Reversible Magnesium Intercalation into a Layered Oxyfluoride

Dec 18, 2015 - Advanced Energy Materials 2016 6 (14), 1600140 ... Advanced hybrid battery with a magnesium metal anode and a spinel LiMn 2 O 4 cathode...
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Reversible Magnesium Intercalation into a Layered Oxyfluoride Cathode Jared T. Incorvati,†,∥ Liwen F. Wan,‡,∥ Baris Key,§,∥ Dehua Zhou,§ Chen Liao,§,∥ Lindsay Fuoco,†,∥ Michael Holland,† Hao Wang,§,∥ David Prendergast,‡,∥ Kenneth R. Poeppelmeier,†,§,∥ and John T. Vaughey*,§,∥ †

Department of Chemistry, Northwestern University, Evanston, Illinois 60208, United States The Molecular Foundry, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States § Chemical Sciences and Engineering, Argonne National Laboratory, Lemont, Illinois 60439, United States ∥ The Joint Center for Energy Storage Research, Argonne National Laboratory, Lemont, Illinois 60439, United States ‡

S Supporting Information *

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compounds for medical battery applications, we studied the mild fluorination of α-MoO3 to slightly reduce the framework and increase the electronic conductivity.15,16 Of the materials isolated, MoO2.8F0.2 was the only one that maintains the layered structure of the parent oxide.17−19 In this report, we investigate its stability as a Mg battery cathode, the mechanism of insertion, and model its structure to show that by moving to a second row transition metal, oxide abstraction has been eliminated as a failure mechanism and a potential new class of Mg battery cathodes has been identified. Synthetic procedures are included in the Supporting Information. Phase purity was determined by powder X-ray diffraction (PXRD) using a Rigaku Ultima diffractometer. All products were found to be phase pure by powder X-ray diffraction methods (PXRD) when compared to literature reports (see SI, Figure S1). PXRD diffractograms of the as synthesized MoO2.8F0.2 were indexed on the basis of an orthorhombic cell with a = 3.877, b = 14.043, c = 3.724 Å. Small variations in unit cell volume and some peak broadening in larger sample batches may be evidence of a small variability in the material’s fluoride content. Although this will likely not affect the electrochemical properties of the material, it may have a small effect on the electronic conductivity and any possible supercell that forms on Mg insertion. Scanning electron microscopy images show MoO2.8F0.2 crystals form as plates tens of micrometers across (see SI, Figure S2). Cathode and cell preparation are described in the Supporting Information. Galvanostatic charge/discharge cycling is shown in Figure 1 and shows the cycling capacity of MoO2.8F0.2 using a 0.2 M Mg(TFSI) 2 in propylene carbonate (PC) electrolyte. MoO2.8F0.2 shows dramatically higher capacity and better capacity retention when compared to the isostructural parent compound α-MoO3. Similar cycling capacities were found for 0.2 M Mg(TFSI)2 and Mg(Triflate)2 in either PC, diglyme or dimethylformamide. The only exception noted was cells based on acetonitrile, as they showed significant cycling inefficiencies and poor cycling symptomatic of solvent stability issues. The observed capacities in Figure 1 correspond to approximately

odern portable devices play a large role in everyday lives, and electric vehicles are becoming more commercially viable with every passing year. Although lithium-ion batteries have been the premier technology in these and other secondary battery applications, they have certain limitations.1,2 Lithiumion batteries use carbon anodes, rather than metallic lithium, which limit full cell lithium-ion energy density. In general, metal anodes make cell design and manufacture easier; however, a generally accepted strategy for using metallic lithium anodes does not exist for several reasons including surface instabilities and dendrite formation on cycling, which combine to limit overall cell safety and cycle life.3,4 Because Mg metal does not have the propensity to form dendritic structures on electrodeposition, is more abundant than lithium, has a reduction potential of −2.372 V vs SHE, and has a higher volumetric capacity than lithium metal, alternative energy storage systems based on magnesium should have certain advantages.5 Since the discovery by Aurbach et al.6 of sulfide-based Chevrel Phase cathodes compatible with electrolytes based on organo-magnesium halide complexes, work with Mg-ion batteries has been focused on identifying higher voltage cathodes and more stable electrolytes. Although electrolyte development has focused on removing chloride ions in order to support reversible electrochemistry at the higher voltages, multivalent cathode development has focused on oxide cathodes.7 Several materials, including MnO2, V2O5, and other traditional lithium-ion oxide cathodes, have been reported but the transition to magnesium has been complicated for several reasons including higher self-discharge, irreversible Mg insertion, and instability toward electrolytes.8−11 Many of these failure mechanisms involve side reactions that rely on oxide abstraction as a pathway to render the cathode material inactive. Because it is well established that first row transition metal oxides are more prone to oxygen loss than second row metals as a form of charge compensation on reduction, we undertook a study to evaluate second row metal oxides as cathode materials.12,13 An early candidate is the layered metal oxide α-MoO3. Although it is known in lithium-ion systems to be electrochemically active, earlier studies in Mg-based cells showed that it had limited stability and cyclability, even at elevated temperatures.14 Using synthetic strategies previously employed to synthesize several new metal oxyfluoride © XXXX American Chemical Society

