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Novel Complex Stacking of Fully-Ordered Transition Metal Layers in Li4FeSbO6 Materials Eric McCalla, Artem Abakumov, Gwenaelle Rousse, Marine Reynaud, Moulay Tahar Sougrati, Bojan Budic, Abdelfattah Mahmoud, Robert Dominko, Gustaaf Van Tendeloo, Raphael P Hermann, and Jean-Marie Tarascon Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/cm504500a • Publication Date (Web): 12 Feb 2015 Downloaded from http://pubs.acs.org on February 17, 2015

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Chemistry of Materials

Novel Complex Stacking of Fully-Ordered Transition Metal Layers in Li4FeSbO6 Materials

Eric McCallaa,b,c,d, Artem Abakumove, Gwenaelle Roussea,b,c,f, Marine Reynaudg, Moulay Tahar Sougratib,c,h, Bojan Budicd, Abdelfattah Mahmoudi, Robert Dominkod, Gustaaf Van Tendelooe, Raphael P. Hermanni and Jean-Marie Tarascona,b,c*

a

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

75231 Paris Cedex 05, France b

c

ALISTORE-European Research Institute, FR CNRS 3104, 80039 Amiens, France

Réseau sur le Stockage Electrochimique de l’Energie (RS2E), FR CNRS 3459, France

d

National Institute of Chemistry, Hajdrihova 19, SI-1000 Ljubljana, Slovenia

e

EMAT, University of Antwerp, Groenenborgerlaan 171, B-2020, Antwerp, Belgium

f

Sorbonne Universités - UPMC Univ Paris 06, 4 Place Jussieu, F-75005 Paris, France.

g

CIC Energigune, Parque Tecnologico de Alava, Calle Albert Einstein 48, 01510 Miñano

(Álava), Spain h

Institut Charles Gerhardt, CNRS UMR 5253, Université Montpellier 2, 34 095 Montpellier,

France i

Jülich Centre for Neutron Science JCNS and Peter Grünberg Institut PGI, JARA-FIT,

Forschungszentrum Jülich GmbH, D-52425 Jülich, Germany

* Corresponding author: [email protected]

ACS Paragon Plus Environment


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Abstract: As part of a broad project to explore Li4MM’O6 materials (with M and M’ being selected from a wide variety of metals) as positive electrode materials for Li-ion batteries, the structures of Li4FeSbO6 materials with both stoichiometric and slightly lithium-deficient are studied here. For lithium content varying from 3.8 to 4.0 the color changes from yellow to black, and extra superstructure peaks are seen in the XRD patterns. These extra peaks appear as satellites around the four superstructure peaks affected by the stacking of the transition metal atoms. Refinements of both XRD and neutron scattering patterns show a nearly perfect ordering of Li, Fe, and Sb in the transition metal layers of all samples, though these refinements must take the stacking faults into account in order to extract information about the structure of the TM layers. The structure of the most lithium rich sample, where the satellite superstructure peaks are seen, were determined with the help of HRTEM, XRD and neutron scattering. The satellites arise due to a new stacking sequence where not all transition metal layers are identical but instead two slightly different compositions stack in an AABB sequence giving a unit cell that is four times larger than normal for such monoclinic layered materials. The more lithium deficient samples are found to contain metal site vacancies, based on elemental analysis and Mössbauer spectroscopy results. The significant changes in physical properties are attributed to the presence of these vacancies. This study illustrates the great importance of carefully determining the final compositions in these materials as very small differences in compositions may have large impacts on structures and properties.


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Introduction: With the critical role that Li-ion batteries are currently playing in the emerging electric vehicle market, there is an ever-growing need for positive electrode materials with higher energy density. In the search for such material, Li-rich layered oxides have emerged as one of the most promising classes of materials.

In particular, Li-rich Ni-Mn-Co oxides (e.g.

Li1.2Mn0.54Co0.13Ni0.13O2), denoted hereafter NMC, show very high capacities but a high irreversible capacity and significant voltage fade with cycling.1-3 These compounds have a layered rocksalt structure built on an ABCABC stacking of close-packed oxygen atoms and with all metals in octahedral sites. The layered structures have metallic layers (alternating between pure lithium layers and layers with the transition metals and any remaining excess lithium) separated by oxygen layers. For example, in Li-Ni-Mn-O materials, the transition metal layer typically shows ordering of manganese on the √3 x √3 sublattices, but Ni2+ and Li+ have a tendency to occupy the same sites given their similar ionic radii (in fact nickel can occupy as much as 10% of the sites on the lithium layers4). As a result, any superstructure peaks seen in XRD are often rather weak and show broadening attributed to stacking faults. More recently, Sathiya et al.5-7 have shown that the family of compounds with this high capacity is quite large, and has called for the systematic study of materials of the form Li4MM’O6. Materials that in theory have the stoichiometry to allow perfect ordering of lithium on the layers primarily occupied by the transition metal atoms (called “TM layers” here). In particular, Sathiya et al.5 called for studies of pairs of metals where the oxidation states of the M and M’ metals was not just 4+, but pairs averaging to 4+ such as 3+/5+ or 2+/6+ (e.g. Li4NiTeO6 in Ref. 6), which could show quite different behavior worthy of study. This study is one such example where Li4FeSbO6 has Fe3+ and Sb5+ as confirmed by



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Mössbauer spectroscopy. As will be demonstrated here, the transition metal layer ordering is far more extensive than in NMC materials, essentially due to the fact that iron and lithium do not mix on the same sites and this leads to new satellite peaks in the superstructure region of the XRD patterns, never seen before. This warrants careful study such that the structures of the pristine materials will be examined in detail here, while the electrochemical performance, which is rather complex, will be presented extensively in a future article. Figure 1 shows the structure reported in the literature for Li4FeSbO6.8 The structure has Sb partially ordered on the TM layer with some mixing with Fe (typically 25%), represented as α sites in Fig. 1, and Li/Fe/Sb all occupy the other two sublattices randomly (β and γ sites).

