Electrolyte Interfaces Employing

Dec 17, 2014 - Abstract Image. Chemical degradation at electrode/electrolyte interfaces in high-energy storage devices, such as Li-ion batteries, impo...
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Stabilization of Battery Electrode/Electrolyte Interfaces Employing Nanocrystals with Passivating Epitaxial Shells Chunjoong Kim,†,¶ Patrick J. Phillips,‡ Linping Xu,¶,# Angang Dong,§,∇ Raffaella Buonsanti,§,⊥ Robert F. Klie,‡ and Jordi Cabana*,†,¶ †

Department of Chemistry, ‡Department of Physics, University of Illinois at Chicago, Chicago, Illinois 60607, United States Environmental and Energy Technology Division, §The Molecular Foundry, Materials Sciences Division, and ⊥Joint Center for Artificial Photosynthesis, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States # Cristal Global, Glen Burnie, Maryland 21060, United States ∇ Department of Chemistry, Fudan University, Shanghai 200433, China ¶

S Supporting Information *

ABSTRACT: Chemical degradation at electrode/electrolyte interfaces in high-energy storage devices, such as Li-ion batteries, imposes durability challenges that affect their life and cost. In oxide electrodes, degradation is linked to the presence of redox active transition metals at the surface. Here, we demonstrate a strategy toward the stabilization of interfaces using core−epitaxial shell nanocrystals. The core of the nanocrystal is composed of an electroactive oxide, which is passivated by an ultrathin epitaxial oxide shell enriched in a redox inactive ion. This approach imparts interfacial stability while preserving the high storage capability and fast carrier transport of the material, compared to unmodified versions. The validity of the concept is proved with Li1+xMn2−xO4 nanocrystals with a 1−2 nm Al-rich shell, which showed reduced sensitivity to harsh environments, compared to bare counterparts. The approach is generalizable to any transition-metal-based battery system where electrode−electrolyte interactions must be controlled.

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INTRODUCTION Controlling the reactions that occur at electrode/electrolyte interfaces is key to long-lasting energy storage technologies, such as Li-ion batteries, especially when pushing voltages of operation to increase their energy density.1 Extreme potentials reached at the electrodes in a fully charged battery render them too reactive toward the electrolyte, triggering deleterious decomposition reactions. In positive electrodes, it has been established that redox-active transition-metal ions at the surface have a role in the oxidation of the electrolyte.2,3 A second mechanism of degradation in oxide electrodes is through dissolution of metals such as Mn, Ni, and Co in the mildly acidic environments commonly created by electrolyte impurities.4,5 These dissolved species migrate across the cell and poison the negative electrode.6,7 Undesired interfacial processes aggravate when the battery is used at temperatures above standard conditions, which occurs in real-life applications. The use of nanoparticles, typically championed to achieve highpower-density devices,8−13 also aggravates the problem, because of their high surface-to-bulk ratio. Durability issues require overdesigning batteries for the sake of calendar life, imposing a quantifiable cost penalty.14 Reducing the transition metal content by introducing inactive ions improves interfacial stability,6 but must be implemented in © 2014 American Chemical Society

the form of thin passivating barriers to avoid unacceptable losses in storage capacity. However, post-synthetic coating of electrode-active material powders with phases such as Al2O315,16 or MgO,3 which is a popular approach, makes it difficult to access buried interfaces, because of particle agglomeration.17 Interestingly, evidence exists that coating functionality stems from the incorporation of the inactive ion into the surface structure of the active oxide material during processing.3 Here, we introduce a new strategy toward effectively stabilized electrode/electrolyte interfaces using nanocrystals (NCs) of an active material, where each individual NC presents a conformal ultrathin oxide shell grown epitaxially, which contains a high concentration of an inactive ion that is essentially absent in the core. Our goal was to reduce the amount of transition metal at every single interface with the electrolyte, while preserving high charge storage and fast transport using a nanoscale architecture. For this purpose, this approach is better suited than core−shell structures designed at the secondary, micrometric particle level, where nanoscale primary particles are homogeneous but have a different Received: September 30, 2014 Revised: November 19, 2014 Published: December 17, 2014 394

