Singular Structural and Electrochemical Properties in Highly Defective

Article pubs.acs.org/cm

Singular Structural and Electrochemical Properties in Highly Defective LiFePO4 Powders Robin Amisse,†,‡,§ Moulay Tahar Sougrati,§,∥ Lorenzo Stievano,§,∥ Carine Davoisne,† Goran Dražič,‡ Bojan Budič,‡ Robert Dominko,‡,§ and Christian Masquelier*,†,§ †

Laboratoire de Réactivité et Chimie des Solides, CNRS UMR 7314, Université de Picardie Jules Verne, 33 rue Saint Leu, 80039 Amiens Cedex, France ‡ National Institute of Chemistry, P.O.B. 660, SI-1001 Ljubljana, Slovenia § ALISTORE-ERI, FR CNRS 3104, 80039 Amiens Cedex, France ∥ Institut Charles Gerhardt, CNRS 5253, 1919 route de Mende, 34090 Montpellier, France S Supporting Information *

ABSTRACT: Highly defective LiFePO4 powders were synthesized via a modified version of the coprecipitation in aqueous medium method using oxidizing experimental conditions. A pure olivine phase containing 44 at. % of Fe3+ was obtained after only 10 min at 108 °C, and the evolution of the structure and purity was followed during reaction. The nature of the native defects and their influence on the crystallographic structure and on the electrochemical reaction mechanisms were thoroughly studied by a combination of ex situ and in situ methods, using high-resolution transmission electron microscopy, inductive coupled plasma, X-ray diffraction, and Mössbauer spectroscopy. The high concentration of defects induced a unit-cell volume 4 Å3 smaller than that of stoichiometric LiFePO4, with a complete cationic redistribution over the M1 and M2 crystallographic sites, as well as a completely new electrochemical signature. A precise structural model during electrochemical operation of the pristine defective LiFePO4 was built.

1. INTRODUCTION Since the pioneering works of Padhi et al.1 on olivine phosphates for Li-ion battery positive electrode materials, a great effort has been put into the optimization of LiFePO4, leading in particular to a large variety of studies focused on the comprehension of the electrochemical reaction mechanisms. On one side, a deep understanding of the structure of this material has been achieved, yet strong contradictions remain between the features predicted by ab initio calculations and experimental values, mainly for the description of the Li+ diffusion paths inside the structure. It has also been demonstrated that the presence of structural defects and surface effects strongly influence these values.2 Indeed, even though LiFePO4 powders have been synthesized with a great variety of different routes, most of them imply a post-treatment at elevated temperature in reducing atmosphere or in the presence of a reducing agent under inert atmosphere (carbon sources for carbon-coating, etc.) in order to remove traces of Fe3+ in the final compound. The native Fe3+ defects in LiFePO4 have thus been subject to several studies that have determined their naturemainly surface ferric films or antisite Fe3+ in the bulk.3−12 With regard to electrochemistry, ab initio studies13,14 supported by experiments have shown the detrimental effect of Fe3+ atoms in the Li+ site (antisite), since they block the Li+ diffusion along the channels [010]. However, high amounts of these defects lead to new structural features and electro© 2015 American Chemical Society

chemical properties that require further analysis, as presented in this work. This study focuses on the nature and properties of structural defects present in high amounts in LiFePO4 powders, enhancing the features already observed for slightly defective powders but also causing new and drastic changes in the structural characteristics and in the reactivity of this material toward Li.

2. EXPERIMENTAL SECTION Pristine powders of defective LiFePO4 were prepared through a direct H2O-based precipitation method adapted from previous works.15−17 A solution of FeSO4·7H2O (Sigma-Aldrich) was prepared in water and stirred continuously for 20 h prior to the reaction in order to oxidize part of the Fe2+ into Fe3+. A stoichiometric amount of orthophosphoric acid (H3PO4, 85 wt %, Sigma-Aldrich) was then added, followed by dimethylsulfoxide (DMSO) for a resulting solvent ratio of 1:1 in DMSO/H2O. The concentrations of each precursor were chosen so that the final concentration in both FeSO4·7H2O and H3PO4 was 0.3 mol·L−1 in 300 mL of solution. In a second step, an aqueous solution of 3.15 mol·L−1 of LiOH·H2O (Alfa Aesar) was added dropwise under vigorous stirring at room temperature, followed by in situ pH measurement and steady uptakes (with total Vtaken < 10% of Vsolution). The resulting solution was brought to its boiling point of Received: February 6, 2015 Revised: May 12, 2015 Published: May 12, 2015 4261

DOI: 10.1021/acs.chemmater.5b00470 Chem. Mater. 2015, 27, 4261−4273

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Chemistry of Materials ∼108 °C and kept under reflux for 90 h. A small portion of the solution was taken at different reaction times during the process, and each uptake was filtered and centrifuged (5000 rpm; 10 min) several times with distilled water and once with acetone, then dried overnight at 70 °C. The resulting pristine powders did not undergo any additional treatment prior to the characterizations. 57 Fe transmission Mössbauer spectra were measured with a source of 57Co in rhodium metal. During the measurements, both the source and the absorber were kept at ambient temperature (294 K). The spectrometer was operated in the constant acceleration mode with a triangular velocity waveform. The velocity scale was calibrated with the magnetically split sextet spectrum of a high-purity α-Fe foil as the reference absorber at room temperature. The powdered samples (50 mg) were mixed with boron nitride to make absorbers with a section of 2 cm2. For operando experiments, a specific in situ cell was used.18 Simultaneous operando Mössbauer and X-ray powder diffraction (XRD) analyses were carried out using the setup described elsewhere.19 The spectra were fitted to appropriate combinations of Lorentzian profiles representing hyperfine magnetic sextets and quadrupole doublets by least-squares methods using the program PC-Mos II.20 In this way, spectral parameters such as the quadrupole splitting (QS), the isomer shift (IS), and the relative resonance areas of the different spectral components were determined. Isomer shifts are given relative to α-Fe metal. XRD diagrams were collected from Bruker D8 diffractometers (40 kV, 40 mA), using either the Co Kα or Cu Kα radiation, mounted in θ−θ configuration. High-quality diffraction patterns were recorded overnight between 2θ = 10° and 2θ = 100° by steps of 10 s/0.02° or during 3 h between 2θ = 10° and 2θ = 70° by steps of 5 s/0.02°. The neutron powder diffraction diagrams were measured inside vanadium cylindrical containers on the high-resolution powder diffractometer D2B at Institut Laue-Langevin (ILL, Grenoble, France) in the 8− 158°(2θ) range, and with a λ = 1.594 Å wavelength. The microscope used for the local characterization of the particles was a probe Cs corrected Jeol ARM 200 F equipped with a cold field emission gun (FEG) and operated at 200 kV. The presence of Fe, and its valence state, was followed by an electron energy loss spectroscopy (EELS) system GIF Quantum ER with dual EELS from Gatan. Some of the experiments were done at −175 °C in Gatan 636 double-tilt liquid-nitrogen cooling holder. The morphology, size, and elemental compositions of the samples were characterized by scanning electron microscopy (SEM) using a FEI Quanta 200F microscope equipped with a Link Isis apparatus (Oxford) for energy dispersive X-ray spectroscopy (EDX). Fe, Li, and P concentrations were determined using inductively coupled plasma optical emission spectrometry (ICP OES, Varian, model 715 ES). The electrochemical tests were carried out on a Biologic Macpile or a Biologic VMP3, using Swagelok-type cells. Galvanostatic cycling with potential limitations (GCPL) was performed between 2.0 and 4.2 V vs Li+/Li0. The positive electrode was built by mixing 83 wt % of active material and 17 wt % of Timcal super P conductive carbon. Glass fiber separators and an electrolyte made of LiPF6 in ethylene carbonate/ dimethyl carbonate (1:1) were used at the interface between the positive electrode and negative lithium foil (∼500 μm thick) electrode. The structural evolution of stoichiometric and defective LiFePO4 during Li+ extraction/insertion was followed by in situ Mössbauer spectroscopy coupled with X-ray diffraction. The sample was placed in a specific electrochemical cell18 right behind an X-ray-transparent beryllium window also acting as the current collector. Alternate in situ Mössbauer spectra and XRD patterns were collected (by steps of 2 and 1 h, respectively, repeated 8 times at each floating step) during a complete charge−discharge cycle versus lithium metal as the negative electrode, in floating mode at selected potentials and at a rate of C/40 during cycling.

