NaV1.25Ti0.75O4: A Potential Post-Spinel Cathode Material for Mg

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NaV1.25Ti0.75O4: a potential post spinel cathode material for Mg batteries Xiaoqi Sun, Lauren Blanc, Gene M. Nolis, Patrick Bonnick, Jordi Cabana, and Linda F. Nazar Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.7b03383 • Publication Date (Web): 08 Nov 2017 Downloaded from http://pubs.acs.org on November 9, 2017

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Chemistry of Materials

NaV1.25Ti0.75O4: a potential post spinel cathode material for Mg batteries Xiaoqi Sun,a Lauren Blanc,a Gene M. Nolis,b Patrick Bonnick,a Jordi Cabana,b Linda F. Nazara,* a

Department of Chemistry, University of Waterloo, Canada

b

Department of Chemistry, University of Illinois at Chicago, USA

ABSTRACT Rechargeable Mg batteries are promising candidates for high energy density storage in theory, when a Mg metal anode is combined with an oxide cathode material. Despite the widely observed sluggish Mg2+ diffusion in most oxide lattices, recent first principles calculations predicted low diffusion barriers in the calcium ferrite (CF) type structures. In the present work, we experimentally examine the prospect of CF-type NaV1.25Ti0.75O4 as a Mg cathode. The Na+ ions, which lie in the ion migration pathway, need to be removed or exchanged with Mg2+ to allow Mg2+ de/intercalation. Partial desodiation was achieved through chemical and electrochemical methods, as proven by X-ray diffraction (XRD) and X-ray absorption spectroscopy (XAS), but deep desodiation was accompanied by partial amorphization of the material. Mg2+ ion exchange at moderate temperature (80 °C) resulted in the formation of Na0.19Mg0.41V1.25Ti0.75O4; however, phase transformation was observed when higher temperatures were applied to attempt complete ion exchange. Such phenomena point to the instability of the CF lattice when the tunnel is empty or occupied by a small ion (Mg2+). Thus while the low migration barrier predicted by computation is partly based on the relative metastability of the theoretical CF-MgxV1.25Ti0.75O4 lattice, the difficulty in stabilizing it also

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renders the material synthetically non-accessible, hindering this post-spinel’s application as an electrode material.


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INTRODUCTION Rechargeable Mg batteries show potential as high energy density storage systems. The Mg metal anode provides high volumetric capacity (3833 mAh mL-1) and low redox potential (2.37 V vs. S.H.E.), while long term application, good safety and low price are ensured by its limited dendritic growth during electrochemical deposition, its stability under ambient conditions and a high abundance on earth. However, the divalent charge on Mg2+ ions hinders its diffusion in solid state lattices, resulting in limited discovery of functional cathode materials to date. After the seminal Mo6S8 Chevrel phase reported by Aurbach et al. in 2000,1 it was not until last year that two titanium sulfides were identified as the second family of Mg insertion cathodes.2,3 These materials contain soft, polarizable anions in the lattices, reducing the interaction with the Mg2+ ion and allowing its facile diffusion. On the other hand, transition metal sulfides provide a relatively low redox voltage, which coupled with a high mass, results in low to moderate energy densities. One solution is to shift from sulfides to oxides. However, Mg2+ diffusion in oxide lattices is typically sluggish due to its strong ionic interaction with the more highly polarizing oxide ions.4,5 In addition, there is a strong driving force towards conversion reactions6,7 as well as a high energy penalty for desolvation at the liquid-solid interface.8 Finding a transition metal oxide that overcomes these barriers is very challenging, but at the same time, necessary to advance the field. First principles calculations predict that the calcium ferrite (CF) structure (Fig. 1 inset), which belongs to the post spinel structural family, allows relatively facile Mg2+ diffusion. Mg2+ migration barriers below 400 meV were calculated for the CF MgMn2O4 and NaV1.25Ti0.75O4 materials.9,10 However, the former, along with many other post spinel materials, can only be 3 ACS Paragon Plus Environment


