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Layered-Layered-Spinel Cathode Materials Prepared by a High-Energy Ball-Milling Process for Lithium-ion Batteries Soo Kim, Jae-Kyo Noh, Muratahan Aykol, Zhi Lu, Haesik Kim, Wonchang Choi, Chunjoong Kim, Kyung Yoon Chung, Chris Wolverton, and Byung Won Cho ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.5b08906 • Publication Date (Web): 08 Dec 2015 Downloaded from http://pubs.acs.org on December 10, 2015

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ACS Applied Materials & Interfaces

Layered-Layered-Spinel Cathode Materials Prepared by a High-Energy Ball-Milling Process for Lithium-ion Batteries Soo Kim,† Jae-Kyo Noh,‡,ǁ Muratahan Aykol,†,ǁ Zhi Lu,† Haesik Kim,‡ Wonchang Choi,‡ Chunjoong Kim,¶ Kyung Yoon Chung,*,‡ Chris Wolverton,*† and Byung-Won Cho‡ †

Department of Materials Science and Engineering, Northwestern University, 2220 Campus Drive, Evanston, Illinois 60208, United States

‡

Center for Energy Convergence Research, Korea Institute of Science and Technology (KIST), Hwarang-ro 14-gil 5, Seongbuk-gu, Seoul 136-791, Republic of Korea ¶

Department of Materials Science and Engineering, Chungnam National University, 99 Daehak-ro, Yuseong-gu, Daejeon 305-764, Republic of Korea

KEYWORDS: Lithium-ion battery, three-component electrode, layered-layered-spinel cathode, high-energy ball-milling process, nanocomposite

ABSTRACT:

In

this

work,

we

report

the

electrochemical

properties

0.5Li2MnO3·0.25LiNi0.5Co0.2Mn0.3O2·0.25LiNi0.5Mn1.5O4 0.333Li2MnO3·0.333LiNi0.5Co0.2Mn0.3O2·0.333LiNi0.5Mn1.5O4

of and

layered-layered-spinel

(L*LS)

cathode materials prepared by a high-energy ball-milling process. Our L*LS cathode materials can deliver a large and stable capacity of ~200 mAh/g at high voltages up to 4.9 V, and do not

ACS Paragon Plus Environment

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show the anomalous capacity increase upon cycling observed in previously reported threecomponent cathode materials synthesized with different routes. Furthermore, we have performed synchrotron-based in situ X-ray diffraction measurements and found that there are no significant structural distortions during charge/discharge runs. Lastly, we carry out (opt-type) van der Waals-corrected density functional theory (DFT) calculations to explain the enhanced cycle characteristics and reduced phase transformations in our ball-milled L*LS cathode materials. Our simple synthesis method brings a new perspective on the use of the high-power L*LS cathodes in practical devices.

1. Introduction More advanced cathode materials for Li-ion batteries that can operate at a high voltage and deliver a large capacity are desired for emerging energy applications.1-12 Recently, stabilizing the hexagonal R3ത m LiM'O2 (M' = Ni, Co, Mn) with the monoclinic C2/m Li2MnO3 to form an integrated layered-layered structure cathode has been the subject of numerous studies.1-12 Furthermore, adding a third spinel component to this integrated nanocomposite cathode has been shown to be a viable option to achieve faster 3-dimensional Li+ diffusion and to operate at a higher voltage (>4.5 V).2-5,12 These pioneering concepts were initially proposed by Thackeray and co-workers.1-4,12 The layered-layered-spinel (L*LS) cathode materials consist of C/2m Li2MnO3 (L*), R3തm LiM'O2 (M' = Ni, Co, Mn) (L), and Fd3തm LiM2''O4 (M'' = Ni, Mn) (S), and they are typically prepared using a co-precipitation method by adding excess Li-salt to transitionmetal-precursors.1-5,12 Due to structural complexities of the resulting composite, it is very difficult to verify the presence and quantity of each component or to examine transition metal distributions in L*LS using X-ray diffraction (XRD), high-resolution transmission electron microscopy (HRTEM), or X-ray absorption spectroscopy (XAS).7-9,11 For the two-component


