materials for LIB. The vulnerability of LiCoO2 at the high voltages is mainly attributed to the two factors of (1) oxygen loss caused by the side reactions with electrolytes and (2) irreversible structural transformation at higher voltages than 4.5 V (vs Li+/Li). Hence, Chen and Dahn suggested that coating with metal oxide materials under heat treatment would remarkably improve the © 2016 American Chemical Society

retention property of LiCoO2 by suppressing the side reactions when charging to 4.5 V.8 This surface modification by metal oxide coating to improve electrochemical performances of LiCoO2 was originally proposed by Cho et al.9−11 and many researchers confirmed the successful improvement with various metal oxides as coating materials.12−18 However, because most of the metal oxides, such as, MgO, ZrO2, Al2O3, and ZnO, etc., are the insulators for lithium ions, the coating materials may work as obstacles for ionic conduction, which is one of the main reasons for high resistivity at the interface between electrode and electrolytes. Therefore, in order to improve the ionic conductivity of the coating layer, recently, lithium ion conductors have been introduced as alternative coating materials. For example, Lu et al. proposed that monoclinic Li2TiO3 would be a good candidate.19 They reported that Li2TiO3-coated LiMO2 (M = Ni, Co, and Mn) nanobelts exhibited very good rate capability and cycling property. In addition, based on the fact that Li4Ti5O12 does not suffer from significant volume change during lithium cycling, which is another example of lithium ion conductor and known as “zerostrain material” on lithiation/delithiation, several researchers have tested it as a coating material for the layer-structured cathode materials.20−24 Even though the above studies reported that the lithium ion conductor-coated cathode materials demonstrated considerable successes in half-coin-cells, they failed to be employed in the battery industry because they Received: March 13, 2016 Accepted: April 29, 2016 Published: April 29, 2016 12205

DOI: 10.1021/acsami.6b03113 ACS Appl. Mater. Interfaces 2016, 8, 12205−12210

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Mahwah, NJ, USA) equipped in the SEM or TEM. The EELS spectrum was obtained in the STEM mode with a total beam current of 0.2 nA for a probe size of 0.5 nm (Gatan GIF200 spectrometer). The electrical conductivity of the materials was analyzed by a C-AFM (Danish Micro Engineering, Herlev, Denmark) in the direct contact mode. Following the sample preparation by helium ion microscope (Carl Zeiss Microscopy GmbH), the measurements were performed in an ultrahigh vacuum (10−8 Torr) to avoid contamination. To evaluate the electrochemical performances, 2032-type coin-cells were fabricated. For the preparation of the cathode, the active material (LATP-LCOMg, LATP-LCO, or uncoated-LCOMg), conducting material (Super-P carbon black), and binder (polyvinylidene fluoride, PVDF) were blended (96:2:2 in weight ratio) in N-methyl-2pyrrolidone. Then, the slurry was cast on an aluminum foil and dried at 110 °C for 10 h in a vacuum oven. The electrolyte solution was 1.15 M LiPF6 in ethylene carbonate/dimethyl carbonate/ ethylmethyl carbonate mixture (3:3:4 in volume ratio). The lithium foil was used as the anode, and the capacity of the manufactured cells was 1.5 mA h/cm2. The cycling capacities were measured between 3.0 and 4.5 V. In order to see if this new material has the potential for industrial use, prismatic full cells of 2900 mA h were manufactured. The cathode was composed of the active material, conducting material, and binder (97.4:1.3:1.3 in weight ratio), in which the contents of the latter two components were reduced to increase the capacity per volume of a full cell. The natural graphite (BSG-L, BTR, Shenzhen, China) anode and the electrolyte of 0.9 M LiPF6 solution in ethylene carbonate/ ethylmethyl carbonate/diethyl carbonate mixture (3:5:2 in volume ratio) were used with several minor additives (fluoroethylene carbonate, vinyl ethylene carbonate, propylsultone, and LiBF4). The electrolyte composition of the full cells was designed differently from that of the coin-cells for common use in various product lines under consideration of the effects from the graphite anode. The cells were cycled between 3.0 and 4.4 V with the rate of 1 C.