This structure, refined in the C2/c monoclinic space group with lattice

parameters a = 5.1706 Å, b = 8.9382 Å, c = 5.1635 Å, and β = 109.49° (Volume = 224.96 Å3), gives an excellent refinement of XRD patterns, excluding the region from 20-33° where peaks due to the superstructure within the honeycomb lattice are seen. Table S.I shows that the metal-oxygen bond lengths for this structure are consistent with typical values from the literature. As such, there is nothing clearly incorrect about the published structure. Five such peaks are typically present: (020), (110), (11-1), (021), (111); clearly illustrated in the scan of sample B10 in Figure 2 to be discussed in the results section. Only the (020) reflection is not affected by stacking faults, while the other four broaden significantly. As Rietveld refinement softwares are currently unable to model stacking faults, these peaks are typically ignored during refinement and as such much information about the local ordering on the TM layers is lost. Here, we utilize a new refinement program, called FAULTS, which enables refinements of experimental X-ray and neutron powder diffraction patterns of layered systems with any type of planar defects, such as stacking faults.9,10


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An interesting aspect of the current study is that satellite peaks around the superstructure peaks affected by stacking faults were seen in certain samples. To the best of our knowledge, there are currently no reports of such satellite peaks in lithium-rich layered oxides. Satellite peaks have been commonly reported in epitaxial thin-films where periodic composition modulations were found to give satellites around the central peaks expected based on the average composition.11

This has led us to consider periodic composition

variation to understand the satellite peaks seen here. Another aspect of this study that is of current interest is that the materials are found to exist over a range of lithium content without changing the transition metal ratios nor changing the oxidation states of the transition metal atoms. This was seen recently in Li-Ni-Mn-O materials where lithium deficiency resulted in the creation of metal site vacancies.12 This situation occurs when the composition of the metals, as determined by elemental analysis, requires more oxygen than metals (i.e. O > Li + Mn + Ni, in that case) in order to charge balance.

In the Li-Ni-Mn-O system, one such vacancy material has been found to be

particularly stable versus the electrolyte at high potentials such that these materials may have interesting properties for some battery applications.13 In the literature, there are no careful studies examining the effect of lithium content in Li4FeSbO6 and in fact no elemental analysis performed on final materials whatsoever. We present both of these aspects here and the consequences are significant.

Experimental section: Synthesis was carried out in two different ways. Synthesis route A was carried out as described in Ref. 8. First, an FeSbO4 precursor was synthesized by mixing Fe2O3 (Alfa Aesar, 99.5%) and Sb2O5 (Alfa Aesar, 99.998%) in a SPEX ball mill for 40 minutes and then heating



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at 800°C overnight. Then stoichiometric amounts of Li2CO3 (Sigma-Aldrich, 99%) and the obtained FeSbO4 precursor were mixed together for 40 minutes in a SPEX ball mill. The mixture was then pelletized and heated at 5 °C/min up to 1000 °C for 15 hours. The sample was then cooled to room temperature by simply turning off the furnace. This yielded sample A0 (synthesis route A, 0% excess lithium). The procedure was repeated with 5, 10, 15, and 20% excess lithium in order to give samples A05, A10, A15 and A20. A second synthesis route was tested to see if the precursor was a necessary part of the synthesis. This route B involved mixing Fe2O3, Sb2O5 and Li2CO3 in the ball mill for 40 minutes. The same thermal treatments were applied as in route A. In this way samples B0, B10, B20 and B30 were synthesized. For all of these samples, the mass of material prepared was roughly 1g. One B0 sample was also heated for 50 hours instead of 15, and is labelled B0*. In order to perform the neutron scattering measurements, larger samples were prepared. As shown in the results section, the materials prepared by route A and B are identical. Thus, the simpler synthesis method, B, was selected to prepare the samples for neutron measurements. These samples, labelled B0N and B20N, had massed of about 5g and as such lost slightly less lithium during synthesis, as confirmed by elemental analysis in Table I. X-ray diffraction (XRD) was performed in the laboratory using a Bruker D8 advance equipped with a copper target X-ray tube and a LynxEye XE detector. A single sample A20 was also sent to Argonne National Laboratories to measure the scattering on the Synchrotron beamline 11-BM, using the mail-in service with λ = 0.413843 Å. Neutron scattering was performed at the Paul Scherrer Institute at SINQ on the HRPT beamline. A wavelength of 1.494 Å was used. FullProf14 and Rietica were both used for Rietveld refinement of the various XRD and neutron scattering patterns.

Neutron scattering complements XRD

measurements nicely for this system as the scattering lengths (Li: -1.90 fm, Sb: 5.57 fm, Fe: 9.45 fm, O: 5.80 fm) result in a greater sensitivity to oxygen and lithium than in XRD where