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layered-layered materials, an alternate synthesis strategy was recently proposed, whereby Li2MnO3 and LiM'O2 are prepared separately and mechanochemically-alloyed with a high-power ball-milling process.8,9,11 This synthesis route yields the same crystallographic features, and comparable electrochemical performance as the layered-layered structure cathodes produced with co-precipitation.7-9,11 However, our synthesis route provides a new strategy where the mole ratios of the various electrode compounds can be precisely controlled to form a closelyconnected nanocomposite.8-11 Here, we present L*LS cathode materials consisting of Li2MnO3 (L*), LiNi0.5Co0.2Mn0.3O2 (L), and LiNi0.5Mn1.5O4 (S), synthesized via a robust and efficient solid-state high-energy ball-milling method for the first time. We prepare the L*LS integrated cathodes with two different compositions: 0.5L*·0.25L·0.25S (denoted as 211 L*LS) and 0.333L*·0.333L·0.333S (denoted as 111 L*LS), along with their heat-treated counterparts 211 L*LS-H and 111 L*LS-H (denoted with -H). The structure and morphology of our prepared samples are characterized by the x-ray diffraction (XRD) methods (both ex situ and in situ techniques), scanning electron microscopy (SEM), and high resolution transmission electron microscopy (HRTEM). Furthermore, the electrochemical properties of all our cathode powders are examined with galvanostatic chargedischarge cycling test and cyclic voltammetry (CV). Lastly, we perform density functional theory (DFT + U) calculations with van der Waals (vdW) corrections to further understand the complex behaviors of the integrated cathode system. 2. Experimental and Computational Methods Li2MnO3 powders were synthesized at 400 °C in a box furnace (32 hr) from LiOH·H2O (Aldrich) and MnCO3·xH2O (Aldrich).6,8-11 High-energy ball-milling process (350 rpm, 3 hr) was carried out in a wet-milling basis with acetone as the solvent with the Li2MnO3 powder


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prepared as above, LiNi0.5Co0.2Mn0.3O2 (LNF Chemicals) powder, and LiNi0.5Mn1.5O4 (Samsung Advanced Institute of Technology) powder to prepare 211 L*LS and 111 L*LS, similar to our previous studies.8-11,13 Additional heat treatment at 750 °C (12 hr) was chosen to prepare 211 L*LS-H and 111 L*LS-H samples, which ensures that the LiNi0.5Mn1.5O4 phase does not decompose into other rocksalt phases or release oxygen gas. The scanning electron microscopy (SEM, FE-SEM, NOVA NanoSEM200, SEI) and high resolution transmission electron microscopy (HRTEM, TITAN, FEI) were used to characterize the morphology of various powered samples. Rigaku X-ray Diffractometer using a monochromatic Cu-Kα line (40 kV, 100 mA) was used for the XRD measurements with a scan rate of 2° min-1. For electrochemical evaluations, active materials were mixed with Ketjen black (KB-500JD) and PVDF (SOLEF5130, Solvay) in N-methyl-2-pyroolidone with a weight ratio of 93:3:4. Slurry was mixed in a homogenizer for 1 hr then casted onto an Al foil using a doctor blade. The average loading level of the electrodes was 3.5 mg cm-2 with an initial thickness of ~90 µm. The electrodes were rolled with a calendar press to a packing density of 3.0 g cm-3. 2032-type coin cell using a Li foil as the counter and reference electrode with a layer of separator (Tonen, Toray) was assembled in a dry room. An electrolyte with 1 M LiPF6 in EC : DMC : DEC = 1 : 1 : 1 by volume was used. The charge/discharge cycling test was performed in a multichannel battery tester (Model 4000, Maccor Inc.) in a potential range of 2.0 to 4.9 V with a test current of 10 mA g-1, unless noted otherwise. Synchrotron based in situ XRD was taken at beamline 1D and 5A at Pohang Accelerator Laboratory (PAL). A modified 2032-type coin-cell was used as an in situ XRD cells. A 3-mm-diameter-hole at the center of the cell with Kapton window serves as a beam path. The 2θ angles of all the XRD spectra have been converted to the 2θ of the conventional X-ray tube corresponding angles for λ = 1.54 Å, which is the wavelength source with a Cu-Kα radiation.