could not overcome a serious capacity fade in full cells due to their poor electrical conductivity and/or structural instability during cycling. In fact, the most available application of the lithium ion conductors is a solid electrolyte for all-solid-state batteries. Among them, Li1+xAlxTi2−x(PO4)3 (LATP) of the NASICON structure exhibits excellent ionic conductivity at room temperature,25 and, therefore, it is one of the most promising solid electrolyte materials for all-solid-state LIBs.26−28 In a few recent studies, however, LATP was tested not as a solid electrolyte but as a minor constituent for cathode composite (LiCoO2 and LATP). Morimoto et al. prepared this composite by simple mechanical mixing and reported high capacity and good cycling property with a high-cutoff potential (4.5 V) in a liquid electrolyte LIB.29 Meanwhile, Nakamura et al. obtained this material by a similar method and used it as a cathode for allsolid-state LIB.30 They showed that this material exhibited much better electrochemical performances than the LiCoO2 cathode without LATP. In this study, we use LATP (x = 0.3 in this work) as a coating material for the LiCoO2 cathode and, consequently, improve lithium ionic conductivity at the interface between the active materials and electrolytes. Moreover, as shown in our previous study,24 it is possible to increase electrical conductivity of the coating layer by diffusion of dopant ions (Mg2+ in this work) from the active materials. As a result, we obtain mixed electronic and ionic conductor-coated LiCoO2 with Mg doping. We thoroughly characterize the coating layer by various analysis techniques including scanning electron microscopy (SEM), transmission electron microscopy (TEM) with electron energy loss spectroscopy (EELS), and conductive atomic force microscopy (C-AFM), and understand the mechanism of the excellent electrochemical performances of this material. Finally, this material is adopted for a prismatic 4.4 V full cell with a graphite anode, which is corresponding to the 4.5 V cathode vs Li+/Li, and exhibits very good capacity retention over 700 cycles. This work provides a very informative guide to the development of good cathode materials by introducing mixed electronic and ionic conductive coating layers.

3. RESULTS AND DISCUSSION The morphologies and internal structures of uncoated-LCOMg and LATP-LCOMg were characterized by SEM and TEM. The typical dimension of the synthesized particles of both samples is ∼15 μm. The LATP coating layer seems to be inhomogeneous in a relatively large scale (Figure 1b); however, it is continuous and evenly thick in a sub-micrometer scale as shown in Figure 1c. The selected area electron diffraction (SAED) pattern of the coating layer in Figure 1c supports that it has the NASICON hexagonal structure of Li1.3Al0.3Ti1.7(PO4)3 with parameters of a = 8.5 Å and c = 20.8 Å.31 Figure 1d shows the elemental distributions of cations from EDS analysis. The Co and Al ions are located in the regions where they are expected to be Co in the active layer and Al in the coating layer. Meanwhile, the Mg and Ti ions are observed in both the coating layer and the LiCoO2 region. Note that Mg was used as a dopant in the LiCoO2 region and Ti was used as one of the sources for the coating layer. The diffusion of Mg ions from the active material to the coating layer was previously reported, where the coating layer containing the diffused Mg ions exhibited improvement of the electrical conductivity and, consequently, the enhanced electrochemical performances.24 We observed similar behaviors in this study, which will be shown later. The Ti ions migrated to the active material are mixed with the Co ions in the transition metal (TM) layers. For identification of the TM ions, the EELS spectra were obtained along the TM layers (yellow dotted arrows in Figure 1e). Even though only the Co ions were detected in most of the measured area (spectra B), a few TM ions were identified as the Ti ions (marked in a red dotted circle, spectra A), which indicates that the small amounts of Co ions occupied on arbitrary sites are replaced by the Ti ions.

2. EXPERIMENTAL SECTION Prior to the LATP coating, Mg-doped LiCoO2 (LiMg0.03Co0.97O2) was prepared by the solid-state reaction as in our previous study.24 Then, the sol−gel method was employed for the LATP coating. First, 1.555 g of Ti(OBu)4 was dissolved in 19 mL of ethanol. In a separate container, 0.549 g of LiNO3 and 0.900 g of NH4H2PO4 were dissolved in 8 mL of water. After completely mixing the two solutions, 0.296 g of Al(NO3)3·9H2O was added. Then, this final solution and 50 g of the active material were mixed by ball-milling for 1 h (zirconia ball, 150 rpm). Following evaporation of the solvent on a hot plate at 110 °C, the mixture was dried in a convection oven at 120 °C. Then, the obtained powder was heat-treated at 750 °C for 6 h in a furnace and naturally cooled to room temperature. This product is hereafter denoted by LATP-LCOMg. In order to demonstrate the excellence of this material, we prepared two references for comparison. The one is LATP-coated but Mg-undoped LiCoO2 (LATP-LCO), the coating procedure of which was the same as described above; the other one is Mg-doped LiCoO2 without the coating (uncoated-LCOMg). The morphology of the materials was investigated by a field emission SEM (Magellan XHR, FEI Co., Hillsboro, OR, USA). The scanning TEM (STEM) measurements were performed using a Cscorrected STEM (JEM-2100F, JEOL, Co., Tokyo, Japan) at 200 kV, which is equipped with a spherical aberration corrector (CEOS Gmbh, Heidelberg, Germany). The elemental distribution of the materials was characterized by the energy dispersive spectrometer (EDS, EDAX, 12206