oxidized Fe precursor was implemented to the synthesis by direct precipitation in aqueous medium previously developed by Delacourt et al.16,21 A solution of FeIISO4·7H2O in water was aged under magnetic stirring for 20 h in air prior to the reaction. During the process, the color of the solution turned from light green, characteristic of Fe2+ in water, to orange, characteristic of Fe3+. The direct chemical titration of the airaged FeSO4·7H2O solution used for the reaction with KMnO4 revealed the presence of ∼10 at. % of Fe3+. A solution of LiOH in water was added dropwise at room temperature to the solution containing FeSO4·7H2O and H3PO4 in H2O/DMSO solvent so as to provide Li+ ions and increase the pH up to a value of 7 in ∼40 min. The pH evolution was followed in situ and is presented in Figure 1 as a

Figure 1. In situ pH monitoring during the drop-by-drop LiOH·H2O introduction into the solution of precursors. Inset: X-ray diffraction patterns of the precipitated powders taken at selected [LiOH] introduced values, prior to the reaction.

function of the concentration in LiOH. During the LiOH introduction, uptakes were performed at steady intervals in order to follow the evolution of possible crystalline phases (inset of Figure 1). Three main regions can be identified. (a) In the first one, the pH remains stable at a value of 3 and an amorphous dark green precipitate of vivianite Fe3(PO4)2· 8H2O or other (hydroxi-)phosphates is formed in agreement with the formation of the (hydroxi-)phosphates reported by Jensen et al.22 (b) At concentrations of LiOH introduced above 0.55 mol· L−1, the pH increases. A first crystalline phase appears, named phase 1, and disappears in the following region. (c) A more pronounced slope is observed from concentrations above 0.8 mol·L−1 up to pH 7. This last region is characterized by the presence of three crystalline phases, phases 2 and 3 and Li2SO4·H2O. Despite the limited number of species involved in this reaction, phases 1−3 could not be identified, and they only bring information on the apparent complexity of the process. Their nature is most likely similar to the intermediate species described by Jensen et al.22 The presence of crystalline Li2SO4· H2O detected by X-ray diffraction even after several washings of the precipitate with H2O accounts for its large excess in solution. This feature highlights a significant difference between this synthesis and previously reported direct precipitation21 or

3. RESULTS AND DISCUSSION 3.1. Preparation of Defective LiFePO4 Powders by Direct Precipitation in DMSO−H2O Solution. In the search for a highly defective pristine material, the use of a partially 4262


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Chemistry of Materials hydrothermal23 methods in which the ratio Li/Fe/P was set to 3:1:1. After the obtained solution reached its boiling temperature, the mechanism of formation of defective LiFePO4 powders was followed through regular uptakes of powder samples, collected from 5 min to 90 h of reaction under reflux. The reaction time was set to zero with the first drop of reflux upon heating, and a first sample was taken after 5 min of reaction. The corresponding X-ray diffraction pattern is very similar to the one at the end of the LiOH introduction, as the reaction had not proceeded yet. After ∼10 min of reaction under reflux, a sudden change in the color of the blend was observed, from dark green to grayish. The XRD pattern corresponding to the 10-min uptake is presented in Figure 2a and can be fully

grows progressively until 90 h of reaction. A straightforward Rietveld refinement on the X-ray diffraction patterns using LiFePO4 and LiFePO4OH phases revealed the presence of ∼20% of tavorite by the end of the reaction. This result is in accordance with previous studies on the aging mechanisms of LiFePO4 in water or moisture.24−28 The SEM micrographs of the uptakes after 3 and 90 h in Figure 2b also provide useful information on the reaction and the products formed over time. After 3 h of reaction, particles similar to the ones formed after 10 min can be observed, with a rhombohedral shape and average 500 × 300 × 200 nm dimensions. The growth of the particles is limited during the whole reaction, with only a ∼50% size increase capped after 60 h. In parallel, the formation of a tavorite phase after long reaction times is witnessed in these images by the modification of the surface of the particles. The rhombohedral shape is altered in the 90 h powder, with the presence of particles with rectangular faces (Figure 2b). A further examination of the X-ray diffraction patterns reveals a significant shift of the 2θ positions of the diffraction peaks, which denotes a clear evolution with the reaction time. The (210) and (011) reflections in Figure 2a move toward lower angles as the reaction proceeds, owing to an increase in the unit-cell parameters. The unit-cell volume, after 10 min of reaction, is extraordinarily low (286.8 Å3), and progressively increases up to 290.1 Å3 at the end of the reaction (90 h). One should note that this volume is still significantly lower than the reference one of 291.3 Å3.29 In parallel, the reflections of the successive X-ray diffraction patterns exhibit a progressive decrease in the full width at half maximum (fwhm). This trend accounts for an increase in the size of the coherent crystalline domains, i.e., the global crystallinity of the powder. Previous observations on the unit-cell volume of LiFePO4 powders synthesized via direct precipitation have suggested a decrease of the unit-cell volume with increasing amounts of structural defects. This regular increase of the unit-cell volume with the reaction time may thus be an indicator of the progressive ordering of the structure. Despite the oxidizing character of the reaction medium, the thermodynamic stability of a less disordered LiFePO4 can be the driving force for this evolution. To answer this conjecture, 57Fe Mössbauer spectroscopy analyses were performed on the different powder uptakes, and the spectra of the 3 and 90 h powders are shown in Figure 2c. These spectra were fitted with a combination of quadrupole doublets for paramagnetic Fe contributions in octahedral environments. The details about the choice of these environments and their precise structural meaning will be discussed thereafter, and all the Mössbauer parameters are gathered in Table S1 in the Supporting Information. (a) The powder obtained after 10 min possesses a remarkably high amount of Fe3+ (44 at. %). Because no crystalline impurity is detected in the powder by XRD, this proportion of Fe3+ corresponds to actual structural defects, which proves that the olivine LiFePO 4 structure can accommodate a considerable amount of defects. The successive spectra recorded for longer reaction times denote a progressive decrease of the proportion of Fe3+. After 3 h, about 34 at. % of Fe3+ is present, and 29 at. % is present after 90 h (Figure 2c). This trend follows the increase in the unit-cell volume mentioned above and supports a progressive ordering of the LiFePO4 structure upon reaction, despite the oxidizing medium in which it takes place.

