High voltage sulphate cathodes Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co): atomic-scale studies of lithium diffusion, surfaces and voltage trends

The search for high voltage cathodes for lithium-ion batteries has led to recent interest in the monoclinic Li 2 Fe(SO 4 ) 2 material which has a voltage of 3.83 V vs. lithium, the highest recorded for a ﬂ uorine-free iron-based compound. Here we investigate the defect, surface and lithium migration properties of the Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co) materials using combined atomistic modelling and density functional theory (DFT) techniques. All intrinsic defect types including Li/M antisite disorder are found to be of high energy, suggesting insigni ﬁ cant concentrations. Low activation energies are found for lithium migration along the a -axis channels giving rise to long-range 1D di ﬀ usion, which are supported by molecular dynamics (MD) simulations. For the crystal morphology a signi ﬁ cant surface area is exposed to these 1D di ﬀ usion channels, which would allow facile Li insertion and extraction. Using DFT simulations we reproduce the high voltage of the Li 2 Fe(SO 4 ) 2 material in accord with electrochemical data and also examine local structural distortions on lithium extraction.


Introduction
The search for alternative cathode materials to replace the layered LiCoO 2 system conventionally used within rechargeable lithium batteries has generated considerable research activity. [1][2][3][4][5][6][7] The Co-based materials pose problems associated with cost, safety and sustainability, particularly for large-scale applications (such as electric vehicles and grid storage). An avenue that has been heavily probed in the quest for new materials involves the combination of abundant Fe and stable phosphate groups (PO 4 3À ), with most interest focused on olivine-structured LiFePO 4 . [1][2][3][4][5]8 In recent studies other polyanion framework compounds have been investigated as alternative cathode materials including the following: Li 2 FeSiO 4 , 9,10 LiFeBO 3 , 11 LiFeSO 4 F, [12][13][14][15] Li 2 FePO 4 F, 16 Li 2 FeP 2 O 7 , 17 and LiFeSO 4 OH. 15,[18][19][20] More recently, the newly synthesised lithium iron sulphate, with the composition Li 2 Fe(SO 4 ) 2 , was found to display a theoretical capacity of 102 mA h g À1 and a high potential of 3.83 V vs. Li/Li + for the Fe 2+ /Fe 3+ redox couple, which is the highest reported voltage for a uorine-free iron-based compound. 21 Further studies of the structural and transport properties will provide useful insights into this signicant iron-sulphate system and may subsequently provide a new platform for related lithium battery electrode research. There is also interest in both the previously reported Li 2 Co(SO 4 ) 2 phase 21 and the recently synthesized Li 2 Mn(SO 4 ) 2 phase. 22 To fully understand the factors inuencing the electrochemical behaviour of the Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co) electrode materials, it is clear that knowledge of the fundamental solid-state properties is needed on the atomic scale. The present study uses well-established atomistic simulation and density functional theory (DFT) methods to investigate key issues related to their defect chemistry, lithium diffusion, surface structures and voltage trends. As shown in our previous studies of other lithium battery cathode materials, atomistic simulation is well suited to treating the extensive lattice relaxation around charged defects, surfaces and migrating ions in polar inorganic solids. [23][24][25][26][27][28][29] DFT methods [30][31][32][33][34] are used to provide insights into the high voltages and the local structural changes observed during electrochemical cycling of these sulphate materials.

Methodology
As these atomistic and DFT techniques are described in detail elsewhere, [35][36][37] only a general outline will be given here. For the atomistic simulations, the interatomic interactions consist of a long-range Coulombic term and a short-range component representing electron-electron repulsion and van der Waals interactions. The short-range interactions were modelled using the two-body Buckingham potential. 35  ) component, the interatomic potential model successfully formulated to simulate M 2 SO 4 (M ¼ Na, K, Rb and Cs) and XSO 4 (X ¼ Sr, Ca, Ba) was used. [38][39][40][41] Table S1 (ESI †) lists the interatomic potential parameters used in this study.
