Synthesis and Electrochemical Performance of the Orthorhombic

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Synthesis and electrochemical performance of the orthorhombic Li2Fe(SO4)2 polymorph for Li-ion batteries Laura Lander, Marine Reynaud, Gwenaelle Rousse, Moulay Tahar Sougrati, Christel Laberty-Robert, Robert J. Messinger, Michael Deschamps, and Jean-Marie Tarascon Chem. Mater., Just Accepted Manuscript • Publication Date (Web): 10 Jun 2014 Downloaded from http://pubs.acs.org on July 2, 2014

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Chemistry of Materials

Synthesis and electrochemical performance of the orthorhombic Li2Fe(SO4)2 polymorph for Li-ion batteries Laura Landera,b,c, Marine Reynauda,d, Gwenaëlle Rousseb,c, Moulay T. Sougratie, Christel Laberty-Robertf, Robert J. Messingerg, Michaël Deschampsg, and Jean-Marie Tarasconb,*

a

Laboratoire de Réactivité et Chimie des Solides (LRCS), Université de Picardie Jules Verne (UPJV),

CNRS UMR 7314, 33 rue Saint Leu, 80039 Amiens cedex, France b

FRE 3677 Chimie du Solide et de l’Energie, Collège de France, 11 place Marcelin Berthelot, 75231

Paris Cedex 05, France and RS2E, CNRS FR 3459, France c

Institut de Minéralogie, de Physique des Matériaux, et de Cosmochimie (IMPMC), Sorbonne

Universités - UPMC Univ Paris 06, UMR CNRS 7590, Muséum National d’Histoire Naturelle, IRD UMR 206, 4 Place Jussieu, F-75005 Paris, France. d

CIC Energigune, Albert Einstein 48, 01510 Miñano (Álava), Spain

e

Institut Charles Gerhardt – Laboratoire des Agrégats, Interfaces et Matériaux pour l’Energie, Université

de Montpellier II, CNRS UMR 5253, 34095 Montpellier Cedex 5, France f

Laboratoire de Chimie de la Matière Condensée de Paris, UPMC-UMR7574, Collège de France, 11

Place Marcelin Berthelot, 75005 Paris Cedex 05, France g

CNRS, CEMHTI, UPR 3079, Univ. Orléans, F-45071 Orléans, France

* Corresponding author: [email protected].

MANUSCRIPT

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Abstract To enhance the safety, cost, and energy density of Li-ion batteries, significant research efforts have been devoted to the search for new positive electrode materials that exhibit high redox potentials and are composed of low-cost, earth abundant elements. Sulfate chemistry has yielded promising results for ironbased polyanionic electrode materials using the FeIII+/FeII+ redox couple, including the recent discovery of a monoclinic marinite Li2Fe(SO4)2 phase (3.83 V vs. Li+/Li0). Here, we report the ball-milling synthesis and electrochemical properties of a new orthorhombic polymorph of Li2Fe(SO4)2, which rapidly and reversibly reacts with lithium through a two-step redox process (3.73 V and 3.85 V vs. Li+/Li0) with an overall sustained capacity of about 90 mAh/g. Using similar synthesis conditions, the cobalt-, zinc-, magnesium- and nickel-based orthorhombic analogues were also obtained, though no electrochemical activity was observed for these phases. Overall, our results demonstrate that polymorphism can play a crucial role in the search for new battery electrode materials and emphasize the need to understand and master synthetic control.


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Introduction Lithium-ion batteries are among the most successful rechargeable battery technologies because of their superior energy densities and cycle life times. They have dominated the portable electronics market, are strong candidates to power electric vehicles, and are serious contenders to other promising grid storage technologies (e.g. Na/S batteries). Both of these latter applications require the development of new sustainable, low-cost, and safe positive electrode materials with improved electrochemical properties. The importance of these criteria are reflected in the great interest that iron-based polyanionic compounds have garnered, among which LiFePO41 is currently the most praised electrode material for large volume applications, despite its relatively low operating voltage (3.45 V vs. Li+/Li0). Over the last 15 years, there has been an intensive search for new iron-based polyanionic compounds demonstrating enhanced FeIII+/FeII+ redox voltages.2 Such efforts have resulted in the discovery of new polyanionic compounds such as borates (LiFeBO3)3–5 and silicates (Li2FeSiO4)6–9, which show larger capacities but display lower redox voltages of ~2.8 V vs. Li+/Li0. Knowing that the redox potential of an intercalation compound correlates with the iono-covalency of the metal–anion bonding, we have recently reported the preparation of a series of fluorosulfates compounds LiMSO4F (M = Fe, Mn, Co, Ni), which combine the electronegativity of the fluorine anion (F–) and the inductive effect of the sulfate polyanion (SO42–).10–14 We have shown that, depending upon the synthesis conditions, LiFeSO4F can either crystallize into a tavorite or a triplite-type structure, with redox potentials of 3.6 V or 3.9 V vs. Li+/Li0, respectively, with the latter being the highest voltage thus far reported for the FeIII+/FeII+ redox couple in an inorganic compound. Such a high redox-voltage (3.9 V vs. Li+/Li0) was recently observed in LiFe1-xMnxP2O7 pyrophosphates, but was responsible for only 30 % of the overall capacity, while the remaining capacity derived from a 3.7 V redox couple.15,16 Searching further for new battery materials that can circumvent the safety aspects associated with the use of fluorine while preserving a high-voltage redox potential, our group recently isolated a new fluorine-free lithium iron sulfate Li2Fe(SO4)2.17 With a potential of 3.83 V vs. Li+/Li0, this new monoclinic marinite phase Li2Fe(SO4)2 displays the highest potential ever observed for the FeIII+/FeII+ redox couple in a fluorine-free iron-based compound, rivaled only by the triplite form of LiFeSO4F (3.9 V vs. Li+/Li0). Other analogous Li2M(SO4)2 phases (M = Co, Mn, Zn, Mg and Ni) were prepared using similar solid-state routes with the cobalt-, manganese-, and magnesium-based phases being isostructural to the iron-based phase,17,18 while the Ni-based material crystallizes in an orthorhombic structure.19,20 In contrast, polymorphism was discovered for the Zn-based phase, as it could be stabilized into either the orthorhombic or the monoclinic marinite structures, depending on the annealing temperature.21,22 Recalling that such polymorphism was previously encountered with other sulfate-based