Figure 2. (a) X-ray diffraction patterns of the powders taken after 10 min, 3 h, and 90 h of reaction, with the theoretical Bragg positions of LiFePO4 and LiFePO4OH as a reference. (b) SEM images of the synthesized powders after 3 and 90 h of reaction and (c) the corresponding fitted Mössbauer spectra.

indexed using the Pnma space-group. In addition, and to detect a possible amorphous FePO4·nH2O phase in the powder, a thermal treatment under reducing atmosphere (500 °C, H2/Ar 10:90, 3 h) was performed, and no crystalline Fe2P2O7 phase appeared in the X-ray diffraction pattern. This early result proves the feasibility of synthesizing pure LiFePO4 powders with a very fast and low-temperature route. The X-ray patterns recorded for three selected uptakes in Figure 2a outline the evolution of the crystalline phases formed during the process. A single triphylite phase is present until 3 h of reaction. After 6 h, a tavorite LiFePO4OH phase appears and 4263


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Chemistry of Materials (b) After 90 h of reaction, a new Fe3+ contribution appears with hyperfine parameters corresponding to those of the tavorite LiFePO4OH (IS = 0.40; QS = 0.54 mm·s−1), in agreement with the values reported by Marx et al.30 It accounts for ∼14 at. % of the total Fe and further supports the aging mechanism proposed before. Besides, 15 at. % of Fe3+ are still present within the structure. The presence of remaining structural defects may be responsible for the slightly lower unit-cell volume than the one of stoichiometric LiFePO4. To understand the nature and characteristics of the highly defective powders formed during short reaction times, the powder obtained after 3 h of reaction was chosen for further investigations and was labeled LFP-3h. A reference stoichiometric LiFePO4 was obtained by the post-treatment in a reducing atmosphere (10% H2 in Ar) of slightly defective powders and was labeled LFP-ref. Figure 3 presents the

from surface sites intensively reported in the literature, two points need to be mentioned. First, and as noted before, the powder is composed of submicrometric particles without internal porosity, in which the surface effects are negligible compared to the bulk. Second, the presence of amorphous phosphates FeIIIPO4·nH2O, often present as an impurity in lowtemperature syntheses and invisible by XRD, can be discarded because the annealing in reducing atmosphere at 500 °C did not trigger the crystallization of Fe2P2O7, which occurs when Li-deficient phases are present in the pristine material. The peculiarity of this new defective LiFePO4 powder called for an in-depth characterization, which was carried out using various complementary characterization techniques. Within the variety of possible structural defects in LiFePO4 powders, the most energetically favorable was calculated to be the antisite Fe.10,31,32 Its presence in the powders synthesized in this work was preliminarily evidenced by Mössbauer spectroscopy and X-ray diffraction in the powders formed at different reaction times. A direct observation of this antisite Fe was carried out with the high-angle annular dark field scanning transmission electron microscopy (HAADF-STEM) analysis of the near-edge region of several LiFePO4 particles. The methodology used in this experiment was directly adapted from the work of Chung et al.4 The difference in contrast between Fe and Li atoms in HAADF-STEM images was used to detect the presence of Fe in the Li site. This intensity is proportional to the Z7/8 of the atom, so that the Fe atoms appear clearly in LiFePO4 with a bright contrast, whereas Li is invisible. Besides, O atoms are not detected either, and P atoms provide a slight contrast in the image. The crystals were oriented along the [010] direction, so that the Li channels face the electron beam. In this configuration, the presence of Fe in these channels could be easily evidenced, and a structural model of this configuration is schematized in Figure 4c. This specific orientation is also favorable to the analysis, because the distinct rows of FeO6 octahedra make the identification of the atoms/orientation straightforward from the image. A crystal from the LFP-ref powder was used as a reference, and the HAADF-STEM image is shown in Figure 4b. The Fe atoms can be clearly distinguished, and the filtered image highlights the presence of P atoms in their vicinity, which form the asymmetric contour in the bright dots. This image can be directly related to the atomic model in Figure 4c. The integrated intensities, recorded in between two rows of Fe, are presented in Figure 4d. The signal is only consists of background noise, which suggests that no or little Fe would be present in this region, where the M2 site is present. On the other hand, the near-edge region of a crystal of LFP-3h with the same orientation is pictured in Figure 4a. After only a few scans, the crystal becomes amorphous; thus, a long acquisition time or a fine adjustment of the beam is not possible. The filtered image in the inset of Figure 4a possesses bright dots in between the rows of Fe atoms. They highlight the presence of Fe at these positions, which is further confirmed by the integrated intensities in the region marked in blue and plotted in Figure 4d. The curve possesses several maxima, regularly spaced by ∼4.5 Å. This distance corresponds to the theoretical Li−Li distance in the LiFePO4 structure in Figure 4c. In summary, the HRSTEM-HAADF images show the presence of antisite iron in the defective powders, which can be observed as far as ∼15 nm from the surface, i.e., deeper than in previous studies.33 Further observations within the full bulk were made impossible because of the sample thickness and the increased reactivity of the

Figure 3. Full-profile matching refinements performed on X-ray diffraction patterns of as-synthesized (a) and stoichiometric (b) LiFePO 4 powders using the Pnma space group, with their corresponding unit-cell parameters. (c) Comparison of the relative intensities and positions of selected diffraction peaks between the two powders.