As argued previously, employing these interatomic potential methods are assessed primarily by their ability to reproduce observed crystal properties. Indeed, they are found to work well, even for phosphate and silicate cathodes [23][24][25][26][27][28][29] where there is undoubtedly a degree of covalency. The lattice relaxation about defects (such as Li vacancies) and migrating ions was treated by the Mott-Littleton scheme incorporated in the GULP code (v4.0), 43 which partitions the crystal lattice into two regions, where ions in the inner region immediately surrounding the defect ($1000 ions) are relaxed explicitly. Hence, the polyanionic framework is not treated as a rigid lattice.
Molecular dynamics (MD) simulations were performed with DL_POLY (v2.2) 44 using the same interatomic potentials and a supercell consisting of 8 Â 4 Â 4 unit cells (3328 atoms) with 10% delithiation in which the lithium vacancies were initially distributed randomly. The systems were rst equilibrated for at least 100 000 steps (with a time step of 0.5 fs). The main simulation runs of 2 million steps were then performed with an NPT ensemble (at T ¼ 473 K) to give a long simulation time of 1 ns. NPT was used to link up with experimental conditions and slight volume changes. However, it should be noted that the large supercells and long time-scales required are currently not accessible by ab initio methods.
The surface structures in this study were simulated with potentials-based methods using the METADISE code. 45 Atoms near to the surface (region I) are allowed to relax while bulk atoms (region II), which represent the rest of the crystal, are kept xed and ensure that all the ions in region I experience the forces associated with the rest of the crystal and that the energies are fully converged. The surface energy, g, of a crystal face is dened as the excess in energy of a surface simulation over the energy of a bulk system containing the same number of atoms, per unit surface area. The equilibrium morphology of the crystal can be obtained from the surface energies using a Wulff construction in which the equilibrium form of a crystal should possess minimal total surface energy for a given volume. The surface area of a facet is not directly proportional to the surface energy. The shape would be such that the normal vector to the face from a point within the crystal would be proportional to the surface energy of that face.
DFT calculations were performed using the plane wave code VASP (v5.4). 46 Since we required optimised lattice parameters, the basis set was converged against the stress which is more sensitive to an under-converged basis set than the forces. A cutoff energy of 750 eV with a k-point mesh density of at least 0.04Å À1 was needed to adequately converge the stress (6 Â 4 Â 4 grid). PAW potentials 47,48 and the PBE exchange and correlation functional 49 were used. An antiferromagnetic ordering of the moments on all transition metal (Fe, Mn and Co) atoms was found to be lower in energy than a ferromagnetic ordering, in line with magnetic measurements (T N < 8 K for Li 2 M II (SO 4 ) 2 with M ¼ Co, Mn, Fe, and T N ¼ 35 K for LiFe III (SO 4 ) 2 ) and with the magnetic structures deduced at 2 K from neutron powder diffraction. 22 DFT + U methodology was used with an effective Hubbard U eff ¼ U À J ¼ 4.0, 3.9 and 5.7 eV (J ¼ 1.0 eV) for Fe, Mn and Co respectively; these values are in agreement with the U values derived for Fe-, Mn-and Co-based cathodes. 50 Previous DFT studies on a variety of oxide electrode materials 31,51,52 have shown that such methods are well suited to probing lithium insertion/extraction properties and to predicting precise trends in cell voltages. Using the diffraction rened structural data for the lithiated (Li 2 M(SO 4 ) 2 materials) and delithiated (LiFe(SO 4 ) 2 ) structures we have calculated the open circuit voltages using: where 3{Y} is the total energy of Y and x is the number of lithium atoms per formula unit removed. In practice we have removed one lithium atom per formula unit to produce the end member LiM(SO 4 ) 2 (M ¼ Fe, Mn, Co). Bulk lithium was used to calculate the chemical potential of a single lithium atom m{Li}, which is standard practice for cell voltage calculations.

Structures and intrinsic defects
The structures of the monoclinic Li 2 M(SO 4 ) 2 (M ¼ Co, Mn, Fe) compounds have previously been reported, 21,22 and the atomic positions and unit cell parameters are collected in Table S2. † This structure comprises isolated MO 6 octahedra linked through shared oxygen vertices with the surrounding SO 4 tetrahedra. Each octahedron is linked to six SO 4 tetrahedra which are orientated in a star or pinwheel pattern when viewed along the b-axis. Conversely, each SO 4 is only bound to three MO 6 octahedra and the fourth unshared corner of the SO 4 tetrahedron points into an open channel where the lithium resides, as shown in Fig. 1.