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Fe and Sb completely dominate. In order to refine the structures of the materials with stacking faults, the new FAULTS program was used as described below.9,10 This program was also used to model various stacking sequences in order to refine the structure of the samples made with 20% excess lithium as discussed extensively in the results section. Since three different software packages were used for refinement, it was necessary to calculate quality factors for fits manually to ensure all weightings are consistent for all models to allow comparison between the various models. The factors chosen were RP and R1P (RWP with all points weighted with unity). The FAULTS program requires some introduction, as it involves describing structures as being made up of a set of layers, and not a stacking of unit cells.15 This software is based on the DIFFaX code, which has been used effectively to model stacking faults in such systems in the past (e.g. Ref. 16-20), and on several modules of the CrysFML Fortran 95 crystallographic library for the refinement of the structural model and for isotropic size broadening treatment.21,22 FAULTS refinements require modeling the structure as being made up of a set of layers following a particular stacking sequence. Thus, to define such layers we had to describe the structure in a new orthogonal set of basis vectors a, b, and cF as illustrated in Figure 1. Figure 1 also shows the three sites representing each of the three √3 x √3 sublattices (labelled α, β, and γ). Two atoms occupy each sublattice within the unit cell (in all refinements, the positions of the two sites were locked together such that each sublattice could not be destroyed). Thus, in FAULTS, two layers are considered: the MO2 (M represents the metals on the T.M. layer) slabs form one layer while the lithium layer makes up the other. Here, the monoclinic distortion is taken into account by how the layers are stacked using the vectors S1 and S2. S1 therefore places a lithium layer above an MO2 layer and S2 does the opposite. For a perfect hexagonal layered material, such as for LiCoO2, the stacking vectors would be S1 = S2 = (-1/6, 0, 0.5). Here, for monoclinic structures, a perfect material without



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stacking faults would have stacking vectors (-Sx, 0, 0.5). Refining the value Sx therefore allows for monoclinic distortions in the material. To model the stacking faults, a simple approach is taken. For all cases, S1 = (-Sx, 0, 0.5), i.e. the lithium layers are not affected by the stacking faults. Furthermore, S2 = (-Sx, 0, 0.5) occurs with a probability 1-P corresponds to the fraction of layers that stack perfectly while S2 = (-Sx, 1/3, 0.5) and S2 = (-Sx, -1/3, 0.5) represents stacking which aligns each of the other two sublattices above the first one along the c-axis and these each occur with probability P/2.

In this simple model, P is then the

probability of two TM layers being separated by a stacking fault. The A20 sample was also investigated with transmission electron microscopy. The sample was handled in an Ar-filled glovebox and crushed in a mortar with hexane. Drop of the suspension was then deposited onto holey carbon grid and the grid was transported to the microscope without contact with air. Electron diffraction patterns (ED), high resolution transmission electron microscopy (HRTEM), and high angle annular dark field scanning TEM (HAADF-STEM) images were obtained using a Tecnai G2 electron microscope operated at 200kV. 57

Fe and

121

Sb Mössbauer spectra were collected in order to determine the oxidation

states of the metals.

The Fe spectra were collected using a

57

Co(Rh) source at room

temperature. The Sb spectra were collected at various temperatures using a close-cycle cryostat (Janis SH-850), and a Ca119mSnO3 source at room temperature was used. All spectra were recorded in the transmission geometry and were fitted using Lorentzian peak shapes. Isomer shifts are reported relative to alpha-iron and CaSnO3, for iron-57 and antimony-121 Mössbauer spectra, respectively. Elemental analysis was performed on selected samples using Inductively Coupled Plasma – Optical Emission Spectroscopy (ICP-OES). Samples were first consumed in Aqua Regia then analyzed for lithium, iron and antimony. The combination of


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elemental analysis and oxidation states from Mössbauer spectroscopy allowed for the calculation of oxygen content based on charge balancing. Electrochemical cycling was performed in Swagelok style cells where the working electrode was a mixture of 80% active material and 20% SP carbon black, previously ball milled for 15 minutes under argon. The separator was a Whatman GF/D borosilicate glassy fiber sheet saturated with LP-30 electrolyte (Merck KGaA).

Galvanostatic cycling was

performed at room temperature between 2.5 and 4.6 V at a rate of C/10 (C-rate here is defined as the rate needed to remove 1 Li from Li4FeSbO6 in one hour).

Results and Discussion Figure 2 shows the XRD patterns for samples B0, B10, B20 and B30 obtained through all solid-state synthesis. The first three are single-phase and will be discussed extensively below, while the sample made with 30% excess lithium shows extra peaks due to a contaminant phase (most likely another layered lithium rich oxide). The pattern of B0 is very similar to that obtained by Zvereva et al.8 where the first superstructure peak (020) at 20° is sharp while the others all show broadening due to stacking faults. It should be noted that very small contaminant peaks are visible here. Such peaks were not seen in the A0 phase, and this constitutes the only differences seen between the two synthesis methods. Furthermore, the peak at 43° (indexed by the reflections (-202) and (131)) appears to be a single peak with a slight asymmetry attributed to monoclinic distortions. As the lithium content increases from B0 to B30, the superstructure peaks all sharpen and the monoclinic distortion becomes far more severe. For 20% excess lithium (B20), all superstructure peaks are equally sharp suggesting there are no stacking faults, but each of the four peaks associated with the stacking of the layers now shows a pair of satellite peaks. This implies that the layers do not stack



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perfectly as illustrated in Figure 1 but instead a new stacking sequence is seen. Figures 3 and S.1 shows the same XRD patterns are obtained with the other synthesis route with the FeSbO4 precursor. The satellite peaks can be clearly seen in both the 15 and 20% excess lithium materials and to a much lesser degree in the 10% excess samples, A10 and B10, which show the first onset of this extra ordering. The monoclinic distortion at 43° now resolves into three distinct peaks for these two most lithium rich samples. ICP results for the concentrations of Li, Fe and Sb in the compounds are given in Table I. It should be noted that 20% excess lithium is needed to reach the stoichiometric composition Li4FeSbO6. Figure S.2 shows the 57Fe and