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iron compounds (e.g., tavorite/triplite LiFeSO4F10–14, tavorite/layered LiFeSO4OH23,24, the α and β polymorphs of Na2Fe(SO4)222,25–28), we hypothesized that Li2Fe(SO4)2 could be stabilized into different crystal structures as well. To answer this question, we revisited the preparation of Li2Fe(SO4)2 by exploring various synthetic approaches (ionothermal, ceramic, and ball-milling). With regards to polymorphism, ball-milling is known to favor the growth of the denser polymorph, as demonstrated by the formation of triplite LiFeSO4F, which has a smaller unit cell volume than its tavorite counterpart.13,14 Knowing that the monoclinic marinite Li2Zn(SO4)2 presents a lower density than its orthorhombic equivalent was an impetus to implement our ball-milling approach to the synthesis of a denser orthorhombic Li2Fe(SO4)2. We demonstrate herein the feasibility of stabilizing Li2Fe(SO4)2 in the orthorhombic structure by using mechanical milling, as opposed to a ceramic route, and demonstrate its improved electrochemical properties compared to the monoclinic marinite polymorph. Synthesis approach The first hint for the existence of polymorphism in the Li2M(SO4)2 series of compounds (M = Mn, Fe, Co, Zn, Mg, Ni) was unveiled recently during our careful investigations of the lithium-extraction/insertion mechanism into the monoclinic marinite phase of Li2Fe(SO4)2, obtained from solid-state route.17,22 The electrochemical curve of this electrode material displayed a main plateau at 3.83 V vs. Li+/Li0, while a sloping contribution at around 3.6-3.7 V typically appeared at the early stage of the oxidation process.17 Further investigations demonstrated that the magnitude of this sloping contribution increased with the electrode preparation ball-milling time (Figure 1a).22 Whereas such a feature is usually ascribed to the presence of secondary phases, no changes were observed by X-ray diffraction (XRD) patterns of the electrode material, neither after its preparation, nor during the first stage of the oxidation process. This observation led us to hypothesize that an amorphous phase had formed during the electrode preparation using high-energetic mechanical milling. To search for the existence of amorphous domains, we characterized samples prepared with different ball-milling times with

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Fe Mössbauer spectroscopy

(Figure 1b), and solid-state 7Li nuclear magnetic resonance (NMR) measurements (Figure 1c) acquired at 7.05 T and 62.5 kHz magic-angle-spinning (see Supporting Information for experimental details). Both Mössbauer and solid-state 7Li NMR measurements established the presence of additional iron and lithium environments in the material, respectively, where the relative populations of these environments increased with longer ball-milling times. Combining quantitative analyses of both techniques led us to the conclusion that an additional amorphous phase had formed with the same chemical composition as the primary monoclinic marinite phase, Li2Fe(SO4)2.22 Thus, we decided to pursue a ball-milling synthesis of this amorphous phase, with the goal that longer milling times would promote its crystallization.


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To this end, iron sulfate FeSO4 and lithium sulfate Li2SO4 precursors, used in stoichiometric ratios to obtain samples with the nominal composition “Li2Fe(SO4)2”, were ball-milled under argon with a planetary miller for times ranging from one to ten hours, and the formation of the targeted phase was followed by XRD, Mössbauer, and solid-state NMR measurements. We noticed the appearance of weak diffraction peaks after 30 minutes of mechanical milling, which increased in intensity with longer milling times. After 5 to 10 hours of milling, we obtained a well-crystallized sample that contained no traces of the sulfate precursors (Figure 2a).

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Fe Mössbauer and solid-state 7Li NMR measurements performed on

the sample obtained after 10 hours of mechanical milling (Figure 2b and c) confirmed that we synthesized the crystalline analogue of the amorphous phase that was observed when ball-milling the marinite monoclinic phase. We show below that the symmetry of this new crystalline phase is orthorhombic Li2Fe(SO4)2. The 57Fe Mössbauer spectrum of this phase (Figure 2b) shows two types of doublets associated with FeII+ (Table SI-1), where the intense component (~67 % of the total iron content) that exhibits the largest quadrupole doublet and narrowest line width is attributed to the crystallized Li2Fe(SO4)2 phase, whereas the second component is ascribed to a less crystallized region of the sample (~33 % of the total iron content). This latter contribution can be decreased by further annealing the samples under vacuum at 200°C after ball milling, which enhances the crystallinity of the material. The solid-state 7Li NMR spectrum of the orthorhombic Li2Fe(SO4)2 phase (Figure 2c), and its comparison to otherwise identical NMR spectra acquired on the monoclinic Li2Fe(SO4)2 subjected to different ballmilling times (Figure 1c), yield insight into the local lithium environments characteristic of orthorhombicand monoclinic-type Li2Fe(SO4)2 structures. The solid-state 7Li NMR spectrum of the orthorhombic Li2Fe(SO4)2 polymorph (Figure 2c, top) exhibits a broad spinning-side-band manifold, consistent with strong electron-nuclear interactions associated with unpaired electrons from the paramagnetic iron atoms. The experimental NMR spectrum (Figure 2c, black lines) is modeled with excellent agreement (Figure 2c, red lines) using only a single, broad 7Li signal with an isotropic shift of -12 ppm (Figure 2c, bottom) and one set of 7Li paramagnetic shift anisotropy (PSA) parameters (Table SI-2) for the entire spinning-side-band manifold (asterisks in Figure 2c, top). The lithium atoms experience broad distributions of local environments throughout the material (full-width-half-maximum = 77 ppm), consistent with the large difference between the coherence lifetime (T2’ = 970 s) and the characteristic decay time of the free-induction decay (T2* = 36 s). Interestingly, the additional amorphous Li2Fe(SO4)2 component within the monoclinic marinite Li2Fe(SO4)2 subjected to increased ball-milling times (Figure 1c, green lines) yields the same isotropic 7Li shifts and PSA parameters as those used for the orthorhombic Li2Fe(SO4)2 (Figure 2c, red lines). These isotropic 7Li shifts and PSA parameters are


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furthermore distinct from those characteristic of the monoclinic marinite Li2Fe(SO4)2 (Figure 1c, blue lines; Table SI-2). Thus, the isotropic 7Li NMR shifts and PSA parameters are a signature of the characteristic local environments around the lithium atoms, and while both ball-milling approaches appear to drive the formation of local orthorhombic-type Li2Fe(SO4)2 structures, only the longer planetary milling times promote its crystallization. Consequently, we employed a similar protocol to prepare powders of the analogous cobalt-, nickel-, zincand magnesium-based Li2M(SO4)2 phases, but synthesized them under air and using different ball-milling times depending on the transition metal. The targeted phases were obtained after one hour of ball-milling for Co (divided into 30-min steps with 15-min pauses), 5 h for Ni, and 7 h for Zn. Occasionally, and when specified, the resulting ball-milled powders were pressed into a pellet and annealed in normal atmosphere (M = Co, Ni, Zn) at around 200°C to improve the crystallinity of the samples. For the case of magnesium, 18 hours of ball-milling resulted in a mixture of the targeted orthorhombic compound and the previously reported Li2Mg2(SO4)3 phase29. Curiously, all our attempts to stabilize a manganese-based orthorhombic analogue through mechanical-milling remained unsuccessful, but systematically lead to the formation of the monoclinic marinite polymorph. To further understand this polymorphism, we returned to the marinite polymorph and successfully prepared the various Li2M(SO4)2 phases (M = Fe, Co, Mn, Mg, Zn, Ni) via a ceramic process that involved two key steps: (i) stoichiometric ratios of the sulfate precursors Li2SO4 and MSO4 were ballmilled for twenty minutes using a Spex 8000 vibratory miller, and (ii) the resulting mixtures were pressed into pellets and annealed at different temperatures and for different times, depending on the nature of the metal. Given the propensity of the FeII+ species to oxidize, the second step was conducted in a quartz tube sealed under vacuum at 320 °C for at least 48 hours. The monoclinic marinite cobalt and manganese phases were obtained more easily through a heat treatment of 320 to 400 ºC under air for one night, while the isostructural Li2Mg(SO4)2 was stabilized after a longer annealing of four weeks. Likewise, the orthorhombic nickel-based compound formed after one night at 500 ºC under air. Finally, Li2Zn(SO4)2 crystallized either into the orthorhombic structure if annealed for one night under air at 400 to 450 ºC, or into the monoclinic marinite phase if the temperature was raised to 480 ºC. Energy Dispersive X-ray spectroscopy (EDX) of the orthorhombic and monoclinic Li2Co(SO4)2 phases confirmed the atomic ratio of Co:S as being 1:2, consistent with the expected stoichiometry. No other elemental impurities were detected in the different samples. Furthermore, characterization by transmission electron microscopy (TEM) of the orthorhombic and monoclinic Li2Co(SO4)2 phases showed the formation of particle aggregates with crystallite sizes in the sub-micrometer range (Figure SI-1).


