Hierarchically Structured CoreâShell Design of a Lithium Transition

Subscriber access provided by Iowa State University | Library

Energy, Environmental, and Catalysis Applications

Hierarchically-Structured Core-Shell Design of a Lithium Transition Metal Oxide Cathode Material for Excellent Electrochemical Performance Jae-Hyun Shim, Young-Hoon Kim, Han-Sol Yoon, Han-A Kim, JiSoo Kim, Jongsik Kim, Namhee Cho, Young-Min Kim, and Sanghun Lee ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b19902 • Publication Date (Web): 04 Jan 2019 Downloaded from http://pubs.acs.org on January 6, 2019

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

ACS Applied Materials & Interfaces

Hierarchically-Structured Core-Shell Design of a Lithium Transition Metal Oxide Cathode Material for Excellent Electrochemical Performance Jae-Hyun Shim†,∥, Young-Hoon Kim‡,∥, Han-Sol Yoon§, Han-A Kim⊥, Ji-Soo Kim¶, Jongsik Kim⊥, Nam-Hee Cho§, Young-Min Kim*,‡,∆, Sanghun Lee*, #

†Department

of Advanced Materials and Energy Engineering, Dongshin University, Naju 58245, Republic of Korea.

‡Department

of Energy Science, Sungkyunkwan University (SKKU), Suwon 16419, Republic of Korea.

§Department

of Materials Science and Engineering, Inha University, Incheon 22212, Republic of Korea.

⊥Department

¶Gumi

of Chemistry, Dong-A University, Busan 49315, Republic of Korea.

Electronics and Information Technology Research Institute, Gumi 39171, Republic of Korea.

#Department

∆Center

of Chemistry, Gachon University, Seongnam 13120, Republic of Korea.

for Integrated Nanostructure Physics, Institute for Basic Science (IBS), Suwon 16419, Republic of Korea.

Keywords: Li ion battery, lithium mixed transition metal oxides, core shell, Li(NixCoyMnz)O2, STEM-EDX/EELS, electron tomography 1

ACS Paragon Plus Environment

ACS Applied Materials & Interfaces 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Abstract Tuning geometrical parameters of lithium-mixed transition metal oxide (LiTM) cathode materials is a promising strategy for resource-efficient design of high-performance Li ion batteries. In this paper, we demonstrate that simple and facile geometrical tailoring of the secondary microstructure of LiTM cathode materials without complex chemical modification or heterostructure engineering can significantly improve overall electrochemical performance of the active cathode materials. An optimized LiTM with a bimodal size distribution of primary particles inside the secondary particles exhibits a 53.8 % increase in capacity at a high discharge rate (10C) compared to a commercial structure, as well as comparable rate capability after 100 cyclic charge/discharge cycles. The key concept of this approach is to maximize the beneficial effects arising from the controlled sizes of primary particles. Multimodal/multiscale microscopic characterizations based on electron tomography and scanning transmission electron microscopy, combined with electron energy loss spectroscopy and energy-dispersive X-ray spectroscopy from the atomic level to the microscale level, were employed to elucidate structural origins of enhanced battery performance. This study paves the way for the resourceefficient microstructure design of LiTM cathode materials to maximize capacity and stability via simple adjustment of processing conditions, which is advantageous for mass production applications.

2


Page 2 of 37

Page 3 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


1. INTRODUCTION Since the commercialization of lithium ion batteries by Sony in 1991, the use of LIBs has steadily expanded to a plethora of energy storage applications for mobile phones, laptop computers, electric vehicles, and massive energy storage systems. In particular, to adopt LIBs in EVs, higher energy density must be achieved and secured for long-term operation. To this end, many researchers have made great efforts to find promising cathode materials that can replace LiCoO2, which has been one of the most popular industrial cathode materials for more than two decades, despite its shortcomings in terms of cost, environmental risk, limited capacity, and other issues.1 Among many potential candidates, recently developed LiNixCoyMn1-x-yO2 (hereafter referred to as NCM) compounds have exhibited well-balanced electrochemical performance and attracted significant attention as promising materials for high-capacity applications.2,3,4,5 In particular, Ni-rich NCM materials (x > 0.5), which characteristically have 40% higher capacity compared to LiCoO2, have been extensively studied in recent years.6,7 However, such Ni-based materials show significant structural degradation during cycling based on cation mixing, collateral reactions on surfaces, and poor thermal stability at high temperatures, eventually leading to capacity reduction.7 In addition to atomistic cation mixing that can be tuned via cation engineering,8 these unwanted phenomena are predominantly affected by the geometrical parameters of particles, such as grain size, porosity, and mesoscale compositional distribution inside secondary particles.9,10 To tackle these technical issues, coating active materials with various passivating substances has been proposed as an effective measure.11,12 Various oxides, such as Al2O3,13 ZnO,14 and MgO,15 can be utilized as coating materials, but the loss of capacity per unit mass of coating material is problematic. Based on core-shell design 3



Page 4 of 37

concepts, Sun et al. demonstrated that a core-shell cathode material consisting of a highcapacity

Ni-rich

core

(LiNi0.8Co0.1Mn0.1O2)

with

a

relatively

stable

shell

(LiNi0.5Mn0.5O2) shows excellent cyclability and thermal stability.16 This approach has been extended to several other systems, including Li-rich materials, to enhance the electrochemical performance of original core materials.17–20 However, when utilizing this typical core-shell method, the occurrence of structural defects at interfaces cannot be avoided because of crystallographic mismatches between the core and shell materials. For example, poor adhesion between the two different materials often leads to the formation of large voids on the order of several nanometers along the interface, which can severely inhibit the transport of lithium ions and electrons during cycling. As a countermeasure, the concept of a concentration gradient structure, where the concentration of transition metal cations gradually decreases or increases from the surfaces to the interiors of particles without a sudden change in the crystal lattice framework, has been proposed.21,22,23 This idea has been adopted by several other research groups to develop various core-shell structured cathode materials for highperformance Li ion batteries.24–28 The primary particles synthesized by either the coreshell or concentration gradient method often exhibit uniform and large sizes on the order of micrometers, which suppresses the formation of a solid electrolyte interface layer during charging/discharging cycles, which can lead to long life stability. Particle size is one of the critical structural parameters that determines the overall electrochemical performance of active materials because capacity and stability strongly depend on particle size configuration, even if they have a particular theoretical capacity calculated based on chemical composition.29–31 Generally, the smaller the particle size, the shorter the paths for Li+ ions and electrons, which improves rate 4