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synthesized under high pressure (a few Gpa),11-13 thus limiting accessibility from both a research laboratory and practical application point of view. NaV1.25Ti0.75O4, on the other hand, can be obtained by conventional solid state synthesis at ambient pressure.14 Despite its relatively low predicted voltage of 1.5 V (vs. Mg),10 it offers an easily accessed platform to experimentally investigate Mg2+ diffusion in the CF structure. Herein, we examine the possibility of utilizing CF-type tunnel-structured NaV1.25Ti0.75O4 obtained through a solid state synthesis method - as a cathode material for Mg batteries. In order to induce Mg2+ de/intercalation, the Na+ cations in the tunnels were either extracted or exchanged with Mg2+ ion through chemical or electrochemical methods. Although the Na+ could be partially removed, our results showed a strong driving force towards structural degradation of the lattice upon full desodiation, demonstrating the instability of the CF phase when the tunnels are empty. Mg insertion in partially desodiated materials (either by chemical or electrochemical means) showed some (albeit limited) reversible capacity, suggesting Mg2+ mobility within the CF lattice. Thus, the approach to prepare metastable materials with low migration energy barriers shows potential in the search for Mg intercalation cathode candidates, but must be balanced with targeting end members that are just on the edge of thermodynamic stability. EXPERIMENTAL SECTION Material synthesis, chemical oxidation and Mg ion exchange reactions NaV1.25Ti0.75O4 was obtained from a small modification in the procedure of a previously reported synthesis.14 Typically, Na2CO3, V2O3 and TiO2 were ball milled in a 2.2:2.5:1.5 molar 4 ACS Paragon Plus Environment

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ratio, and the mixture was pelletized and heated at 600 °C for 42 hours under Ar flow. The resulting powder was re-ground and re-heated at 650 °C for 12 hours under Ar flow. Alternatively, the precursors were ground with a mortar and pestle followed by a heat treatment, and the product was then ball milled to reduce the particle size. This resulted in the same NaV1.25Ti0.75O4 product but with different morphology. This material was used for the chem-ox-2 reaction. The different morphology in the pristine material did not affect the overall results of chemical oxidation or ion exchange reactions. Chemical oxidation reactions were carried out in FeCl3/water or I2/acetonitrile (ACN). Typically, the chem-ox-1 sample was obtained by stirring NaV1.25Ti0.75O4 in 2 M FeCl3 in aqueous solution at room temperature for 6 days, with an intermediate exchange of fresh FeCl3 solution. The mixture was filtered and washed with water and ethanol. The chem-ox-2 sample was obtained by mixing NaV1.25Ti0.75O4 in 1 M I2 in acetonitrile (ACN) and sealing the mixture in a hydrothermal bomb, which was placed in an 80 °C oven for 9 days. The product was collected by filtration and washed with ACN. The chem-ox-1 reaction was carried out in ambient atmosphere while the chem-ox-2 reactions were performed in an Ar filled glovebox. The Mg ion exchange reaction was carried out by stirring NaV1.25Ti0.75O4 in 1 M MgSO4 in aqueous solution at 80 °C for 4 days. The product was filtered and washed with water and ethanol. High temperature reactions were carried out in a 1:2 molar ratio MgCl2:KCl eutectic mixture, the melting point of which is 430 °C.15 NaV1.25Ti0.75O4 was mixed with the eutectic in 1:14 and 1:2 weight ratios for low and high concentration reactions, respectively, pelletized, sealed in quartz tube under vacuum, heated at 500 °C for 1 hour and cooled down naturally in



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the furnace. The tube was then opened in air and water was added to transfer the powder to a funnel. The product was filtered and washed with water and ethanol. Materials characterization X-ray diffraction (XRD) analysis was carried out on a PANalytical Empyrean diffractometer with Cu-Kα radiation in Debye-Sherrer geometry. The sample was mixed with standard silicon (325 mesh, Sigma-Aldrich, 99%) and sealed in a 0.3 mm capillary to evaluate the crystalline fraction (Table S2). Rietveld refinement16 was performed using the FullProf suite17. Scale factor, zero point, background, lattice parameters, fractional coordinates, and thermal parameters were sequentially refined. Vanadium and titanium have very similar X-ray scattering factors and cannot be distinguished by XRD. Thus, the occupancies were fixed at values reported in the initial reference from neutron diffraction.14 Trying to refine Na occupancy and thermal parameter in the chemically oxidized and Mg ion exchanged materials resulted in unstable values, so the thermal parameter was fixed as the pristine NaV1.25Ti0.75O4 and the occupancy was refined. The morphologies and elemental ratios of the materials were studied with a Zeiss Ultra field emission scanning electron microscope (SEM) equipped with an energy dispersive X-ray spectroscopy (EDX) detector. X-ray photoelectron spectroscopy (XPS) was carried out on a Thermo VG Scientific ESCALab 250 instrument. Spectra were processed with GaussianLorentzian functions and a Shirley type back ground with CasaXPS software, and referenced to adventitious carbon at 285.0 eV. Electrochemical studies 6 ACS Paragon Plus Environment