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cm2, followed by drying under vacuum at 110 °C overnight. Dried electrodes were punched with a diameter of 1/2 in. and assembled into 2032 coin-type cells in an Ar-filled glovebox (water and oxygen, ≤0.1 ppm) with a high-purity lithium foil (Product No. 10769, Alfa Aesar) as a counter/reference electrode, 25-μm-thick polypropylene membrane (Celgard 2400) as a separator, and the electrolyte composed of 1 M LiPF6 in a mixture of ethylene carbonate (EC) and diethyl carbonate (DEC) (1:1 (v/v)) (Novolyte Technologies). Charge− discharge measurements were performed on a Biologic VMP3 potentiostat at room temperature (RT) or 50 °C under different rates with a discharge cutoff voltage of 3.5 V and a charge cutoff voltage of 4.3 V vs Li+/Li0. The rate, C/n, was defined as the current density required to achieve a theoretical capacity of C = 148 mAh/g, in n hours. The rate capability of the samples was measured by fixing the charge current to C/10 while discharge currents were sequentially varied from C/10 to 20 C. The cycling tests were replicated three times. The values provided constitute an average of the measurements. The results were found to be reproducible and significant.

composition in the core, compared to the shell.18,19 In these secondary structures, buried interfaces in the core can be accessed by the electrolyte either through pores needed to ensure fast ion transport in the architecture or through particle cracking resulting from mechanical strain during electrochemical cycling,17 especially under harsh conditions.20 We exemplify the approach with NCs of spinel-type Li1+xMn2−xO4 (0 ≤ x ≤ 0.1), a primary component of batteries in electric drive vehicles in the market today,14 engineered with a 1−2 nm shell that was enriched in Al at the surface, as visualized by scanning transmission electron microscopy (STEM). The coreepitaxial shell (C-ES) NC architecture showed increased durability upon cycling and reduced sensitivity to acidic environments, compared to the unmodified counterpart, even at high temperature.

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EXPERIMENTAL SECTION

RESULTS AND DISCUSSION A scheme depicting the procedure to synthesize the NCs, with and without epitaxial shells, is represented in Figure 1. First,

Synthesis of Nanocrystals. The nanocrystals (NCs) were prepared following a colloidal synthesis method. First, 5 mM manganese(III) acetate dihydrate (Product No. 215880, Sigma− Aldrich) was dissolved in a 20-mL oleylamine (OAm) solution (Product No. O7805, Sigma−Aldrich) at room temperature (RT). The solution was transferred to the oil bath at 140 °C and then kept under magnetic stirring for 30 min. A mild oxidizing agent, 15 mM trimethylamine N-oxide (No. 317594, Sigma−Aldrich), was added to the solution and then kept under magnetic stirring for another 60 min. The resulting Mn3O4 nanoparticles were collected by centrifugation for 10 min at 10 000 rpm, redispersed in hexane and washed with ethanol several times to remove the excess surfactant, reaction byproducts, and unreacted precursors. The clean Mn3O4 NCs were redispersed in hexane and then added to 20 mL of OAm solution at 220 °C. Al2O3 shell layers were formed on the surface of Mn3O4 particle by a 60-min dropwise addition of a 0.5 mM aluminum acetylacetonate (Product No. 208248, Sigma−Aldrich) solution in 10 mL of OAm at 220 °C. Mn3O4/Al2O3 core−shell NCs were separated by centrifugation and washed thoroughly with ethanol. After the cleaning process, the NCs were dried overnight at 60 °C and heated at 300 °C to remove any possible organic residuals.21 Mn3O4 and Mn3O4/Al2O3 core−shell NC precursors were then thoroughly mixed with LiOH·H2O (Product No. 43171, Alfa Aesar) and annealed at 500 °C for 4 h in air to produce Li1+xMn2−xO4 NCs without an epitaxial shell (bare NCs) and with an epitaxial shell (C-ES NCs). The Al/Mn ratio in C-ES NCs was found to be 0.05 by chemical analysis. All chemicals were used without any purification process. STEM and EDX Analysis. STEM images were acquired using an aberration-corrected STEM (JEOL, Model JEM-ARM200CF) operated at 200 kV, which can achieve a spatial resolution of ∼73 pm. Energy-dispersive X-ray (EDX) data were collected using a high-solidangle Oxford X-MaxN 100TLE silicon drift detector at 80 kV. EDX line scan data were smoothed from 30 data points using the Savitzky− Golay method and a reflection boundary. Smoothed line scans are plotted in Figures 3b and 3d (presented later in the paper), using the line and symbols, which present actual data points but symbols are shown with every 30 or 40 points to guide the eye. Only compositions of aluminum and manganese were shown, for the sake of simplicity. Transmission electron microscopy (TEM) images from different regions of NCs were collected and used to analyze the particle size distribution. Size analysis from ∼300 particles revealed average diameters of 13.5 ± 2.7 and 18.2 ± 3.1 nm in precursors and final Li-containing spinel NCs, respectively (see Figures 2c and 2d, presented later in this work). These values were consistent with the estimates derived from fittings of peak broadening in (111) peaks. Electrochemical Measurement. The working electrodes were prepared by mixing either bare or core/epitaxial-shell NCs, carbon black (Denka), and polyvinylidene difluoride (PVDF) (Kynar) in Nmethylpyrrolidone (NMP) (Product No. 270458, Sigma−Aldrich) (80:10:10 (wt %)), which were then cast on an electrochemical-grade aluminum foil using a doctor blade to have a loading level of ∼3 mg/