comparison between the XRD patterns of (a) LFP-3h and (b) LFP-ref. The two phases crystallize in the same space group, and no crystalline impurity can be detected. The full-profile matching using Le Bail method and the Pnma space group indicates a unit-cell volume as low as 287.3 Å3 for the defective powder, which is 4 Å3 lower than the value reported in the literature.1,29 More precisely, the a and b parameters are mainly responsible for this difference, while c seems to be unaffected: a noticeable deviation of the (020) plane reflection versus the (211), inherent to the strong decrease in the b parameter of 0.9%, is clearly visible in Figure 3c. This strong difference in unit-cell implies the presence of many defects, vacancies, local distortions, and/or cationic redistribution. It is further supported by the difference in the relative intensities of each reflection, as detailed in Figure 3c, with discrepancies ranging from −12% to +38% between 20° and 30° (2θ). To discriminate between defects attributed to the bulk and those 4264


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At first, an extensive bibliographic study was performed to extract general trends on the hyperfine parameters. Starting from the early work of Andersson et al.,35 relevant publications on the study of LiFePO4 were selected, and the IS and QS values were collected for each Fe contribution. As a result, the values of QS were plotted in Figure 5b as a function of the

Figure 4. (a) HAADF-STEM images of LFP-3h in the near-edge region of a particle oriented along [010]. Inset: corresponding magnified and filtered image (b) HAADF-STEM images of LFP-ref in the near-edge region of a particle oriented along [010]. Inset: corresponding magnified and filtered image superimposed with the according structure with the Fe atoms in green, the P atoms in pink, the Li atoms in yellow, and the O atoms in dark gray. The brighter areas have a triangular shape provoked by the influence of different atoms: iron, which has the highest atomic number, and then phosphorus and oxygen. To highlight this fact, the atomic projection along [010] has been superimposed to the filtered image in the inset of (b) showing the agreement between the atomic arrangement and the image contrast. (c) Schematics of the LiFePO4 structure in the same configuration. The average Li−Li distance is 4.7 Å. (d) Integration of the intensities recorded along the blue and red zones in, respectively, LFP-3h and LFP-ref, between the rows of Fe.

Figure 5. (a) Mössbauer spectrum of the pristine defective LiFePO4 fitted with four Fe environments. Their relative population and hyperfine parameters are listed on the left. (b) Quadrupole splitting vs isomer shift for different octahedral iron environments of olivine-type LiFePO4 reported in the literature, compared with the values for the pristine defective LiFePO4. The ellipses are inclusive.

sample under the electron beam (amorphization) for highly defective samples. The quantification of this Fe in the M1 antisite was performed using Mössbauer spectroscopy on the pristine powder, as a complement to the previous techniques and to refine the structural model. The refinement of a model was not straightforward, and for each valence state several populations of iron had to be taken into account and rigorously compared with the hyperfine parameters of LiFePO4 reported in the literature. Over the past 15 years, 57Fe Mössbauer spectroscopy has been occasionally used to characterize LiFePO4 materials for Li-ion batteries, with its use being often limited to the detection of FeIII impurities in the pristine powders or after aging. During in situ measurements,34 it allows the precise determination of the FePO 4 and LiFePO 4 respective proportions. The hyperfine parameters remain nevertheless constant during the electrochemical reaction process, which greatly limits the impact of this technique on the understanding of complex mechanisms. In this work, special attention was ascribed to the analysis of Mössbauer data, from both the literature and the presented experiments, in order to extract useful information from the calculated IS and QS values. Indeed, the presence of structural defects in LiFePO4 powders significantly affects those values. As a consequence, new Fe environments can be foreseen and later correlated with results from other characterization techniques.

values of IS to create a landscape of the different Fe populations in LiFePO4. From this summary, three main regions could be isolated and are represented by the ellipses. The main one (blue ellipsis7,25,34−53) includes the doublets with IS and QS values around 1.22 and 2.95 mm·s−1, respectively. This set of parameters corresponds to high-spin FeII in an octahedral site, namely, the regular M2 site in olivine LiFePO4. The very small discrepancies found among the different reports are probably due to the lower sensitivity to lattice distortions of FeII than FeIII, as already pointed out by Perea et al.34 As a consequence, the QS parameter, which directly accounts for this feature, is almost identical in all reports. Note that some small discrepancies on IS data may originate from spectrometer calibration considerations. Fewer reports have characterized the environment of FeIII in heterosite FePO4 by Mössbauer spectroscopy. All the results are included in the purple ellipsis,34,35,39,44,45,52 characterized by a narrow range of QS values around 1.5 mm·s−1 and IS values between 0.35 and 0.45 mm·s−1. With those two regions, the identification of the possible Fe environments in stoichiometric LixFePO4 (0 ≤ x ≤ 1) is straightforward. However, the presence of structural defects, surface impurities, and other 4265


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Chemistry of Materials Table 1. Atomic Positions of LFP-3h Determined by Rietveld Refinement from Neutron Powder Diffraction Data site

Wyck.

x

y

z

atom

occ.

Biso

Li

4c

Fe

4c

P1 O1 O2 O3

4c 4c 4c 8d

0 0 0.28025(23) 0.28025(23) 0.09523(43) 0.09696(44) 0.45311(34) 0.16693(31)

0 0 0.25 0.25 0.25 0.25 0.25 0.04310(43)

0 0 0.97660(55) 0.97660(55) 0.41407(63) 0.74493(83) 0.20857(76) 0.27882(53)

Li Fe Fe Li P O O O

0.853(2) 0.147(2) 0.896(2) 0.092(2) 0.923 1 1 2

1.18 1.18 1.187(5) 1.187(5) 1.135(63) 1.627(65) 1.549(78) 1.512(48)

regular stoichiometry of phosphorus in LiFePO4. However, this assumption, together with the presence of ∼34 at. % of Fe3+ in the pristine powder as revealed by Mössbauer spectroscopy (Figure 2c), would induce two major contradictions. First, the crystallographic sites available for the iron atoms, namely, M1 and M2 in the respective Wyckoff positions 4a and 4c, would be occupied at >100%. Second, the excess of Fe combined with its partial oxidation to the trivalent state would induce a large excess of charge in the global composition, in contradiction with the electroneutrality of the powder. Consequently, a slight under-stoichiometry in Li and P atoms was taken into account in the refinement of the crystal structure using neutron diffraction data. The results of the powder neutron diffraction analysis of LFP-3h are presented in Figure 6. A single olivine phase was