Despite the isostructural nature of the three Li 2 M(SO 4 ) 2 phases, only the Fe analogue is observed to be electrochemically active. 21,22 Upon cycling, the Li 2 Fe(SO 4 ) 2 material transforms to the delithiated phase LiFe(SO 4 ) 2 through a biphasic mechanism, as suggested by the plateau observed at 3.83 V vs. Li/Li + on the electrochemical curve and further conrmed by in situ XRD measurements. 21 The delithiated phase LiFe(SO 4 ) 2 can also be obtained through chemical oxidation of Li 2 Fe(SO 4 ) 2 . It was found that the delithiated phase maintains the general 3D framework of Li 2 Fe(SO 4 ) 2 , 22 but the FeO 6 octahedra and SO 4 tetrahedra undergo slight rotations upon lithium removal. As a result, each octahedron is still linked to six SO 4 tetrahedra with the fourth unshared corner of the SO 4 tetrahedra pointing at the open channels where the lithium resides. However, in contrast to the lithiated phase, the lithium ions of the oxidized LiFe(SO 4 ) 2 were localized at the centre of the channels, in a halfoccupied general position in the vicinity of the (1/2 0 1/2) site. 22 The lower panel of Table S2 † presents the precise atomic positions obtained from a joint renement of the synchrotron XRD and NPD data, 22 and Fig. 2 shows the resulting crystal structure.
The starting point of the simulation study was to reproduce the experimentally observed crystal structures exhibited by the Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co) electrode materials. A direct comparison of their calculated and experimental structures is given in Table 1, which shows that the unit cell parameters calculated by interatomic potentials deviate from the experiment by at most 0.1Å, and in most cases by much less; the same is found for the individual bond lengths with mean deviations less than 0.03Å. The successful reproduction of these complex sulphate structures gives us condence that the potential models can be used reliably for subsequent defect, diffusion and surface calculations.
As noted, insight into the defect properties of cathode materials is crucial to the full understanding of their electrochemical behaviour. Isolated point defect (vacancy and interstitial) energies were calculated for these sulphate systems, which were combined to determine the formation energies for Frenkel-and Schottky-type intrinsic defects. The following equations represent the reactions involving these defects (using Kröger-Vink notation and where M ¼ Fe, Mn or Co): We also examined the Li/M "anti-site" pair defect, which involves the exchange of an Li + ion (radius 0.76Å) with an M 2+ ion (Fe 2+ radius 0.78Å, Mn 2+ radius 0.83Å, and Co 2+ radius 0.75 A), according to: This type of defect is worth investigating since Li/M anti-site or cation exchange effects have been found in other polyanionic systems including LiMPO 4 (M ¼ Mn, Fe, Co, Ni) 23,24 and Li 2 FeP 2 O 7 . 29 Examination of the resulting defect energies listed in Table 2 reveals two main predictions. First, the formation of all Frenkel and Schottky defects is unfavourable in all Li 2 M(SO 4 ) 2 structures. This suggests that the creation of vacancies (especially oxygen vacancies) and interstitial defects are unfavourable on energetic grounds. Second, the antisite energies are also relatively high, which suggests that there would be no signicant concentration of M on Li sites at typical operating temperatures in these uorine-free sulphate materials, and is consistent with  Table S2. †]  Table S2. †] Table 1 Comparison of the calculated (potentials-based) and experimental structural parameters of Li 2 M(SO 4

Lithium diffusion paths
To fully understand charge/discharge rates of potential cathodes, it is important to examine the energy barriers and dimensionality of lithium diffusion. Using atomistic simulation techniques, it is possible to examine various possible Li + transport paths, which are oen difficult to probe on the atomic scale by experiment alone. Five possible migration paths in the sulphate structures were examined, including all possible jumps between adjacent lithium sites, which are symmetry equivalent; these paths are labelled L1-L5 in Fig. 3. Activation energies for these migration paths can be derived by calculating the energy of the migrating ion along the diffusion path, and are listed in Table 3.