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Structural characterization The XRD patterns of the as-prepared Li2M(SO4)2 phases (M = Fe, Co, Zn, Mg and Ni) phases were recorded using either a Bruker D8 diffractometer, equipped with a copper source (λCu-Kα1 = 1.54056 Å, λCu-Kα2 = 1.54439 Å) and a Vantec detector, or a Panalytical X'Pert Pro MPD diffractometer supplied with a cobalt source (λCo-Kα1 = 1.78897 Å, λCo-Kα2 = 1.79285 Å) and an X’Celerator detector. The patterns were refined starting from the structural model proposed by Isasi et al. for the orthorhombic Li2Ni(SO4)219, and using the Rietveld method30 as implemented in the FullProf program31,32. The results of these refinements are presented in Figure 3 and Table 1 for the well-crystallized cobalt-based compound, and in Figures SI-2 to SI-5 and Tables SI-3 to SI-6 for the other phases. The orthorhombic structure adopted by the Li2M(SO4)2 phases (M = Fe, Co, Zn, Mg, and Ni) consists of isolated MO6 octahedra linked through SO4 tetrahedra by oxygen vertices, as illustrated in Figure 4. Each octahedron is linked to six SO4 tetrahedra, whereas each SO4 tetrahedron is only bound to three MO6 octahedra. The unshared fourth corner of each SO4 tetrahedron are directed towards the open channels, where the lithium cations reside. The lithium cations occupy distorted octahedral sites within these tunnels, thus forming zig-zag chains of edge-sharing LiO6 octahedra along the b-axis. In addition, because such tunnels are favorable for lithium diffusion, we were encouraged to explore the electrochemical and conduction properties of the orthorhombic Li2M(SO4)2 phases (M = Fe, Co, Zn, Mg and Ni). Electrochemistry Bearing in mind the high redox potential observed for the monoclinic marinite iron analogue (3.83 V vs. Li+/Li0),17 we first probed the electrochemical response of the iron-based compound. Electrochemical characterization were performed with Swagelok-type cells, using lithium metal as the negative electrode, a composite made of active material and carbon SP (80:20 wt%) as the positive electrode, and 1M LiClO4 in PC or LiPF6 in EC:DMC as the electrolyte. A typical voltage-composition trace obtained for the orthorhombic Li2Fe(SO4)2 cycled at a rate of C/20 is shown in Figure 5a. It reveals the presence of two successive plateaus centered at 3.73 V and 3.85 V vs. Li+/Li0, as corroborated by the corresponding dx/dV derivative curve (Figure 5b) and GITT (Galvanostatic Intermittent Titration Technique) experiments (Figure SI-6). During the first charge to 4.5 V vs. Li+/Li0, about 0.95 Li ions are removed from Li2Fe(SO4)2, while 0.85 Li ions are reinserted upon the subsequent discharge (Figure 5a), resulting in a reversible capacity of around 91 mAh·g-1 (theoretical specific capacity: 102 mAh·g-1). Subsequent charge/discharge cycles superimpose well over more than 10 cycles, as seen from Figure 5c. Moreover, note that under similar electrode preparation and testing conditions, electrodes made of orthorhombic Li2Fe(SO4)2 repeatedly showed less polarisation (~100 mV) than the monoclinic one (~300 mV), suggesting better insertion/desertion kinetics, though a thorough optimization of the electrode


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composition was not performed. This was further confirmed by determining its rate capability, which shows that 90 % of the initial capacity is maintained at a 1 C rate as compared to only 80 % for the marinite polymorph (Figure SI-7). Beside Li2Fe(SO4)2, the electrochemical responses of the other orthorhombic phases Li2M(SO4)2 (M = Co, Zn, Mg and Ni) were probed as well; however, none of them showed any electrochemical activity within the stability range of the electrolyte. These results do not come as a surprise given (i) that zinc and magnesium cations are not electrochemically active and (ii) that the higher potentials observed for the CoIII+/CoII+ and NiIII+/NiII+ redox couples with respect to the FeIII+/FeII+ couple (~1.3 and 1.6 V increase, respectively) in other polyanionic compounds would result in potentials that cannot be reached with current electrolytes. The intriguing presence of two definite plateaus in the charge/discharge curves of orthorhombic Li2Fe(SO4)2, presumably indicative of two subsequent biphasic de/intercalation processes, called for further examination of the lithiation/delithiation mechanism. PITT (potentiostatic intermittent titration technique) experiments (Figure SI-8) show that the current decay curves do not follow a simple diffusion-driven Cottrell-type law in the whole composition range but indicate the succession of two reversible biphasic deinsertion/insertion processes close to 3.73 V and to 3.85 V, in agreement with the OCV limits determined previously from the GITT experiment in the corresponding composition ranges. This mechanism was further confirmed by performing in operando XRD measurements (Figure 6). During the first half of the charge, the diffraction peaks of the pristine electrode material (Figure 6, green pattern) progressively disappear, while a new set of diffraction peaks appear at higher 2θ angles that becomes unique at Δx ~ 0.5. These results indicate the existence of an intermediate phase, “Li1.5Fe(SO4)2” (Figure 6 blue pattern). During the second part of the charge, new diffraction peaks once again appear, while the peaks associated with the intermediate phase disappear, yielding a new powder pattern at Δx ~ 1.0 (Figure 6, orange pattern). Thus, these measurements also reveal the existence of a delithiated phase, “Li1.0Fe(SO4)2”. Upon discharge, the peak amplitudes change in the opposite direction, implying two subsequent bi-phasic transitions, as observed during charge. Moreover, the similarity between the XRD powder patterns of the fully discharged and pristine samples indicates a fully reversibility process. To isolate the delithiated “Li1.0Fe(SO4)2” phase, the pristine orthorhombic Li2Fe(SO4)2 was chemically oxidized using NO2BF4 dissolved in acetonitrile as an oxidizing agent. The XRD pattern obtained for this chemically delithiated phase was in excellent agreement with the pattern recorded at the end of the first charge during the in operando XRD experiment described above. The Bragg reflections observed for “Li1.0Fe(SO4)2” could be indexed into the orthorhombic space group Pbca with the lattice parameters a = 9.1652(4) Å, b = 8.9304(4) Å, c = 13.4532(8) Å and V = 1101.12(9) Å3. We carried out a Rietveld refinement of these data, starting with the structure established for the mother phase Li2Fe(SO4)2, and