foreign phases (doping, off-stoichiometry) gives rise to a third region noted by the orange ellipsis.7,25,37,43,47−49,54−56 This region accounts for FeIII contributions in Li1−εFePO4 with a distorted/frustrated octahedral local environment, such as the antisite/surfacic FeIII, or the FeIII from other impurity phases (e.g., tavorite LiFePO4OH). In addition, some polymorphs of57−60iron oxyhydroxides57−60 or LiFeO261−65 may exhibit doublets with similar IS−QS values. With this knowledge, the Mössbauer spectrum of LFP-3h was carefully analyzed, and the refinement of the different contributions from the model proposed is plotted in Figure 5a. As expected from the complexity of the X-ray diffraction data and the high amount of structural defects, the refinement of a model was not straightforward, and several populations of iron for each valence state had to be taken into account and rigorously compared with the hyperfine parameters of LiFePO4 reported in the literature. The major contribution (blue) arises from a FeII population of ∼54 at. % with an IS of 1.23 mm·s−1 and a QS of 2.94 mm·s−1. These values are characteristics of FeII in the regular M2 site in the lithiated form of LiFePO4, as reported by more than 20 groups of researchers.7,25,34−53 This first result implies that the remaining 46 at. % of Fe in the structure originates from defects and/or local distortions of the lattice. Among them, another population of ∼12 at. %, marked in green, possesses a similar value of IS but a lower QS, accounting for a distorted local environment in the M2 site. The remaining 34 at. % of Fe was previously attributed to Fe in the trivalent state but can be divided into two populations. The first one in brown accounts for 14 at. % and possesses hyperfine parameters of IS = 0.44 mm·s−1 and QS = 0.7 mm·s−1. These values were previously reported43,66 and ascribed to Fe in the M1 (anti)site. In addition, this proportion is much higher than the amount of antisite Fe that can be found in the literature, typically 5 at. % or less. The last FeIII population of 20 at. %, marked in orange, is characterized by a high QS value compared to other defect-type FeIII. Similar values have been reported in the case of LiFe1−yMnyPO4 powders studied in situ during electrochemical operation.34,67 In the intermediate region between Mn and Fe operating potentials, the Mn2+ cations were found to induce a local distortion in the vicinity of Fe3+, causing a decrease of the QS from regular values (1.3 mm· s−1, within the purple ellipsis) down to 1.1 mm·s−1. In our case, this distortion would originate not from the presence of Mn in the structure but most probably from the high amount of native structural defects. Another analogy could be drawn, in some aspects, with the ferrisicklerite (Li,Na)y(Mn,Fe)zPO4 in which the large excess of Mn over Fe has a noticeable effect on the hyperfine parameters of Fe.68 An inductively coupled plasma mass spectrometry (ICP-MS) elemental analysis of LFP-3h was carried out and converged into a relative ratio in Li/Fe/P equal to 1:1.1:1, considering a

Figure 6. Neutron diffraction pattern of LFP-3h, with the parameters and formula calculated from the refinement coupled with ICP results. Inset: structural model of the unit-cell.

refined as no extra peak can be detected. The calculated unitcell parameters confirm the previous observations, as the low unit-cell volume suggests a defective structure. The Rietveld refinement of the structure was carried out starting from stoichiometric LiFePO4 and using the Pnma space group (see Table 1). At first, the occupancy of Fe on the Li (M1) site was refined independently and converged to y ≈ 0.14 in (LiFey)M1(Fez)M2PO4. This value is in perfect agreement with the results from the Mössbauer spectroscopy analysis of LFP3h, which determined 14 ± 1 at. % of antisite Fe. The 4266


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Chemistry of Materials remaining z = 0.9 Fe (from ICP results) were set to belong in the regular M2 site, as no evidence of Fe in another crystallographic site was found. The refinement of the structure was then carried out with two constraints, namely, the ratio of cations dictated by ICP analyses and the global electroneutrality. In addition, the total occupancy of both M1 and M2 (Li + Fe) was also capped to 1. The occupancy of P was refined and spontaneously converged to a value close to 0.9. This outcome supports the possibility of a P deficiency discussed above. Owing to that, this occupancy was set to 0.923, as deduced from the ICP analyses. To further confirm this hypothesis, the average P−O bond lengths were compared to those of a stoichiometric LiFePO4. In the reference material, the average P−O bond length is 1.54 Å. The contracted unitcell of LFP-3h would tend to possess slightly shorter bonds. However, the calculated value is 1.56 Å in this case. This increase can be attributed to the electrostatic repulsion of the oxygen atoms in the presence of a P deficiency. Although this phenomenon has not been reported to our knowledge, all the previous observations lead to this conclusion. The final formula [Li0.85Fe3+0.15](M1)[Fe2+0.71Fe3+0.19Li0.09□0.01](M2)P0.92O4 is hence proposed. It implies the full occupancy of the M1 site, and only 1 at. % vacancies on the M2 site. The model still presents discrepancies with the measured data (χ2 = 9.70 and RBragg = 3.47), but no further improvement could be provided by the current knowledge of the material. On the other hand, the presence of Fe3+ in the structure was clearly identified by Mössbauer spectroscopy. The smaller ionic radius of Fe3+ affects the Fe−O bond length in the M2 site. While the average Fe−O length is 2.16 Å in stoichiometric LiFePO4, it is only 2.13 Å in LFP-3h. This difference can be the cause of the lower unit-cell volume observed, which suggests that the presence of Fe3+ in defective LiFePO4 tends to decrease the overall unit-cell dimensions. In parallel, the similar ionic radii of Li+ and Fe3+ would imply that the presence of antisite Fe3+ does not affect these dimensions. In summary, these highly defective LiFePO4 powders can be described with an olivine structure affected by the presence of high amounts of Fe3+, which triggers a strong cationic redistribution over the M1 and M2 sites, together with lattice distortions also caused by the vacancies in Li and P. 3.2. Electrochemical Signatures of Highly Defective LiFePO4 Probed by Operando X-ray Diffraction. The presence of defects in LiFePO4 powders is reported to have a significant impact on their electrochemical signature, either beneficial or detrimental.8,11,32,43,69,70 In the case of LFP-3h, the existence of a high amount of defects/disorder is responsible for major changes in the electrochemical reaction mechanism. Figure 7 displays the galvanostatic charge/discharge cycles of LFP-3h at a C/50 rate. This low rate was chosen so as to reduce the kinetics effects and approach a thermodynamic mechanism of lithium exchange. During the first charge, only ∼62% of the specific capacity is achieved. This observation is in full agreement with the previous results from Mössbauer analysis denoting the presence of 34 ± 2% of Fe3+ in the pristine powder. To further support this idea, the same experiment was carried out with a discharge first (not shown here), and ∼33% of the specific capacity was achieved. After the first charge, a reversible capacity of 130 mAh·g−1 is obtained, with a mean Coulombic efficiency of 99.5% over cycles 2−20. This LFP-3h thus exhibits a good cycling stability, despite the high amount of structural disorder and defects. One should note that the specific capacity could be limited here by the

Figure 7. Potential profile of LFP-3h measured in galvanostatic mode, at a C/50 rate. Upper inset: evolution of the specific capacity upon cycling. Lower inset: dq/dU derivative in oxidation and reduction.