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Cathodes were prepared by mixing the pristine material with Super P and polyvinylidene fluoride (PVDF, Sigma-Aldrich, average Mw ~ 534,000) at 8:1:1 weight ratio in N-methyl-2pyrrolidone (NMP, Sigma-Aldrich, 99.5%) and casting the slurry on stainless steel foil. An electrolyte comprised of 0.5 M Mg(CB11H12)2 in tetraglyme18 and a Mg metal anode were used for Mg cells, while an electrolyte comprised of 1 M NaClO4 in PC with 2 vol% FEC, and a Na metal anode were used for Na cells. Coin cells (2325) were assembled for cycling studies. Electrochemical studies were performed with a VMP3 potentiostat/galvanostat (Bio-logic), with cells held in a thermostatted oven at designated temperatures. In situ XRD The cathode slurry was cast on a glassy carbon electrode and assembled in a homemade in situ cell with 1 M NaClO4 in PC with 2 vol% FEC electrolyte and Na metal anode. The cell was cycled at 5 mA g-1 current density at room temperature, with XRD scans recorded operando. Each scan was half-hour long, corresponding to 2.5 mAh g-1 capacity or Δx = 0.0174 in NaxV1.25Ti0.75O4. Cell parameters were exacted by Le Bail fitting19 using the FullProf suite.17 X-ray absorption spectroscopy (XAS) X-ray absorption spectroscopy measurements were conducted at beamline 6.3.1.2, Advanced Light Source, Lawrence Berkeley National Laboratory. Spectra were collected simultaneously in both the total-electron-yield (TEY) and total-fluorescence yield (TFY) mode to make direct surface to bulk comparisons. Data were obtained at a spectral resolution of ~0.2 eV, with a 2 sec dwell time. At least 2 scans were performed on each sample, at each absorption edge, and scans were averaged in order to maximize the signal to noise ratio. V L2,3- and O K-



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edges were collected simultaneously at 505 – 600 eV, and Ti L2,3-edges were scanned at 445 480 eV. The current corresponding to I0 in electron yield was unstable at the energy of the Ti L2,3-edges. Since I0 did not change in the short timeframe between samples, the measured yields of the spectra were not divided by it. RESULTS AND DISCUSSION NaV1.25Ti0.75O4 was synthesized following the previously reported procedure14 with minor modifications (see Experimental Section for details). The material was at least 93 wt% crystalline as determined by comparison with a standard, and the crystal structure was confirmed by Rietveld16 refinement (Fig. 1 and Table S1a,2). SEM images show that the crystallites are about one micron in dimension (Fig. 1 inset). In order for the material to

Figure 1. Rietveld refinement fit of pristine NaV1.25Ti0.75O4 with 16 wt% standard Si addition (black crosses – experimental data, red lines – fitted data, blue line – difference map between observed and calculated data, green ticks – the Pmnb CF phase, magenta ticks – the Fd-3m Si phase; χ2 = 4.20, RBragg(CF) = 3.37, RBragg(Si) = 3.44. Insets show the crystal structure of the CF phase (yellow circles = Na, blue octahedra = V/TiO6) and SEM image of pristine NaV1.25Ti0.75O4. 8 ACS Paragon Plus Environment

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function as a Mg cathode, the Na+ ions in the structure must either be removed or exchanged with Mg2+. We examined Na+ extraction using chemical oxidation. Two different oxidizing agents, FeCl3/water and I2/acetonitrile (ACN) were employed, providing products that were labeled as chem-ox-1 and chem-ox-2, respectively. The XRD patterns of both materials show a clear shift in d-spacing and intensity change compared to the pristine material (Fig. 2a), with chem-ox-2 showing a larger difference that is indicative of a higher degree of Na+ extraction. The most significant decrease is in the b lattice parameter, as summarized in Table 1. Rietveld refinement reveals a decrease of