Page 5 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


capability. Furthermore, smaller particles promote an increase in the volumetric mass density of an active material because they can be more efficiently packed in a cell and leave space for fewer structural defects. However, the disadvantage of a small particle size arises from the fact that cycle life can be reduced because the surface area on which side reactions can take place is relatively large. Significant efforts have been devoted to synthesizing diverse hierarchical structures for secondary particles, such as cubes,32,33 hollows,34,35 dumbbells,36 and microrods,37 with the goal of alleviating the undesirable side effects resulting from the downsizing or nanosizing of primary particles. Although hierarchical structure engineering has been somewhat successful, the associated synthesis methods are relatively complicated because more elaborate processing steps are required. Additionally, given that all the primary particles forming secondary structures in such studies are similar in size with unimodal distributions, there is still room for further improvements to optimize such approaches from the perspective of primary particle size. In this work, we present a facile and resource-efficient synthesis method to form a hierarchically core-shell-type secondary microstructure for lithium transition metal oxide (LiNi0.6Co0.2Mn0.2O2, hereafter referred to as NCM622) cathode materials with a bimodal size distribution of primary particles. The key concept is to maximize the beneficial effects arising from specific sizes of primary particles by engineering small particles in the core structure for enhanced capacity and rate capability, and large particles in the shell structure for improved cycle life. Our study revealed that the proposed microstructure design can significantly boost the capacity and stability of NCM622 active materials compared to its commercial counterpart. Furthermore, because there is no need for either chemical modification or heterostructure engineering, 5



in contrast to conventional core-shell cathode materials, the proposed high-performance design can be implemented by simple tuning of the parameters for the traditional coprecipitation process and subsequent microwave calcination, meaning the proposed design is readily applicable to the mass production of such materials. To elucidate the structural origins of enhanced battery performance, we examined multiple aspects of material structure from the atomic scale to the microscopic scale based on a multimodal microscopy approach utilizing electron tomography and scanning transmission electron microscopy (STEM) with electron energy loss spectroscopy (EELS) and energy dispersive X-ray spectroscopy (EDX), as well as conventional characterization tools, such as X-ray diffraction (XRD).

2. RESULTS AND DISCUSSION During the stage of precursor synthesis, the stirring rate directly affects the size of primary particles.38 Increasing the stirring rate can intensify collisions between particles, thereby inhibiting continuous grain growth. As shown in Figure S1, the size the of precursor crystallites synthesized at a stirring rate of 1500 rpm is very small (~50 nm) compared to that at a stirring rate of 1000 rpm (~150 nm). Additionally, the shape of the crystals is strongly dependent on the stirring rate. Fast stirring (1500 rpm) creates needle-like crystallites, whereas slow stirring (1000 rpm) creates plate-like crystallites. By exploiting these phenomena, we can control the size of the primary precursor particles for the large secondary particles as a function of stirring rate. Therefore, we utilized a high stirring rate for the formation of small particles for the core structure during the initial period, followed by a low stirring rate for the formation of large particles for the shell structure of the NCM622 precursor. Three representative precursor particles for active material synthesis were obtained, as shown in Figure S2. The sample labeled 6


Page 6 of 37

Page 7 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


as P1 is a conventional precursor for reference, which was produced by utilizing a constant stirring speed of 1000 rpm for 34 h. The two samples with bimodal core-shell structures labeled as P2 and P3 were synthesized at a high stirring rate of 1500 rpm for 0.5 and 2 h, respectively, for the growth of the core structure. Next, the same stirring rate as that for P1 was applied to samples P2 and P3 for 33.5 and 32 h, respectively. The total stirring time was fixed as 34 h for all samples to create secondary particles with similar shapes and sizes. However, based on the different stirring rates applied during the initial growth period, the internal distributions of the primary particles differ significantly from each other. In contrast to P1, which has primary particles similar to the surrounding secondary particles, the core regions of P2 and P3 are less dense and filled with relatively small primary particles compared to their secondary particles. However, the inner portion of P3 is approximately twice as large as that of P2. This result indicates that the duration of the initial fast stirring period directly determines the final size of the primary particles in the core structure. High-resolution transmission electron microscopy (HRTEM) and selected area electron diffraction (SAED) analysis results revealed that the primary particles have the same crystal structure of Ni0.6Co0.2Mn0.2(OH)2 (P𝟑m1), regardless of their location in the core or shell, but the core region consists of needle-like fine grains with a width of ~50 nm or less, whereas the grain width in the shell region is greater than ~100 nm (Figure S3). This observation was supported by XRD analysis in which all three precursors (P1, P2, and P3) were identified as trigonal Ni0.6Co0.2Mn0.2(OH)2, which is similar to a pure Ni(OH)2 structure (Figure S4).39,40 This result indicates that Co2+ and Mn2+ ions partially replaced Ni2+ in the Ni(OH)2 structure.40 The mixture of Ni0.6Co0.2Mn0.2(OH)2 and LiOH·H2O was irradiated with microwaves to synthesize NCM622. Hereafter, the calcined particles resulting from the microwave heating of P1, P2, and P3 are referred to as C1, C2, and C3, respectively. Figure 1 presents cross7



sectional SEM images of the three NCM622 secondary particles. The C1 sample exhibits the typical appearance for NCM622 particles in commercial use, with a large primary particle size of more than 500 nm (Figure 1a). The C3 sample appears to have a hollow structure (Figure 1c) based on the very small primary precursors with large empty spaces formed by the high stirring rate during the initial period. In contrast, the C2 sample shows a well-defined coreshell-type bimodal microstructure with small primary particles (~100 nm or smaller) in the core and large primary particles (~500 nm or larger) in the shell (Figure 1b). SEM-EDX elemental analysis of the C2 sample revealed that the elemental composition of Ni, Co, and Mn is nearly constant across all locations inside the secondary particles (Figure S5). It should be noted that although their internal structures significantly different, the surface particle morphologies of the three secondary particles are very similar. (Figure S6). Through XRD analysis (Figure S7), we confirmed that all three samples have the same layered hexagonal α-NaFeO2 structure (R𝟑 m) without the formation of any secondary phase. Furthermore, we observed clear splits between the (006)/(102) fine peaks in their XRD patterns. This indicates that the samples have similar highly ordered structures.41,42 Figure 2a presents an annular bright field (ABF) STEM image of a C2 particle aligned with the [010] zone axis. One can see that all atomic columns of Li and O, as well as those of transition metals (TM), are well separated in this imaging mode, where the weakly scattered signals from light elements are effectively collected.43,44 The ABF-STEM image reveals an excellent match with the atomic model of the hexagonal layered structure discovered by XRD analysis and no evidence of cation disorder, suggesting that the sample structure is highly ordered. To directly probe the elemental distribution inside the ordered structure, we examined the C2 sample utilizing the atomic-scale STEM-EDX technique.45-47 The atom-resolved EDX maps (Figures 2b–e) for Ni (yellow), Co (green), Mn (blue), and O (red), which were obtained 8


Page 8 of 37

Page 9 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


along with the ABF-STEM image, clearly indicate that the doping elements of Co and Mn reside exclusively in the Ni sites without any cation intermixing with Li. This elemental order can be more clearly visualized in the color composite map with an overlaid atomic model (Figure 2f). To obtain information regarding the 3D nature of the three samples (C1, C2, and C3), we utilized the serial section tomography technique, which collects stacks of 2D images that can be aligned to construct a 3D volume for a sample.48 This technique is useful for visualizing the overall 3D morphology of secondary particles and analyzing the 3D distributions of geometrical parameters, such as local primary grain size and porosity. Figures 3a–c present the sectional 3D volumes of C1, C2, and C3, respectively. One can see that the three samples clearly differ in their distributions of primary grain size inside the secondary particles, similar to Figure 1. By performing segmentation analysis on the reconstructed volumes of the secondary particles, the total porosities (V) of the three samples were estimated to be VC1 = 17%, VC2 = 23%, and VC3 = 28%, which indicates that total porosity tends to increase as the stirring rate and duration during the stage of precursor preparation increase. The total porosity was calculated as the ratio of empty space over the volume filled by primary grains in the reconstructed volumetric data. The cross-sectional views of the three reconstructed volumes of the secondary particles clearly show different multigrain structures for the three samples (Figures 3d–f). The C1 particle has a typical random and dense multigrain structure and the C3 particle has a large void in the center. In contrast, the C2 particle has a bimodal multigrain structure with fine particles in the core and large particles in the shell. From the perspective of internal pore distribution, the different grain size distributions are more clearly visualized in Figures 3g–i. The C1 particle has an randomly distributed pore structure (Figure 3g). However, the pore structures of samples C2 and C3 are bimodally distributed with a larger number of 9



