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July 14, 2015

C 2015 American Chemical Society

Nanostructured Layered Cathodefor Rechargeable Mg-Ion BatteriesSanja Tepavcevic,*,† Yuzi Liu,† Dehua Zhou,‡ Barry Lai,§ Jorg Maser,§ Xiaobing Zuo,§ Henry Chan,^

Petr Kral,^, ) Christopher S. Johnson,‡ Vojislav Stamenkovic,# Nenad M. Markovic,# and Tijana Rajh*,†

†Center for Nanoscale Materials, ‡Chemical Sciences and Engineering Division, §X-ray Science Division, and #Materials Science Division, Argonne National Laboratory,9700 S. Cass Avenue, Argonne, Illinois 60439, United States and ^Chemistry Department and )Physics Department, University of Illinois at Chicago,Chicago, Illinois 60607, United States

In recent years, significant attention hasbeen devoted to the development ofrechargeable batteries beyond Liþ-ion

systems, in order to alleviate energy storageneedsassociatedwith the anticipatedgrowthin renewable energy systems.1 Developmentof batteries that implement multivalenttransporting ions such as environmentallyfriendly (nontoxic) and highly abundantmagnesium (2.9% as compared to 0.002%for Li) is an exciting opportunity because atwo-electron Mg/Mg2þ redox reaction canpotentially lead to a high theoretical volu-metric capacity (3832 mAh cm�3), thusrivaling that of Li-metal batteries.2 Contraryto lithium batteries, however, the develop-ment of Mg rechargeable systems (or anyother multivalent system) did not makesignificant progress in the past decadedue to mainly two reasons: (i) kineticallyhindered intercalation and diffusion ofMg ions within cathode materials and (ii)incompatibility of high-performance elec-trolytes with metallic anodes and/or high-voltage cathodes.3�5

In 2008, Aurbach et al.6 demonstrated thefirst viable rechargeable magnesium bat-tery technology with an electrolyte solutionbased on Mg organohaloaluminate salts, amagnesium metal anode, and a MgxMo3S4(Chevrel) cathode. Magnesium electrodesbehave reversibly in Grignard reagent solu-tions in ethers (RMgX, R = alkyl, aryl groupsand X = halide, Cl, Br, etc.) likely due tomagnesium inactivity toward ethers andRMgX compounds.7 However, despite thehigh reversibility ofMgelectrodes inGrignard/ether solutions, organohaloaluminate/etherelectrolytes exhibit a relatively narrow electro-chemical stability window (up to 2.2 V vsMg),thus greatly limiting the choice of cathodes forMg batteries.5 In order to improve Mg batteryperformance, new electrolyte solutions witha wide electrochemical window8 as well asnovel polymeric gel electrolytes9 are begin-ning to emerge.10 However, all currently pro-posed electrolytes are based on ethers, andfrom the viewpoints of safety, reliability, andtheir compatibilitywithbothanodesandhigh-voltage cathodes, they are still not adequate.

* Address correspondence toT. Rajh (rajh@anl.gov);S. Tepavcevic (sanja@anl.gov).

Received for review April 23, 2015and accepted July 14, 2015.

Published online10.1021/acsnano.5b02450

ABSTRACT Nanostructured bilayered V2O5 was electrochemically deposited

within a carbon nanofoam conductive support. As-prepared electrochemically

synthesized bilayered V2O5 incorporates structural water and hydroxyl groups,

which effectively stabilizes the interlayers and provides coordinative preference to

the Mg2þ cation in reversible cycling. This open-framework electrode shows

reversible intercalation/deintercalation of Mg2þ ions in common electrolytes such

as acetonitrile. Using a scanning transmission electron microscope we demonstrate that Mg2þ ions can be effectively intercalated into the interlayer

spacing of nanostructured V2O5, enabling electrochemical magnesiation against a Mg anode with a specific capacity of 240 mAh/g. We employ HRTEM and

X-ray fluorescence (XRF) imaging to understand the role of environment in the intercalation processes. A rebuilt full cell was tested by employing a high-

energy ball-milled Sn alloy anode in acetonitrile with Mg(ClO4)2 salt. XRF microscopy reveals effective insertion of Mg ions throughout the V2O5 structure

during discharge and removal of Mg ions during electrode charging, in agreement with the electrode capacity. We show using XANES and XRF microscopy

that reversible Mg intercalation is limited by the anode capacity.

KEYWORDS: nanostructured electrodes . electrochemical synthesis . bilayered V2O5. hydrated oxide . magnesium ion battery .

XRF mapping of transporting ions . HAADF

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The use of ionic materials for the intercalation ofmagnesium ions with a large charge density, however,causes a substantial polarization on the cathode. Inmetal oxide electrodes, intercalation of ions with alarge charge density leads to relocalization of theelectron density in the host material from collectiveorbitals onto oxygen p-orbitals. As such, the resultingpolarized host electrode is not well suited to facilereversible intercalation processes. In particular, densecrystalline structures of metal oxides limit the diffusionof Mg ions. The common approach to bypass thislimitation is to use materials with “soft-bonding” ionchannels able to modulate polarization effects. Forexample, chevrel phases (CP) MxMo6T8 (M = metal,T = S, Se)6 in their microcrystalline state or TiS2nanotubes11 or MoS2 nanoribbons12 allow for fastand reversible Mg insertion at ambient temperatures.However, a fundamental disadvantage is their lowaverage voltage of ∼1.1 V vs Mg/Mg2þ. On the otherhand, polyanion framework electrode materials, suchas strongly bonded Mg1.03Mn0.97SiO4, have a highredox voltage, but they suffer from sluggish Mg dif-fusion kinetics, caused by both strong electrostaticinteractions and low electronic conductivity, resultingin poor charge- and discharge-rate capability.13