Table 1. Summary of Rietveld refinement and chemical composition results of the desodiated and magnesium ion exchanged materials. Lattice parameters (Å) Sample

Xtal wt%

Overall

(Rietveld Si

x(Na)y(Mg)

std. addt’n)

/V1.25Ti0.75O4

Rietveld

(S.G. = Pmnb, #62)

x(Na) a

b

c

Pristine

2.947

9.167

10.670

1 (fixed)

94(5)%

0.87Na (EDX)

Chem-ox-1

2.937

8.954

10.649

0.36

99(5)%

0.24Na (EDX)

Chem-ox-2

2.939

8.883

10.643

0.25

64(4)%

0.14Na (EDX)

Echem-ox-RT

2.936

9.058

10.634

―

―

0.44Na (Echem)

2.933

8.996

10.649

―

―

Echem-ox-

0.02Na (Echem)

90°C Mg-ex

0.44Na (EDX) 2.932

9.070

10.661

0.19 (Na)

92(5)%


0.39Na, 0.22Mg


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0.41 (Mg)

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(EDX)

Na occupancy in the chemically oxidized materials (Fig. 2b,c and Table S1b,c), demonstrating that Na+ was indeed removed from the structure. The lowest Na+ content of 0.25 was achieved in chem-ox-2; i.e., a formula of Na0.25V1.25Ti0.75O4, whereas chem-ox-1 yielded Na0.36V1.25Ti0.75O4. The change in the Na content is in accord with EDX results (Table 1).


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Figure 2. (a) XRD patterns in a selected range to show shifts in d-spacing, (b)(c) Rietveld refinement fits with standard Si (34 wt% for chem-ox-1 and 17 wt% for chem-ox-2) addition, and (d) SEM images of the two chemically oxidized samples (black crosses – experimental data, red lines – fitted data, blue line – difference map between observed and calculated data, green ticks – the Pmnb CF phase, magenta ticks – the Fd-3m Si phase. The different particle size in chem-ox-2 is a result of slightly different synthesis conditions being used to generate the pristine materials). 11 ACS Paragon Plus Environment


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While Na can be chemically extracted from NaV1.25Ti0.75O4, upon deep desodiation, simultaneous amorphization occurs. The crystalline weight percentages for chem-ox-1 and chem-ox-2 were 99% and 64%, respectively, suggesting a crystalline limitation of ~ Na0.25V1.25Ti0.75O4 for the CF phase, where amorphization is initiated by further desodiation (Table S2). We previously observed a similar amorphization for Mg2Mo3O8 upon Mg2+ extraction.20 We believe that an inhomogeneous Na+ extraction process, as indicated by the broadening of XRD peaks after chemical desodiation, is responsible for the partial amorphization. The portion with low Na+ content has poor structural stability and collapses. The amorphous phase in chem-ox-2 is evident in the SEM images, appearing as small particles attached to the surface of the material (Fig. 2d).


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Electrochemical desodiation of NaV1.25Ti0.75O4 was also examined in a Na cell, while in situ XRD was carried out to examine the evolution in the lattice parameters. Fig. 3a-c shows the electrochemical profiles on charge together with the lattice parameters extracted by a Le Bail fit.19 Similar to chemical oxidation, electrochemically charging NaV1.25Ti0.75O4 also results in significant decrease of the b lattice parameter, while minor changes are observed for the a and



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Figure 3. (a)-(c) Lattice parameter evolution of NaV1.25Ti0.75O4 during first charge examined in an in-situ cell with a Na anode using an electrolyte comprised of 1M NaClO4 in PC with 2 vol% FEC. The electrochemical charge profiles of the corresponding cells (run at a current density of 5 mA g-1 at room temperature) are also depicted. (d) Comparison of first charge profiles from cells run at room temperature and 90 °C, both at 5 mA g-1. c parameters (Table 1). The final composition, Na0.44V1.25Ti0.75O4, shows a higher Na content than obtained from the chemical method (i.e., 0.25 in chem-ox-2). The difficulty in removing the last half of the Na+ is in good agreement with first principles calculations, which suggest a Na+ migration barrier that is about 1.5 times higher in an empty lattice compared to a full lattice (600 meV vs. 400 meV).10 Even at an elevated temperature of 90 °C, the Na+ content at 14 ACS Paragon Plus Environment