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using soft constraints for the S–O distances and fixed isotropic temperature factors Biso. The results of this Rietveld refinement are shown in Figure 7a and Table SI-7, and the goodness of this fit suggests that the general structural framework of the pristine orthorhombic Li2Fe(SO4)2 phase is maintained upon delithiation, which is merely accompanied by a volume reduction of about 5.7 %. Note, however, that the positions of the lithium cations could not be determined from our laboratory XRD data; neutron powder diffraction experiment will be performed in the near future on delithiated samples to determine more precisely the structure of this “Li1.0Fe(SO4)2” phase. 57

Fe Mössbauer spectroscopy performed on the delithiated orthorhombic “Li1.0Fe(SO4)2” (Figure 7b)

confirmed the complete oxidation of FeII+ to FeIII+. However, the spectrum showed two different environments of FeIII+ with a ratio of 78/22 (orange and purple lines, respectively), as a result of the oxidation of the two environments of FeII+ previously observed in the mother phase Li2Fe(SO4)2 (Figure 2b). A solid-state 7Li NMR spectrum was acquired on the chemically oxidized orthorhombic “Li1.0Fe(SO4)2” phase (Figure 7c) to characterize the local lithium environments, which shows an intense 7

Li signal at 28 ppm with a broad spinning-side-band manifold (Figure 7c, top) associated with lithium

within well-crystallized “Li1.0Fe(SO4)2” domains. The different isotropic 7Li shift in this structure compared to the pristine orthorhombic Li2Fe(SO4)2 phase (-12 ppm, Figure 2c) indicates that the lithium atoms experience different average local environments after chemical oxidation, consistent with the slight structural distortions and volume changes discussed above. The narrower line width (full-width-halfmaximum = 14 ppm) compared to the pristine phase (77 ppm, Figure 2c) is likely due to enhanced lithium mobility within the crystal structure upon partial lithium removal, where faster lithium motions partially average the electron-nuclear interactions that broaden the 7Li NMR linewidths. Interestingly, the 7

Li paramagnetic shift anisotropy parameters are very similar between the chemically oxidized

“Li1.0Fe(SO4)2” and pristine orthorhombic Li2Fe(SO4)2 phases (Table SI-2), suggesting that the lithium cations experiences similar anisotropic interactions within both structures (e.g., dipole-dipole couplings between 7Li spins and unpaired electron spins), despite their different isotropic 7Li shifts (dominated by Fermi-contact interactions). Quantitative deconvolutions of the solid-state 7Li NMR spectrum establish that 86 % of total lithium content is associated with the orthorhombic “Li1.0Fe(SO4)2” phase, whereas 12 % is associated with less ordered domains (broad 7Li signal at 52 ppm) and 2 % is due to an Li2SO4 impurity (diamagnetic 7Li signal at 0 ppm). Conductivity measurements The better insertion/deinsertion kinetics obtained for the orthorhombic as compared to the monoclinic Li2Fe(SO4)2 polymorphs suggests it has enhanced ionic and electronic conducting properties. To confirm this assumption, we measured the conductivity of dense pellets made of the orthorhombic and the


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monoclinic Li2Co(SO4)2 phases in the temperature range 200-400 °C with impedance spectroscopy using a Solartron Analytical Modulab unit, using ionically blocking gold electrodes. Figure 8 shows the evolution of the temperature dependence of the a.c. conductivity for the two aforementioned phases, showing a difference of 1.5 orders of magnitude between the monoclinic and orthorhombic polymorphs (blue and green curves in Figure 8, respectively). The experimental curves were fitted with the Arrhenius equation T = 0·exp(-Ea/kBT), where T is the conductivity at the temperature T, 0 is a pre-exponential factor, Ea the apparent activation energy for Li+ migration, and kB the Boltzmann constant. The Arrhenius model yields a good fit to the experimental data, revealing an apparent activation energy of 1.54 eV for the monoclinic phase of Li2Co(SO4)2 and 1.19 eV for its orthorhombic counterpart. Moreover, linear extrapolations of the curves yield room temperature a.c. conductivity values of 2.6·10-18 S/cm for the former and 2.2·10-14 S/cm for the latter. Note that these conductivity values are not corrected by pellets’ porosities; the values are thus underestimated. The poor conductivity figures obtained for the monoclinic phase are in agreement with the ones previously reported for the marinite monoclinic Li2Fe(SO4)2 phase: Ea ~ 1.1 eV and RT ~ 10-16 S/cm.33 For reference, better conductivity values were reported for one of the most praised materials for electric vehicle applications, namely LiFePO4: Ea ~ 0.6 eV and RT ~ 109

S/cm.34

Further heating of the Li2Co(SO4)2 orthorhombic compound revealed an abrupt drop of the conductivity after 360ºC, as shown in the inset of Figure 8, indicative of a possible phase transition at this temperature, as had previously been observed for the zinc analogue.22 Such a finding was a further impetus to embark in a thorough study of the thermal stability of the newly synthesized orthorhombic Li2M(SO4)2 phases (M = Co, Fe, Zn, Ni) with respect to the corresponding monoclinic ones; the latter being found to preserve their structure from 2 K18 to their respective decomposition or melting temperature (from 400 to 600 °C depending on the nature of the divalent metal)17,22. Polymorph stability Differential Scanning Calorimetry (DSC) was used to determine the thermal stabilities of the different polymorphs and to search for possible transitions between each polymorph. The DSC curve obtained for the orthorhombic Li2Co(SO4)2 sample is shown in Figure 9. The endothermic peak at the beginning of the measurement between 100ºC and 200°C is associated with a loss of water molecules, presumably from moisture adsorbed to the particle surfaces. The exothermic peak at 420°C, however, could be ascribed to the orthorhombic/monoclinic phase transition. A subsequent second heating cycle of the same sample shows no feature on the DSC curve, which confirms the thermal stability of the monoclinic phase as opposed to the orthorhombic phase. From the DSC curve, we can deduce the enthalpy of the phase