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Synchrotron-based in situ XRD scans were continuously collected at 20 mA g-1. Density functional theory (DFT) calculations within GGA + U approximation14-16 were carried out using the U values of 3.3, 6.4, and 4.5 eV for Co, Ni, and Mn, respectively. We used a cutoff energy of 520 eV for the plane-wave basis set in all calculations with at least 8,000 k-points per reciprocal atom. We further considered van der Waals (vdW) functionals in DFT + U calculations to more accurately describe structural properties of the integrated cathode system. 3. Results and Discussion Figure 1 displays the XRD profiles of Li2MnO3, LiNi0.5Co0.2Mn0.3O2, LiNi0.5Mn1.5O4, 211 L*LS, 111 L*LS, 211 L*LS-H, and 111 L*LS-H [see Figure S1 in supporting information (SI) for magnified XRD profiles of Li2MnO3, LiNi0.5Co0.2Mn0.3O2, and LiNi0.5Mn1.5O4]. Superlattice XRD peaks between 21 and 25° characterize the cation ordering of Li and Mn atoms occurring in the transition-metal layer of the Li2MnO3 component.8-11 While it was difficult to confirm the exact stoichiometric yields in previous L*LS cathode materials1-5 due to the similarities in XRD patterns of Li2MnO3, LiM'O2, and LiM''2O4 cathode materials as shown in Figure 1, synthesis via high-energy ball-milling process has allowed us to prepare 211 L*LS and 111 L*LS cathode powders starting with precisely measured amounts of each cathode component. After heat treatment, we find that the intensity of weak Li2MnO3 superlattice peaks (21 to 25°) become more apparent in 211 L*LS-H and 111 L*LS-H and new double peaks appear at 65° and 70° with increasing temperature, compared to the pristine 211 L*LS and 111 L*LS samples. We note the peak splitting at ~65° is also observed for LiNi0.5Co0.2Mn0.3O2 cathode materials [see Figure S1b in SI]. The large peak broadening observed in the as-milled samples (i.e. 211 L*LS and 111 L*LS) due to decreased crystallite size of cathode components vanishes with the secondary heat treatment, where the peaks become sharp and narrow.7-9,11 This indicates that our heat-treated


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samples re-crystallize upon annealing7 and improve their crystallinity.8,9,11 It has been previously shown that the increase in crystallinity with the secondary heat treatment directly contributes to enhanced electrochemical performance of the Li2MnO3-stabilized R 3ത m cathode materials prepared by a high-power ball-milling process.8,9,11

Figure 1. XRD patterns of single- and three-component electrode materials. Figures 2a–d show the scanning electron microscopy (SEM) images of the raw Li2MnO3, LiNi0.5Co0.2Mn0.3O2, and LiNi0.5Mn1.5O4 powders (Figures 2a–c, respectively) used to prepare our L*LS cathode materials (Figure 2d). The HRTEM image of 111 L*LS-H is provided in Figure 2e with three distinct crystal structures of cathode components verified by the fast Fourier transform (FFT) patterns.8,9,11,17 The Fd3തm spinel phase can be identified from our HRTEM image similar to Ref. 17, while the layered-layered (C2/m and and R3തm) structures are observed in Figure 2e similar to previous studies in Refs. 8, 9, and 11. The HRTEM images of 211 L*LS, 111 L*LS, and 211 L*LS-H cathode materials are also shown Figure S2 (SI). The as-milled samples have very small crystalline sizes, and they show a weak tendency to separate along the


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grain boundaries. The secondary heat treatment aids the intermixing of each cathode component, as observed in Figure 2e and Figures S2c (SI).