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the electrochemical performances is fairly limited without the Mg doping. As shown earlier, the coating layer of LATPLCOMg contains a considerable amount of Mg ions. We suppose that the migrated Mg ions improve the electrical conductivity of the coating layer and, as a result, the cells from LATP-LCOMg exhibit much better performances than the ones from LATP-LCO. Experimental evidence supporting this assertion will be presented later. Sometimes, differential capacity (dq/dv) vs cell voltage plots provide very useful information on the phase transitions of electrode materials during lithiation or delithiation. Panels c and d of Figure 2 show the dq/dv plots of uncoated-LCOMg and LATP-LCOMg, respectively, as a function of the number of cycles. In these measurements, charging and discharging were performed within a range of 3.0−4.6 V. For both samples, several peaks indicating phase transitions are clearly recognized in the curves of the first cycle. As previously reported,32−35 the peak of “a” at 3.92 V corresponds to the transition between the two O3 phases and the “b” (4.09 V) and “c” (4.18 V) peaks are assigned to the order−disorder (hexagonal to monoclinic and monoclinic to hexagonal) transitions. In addition, the forth peak of “d” at 4.53 V is also shown, which corresponds to the transition from the O3 phase to the H1−3 phase. The most remarkable difference between the two samples is the variation of the “a” peak with continuing cycles. In the case of uncoatedLCOMg, as the charging/discharging repeats, it shifts markedly to high voltages along with significant intensity reduction; however, in the case of LATP-LCOMg, the extent of variation of the “a” peak is relatively small. This observation implies that the structural transition indicated by the “a” peak of uncoatedLCOMg significantly loses its reversibility as the cycle repeats whereas LATP-LCOMg relatively well maintains its original structure even after cycles. The behaviors of the “b” and “c” peaks, which are shown in the enlarged insets, also support this argumentation. It is well-known that the structural degradation of LiCoO2 with cycling is driven by oxygen loss due to detrimental reactions of the active materials with electrolytes. In this point of view, it is considered that the LATP coating layer successfully works as a protection layer to reduce the side reactions. To reduce the reactions with electrolytes at the electrode surface, it is necessary to obtain low resistivity, i.e., high electrical conductivity of the materials. From the electrochemical performance results in the above, we supposed that the Mg ions from the active material would improve electrical conductivity of the LATP coating layer. To support this presumption, we investigated the electrical conductivity of LATP-LCOMg and LATP-LCO using the C-AFM, which was equipped with the SEM to find a proper area to be measured in the very small scale.36 A cross-section near the interface between the coating layer and the active material of a cleaved particle was characterized. The surface topography and the corresponding current image were simultaneously obtained by applying an electrical bias of 4 V. As shown in Figure 3a, LATPLCOMg has a relatively uniform height throughout the entire area while its coating layer displays much higher conductivity than the active material. Meanwhile, in the case of LATP-LCO, the cross-sectional topography of which is very close to that of LATP-LCOMg, the LATP coating layer is almost electrically nonconducting (Figure 3b). This observation verifies that the diffused Mg ions from the LiCoO2 region to the LATP layer noticeably raise its conductivity and, consequently, create a synergistic positive effect on the cell performances.

Figure 1. SEM images of (a) uncoated-LCOMg and (b) LATP-LCOMg. (c) Bright-field TEM image and SAED pattern of LATP coating layer. (d) STEM image and elemental distributions from EDS of LATP coating layer. (e) Enlarged STEM image and EELS spectrum of active material of Mg-doped LiCoO2.

In Figure 2a,b, the cycling performance at 1 C and the C-rate property of the materials are respectively shown. The capacities were measured within a range of 3.0−4.5 V vs Li metal anode and two cycles were performed at 0.1 C for the formation before the measurements. Compared to uncoated-LCOMg (167.6 mA h g−1), LATP-LCOMg exhibits a significantly improved discharge capacity at the first cycle (185.0 mA h g−1). Even more remarkable is the enhanced cycling ability of LATP-LCOMg. As shown in Figure 2a, uncoated-LCOMg demonstrates a very sharp discharge capacity fade, particularly, after 30 cycles, whereas LATP-LCOMg shows an extremely improved cycling ability; its capacity after 100 cycles (176.8 mA h g−1) is higher than 95.5% of its initial capacity. In addition, the LATP coating also offers superiority in the C-rate capability of the cathode materials. As shown in Figure 2b, the positive influence of the coating on the discharge capacity becomes greater as the C-rate of cycling increases; the capacity of LATPLCOMg is 1.07 and 1.46 times higher than that of uncoatedLCOMg at 0.1 and 3.0 C, respectively. This improvement of the cell performances with the LATP coating is even more noteworthy when it is applied to the prismatic full cells, which will be shown later. Another important point which deserves attention in Figure 2a,b is that the coating effect on 12207