121

Sb Mössbauer spectra for the A0,

A10 and A20 samples while Table S.II and S.III show the results for the fitting parameters. In all cases, Fe is found to be in the 3+ oxidation state (isomer shift of 0.33 mm/s is typical for Fe3+ in octahedral oxygen environments). In the more lithium deficient sample, A0, there are two possibilities to charge balance given the lithium shortage: (1) one of the metals is in a lower oxidation state (since Fe Mössbauer clearly shows Fe3+, the only option is that some Sb is reduced from its more common 5+ oxidation state down to 3+), or (2) the oxygen content is non-stoichiometric (i.e. there is more oxygen than total metal atoms). To decide between these two possibilities, Sb Mössbauer measurements were performed on the A0 sample. Figure S.2(b) and Table S.III clearly show that Sb is still in the 5+ state. As such, in Table I, vacancy and oxygen contents were calculated by charge balancing with Li+, Fe3+ and Sb5+. In all cases, there is a slight Sb excess as compared to Fe (average ratio is 0.52:0.48, instead of 0.5:0.5 as mixed initially). The most likely source for this imbalance is a small amount of iron loss as it reacts with the alumina crucible during synthesis. More importantly, the ICP results show that these materials do not have excess lithium as compared to the target Li4FeSbO6, but in fact only the materials made with 20% excess lithium attain this concentration. Thus, during synthesis the material saturates at a lithium concentration of 4. For samples with


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less than 20% excess lithium, the samples end up slightly lithium deficient giving rise to metal site vacancies, as discussed in the introduction. The sample made by heating for 50 h, B0*, shows the smallest amount of lithium and as such the highest concentration of vacancies, 5.8 % of all metal sites. Though not shown here, the XRD pattern looks identical to B0, with slightly different lattice parameters as discussed below. It should also be noted that these values for vacancy concentrations are all below the maximum value of 8.3% obtained for LiNi-Mn-O materials.12 Figure 4 shows the electrochemical data for, and images of, each of the B0, B10 and B20 samples. The electrochemical curves each show the same features: an oxidation plateau at 4.2 V, some activity above 4.2 V which becomes more significant with cycling for B0 and B10, a discharge plateau at 4.0 V which shrinks with cycling and a second below 2.6 V which grows with cycling. Clearly the cycling for B20 is different from the other two, particularly above 4.2 V on the first charge, but the main features are all identical it is only the relative amounts of each plateau that varies, and in all three cases the later cycles look quite similar. The electrochemical performance of these materials involves Fe oxidation, peroxo-group formation, oxygen gas release and the materials reacting with electrolyte at low potential. As such, it has proven quite complex to explain, and will be the subject of an upcoming paper. Here, it is sufficient to demonstrate that there are significant variations in the properties of these materials even though the changes in composition are very minor. Insets in Fig. 4 also show a dramatic color change from a bright yellow for the material with 0% excess lithium (contains 5% vacancies on TM layer) to a dark black for the 20% excess lithium material which contains no vacancies. The material made with more lithium is also much harder and very difficult to grind manually in a mortar. Clearly, the physical and electrochemical properties of the material change significantly with the very small compositional change seen with ICP. Thus, a detailed study of the structural changes with composition is justified.



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The XRD patterns were refined both using Rietveld refinement (with the model from Ref. 8 allowing for mixing of Sb/Fe on α sites and Li/Fe/Sb on the other two) using Rietica and with a model taking stacking faults into account with the FAULTS program where each of the three sublattice sites were distinct so could be ordered or not. The overall compositions were set to match the ICP results (or the extrapolated values for the 5, 15% samples). Figure 3 shows the results of these refinements. Clearly, the model with stacking faults yields a far better fit for the samples with 0-10% excess lithium, as expected. This is confirmed with the RP values reported in Table II where FAULTS give values in the range 11 – 15% while Rietveld refinement (without taking stacking faults into account) gives values of 25 – 30 %. However, for 15 and 20% the fit is not as good as that obtained with Rietveld refinement, showing that the FAULTS program has a tendency to over-estimate the stacking faults for systems with a low concentration of faults. The resulting lattice parameters are plotted in Figure 5. Once again, there is nothing to distinguish synthesis route A and B here, and from now on no reference will be made to which route was used as it has no influence on the final structure. It is of interest that both a and c increase with excess lithium while b decreases, suggesting that there is a rearrangement of atoms taking place and not simply a random substitution of lithium replacing transition metal atoms. It is also of note that the lattice parameters for 20 and 30% match each other well, confirming that the maximum lithium content is in fact obtained for 20% and that the 30 % material was made up of this saturated composition co-existing with another phase. The trend in the β angle also shows that the monoclinic distortion increases as lithium content increases. Figure S.3 shows the Sx and cF parameters obtained with FAULTS (note, the a and b values were identical to those obtained with Rietveld refinement). The cF values show that the spacing between layers stays roughly constant; it is only the shearing between layers, Sx, that increases as the monoclinic distortions become more severe. The lattice parameters for the B0* sample (heated for 50 h, lowest Li


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content value of 3.55) are a = 5.167(7)Å, b = 8.959(1) Å, c = 5.168(3) Å, and β = 109.41(1)°, or for the layered notation used here: cF = 4.875 Å and Sx = 0.1662 showing stacking very close to that of perfect O3 hexagonal structures. Thus, for all lattice parameters, the values for B0* are qualitatively correct based on the trends in Figure 5. Thus, samples B0 and A0 should not be considered the endpoint of the single-phase region. Table II shows the parameter values resulting from fitting with stacking faults taken into account. As expected from visual inspection of the XRD patterns, the stacking fault percentages steadily diminish as lithium content increases.

Of greater significance, the

program always converged to an Sb occupancy of 1 on the α-sites and a very high degree of ordering of Fe and Li on the β and γ sites, respectively. This is in stark contrast to the typical results obtained with Rietveld refinement performed ignoring the superstructure peaks where typically 25% Fe is seen on the α site and the other two sites are randomly occupied. This is not surprising, ignoring the superstructure peaks results in much of the information regarding nearest neighbors on the transition metal layers being ignored. As such, refining while taking the stacking faults into account does not only allow for the determination of the concentration of the stacking faults, but it also allows for an improvement in the determination of the structure on the transition metal layers. Here, the fact that Li+, Fe3+ and Sb5+ all want to order is not surprising as Sb-Sb and Fe-Fe nearest neighbors are energetically costly based on Coulombic interactions which have been demonstrated to be the driving force in the ordering on the transition metal layers in other Li-rich oxides.23

This complete ordering on the

transition metal layer was never determined before, and certainly not from XRD patterns alone. Table II shows a correlation between stacking faults and the vacancy concentration on the γ-sites. It is therefore feasible that the disorder caused by the vacancies results in the stacking faults between the layers. Thus, in the 20% materials where there are no vacancies, the XRD patterns show no evidence for stacking faults.