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the end of charge is about 0.44 as revealed by EDX. Electrolyte decomposition can take place below 4V at such temperatures, adding slightly to the total capacity (Fig. 3d). The standard addition method to evaluate the weight percent of a crystalline phase in a cathode material using ex situ XRD data would result in large error due to the low cathode mass (a few milligrams), while reliable Rietveld refinement is not possible with in situ laboratory-XRD data due to the thinness of the electrode. Nevertheless, a rough estimation of any change in crystallinity on charging can still be made from in situ XRD peak areas. The total peak area of the charged material is about 93% of the pristine, showing that only minor amorphization, if any, of the electrochemically charged material takes place. This suggests that amorphization either only starts to take place at lower Na content, and/or is primarily induced by the more forcing potential provided by chemical oxidation (I2 at elevated temperature), which is generally a less gentle oxidation method. X-ray Absorption Spectroscopy (XAS) of the pristine and desodiated samples measured for various materials at the V L3- Ti L2, 3- and O K-edges, with total electron and fluorescence yields, are shown in Fig. 4 and Fig. S1, respectively. Electron yield signals correspond to the chemical state of the surface layer of the material, while fluorescence yield signals correspond to bulk properties. The trends discussed hereafter were very similar in both modes, indicating that surface and bulk undergo similar changes. The complete V L2, 3 spectra can be found in Fig. S2. The features correspond to transitions from transition metal 2p3/2 (L3) and 2p1/2 (L2) states directly to the unoccupied 3d states,21 making them very sensitive to redox changes and local transition metal coordination. The changes at the L3 edge, shown in Fig. 4a, are representative of both edges. The spectrum of pristine NaV1.25Ti0.75O4 (#4, red) presents an intense feature at 15 ACS Paragon Plus Environment


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Figure 4. X-ray near edge absorption spectroscopy (XANES) normalized fluorescence yields for the materials examined in this study: (a) V L and O K edges between 513 – 522 eV; (b) Ti L edge between 455 – 470 eV; and (c) V L and O K edges between 525 – 550 eV. XANES data for standards correspond to: (1) VO2; (2) V2O5; (3) TiO2 anatase; (4) pristine NaV1.25Ti0.75O4. XANES data for NaV1.25Ti0.75O4 subjected to various treatments are shown as follows: ((5) Chem-ox-1, (6) Chem-ox-2)), (7) electrochemically charged and (8) chemically magnesiated. Full spectra are shown in the Supporting Information. 518.3, with a shoulder at 517.1 eV, and a tail to lower energies. The spectrum is located at significantly lower energies than VO2 (#1, black) and V2O5 (#2, grey), which show intense features around 519 eV, in agreement with literature reports.22 The slightly lower position of the center of gravity of VO2 compared to V2O5 is consistent with the lower formal oxidation state. The spectrum of NaV1.25Ti0.75O4 is close to V2O3 measured by Abbate,22 suggesting that V is in a majority (III) formal state in the compound as synthesized. This is in accord with previous magnetic susceptibility measurements of NaV1.25Ti0.75O4 that reported an average oxidation state of V+3.2, assuming a Ti(IV) valence.14 Our Ti 2p XPS data of the pristine material confirms the latter state (Fig. S3). The Ti L2, 3-edge spectra of all NaV1.25Ti0.75O4 samples (Fig. 4b) were consistent with a formal Ti(IV) state, indicating that redox takes place solely at the vanadium center. Slight variations in the symmetry of the coordination shell affect the multiplet splitting