The calculated activation energies reveal two main points. First, for all sulphates, the paths with lowest migration energies are found to be the L1 and L4 jumps. The results indicate that combinations of the L1 and L4 pathways consists of continuous diagonal or zigzag jumps enabling long-range diffusion along the [100] direction, i.e. along the a-axis channels ( Fig. 4a and b). The net diffusion barriers are 0.47 eV, 0.54 eV and 0.42 eV for Li 2 Fe(SO 4 ) 2 , Li 2 Mn(SO 4 ) 2 , and Li 2 Co(SO 4 ) 2 respectively. We note that as with other battery cathode materials, there is no clear correlation between the Li-Li jump distances (Table S3 †) and the migration barriers are largely due to the different local structural and steric factors at the saddle-point conguration.
Secondly, we note that the remaining Li migration paths considered (L2, L3 and L5) were found to have high and unfavourable activation energies ($1.2 eV); this is probably due to the migrating Li ion coming into close proximity of the MO 6 octahedra. Hence it would appear that the Li 2 M(SO 4 ) 2 materials exhibit effective 1D transport of Li ions. We have determined experimentally the conduction activation energy in Li 2 Fe(SO 4 ) 2 by AC-conductivity performed on a sintered pellet using blocking electrodes within the 200-350 C temperature range and found from linear extrapolation a value of $1.1 eV. This   Table S2 †). Li-Li jumps considered are labelled L1-L5 (with Li-Li separations found in Table S3 †).  corresponds to the total conductivity of pristine stoichiometric Li 2 Fe(SO 4 ) 2 and will comprise both ionic and electronic conductivity terms. Hence direct comparison between this measured conductivity value and the calculated migration energies for Li ion conduction (using Li vacancy sites) is not straightforward. Despite the difference, our simulations have elucidated the pathways and dimensionality of Li ion diffusion, which are the main aims of this study.
To complement the energy minimisation calculations, molecular dynamics (MD) simulations were performed to examine the Li ion diffusion pathways and dimensionality within the sulphate structures. Scatter or density plots of Li ion coordinates over the simulation timescale enable the migration paths and the regions of the structure most frequently traversed to be visualized. Fig. 4c and d show the location of the Li density found from the simulation runs; this conrms the prediction of our static lattice calculations that Li diffusion will occur via a combination of the L1 and L4 jumps and thus allow long-range 1D diffusion along the [100] direction. As noted, Li/M antisite disorder is energetically unfavourable in these sulphates, and hence the 1D Li + diffusion is much less likely to be hindered by such defects. The MD simulations also conrm that Li diffusion involves conventional hopping between Li sites with no evidence of cooperative 'knock-on' mechanisms.

Surfaces and particle morphology
The particle morphology is important because for facile lithium insertion/extraction the predicted 1D conduction channels must terminate at a low energy surface facet. However, direct experimental measurements on surface structures of the sulphates have not been carried out.
We initially focus on the Fe-based material as a model system and due to its superior electrochemical activity. Thirteen low index crystal facets and their symmetry related equivalent surfaces of Li 2 Fe(SO 4 ) 2 have been modelled for the rst time and the resulting surface energies are listed in Table 4. Due to the highly interconnected nature of the FeO 6 and SO 4 polyhedra, there are no simple planar cuts that are non-polar. In all cases the layers must be reconstructed in order to remove the dipole.
Using these calculated surface energies a Wulff construction of the equilibrium morphology has been made which reveals a rhombohedral-like shape (Fig. 5). For small crystals where rearrangement of the crystal at each stage of the growth process is possible, a morphology close to the equilibrium form would be expected. However, there are no published indexed TEM images of Li 2 Fe(SO 4 ) 2 for direct comparison; indeed, such comparison is not trivial due to different experimental synthesis methods and possible non-equilibrium conditions. Nevertheless, our focus here is to reveal new information on the surface facets that are exposed, which are signicant for lithium extraction/insertion processes.
The {011} facet has the lowest surface energy at 0.68 J m À2 and makes up two thirds (65.5%) of the overall surface area of the particle (shown in blue in Fig. 5a). The {210} surface is the next most dominant surface comprising 27% of the surface area (shown in red in Fig. 5b). Earlier in this study we showed that the facile direction for lithium migration is along the channels in the [100] direction. Both the {210} and {102} facets expose these channels (Fig. 5c). It is well known that the growth environment can have a signicant effect upon the morphology of a particle and it may be possible to increase the surface area that allows facile Li access by alteration of the synthesis conditions.