pores in the center of the core compared to the outside of the core (dotted circles in Figures 3h and 3i, respectively). Live rotation views of the three reconstructed pore structures are provided in supporting movies S1–3. To evaluate the geometrical parameters of the large secondary particles, such as local porosity and grain size distribution, we segmented a rectangular block with dimensions of 5.6 (X) × 4.2 (Y) × 8.8 (Z) μm3 for each reconstructed secondary particle volume, then digitally disjoined the block into pores (empty spaces) and rendered grains (primary particles). Next, we sliced the segmented volume data by utilizing an arbitrary unit slice as a function of relative radius (from the center (0) to the radius (R)) to examine the variations in local primary grain size (Figure 4a) or as a function of distance from the center (Figure 4b) to investigate local changes in porosity inside the secondary particles. For the local variation of primary grain size (Figure 4a), sample C1 showed a nearly uniform distribution of primary particle sizes (0.746 μm on average) throughout the entire secondary particle. In contrast, samples C2 and C3 show significant variation in primary particle size inside the secondary particle. Their primary particle sizes substantially increased from the core (0–0.25R) to the shell (0.5–0.75R), where R is the radius of the secondary particle. It should be noted that nanograins less than 300 nm in size fill the core of sample C2, but there is substantial empty space in the core of sample C3. Based on the 3D volume data in Figures 3d-f, one can see that the sizes of the primary particles of the two samples in the shell region are relatively large and elongated in the radial direction of the secondary particles. Figure 4b presents the local changes in porosity per unit slice (set as 5.6 (X) × 4.2 (Y) × 0.15 (Z) μm3) when moving from the center to the outside of a secondary particle. The commercial C1 sample shows a random distribution of local pores (black line). In contrast, the local porosity profile of sample C3 (blue curve) indicates that local pores are clustered to form a large void at the center of the sample, which can lead to a loss of capacity 10


Page 10 of 37

Page 11 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


and structural stability. In the profile of sample C2 (red curve), one can see that the pore distribution is similar to that of sample C3 sample, but there is only a difference of less than 5% in terms of local porosity between the core and shell regions, suggesting that the small dense pores (or particles) in the center are closely connected. This result indicates that sample C2 has a well-defined bimodal microstructure with fine nanoparticles in the core. Through the simple alteration of stirring rate during the synthesis process, we have demonstrated that the internal microstructures of secondary particles can be readily tailored and optimized to create a bimodal microstructure without chemical alteration. Given the geometrically optimized bimodal microstructure of the C2 sample, we anticipate major improvement in battery performance in terms of capacity and stability because the outer shell with a large grain size can effectively inhibit the continuous outbreak of side reactions with electrolytes on the surface and the fine nanoparticles in the core promote Li ion diffusion during charging and discharging processes. To evaluate the rate performance of NCM622 cathode materials, cells manufactured from samples C1, C2, and C3 were tested with increasing current densities (0.1−10.0C, 1C = 165 mAh/g). For samples C1 and C3, as shown in Figure 5a, the measured capacities were smaller than that of C2 at a low discharge rate of 0.1C and degraded significantly as the discharge rate increased. The C2 sample outperformed the other two samples in terms of total capacity. The initial capacity of sample C2 was 182.8 mAh/g at 0.1C and gradually fell to 180.5, 175.1, 169.9, and 159.5 mAh/g as the discharge rate increased from 0.5C to 10.0C, but maintained the highest value among all samples. The measured capacities of the three samples are tabulated in Table S1 for comparison. The capacity reduction at a discharge rate of 1.0C as a function of the number of cycles was measured and plotted in Figure 5b. Once can see that the initial capacity of sample C2 sample (175.1 mAh/g) gradually decreased to 145.8 mAh/g after 100 cycles. Compared to the values of samples C1 (123.9 11



mA/g) and C3 (123.5 mAh/g) after 100 cycles, sample C2 provides an increase of approximately 18% in terms of restored capacity. Overall, sample C2 showed 53.8% greater capacity at a high discharge rate (10C) and comparable capacity retention after 100 charge/discharge cycles (1.0C discharge rate) compared to the commercial sample (C1). It should be noted that in terms of capacity retention ratio, sample C2 sample (83.3%) is similar to samples C1 and C3 (80.9% and 78.0%, respectively). We also measured the capacity stability of the three samples at a low discharge rate of 0.1C and the sample C2 showed slightly better stability (95.7%) as compared to the other samples C1 and C3 (89.4% and 88.4%, respectively) (Figure S8). This result can be attributed to similarities in the sizes of secondary particles and relative surface areas between the three samples because the ratio of capacity retention is largely dependent on surface morphology and the relative surface area of active materials with which electrolyte side reactions can occur. If further progress is made in the development of surface passivation techniques to suppress unwanted side reactions, we believe that capacity retention for long-term cyclic performance can be significantly improved. Given the homogeneity in secondary particle size, chemical composition, and crystal structure between the three samples, the superiority of sample C2 in terms of electrochemical performance can be attributed to the small primary particles connected throughout the core regions of the secondary particles. Figure 6 presents cyclic voltammograms of samples C1 and C2 at different scan rates ranging from 0.01–1.00 mV s-1. As reported in previous studies on lithium-mixed transition metal oxides,34,49,50 all curves exhibited only one redox pair in the range of 3.4−4.4 V, indicating that our NCM622 materials experienced one structural transformation (i.e., hexagonal to monoclinic) during the charge/discharge cycles. One can clearly see that the hysteresis between the oxidation and reduction of C1 is more significant than that of C2. This 12


Page 12 of 37

Page 13 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


characteristic suggests that sample C2 provides better reversibility of the Li+ intercalation/deintercalation process and smaller electrode polarization.51 With an increasing scan rate, the reduction peak shifts to the low-potential region, whereas the oxidation peak moves to the high-potential region. To estimate the diffusion coefficient of the lithium ions in the cathode materials, the Randles–Sevcik equation was utilized52