Figure 2. SEM images of: (a) C2/m Li2MnO3 (L*), (b) R3തm LiNi0.5Co0.2Mn0.3O2 (L), (c) Fd3തm LiNi0.5Mn1.5O4 (S), and (d) 111 L*LS. The high-energy ball-milling process yields fine powders of L*LS cathodes as shown in Figure 2d. (e) HRTEM image of 111 L*LS-H.


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Figure 3a shows the 1st charge/discharge profiles of our 211 L*LS and 111 L*LS samples, where the initial discharge capacities are 200.0 and 177.3 mAh g-1, respectively. With the secondary heat treatment, the 1st discharge capacities for 211 L*LS-H and 111 L*LS-H are improved to 257.3 and 226.1 mAh g-1, respectively, as shown in Figure 3b. If the Li2MnO3 component evolves oxygen during the 1st charging process, it can lead to very low columbic efficiency.6,10 We find that the initial columbic efficiency has significantly improved for the heattreated samples (211 L*LS-H and 111 L*LS-H) in Figure 3b, when compared with a single Li2MnO3 cathode or 211 L*LS and 111 L*LS samples in Figure 3a. Three distinct L*LS voltage regions can be clearly identified during the 1st charging process in Figures 3a–b. By comparing with known voltages of the various layered and spinel components, we can identify the phases responsible for each region of the charge/discharge profile. At the onset of charge, Li+ ions are mostly extracted from LiNi0.5Co0.2Mn0.3O2. The subsequent plateau near 4.5 V (vs. Li/Li+) corresponds to the activation of Li2MnO3 during the 1st charging run.1-12 Since 211 L*LS-H has more Li2MnO3, a wider activation plateau at 4.5 V is observed. The highvoltage LiNi0.5Mn1.5O4 spinel phase in L*LS shows a redox reaction of Ni2+/Ni4+ at a plateau near 4.7 V. It is interesting to note this spinel component contributes ~80 mAh g-1 of capacity in 111 L*LS-H and only ~10 mAh g-1 in 211 L*LS-H based on the plateau near 4.7 V upon the 1st charging process in Figure 3b. During the 1st discharge run in Figures 3a–b, the LiNi0.5Mn1.5O4 spinel phase also contributes to voltage plateaus around 2.7 V. We show the 2nd charge/discharge profiles for 211 L*LS-H and 111 L*LS-H in Figures 3c–d, where the irreversible capacities of the heat-treated samples have decreased after the 1st cycle. During the 2nd charging run in Figures 3c–d, both the high-voltage spinel redox plateau at 4.7 V and the Li2MnO3 activation plateau at 4.5 V are no longer observed in either heat-treated