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of the final 3d states.23-25 The L3 peaks, between 455 and 462 eV, are most sensitive to these changes in symmetry.23 The subtle differences between NaV1.25Ti0.75O4 (#4, red) and a standard spectrum for TiO2 (#3, olive) can therefore be ascribed to the specific geometric distortions of the octahedral environment of Ti(IV). All O K-edge features originate from O 1s -> O 2p transitions, where the states corresponding to O 2p orbitals hybridized with V and Ti 3d and 4s, p orbitals, appear roughly below and above 535 eV, respectively (Fig. 4c).22 The pre-edge features below this energy are again the most sensitive to redox changes. The O pre-edge of pristine NaV1.25Ti0.75O4 is composed of a broad, multi-component signal starting at 529.6 eV and ending with a maximum at 531.9 eV. The shape, energy and distribution of intensity were all different from VO2, V2O5 and TiO2, which show characteristic t2g and eg doublets around 530 and 532 eV,22,25,26 respectively, with the t2g absorption event being the most intense of the two. The broad preedge with higher intensity at lower energy in NaV1.25Ti0.75O4 again compares well with V2O3,22 consistent with a dominant V(III) formal state. Chemical oxidation induced changes in all spectra. In chem-ox-1 (#6, green), the most intense V L3 feature remains centered at 518.3 eV (Fig. 4a), but the center of gravity is shifted to higher energies driven by the disappearance of the shoulder at 517.1 eV. This subtle change would be consistent with the slight oxidation of V. In the case of chem-ox-2 (#5, wine), a significant shift of the most intense V L3 feature to 518.7 eV was observed. The resulting overall lineshape resembled the spectrum of VO2 (#1, black), indicating that chem-ox-2 was significantly more oxidized than chem-ox-1. This observation is consistent with the XRD analysis, and suggests that the amorphous fraction resulting from the chemical treatments is 17 ACS Paragon Plus Environment


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significantly more oxidized than the crystalline oxide. The oxidation of the compound is accompanied by a subtle change in position of the Ti L2, 3-edge spectra (e.g., around 460 and 465 eV, Fig. 4b), indicative of changes in the local coordination of Ti(IV)23 due to the structural changes described by XRD. The formation of V(IV) in chem-ox-2 changed the hybridization with O, resulting in an increase of the t2g feature at 530 eV in the pre-edge doublet (Fig. 4c). The resulting O K-edge was closer to VO2. In the case of chem-ox-1, the pre-edge doublet was washed out, indicating a variety of V-O interactions at slightly different geometries. Its intensity increased with respect to the O 2p-V 4sp signals at the edge compared to the pristine state. Since XAS probes unoccupied electronic states, the increase in intensity is consistent with the decrease in electron occupancy of the O 2p-V 3d expected from oxidation. The changes at the V L3 and Ti L- were small in Echem-ox-90°C compared to the pristine state. However, the O K-edge spectrum was comparable to Chem-ox-1, suggesting some oxidation of the compound. These trends agree with the XRD and EDX data. The electrochemical behavior of both chemically and electrochemically desodiated materials was evaluated in Mg cells, with a Mg monocarborane electrolyte (Mg(CB11H12)2 in tetraglyme)18 and Mg metal anode. At 60 °C, the chemically desodiated chem-ox-2 material (Na0.25V1.25Ti0.75O4) shows an initial sloping discharge curve at an average voltage of 1.5 V (Fig. 5a), in good agreement with calculation.10 This trend continues up to a capacity of 50 mAh g-1, followed by a region of steeper voltage profile. A capacity of 80 mAh g-1 is reached at the end of discharge. The following charge only retrieves half of the initial capacity and a large overpotential is observed, showing moderate reversibility and kinetics of the process. The reversible capacity is retained on subsequent cycles, however. EDX revealed 0.12(2) Mg2+ is 18 ACS Paragon Plus Environment

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present in the chem-ox-2 cathode at the end of the 5th discharge and demonstrates Mg intercalation takes place. The electrochemically desodiated material (Na0.4V1.25Ti0.75O4) exhibits poorer behavior upon Mg2+ intercalation; about two-fold less capacity is reversible (Fig. 5b). This is may be due to its greater Na-ion content which could affect Mg ion migration pathways. The relatively poor electrochemistry of the chemically desodiated material (chem ox-2) could result from the amorphous phase that is likely electrochemically inactive, and resides on the surface. The electrochemically desodiated material, on the other hand, is still active as demonstrated by its low degree of amorphization (see discussion of Fig. 3) as well as by its subsequent successful reversible cycling of sodium ions in a Na cell (Fig. S4). Hence we believe that the difficulty in inserting Mg2+ is due to the strong driving force towards structural degradation (in accord with calculations that predict instability of the Mg-inserted lattice) which is not predicted to occur in the case of Na+ insertion.10