Among various cathode materials considered forMg-ion batteries, vanadium-based oxides have beenintensively studied for many decades, due to theavailability of multiple valence states (V5þ f V3þ),offering the potential for high energy density asso-ciated with multiple electron transfer.14 The usualsolid-state synthesis of vanadium pentoxide (V2O5)via the hydrothermal heating of aqueous solutionsleads to the well-known orthorhombic V2O5. Manyother structures based on vanadium oxide have beendescribed in the literature, having different types ofvanadium coordination polyhedra.2 Previous electro-chemical studies included V2O5/carbon composites,15

thin-film V2O5 electrodes,16 VOx nanotubes,

17 and Mgintercalated vanadium oxides MgxV2O5,

18 albeit withfading capacities upon prolonged cycling.In this work, we have used electrochemical synthe-

sis, which allows formation of unique nanostructuredbilayered V2O5 structure with wide and adjustableinterlayer spacing. This method also provided effectiveincorporation of defects, water, and hydroxyl groups,which, in turn, fostered electron transfer and (de)-intercalation of highly charged Mg2þ ions. We hypothesize that the presence of adsorbed and hydrogen-bonded water is important for a double-layer ca-pacitance behavior, while the presence of structuralhydroxyl groups is critical for the intercalation pro-cesses.We demonstrated that electrochemically syn-

thesized nanostructured V2O5 can effectively inter-calate Mg ions between its bilayers and can bedischarged against a Mg anode to create a fully

loaded Mg-containing cathode in the discharged state(240 mAh/g). We employed HRTEM and X-ray fluo-rescence (XRF) imaging to understand the role ofenvironment in the Mg-ion intercalation processes.Classical molecular dynamics simulations were usedto understand the structure of a Mg-intercalated V2O5

cathode. This cathode was successfully coupled in arebuilt full cell with a high-energy ball-milled Sn/Ccomposite anode in high-voltage/conductivity electro-lytes, such as acetonitrile withMg(ClO4)2 salt, to reach aconsiderable reversible capacity that improves withcycling at elevated voltages. We have also shown,using X-ray absorption near-edge structure (XANES)and XRF, that reversible Mg intercalation is limited bythe anode capacity.

RESULTS AND DISCUSSON

Electrochemical Preparation of Nanostructured V2O5 Electro-des. One of the main issues in the application of nano-structured materials for energy storage is the integra-tion of nanoscale materials into mesoscale hierarchicalarchitectures used in real devices. Electrochemicalsynthesis intrinsically results in nanoscale architectureselectronically coupled to conductive supports, allow-ing their integration with other battery components,without the need for further post-treatments beforeutilization.19 The electrons that drive electrochemicalsynthesis create reactive species that propagate che-mical reactions fostering electrical connectivity to theextent that electrons can penetrate the structure.Therefore, electrodes grown by an electrochemicalsynthesis naturally evolve toward an optimized struc-ture, which can achieve a high power, energy density,and stability, needed in the next generation of hybridsystems. On the other hand, an electronically inter-connected nanoporosity allows short diffusion lengthsof transporting ions and a full participation of thewhole cathode in achieving its theoretical capacity.

In our earlier work,19 we synthesized nanostruc-tured V2O5 by electrodeposition from vanadyl sulfateon a Ni substrate and obtained a thin film of nano-structured V2O5. Herein, we developed a new proce-dure for electrodepositing bilayered V2O5 within ahighly conductive commercially available substrate:carbon nanofoam (CNF). Figure 1 shows an HRTEMimage of an “as-prepared” individual V2O5 nanoribbonin conjunction with a SEM image of an interconnectednetwork of nanoribbons within a CNF framework(inset, Figure 1). Because of the electronic conductionand high porosity of CNF, it is possible to increase themass loading of an active material up to 30 mg/cm2,without any mechanical problems. This easy bulkelectrodeposition method allows straightforward scal-ing up of electrodes for a large-scale production.Importantly, this synthetic approach provides an in-timate direct contact of bilayered V2O5 with the carbonsupport, without the addition of binders or conductive

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high surface area carbon diluents. The electrochemicalsynthesis within a conductive CNF developed in thiswork provides an exceptionally large interlayer dis-tance of 13.5 Å of V2O5, compared to d = 8.8�11 Å19

obtained by other methods. This large interlayer dis-tance was found to be crucial for an efficient cycling oflarge Na ions, due to balancing electrostatic forcespresent during cycling. We hypothesize that the largespacing and structural adaptability of V2O5 might leadto reversible cycling of small but polarizing Mg ions, asboth the dipole moment interaction � R�2 and thepolarizability component� R�4 that constitute materi-al polarization rapidly decaywithin the large interspacedistance between the V2O5 layers.