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composite. After the 2nd charge/discharge runs, there are no dramatic changes observed in the voltage vs. capacity curves [see Figure S3 in SI], which we attribute to further stabilization of the layered components by the spinel components.1-4 The average discharge voltages of the L*LS-H cathode system after the 2nd charge/discharge runs are observed at ~3.5 V. Li2MnO3 may transform into a spinel-like phase during charging process.1,5,6 Figures 3e–f show the cycling performance of the prepared samples depends strongly on the amount of Li2MnO3 and the post heat-treatment process. The as-milled samples show significant capacity loss during the first few cycles. The heat-treated cathode powders (211 L*LS-H & 111 L*LS-H) display enhanced cycle performances compared with 211 L*LS and 111 L*LS. This improvement agrees well with previous work,8,9,11 where it was shown that the secondary heat treatment at high temperature helps to improve the crystallinity and mixing in the layeredlayered integrated cathodes, and subsequently enhances the electrochemical performance of samples prepared with high-energy ball-milling. Among the heat-treated samples, while 211 L*LS-H achieves a higher initial capacity, 111 L*LS-H is more promising for long-term use with better cyclability [see Figure S4 in SI for rate capabilities of 211 L*LS-H and 111 L*LS-H]. One of the most distinct electrochemical behaviors that differentiate our L*LS cathode prepared in this work from similar L*LS cathodes reported in previous studies1-5,12 is that during cell cycling we observe that the capacity decreases gradually as expected, as shown in Figures 3e–f. Although this “normal” cycling characteristic (i.e. gradual capacity decrease with cycling) is commonly observed for almost all other cathode electrode materials such as LiCoO2, LiNi0.5Co0.2Mn0.3O2, LiMn2O4, LiNi0.5Mn1.5O4, and LiFePO4, previous reports of L*LS cathodes display the unique phenomenon of a gradual increase in capacity upon cycling.2,3,5,12 To explain the capacity increase with cycling in their L*LS cathodes, Lee et al.5 proposed that Li+ ions


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cannot re-insert back to the layered-layered cathode and they enter the empty 16c octahedral sites of the spinel phase. Our results indicate that the synthesis conditions significantly affect the final electrochemical performance in these L*LS type cathodes, as this capacity increase is not observed in our samples. While the capacity increase with cycling could offer a higher energy and power density,2,3,5,12 it would require very careful engineering design for practical applications. For instance, this capacity increase would make it difficult to optimize the ratio of the amount of negative and positive electrodes to be incorporated in a full cell configuration. The gradual increase in capacity would result in excess lithium ions generated from the positive electrodes, which could plate on the nearest section of the current collector and lead to safety issues.18


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Figure 3. The 1st cycle charge/discharge profiles of: (a) L*LS and (b) L*LS-H; and the 2nd cycle charge/discharge profiles of (c) 211 L*LS-H and (d) 111 L*LS-H. Cycle characteristics of: (e) 211 L*LS and 211 L*LS-H and (f) 111 L*LS and 111 L*LS-H.


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The cyclic voltammetry (CV) results for 211 L*LS-H and 111 L*LS-H are provided in Figures 4a–b. The hysteresis of the Li2MnO3 component is observed from the CV analysis at 4.4 V (charging) and 3.2 V (discharging) in Figures 4a–b, similar to Refs. 19 and 20. We also find that that the redox peak of the high-voltage spinel phase has slightly shifted from 4.7 V to 4.6 V. Transition metal mixing between Li2MnO3 and LiNi0.5Mn1.5O4 during the secondary heat treatment could slightly change the redox voltage. To clarify such mixed potential effects, we also report the average voltage for the redox reaction of LiM''2O4-M''2O4 using density functional theory (DFT + U) calculations in Figure 4c. We obtain the lowest energy LiNi0.5Mn1.5O4 spinel structure from the cluster expansion method (Li6Ni3Mn9O24 supercell).21 Starting with this structure, we tested further Mn additions by considering all possible transition metal mixing arrangements in LiNi0.5Mn1.5O4 by swapping Ni atoms by Mn atoms, one at a time. Since both L*LS cathode materials are Mn-rich (i.e. ≥70% Mn in transition metals), we choose to mix the Mn ions into Ni sites starting from LiNi0.5Mn1.5O4 composition. The calculated DFT voltages indicate that a considerable drop in voltage of LiNi0.5Mn1.5O4 can be reasonably expected in the CV analysis for this high-voltage spinel phase as it further enriched in Mn during heat treatment.


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Figure 4. Cyclic voltammetry (CV) analysis of: (a) 211 L*LS-H and (b) 111 L*LS-H. Mixed potential effects are observed in Figures 4a – b, where the redox peak of the high-voltage spinel phase is slightly shifted from 4.7 V to 4.6 V. In Figures 4a – b, L*, L, and S correspond to the redox reaction of the Li2MnO3, LiNi0.5Co0.2Mn0.3O2, and LiNi0.5Mn1.5O4, respectively. (c) Mixed potential effects in LiM''2O4 (M'' = Ni and Mn) electrodes. DFT + U calculations were carried out to examine the effect of transition metal mixings (occurred during the secondary heat treatment) on the redox voltage.