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transformation as being approximately -2.36 kJ/mole for Li2Fe(SO4)2 and -1.06 kJ/mole for Li2Co(SO4)2, which are in the range of expected values for polymorphic materials. To confirm the phase transition observed in the DSC curve, an in situ high-temperature XRD experiment was performed on the orthorhombic Li2Co(SO4)2 phase (Figure 10a). Up to 320°C, the orthorhombic phase (green) did not show any significant structural changes, except for improved crystallinity. Between 340°C and 360°C, the peaks of the monoclinic phase began appearing at the expense of the peaks of the pristine orthorhombic phase (pink), yielding the pure monoclinic Li2Co(SO4)2 phase (blue) at 380°C, which remained stable till the end of the measurement at 480°C (light blue pattern). The results are in good agreement with the DSC measurements, with the exception of a difference in the transition temperature. However, this difference is simply a result of the use of different heating rates (10°C/min for DSC, 2-hours steps for each temperature for XRD). Similar results were observed for the iron- and zincbased materials, with transition temperatures of approximately 360ºC and 480ºC, respectively (see Figure SI-10). From these variable-temperature measurements, we can conclude that the monoclinic marinite phases are more stable under heating than their orthorhombic counterparts. To further probe the stability of the two structures, and bearing in mind that the monoclinic marinite Li2Fe(SO4)2 was altered by ball-milling, the various Li2M(SO4)2 monoclinic marinite phases were subjected to planetary milling trials using varying ball-milling times. In all cases, we observed by X-ray diffraction the progressive alteration of the marinite phase and, for extended period of milling, the crystallization of the orthorhombic analogue, as shown in Figure 10b for the cobalt system. Therefore, considering that mechanical milling is associated with local increases of both temperature and stress, we can speculate that the orthorhombic phases are favored at higher pressures as compared to the monoclinic marinite compounds. For the sake of conciseness regarding the polymorphs stability, we summarize within Figures 11 and 12 the relative stability of the different Li2M(SO4)2 materials depending upon the nature of the synthesis approach. Figure 11 shows the formation of the polymorphs depending on the synthesis approach, starting from the sulfate precursors, while Figure 12 shows the transformation from one polymorph into the other, either by heating the orthorhombic polymorph or by ball-milling the monoclinic polymorph. While the ball-milling and ceramic processes usually lead to the orthorhombic and monoclinic polymorphs, respectively, there are a few exceptions worthy of comments. For example, the orthorhombic polymorph of Li2Mn(SO4)2 could not be stabilized despite many attempts, nor could the monoclinic marinite phase of the nickel-based compound. It is quite likely that such an effect is linked to the size of the transition metals (Figure 13), where the orthorhombic phase is favored by smaller transition metal


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cations (NiII+) and the marinite phases by larger cations (MnII+). This remains to be rationalized by both thermodynamic measurements, as recently being done by our group for other systems35,36 as well as by theoretical calculations37. Discussion In summary, through the use of both ball-milling synthesis and low temperature solid-state reactions, we have unveiled polymorphism in the Li2M(SO4)2 compounds (M = Mn, Fe, Co, Zn, Mg, Ni), which can crystallize into either an orthorhombic or a monoclinic structure. These new orthorhombic polymorphs exhibit interesting differences in their structural, conducting, and electrochemical properties with respect to the marinite monoclinic polymorphs. From a structural point of view, the orthorhombic and the monoclinic marinite structures exhibit similar structural motifs but exhibit one key difference. Both structures consist of a 3D framework of MO6 and SO4 groups, where the MO6 octahedra are isolated from each other and only linked through the SO4 tetrahedra via oxygen corner-sharings. In particular, each MO6 octahedron is surrounded by six cornersharing SO4 groups, and each SO4 tetrahedron is connected to only three corner-sharing MO6 octahedra. The main difference between the orthorhombic and monoclinic structures of Li2M(SO4)2 is the way in which the MO6 and SO4 polyhedra are interconnected, as seen in Figure 14. As a consequence, the M–M distances in the orthorhombic phases are shorter compared to those in the monoclinic marinite counterparts, while its density is greater (ρortho ~ 2.97-3.14 g/cm3 vs. ρmono ~ 2.55-2.99 g/cm3). Compared to its monoclinic analogue,

57

Fe Mössbauer spectroscopy shows that the orthorhombic form

exhibits a larger quadrupole splitting as compared to the monoclinic marinite structure (Table SI-1), likely due to the shorter distance between the iron centers. Such a larger quadrupole splitting has been observed for many inorganic compounds with oxygen-coordinated octahedral FeII+ environments and was found to be related to a complex combination of geometrical, electronic and nuclear parameters.38,39 For sulfates such as FeSO4·nH2O (n = 0 and 7), large quadrupole splitting (> 3 mm/s) have been reported and were attributed to a combined action of axial and rhombic crystalline fields and the spin-orbit interaction.40,41 Interestingly, solid-state 7Li NMR measurements reveal that the orthorhombic Li2Fe(SO4)2 exhibits a negative isotropic 7Li NMR shift (-12 ppm), compared to a positive isotropic 7Li shift (32 ppm) observed for its monoclinic analogue. This difference in the sign of the 7Li shifts, which are dominated by Fermicontact interactions between the 7Li spins and unpaired electron spins, indicates a different Fermi-contact mechanism: the delocalization mechanism occurs when there is a direct overlap between the Li–O–Fe molecular orbitals, resulting in a transfer of positive electron spin density to the lithium (and hence a


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positive shift), whereas the polarization mechanism occurs when there is no direct overlap between these orbitals, resulting in the transfer of negative electron spin density to the lithium (and hence a negative shift).42 Thus, the two polymorphs exhibit Fermi-contact interactions that are dominated by different mechanisms, which is a direct consequence of their different Li–O–Fe bonding geometries. Turning to the electrochemical properties, while the shape of the voltage-composition curve changes between the polymorphs, the potential of the FeIII+/FeII+ redox couple remains nearly the same. Such results are in sharp contrast with the polymorphism displayed by either LiFeSO4F or LiFeSO4OH, for which we have measured important differences in the FeIII+/FeII+ redox potential between two polymorphs10–12,23,24. One explanation for this comparable redox potential might be due to the close structural similarities between the orthorhombic and monoclinic Li2Fe(SO4)2 polymorphs, since polymorphs showing a pronounced structural difference such as triplite and tavorite LiFeSO4F or tavorite and layered LiFeSO4OH show FeII+/FeIII+ redox couples differing by nearly 300 mV. Another interesting finding is the appearance of two subsequent plateaus/biphasic processes for the orthorhombic Li2Fe(SO4)2 polymorph, which occur at redox potentials that are close in value (3.73 V vs. 3.85 V), but are nevertheless distinct. This stair-case voltage plateau usually suggests a poor structural elasticity of the precursor phase, which cannot accommodate the volume change associated with the removal of lithium, and thus an intermediate phase forms that acts as a buffer.43 Therefore, the minor voltage difference between the two plateaus is indicative of minor structural differences between the Li2-xFe(SO4)2 (x = 0, 0.5 and 1) phases. This is consistent with our preliminary structural refinements of the delithiated phases, which indicate a structural transformation involving neither bond breakings nor anisotropy changes of the unit cell. Whatever the exact Li-insertion/deinsertion structural mechanism in the orthorhombic Li2Fe(SO4)2 polymorph, the small mechanical stresses associated with its limited volume change (ΔV/V ~ 5.7 %) are consistent with its better capacity retention upon cycling, as compared to the monoclinic polymorph, which exhibits a greater volume change (ΔV/V ~ 12 %). Lastly, returning to the synthesis processes, it is interesting to note the important role of mechanical milling in the discovery of new polymorphs among the sulfate-based family of polyanionic compounds (tavorite/triplite LiMSO4F13, tavorite/layered LiMSO4OH24, and monoclinic marinite/orthorhombic Li2M(SO4)2). Such a success is directly linked to the relatively low melting temperature of sulfates, such that the local frictional heating generated at the particle level due to ball milling is sufficient to reach temperatures at which mass transport is sufficient for these phases to form. In all cases, the polymorph that forms through mechanical milling is always the denser one, as illustrated in Figure 13 for the Li2M(SO4)2 series. This result indicates that besides temperature, the local stresses associated with the mechanical milling process are also essential to stabilize the orthorhombic polymorph. This is not