20

Mg2þ Enrichment. Electrochemical synthesis facili-tates also Mg enrichment of as-prepared V2O5. Bi-layered V2O5 electrodes are originally synthesized inthe charged state; in order to operate effectively in theMg-ion full cell, it is necessary to introduce the sourceof magnesium into the cathode. An electrochemicalpreconditioning approach was used to incorporate Mgions in the cathode: V2O5 half-cells were dischargedgalvanostatically at 20 μA to the potential 0.2 V versus

Mg/Mg2þ. The first galvanostatic discharge voltageprofile of the Mg half-cell (with polished Mg foil as ananode) is shown in the inset of Figure 2. The voltagestarts out at an OCV of ∼1.7 V and then progressivelyis driven to a lower cutoff potential of 0.2 V vs theMg/Mg2þ anode. The specific capacity of 240 mAh/gwas obtained in this preconditioning procedure,suggesting reduction of vanadium to V4þ. This oxida-tion state was also confirmed by direct measure-ments of the V oxidation state using X-ray absorption

spectroscopy (XAS),21 as the main edge position ofthe discharged sample was close to that of VO2 (V

4þ

oxidation state standard). It should be mentioned thatthis oxidation state also suggests MgV2O5 stoichiome-try (one Mg per V2O5 molecule). Furthermore, theintensity of the V K-edge pre-edge peak, related tothe probability of the forbidden sf d dipole electronictransition of vanadium, decreases with the insertion ofMg ions during preconditioning. The lower intensity ofthis peak reflects the formation of a centrosymmetricoctahedral symmetry of vanadium sites, suggestingthe ordering of a local vanadium environment uponinsertion of Mg.

The distribution of Mg transporting ions withinV2O5 was furtherer investigated using XRFmicroscopy.The importance of direct imaging of the distribution oftransporting ions in conjunction with elemental map-ping of the electrode material is that element colo-calization studies directly report the mechanism ofcharge compensation. Transporting ions such as Liþ

or Naþ have too low fluorescence yields and efficien-cies for detection; however, Mg2þ ions aremore readilymapped within the matrix of the host V2O5. Figure 3shows a fluorescence image of Mg and V, in a Mg-enriched rod-like particle extracted from CNF. Thelinear color scale ranges from black (low intensity) tored (high intensity). The resulting fluorescence inten-sity map provides information on the total Mg contentin the volume illuminated by the beam at each scan-ning step. The images show that distributions of bothMg and V mostly follow intensities expected for a rod-like object: intensities peaked in the middle, where thesample was thickest, and decreased toward the edges,where the sample was thinnest (the thickness of thesample is shown by X-ray transmittance, Figure S1),suggesting that both transporting ions (Mg) and hostelements (V) are colocalized throughout the structure,confirming intercalation of Mgwithin the V2O5 sample.

Figure 2. Half cells: Normalized V K-edge XANES for as-prepared, bilayered V2O5 (V

5þ), V2O5 after Mg enrichment(V4þ), V2O5 after Na enrichment (V4þ and V3þ), and twostandards (V5þ and V4þ). Inset: Chronopotentiometric curvesofMg insertion intonanostructuredV2O5 in 1MMg(ClO4)2/AN.

Figure 1. HRTEM image of a nanoribon of bilayered V2O5

electrochemically grown on a carbon nanofoam (CNF) sub-strate. Inset shows a SEM image of interconnected nano-ribon architecture electrochemically grown on CNF.

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While imaging of the Mg distribution providesdirect information about the distribution of Mg afterinsertion within the sample, the quantitative interpre-tation of the results needs further considerations. Mg isa light element, with a low energy of ionization, andemits low-energy X-ray fluorescence. As a conse-quence, the fluorescence radiation of Mg experiencesstrong self-absorption characteristic for imaging of alllight elements due to nonspecific absorption of low-energy X-rays through the sample.22 In our measure-ments the detector was positioned at a 15-degreetakeoff angle (right), and the intensity from the oppo-site (left) side of the rod was reduced due to enhancedself-absorption associated with larger path lengthsthrough the sample (Figure 3A). However, in contrasttoMg, the V X-ray fluorescence is not nearly as affectedby self-absorption due to higher energy X-ray fluores-cence, and the V image in Figure 3B is more consistentacross the sample. The overlay of the intensity map ofMg (green) and V (red) in Figure 3C reflects the signaldetection inhomogeneity arising from Mg fluo-rescence self-absorption. By carefully modeling theabsorption of low-energy Mg emission through a VO2

rod (SI, Figure S2 and Table S1), one obtains anabsorption correction factor and Mg/V ratio based onthe experimental geometry. The corrected molar ratioreveals a molecular formula of Mg0.7V2O5, suggestingan oxidation state of V close to þ4.3. This value is invery good agreement with sample average measure-ments of the Vþ4 oxidation state obtained using XANES(Figure 2) and reflects some distribution of the oxida-tion states at different parts of the V2O5 electrode.

Mg2þ-Ion Charge, Role of Water, and Solvent Co-intercalation.Relatively slow transport of Mg2þ ions in common hostswas shown to be related to their divalent character.

Electrochemically synthesized nanostructured V2O5 hasa very unique morphology and structure compared tosimilar structures createdby sol�gelmethods.23Uniqueto electrochemical preparation is formation of a highlyhydrated sample (Figure S3), with an exceptionallywide interlayer spacing of 13.5 Å due to incorpora-tion of structural water between the layers. Novak andDesilvestro24 studied electrochemical intercalation ofMg2þ into conventional orthorhombic V2O5 electrodesin an electrochemical cell using an electrolyte com-posed of 1 M Mg(ClO4)2/CH3CN (AN) with water addedto AN. Using cycling voltammetry the authors foundthat intercalation of Mg ions coupled to reduction ofV2O5 improved with the addition of water to theelectrolyte. They speculated that water molecules co-intercalate with the Mg2þ cations and provide a smallersolvation shell than the organic solvent molecules. Theyhave also shown that reduction of V2O5 correspondingto capacities of ∼170 mAh/g could be achieved; how-ever, the capacity drastically decayed upon furthercycling. The authors have shown that added water,essential for reversible intercalation ofMg2þ, is removedduring the cycling and released back to the electrolyte,resulting in the decay of electrode capacity.25

Electrochemical synthesis of bilayered V2O5 results,however, in incorporation of three types of water: (1)mostly free water reversibly adsorbed between theV2O5 layers; (2) bonded water incorporated into thelayers cross-linked through hydrogen bonds; and (3)structural hydroxyl groups typical for electrochemicallysynthesized oxides.26 We have shown using thermo-gravimetric analysis that bound water (reversibly ab-sorbed and hydrogen-bonded water) is removed uponannealing of as-prepared samples at 120 �C undervacuum (Figure S3). Removal of weakly bound water

Figure 3. XRFmaps of Mg (a) and V (b) in aMg-enriched sample discharged at 0.02 V vs theMg anode obtained using 5.45 nd10.1 keV X-rays, respectively. The images were derived from detectors placed on the right with a 15-degree takeoff. Blackpixels represent the lowest fluorescence intensity, while red pixels correspond to the highest. (c) Overlay of the Mg andV maps shows the effect of self-absorption of Mg fluorescence.