absence of carbon-coating on the particles or agglomerated particles, yet 76% of the theoretical capacity was recovered. The major feature here is the presence of a low-potential contribution, which corresponds to the redox activity of the iron originally present as Fe3+ before cycling. This simple observation implies that part of Fe3+ originally present in the distorted M2 site (in orange in Figure 5) operates at a lower potential than the regular Fe in the M2 site. This effect may originate from the variation in the local geometry of the redox center, affecting the operating potential beyond the inductive effect, as proposed recently by Saubanère et al. using density functional theory (DFT) calculations.71 In fact, the quadratic elongation (QE) and bond angle variance (BAV) differ between LFP-3h and a stoichiometric LiFePO 4 , with, respectively, QE = 1.038 against 1.045 and BAV = 129.8 square-degrees against 155.6 square-degrees. The other population of Fe3+ attributed to Fe in the M1 site (antisite) was previously reported to operate in this range of potentials.43 In a more general view, this low-potential contribution presents a sloping shape between 2 and 3.2 V vs Li+/Li, which suggests a single-phase mechanism upon Li insertion/removal, further supported by the very broad peaks in the dq/dU derivative during both oxidation and reduction in Figure 7, bottom inset. The high polarization of ∼300 mV associated with this process may originate from kinetic limitations or potential hysteresis and will be discussed in light of in situ X-ray diffraction and high-rate galvanostatic measurements. On the other side, the potential profile at higher potential values, ∼3.5 V vs Li+/Li, is comparable with the reported data on stoichiometric LiFePO4. However, some differences remain: the polarization of only 20 mV, illustrated in the dq/dU derivative (Figure 7, bottom inset), is remarkably low for an electrode formulated without carbon-coating. Furthermore, the presence of structural defects in LiFePO4 was reported multiple times to hinder the Li diffusion along the [010] channels,32 which does not seem to apply in this case. Finally, the dq/dU derivative exhibits broad peaks during both oxidation and reduction, even at this low C-rate, which would suggest a partial solid solution reaction. A further insight into the presumed 4267


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Figure 8. (a, b) Evolution of the unit-cell parameters of the defective LiFePO4 annealed in air at different temperatures. (c) Corresponding X-ray diffraction patterns recorded at room temperature, with highlight on the diffraction peaks of the planes (020) and (002).

disappearance of peaks can be detected. The successive refinements of the structural model using a single olivine phase offer a very good agreement with the experimental data, with an average of χ2 = 1.7 and Rbragg = 0.3. With this model, the corresponding calculated unit-cell parameters are gathered in Figure 8c. A first observation reveals that the c parameter generally increases during the delithiation and decreases during the lithiation, in an opposite fashion to a and b. The two different regions present in the galvanostatic cycling match the two different slopes observed for the variations of the unit-cell parameters. This suggests that at least two distinct iron environments are involved in the process, as will be discussed later in light of the in situ Mössbauer experiment. More precisely, an inversion of the slope in the evolution of the c parameter seems to occur at the very end of discharge, which would correspond to the electrochemical reaction of Fe3+ in the M1 site (antisite). It should also be added that the reaction follows the same path upon insertion and removal of Li. A complete solid-solution Li insertion/removal mechanism during electrochemical process is hence suggested; nevertheless, the possible existence of a limited biphasic reaction in the middle of the 3.5 V pseudoplateau cannot be excluded owing to the temporal and spatial limitations of the measurement, as well as to a minor peak broadening in the X-ray diffraction patterns. Its existence, however, has not been evidenced by any other characterization technique.

solid solution character of this reaction has been provided by galvanostatic intermittent titration technique (GITT) and potentiostatic intermittent titration technique (PITT) measurements (cf. Supporting Information, Figure S2). To further characterize the electrochemical performances of this material, a galvanostatic measurement at 1 C rate has been performed (Supporting Information, Figure 3). The cell could deliver 65 mAh·g−1, or 50% of the capacity obtained at a C/50 rate, stable over 100 cycles, and the corresponding Coulombic efficiency reaches 99.6%. These results demonstrate a favorable Li diffusion within the structure. Contrary to several other studies on the defects in LiFePO4,8,9,32 the presence of these defects does not have a huge detrimental effect on the overall electrochemical performances. All these observations reveal a completely new and peculiar reaction mechanism versus Li, which was further studied with the help of in situ XRD and Mössbauer spectroscopy during the electrochemical reactions. Figure 8a displays the potential versus composition signature of LFP-3h during charge and discharge. The first charge exhibits a capacity of 82 mAh·g−1, compared to 128 mAh·g−1 for the following charge/discharge cycles. It corresponds to 64% of the reversible capacity and only involves the upper-potential contribution, as observed earlier. A detailed view of four selected reflections in the corresponding X-ray diffraction patterns is presented in Figure 8b. A progressive shift of all these reflections occurs during all the charges and discharges. In addition, no appearance or 4268


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Figure 9. (a, b) Mössbauer spectra measured in situ during electrochemical operation of LFP-3h at (a) 3.90 V and (b) 2.00 V vs Li+/Li0, fitted with four Fe environments. (c) Successive potential steps recorded in floating mode and separated by galvanostatic charge/discharge at C/20 rate. (d) Evolution of the relative population of each Fe environment during the electrochemical operation. (e) Hyperfine parameters of the different Fe environments and the corresponding crystallographic configuration models.

structure (mainly Li excess and P deficiency) may have a compensative effect in terms of unit-cell dimensions. From these considerations, we conclude that the presence of defects in the structure is thus responsible not only for the appearance of a new phenomenon at potentials lower than 3.5 V but also for a complete change in the electrochemical reaction mechanisms as compared to that of stoichiometric LiFePO4. This unusually large single-phase region is present at room temperature, for particles that are not nanometer-sized, and at a slow C-rate. Owing to these features, the surface and kinetic effects can be neglected, and this phenomenon can solely be ascribed to the presence of intrinsic structural defects. 3.3. Electrochemical Signatures of Highly Defective LiFePO4 Studied by Operando Mössbauer Spectroscopy. The 57Fe Mössbauer results obtained during the coupled in situ XRD/Mössbauer experiment at selected potentials18 are gathered in Figure 9. Four spectral components were used in the fitting, two for Fe2+ and two for Fe3+. However, while the

The overall unit-cell volume variation of this defective powder upon Li+ extraction/insertion was found to be lower than for stoichiometric LiFePO4 (ΔV/V = 5.7% against 6.6%). On the basis of simple geometrical considerations, this lower volume swelling would allow an easier accommodation of the strains during electrochemical operation, thus favoring a singlephase reaction. The value at the end of charge is significantly higher than the one reported for stoichiometric heterosite FePO4 (274.3 vs 272.0 Å3), caused either by the presence of Fe3+ remaining in the Li (M1) site or by the remaining Li+ that could not be extracted from the structure, as only ∼66 at. % of Fe was originally divalent. On the other hand, the fully discharged material would have, in theory, a composition close to Li1.34Fe1.05P0.92O4, which means that the structure has to accommodate an excess of Li to be located in crystallographic sites other than M1 and M2. The unit-cell volume at the end of discharge is surprisingly close to the one of LiFePO4 (290.8 vs 291.3 Å3). As a consequence, the many singularities in the 4269