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Figure 6 shows the neutron scattering pattern of the samples made with 0 and 20% excess lithium. The B0N sample was once again refined both with and without taking faults into account. Figure S.4 shows two fits obtained with Rietveld refinement ignoring the superstructure peaks. The first has Sb predominantly on the 2a sites as suggested by Zvereva et al.8 and it results in very weak superstructure peaks. The second has Fe predominantly on the 2a sites and is a far better fit. Therefore, to fit the XRD patterns well Sb must be ordered on the TM layer, and to fit the neutron data well Fe must be ordered. This implies that all three must be ordered. This is in fact the result obtained when the pattern is fit using the FAULTS program allowing for ordering of all three sites on the TM layer. Table II shows that there is negligible mixing between the sites, with only 0.2% lithium on the β-site. The structure obtained by fitting the neutron data with faults is given in Table S.IV and is presented as the structure for the lithium deficient materials where stacking faults are present in significant concentrations.

Though some distortions were allowed, the metal atoms

remained very near to the ideal positions. Finally, let us discuss the structure of the materials made with 20% excess lithium. The structure was determined using a combination of XRD, TEM and neutron scattering. Figure 7 shows the ED patterns for the A20 material. The [001] pattern is typical for a Li2MO3 like material.24 By contrast, the [110] pattern shows the superstructure peaks that can be indexed with a propagation vector q = 0.25c*, implying a c-axis that is 4 times larger than that of the regular unit cell discussed up to now. It should be noted that the brightest satellite peaks, spaced by 4q, attributed to the reflections arising from the honeycomb cation ordering on the transition metal layers in typical Li-rich materials22 that do not show the satellite peaks in the XRD patterns. Thus, the ordering is not due to the changing of the stacking sequence of the oxygen close packed layers, but solely to the occupational ordering within the transition metal layers.


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Two models have been generated which both show the satellite peaks in the XRD patterns. Figure 8 shows these two models side-by-side. The first, called the alternating layers model, can be described in the C2/c space group (Table S.V shows the atomic coordinates for this model). This model involves three different transition metal layers with Fe, Sb and Li all ordered on the honeycomb lattice on layer L1.

L1’ has this same

composition but Fe and Li exchange positions. Finally, the L2 layers have Sb ordered and the other two sites randomly occupied. The layers are stacked such that Sb stacks perfectly along the c-axis. The satellite peaks then arise from the Li and Fe inversion between the L1 and L1’ layers (ultimately it is the presence of the L1’ layers that gives the 4c periodicity seen in the TEM). Figure 9 shows the Synchrotron XRD pattern from sample A20 along with the fit obtained from this alternating layers model using FullProf. All peaks index very well and the refinement gives excellent values for the quality parameters RP = 13.1% and RWP = 15.8%. The refined lattice parameters are a = 5.17596(7)Å, b = 8.9289(1) Å, c = 20.6832(3) Å, and β = 109.878(1)°, consistent with results from Figure 5 with a four times larger c value. Despite the success of the alternating layers model in fitting the XRD pattern, it is important to determine whether or not the satellite peaks can be obtained from the stacking of perfectly ordered layers only and not a combination of ordered and partially ordered layers. This is particularly important since the current study has already demonstrated perfect ordering in the layers with lithium deficiency, such that it would be surprising to obtain less ordering without the deficiencies. To explore the different ways to stack the ordered layers, the FAULTS program was used once again. Assuming Sb continues to stack perfectly as in the alternating layers model, we need now only consider the L1 and L1’ layers. There are therefore only two possible sequences giving a 4c lattice parameter: L1-L1-L1-L1’ and L1L1L1’L1’. Figure S.5 shows that only the second sequence gives the correct satellite peaks in the XRD pattern, such that the second model explored here, and shown in Figure 8 is based



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on an L1-L1-L1’-L1’ (or AABB) stacking of the transition metal layers and is called the paired layers model. This can be indexed in the C2 space group (with same lattice parameters as the previous C2/c one), and Table S.VI shows the atomic coordinates obtained for this model. Refinements of XRD patterns proved insufficient to distinguish the two models. As neutron scattering is far more sensitive to lithium than XRD, it proved to be able to determine which model describes the structure the best. Figure 6 and Figure S.4 shows the neutron scattering pattern obtained for sample B20N with the fits obtained for each model. Overall, both fits are good as indicated by the quality factor values in Table III. Both models give lattice parameters of a = 5.18061(8) Å, b = 8.92649(13) Å, c = 20.7228(3) Å, β = 109.9901(9)°. However, zooming in on the superstructure region in Figure 10 clearly shows an improvement in the fit for the paired layers model and this is reflected in a significant reduction in the quality factors over the restricted range 18.5-33° given in Table III. Comparing the XRD pattern to the neutron pattern shows how neutron scattering is far more sensitive to the satellite peaks than XRD is and so it is not surprising that it was required to distinguish the two models proposed here. Once again, the results show that Li, Fe and Sb order completely on the transition metal layers as was seen in the lithium deficiency sample B0N. The refined values for the occupations on the L1 layer are Li, Fe, Sb on α, β, and γ respectively, while on L1’ they are 0.98Sb/0.02Fe, 0.858Fe/0.142Li, 0.998Li/0.002Fe (i.e. all attempts to mix Sb and Fe converged strictly to zero and no mixing on L1 was found to improve the fit). This strongly suggests that the stacking change between the L1 and L1’ layers arises from the disorder caused by the Li on the primarily Fe-occupied β site on the L1’ layers. Figure 11 shows the structure obtained from fitting the neutron pattern and the corresponding atomic coordinates are given in Table S.VI. Table S.I also shows that the metal-oxygen bond lengths for this new structure are once again consistent with values