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only mildly alters the structure with contraction of theinterlayer spacing to a distance of 12 Å,27 while struc-tural hydroxyl groups28 still remain in the structure(Figure 4A) with a stoichiometry of V2O5 3 0.6H2O.

29 Thestructure becomes significantly altered, however, afterremoval of structural hydroxyl groups (achieved onlyupon annealing above 300 �C), when the bilayeredstructure collapses into crystalline orthorhombic V2O5

with poor electrochemical performance.We have used computational methods to simulate

V2O5 bilayered structures with different contents ofwater and obtained that the spacing between bilayerscorresponds to the spacing observed in X-ray diffrac-tion measurements (Figure S4) and HRTEM images(Figure 1). The amount of water that swells V2O5 toobserved interlayer spacing was found to correspondto ∼2 mol of H2O per mol of V2O5 and was confirmedby thermogravimetric analysis. We believe that thislarge amount of water can facilitate the intercalation of

Mg ions. It has been shown previously that solvation ofMg ionswithwater is energetically favorable in amixedsolvent. It was found that water molecules substituteacetonitrile in the first solvation shell in a linear fashionwith up to six water molecules per Mg ion.30 Althoughpreferentially coordinated with six molecules of water,the energies of water solvation quickly decrease withincreasing water content, forming a stable configura-tion with four water molecules in the first solvationshell.31 Due to the preferential solvation of theMg2þbywater compared to acetonitrile and greater affinity ofacetonitrile to water already coordinated to Mg2þ, thereversible co-intercalation of already hydrated Mg ionscan be expected.30 A water solvation shell partiallyshields the charge of the Mg2þ cations, decreasingpolarization effects and the strength of their interac-tion with apical oxygens and hydroxyl groups. Indeed,unlike added andweakly boundwater, strongly boundstructural hydroxyl groups remain in the structure

Figure 4. XPS spectra of bilayeredV2O5 before (A) and after 10 cycles of chargingwithMg2þ ions (B) in theO 1s core level regionin conjunctionwithmolecular dynamics simulation of the V2O5 bilayer immersed inwater (C) and in the presenceofMg ionswithV in theþ4 state in the presence of water (D). HRTEM image of Mg-enriched V2O5 (E) in conjunction with anHAADF image (F), inwhich the relativeposition of V andMgatomsobtainedwith EDS is shown in the inset (V: red,Mg: green). The latticemodel in theviewing direction [102] is also shown on the HAADF image, in which bright lines depict the V plane.

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during cycling (Figure 4B) and act as a lubricant forreversible (de)intercalation of solvated Mg2þ ions be-tween enlarged galleries of electrochemically synthe-sized V2O5. In contrast to water added to electrolytethat can enter and swell V2O5 but easily comes out ofthe structure upon cycling due to its weak interactionwith the metal oxide framework,24,25 structural hydroxylgroups are covalently linked toV2O5and thereforedonotget depleted during cycling. On the contrary, stronglybound hydroxyl groups can participate in hydrogenbonding of water molecules that hydrate Mg2þ ions,facilitating intercalation of integrated solvated Mg2þ.Thus, the interaction of hydrated Mg2þ with V2O5 willdepend on the relative ratio between the energies forfurther solvation of hydrated Mg2þ with acetonitrile andthe energy of interaction with structural hydroxyl groupsin abilayered V2O5 structure, and this ratiowill determinethe ability for reversible Mg2þ ion intercalation.

Similar reasoning can also account for previouslyshown behavior of microcrystalline V2O5, where addi-tion of water to the electrolyte solution improvedinsertion (albeit limited) of Mg ion.4 Indeed, we findusing molecular dynamics (MD) simulation of a V2O5

model slab composed of three bilayers immersed inwater that water swells to a structure with an interlayerspacing of ∼12 Å by taking up water to screen allexposed surfaces, reaching two layers of water perV2O5 slab (Figure 4C). MD simulations suggest thatupon insertion of Mg ions into the V2O5 structure, thespacing between bilayers narrows to∼11 Å due to theinteraction of Mg ions with bilayer apical oxygensand structural hydroxyl groups (Figure 4D). Althoughthe spacing decreases, MD simulations show that asignificant amount of water remains in the structure,solvating inserted Mg ions. Figure 4E presents theHRTEM image of the well-crystallized area showingthe lattice fringes of the layered V2O5 structure afterMg ions were intercalated in the structure featuring aspacing of 6.0 to 6.4 Å. It can be seen from the imagethat the bilayer distance is not uniform across thewhole imaged area, which is consistent with the broadpeak observed in the XRD measurements correspond-ing to the interlayer spacing. On close inspectionof the fringes by using a high-angle annular darkfield (HAADF) image using scanning transmissionmicroscopy (STEM) (Figure 4F) of the selective area(marked in Figure 4C with a white square) one can seethat each of the widely spaced fringes is actuallycomposed of two closely spaced bright lines of Vplanes with a spacing of 1.87 Å and two V planesseparated by 3.40 Å. These spacings are consistent withthe structure proposed by MD simulations shownin a viewing direction of [1 0 2] (Figure 4F), whichdepicts also a lattice model of monoclinic bilayeredbase-facing square-pyramidal V2O5 with lattice para-meters a = 11.65 Å, b = 3.62 Å, and c = 11 Å with β =88.63�. The signal intensity of V and Mg from an EDX