Received: July 28, 2015 Revised: December 7, 2015

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DOI: 10.1021/acs.chemmater.5b02746 Chem. Mater. XXXX, XXX, XXX−XXX

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(200) peak suggests Bragg reflection from molybdenum atoms. This peak indicates a symmetry-breaking shift in the relative ordering of MoO2.8F0.2 layers. The discharged phase was indexed using JADE software package’s whole pattern fitting. Figure 2 highlights that at very low levels of insertion, a twophase insertion is evident as the middle pattern (approximately 0.05 Mg2+ inserted) shows peaks from the initial host material and the 0.09 Mg2+ endmember pattern in an approximate 1:1 ratio. This indicates that the solid solution indicated by the electrochemical evaluation may have Mg0.09[MoO2.8F0.2] as an endmember or that the Mg insertion kinetics are very slow at the rates the samples were evaluated (2 μA/mg). In light of the structural reversibility and the slow rate used, we infer that full range cycling shown in Figure 1 includes a region where twophases coexist above 0.09 Mg2+ inserted and is the active phase for the higher insertion levels. In addition to the PXRD and electrochemical studies, we utilized 25Mg MAS NMR spectroscopy on the discharged cathode samples to get a better understanding of the Mg environment in the host lattice. In Figure 3, 25Mg NMR spectra

Figure 1. (A) Voltage Profile for MoO2.8F0.2 over the first 18 cycles and (B) capacity versus cycle number. Electrolyte is 0.2 M Mg(TFSI)2 in PC.

0.25Mg2+ inserted into the host lattice per formula unit. The sloping discharge curves are indicative of a single phase solid solution process at lower levels of insertion. Powder X-ray diffraction (PXRD) studies were performed on laminated cathodes before and after cycling, as shown in Figure 2. After processing to make an electrode, the laminates of Figure 3. 25Mg MAS NMR spectra of the solid state synthesized Mg2Mo3O8 (red and magenta) and the electrochemical cycled sample MgxMoO2.8F0.2 (blue) collected at 11.7 T (red) and 19.89T (blue and magenta). Signal is in reference to saturated MgCl2(aq) at zero ppm and a recycle delay of 1s is used. Spinning speeds of 16, 20, and 14 kHz are used respectively for blue, red and magenta spectral. The spectra intensities are not normalized. * spinning sidebands.

of electrochemically magnesiated molybdenum oxyfluoride is compared to a magnesium molybdenum oxide model compound (Nolanite-type Mg2Mo3O8) with a symmetric tetrahedral and an asymmetric Mg site. The small sharp peaks at higher frequency in Mg2Mo3O8 spectra (26 ppm) are due to MgO from exposure to air. The NMR signal of intercalated Mg is expected to shift to much higher or lower resonance frequencies owing to increased paramagnetic contribution from Mo-d electrons (i.e., Fermi-contact shifts due to Mo4+, d2 for Mg2Mo3O8). However, in the 25Mg NMR spectrum of the model compound Mg2Mo3O8, the quadrupolar 25 Mg NMR lineshapes for the two lattice Mg sites appear at around −60 ppm and are due Mg ions coordinated to the molybdenum oxide structure. For the electrochemically cycled sample Mgx[MoO2.8F0.2] (x ∼ 0.16, from electrochemical testing), a broad peak appears at a similar shift of −50 ppm, which is most likely due to disordered Mg intercalation in the MoO2.8F0.2 structure. The small peak at 26 ppm was assigned to MgO impurities and a large sideband envelope in the

Figure 2. PXRD of pristine, synthesized MoO2.8F0.2, MoO2.8F0.2 laminate, laminate cathodes cycled 12 times to 15 and 30 mAh/g, and a recharged cathode. * indicates an Al substrate peak.