𝐼𝑝 =

1 3 1 1 𝐹3 2 2 2 0.4463(𝑅𝑇) 𝑛 𝐴𝐷 𝐶0𝜈 2

······································ (1)

where Ip is the peak current, R is the gas constant, T, is the absolute temperature, F is the Faraday constant, n is the number of electrons transferred per molecule, A is the active surface area of the electrode (0.50 cm2), C0 is the concentration of Li+ ions in the cathode (9.8 × 10−3 mol/cm3), D is the apparent ion diffusion coefﬁcient, and ν is the scan rate. From the slope of plots of Ip and ν1/2, the apparent diffusion coefﬁcient D can be calculated as shown in Figure S9. We found that D value of sample C2 is 3−4 times larger than that of the commercial sample C1. The faster lithium ion diffusion of sample C2 is directly related to its superior electrochemical performance in terms of capacity and rate capability. The active lithium content participating in discharging (lithiation) or charging (delithiation) determines real battery capacity and the oxidation of TMs varies accordingly during the battery operation process, as shown in equation (2) and (3). The oxidation number of TMs increases during charging (delithiation) because the increased oxidation number is equal to the number of lithium ions (or electrons) that escape from the active material. 𝐿𝑖𝑇𝑀𝑛𝑎𝑂𝑏 = 𝐿𝑖1 ― 𝑥𝑇𝑀𝑛𝑎 + 𝑥𝑂𝑏 + 𝑥𝐿𝑖 + + 𝑥𝑒 ― ······························· (2) 𝐶𝑃 = 𝑥𝐹·················································· (3) where, Cp is the actual capacity of the cell, F is the faraday constant, and x is the number of 13



moles of electrons produced by the reaction during discharging. Given the above equations, a larger change in the oxidation number during the reaction indicates that an active material has a greater capacity. Therefore, by visualizing the changes in the oxidation states of TMs in the primary particles of samples C1 and C2, we can directly explain why sample C2 shows superior electrochemical properties. To monitor changes in the oxidation states of TMs, such as Ni, Co, and Mn in the primary particles, we performed EELS spectrum imaging (SI) utilizing a STEM imaging mode for samples C1 and C2 in their charged (delithiated) and discharged (lithiated) states. In the EELS SI data, the shape and position of the L peaks of TMs that originate from transitions of the 2p core electrons to their final states with s and 3d orbital components vary sensitively with changes in the oxidation state and local bonding environment.53-55 This characteristic can be utilized as an indicator of the charge state change in battery materials.47,56 To map the spatial distribution of the oxidation states of the TMs in the primary particles of samples C1 and C2, we measured the intensity ratio of the L3 and L2 peaks of the L edges of Ni, Co, and Mn from the EELS SI data cube,57 which can be performed utilizing an open-source software package called “Oxide Wizard.”58 The reliability of the charge state measurements was verified by performing the same EELS analysis routine on known samples.47 The representative EEL spectra of Ni, Co, and Mn for samples C1 and C2 in their charged/discharged states are presented in Figure S10 for the sake of clear comparison. The resulting valence state maps are presented in Figure 7. As expected, prior to charging, the oxidation state of each transition metal for the two samples was found to be Ni2+, Co3+, and Mn4+ throughout the entire region of interest. In contrast, differences between the two samples emerged in post-charge measurements. For sample C1 (Figure 7a), the Ni shows a mixed valence state of 3+ and 4+ and the oxidation states of Co and Mn were nearly unchanged on average, remaining as Co3+ 14


Page 14 of 37

Page 15 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


and Mn4+, although a slightly higher oxidation state of Co3.5+ is detected in the region near the surface (denoted by the arrow in Figure 7a). In contrast, the charge states of the TMs in sample C2 are significantly differently (Figure 7b). The oxidation state of the Ni ions increased significantly to a final value of 4+ and the Co ions were oxidized to 3.5+ on average (see the maps outlined with green squares in Figure 7b), whereas the Mn ions exhibited no discernable change in their valence state over the entire particle, in contrast to sample C1. These results demonstrate that sample C2 has a greater charge density compared to sample C1 because the amount of extracted lithium ions is proportional to the increase in the charge state of the TMs. To directly compare the extent of lithium extraction between samples C1 and C2 in the charged state, we employed atomic-resolution high-angle annular dark field (HAADF) and ABF-STEM imaging modes. By utilizing both imaging modes simultaneously, we can clearly visualize the entire atomic configurations of the samples based on the atomic number dependency of the contrast in the HAADF imaging mode and imaging capability for light elements of ABF imaging modes.43,44 The significant difference in capacity between samples C1 and C2 can be also corroborated by tracking Li ions before and after charging from both atomic-scale STEM images. Figures 8a and 8b present a pair of HAADF (left) and ABF (right) images of samples C1 and C2, respectively, aligned with the [010] zone axis prior to charging. From these STEM images, we confirmed that the two samples have the same layered hexagonal α-NaFeO2 structure (R𝟑m) without cation disorder and show partial site exchange between Li and TM ions. All Li and TM ions are located in a cubic close-packed lattice framework of oxygen atoms, forming ordered octahedral sites in the corresponding slabs. The overlaid atomic model for the [010] orientation represents an excellent match to the STEM images. After charging, notable differences between the two samples can be observed in the STEM images, as shown in Figures 8c and 8d. Regarding the TM lattice framework, both samples still exhibit 15



a well-defined layered structure without cation intermixing between Li and TM ions (see left HAADF images). However, in the ABF STEM images, one can see different image contrast levels between the Li layers of both samples. The contrast between Li sites appears distinctively weak in sample C2 (Figure 8d, right panel) compared to that before charging (Figure 8b, right panel). However, strong contrast can still be observed for sample C1 after charging (Figures 8a and 8c, right panels). This result indicates that delithiation (charging) is much facile in the core-shell-type sample C2 compared to the commercial sample C1. The (reversed) intensity profiles of the Li sites can more clearly visualize the different delithiation behaviors, as shown in Figures 8e and 8f. Given that the intensities of most Li sites in sample C2 are significantly reduced after charging (Figure 8f), it is clear that Li migration in the C2 primary particles occurs more readily during charging compared to sample C1, which leads to enhanced charge capacity and cyclability. Furthermore, as indicated by the 3D tomographic reconstruction (Figure 3), we expect that this advantageous fast reaction in sample C2 can be sustained for many cyclic operations, despite the small size of the primary particles. This is because empty space in the core is minimized and the small primary particles are densely interconnected by shell structures containing relatively large primary particles. It should be noted that there is no one-to-one correspondence between Li sites before and after charging because we observed different sample regions between processes. The valence state of the TMs can also be changed via cation mixing and Li extraction.47 The cation mixing of Ni and Co into the Li layer can lead to phase transition to an ineffective and impotent rock-salt or spinel structure for further Li insertion/extraction, which has been recognized as a harmful phenomenon, rather than simply degrading electrochemical performance. However, from the STEM-EDX mapping, we observed no discernable cation disorder inside the primary particles of the two samples after charging (see Figure S11). 16


Page 16 of 37

Page 17 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


3. CONCLUSION In this work, we demonstrated that the capacity and cyclability of cathode materials can be significantly improved through the simple geometrical design of NCM cathode materials without the need for complicated chemical modification or heterostructure engineering. By synthesizing hierarchically-structured secondary NCM622 particles with a bimodal size distribution of primary particles, the electrochemical performance was vastly superior to that of a commercial NCM622 material. This resource-efficient approach can be easily implemented through simple adjustments of stirring rates during the early stages of precursor synthesis and subsequent microwave calcination, which is advantageous for mass-production applications. Overall, we determined that the cell performance of the modified NCM622 sample (C2) could be attributed to the small size of the primary particles in the core, which allow for facile and fast lithium insertion/extraction through the layered structure based on shorter migration paths. Our study clearly demonstrates that geometrical parameters are critical to the electrochemical performance of NCM cathode materials and provides novel insights regarding the resource-efficient design of high-performance Li ion batteries at a mesoscale.