line scan along the blue line in the HAADF image isshown as the inset. We can see that the higher Mgintensity appears in conjunction with the lower Vintensity and vice versa. This indicates that the Mg ionsare inserted in between the V layers. The overall lowerintensity of Mg compared to V is again the conse-quence of self-absorption of low-energy X-raysemitted by Mg in a low takeoff angle in which thedetector is configured in the microscope.

Our MD simulations suggest that Mg ions areplaced close to the oxygen atoms that terminatebilayers (Figure 4D). The insertion of Mg ions in theV2O5 bilayer structure does not cause removal ofthe first coordination water shell that solvates theions. Although accompanied by closing of the gapbetween the bilayers, intercalation does not result inthe collapse of the bilayered structure. Upon inter-calation, some of the coordination sites of the Mg ionsare replaced with an interaction with terminal hydro-xyl groups from the bilayer. The interlayer water playstwo important roles: it maintains a sufficient inter-layered space to allow the physical diffusion of sol-vated Mg cations (effective size of the Mg ion withthe first shell water solvent is at least 6 Å)32 and assistsin stabilizing intercalated Mg2þ through dipole inter-actions. Removal of excess water reduces thesetwo functions and results in poor electrochemicalperformance of highly crystalline orthorhombic V2O5.

XRD and WAXS studies show that the vanadiumatomic environment within the bilayer is not signifi-cantly altered by the Mg insertion. Figure 5 showssmall-angle (SAXS) and wide-angle X-ray scattering(WAXS) of V2O5 before and during cycling in thed-spacing range from 20 to 9 and 6 to 1 Å, respectively.The spectrum in the SAXS region confirms MD simula-tions and the decrease of the interlayer spacing from13.1 Å to 11.0 Å upon intercalation of Mg ions. Whilediffraction features in the WAXS range show that theintralayer structure experiences changes of a tenth ofan angstrom, close inspection of the range 3.5�6 Åshowsmore significant changes withMg insertion. Thediffraction peak at 4.7 Å disappears upon Mg insertionin conjunction with the sharpening of the feature at3.96 Å.Modeling of X-ray diffraction using amonoclinicbilayered V2O5 structure indicates that these changesare consistentwith contractionof interlayer spacing andcolocalization ofMg ions near terminal hydroxyl groups.Insertion of Mg ions in the center of the interlayer gapwould result in the enhancement of a diffraction peakrepresenting half of the interlayer distance (5.9 Å [002])and this peakwould become themost intense feature inthe spectrum (as observed previously by intercalation ofNa ions that were placed in the middle of the interlayergap19). In the dischargedMgV2O5 sample, however, the[002] diffraction peak has a smaller intensity than [202],and both of these features have a smaller intensity thanthepeak at 3.48Å ([110]), which is themost intense peak

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in the spectrum, suggesting localization ofMg ions nearterminal hydroxyl groups.

The a and b lattice parameters of the unit cell, whichdetermine the intralayer dimensions, do not changesignificantly upon insertion, and all the bond lengthswithin the bilayer structure remain the same. The onlyparameter that is changed is the interlayer spacing, asdefined by the lattice parameter c, which decreasesupon intercalation of Mg2þ most probably due to thestrong interaction of the small doubly chargedMg ionsand terminal hydroxyl groups. Observed changes arevery different from the behavior found for intercalationof large Na atoms, where the size of the crystalline unitcell and the bond lengths were enlarged in order toaccommodate for the intercalation of a large Na ion.19

Full Mg Cells. Identification of a suitable negativeelectrode for cycling of magnesium is a critical issuefor the future development of magnesium ion energystorage devices. It has been shown that a Mg anodedoes not show good stability and reversibility due tothe formation of a nonionically conductive solid elec-trolyte interphase. For example, in the coin cell com-posed of V2O5 grown on CNF (∼10mg of V2O5) andMgfoil in 1 M Mg(ClO4)2/CH3CN electrolyte, Mg ionscannot be redeposited on the Mg anode after the firstdischarge, preventing further cycling (Figure S5). Whenthe Mg anode is, instead, replaced with graphite, Mgions could effectively be inserted into the graphiteanode but only irreversibly: they could not be removedin the following discharge step. When CNF (a hardcarbon disordered material) was used as the anodeitself, a reversible cycling of the fullMg cell was achieved(Figure 6A, green curve). Although the capacity of thecell was low, it was consistent with the limiting capacityof the CNF anode. Significant improvements of thebattery capacity were obtained when the anode waschanged to pressed nanocrystalline Sn: charge storage

was increased by 10 times. This clearly shows that thecapacity of the full cell in this case was strictly limited bythe anode capacity. Previously it has been also shownthat Sn readily forms alloys with metallic Mg, and thesealloys are used formany applications ranging from lead-free soldering to anodes in lithium ion batteries.33,34