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We examined our 111 L*LS-H powder using synchrotron-based in situ XRD, performed at beamline 1D and 5A at the Pohang Accelerator Laboratory (PAL) and show the change in the diffraction pattern and the lattice parameters of spinel and layered components as a function of voltage in Figure 5a and Figure 5b, respectively. At open-circuit voltage, L*(002), L(003), and S(111) peaks overlay at 18.7° 2θ because of their structural similarities, as shown in Figure 5a. At 4.5 V, the L(003) peak starts shifting to lower angles, which indicates an expansion of c-axis of the layered structure. The S(111) peak also shifts to higher angles at ~4.5 V, before its typical redox plateau at 4.7 V. The mixed potential effects discussed earlier in Figure 4c (i.e. Mn addition to the high-voltage spinel during the secondary heat treatment) are supported by the S(111) peak shifts at 4.5 V) where the layered C2/m and R3തm normally degrade, it effectively stabilizes these layered components (analogous to Li2MnO3 stabilizing LiM'O2 in Ref. 1) and in turn helps


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inhibiting the decay in electrochemical performance. In Figure 5a, the spinel peaks, S(511), S(440), and S(531), shift to the higher angle positions during charging process at 58.8, 64.5, and 67.9°, respectively.24,25 Figures S5a–b in SI show magnified images of Figure 5a; and the 2nd charging/discharging in situ XRD runs are provided in Figures S5c–d in SI. In Figure S5c (SI), it is possible to reconfirm that the shift of the S(111) peak occurs at a voltage slightly lower than 4.7 V. In addition, we show that the single Li2MnO3 cathode materials do not show any major peak shifts during the 1st charging and discharging runs in Figure 6, consistent with a previous study.4 Thus, we conclude that no significant peak shifts in Figure 5a are originated from the Li2MnO3 component.


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Figure 5. (a) In situ XRD patterns during 1st charge/discharge for 111 L*LS-H. (b) Lattice parameter changes in layered (L) and spinel (S) phases. Al** peaks originate from the aluminum current collector.


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Figure 6. In situ XRD measurement of Li2MnO3 cathode materials during 1st charge/discharge and 2nd charge/discharge cycles. No major peak shifts are being observed from the single phase Li2MnO3 component from our synchrotron-based in situ XRD runs, similar to Ref. 4. We carried out first-principles calculations to investigate the extent of structural stabilization of the layered R3തm structures by the Li2MnO3 component. First we examine the lattice parameter and volume changes of the R3തm LiM'O2 structures with a complete delithiation process. For our DFT calculations of structural parameters, we used the opt-type van der Waals density functional (vdW-DF).26-28 We have recently shown that standard DFT with (semi)-local exchange correlation functionals are not sufficient to describe the structural changes in layered R3ത m structures upon delithiation, and the “opt” type vdW-DF gives a more accurate lattice parameter, crystal volume, and interlayer spacing by incorporating dispersion interactions.28 The current optB88 (+U) results in Figure 7a show that the lattice parameter c of a single R3തm LiM'O2 component significantly decreases upon delithiation; i.e. these layered structures significantly distort (“collapse”) after the complete delithiation process. The corresponding volume changes