4. METHODS Materials Synthesis. First, a Ni0.6Co0.2Mn0.2(OH)2 compound was synthesized as a precursor via the co-precipitation method. NiSO4·6H2O, CoSO4·7H2O, and MnSO4·H2O were utilized as starting materials and dissolved in distilled water with a molar ratio of Ni:Co:Mn = 6:2:2. The total concentration of transition metal ions in the solution was set to 1.5 M. Under a nitrogen atmosphere, the solution was slowly pumped into a tank reactor under continuous stirring at 50 °C. Simultaneously, an aqueous NaOH solution (4 M) was injected into the reactor to maintain 17



the pH of the solution within a certain range (10–11). Additionally, an aqueous NH4OH solution was also pumped in as a chelating agent at a rate of 0.96 L/h. The retention time in the reactor, which determines the degree of sphericity and growth rate of the precursor particles, was 34 h. The size of the is primary particles is determined by the stirring rate. To prepare the precursor for the reference material, the stirring rate was maintained at 1000 rpm (agitation power of 2.5 kW) throughout the entire reaction. To decrease the primary particle size in the core, one can speed up the initial stirring rate. The optimized stirring rate was found to be 1500 rpm (3.0 kV) for the formation of the core of the precursor. The stirring rate was then reduced to the same rate as that for the reference condition (1000 rpm) during the second stage to form a shell. The duration of the initial high-speed stirring can be adjusted (0.5–2 h). After filtering and washing the solution to remove residual ions, such as Na+ and SO42−, the precipitate was dried at 80 °C for 12 h. The obtained precursors of Ni0.6Co0.2Mn0.2(OH)2 and LiOH·H2O were thoroughly mixed with a molar ratio of 1:1.02 (TM:Li) in an agate mortar. For calcination of the powder mixtures, which were loaded onto an alumina boat, we employed the microwave irradiation method, which allowed us to synthesize the materials quickly and efficiently, despite the relatively low crystallinity of the calcined products.59 The temperature was raised to 600 °C within 5 min (> 100 °C min-1), then maintained at 600 °C for 10 min under the microwave irradiation (2.45 GHz and 1200 W). The calcined powder was then rapidly cooled to room temperature in the microwave furnace at a rate of 100 °C min-1.

Characterizations of Microstructures. The chemical compositions of the samples were analyzed via inductively coupled plasma-atomic emission spectrometry (ULTIMA2, HORIBA 18


Page 18 of 37

Page 19 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


JOBIN YVON, France). XRD analysis was performed utilizing a PANalytical X’pert MPD diffractometer with copper Kα radiation (λ = 1.54056 Å). The morphologies of the samples were observed via field-emission SEM (Magellan XHR, FEI Co., Hillsboro, OR, USA). To create thin cross-sectional TEM samples for atomic-scale STEM analysis, the focused ion beam (FIB, Auriga CrossBeam Workstation, Carl Zeiss) milling and lift-out techniques were utilized. Subsequently, low-kV Ar surface milling at 700 V (Fischione Model 1040 NanoMill) was performed for 15 min to remove damaged surface layers created by heavy Ga ion beam milling in the FIB system. Atomic structure images of the samples were captured by an 80-kV aberration-corrected STEM (JEM-ARM200CF, JEOL) with a probe forming semi-angle of ~23 mrad. The detector angle ranges of the HAADF and ABF imaging modes were 70-175 and 7.5-17 mrad, respectively. Background noise in the acquired STEM images was reduced by the Wiener filtering method implemented in a commercial software package (HREM Filter Pro, HREM research Ltd.). Atomic-resolution elemental maps of the samples were acquired by an EDX spectrometer (JED-2300T, Jeol Ltd.) equipped with a dual-type detector with a large effective solid angle (~1.2 sr) and highly focused electron probe (~1.1 Å) in the STEM imaging mode. EELS data were acquired by a Gatan Quantum spectrometer (ER965) attached to the microscope. The EELS SI data for the samples were acquired in dualEELS acquisition mode to obtain full energy loss information over both the low-loss and highloss ranges with an energy dispersion of 0.25 eV/ch and probe dwell time of 2.00 s/pixel. The average thickness of samples was measured to be ~58 ± 1.2 nm by the EELS log-ratio method.38 To perform 3D tomography on the samples, a serial section imaging technique was adopted utilizing a Ga+ FIB-SEM (Auriga CrossBeam Workstation, Carl Zeiss) at 30 kV. The thickness interval for serial sectioning was set to 30 nm and 400 slices were recorded to form a 3D volumetric dataset for each sample. 3D visualizations of the serial section datasets were 19



created utilizing an open-source software package called Tomviz60 and quantitative analysis was performed to obtain the local porosity and grain size distributions of the 3D tomography datasets utilizing a non-commercial version of a software package called Dragonfly 3.5.61

Cell Performance Test. To evaluate the electrochemical performance of the synthesized materials, 2032-type coin-cells were prepared in a dry room. The active materials, conducting materials (Super-P carbon black), and binders (polyvinylidene fluoride) were blended (96:2:2 weight ratio) in N-methyl-2-pyrrolidone. The slurry was then coated onto an aluminum foil and dried at 110 °C for 10 h in a vacuum oven. The electrolyte solution was 1.15-M LiPF6 in an ethylene carbonate/dimethyl carbonate/ethylmethyl carbonate mixture (3:3:4 volume ratio). Lithium foil was utilized as a counter-electrode and the capacity of the manufactured cells was evaluated to be 1.5 mAh/cm2. The capacities of all samples were estimated to be between 3.0 and 4.3 V. The cells were charged galvanostatically (constant current mode) in the first cycle, then cycled repeatedly in constant voltage mode. The loading density of the cathode materials was 20.0 ± 2.0 mg/cm2.

20


Page 20 of 37

Page 21 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


ASSOCIATED CONTENT Supporting information The Supporting Information is available free of charge on the ACS Publications website. Figures S1-S10 (PDF), Movie S1-S3.

AUTHOR INFORMATION Corresponding Author *E-mail (Y.-M. Kim): [email protected] *E-mail (S. Lee): [email protected] Author Contributions ∥J.-H.

S. and Y.-H. K. contributed equally to this work.

Notes The authors declare no competing financial interest.

AKNOWLEDGEMENTS S. Lee and J.-H. Shim acknowledge financial support from the National Research Foundation of Korea (NRF-2016R1D1A1B03932245) and (NRF-2017R1D1A1B03028199), respectively. Y.-M.K. acknowledges financial support by the Institute for Basic Science (IBS-R011-D1) and Creative Materials Discovery Program (NRF-2015M3D1A1070672) through the NRF grant.