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Figure 2. (a) Cycling performance at 1 C and (b) rate capability of half-coin-cells in the range of 3.0−4.5 V. Differential capacity vs voltage (dq/dv) curves of (c) uncoated-LCOMg and (d) LATP-LCOMg cells.

Figure 3. Height and current maps of (a) LATP-LCOMg and (b) LATP-LCO obtained from C-AFM.

To examine the industrial applicability of this material, prismatic full cells (3.0−4.4 V, 2900 mA h) were manufactured. It should be noted that the cells were cycled at quite rapid charging/discharging rate of 1 C to investigate the cycling

ability. In Figure 4a, the capacity retention at a certain cycle represents the ratio of the discharge capacity at the corresponding cycle to that at the first cycle. The cells with LATP-LCOMg show very stable retention characteristics (85% 12208

DOI: 10.1021/acsami.6b03113 ACS Appl. Mater. Interfaces 2016, 8, 12205−12210

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reactions allowed by the absence of the protection layer, while those with LATP-LCO show a sudden catastrophe in the capacity after ∼200 cycles even though they demonstrate quite stable performances in the beginning of the cycling tests. To compare the deteriorated states of LATP-LCOMg and uncoated-LCOMg, the SEM images and the elemental distributions from the EDS of the cross-sectioned particles after cycling are shown in Figure 4b,c. In the EDS map of uncoated-LCOMg, a large amount of F is observed between the particles, which is obviously a consequence of the side reactions with LiPF6 salts in the electrolyte whereas, in the map of LATP-LCOMg, the F atoms are scarcely recognized. This observation is indicative of the fact that the LATP coating layer successfully works to suppress the side reactions even in the prismatic full cells for practical application. For further investigation, the surface area of LATP-LCOMg after 700 cycles was measured by the TEM. Because the particles are not completely surrounded by the LATP layer, which means that coated and noncoated areas coexist, the effects of coating can be clearly determined in one image. In Figure 5, the SAED pattern of the area covered by the LATP layer (denoted as B) proves that the area preserves its original structure of LiCoO2 (R3̅m) after cycling. However, the pattern of the noncoated surficial area (denoted as A), which is a typical pattern of the spinel structure (Fd3̅m), indicates that the noncoated region undergoes structural transformation during cycling. The STEM images in the high resolution are also consistent with this explanation. The spinel phase is clearly shown in the image of A region whereas only the layered structure of LiCoO2 is observed in the B area. Figure 4. (a) Retention capacity of 2900 mA h 4.4 V prismatic full cells with graphite anode. SEM images and elemental distributions obtained from EDS of cross-sections of (b) LATP-LCOMg and (c) uncoated-LCOMg particles. For the measurements, the cells were disassembled after 700 and 400 cycles, respectively.

4. CONCLUSIONS We have prepared and characterized LATP-coated LiCoO2 with Mg doping for 4.5 V high-performance LIBs. From the elemental distribution analysis, it has been revealed that some amounts of Mg ions, which are intended to be used as a dopant for LiCoO2, are diffused to the coating layer. As a result, the lithium ionic conductor, LATP, also obtains high electrical conductivity and creates a positive effect on the electrochemical performances of the cycling property and C-rate capability. In addition, we have proved that the excellence of this material is

after 500 cycles and more than 60% even after 700 cycles) at room temperature. On the contrary, the references for comparison show poor performances. The cells with uncoated-LCOMg undergo a significant capacity drop at the early cycles (up to ∼100 cycles), which is caused by the side

Figure 5. Bright-field TEM image in the interface of LATP-LCOMg after 700 cycles. The SAED patterns and STEM images of A and B regions are also shown. 12209

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preserved in the industrial application for the 4.4 V prismatic full cells. The cells show excellent capacity retention of 85% after 500 cycles and more than 60% after 700 cycles, and the cathode material demonstrates a strong structural tolerance to the cycling degradation.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (J.-H.L.). *E-mail: [email protected] (S.L.). Notes

The authors declare no competing financial interest.

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REFERENCES

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