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for LiCoO2, LiNiO2, and LiMnO2 are found to be -5.84, -12.54, and -13.21 %, respectively in Figure 7b. Thus, if these large volume changes are not suppressed, lithium ions would have larger kinetic barriers to re-insert into these layered oxides.5,9 Next, we further test composite Li2MnO3-LiCoO2, Li2MnO3-LiNiO2, and Li2MnO3-LiMnO2 DFT supercells, where we removed all lithium atoms from the R3തm component, as shown in Figures S6a–b (SI). Here, we use a simplified DFT model not accounting for the release of oxygen or the gradual Mn migration in Li2MnO3 component. In contrast to the pure layered phase calculations, these supercell calculations show the Li2MnO3 stabilization effects1 (Figure 7a), where the lattice parameter c increases very slightly and there are negligible changes in the lattice parameter a. These lattice parameter changes are also consistent with the XRD data in Figure 5. Our DFT calculations indicate that the volume changes are significantly reduced for Li2MnO3-LiCoO2 (2.17 %), Li2MnO3-LiNiO2 (0.33 %), and Li2MnO3- LiMnO2 (2.53 %), respectively. Therefore, these admittedly simplified DFT models with opt-functionals support the hypothesis of effective structural Li2MnO3 stabilization in our L*LS integrated cathode materials synthesized with the high-energy ball-milling process.


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Figure 7. (a) Calculated lattice parameters a and c of the R 3ത m structures before/after the delithiation. Here, L*-LCO, L*-LNO, and L*-LMO correspond to Li2MnO3-stabilized LiCoO2, Li2MnO3-stabilized LiNiO2, and Li2MnO3-stabilized LiMnO2 electrode materials, respectively. (b) Calculated volume changes (%) during the R3തm delithiation. The y-axis in Figure 7 (b) is provided as an absolute value to compare the volume changes with the R3തm delithiation.


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The anomalous capacity increase with cycling is not observed for the pure high-voltage spinel [see Figure S7 in SI] or the layered-layered cathode materials. Previous work has suggested a connection between this capacity increase and an overlithiation of the spinel component in L*LS materials.5 We show the conventional cell of the ordered LiNi0.5Mn1.5O4 structure adapted from Ref. 29 in Figure S8 (SI); Ni and Mn occupy on the 4b and 12d sites of the spinel structure. Figure S8b (SI) shows the chemical structure of Li1+xNi0.5Mn1.5O4, where a Li atom occupies an empty 16c site. The overlithiation of the spinel LiNi0.5Mn1.5O4 phase typically occurs at ~2.7 V, as also experimentally observed in Figures 3–4. We believe that the anomalous capacity increase in previous studies2,3,5,12 is too large to be solely explained by the overlithiation of the spinel at ~2.7 V, as suggested in Ref. 5. The anomalous capacity increase with cycling may involve more complex chemistry (i.e. possibly, a phase transformation), and some L*LS samples could be more prone to these phase transformations. For example, recent efforts by Long et al. successfully controlled the spinel content in the L*LS to obtain a steady capacity of 190 mAh g-1 with little voltage fade.30 Our method that accurately controls the three-component L*LS system also does not display any anomalous capacity increases. Since L*LS composites are quite complex systems, there is likely more than one factor associated with this behavior.2,3,5,12,30 The major challenges associated with the Li-rich cathode materials include cycling stability, phase transformation upon cycling, low 1st cycle columbic efficiency, and poor rate capability. In this work, we have attempted to integrate LiNi0.5Mn1.5O4 in layered-layered cathode system using high-energy ball-milling for the first time. We are able to enhance the cycling behavior, to reduce phase transformations during cycling, and to improve columbic efficiency. More future work is essential in order to optimize the L*LS compositions and synthesis conditions as well as to improve the rate capabilities for further commercial use.