21



REFERENCES (1) Fergus, J. W. Recent Developments in Cathode Materials for Lithium Ion Batteries. J. Power Sources 2010, 195, 939−954. (2) Wang, L.; Li, J.; He, X.; Pu, W.; Wan, C. Jiang, C. Recent Advances in Layered LiNixCoyMn1-x-yO2 Cathode Materials for Lithium Ion Batteries. J. Solid State Electrochem. 2009, 13, 1157−1164. (3) Robert, R.; Villevieille, C.; Novák, P. Enhancement of the High Potential Specific Charge in Layered Electrode Materials for Lithium-Ion Batteries. J. Mater. Chem. A 2014, 2, 8589−8598. (4) He, P.; Yu, H.; Li, D.; Zhou, H. Layered Lithium Transition Metal Oxide Cathodes Towards High Energy Lithium-Ion Batteries. J. Mater. Chem. 2012, 22, 3680−3695. (5) Liu Y.; Ning D.; Zheng L. Zhang Q.; Gu L.; Gao R.; Zhang J.; Franz A.; Schumacher G.; Liu X. Improving the electrochemical performances of Li-rich Li1.20Ni0.13Co0.13Mn0.54O2 through a cooperative doping of Na+ and PO43- with Na3PO4. J. Power Sources 2018, 375, 1-10. (6) Xu, J.; Lin, F.; Doeff, M. M.; Tong, W. a Review of Ni-Based Layered Oxides for Rechargeable LiIon Batteries. J. Mater. Chem. A 2017, 5, 874−901. (7) Liu, W.; Oh, P.; Liu, X.; Lee, M.-J.: Cho, W.; Chae, S.; Kim, Y. Cho, J. Nickel-Rich Layered Lithium Transition-Metal Oxide for High-Energy Lithium-Ion Batteries. Angew. Chem., Int. Ed. 2015, 54, 4440−4457. (8) Kim, N. Y.; Yim, T.; Song, J. H.; Yu, J.-S.; Lee, Z. Microstructural Study on Degradation Mechanism of Layered LiNi0.6Co0.2Mn0.2O2 Cathode Materials by Analytical Transmission Electron Microscopy. J. Power Sources 2016, 307, 641-648. (9) Yan, P.; Zheng, J.; Gu, M.; Xiao, J.; Zhang, J.-G.;Wang, C.-M. Intragranular Cracking as a Critical Barrier for High-voltage Usage of Layer-structured Cathode for Lithium-ion Batteries. Nature Comm.

22


Page 22 of 37

Page 23 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


2017, 8, 14101-14109. (10) Hausbrand, R.; Cherkashinin, G.: Ehrenberg, H.:Gröting, M.; Albe, K.: Hess, C.; Jaegermann, W. Fundamental Degradation Mechanisms of Layered Oxide Li-ion Battery Cathode Materials: Methodology, Insights and Novel Approaches. Mater. Sci. Eng. B 2015, 192, 3-25. (11) Nitta, N.; Wu, F.; Lee, J. T.; Yushin, G. Li-Ion Battery Materials: Present and Future. Materials Today 2015, 18, 252−264. (12) Ellis, B. L.; Lee, K. T.; Nazar, L. F. Positive Electrode Materials for Li-Ion and Li-Batteries. Chem. Mater. 2010, 22, 691−714. (13) Myung, S.-T.; Izumi, K.; Komaba, S.; Sun, Y.-K.; Yashiro, H.; Kumagai, N. Role of Alumina Coating on Li-Ni-Co-Mn-O Particles as Positive Electrode Material for Lithium-Ion Batteries. Chem. Mater. 2005, 17, 3695−3704. (14) Guo, R.; Shi, P.; Cheng, X.; Sun, L. Effect of ZnO Modification on the Performance of LiNi0.5Co0.25Mn0.25O2 Cathode Material. Electrochim. Acta 2009, 54, 5796−5803. (15) Yoon, W.-S.; Nam, K.-W.; Jang, D.; Chung, K. Y.; Hanson, J.; Chen, J.-M.; Yang, X.-Q. Structural Study of the Coating Effect on the Thermal Stability of Charged MgO-Coated LiNi0.8Co0.2O2 Cathodes Investigated by in situ XRD. J. Power Sources 2012, 217, 128−134. (16) Sun, Y.-K.; Myung, S.-T.; Kim, M.-H.; Prakash, J.; Amine, K. Synthesis and Characterization of Li((Ni0.8Co0.1Mn0.1)0.8(Ni0.5Mn0.5)0.2)O2 with the Microscale Core-Shell Structure as the Positive Electrode Material for Lithium Batteries. J. Am. Chem. Soc. 2005, 127, 13411−13418. (17) Myung, S.-T.; Lee, K.-S.; Kim, D.-W.; Scrosati, B.; Sun, Y.-K. Spherical Core-Shell Li((Li0.05Mn0.95)0.8(Ni0.25Mn0.75)0.2)2O4 spinel as high performance Cathodes for Lithium Batteries. Energy Environ. Sci. 2011, 4, 935−939. (18) Yang, X.; Wang, X.; Hu, L.; Zou, G.; Su, S.; Bai, Y.; Shu, H.; Wei, Q.; Hu, B.; Ge, L.; Wang, D.;

23



Liu, L. Layered Li(Ni0.5Co0.2Mn0.3)O2-Li2MnO3 Core-Shell Structured Cathode Material with Excellent Stability. J. Power Sources 2013, 242, 589−596. (19) Li, J.; Camardese, J.; Shunmugasundaram, R.; Glazier, S.; Lu, Z.; Dahn, J. R. Synthesis and Characterization of the Lithium-Rich Core-Shell Cathodes with Low Irreversible Capacity and Mitigated Voltage Fade. Chem. Mater. 2015, 27, 3366−3377. (20) Cho, Y.; Lee, S.; Lee, Y.; Hong, T.; Cho, J. Spinel-Layered Core-Shell Cathode Materials for LiIon Batteries. Adv. Energy Mater. 2011, 1, 821−828. (21) Sun, Y.-K.; Myung, S.-T.; Park, B.-C.; Prakash, J.; Belharouak, I.; Amine, K. High-Energy Cathode Material for Long-Life and Safe Lithium Batteries. Nat. Mater. 2009, 8, 320−324. (22) Sun, Y.-K.; Chen, Z.; Noh, H.-J.; Lee, D.-J.; Jung, H.-G.; Ren, Y.; Wang, S.; Yoon, C. S.; Myung, S.-T.; Amine, K. Nanostructured High-Energy Cathode Materials for Advanced Lithium Batteries. Nat. Mater. 2012, 11, 942−947. (23) Zhang J.; Yang Z.; Gao R.; Gu L.; Hu Z.; Liu X. Suppressing the Structure Deterioration of NiRich LiNi0.8Co0.1Mn0.1O2 through Atom-Scale Interfacial Integration of Self-Forming Hierarchical Spinel Layer with Ni Gradient Concentration. ACS Appl. Mater. Interfaces 2017, 9, 29794-29803. (24) Yang, X.; Wang, X.; Zou, G.; Hu, L.; Shu, H.; Yang, S.; Liu, L.; Hu, H.; Yuan, H.; Hu, B.; Wei, Q.; Yi, L. Spherical Lithium-Rich Layered Li1.13(Mn0.534Ni0.233Co0.233)0.87O2 with ConcentrationGradient Outer Layer as High-Performance Cathodes for Lithium Ion Batteries. J. Power Sources 2013, 232, 338−347. (25) Hou, P. Y.; Zhang, L. Q.; Gao, X. P. a High-Energy, Full Concentration-Gradient Cathode Material with Excellent Cycle and Thermal Stability for Lithium Ion Batteries. J. Mater. Chem. A 2014, 2, 17130−17138. (26) Wei, Q.; Wang, X.; Yang, X.; Ju, B.; Hu, B.; Shu, H.; Wen, W.; Zhou, M.; Song, Y.; Wu, H.; Hu, H. Spherical Concentration-Gradient LiMn1.87Ni0.13O4 Spinel as a High Performance Cathode for