Having predicted the dominance of the {011} and {210} surfaces in the morphology, we present details of both relaxed surface structures in Fig. 6. A useful quantity here is the relaxation energy, which is the difference between the unrelaxed surface energy of the crystal cleaved from the bulk and the nal relaxed surface energy. This quantity gives a measure of the degree of structural change in the surface region as the surface is formed. Structural relaxations in the {011} surface are small which is reected in the relaxation energy of 0.16 J m À2 . The largest relaxations occur in the topmost layer and these are associated with a small rotation of the terminal SO 4 units, which causes an inward displacement of 0.43Å of the oxygen  ion at the outer vertex and an inward displacement of 0.13Å of the Li atom that is adjacent to this. There are 50% lithium vacancies in the topmost layer and half of the SO 4 tetrahedra are absent which gives rise to a degree of rumpling. By contrast there is signicant relaxation at the {210} surface with a large relaxation energy of 1.79 J m À2 . Some loss of coordination is present in all of the terminal FeO 6 octahedra and the terminal SO 4 units undergo signicant tilting. The near surface Li move outwards by up to 1Å to form part of the terminal layer. As indicated in Fig. 6b, the prominence of the {210} surface is signicant since it exposes the facile pathway for lithium ion conduction (along the [100] channel), and hence is important for the reversible insertion/de-insertion of lithium ions.

Cell voltage trends
As with the potentials-based calculations, the DFT simulations have reproduced the Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co) experimental structures and also the delithiated LiFe(SO 4 ) 2 observed structure with a high level of accuracy (Tables S4 and S5 †). The computed cell voltage relies on a knowledge of the structure of the delithiated end member LiFe(SO 4 ) 2 . Typically for cell voltage calculations the relevant lithium atoms are removed directly from the lithiated structure, which is then energy minimised to form the delithiated structure. Numerous congurations of differing lithium ordering schemes were considered, and in our calculations we used the lowest energy structure. However, this energy minimisation may only nd the local energy minimum and it is entirely possible that a phase change can occur on lithium extraction.
The cell voltage computed directly in this way for Li 2 Fe(SO 4 ) 2 is 3.99 V which is in reasonable agreement with the measured voltage of 3.83 V. 21 However, using the joint Synchrotron XRD/neutron rened LiFe(SO 4 ) 2 structure (Table  S2 †), it is clear that the main difference between the lithiated and delithiated materials is that in the Li 2 Fe(SO 4 ) 2 structure there are two rows of lithium in a zig-zag arrangement whereas in the LiFe(SO 4 ) 2 structure (Fig. 2) there are halfoccupied Li sites in the channel. This subtle change in the structure results in a computed cell voltage of 3.91 V, which is in better agreement with that measured experimentally (Table  5). This result highlights the effect of structural changes on lithium removal and their importance in cell voltage computations.
Having isolated the delithiated (LiFe(SO 4 ) 2 ) structure we have subsequently calculated voltages for the Mn-and Co-based sulphates by setting both LiMn(SO 4 ) 2 and LiCo(SO 4 ) 2 initial structures equivalent to LiFe(SO 4 ) 2 prior to structural relaxation. The voltages calculated for Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co) are presented in Table 5.
Our calculations suggest that both Li 2 Mn(SO 4 ) 2 and Li 2 Co(SO 4 ) 2 have higher cell voltages than that of Li 2 Fe(SO 4 ) 2 as typically observed for polyanion cathodes. The value for the Cobased sulphate is above the electrolyte stability window. Such a signicant shi upwards in voltage is found for other Co-polyanion compounds, such as LiCoPO 4 for which a voltage of 4.8 V vs. Li/Li + has been observed 3,55 in comparison to 3.4 V for LiFePO 4 .