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reported in the literature for similar structures. This structure is presented here as the best model based on the current data. Interestingly, the lithium layers that lie between the L1 and L1’ layers are distorted while those lying between two identical layers are not. The crystal structure of the A20 sample was verified with the HAADF-STEM and HRTEM imaging. The [110] HAADF-STEM image (Fig. S6 of Supporting information) reveals perfect stacking of the honeycomb-like layers, which are only occasionally violated by rare stacking faults. No systematic difference in the brightness of the dots of the projected cationic columns was observed indicating that the occupational ordering only weakly alters the projected potential along this viewing direction. Contrarily to that, the [110] HRTEM image in Fig. 12, taken from the same crystallite, clearly shows the 4c superstructure visible as a pattern of alternating pairs of brighter and darker dot rows (clearly shown in the left half of Fig. 12). According to the calculated HRTEM images, the dots correspond to the regions of low scattering potential in the crystal. Thus, the atomic displacements associated with the occupational ordering primarily impact the superstructure contrast in the HRTEM image. The HRTEM image calculated with the paired layers model indexed in C2 is in good correspondence with the experimental one, correctly reproducing the patterns of alternating pairs of blurred and brighter rows. In contrast to that, the alternating layers model indexed in C2/c model, although reproducing the main motif of the image, shows alternation of a single blurred row and three brighter rows. This discrepancy provides additional argument in favor of the paired layers model, consistent with the neutron scattering results. Though it should be pointed out that this effect is quite weak, and in fact the patterns of bright and blurred rows cannot be discerned at all in the right side of Fig. 12. As such, the current study demonstrates how vital neutron scattering is in determining the ordering in such Li-rich oxides as it was the only technique able to clearly differentiate between the two models.



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Conclusion: A solid-solution region was observed near Li4FeSbO6 where lithium deficiency was accommodated for by the creation of metal site vacancies. Though the range of this solid solution was small (from lithium content of 3.55 to 4.07), the properties varied dramatically with a color change from bright yellow for a lithium content of 3.8 to black for 4.05, which was accompanied with an increase in hardness. Structurally, all samples were found to be made up of transition metal layers where Fe, Sb and Li all order perfectly on the honeycomb lattice; which is in contrast to previously published structures. The previous results were based on refinement in space groups that did not distinguish between all three √3 x √3 sublattices, and superstructure peaks where stacking faults gave broadening were ignored.8 Here, XRD and neutron patterns were fit taking stacking faults into account, and the results demonstrate the usefulness of this approach such that the extent of ordering in the layers can be determined even from laboratory XRD patterns. In the lithium deficient materials where defects in the form of metal site vacancies were seen, the resulting scattering patterns showed stacking faults occurring with probabilities as high as 18.4%. The color change seen in these materials may be due to the changing concentration of metal site vacancies (from roughly 10% of the transition metal layer to 0% in the sample with highest lithium content). Such a small quantity of defects can have significant impact on the optical properties of the material. In the more lithium-rich sample, extra satellite peaks were seen around the superstructure peaks in the XRD patterns. This was determined to be due to a new stacking sequence of pairs of ordered transition metal layers (2 with perfect stoichiometry, 2 with slight lithium excess) stacking in an AABB sequence resulting in a unit cell four times larger than that seen previously. The current study is presented to the lithium-ion battery research community in part as a cautionary tale. The structures and properties of these layered materials vary significantly


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with minor changes in composition such that the importance of knowing the final composition of the materials after synthesis cannot be over-stated.

The XRD patterns alone cannot

confirm that the expected phase is obtained, elemental analysis must be performed. The consequences of such small changes in composition must then be studied with very great care, as illustrated here with the combination of TEM, XRD and neutron scattering. The fact that large amounts of excess lithium were used in this study and yet the final products were in fact only stoichiometric or slightly lithium deficient proves the necessity of carefully determining the compositions of the final products before making any structural determinations.

Acknowledgments: EM thanks the Fonds de recherche du Québec – Nature et Technologies and ALISTORE – European Research Institute for funding this work, as well as the European community I3 networks for funding the neutron scattering research trip. The authors also thank RS2E for funding this research and V. Pomjakushin for his help in performing the neutron experiments. AM thanks the Forschungszentrum Juelich for an incoming international postdoctoral grant.

Supporting Information Available: The supporting information referred to throughout the text is available free of charge via the Internet at http://pubs.acs.org/.

References: (1) Lu, Z.; Beaulieu, L. Y.; Donaberger, R. A.; Thomas, C. L.; Dahn, J. R. J. Electrochem. Soc. 2002, 149, A778. (2) Thackeray, M. M.; Johnson, C. S.; Vaughey, J. T.; Hackney, S. A. J. Mater. Chem.