24


Page 24 of 37

Page 25 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Lithium Ion Batteries. J. Mater. Chem. A 2013, 1, 4010−4016. (27) Du, K.; Hua, C.; Tan, C.; Peng, Z.; Cao, Y.; Hu, G. a High-Powered Concentration-Gradient LiNi0.85Co0.12Mn0.03)O2 Cathode Material for Lithium Ion Batteries. J. Power Sources 2014, 263, 203−208. (28) Hou, P.; Zhang, H.; Zi, Z.; Zhang, L.; Xu, X. Core-Shell and Concentration-Gradient Cathodes Prepared via Co-Precipitation Reaction for Advanced Lithium-Ion Batteries. J. Mater. Chem. A 2017, 5, 4254−4279. (29) Drezen, T.; Kwon, N.-H.; Bowen, P.; Teerlinck, I.; Isono, M.; Exnar, I. Effect of Particle Size on LiMnPO4 Cathodes. J. Power Sources 2007, 174, 949−953. (30) Okubo, M.; Hosono, E.; Kim, J.; Enomoto, M.; Kojima, N.; Kudo, T.; Zhou, H.; Honma, I. Nznosize Effect on High-Rate Li-Ion Intercalation in LiCoO2 Electrode. J. Am. Chem. Soc. 2007, 129, 7444−7452. (31) Li, J.; Daniel, C.; Wood, D. Materials Processing for Lithium-Ion Batteries. J. Power Sources 2011, 196, 2452−2460. (32) Wu, Y.; Cao, C.; Zhu, Y.; Li, J.; Wang, L. Cube-Shaped Hierarchical LiNi1/3Co1/3Mn1/3O2 with Enhanced Growth of Nanocrystal Planes as High-Performance Cathode Materials for Lithium-Ion Batteries. J. Mater. Chem. A 2015, 3, 15523−15528. (33) Wu, Y.; Zhang, J.; Cao, C.; Khalid S.; Zhao, Q.; Wang, R.; Butt, F. K. LiNi0.5Mn1.5O4 NanoSubmicro Cubes as High-Performance 5V Cathode Materials for Lithium-Ion Batteries. Electrochim. Acta 2017, 230, 293−298. (34) Li, J.; Cao, C.; Xu, X.; Zhu, Y.; Yao, R. LiNi1/3Co1/3Mn1/3O2 Hollow Nano-Micro Hierarchical Microspheres with Enhanced Performances as Cathodes for Lithium-Ion Batteries. J. Mater. Chem. A 2013, 1, 11848−11852.

25



(35) Lou, M.; Zhong, H.; Yu, H.-T.; Fan, S.-S.; Xie, Y.; Yi, T.-F. Li1.2Mn0.54Ni0.13Co0.13O2 Hollow Hierarchical Microspheres with Enhanced Electrochemical Performances as Cathode Material for Lithium-Ion Battery Application. Electrochim. Acta 2017, 237, 217−226. (36) Ryu, W.-H.; Lim, S.-J.; Kim, W.-K.; Kwon, H. 3-D Dumbbell-like LiNi1/3Mn1/3Co1/3O2 Cathode Materials Assembled with Nano-Building Blocks for Lithium-Ion Batteries. J. Power Sources 2014, 257, 186−191. (37) Yang, Z.; Lu, J.; Bian, D.; Zhang, W.; Yang, X.; Xia, J.; Chen, G.; Gu, H.; Ma, G. Stepwise Coprecipitation to Synthesize LiNi1/3Co1/3Mn1/3O2 One-Dimensional Hierarchical Structure for Lithium Ion Batteries. J. Power Sources 2014, 272, 144−151. (38) Liang, L,; Du, K.; Peng, Z.; Cao, Y.; Duan, J.; Jiang, J.; Hu, G. Co-Precipitation Synthesis of Ni0.6Co0.2Mn0.2(OH)2 Precursor and Characterization of LiNi0.6Co0.2Mn0.2O2 Cathode Material for Secondary Lithium Batteries. Electrochim. Acta 2014, 130, 82−89. (39) Kim, M-H.; Shin, H.-S.; Shin, D.; Sun, Y.-K. Synthesis and Electrochemical Properties of Li(Ni0.8Co0.1Mn0.1)O2 and Li(Ni0.8Co0.2)O2 via Co-Precipitation. J. Power Sources 2006, 159, 1328−1333. (40) Ying, J.; Wan, C.; Jiang, C.; Li, Y. Preparation and Characterization of High-Density Spherical LiNi0.8Co0.2O2 Cathode Material for Lithium Secondary Batteries. J. Power Sources 2001, 99, 78−84. (41) Woo, S.-U.; Yoon, C. S.; Amine, K.; Belharouak, I.; Sun, Y.-K. Significant Improvement of Electrochemical Performance of AlF3-Coated Li(Ni0.8Co0.1Mn0.1)O2 Cathode Materials. J. Electrochem. Soc. 2007, 154, A1005−A1009. (42) Li, J.; Wang, L.; Zhang, Q.; He, X. Synthesis and Characterization of LiNi0.6Mn0.4-xCoxO2 as Cathode Materials for Li-Ion Batteries. J. Power Sources 2009, 189, 28−33. (43) Wen, Y.; Shang, T.; Gu, L. Analytical ABF-STEM Imaging of Li Ions in Rechargeable Batteries. Microscopy. 2017, 66, 25−38.

26


Page 26 of 37

Page 27 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


(44) Pennycook, S. J.; Li, C.; Li, M.; Tang, C.; Okunishi, E.; Varela, M.; Kim, Y.-M.; Jang, J. H. Material structure, properties, and dynamics through scanning transmission electron microscopy. J. Anal. Sci. Tech. 2018, 9, 11-24. (45) Yan, P.; Zheng, J.; Lv, D.; Wei, Y.; Zhen, J.; Wang, Z.; Kuppan, S.; Yu, J.; Luo, L.; Edwards, D.; Olszta, M.; Amine, K.; Liu, J.; Xiao, J.; Pan, F.; Chen, G.; Zhang, J.-G.; Wang, C.-M. AtomicResolution Visualization of Distinctive Chemical Mixing Behavior of Ni, Co, and Mn with Li in Layered Lithium Transtion-Metal Oxide Cathode Materials. Chem. Mater. 2015, 27, 5393−5401. (46) Yan, P.; Zheng, J.; Zhang, J.-G.; Wang, C.-M. Atomic Resolution Structural and Chemical Imaging Revealing the Sequential Migration of Ni, Co, and Mn upon the Battery Cycling of Layered Cathode. Nano lett. 2017, 17, 3946−3951. (47) Shim, J.-H.; Im, J.-S; Kang, H.; Cho, N.; Kim, Y.-M.; Lee, S.; Implications of Cation-Disordered Grain Boundary on the Electrochemical Performance of LiNi0.5Co0.2Mn0.3O2 Cathode Material for Lithium Ion Batteries. J. Mater. Chem. A 2018, 6, 16111-16120. (48) Burnett, T. L.; Kelly, R.; Winiarski, B.; Contreras, L.; Daly, M.; Gholinia, A.; Burke, M. G.; Withers, P. J. Large volume serial section tomography by Xe plasma FIB dual beam microscopy. Ultramicroscopy 2016, 161, 119-129. (49) Wang, D.; Li, X.; Wang, Z.; Guo, H.; Xu, Y.; Fan, Y. Co-Modification of LiNi0.5Co0.2Mn0.3O2 Cathode Materials with Zirconium Substitution and Surface Polypyrrole Coating: Towards Superior High Voltage Electrochemical Performances for Lithium Ion Batteries. Electrochim. Acta 2016, 196, 101−109. (50) Li, J.; Yao, R.; Cao, C. LiNi1/3Co1/3Mn1/3O2 Nanoplates with {010} Active Planes Exposing Prepared in Poylol Medium as a High-Performance Cathode for Li-Ion Battery. ACS Appl. Mater. Interfaces 2014, 6, 5075−5082. (51) Heinze, J. Cyclic Voltammetry−“Electrochemical Spectroscopy”. Angew. Chem., Int. Ed. 1984,