Formation of chemical complexes such as Mg2Sn havebeen found to exist in the Mg�Sn melts, but thebonding was found to be only weakly interacting.35 Inthis work we used nanocrystalline Sn obtained by high-energy ball-milling of bulk metallic tin.36 Using Sn as ananode we demonstrated reversible charging of Mg-ionfull cells in common electrolytes such as acetonitrile andachieved a reversible capacity of 150 mAh/g at C/15rate. Our cycling voltammetry study (Figure 6B) ofnanostructured V2O5 showed pronounced peaks involumetric sweeps indicating reversible Faradaic reac-tion,while the featureless, rectangular shape for the CNFcurve confirmed the capacitive origin of the CNF signal.Increasing the cycling rate to C/8 initially decreases thecapacity to 100 mAh/g. Upon extended cycling, how-ever,weobserve self-improvingof theelectrode specificcapacity previously observed in nanocrystalline systems(Figure 6C and E, 50 cycles). Figure 6D shows thecorresponding dQ/dV plot of the third charge anddischarge process with two distinctive and highly re-versible main peaks in the range 0.0�2.0 V; one islocated at about 0.95 V, and the other is located atabout 1.25 V. Two peaks observed in the differentialcapacity curve are indicating that there are twoavailablesites for Mg occupancy at these voltages, and theirheight is proportional to their availability. Furthermore,we demonstrated single-phase solid-state intercalationwith the V2O5 cathode in a wide concentration range ofthe Mg2þ ions (Figure 7A).

The pre-edge structure in the XANES spectrumobtained during cycling shows that upon first insertion

Figure 5. SAXS and WAXS spectra of “as-prepared” (emerald) sample, Mg-enriched sample obtained by first discharge vsmagnesium foil (red), and Mg-depleted sample (black) in conjunction with corresponding structures obtained by moleculardynamics simulations. The annotation of the scattering peaks was performed using diffraction spectra simulations usingcorresponding structures.

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ofMg into the bilayered V2O5 structure, the presence ofMg ions causes ordering of bilayers that is reflectedby a decrease in the intensity of the Mg pre-edge(Figure 7B). Selected area electron diffraction (SAED)images also confirm ordering in the presence of Mgions (discharged state) where the whole grain showspolycrystalline morphology, in contrast to the chargedsample, where in the absence of Mg only a fewscattered larger size crystallites were observed at theedges of the sample (Figure 7C, left and center).Elemental analysis using the TEM EDAX detector showsthat an approximately 10 times larger concentration ofMg ions is in the discharged sample compared to thatin the charged one (Figure 7C, right). It is worthmentioning that no chlorine atoms were detected inthe elemental composition of the discharged (andcharged) sample, confirming that the presence of Mgin the V2O5 compound was due to the charge compen-sation of the reduced V4þ ion, rather than the residualMg(ClO4)2 salt used in theelectrolyte. Thepositionof the

main edge in the XANES spectrum shows the averagevalence state upon repeated discharge and charge ofthe sample from the þ4 to the þ4.8 oxidation state,respectively. These results suggest either that the capa-city was limited by the capacity of the Sn anode or thatsomeof theMg ionswere retained in the structure uponrepeated cycling (10 cycles). The latter case would notbe necessarily surprising given that polarizing Mg ionscould be retained in a highly polarized V2O5 environ-ment, requiring still significant activation energy inorder to be removed from the structure.

X-ray fluorescence microscopy studies also supportreversibleMg ion intercalationwith respect to V chargecompensation. Figure 8 shows the change of therelative concentration between Mg and V upon cy-cling. While in the initial Mg-enriched dischargedsample the concentration of Mg vs V reaches 1:3, theconcentration of Mg subsequently decreases uponcharging. The relative concentration of Mg vs V afterdeintercalation (during charging) was found to be 1:8,

Figure 6. Bilayered V2O5 cathode electrochemical performance. (A) First three charge�discharge cycles of Mg-enrichedbilayered V2O5 vs nanocrystalline Sn electrode. Charge�discharge curve of the same cathode vs CNF is shown in green forcomparison. Both cells were cycled at 20mA/g, at C/15 rate, within the potential window of 2.2�0.0 V (vsMg/Mg2þ) from 1MMg(ClO4)2 in acetonitrile. (B) Cyclic voltammetry of bilayered V2O5 grown on CNF (blue) and CNF in 1 M Mg(ClO4)2 inacetonitrile at 0.02mC/s scaning rate. (C) Electrode stabilitymeasured for 50 cycles. (D) Charge�discharge curves of the samecathode cycled at the C/10 rate shown at the 20th and 50th cycles. Inset shows electrode stability study. Bottom right:corresponding differential capacity (dQ/dV) plot for the V2O5/Sn full cell. (E) Charge�discharge curves of the 20th and 50thcycle. Note the increase in the capacity in the 50th cycle.

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suggesting a stoichiometry ofMg0.25V2O5 (Figure 8 andFigure S3, Table S2). This stoichiometry is in goodagreement with the determined average oxidationstate of þ4.8 obtained in our in situ XANES measure-ments. Although in very small concentrations, it seemsthat the map of Mg in the charged sample does notfollow the distribution of V or the sample thicknessmap; rather it shows diminished fluorescence inten-sity from top-protruding facets, suggesting a lowerdiffusion energy barrier for ions inserted in exposedterraces. We also show using XRF microscopy thatwhen nanocrystalline Sn is used as the anode, arelatively large concentration of Sn ions is transferredto the cathode area during cycling. While particlesfrom the V2O5 cathode cycled against Mg foil (1 cycle)did not contain any impurities, the same particles

cycled for 10 cycles against nanocrystalline Sn con-tained up to 100 μg/cm2 of Sn and Cu in their surfacelayers (Figure 8, right panel). Both Sn and Cu areelements that are used in making the anode sample,Sn is used as an active anode material, and Cu is usedas the electrode support, suggesting dissolution ofthe Sn anode in the acetonitrile electrolyte duringdischarging of the battery. These large ions, however,do not intercalate into the structure but concentrateat the nanoparticle interface, possibly contributing tothe capacitance behavior of the electrode.