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Figure 5. Electrochemistry of (a) chem-ox-2 and (b) electrochemically desodiated material (i.e., Na0.4V1.25Ti0.75O4) with a Mg anode and 0.5 M Mg(CB11H12)2 in tetraglyme electrolyte at 5 mA g-1 current density and 60 °C. We also explored altering the strategy to directly exchange the Na+ ions with Mg2+. An


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ion exchange reaction was carried out by stirring the Na1.0V1.25Ti0.75O4 material in aqueous MgSO4 at 80 °C. The XAS spectra of the Mg-ex sample were essentially the same as in pristine NaV1.25Ti0.75O4, except for a shift of the peak at 460.4 eV to slightly lower energies (Fig. 4). These similarities denote the preservation of the redox states in the compound upon Na+/Mg2+ exchange, but hint at a small distortion of the local Ti environment due to the exchange. The XRD pattern shows a peak shift in the ion exchanged material (labeled as Mg-ex), while the



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intensity change is very small as expected (Fig. 6a). X-rays cannot distinguish between Na and Mg due to their extremely similar scattering factors. Therefore, after initially refining the occupancy on the sodium site, we ensured charge balance was maintained in the final composition by assuming no oxidation state change on the transition metals arose during the ion exchange process (based on the XAS results); namely that the charge on that site (a combination of Na+ and Mg2+) was still equal to one. The refined Na and Mg occupancies (Table

Figure 6. (a) XRD of the Mg ion exchanged sample, showing a shift of XRD peaks. (b) Rietveld refinement fit the Mg ion exchanged sample with 17 wt% standard Si addition (Black crosses – experimental data, red lines – fitted data, blue line – difference map between observed and calculated data, green ticks – the Pmnb CF phase, magenta ticks – the Fd-3m Si phase, orange ticks – the R-3c V2O3 impurity phase which also presents in the pristine material used in this ion exchange reaction due to slightly different synthesis conditions). Inset shows the SEM image.


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1, Fig. 6b and Table S1d) yield a composition of Na0.19Mg0.41V1.25Ti0.75O4 that is fully crystalline (see Table S1d for refinement details). A second ion-exchange process did not result in any change in either the Na/Mg composition or the crystallinity. This demonstrates that while some Mg2+ mobility indeed exists in the CF lattice, not all of the Na+ ions can be exchanged with Mg2+.



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Alternatively, we carried out Mg ion exchange with a MgCl2/KCl-1:2 eutectic which melts at 420 °C. Higher temperatures enhance ion diffusion in solid state lattices and would promote a deeper degree of ion exchange. Two different ratios of NaV1.25Ti0.75O4: eutectic were examined at 500 °C (see Experimental Section for details), but both resulted in complete structural transformation. At a low CF ratio, MgM2O5 (M = V, Ti) is obtained as determined by XRD measurements (Fig. 7a), which exhibits an XRD pattern identical to MgTi2O5 (ICSD


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#157236). This phase exhibits a similar tunnel structure to that of the CF phase, but the size of the tunnel is smaller owing to a different connectivity of the octahedral motifs (Fig. 7a inset). This phase is presumably more thermodynamically stable at the higher temperatures employed in the reaction owing to the smaller size of the Mg2+ cation compared to Na+; hence it is stabilized with shorter bonds. At a high CF ratio, a layered MgMO3 (M = V, Ti) structure is formed (Fig. 7b), which is the same as MgTiO3 (ICSD #164766). We note that for both new



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Figure 7. XRD and SEM images of (a)(c) low concentration and (b)(d) high concentration CF in eutectic products, showing MgTi2O5 (red phase marks) and MgTiO3 (orange phase markers) structures, respectively. phases, EDX shows that the same V:Ti ratio is preserved as in the starting NaV1.25Ti0.75O4 phase. As both phases require a complete rearrangement of the CF structure, a dissolutionreprecipitation process must be involved in the reactions in the eutectic melt. This is in agreement with the new morphologies of the products shown in the SEM images (Fig. 7c,d), that are either square platelets (MgTi2O5) or sub-micron crystallites (MgTiO3).