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surprising, as numerous compounds that require high hydrostatic pressure during their syntheses have also been obtained at room temperature by ball milling. A notable example is LiC2, the high pressure member of the LixC6 family that usually requires 200 kbar, was synthesized by ball milling in three hours.44 This example shows why reactive ball milling should be more intensively exploited in the search for new battery electrode materials. Conclusion We have reported the synthesis and electrochemical properties of a new polymorph of the recently reported monoclinic marinite structure Li2Fe(SO4)2, which crystallizes in an orthorhombic Pbca structure. Electrochemically, this new orthorhombic Li2Fe(SO4)2 phase cannot compete in terms of potential and capacity with common polyanionic positive electrode materials, such as LiFePO4. Nevertheless, its stabilization emphasizes the importance of exploring different synthesis methods, and mechanical milling in particular, to fabricate new electrode materials. We therefore predict that understanding and mastering the use of mechanical milling will further fuel the search for new battery electrode materials, including those beyond the sulfate-based polyanionic compounds currently being explored by our group.


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Acknowledgements The authors would like to thank Matthieu Courty (LRCS, UPJV Amiens) for performing ATG and DSC measurements, as well as Ludovic Delbes and Benoît Baptiste (IMPMC, UPMC Paris) for their help in conducting high-temperature XRD experiments. L.L. and M.R. also thank Raphaël Janot and Chinmayee Subban (Collège de France, Paris) for fruitful discussions about mechanical milling, as well as Charles Delacourt (LRCS, UPJV Amiens) for his advice on conducting and analyzing electrochemical and conductivity measurements. The authors also thank Juan Rodríguez-Carvajal for helpful discussions. The authors thank Mohamed Ben Hassine (LRCS, UPJV Amiens) for acquiring TEM images of selected samples. M.R. is grateful to Alexandre Ponrouch (ICMAB-CSIC, Barcelona) for carrying out room-temperature carbon coating on one of the samples. L.L. acknowledges the French Agence Nationale de Recherche via the research project "Hipolite" for her Ph.D. grant. R.J.M. gratefully acknowledges support from a Marie Curie Postdoctoral Fellowship through the European Union Seventh Framework Programme (FP7/2007-2013) under grant agreement 330735.

Author Information Notes * Corresponding author: [email protected]. Authors Contributions: M.R., G.R. and J.-M.T. originated the idea and designed the research approach. L.L. and M.R. carried out the syntheses and electrochemical work. L.L., M.R. and G.R. analyzed the diffraction patterns and crystal structures. M.T.S. and M.R. collected and analyzed the Mössbauer spectra. R.J.M. and M.D. acquired and analyzed all solid-state NMR data. L.L. and C.L. performed the conductivity measurements. L.L., M.R., G.R. and J.-M.T. wrote the manuscript. All authors discussed the results and final manuscript.

Supporting Information Complementary experimental details, structural information, detailed Mössbauer and solid-state NMR analyses, additional electrochemical characterization, SEM images. This material is available free of charge via the Internet at http://pubs.acs.org.


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(27) Cot, L. C. R. Acad. Sci. C 1968, 266, 921. (28) Cot, L. Rev. Chim. Min. 1969, 6, 1041. (29) Touboul, M.; Quarton, M.; Lokaj, J.; Kettmann, V. Acta Crystallographica Section C Crystal Structure Communications 1988, 44, 1887. (30) Rietveld, H. M. J. Appl. Crystallogr. 1969, 2, 65. (31) Rodr gue -Carvajal, J. FullProf Suite. (32) Rodr gue -Carvajal, J. Physica B 1993, 192, 55. (33) Clark, J.; Eames, C.; Reynaud, M.; Rousse, G.; Chotard, J.-N.; Tarascon, J.-M.; Islam, M. S. Journal of Materials Chemistry A 2014, 2, 7446. (34) Delacourt, C.; Laffont, L.; Bouchet, R.; Wurm, C.; Leriche, J.-B.; Morcrette, M.; Tarascon, J.-M.; Masquelier, C. J. Electrochem. Soc. 2005, 152, A913. (35) Radha, A. V.; Furman, J. D.; Ati, M.; Melot, B. C.; Tarascon, J.-M.; Navrotsky, A. J. Mater. Chem. 2012, 22, 24446. (36) Radha, A. V.; Subban, C. V.; Sun, M. L.; Tarascon, J.-M.; Navrotsky, A. J. Mater. Chem. A 2014, 2, 6887. (37) Ben Yahia, M.; Lemoigno, F.; Rousse, G.; Boucher, F.; Tarascon, J.-M.; Doublet, M.-L. Energy Environ. Sci. 2012, 5, 9584. (38) Evans, R. J. Ph.D. dissertation, University of Ottawa: Ottawa, 2006. (39) Benmokhtar, S.; El Jazouli, A.; Chaminade, J. P.; Gravereau, P.; Wattiaux, A.; Fournès, L.; Grenier, J. C.; Waal, D. Journal of Solid State Chemistry 2006, 179, 3709. (40) Ok, H. N. Phys. Rev. B 1971, 4, 3870. (41) Ingalls, R. Phys. Rev. 1964, 133, A787. (42) Grey, C. P.; Dupré, N. Chem. Rev. 2004, 104, 4493. (43) Casas-Cabanas, M.; Roddatis, V. V.; Saurel, D.; Kubiak, P.; Carretero-González, J.; Palomares, V.; Serras, P.; Rojo, T. J. Mater. Chem. 2012, 22, 17421. (44) Janot, R.; Conard, J.; Guérard, D. Carbon 2001, 39, 1931. (45) Brown, I. D.; Altermatt, D. Acta Crystallogr. B 1985, 41, 244. (46) Ponrouch, A.; Goñi, A. R.; Sougrati, M. T.; Ati, M.; Tarascon, J.-M.; Nava-Avendaño, J.; Palacín, M. R. Energy Environ. Sci. 2013, 6, 3363.