The Mn-based sulphate is known to be electrochemically inactive. As the system is delithiated, Mn 3+ (d 4 ) forms which is a Jahn-Teller active valence state. It has been debated whether Jahn-Teller distortion of the MnO 6 octahedra could lead to nonuniform structural distortions which will render the delithiated phase unstable and thus unable to form. Such effects are well known in other Mn-based battery materials. 56,57 GGA + U calculations provide a useful probe of local structural distortions that are not affected by thermal motion and statistical averaging over many unit cells. To check for Jahn-Teller effects we have examined the MO 6 distortion in the Fe, Co and Mn based sulphates, which is presented in Table 6. For this we have used the Baur 58 distortion coefficient, D, which is dened as: where l i is the distance from the central atom to the i th coordinating atom, and l av is the average bond length to n coordinating atoms. The data in Table 6 indicate that all the lithiated structures Li 2 Fe(SO 4 ) 2 , Li 2 Mn(SO 4 ) 2 and Li 2 Co(SO 4 ) 2 possess a similar low   6 and CoO 6 respectively). Since the Mn-based sulphate is electrochemically inactive, the calculated voltage presented in Table 5 assumes the same delithiated starting structure as the Co and Fe based sulphates, but contains cooperative Jahn-Teller distortion aer structure relaxation. Further work to determine if such structural distortion in this phase leads to mechanical instability is on-going including phonon spectrum simulations. A possible source of the different cell voltages for Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co) is in the relative strength of the corresponding Fe 2+ /Fe 3+ , Mn 2+ /Mn 3+ and Co 2+ /Co 3+ redox couples. As with other systems involving M 2+ /M 3+ redox couples we would expect the respective M-O bond lengths to shorten upon Li removal (oxidation) within these sulphate systems. This is indeed found with the Fe-O bond lengths contracting by $6.5% (from 2.15Å to 2.01Å), the Mn-O bond lengths contracting by $6% (from 2.21Å to 2.08Å) and the Co-O bond lengths contracting by $5% (from 2.12Å to 2.01Å).
However, distinguishing among the most reliable parameters to account for voltage variation is not an easy task. The voltage also depends upon other factors such as: inductive effects of SO 4 , PO 4 and SiO 4 ; the degree of ionic or covalent character of the cation-anion bonds; and the Madelung energy of the ionic component of the bonding. As discussed previously, 2,59 another possibility relates to the inductive effect played by lithium. There are obviously hidden issues which remain to be unravelled, and are part of on-going studies.

Conclusions
This systematic survey of the Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co) cathode materials has used a combination of atomistic simulation and DFT techniques to provide insights into the defect chemistry, Li diffusion paths, crystal morphologies and voltage trends. The key results can be summarized as follows: (1) The atomistic simulations of Li 2 M(SO 4 ) 2 (M ¼ Fe, Mn, Co) show good reproduction of their observed structures. The defect calculations indicate that the formation of all Frenkel, Schottky and Li/M anti-site intrinsic defects is unfavourable for all phases. In particular, the results suggest there would be no significant intrinsic anti-site defect concentration of Fe, Mn or Co on Li sites at typical operating temperatures within these sulphates. This is in contrast to the LiFePO 4 material, which has a small amount of Fe on Li sites.
(2) Lithium diffusion simulations reveal low migration energies ($0.4-0.5 eV) along the a-axis channels of Li 2 Fe(SO 4 ) 2 , Li 2 Mn(SO 4 ) 2 and Li 2 Co(SO 4 ) 2 suggesting 1D diffusion and good Li mobility which is important for favourable rate capability. MD calculations verify the 1D migration pathways along the a-axis channels, with no evidence of cooperative mechanisms.
(3) The calculated equilibrium morphology of Li 2 Fe(SO 4 ) 2 adopts a rhombohedral-like shape with {011} and {210} surfaces dominating. The prominence of the {210} surface is signicant since it allows access to the facile pathway for lithium ion conduction (along the [100] channel), and hence is important for the reversible insertion/de-insertion of lithium ions.
(4) The calculated cell voltage (3.91 V) for Li 2 Fe(SO 4 ) 2 is in excellent agreement with that determined experimentally (3.83 V). Characterisation of the experimental structure of the delithiated LiFe(SO 4 ) 2 phase was found to be important in this agreement. Li 2 Mn(SO 4 ) 2 is known to be electrochemically inactive and the predicted voltage for Li 2 Co(SO 4 ) 2 was found to be 5.20 V, which may be beyond the stability range of the current electrolyte.