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Overall, our experiments show that partial Na+ removal or Mg2+ exchange in the CFNaV1.25Ti0.75O4 material can be achieved under certain conditions. The largest change in the unit cell of the structure in both cases is observed in the b lattice parameter, which is perpendicular to the 1D ion tunnel direction, while minor changes are evident in the a and c parameters. Concomitant partial amorphization/phase transformation takes place, due in part to the inhomogeneity of Na+/Mg2+ ion distribution throughout the sample. Compositions with a lower Na+ content and/or more Mg2+ are unfortunately prone to framework collapse driven by thermodynamics. A similar trend towards amorphization/conversion was observed previously in other oxide-based Mg cathodes,

6,8,20,27-29

which strongly limits their electrochemical

behavior.7 Such structural instability of the CF phase is in accord with first principles calculation results published after the start of our experimental study. According to Ceder et al., the Ehull of the empty lattice approaches 80 meV/atom, while Mg2+ insertion results in even higher Ehull values.10 Such values demonstrate a strong thermodynamic driving force towards structural decomposition to more stable phases, resulting in the amorphization or transformation into the MgTi2O5 and MgTiO3 structures that we observed experimentally. The structural instability, on the other hand, lowers the energy difference between the starting and transition point of ion hopping and promotes a lower barrier. In fact, we observed some Mg2+ mobility in the CF structure, based either on the – albeit limited - electrochemical performance of the Mg cells or the fact that we were able to achieve partial ion exchange of Na+ for Mg2+. Nevertheless, stabilizing a structure over a wide compositional range is required for cathode materials, which was not successful for the CF phase studied. 27 ACS Paragon Plus Environment


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CONCLUSIONS In summary, we carried out experimental investigations of calcium-ferrite (CF) type NaV1.25Ti0.25O4 as a potential cathode material for Mg batteries. Upon deep Na+ removal amorphization is observed, demonstrating the instability of the empty lattice, and very limited Mg intercalation was observed into the electrochemically desodiated lattice. Such results are in agreement with the high Ehull values suggested by first principles calculation. The structural instability of end members helps lower the ion migration barrier; however, it is essential that 1) the end members can be kinetically stabilized, and 2) not undergo a high penalty for Mg2+ desolvation at the interface. Finding this “sweet spot” is challenging, but can be potentially achieved by searching for metastable phases that provide a low energy barrier for ion diffusion with structures selected for optimal stability of the end members. AUTHOR CONTRIBUTIONS X.S. and L.F.N conceived the experiments and study; X.S. carried out the materials synthesis, electrochemical studies, and materials characterization/analysis with the help of L.B.; G.N. and J.C. performed the XAS study and contributed its discussion to the text; P.B. participated in the dialogue on divalent cation intercalation; X.S. and L.F.N. wrote the manuscript. Supporting Information Available Refined parameters and crystalline phase weight percentages for the pristine, desodiated, Mg ion exchanged NaV1.25Ti0.75O4 samples; XAS of the pristine, desodiated and Mg ion exchanged samples measured for various materials in the full energy range at the V L3- Ti L2, 3- and O K28 ACS Paragon Plus Environment

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edges with total electron and fluorescence yields; complete V L2, 3 spectra; Ti 2p XPS spectrum; charge/discharge profiles for Na-ion insertion in NaV1.25Ti0.75O4 in a Na cell. ACKNOWLEDGEMENTS This work was supported by the Joint Center for Energy Storage Research (JCESR), an Energy Innovation Hub funded by the US Department of Energy (DOE), Office of Science, Basic Energy Sciences. NSERC is acknowledged by LFN for a Canada Research Chair. This research used resources at the Advanced Light Source, Lawrence Berkeley National Laboratory, a User Facility funded by the Office of Science of the Department of Energy under contract no. DEAC02-05CH11231. Also, we kindly thank Drs. Yi-sheng Liu and Jinghua Guo for providing the XAS spectra for the VO2, V2O5 and TiO2 standards, and Drs. Ka-Cheong Lau and Chen Liao for providing the Mg(CB11H12)/tetraglyme electrolyte.



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Graphic for Table of Contents


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NaV1.25Ti0.75O4: A Potential Post-Spinel Cathode Material for Mg

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