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Tables Table 1: Crystallographic data and atomic positions of the orthorhombic Li2Co(SO4)2 determined from the Rietveld refinement of its XRD pattern. A Bond Valence Sum analysis (BVS) is also indicated.45

Orthorhombic Li2Co(SO4)2 Pbca a = 9.20688 (9) Å Wyckoff Atom position Co 8c

b = 9.10175 (9) Å Occupancy

x/a

RBragg = 1.99 %

χ2 = 1.7

c = 13.71190 (16) Å

V = 1149.04 (2) Å3

y/b

z/c

Biso (Å2)

BVS

1.0

0.8618 (3)

0.6018 (3)

0.3768 (2)

1.43 (10)

2.077 (24)

Li1

8c

1.0

0.483 (2)

0.721 (3)

0.358 (3)

1.5

0.971 (38)

Li2

8c

1.0

0.723 (3)

0.534 (3)

0.628 (3)

1.5

1.013 (43)

S1

8c

1.0

0.6612 (7)

0.8146 (5)

0.5095 (4)

1.36 (11)

5.866 (98)

O11

8c

1.0

0.4981 (10)

0.8008 (10)

0.5261 (8)

1.10 (12)

1.784 (42)

O12

8c

1.0

0.7103 (9)

0.9671 (13)

0.4961 (9)

1.10 (12)

2.035 (58)

O13

8c

1.0

0.6884 (9)

0.7342 (11)

0.4187 (7)

1.09 (12)

2.079 (48)

O14

8c

1.0

0.7420 (7)

0.7547 (11)

0.5931 (11)

1.09 (12)

1.848 (59)

S2

8c

1.0

0.5751 (5)

0.4313 (6)

0.2705 (3)

1.36 (11)

5.521 (83)

O21

8c

1.0

0.4815 (11)

0.5018 (10)

0.3482 (7)

1.09 (12)

1.865 (49)

O22

8c

1.0

0.5232 (10)

0.4614 (9)

0.1707 (7)

1.09 (12)

2.067 (45)

O23

8c

1.0

0.5773 (9)

0.2657 (11)

0.2788 (8)

1.09 (12)

1.938 (50)

O24

8c

1.0

0.7304 (12)

0.4874 (11)

0.2769 (7)

1.09 (12)

1.833 (48)


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Figure captions Figure 1 (a) Voltage-composition traces, (b) 57Fe Mössbauer spectra, and (c) solid-state 7Li NMR spectra (central band) obtained for different electrode materials made of the monoclinic marinite Li2Fe(SO4)2 active material and conductive carbon. The powdered Li2Fe(SO4)2 sample was either carbon-coated at room temperature46 before being mixed with carbon Super P by soft hand-grinding (top), or ball-milled (highenergy impact Spex 8000® miller) with CSP for 30 (middle) and 60 minutes (bottom). Increasing the ball-milling time increases the sloping part (green) of the charge curve at the expense of the plateau (blue), which is correlated with the appearance of different iron and lithium environments (see legend) observed in the 57Fe Mössbauer and solid-state 7Li NMR spectra. Figure 2 (a) Evolution of the XRD patterns of a stoichiometric mixture of Li2SO4 and FeSO4 ball-milled (planetary Retsch PM100 miller) for different times, showing the progressive formation of the orthorhombic Li2Fe(SO4)2 phase. (b) 57Fe Mössbauer spectrum and (c) solid-state 7Li NMR spectrum (top, entire spectrum; bottom, central band) of orthorhombic Li2Fe(SO4)2 obtained after 10 hours of planetary ballmilling. Figure 3 Rietveld refinement of the XRD pattern of the orthorhombic Li2Co(SO4)2 phase. The green crosses, black line, and grey line represent the observed, calculated, and difference patterns, respectively. The positions of the Bragg reflections are shown as vertical black bars. Figure 4 Representation of the orthorhombic structure of the Li2M(SO4)2 (M = Co, Fe, Zn, Ni, Mg) phases, viewed along the (a) [010], (b) [-110] and (c) [100] directions. MO6 octahedra and SO4 tetrahedra are blue and green, respectively. Lithium cations are shown as grey spheres. O atoms are not shown for clarity but define the corners of each polyhedra. Figure 5 Electrochemical characterization of the orthorhombic Li2Fe(SO4)2 material obtained after 5 hours of ball-milling. The voltage composition trace (a) and its derivative (b) show two plateaus at 3.73 V and 3.85 V respectively. The low polarization observed in (a), the good capacity retention after 10 charge-discharge cycles (c), and the Ragonne curve (d) indicate good electrochemical performances of the orthorhombic Li2Fe(SO4)2.


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Figure 6 In operando XRD measurement of the orthorhombic Li2Fe(SO4)2 active material, acquired during the first charge/discharge cycle. (a) Evolution of the XRD patterns along the electrochemical curve. The green, blue, and orange patterns are assigned to the pristine Li2Fe(SO4)2 compound, the half-delithiated intermediate phase “Li1.5Fe(SO4)2”,

and the delithiated phase “Li1.0Fe(SO4)2”, respectively. The

right-hand part displays zoomed-in areas of relevant peaks, which show the two subsequent bi-phasic mechanisms that occur during (b) charge (bottom) and (c) discharge (top). Figure 7 (a) Rietveld refinement of the XRD pattern of the chemically oxidized orthorhombic “Li1Fe(SO4)2” phase. The orange crosses, black line, and grey line represent the measured, calculated, and difference patterns, respectively. The positions of the Bragg reflections are shown as black bars. The blue star indicates a peak assigned to a homemade anoxic chamber. (b) 57Fe Mössbauer spectrum and (c) solid-state 7Li NMR spectrum (top, entire spectrum; bottom, central band) of the chemically oxidized orthorhombic “Li1Fe(SO4)2” phase: Figure 8 Temperature dependence of the a.c. conductivity of the orthorhombic (green) and monoclinic (blue) polymorphs of Li2Co(SO4)2. The inset shows the abrupt decrease of the a.c. conductivity that occurs when heating the orthorhombic phase at temperatures higher than 360ºC. Figure 9 DSC measurement of the orthorhombic Li2Co(SO4)2 heated from room temperature up to 500 °C with a ramp of 10 K/min. The green line corresponds to the first heating cycle of the pristine orthorhombic Li2Co(SO4)2 compound, whereas the blue line corresponds to a subsequent heating cycle of the product obtained after cooling the sample to room temperature. The first peaks between 100°C and 200°C are associated with a loss of moisture initially adsorbed on the particle surface. The exothermic peak at 420°C is assigned to the phase transition from the orthorhombic to the monoclinic structure. Figure 10 (a) In situ high-temperature XRD measurements of Li2Co(SO4)2 upon heating, which follows the orthorhombic-to-monoclinic phase transition. The green, pink, and blue patterns correspond to the orthorhombic polymorph, the co-existence of the orthorhombic and monoclinic phases, and the monoclinic polymorph. The light blue pattern (top) was recorded after cooling the sample to room temperature, showing that the monoclinic structure is preserved. (b) XRD measurements of a Li2Co(SO4)2


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sample, which follows the monoclinic-to-orthorhombic phase transition that occurs upon mechanical milling. Figure 11 Schematic diagram summarizing the nature of the Li2M(SO4)2 polymorph (orthorhombic in green vs. monoclinic in blue) obtained from a stoichiometric mixture of MSO4 and Li2SO4, and using different synthesis conditions (mechanical milling vs. solid-state route) for the six metals M = Ni, Mg, Zn, Co, Fe, Mn. Figure 12 Schematic diagram summarizing the stability of the orthorhombic (green) polymorph upon heating and monoclinic (blue) polymorph upon mechanical milling of Li2M(SO4)2 for the four metals M = Ni, Co, Fe, Mn. Figure 13 Evolution of the volume per formula unit as a function of the ionic radii of the divalent cations MII+ in the Li2M(SO4)2 series (M = Ni, Mg, Zn, Co, Fe, Mn). Blue and green points represent the monoclinic marinite and orthorhombic polymorphs, respectively. Figure 14 Comparison of the MO6 and SO4 connectivity in the 3D frameworks of the (a, c) monoclinic marinite and (b, d) orthorhombic structures. In both structures, the MO6 octahedra (blue and grey) are isolated from each other and are only linked through six corner-sharing SO4 tetrahedra (green and orange). The main structural differences between the two polymorphs are associated with the position of the bridging SO4 groups (orange) around the MO6 octahedra within the chains running along the a-axis in the monoclinic structure (c) and the b-axis in the orthorhombic structure (d).