Figure 1. Scheme for the preparation of core−shell nanometric precursors and resultant core−epitaxial shell nanocrystals with a spinel structure.

colloidal Mn3O4 NCs were prepared following a solution procedure in oleylamine, which is a high-boiling-point solvent that also acts as a surface-stabilizing ligand,22−25 using Mn3+ acetate as the precursor in the presence of a mildly oxidizing agent (trimethylamine N-oxide). The formation of pure spineltype Mn3O4 was confirmed by powder X-ray diffraction (XRD; see Figure S1a in the Supporting Information). In order to build Mn3O4/Al2O3 core−shell NCs, the pristine Mn3O4 NCs were redispersed in oleylamine prior to hot injection of a solution of Al3+ acetylacetonate in oleylamine at 220 °C. The resulting Mn3O4/Al2O3 NCs showed a mean particle size of 13.5 nm with a narrow distribution (see Figures 2a and 2c). Neither new crystalline phases nor shifts of the spinel peaks related to the presence of Al could be detected in the XRD pattern (see Figure S1a in the Supporting Information). In a subsequent step, the oxide precursors were mixed with LiOH and the powder mixture was heated in air at 500 °C to prepare the final lithium-containing NCs. The reagent ratios and annealing temperature were tailored to produce the phase of 395

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Figure 3. (a) EDX mapping and (b) corresponding line scans of Al2O3/Mn3O4 core−shell precursors (Mn (red) and Al (green)). (c) EDX mapping and (d) corresponding line scans of C-ES NCs (Mn (red) and Al (green)). (e) LAADF image of surface-modified C-ES NCs with a [110] zone axis; (f) enlarged view of the area of interest noted in panel (e). (Borders between bulk and surface layers are marked by broken lines. High-resolution HAADF images indicate 1−2 nm thick surface layers formed on the surface with local inhomogeneities.) (g) High-resolution LAADF image from a similar edge region of another particle. Mn atomic positions in the spinel structure are indicated by solid red spheres, while O and Li atomic positions in the spinel structure are represented by blue and yellow hollow spheres, respectively.

Figure 2. (a) Representative low-magnification TEM image and (c) corresponding particle size distribution of Mn3O4/Al2O3 core−shell precursors. (b) Representative low-magnification TEM image and (d) corresponding particle size distribution of C-ES NCs. Growth of the particle size from Mn3O4/Al2O3 precursors (panel (c)) to lithiumcontaining C-ES NCs (panel (d)) was observed.

interest while minimizing crystal coarsening (see Figure S2 in the Supporting Information).26 Indeed, only a slight growth to 18.2 nm was induced on the particles, with little aggregation (see Figures 2b and 2d). The resulting XRD patterns matched the spinel structure of Li1+xMn2−xO4, with no detectable impurities (Figure S1b in the Supporting Information). The composition of the crystals without shells (bare NCs) was estimated to be Li1.05Mn1.95O4, based on the cell parameters.27 The peak positions were slightly shifted toward lower angles for the C-ES NCs, indicative of a slightly larger average unit cell. This result could reflect structural distortions induced by the formation of the shell, which are further discussed below. X-ray spectroscopy data collected for C-ES NCs at the O K-edge and Mn L-edge (see Figure S3 in the Supporting Information), using a surface-sensitive electron yield detector, were found to be consistent with oxides containing Mn in oxidation states higher than 2+,28 and were close to Li1+xMn2−xO4 spectra available in the literature.29 Taken together, these results indicate that lithium quantitatively reacted with Mn3O4 (or Mn2+[Mn3+]2O4). STEM images revealed that the precursors possess a crystalline core of Mn3O4 with an amorphous shell of Al2O3 (Figure S4a in the Supporting Information), whereas the surfaces of the C-ES NCs clearly consisted of thin crystalline layers that could be assigned to the Al-containing shell, as discussed below (see broken lines in Figures S4b−d in the Supporting Information). High-resolution imaging of C-ES NCs revealed that the thickness of the surface layer was ∼1−2 nm. The distributions of Al and Mn were directly confirmed on both the core−shell Mn3O4/Al2O3 precursors and the C-ES NCs by EDX spectroscopy conducted at 80 kV with a probe size of ∼120 pm. The representative EDX map of a Mn3O4/ Al2O3 crystal in Figure 3a shows the higher Al content that is homogeneously distributed on the surface. A more-detailed look was afforded by line scans at the Mn and Al K energies (see Figure 3b). The Al/Mn ratio, depicted in blue in the line scans shown in Figure 3, showed a striking border between