27



23, 831−918. (52) Nicholson, R. S.; Shain, I. Theory of Stationary Electrode Polarography. Single Scan and Cyclic Methods Applied to Reversible, Irreversible, and Kinetic Systems. Anal. Chem. 1964, 36, 706−723. (53) Cavé, L.; Al, T.; Loomer, D.; Cogswell, S.; Weaver, L. A STEM/EELS Method for Mapping Iron Valence Ratios in Oxide Minerals. Micron 2006, 37, 301−309. (54) Tan, H.; Turner, S.; Yücelen, E.; Verbeeck, J.; Van Tendeloo, G. 2D Atomic Mapping of Oxidation States in Transition Metal Oxides by Scanning Transmission Electron Microscopy and Electron EnergyLoss Spectroscopy. Phys. Rev. Lett. 2011, 107, 107602. (55) Varela, M.; Oxley, M. P.; Luo, W.; Tao, J.; Watanabe, M.; Lupini, A. R.; Pantelides, S. T.; Pennycook, S. J. Atomic-Resolution Imaging of Oxidation States in Manganites. Phys. Rev B 2009, 79, 085117. (56) Cosandey, F.; Su, D.; Sina, M.; Pereira, N.; Amatucci, G. G. Fe Valence Determination and Li Elemental Distribution in Lithiated FeO0.7F1.3/C Nanocomposite Battery Materials by Electron Energy Loss Spectroscopy (EELS). Micron 2012, 43, 22−29. (57) Pearson, D. H.; Ahn, C. C.; Fultz, B. White Lines and d-Electron Occupancies for the 3d and 4d Transition Metals. Phys. Rev B 1993, 47, 8471. (58) Yedra, L.; Xuriguera, E.; Estrader, M.; López-Ortega, A.; Baró, M. D.; Nogués, J.; Roldan, M.; Varela, M.; Estradé, S.; Peiró F. Oxide Wizard: An EELS Application to Characterize the White Lines of Transition Metal Edges. Microsc. Microanal. 2014, 20, 698−705. (59) Lee, K.-S.; Myung, S.-T.; Sun, Y.-K. Microwave Synthesis of Spherical Li(Ni0.4Co0.2Mn0.4)O2 Powders as a Positive Electrode Material for Lithium Batteries. Chem. Mater. 2007, 19, 2727−2729. (60) Levin, B.; Jiang, Y.; Padgett, E.; Waldon, S.; Quammen, C.; Harris, C.; Ayachit, U.; Hanwell, M.; Ercius P.; Muller, D.; Hovden, R. Tutorial on the Visualization of Volumetric Data Using tomviz.

28


Page 28 of 37

Page 29 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Microscopy Today 2018, 26(1), 12-17. (61) Dragonfly 3.5 (Computer software). Object Research Systems (ORS) Inc, Montreal, Canada, 2016; software available at http://www.theobjects.com/dragonfly.

Figure 1. Cross-sectional SEM images of three LiNi0.6Co0.2Mn0.2O2 (NCM622) powders prepared from the three precursors (P1, P2, and P3). The three samples are hereafter referred to as (a) C1, (b) C2, and (c) C3, respectively. Each figure panel consists of the overall (left), core (middle), and shell (right) structures of the three secondary particles.

29



Figure 2. STEM-EDX mapping of sample C2 oriented in the [010] direction. The elemental maps obtained for (a) the ABF-STEM image are displayed with different false colors: yellow for (b) Ni L, green for (c) Co L, blue for (d) Mn L, and red for (e) O K peaks. (f) The composite image of all elemental maps clearly corresponds to the overlaid atomic model.

30


Page 30 of 37

Page 31 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 3. 3D tomography of the three samples revealed by the FIB-SEM serial section imaging technique. (a–c) Perspective views of the sectional 3D volumes of samples C1, C2, and C3, respectively. (d–f) Cross-sectional views of the three reconstructed volumes at the centers of the respective secondary particles for C1, C2, and C3, showing their primary particle distributions. (g–i) 3D distribution of internal pores in the three secondary particles of C1, C2, and C3, respectively. Note that the arbitrary rectangular volumes (5.6 (X) x 4.2 (Y) x 0.7 (Z) μm3) in the three samples (marked by yellow dotted rectangles in a–c) were chosen for the comparison of quantitative grain size and pore distributions.

31



Figure 4. Primary grain size and pore distributions inside the reconstructed secondary particles. (a) Plot for the change in primary particle size measured as a function of relative radius (from the center (0) to the radius (R)). (b) Plot of the variation in local porosity inside the secondary particles measured as a function of distance from the center.

32


Page 32 of 37

Page 33 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 5. Cell performance tests utilizing C1, C2, and C3 NCM622 cathode materials. (a) Plots of rate capability at incremental discharge rates ranging from 0.1–10.0 C. (b) Cycling performance of discharge capacity at 1.0 C in the range of 3.0–4.3 V.

33



Figure 6. Cyclic voltammograms of samples (a) C1 and (b) C2 obtained at different scan rates ranging from 0.01–1.00 mV∙s-1.

34


Page 34 of 37

Page 35 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 7. Oxidation state maps of TM ions (Ni, Co, and Mn) of samples (a) C1 and (b) C2 before and after charging. The oxidation state maps of the TMs were derived from the intensity ratios (L3/L2) of the respective Ni, Co, and Mn L edges for the regions corresponding to the yellow boxes in the corresponding ABF-STEM images. In both (a) and (b), the upper array of figure panels presents the results of valence state mapping prior to charging (discharged state) and the lower array of figure panels presents the results after charging (charged state).

35



Figure 8. Comparison of atomic structures between the primary particles of samples C1 and C2 before and after charging. (a, b) A pair of HAADF (left) and ABF (right) STEM images of samples C1 and C2 samples prior to charging, respectively. (c, d) A pair of HAADF and ABF STEM images of the same two samples after charging. (e, f) Comparison of the (reversed) intensity profiles of the Li sites (marked with green dotted rectangles in each ABF-STEM image) of the two samples before and after charging. It should be noted that the intensities of the Li sites in sample C2 sample are significantly reduced compared to those in sample C1 after charging.

36


Page 36 of 37

Page 37 of 37 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Table of Contents (TOC) graphic

37


Hierarchically Structured CoreâShell Design of a Lithium Transition

Recommend Documents

Hierarchically Structured CoreâShell Design of a Lithium Transition