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Figures Figure 1

Caption (a) Voltage-composition traces, (b) 57Fe Mössbauer spectra, and (c) solid-state 7Li NMR spectra (central band) obtained for different electrode materials made of the monoclinic marinite Li2Fe(SO4)2 active material and conductive carbon. The powdered Li2Fe(SO4)2 sample was either carbon-coated at room temperature47 before being mixed with carbon Super P by soft hand-grinding (top), or ball-milled (highenergy impact Spex 8000® miller) with CSP for 30 (middle) and 60 minutes (bottom). Increasing the ball-milling time increases the sloping part (green) of the charge curve at the expense of the plateau (blue), which is correlated with the appearance of different iron and lithium environments (see legend) observed in the 57Fe Mössbauer and solid-state 7Li NMR spectra.


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Figure 2

Caption (a) Evolution of the XRD patterns of a stoichiometric mixture of Li2SO4 and FeSO4 ball-milled (planetary Retsch PM100 miller) for different times, showing the progressive formation of the orthorhombic Li2Fe(SO4)2 phase. (b) 57Fe Mössbauer spectrum and (c) solid-state 7Li NMR spectrum (top, entire spectrum; bottom, central band) of orthorhombic Li2Fe(SO4)2 obtained after 10 hours of planetary ballmilling.


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Figure 3

Caption Rietveld refinement of the XRD pattern of the orthorhombic Li2Co(SO4)2 phase. The green crosses, black line, and grey line represent the observed, calculated, and difference patterns, respectively. The positions of the Bragg reflections are shown as vertical black bars.


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Figure 4

Caption Representation of the orthorhombic structure of the Li2M(SO4)2 (M = Co, Fe, Zn, Ni, Mg) phases, viewed along the (a) [010], (b) [-110] and (c) [100] directions. MO6 octahedra and SO4 tetrahedra are blue and green, respectively. Lithium cations are shown as grey spheres. O atoms are not shown for clarity but define the corners of each polyhedra.


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Figure 5

Caption Electrochemical characterization of the orthorhombic Li2Fe(SO4)2 material obtained after 5 hours of ball-milling. The voltage composition trace (a) and its derivative (b) show two plateaus at 3.73 V and 3.85 V respectively. The low polarization observed in (a), the good capacity retention after 10 charge-discharge cycles (c), and the Ragonne curve (d) indicate good electrochemical performances of the orthorhombic Li2Fe(SO4)2.


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Figure 6

Caption In operando XRD measurement of the orthorhombic Li2Fe(SO4)2 active material, acquired during the first charge/discharge cycle. (a) Evolution of the XRD patterns along the electrochemical curve. The green, blue, and orange patterns are assigned to the pristine Li2Fe(SO4)2 compound, the half-delithiated intermediate phase “Li1.5Fe(SO4)2”,

and the delithiated phase “Li1.0Fe(SO4)2”, respectively. The

right-hand part displays zoomed-in areas of relevant peaks, which show the two subsequent bi-phasic mechanisms that occur during (b) charge (bottom) and (c) discharge (top).


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Figure 7

Caption (a) Rietveld refinement of the XRD pattern of the chemically oxidi ed orthorhombic “Li 1.0Fe(SO4)2” phase. The orange crosses, black line, and grey line represent the measured, calculated, and difference patterns, respectively. The positions of the Bragg reflections are shown as black bars. The blue star indicates a peak assigned to a homemade anoxic chamber. (b) 57Fe Mössbauer spectrum and (c) solid-state 7Li NMR spectrum (top, entire spectrum; bottom, central band) of the chemically oxidized orthorhombic “Li1.0Fe(SO4)2” phase.


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Figure 8

Caption Temperature dependence of the a.c. conductivity of the orthorhombic (green) and monoclinic (blue) polymorphs of Li2Co(SO4)2. The inset shows the abrupt decrease of the a.c. conductivity that occurs when heating the orthorhombic phase at temperatures higher than 360ºC.


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Figure 9

Caption DSC measurement of the orthorhombic Li2Co(SO4)2 heated from room temperature up to 500 °C with a ramp of 10 K/min. The green line corresponds to the first heating cycle of the pristine orthorhombic Li2Co(SO4)2 compound, whereas the blue line corresponds to a subsequent heating cycle of the product obtained after cooling the sample to room temperature. The first peaks between 100°C and 200°C are associated with a loss of moisture initially adsorbed on the particle surface. The exothermic peak at 420°C is assigned to the phase transition from the orthorhombic to the monoclinic structure.


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Figure 10

Caption (a) In situ high-temperature XRD measurements of Li2Co(SO4)2 upon heating, which follows the orthorhombic-to-monoclinic phase transition. The green, pink, and blue patterns correspond to the orthorhombic polymorph, the co-existence of the orthorhombic and monoclinic phases, and the monoclinic polymorph. The light blue pattern (top) was recorded after cooling the sample to room temperature, showing that the monoclinic structure is preserved. (b) XRD measurements of a Li2Co(SO4)2 sample, which follows the monoclinic-to-orthorhombic phase transition that occurs upon mechanical milling.


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Figure 11

Caption Schematic diagram summarizing the nature of the Li2M(SO4)2 polymorph (orthorhombic in green vs. monoclinic in blue) obtained from a stoichiometric mixture of MSO4 and Li2SO4, and using different synthesis conditions (mechanical milling vs. solid-state route) for the six metals M = Ni, Mg, Zn, Co, Fe, Mn.


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Figure 12

Caption Schematic diagram summarizing the stability of the orthorhombic (green) polymorph upon heating and monoclinic (blue) polymorph upon mechanical milling of Li2M(SO4)2 for the four metals M = Ni, Co, Fe, Mn.


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Figure 13

Caption Evolution of the volume per formula unit as a function of the ionic radii of the divalent cations MII+ in the Li2M(SO4)2 series (M = Ni, Mg, Zn, Co, Fe, Mn). Blue and green points represent the monoclinic marinite and orthorhombic polymorphs, respectively.


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Figure 14

Caption Comparison of the MO6 and SO4 connectivity in the 3D frameworks of the (a, c) monoclinic marinite and (b, d) orthorhombic structures. In both structures, the MO6 octahedra (blue and grey) are isolated from each other and are only linked through six corner-sharing SO4 tetrahedra (green and orange). The main structural differences between the two polymorphs are associated with the position of the bridging SO4 groups (orange) around the MO6 octahedra within the chains running along the a-axis in the monoclinic structure (c) and the b-axis in the orthorhombic structure (d).


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TOC abstract graphic


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Synthesis and Electrochemical Performance of the Orthorhombic

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