uncycled MoO2.8F0.2 show significant preferred orientation consistent with the roll pressing used to reduce the porosity of the formed electrode and the material’s plate-like morphology. After cycling, the structure of MoO2.8F0.2 appears to change. Reflections from MoO2.8F0.2 shift to higher d-spacing, which is consistent with insertion of Mg into the interlayer galleries. On discharge, peaks that are symmetry forbidden in the fluorobronze appear, including the (200). The high intesity of the B

DOI: 10.1021/acs.chemmater.5b02746 Chem. Mater. XXXX, XXX, XXX−XXX

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Chemistry of Materials electrochemical sample may be from air exposure during transfers or surface reactions with adventitious water from the electrolyte. The position of the peak is therefore found to be consistent with the location of the Mg-ion in the lattice as identified from literature structures and models.20 19F MAS NMR suggests no coordination changes for lattice fluorines such as formation of MgF2 upon magnesiation (see SI, Figure S3). No intercalation of protons or solvent species was detected via 1H MAS NMR (see SI, Figure S4). It is not clear exactly where fluoride resides in MoO2.8F0.2; however, there are three types of oxygen in the pristine αMoO3 structure, as shown in Figure 4. When fluoride is

Figure 5. Total density of states for single fluorine substituted (αMoO3)3×1×3 structure. Inset renders the charge distribution (yellow region) around the Fermi level. The top frame presents the entire supercell (in ab plane) that shows the charge is localized on a single Mo−O layer where fluorine (shown in blue) is substituted. The bottom frame plots the charge spread in this single Mo−O layer (in ac plane).

fluorine atoms far apart within the supercell, with each occupying an O2 site. Therefore, we conclude that at higher fluorine concentration, such as the experimental MoO2.8F0.2 composition, the fluorine atoms are most likely to be randomly and homogeneously distributed among the O2 sites, with little chance of clustering. This is consistent with the lack of changes observed via 19 F NMR of lattice fluorine sites upon magnesiation of interlayer galleries which should heavily affect O1 sites. In addition, when the F concentration is increased from 0.028 (1 F per (α-MoO3)3×1×3 supercell) to 0.2 (7 F per (α-MoO3)3×1×3 supercell), the density of states within the band gap (shown in Figure 5) grows and broadens significantly, as might be expected for this molybdenum oxide bronze, thereby reducing the optical band gap of the material. On the basis of literature precedent and experimental observations, we have identified several Mg battery cathode material failure mechanisms are attributable to the ease of loss of oxygen (oxide) from the host lattice. By moving to secondrow transitional metals, we can mitigate this as a significant effect. We studied several Mo oxides and have identified MoO2.8F0.2 as a promising new Mg-ion cathode material, as it possesses a layered structure with a measured band gap of ∼0.2 eV.19 A combination of electrochemical testing, X-ray diffraction, 25Mg NMR, and first-principles electronic and atomic structural modeling has been used to show that, unlike previous nonaqueous systems, this system works by an intercalation mechanism versus the usually observed disproportionation or conversion reactions.

Figure 4. Crystal structure of α-MoO3. The three types of oxygens are shown in different colors (O1: light blue. O2: gold. O3: red). The Mo atoms are shown in purple.

introduced, it is expected to substitute for oxygen and as a consequence modify its electronic interactions with the neighboring Mo. Here we perform ab initio electronic structure calculations using density functional theory (DFT)21 to investigate the structural variations of fluorine substituted αMoO3. Calculation methods are described in the Supporting Information. To study the preferential site for fluorine substitution, we place one fluorine atom (at different oxygen sites) into an approximately cubic 3 × 1 × 3 supercell of α-MoO3 and compare the resulting formation energies. The calculated fluorine at the O1 or O3 site will result in at least a 0.5 eV higher energetic cost, significantly reducing their probability (see SI, Table 1). When fluorine is substituted into the O2 site, it induces strong lattice distortions. Unlike the pristine α-MoO3 structure where O2 is shared by two neighboring Mo with dramatically different bond lengths (Mo−O1, 1.70 Å; Mo−O2, 1.73 and 2.225 Å; Mo−O3, 1.95 and 2.34 Å), the Mo−F bond lengths are approximately the same, 2.13 Å. This means the cation−cation distortion along the a-axis is significantly reduced. In addition to the lattice distortion, the introduction of fluorine also liberates an electron that induces defect states inside the electronic band gap as shown in Figure 5. This extra electron is delocalized over the entire Mo−O layer in the ac plane where the fluorine defect resides and fills up Mo dxz states that are hybridized with O2 pz states and O3 px states. It is observed that upon fluorine substitution, the Mo−F bond length is notably longer than the original Mo−O2 bond, which indicates a destabilized π* antibonding interaction between Mo dxz states and anion pz states. In the above simulations, only one fluorine atom is introduced into the 3 × 1 × 3 supercell of α-MoO3. We also test the possibility of forming F−F pairs, i.e., two fluorine atoms within a distance of ∼4 Å. It is found that the lowest energy is always achieved by separating the two