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2004, 15, 2257. (3) Koga, H.; Croguennec, L.; M/n/trier, M. ; Douhil, K. ; Belin, S. ; Bourgeois, L. ; Suard, E. ; Weill, F. ; Delmas, C. J. Electrochem. Soc. 2013, 160, A786. (4) McCalla, E.; Li, J.; Rowe, A.W.; Dahn, J.R. J. Electrochem. Soc. 2014, 161, A606. (5) Sathiya, M.; Rousse, G.; Ramesha, K.; Laisa, C.P.; Vezin, H.; Sougrati, M.T.; Doublet, M.-L.; Foix, D.; Gonbeau, D.; Walker, W.; Prakash, A.S.; Ben Hassine, M.; Dupont, L.; Tarascon, J.-M. Nature Materials 2013, 12, 827. (6) Sathiya, M.; Ramesha, K.; Rousse, G.; Foix, D. ; Gonbeau, D. ; Guruprakash, K.; Prakash, A.S.; Doublet, M.-L.; Tarascon, J.-M. Chem. Commun. 2013, 49, 11376. (7) Sathiya, M.; Ramesha, K.; Rousse, G.; Foix, D. ; Gonbeau, D. ; Prakash, A.S. ; Doublet, M.-L.; Hemalatha, K.; Tarascon, J.-M. Chem. Mater. 2013, 25, 1121. (8) Zvereva, E.A.; Savelieva, O.A.; Titov, Y.D.; Evstigneeva, M.A.; Nalbandyan, V.B.; Kao, C.N.; Lin, J.-Y.; Presniakov, I.A.; Sobolev, A.V.; Ibragimov, S.A.; Abdel-Hafiez, M.; Krupskaya, Y.; Jähne, C.; Tan, G.; Klingeler, R.; Büchner, B.; Vasiliev, A.N. Dalton Trans. 2013, 42, 1550. (9) Casas-Cabanas, M.; Rodríguez-Carvajal, J.; Palacín, M.R. Zeitschrift für Kristallographie, Supplement 2006, 23, 243. (10) FAULTS program is distributed within the FullProf Suite, available at http://www.ill.eu/sites/fullprof/index.html (11) Shinn, M.; Hultman, L.; Barnett, S.A. J. Mater. Res. 1991, 7, 901. (12) McCalla, E.; Rowe, A.W.; Camardese, J.; Dahn, J.R. Chem. Mater. 2013, 25, 2716. (13) Rowe, A.W.; Camardese, J.; McCalla, E.; Dahn, J.R. J. Electrochem. Soc. 2014, 161, A1189. (14) Rodriguez-Carvajal, J. http://www.ill.eu/sites/fullprof/index.html

FullProf

Suite,

available

at:

(15) Casas-Cabanas, M.; Reynaud, M.; Rikarte Ormazabal, J.; Rodríguez Carvajal, J. Submitted to J. Applied Crystallogr. 2015. (16) Bréger, J.; Jiang, M.; Dupré, N.; Meng, Y.S.; Shao-Horn, Y.; Ceder, G.; Grey, C.P. J. Sol. State Chem. 2005, 178, 2575. (17) Ramesh, T. N. Mater. Chem, Phys, 2009, 114, 618. (18) Casas-Cabanas, M. ; Rodríguez-Carvajal, J. ; Canales-Vázquez, J. ; Laligant, Y.; Lacorre, P.; Palacín, M. R. J, Power Sources, 2007, 174, 414.


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(19) Carlier, D., Croguennec, L., Ceder, G., Ménétrier, M., Shao-Horn, Y., & Delmas, C. Inorg, chem, 2004, 43, 914. (20) Treacy, M.M.; Newsam, J.M.; Deem, M.W. Proc. R. Soc. London Ser. 1991, 433, 499520. (21) Rodriguez-Carvajal, J.; Gonzalez-Platas, J. IUCr CompComm Newsletter, 2003, 1, 50. (22) CrysFML repository: http://forge.epn-campus.eu/projects/crysfml/repository (23) McCalla, E.; Lowartz,C.M.; Brown, C.R.; Dahn, J.R. Chem. Mater. 2013, 25, 912. (24) Boulineau, A.; Croguennec, L.; Delmas, C. ; Weill, F. Chem. Mater. 2009, 21, 4216. (25) Skakle, J.M.S.; Castellanos R.; Tovar, S.T.; Fray, S.M.; West, A.R. J. Mater. Chem. 1996, 6, 1939. (26) Kolitsch, U.; Slade, P.G.; Tiekink, E.R.T.; Pring, A. Mineral. Mag. 1999, 63, 17. (27) Menil, F. J. Phys. Chem. Solids 1985, 46, 763. (28) Lippens, P.E. Solid State Comm. 2000, 113, 399.



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Figure captions: Figure 1: Structural model used previously in the literature. (a) The stacking of the TM and lithium layers is shown while (b) shows the composition of a TM layer with the three sublattices defined by the α, β and γ sites. Parameters required to describe stacking faults, cF, S1 and S2, are also shown. Figure 2: XRD patterns for samples made from all solid-state synthesis with 0, 10, 20 and 30% excess lithium.

indicates a contaminant phase present when too much excess lithium is

used, while * indicates the satellite peaks found around four superstructure peaks. Figure 3: XRD patterns from samples A0 to A20 along with fits obtained from FAULTS taking stacking faults into account in red, and Rietveld fits ignoring stacking faults in blue (note that the blue lines are only included in the left panel). Figure 4: Electrochemical cycling data for B0, B10 and B20 samples, along with images for each. The first cycle is shown in red. Figure 5: Monoclinic lattice parameters as a function of excess lithium used during synthesis. Red is synthesis route A with FeSbO4 precursor, while blue is route B. The blue shaded region indicates the sample was no longer single phase, though it was refined as single-phase. Figure 6: Neutron scattering data for the 0% (B0, top) and 20% excess lithium (B20, bottom) samples. In the top panel the fit is obtained using FAULTS. In the lower panel, the refinement is obtained using FullProf with the paired layers model. Figure 7: Electron diffraction patterns of Li4FeSbO6 (sample A20) indexed with the 4c supercell.