CONCLUSIONS

In summary, we have shown efficient reversiblecycling of Mg rechargeable batteries in a commonelectrolyte, acetonitrile, byusingananosizedopen-frame

Figure 7. (a) Charge�dischargeprofile ofMg-enrichedbilayeredV2O5 vsnanocrystalline Sn anode after the 10th cycle, cycledat 20 mA/g, within the potential window of 2.0�0.0 V (vs Mg/Mg2þ) from 1 M Mg(ClO4)2 in acetonitrile. (b) Normalized VK-edge XANES for as-prepared, Mg-enriched, and charged/discharged (after the 10th cycle) bilayered V2O5 in a full cell with aSn anode. (c) Selected area electron diffraction andmagnesiumdistribution during different states of charging in conjunctionwith elemental analysis obtained using an EDAX detector in TEM.

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conformable V2O5 cathode paired with a Mg alloyforming a nanocrystalline tin anode. NanocrystallineV2O5 electrochemically deposited within a porouscarbon nanofoam substrate enables incorporation ofstructural hydroxyl groups and adsorbed water. Whilethe CNF matrix enabled superior conductivity with-out any binders or conductive additives, we believestrongly bound hydroxyl groups play two importantroles in reversible cycling: (i) theymaintain a sufficientinterlayered space to allow diffusion of Mg cations,and (ii) they reduce the symmetry of the V2O5

that contributes to the electron transfer and themagnesium intercalation. We have shown using MDsimulations that the presence of structural waterand ensuing hydration of Mg ions is necessary forreversible cycling of Mg ions. In response to Mgintercalation during the first Mg enrichment, the host,hydrated bilayered V2O5, compresses and orders alongthe c axes. Upon prolonged cycling (10�50 cycles) withMg ions, the host structure becomes more disorderedand polycrystalline but improves capacity for Mg2þ.

Taking advantage of the ability of X-rays to monitorboth the oxidation state of electroactive vanadiumatoms (XAS) and distribution of chemical species(XRF), we were able to correlate introduced chargewith the distribution of transporting ions and determinethe mechanism of charge compensation in Mg re-chargeable systems. By employing in situ XANES weshow that the oxidation state of V changes reversiblyand cycles between an average þ4.0 and þ4.8 oxida-tion states. XRF microscopy, on the other hand, showsthat Mg ions are distributed throughout the cathodein a discharged state in a near-stoichiometric ratio.Atomic precision of the elemental distribution usingHAADF confirms that Mg is uniformly intercalatedbetween two V-containing bilayers. This first demon-stration of direct mapping of transporting ions withinthe host lattice revealed successful intercalation ofMg ions throughout the cathode in a dischargedstate, while only traces of Mg are detected in thecharged cathode, as reversible Mg intercalation islimited by the Sn anode capacity.

METHODS

Synthesis of Nanostructured V2O5 Electrodes. Nanostructured V2O5

was synthesized by electrochemical deposition on commerciallyavailable carbon nano foam (Marcketech International, Inc.).

The electrochemical deposition was carried out in a two-elec-trode cell with CNF as theworking electrode and Ptmesh as botha counter and reference electrode in aqueous 0.1 M VOSO4

solution at a constant potential of 1.5 V. As-prepared bilayered

Figure 8. (a) XRFmaps of (left) Mg andV in a discharged sample and (right) charged sample obtained using 5.45 keV (top) and10.1 keV X-rays (bottom), respectively. The images were derived from detectors placed on the right with a 15-degree takeoff.Black pixels represent the lowest fluorescence intensity, while red pixels correspond to the highest fluorescence. (b) XRFimages of elemental distribution of V (top left), Sn (top right), and Cu (bottom left) and their overlay (bottom right) obtainedusing 10.1 keV X-rays.

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V2O5 electrodes were synthesized by vacuum annealing at120 �C for 20 h. The crystallized orthorhombic V2O5 electrodeswere obtained by annealing of as-prepared V2O5 under an O2

atmosphere at 500 �C for 4 h.Electrochemical Insertion/Extraction of Mg2þ and Characterization.

Magnesium half-cells were assembled in a He-filled dry glove-box into coin-type cells with a Mg foil as the negative electrode,an electrolyte of 1 M Mg(ClO4)2 (Aldrich) in acetonitrile (ACN),and a glassy fiber separator. In order to cycle, a full cell source ofMg in the system is necessary. We first discharged the V2O5

cathode to 0 V vsMg foil and then opened that cell in a gloveboxunder an inert atmosphere. Finally we reassembled this Mg-enrichedMgxV2O5 cathodewith a Sn anode into the full cell. FullMg cells were assembled in coin-type cells (Hohsen 2032) with amagnesium-enriched MxV2O5 nanoribbon as the positive elec-trode and high-energy ball-milled Sn as the negative electrode,a glass fiber separator (Whatman GF/F), and 1MMg(ClO4)2/ACNas electrolyte. The traditional Sn anode was made by mixing84 wt % active material (Sn powder, Aldrich), 4 wt % graphite(Timcal, SFG-6), 8 wt % poly- (vinylfifluoride) binder (Kynar), and4 wt % carbon black (Toka) and pressed on Cu foil. Full Mg cellswere cycled galvanostatically at varying currents between 2.3and 0 V vs Mg/Mg2þ, respectively, using an automated Maccorbattery tester at ambient temperature. All cell assembly anddisassembly operations were performed in an Ar-filled dryglovebox (oxygen level <2 ppm).