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Chemistry of Materials two Fe2+ components represent the two Fe2+ environments in pristine LFP-3h, only one of the two Fe3+ was used to represent the two Fe3+ environments in the original material, whereas the second one was used to represent an additional Fe 3+ environment that appears only after partial delithiation. This new component, shown in red in Figure 9, stands for Fe3+ in Lipoor M2 environment and is characterized by hyperfine parameters similar to those reported in the literature: IS = 0.43 mm·s−1 and QS = 1.47 mm·s−1 (cf. Figure 5b, violet ellipsis). Attempts to use three independent spectral components for the three Fe3+ environments led to inconsistent fitting results, due to strong correlations between the fitted hyperfine parameters. For this reason, it was decided to merge the two original Fe3+ spectral components of pristine LFP-3h into a single one. Moreover, to further reduce possible correlations among the different spectral parameters, the whole series of spectra was treated simultaneously. In this way, both IS and QS of the different components were fitted using common values in all spectra, and only their relative proportions were left free to vary independently in each spectrum. Therefore, the population represented in orange in Figure 9 represents both the Fe3+ in distorted M2 site and the Fe3+ in M1 (antisite), and the resulting quadrupole splitting value (0.89 mm·s−1) is an average between the two former QS (respectively, 1.1 and 0.70 mm·s−1). The width at halfmaximum of the Lorentzian contribution (LW) is also increased, from 0.33 mm·s−1 for each of the two original contributions to 0.55 mm·s−1 for the component including both of them. 3.3.1. Evolution of Divalent Fe during First Charge. In the case of Fe2+, the two environments used for the fitting of the spectra are similar to those of the pristine powder. The population accounting for Fe2+ in the regular M2 site with Lirich environment is marked in blue and possesses regular IS and QS values (IS = 1.21 mm·s−1, QS = 2.92 mm·s−1). However, the hyperfine parameters of the other population (green) are significantly different from those found for pristine LFP-3h. The QS value is 2.23 mm·s−1, compared to 2.64 mm· s−1 found previously. This difference is due to the fact that, similarly to the case of Fe3+, this population contains not only Fe2+ in the distorted M2 site, as in the pristine powder, but also Fe2+ in the M1 site (antisite). Indeed the latter appears only during electrochemical operation, when the electrode is discharged at low potentials. One can conclude that this antisite Fe2+ environment should be characterized independently by a very low QS value. 3.3.2. Evolution of the Trivalent Fe in the First Charge. Prior to cycling, LFP-3h presents, as expected, the same Mössbauer signature as for the ex situ analysis. The red contribution is absent and appears only as the charge proceeds (Figure 9d). This population then grows to represent up to 67 at. % of the total Fe at 3.90 V vs Li+/Li. In parallel, the proportion of orange contribution decreases down to ∼26 at. % at the end of charge. This evolution indicates the relaxation of part of the distorted Fe3+ environment into regular Fe3+ in M2, owing to the partial disappearance of the lattice distortions upon delithiation. The presence of ∼7 at. % of Fe2+ remaining at the end of charge reveals that full oxidation is not reached. This is possibly due to engineering issues with the high electrode material loading of the in situ cell (26.4 mg of powder, mixed with carbon black without binder). Regardless of that, the Li-poor near-heterosite structure at the end of charge still possesses ∼26 at. % of defect-related Fe3+ (orange

contribution), which includes 14 at. % of antisite Fe3+ originally present in the pristine material. This feature may be responsible for the large solid solution domain observed during the electrochemical operation in charge, as it can contribute to reducing the strains/mismatch between the Li-rich and Li-poor domains. 3.3.3. Strong Distortions in Discharge. During the subsequent reduction, a sharp drop in the proportion of the regular Fe3+ in M2 site (red) population occurs, compensated by the growth of the distorted Fe3+ in M2 (orange) population at 3.42 V vs Li+/Li. This is evidence of the reappearance of a very strong lattice distortion during Li insertion, which persists until the end of the discharge, where a negligible or no Fe3+ in regular M2 site is present. In parallel, an increase of the population of regular Fe2+ in M2 (blue) is observed, with no significant change in the proportion of Fe2+ in distorted M2 or M1 (green). This observation is in agreement with previous works on this material34 stating that Fe3+ is more sensitive to lattice distortion than Fe2+, even though the origin of this distortion is due to the second nearest neighbors.34 Below 3.17 V vs Li+/Li, however, the defect-related Fe2+ population (green) increases, as a consequence of the reduction of part of the Fe3+ in the M1 site. The exact quantity cannot be determined precisely though, due to the gathering of two environments in one component assumed earlier. At the end of the discharge, ∼13 at. % of Fe3+ remains. The origin of this is not clear, as it can stem from the absence of appropriate electrode formulation, the impossibility to insert more lithium in the structure, or the presence of structural Fe3+ that cannot react. 3.3.4. Second Charge. The second charge of the electrode unveils two relevant features. At first, the decrease of the defectrelated Fe2+ population (green) occurs at higher potentials than for the discharge. This is in line with the previous electrochemical measurements that show a potential hysteresis for the lower-potential contribution. This would also illustrate a characteristic polarization for the redox reaction of Fe3+/2+ in the M1 site. The second feature is the good reversibility of the electrochemical process, previously observed during galvanostatic measurement and followed in situ by X-ray diffraction. Despite the strong lattice distortions inherent to the Li insertion/removal and the complexity of the system, the two end-of-charge members present a very similar Fe distribution. A subtle difference can be noticed, however, between these two sweeps. The 3.17 V potential step of the second charge presents a small proportion of Fe3+ in the regular Li-poor M2 site, which is not the case at the first measured point (OCV). This feature supports the idea of a solid-solution reaction mechanism, as the redox reaction of the regular Fe in M2 would begin at potentials as low as ∼3.17 V vs Li+/Li. The extrapolation of the dq/dU derivative peak of LFP-3h during electrochemical oxidation in the bottom inset of Figure 7 would also indicate a similar onset potential.