Al2O3 and Mn3O4, thus proving that the Al2O3 layer exists as a shell over the Mn3O4 core. On the other hand, the EDX maps of the C-ES NC (see example in Figure 3c), while confirming the core−shell character of the material, showed a somewhat increased inhomogenity in the distribution of aluminum, which can be attributed to the reconstruction and crystallization of the surface during the transformation induced at 500 °C. Mn and Al line scans revealed an absence of a striking border in Al/Mn and a higher concentration of Mn at the surface, compared to the core−shell precursor (see Figure 3d), resulting from cation migration during the thermal treatment. Figures 3e−g present various high-resolution images of a typical C-ES NC. While the structure of the bulk corresponds to a spinel in both bare and C-ES NCs (see Figure S1 in the Supporting Information), locally, the formation of defect clusters was identified in the LAADF STEM images (see Figures 3f and 3g, as well as Figure S5 in the Supporting Information). These defects were characterized by the anomalous intensity (arrowed in Figure 3f) observed in the center of the “diamond” of the spinel structure at the [110] zone axis, corresponding to the 16c site. Given that Li is too light to be observed in LAADF mode, this intensity must originate from the presence of a transition-metal atom in that site. Such defects have been observed by others at the subsurface of LiMn2O4 upon lithium extraction,30 and bear resemblance to a disordered rock-salt structure. Their exact identity and origin are not clear at this point and will be the subject of future work. Through high/low angle annular darkfield (H/LAADF) and annular bright-field (ABF) STEM imaging, the core and surface crystal registry of the particles was confirmed. This epitaxy should favor the transport of carriers into and out of the crystal by minimizing interfacial resistance. The dashed lines in Figures 3f and 3g approximate the border between the core and the Al-containing surface 396

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presence of the Al-rich epitaxial shell. In discharge rate capability tests (Figure 4c), the ratio of capacity at 20C vs C/ 10 was 80% for bare NCs and 83% for C-ES NCs, which are comparable to the highest values available in the literature.8−10,18,34 This observation indicates that charge transfer kinetics is not handicapped by the Al-rich shell, consistent with the small polarization observed between charge and discharge (see Figure S6b in the Supporting Information).35 Differences in material properties were more apparent at elevated temperature, where side reactions are accelerated. At 50 °C, the bare spinel NCs retained only 68% of their initial discharge capacity at C/10, compared to 83% cycle retention by C-ES NCs (Figure 4d). The Coulombic efficiency decreased in both cases, but appeared to be lower upon cycling for bare NCs, compared to C-ES NCs. Further proof of the improved stability was provided by accelerated tests of Mn leaching from the oxide NCs into dilute acidic solutions, according to an existing literature protocol.36 Solutions exposed to C-ES NCs showed seven times less Mn, compared to bare NCs (0.3 wt % vs 2.1 wt %). These results suggest that the Al-rich shells are effective at passivating the NC surface toward both electrolyte decomposition and corrosion, yet retain excellent electrochemical response by virtue of their nanometric features.

layer. Increased intensity was explicitly observed in certain atomic positions corresponding to the tetrahedral sites in the spinel structure, which are identified with yellow hollow circles in Figures 3f and 3g. This observation is again indicative that strongly scattering ions such as Al or Mn are occupying these sites, which should be occupied by Li only in a defect free Li1+xMn2−xO4 structure. Instead, these defects could be produced by the formation of Li-deficient domains related to Li1−xMn2+xO4,27 which, at low x, preserve the cubic spinel structure, but lead to the presence of Mn ions in the tetrahedral sites.31 These defects could be the result of insufficient Li content provided in the form of LiOH during the synthesis or the formation of related phases, because of the presence of significant fractions of Al at the surface. The fact that Mn was found to be in a high oxidation state by XAS again supports the notion that Li was present in this layer in order to compensate charge in the spinel structure. The performance of C-ES NCs as electrode materials was evaluated in Li metal half cells and compared to bare NCs. The voltage versus specific capacity profiles collected during galvanostatic cycling at RT did not show any obvious differences (see Figure 4a), except from a slight decrease in