Synchrotron XRF Microscopy. XRF was performed at the 2-ID-Dbeamline at the Advanced Photon Source (APS) at ArgonneNational Laboratories, where an undulator source was used tocreate hard X-rays with energies of 1.75�10 keV and focusedusing Fresnel zone plate optics. Emitted X-ray fluorescence wasdetected using an energy-dispersive germanium detector(LEGe detector, Canberra).

Synchrotron SAXS/WAXS Measurements. SAXS/WAXS data werecollected at Beamline 12ID-B of the APS at the ArgonneNationalLaboratory. The X-ray was focused, and the spot size on thesample was ∼50 μm � 50 μm. SAXS and WAXS data werepresented in momentum transfer, q (q = 4π sin θ/λ, where θ isone-half of the scattering angle, and λ = 1.033 Å is thewavelength of the 12 keV energy probing X-ray), measuredin the range 0.01�2.3 Å�1. The charged samples were pre-pared by stripping V2O5 films onto Kapton tape. The dis-charged samples were scratched off and sealed inside a3 mm diameter hole in a piece of aluminum foil by sealingthe Kapton sheet to the foil using epoxy. All cell assembly anddisassembly operations were performed in a He-filled dryglovebox (oxygen level <2 ppm).

XPS. A Scienta hemispherical electron analyzer (SES100)wasused to obtain the XPS/UPS measurements, and total energyresolution of the spectra, including photon energy, was set toless than∼0.1 eV. The acceptance angle for incoming electronsis 5 degrees. All experimental data were taken under a pressureof 2 � 10�10 Torr or less.

Electron Microscopy. Scanning electron microscopy (SEM)images were recorded with a JEOL JSM-7500F field emissionSEMoperating at 30 kV. HRTEM imageswere recorded on a JEOLEM-2100F. HAADF images and line scans were obtained usingJEM-ARM200CF, a probe aberration-corrected 200 kV STEM/TEM with a cold field emission source with a 0.35 eV energyresolution.

XANES. X-ray spectroscopy and extended X-ray absorp-tion fine structure (EXAFS) measurements were performed atPNC-XOR bending magnet beamline (20-BM-B) of APS in Ar-gonne National Laboratory. Measurements at the V K-edgewere performed under transmission mode using gas ionizationchambers to monitor the incident and transmitted X-ray in-tensities. A third ionization chamber was used in conjunctionwith a Ti-foil standard to provide internal calibration for thealignment of the edge positions. The incident beam wasmonochromatized using a Si (111) double crystal fixed exitmonochromator. Harmonic rejection was accomplished using arhodium-coatedmirror. The charged samples were prepared bystripping V2O5 films onto Kapton tape. The discharged sampleswere scratched off and sealed inside a 3 mm diameter hole in apiece of aluminum foil by sealing the Kapton sheet to the foil

using epoxy. All cell assembly and disassembly operations wereperformed in a He-filled dry glovebox (oxygen level <2 ppm).The reference standards (V5þ and V4þ) were prepared byspreading thin, uniform layers of the V2O5 and VO2 powderon Kapton tape and stacking a few layers to attain the desiredabsorption step height. Each spectrum was normalized usingthe data processing software package IFEFFIT.37 The alignmentof each sample reference spectrum with respect to the Vstandard spectrum is within the range (0.03 eV.

MD Simulations. The atomistic model of V2O5 layers wasprepared in reference to crystallographic data.38 Bonding param-eters of the V2O5 layers were assigned based on the same data,while van der Waals parameters were adopted from otherstudies.39 Partial charges of the V2O5 layers were assigned totheir oxidation states. The TIP3Pmodel was used to describe thewater molecules. The CHARMM force field40 was used for theparameters of Mg ions and water molecules. Molecular dy-namics simulations were performed using NAMD41,42 with asimulation time step of 2 fs. The simulations were run under anisothermal�isobaric (NPT) ensemble at a temperature of T =300 K, using a Langevin damping constant of γLang = 0.01 ps�1.Periodic boundary conditions were applied to the system, andthe particle Ewarld summation algorithm43 was used to com-pute long-range electrostatics.

Conflict of Interest: The authors declare no competingfinancial interest.

Supporting Information Available: XRF imaging analysis forself-absorption of Mg in VO2/V2O5 is presented. SynchrotronX-ray diffraction of a bilayered vanadium cathode in thecharged and discharged state in conjunction with large-areaEDAX analysis is also shown. The Supporting Information isavailable free of charge on the ACS Publications website at DOI:10.1021/acsnano.5b02450.

Acknowledgment. This work was performed, in part, at theCenter for Nanoscale Materials, a U.S. Department of EnergyOffice of Science User Facility under Contract No. DE-AC02-06CH11357. Use of the Advanced Photon Source, an Office ofScience User Facility, was supported by the U.S. Department ofEnergy, Office of Science, Office of Basic Energy Sciences, undercontract no. DE-AC02-06CH11357. The work at the Joint Centerfor Energy Storage Research (JCESR), an Energy Innovation Hub,was funded by the Department of Energy, Office of Science,Basic Energy Sciences, is operated under Contract No. DE-AC02-06CH11357. The authors would like to thank Pietro P. Lopes forassistance with FTIR measurements. S.T. prepared all samplesand performed their characterization. D.Z. prepared the high-energy ball-milled Sn anode and performed long cycling of thefull Mg cells. Y.L. performed HRTEM as well as HAADF measure-ments. B.L. and J.M. performed XRF imaging measurements.H.C. and P.K. performed MD simulations, and V.S., N.M., C.J., andT.R. directed the research and performed data analysis.

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