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4. Conclusions Three-component layered-layered-spinel (L*LS) electrode materials were prepared using a solid-state high-energy ball milling process with precisely controlled mole ratios for the first time. Our facile and robust synthesis route represents an important step towards the practical utilization of the high-power L*LS integrated electrode materials for emerging high-energy applications. We have investigated the structural and electrochemical properties of our L*LS cathode and found that the electrochemical performance of L*LS cathode significantly improves with the secondary heat treatment after the high-energy ball-milling process. Synchrotron-based in situ XRD technique is utilized to characterize the 111 L*LS-H cathode powder, where we find that there are no significant structural distortions for either layered or spinel phases during charge/discharge processes for our L*LS three-component cathode. Lastly, we have carried out the opt-type van der Waals corrected DFT calculations to analyze the extent of structural stabilization of the layered R3തm components with the addition of the Li2MnO3 component. Our calculated results are consistent with our in situ XRD results. We finally note that each component has to be prepared separately and integrated together in our new approach in order to prepare three-component cathodes, which might add another process cost in the industrial-scale. Future work is certainly necessary to find the optimum combinations of these systems to develop more reversible and durable high-power composite systems, and at the same time, to reduce the overall manufacturing cost.


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ASSOCIATED CONTENT Supporting Information XRD profiles of single cathode component (Li2MnO3, LiNi0.5Co0.2Mn0.3O2, and LiNi0.5Mn1.5O4), HRTEM images of 211 L*LS, 111 L*LS, and 211 L*LS-H, discharge capacity vs. voltage plots of 211 L*LS-H and 111 L*LS-H, rate capabilities of 211 L*LSH and 111 L*LS-H, in situ XRD analysis of 111 L*LS-H, and chemical structures of Li2MnO3-LiM’O2 (M’ = Ni, Co, Mn) and LiNi0.5Mn1.5O4 cathode materials, cycle characteristic of LiNi0.5Mn1.5O4 cathode AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] *E-mail: [email protected] Author Contributions S.K. and J.-K.N. synthesized and characterized the electrode materials. S.K., M.A., and Z.L. performed the computational analysis. J.-K.N. and H.K. carried out the in situ XRD analysis. W.C. and C.K. helped interpret the experimental results. K.Y.C, C.W., and B.-W.C. supervised this project. All authors were involved in completion of this paper and have given approval to the final version of the manuscript. ǁ These authors contributed equally.

ACKNOWLEDGMENT S.K. was supported by Northwestern-Argonne Institute of Science and Engineering (NAISE) (171-8214100-10031635). M.A., Z.L., and C.W. were supported by The Dow Chemical


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Company (Dow Agmt #223029AS). J.-K.N., H.K., W.C., K.Y.C., and B.-W.C. were supported by

the

National

Research

Foundation

of

Korea

Grant

funded

by

the

Korean

Government(MSIP)"(NRF-2011-C1AAA001-0030538) and by KIST Institutional Program (Project No. 2E25303). C.K. was supported by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy (DOE) under Contract No. DE-AC02-05CH11231, as part of the Batteries for Advanced Transportation Technologies (BATT) Program. We acknowledge the use of Beamline 1D at Pohang Accelerator Laboratory (PAL). This research used resources of the National Energy Research Scientific Computing Center, a DOE Office of Science User Facility supported by the Office of Science of the U.S. Department of Energy under Contract No. DE-AC02-05CH11231. S.K. is grateful for enlightening discussion with Dr. Shiqiang Hao at Northwestern University. The authors are grateful to Dr. Eungje Lee, Dr. Jason R. Croy, and Dr. Michael M. Thackeray at Argonne National Laboratory for fruitful discussions and suggestions. REFERENCES (1) Thackeray, M. M.; Kang, S. -H.; Johnson, C. S.; Vaughey, J. T.; Benedek, R.; Hackney, S. A. Li2MnO3-stabilized LiMO2 (M = Mn, Ni, Co) Electrodes for Lithium-ion Batteries. J. Mater. Chem. 2007, 17, 3112–3125. (2) Park, S. -H.; Kang, S. -H.; Johnson, C. S.; Amine, K.; Thackeray, M. M. Lithiummanganese-nickel-oxide Electrodes with Integrated Layered-spinel Structures for Lithium Batteries. Electrochem. Commun. 2007, 9, 262–268.


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Layered-Layered-Spinel Cathode Materials ... - ACS Publications

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