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CONCLUSIONS

In conclusion, we have successfully demonstrated interfacial stabilization of Li-ion battery electrodes through the concept of core−epitaxial shell oxide NCs. The ultrathin shell has a similar structure and composition as the core, but is rich in a redox inactive ion. This approach was applied to multinary chemical spaces such as spinel-type Li1+xAlyMn2−x−yO4. The concentration of dissolution-prone Mn3+ ions was reduced at the surface at the expense of isovalent Al3+, to avoid leaching by acidic impurities and their participation in the decomposition of the electrolyte.2,3 Through colloidal chemistry processes, the shell was grown conformally, thereby effectively passivating all surfaces. The result was an architecture with increased durability under harsh conditions, with respect to materials where the shell was not present. While core−epitaxial shell nanocrystal architectures have been demonstrated,37,38 this level of chemical complexity has not been achieved with battery materials, which creates opportunities to design electrodes with practical impact.39 Indeed, the nanocrystals presented here could be used as building blocks for secondary architectures that optimize material density without compromising properties.40 Given the similarity in chemistry between multiple oxide electrode systems based on transition metal oxides,38 this approach could be extended to other cases where electrode/electrolyte interactions must be controlled.1,41 The challenges of controlling electrolyte/electrode interfaces will exist, regardless of the technology, from Li-ion to emerging multivalent concepts.42 Therefore, this synthetic strategy holds the promise of broad applicability. However, we note that further tailoring and breakthroughs could be achievable by properly understanding the mechanism of formation of these multifunctional structures, to enable the atomic-efficient design of even more complex structures, thus expanding the ability of materials chemists to create increasingly advanced materials systems.

Figure 4. (a) Voltage versus specific capacity profiles, (b) evolution of specific capacity (solid symbols) and Coulombic efficiency (open symbols) when cycling at C/10, and (c) discharge rate capability performance, measured at room temperature. (d) Evolution of specific capacity (solid symbols) and Coulombic efficiency (open symbols) when cycling at C/10 at 50 °C.

specific capacity from 113 mAh/g to 107 mAh/g after the addition of inactive Al3+ ions. Two reversible processes, centered at 4.02 and 4.15 V vs Li+/Li0 on charge, were observed. These electrochemical results confirm that the bulk of the samples was composed of Li1+xMn2−xO4,32,33 and that the defects observed by STEM did not significantly alter their electrochemical properties. The profiles did not change significantly after extensive cycling (see Figure S6a in the Supporting Information). The specific capacity retention of CES NCs was 94% after 100 cycles, compared to 89% for the bare NCs (see Figure 4b). The Coulombic efficiencies of bare and C-ES NCs were 92% and 94%, respectively, at cycle 1, and stabilized at just over 99% upon cycling in both cases. The somewhat-higher initial Coulombic efficiency of the C-ES, compared to bare NCs, could be indicative of reduced side reactions at the electrode−electrolyte interface, because of the 397

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ASSOCIATED CONTENT

S Supporting Information *

Material characterization and methods (XRD, SEM, TEM, XAS, and ICP). This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Author Contributions

C.K., L.X., A.D., R.B., and J.C. conceived of and planned the experiments. C.K. and L.X. prepared the samples and carried out the measurements for the electrochemical performances. C.K. and P.J.P. analyzed microstructures of precursors and (CES) NCs using STEM, under the supervision of R.F.K. C.K., P.J.P., R.F.K., and J.C. prepared the manuscript, which incorporates critical input from all authors. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The design of C-ES NCs was initiated by L.X., C.K., and J.C. at Lawrence Berkeley National Laboratory, under support by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy under the Batteries for Advanced Transportation Technologies (BATT) program. Portions of this work were carried out at the Molecular Foundry as a user funded by the DOE Office of Science, Office of Basic Energy Sciences. These funding was provided under Contract No. DE-AC0205CH11231. The acquisition of the UIC JEOL JEMARM200CF was supported by a MRI-R2 grant from the National Science Foundation [No. DMR-0959470]. Support from the UIC Research Resources Center is also acknowledged. The electron microscopy portion of the work, carried out at UIC by P.J.P., C.K., R.F.K., and J.C., was performed with support from the Joint Center for Energy Storage Research, an Energy Innovation Hub funded by the U.S. Department of Energy, Office of Science, Basic Energy Sciences. Portions of this research were carried out at the Stanford Synchrotron Radiation Lightsource, a Directorate of SLAC National Accelerator Laboratory and an Office of Science User Facility operated for the U.S. Department of Energy Office of Science by Stanford University. The authors are grateful to Prof. Clare P. Grey (University of Cambridge, U.K.) for valuable discussions.

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