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b02746. S1: Rietveld refinement of MoO2.8F0.2. S2: SEM images of MoO2.8F0.2 before and after grinding. S3: 19F MAS NMR spectra. S4: 1H MAS NMR spectra. Table 1: Formation energies of fluorine substitution. Synthetic C

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cathode material for rechargeable Mg batteries. Electrochem. Commun. 2012, 23, 110−113. (11) Novák, P.; Scheifele, W.; Haas, O. Magnesium insertion batteries  an alternative to lithium? J. Power Sources 1995, 54, 479− 482. (12) Wiley, J. B.; Poeppelmeier, K. R. Reduction chemistry of platinum group metal perovskites. Mater. Res. Bull. 1991, 26, 1201− 1210. (13) Katzke, H.; Schlogl, R. Mechanism of the morphotropic transformation between the rutile and corundum structural types. Acta Crystallogr., Sect. B: Struct. Sci. 2003, 59, 456−462. (14) Spahr, M. E.; Novák, P.; Haas, O.; Nesper, R. Electrochemical insertion of lithium, sodium, and magnesium in molybdenum(VI) oxide. J. Power Sources 1995, 54, 346−351. (15) Sauvage, F.; Bodenez, V.; Vezin, H.; Albrecht, T. A.; Tarascon, J.-M.; Poeppelmeier, K. R. Ag4V2O6F2 (SVOF): A High Silver Density Phase and Potential New Cathode Material for Implantable Cardioverter Defibrillators. Inorg. Chem. 2008, 47, 8464−8472. (16) Donakowski, M. D.; Görne, A.; Vaughey, J. T.; Poeppelmeier, K. R. AgNa(VO2F2)2: A Trioxovanadium Fluoride with Unconventional Electrochemical Properties. J. Am. Chem. Soc. 2013, 135, 9898−9906. (17) Pierce, J. W.; Vlasse, M. The crystal structures of two oxyfluorides of molybdenum. Acta Crystallogr., Sect. B: Struct. Crystallogr. Cryst. Chem. 1971, 27, 158−163. (18) Sleight, A. W. Tungsten and molybdenum oxyfluorides of the type MO3-xFx. Inorg. Chem. 1969, 8, 1764−1767. (19) Pierce, J. W.; McKinzie, H. L.; Vlasse, M.; Wold, A. Preparation and properties of molybdenum fluoro-bronzes. J. Solid State Chem. 1970, 1, 332−338. (20) Wang, H.; Senguttuvan, P.; Proffit, D. L.; Pan, B.; Liao, C.; Burrell, A. K.; Vaughey, J. T.; Key, B. Formation of MgO during Chemical Magnesiation of Mg-Ion Battery Materials. ECS Electrochem. Lett. 2015, 4, A90−A93. (21) Kohn, W.; Sham, L. J. Self-Consistent Equations Including Exchange and Correlation Effects. Phys. Rev. 1965, 140, A1133− A1138.

procedures, elemental analysis, tables of DFT results, and notes on cell design and construction (PDF).

AUTHOR INFORMATION

Corresponding Author

*J. T. Vaughey. Email: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported as part of 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. We acknowledge the use of the Center for Nanoscale Materials, supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. The computational work is performed through a user project at the Molecular Foundry using the local cluster (vulcan and catamount) that is managed by the High Performance Computing Services Group at Lawrence Berkeley National Laboratory supported by the Office of Science of the U.S. Department of Energy under Contract DE-AC0205CH11231. Use of the Advanced Photon Source at Argonne National Laboratory was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-06CH11357. High field NMR access at Environmental Molecular Sciences Laboratory at Pacific Northwest National Laboratory is gratefully acknowledged. This work made use of the J.B.Cohen X-ray Diffraction Facility supported by the MRSEC program of the National Science Foundation (DMR-1121262) at the Materials Research Center of Northwestern University. We thank Dr. A. Lipson and Dr. D. Proffit for their electrochemical insight and Dr. N. Sa and Dr. V. Duffort for providing materials.

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REFERENCES

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DOI: 10.1021/acs.chemmater.5b02746 Chem. Mater. XXXX, XXX, XXX−XXX