4. CONCLUSION This work brings a new perspective to the comprehension of the direct coprecipitation in aqueous medium of defective olivine LiFePO4 powders and provides evidence of the possibility of a fast and low-temperature preparation of this important material. The increase in the reaction time goes with a progressive ordering of the structure despite the oxidizing reaction atmosphere, and the high concentration of defects present at short reaction times has been thoroughly studied. 4270


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(7) Martin, J. F.; Yamada, A.; Kobayashi, G.; Nishimura, S.; Kanno, R.; Guyomard, D.; Dupré, N. Air Exposure Effect on LiFePO4. Electrochem. Solid-State Lett. 2008, 11, A12. (8) Axmann, P.; Stinner, C.; Wohlfahrt-Mehrens, M.; Mauger, A.; Gendron, F.; Julien, C. Nonstoichiometric LiFePO4: Defects and Related Properties. Chem. Mater. 2009, 21, 1636. (9) Badi, S. P.; Wagemaker, M.; Ellis, B. L.; Singh, D. P.; Borghols, W. J.; Kan, W. H.; Ryan, D.; Mulder, F.; Nazar, L. Direct synthesis of nanocrystalline Li0.90FePO4: Observation of phase segregation of antisite defects on delithiation. J. Mater. Chem. 2011, 21, 10085. (10) Hoang, K.; Johannes, M. Tailoring Native Defects in LiFePO4: Insights from First-Principles Calculations. Chem. Mater. 2011, 23, 3003. (11) Kuss, C.; Liang, G.; Schougaard, S. Atomistic modeling of site exchange defects in lithium iron phosphate and iron phosphate. J. Mater. Chem. 2012, 22, 24889. (12) Jensen, K. M.; Christensen, M.; Gunnlaugsson, H.; Lock, N.; Bøjesen, E.; Proffen, T.; Iversen, B. B. Defects in Hydrothermally Synthesized LiFePO4 and LiFe1−xMnxPO4 Cathode Materials. Chem. Mater. 2013, 25, 2282. (13) Fisher, C. A.; Prieto, V. M.; Islam, M. S. Lithium Battery Materials LiMPO4 (M = Mn, Fe, Co, and Ni): Insights into Defect Association, Transport Mechanisms, and Doping Behavior. Chem. Mater. 2008, 20, 5907. (14) Yang, J.; Tse, T. S. Li Ion Diffusion Mechanisms in LiFePO4: An Ab Initio Molecular Dynamics Study. J. Phys. Chem. A 2011, 115, 13045. (15) Delacourt, C.; Wurm, C.; Reale, P.; Morcrette, M.; Masquelier, C. Low temperature preparation of optimized phosphates for Libattery applications. Solid State Ionics 2004, 173, 113. (16) Delacourt, C.; Poizot, P.; Morcrette, M.; Tarascon, J. M.; Masquelier, C. One-Step Low-Temperature Route for the Preparation of Electrochemically Active LiMnPO4 Powders. Chem. Mater. 2004, 16, 93. (17) Delacourt, C.; Poizot, P.; Masquelier, C. Crystalline Nanometric LiFePO4. WO 2007/000251A1, 2007. (18) Leriche, J. B.; Hamelet, S.; Shu, J.; Morcrette, M.; Masquelier, C.; Ouvrard, G.; Zerrouki, M.; Soudan, P.; Belin, S.; Elkaïm, E.; Baudelet, F. An Electrochemical Cell for Operando Study of Lithium Batteries Using Synchrotron Radiation. J. Electrochem. Soc. 2010, 157, A606. (19) Jumas, J. C.; Stievano, L.; Sougrati, M.; Fullenwarth, M.; Fraisse, B.; Leriche, B. Analysis of the features of an electrode material of an electrochemical cell. WO 2014033402A1, 2014. (20) Grosse, G. PC-Mos II, 1.0 edn; Technische Universität München: Munich, Germany, 1993. (21) Delacourt, C.; Poizot, P.; Levasseur, S.; Masquelier, C. Size Effects on Carbon-Free LiFePO4 Powders. Electrochem. Solid-State Lett. 2006, 9, A352. (22) Jensen, K.; Christensen, M.; Tyrsted, C.; Brummerstedt Iversen, B. Real-time synchrotron powder X-ray diffraction study of the antisite defect formation during sub- and supercritical synthesis of LiFePO4 and LiFe1−xMnxPO4 nanoparticles. J. Appl. Crystallogr. 2011, 44, 287. (23) Yang, S.; Zavalij, P. Y.; Whittingham, M. S. Hydrothermal synthesis of lithium iron phosphate cathodes. Electrochem. Commun. 2001, 3, 505. (24) Yang, S.; Song, Y.; Zavalij, P. Y.; Whittingham, M. S. Reactivity, stability and electrochemical behavior of lithium iron phosphates. Electrochem. Commun. 2002, 4, 239. (25) Porcher, W.; Moreau, P.; Lestriez, B.; Jouanneau, S.; Le Cras, F.; Guyomard, D. Stability of LiFePO4 in water and consequence on the Li battery behaviour. Ionics 2008, 14, 58. (26) Porcher, W.; Moreau, P.; Lestriez, B.; Jouanneau, S.; Guyomard, D. Is LiFePO4 Stable in Water? Electrochem. Solid-State Lett. 2008, 11, A4. (27) Kobayashi, G.; Nishimura, S.; Park, M. S.; Kanno, R.; Yashima, M.; Ida, T.; Yamada, A. Isolation of Solid Solution Phases in SizeControlled LixFePO4 at Room Temperature. Adv. Funct. Mater. 2009, 19, 395.

Strong lattice distortions are evidenced, together with a redistribution of the Li and Fe atoms over the M1 and M2 crystallographic sites and a general decrease in the unit-cell volume. These defects, found in the literature to be very detrimental to the overall electrochemical performances of the electrodes at low concentrations, open the way to new reaction mechanisms offering surprisingly stable performances. This modification triggers the appearance of large/complete solidsolution domains and very strong unit-cell volume variations upon cycling in the case of the air-annealed powders. As pointed out in a recent review paper by Whittingham,72 LiFePO4 is indeed a rather flexible structure as far as site substitution and defect formation are concerned, as witnessed by the direct observations by several independent groups. For instance, Meng and co-workers73 detected the presence of sarcopside Fe1.5PO4 regions within LFP single crystals. Recent alternative approaches include partial substitution of Fe by V70 or Zr74 that is proposed to significantly increase the disorder possible on the lithium lattice and, thus, enable a metastable single-phase LixFePO4.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.5b00470.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The ALISTORE European Research Institute (ERI) is acknowledged for supporting R.A. through an international Ph.D. student scholarship shared between UPJV-LRCS Amiens (France) and NIC Ljubljana (Slovenia). The financial supports from the Ministry of Education, Science and Sport of Slovenia and the Slovenian Research Agency are acknowledged. Stéphane Hamelet and Charles Delacourt are gratefully acknowledged for enlightening discussions. B. Fraisse is gratefully acknowledged for technical help in the setting of the simultaneous XRD−Mössbauer experiments.

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REFERENCES

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