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Figure 8: Structural models used to describe the materials made with 20% excess lithium. There are three different TM layers: L1 is completely ordered, as is L1’ though with Li and Fe interchanged as compared to L1, and L2 allows for mixing of Li and Fe on two sublattices. In both models, the c-axis is four times larger than that normally used to describe layered oxides in the C2/m space group. Figure 9: Synchrotron pattern for sample A20 along with the Rietveld fit obtained using the alternating layers model. The insert shows the superstructure peaks in greater detail. Figure 10: Neutron scattering (λ = 1.494Å) data (top) and XRD (λCu ≈ 1.5418Å) data (bottom) for materials made with 20% excess lithium. Left: fits obtained with the paired layers model, right: fits obtained with the alternating layers model. All fits were obtained using FullProf. Figure 11: Final structure for B20N sample obtained by refining neutron data in the C2 space group. Figure 12: [110] HRTEM image of Li4FeSbO6. The inserts show the theoretical HRTEM images calculated using the paired layers model in C2 and the alternating layers model in C2/c (defocus f = -26 nm, thickness t = 2.3 nm).



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Table I : ICP-OES results for a number of samples discussed in the text. A involves the FeSbO4 precursor, while B is all solid-state, 0, 10, 20 refers to the percentage of excess lithium used during synthesis while * indicates the sample was heated 50 hours. Vacancies and oxygen content were calculated using charge balancing with Li+, Fe3+ and Sb5+. B0N and B20N are sample batches specially prepared for the neutron measurements. Sample A0 A10 A20 B0 B10 B20 B0* B0N B20N

Li (±0.20) 3.81 3.86 4.03 3.86 3.93 4.05 3.55 3.92 4.07

Fe (±0.05) 0.97 0.97 0.96 0.97 0.96 0.94 1.00 0.98 0.94

Sb (±0.05) 1.06 1.05 1.01 1.05 1.04 1.01 1.10 1.03 0.99



O

0.16 0.12 0 0.13 0.07 0 0.35 0.07 0

6 6 5.98 6 6 5.96 6 6 5.92


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Table II: Structural results obtained by fitting XRD patterns from Fig. 3 and the neutron pattern of sample B20N from Fig. 10. Refinements take stacking faults into account, except for the Rietveld refinements. For each of A0, A05 and A10 the mixing of Li and Fe converged strictly to zero. In all cases, the mixing of Sb and Fe on the α site converged strictly to zero. The R-values for the Rietveld fit were obtained using the structural model from Ref. 8. Site Occupations β γ

Sample Lab XRD A0

Sb

Fe 0.94 Sb 0.06

A05

Sb

A10

Faults RP R1P

Rietveld RP R1P

Sx

P (%)

Li 0.81 Fe 0.03  0.16

0.1685

12.7

11.4

9.1

29.7

18.4

Fe 0.95 Sb 0.05

Li 0.83 Fe 0.02  0.15

0.1683

12.4

11.8

8.9

29.5

16.7

Sb

Fe 0.95 Sb 0.05

Li 0.86 Fe 0.02  0.12

0.1692

4.5

15.5

12.3

24.9

19.7

A15

Sb

Fe 0.96 Sb 0.03 Li 0.01

Li 0.99 Fe 0.01

0.1696

1.3

15.5

12.7

15.8

10.2

A20

Sb

Fe 0.80 Li 0.19 Sb 0.01

Li 0.84 Fe 0.16

0.1708

0.8

18.4

14.7

18.7

12.8

Neutron B0N

Sb

Fe 0.978 Sb 0.02 Li 0.002

Li 0.918 Sb 0.01 Fe 0.002

0.1681

18.2

6.7

8.2

7.4* 7.1**

9.7* 8.9**

α

* Fit with Sb predominantly on 2a site **Fit with Fe predominantly on 2a site



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Page 26 of 39

Table III: Quality factors for fits of neutron refinement of sample B20N, over either the entire range, or limited to the range where superstructure peaks are seen (18.5 – 33.0°). Model Alternating layers Paired layers

5 - 160° RP R1P 5.5 6.7

18.5 – 33.0° RP R1P 8.3 12.5

5.2

6.3

6.8


4.8

Page 27 of 39

(a)

(b)

S2 c a

α

α β

α β

γ

α b

b

Antimony

γ

β γ

cF

S1

Lithium

α

chex

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Iron

Figure 1.


a c

Oxygen


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** * * * * * * B30

B20

B10

B0

Figure 2.


A20

A15

A10

A05

A0

Figure 3.


(022/220) (202) (131)

(111)

(021)

(111)

(110)

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(020)

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Voltage (V)

B0 4

3

2

Voltage (V)

B10 4

3

2

B20 Voltage (V)

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4

3

2 1.5

2

2.5

3

3.5

x in Lix(Fe0.96Sb1.04)yO6 Figure 4.


4

Page 31 of 39

a (Å)

5.185 5.180 5.175 5.170 8.945

b (Å)

8.940 8.935 8.930 8.925 5.185

c (Å)

5.180 5.175 5.170 5.165

110.0

β (deg)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


109.9 109.8 109.7 109.6

0

5

10

15

20

25

30

Excess lithium (%) Figure 5.



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B0 Neutron: Faults fit

B20 Neutron: Paired Layers Model

Figure 6.


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Figure 7.



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Page 34 of 39

Alternating layers model:

Paired layers model: L1

L2

L1’

L1’

L1’

L2

L1

L1 c

L2

a

TM layer L1

b

c

L1

a

TM layer L1’

b

TM layer L2

a b

Lithium

Antimony

Iron

Figure 8.


c

Oxygen

Page 35 of 39

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Figure 9.



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B20 Neutron – Paired Layers Model

A20 XRD – Paired Layers Model

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B20 Neutron – Alternating Layers Model

A20 XRD – Alternating Layers Model

Figure 10.


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L1’

L1’

L1

L1 L1

L1’

L1

L1’

Figure 11.



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Figure 12.


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Graphical table of contents:

Scattering Angle (deg.)


Novel Complex Stacking of Fully-Ordered ... - ACS Publications

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