The Structural and Electrochemical Impact of Li and Fe Site

Jun 12, 2013 - Mechanistic insights of Li+ diffusion within doped LiFePO4 from Muon Spectroscopy. Ian D. Johnson , Thomas E. Ashton , Ekaterina ...
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The Structural and Electrochemical Impact of Li and Fe Site Substitution in LiFePO4

Fredrick Omenya,† Natasha A. Chernova,‡ Qi Wang,‡,§ Ruibo Zhang,‡ and M. Stanley Whittingham*,†,‡,∥ †

Department of Chemistry, Binghamton University, Binghamton, New York 13902-6000, United States Institute for Materials Research, Binghamton University, Binghamton, New York 13902-6000, United States § Brookhaven National Laboratory, Upton, New York 11973, United States ∥ Northeastern Center for Chemical Energy Storage, Department of Chemistry, Stony Brook University, Stony Brook, New York 11794-3400, United States ‡

ABSTRACT: The crystal structure and delithiation mechanism of Li-site substituted LiFePO4 have been revealed by investigation of supervalent V3+ substitution. The combined X-ray and neutron powder diffraction data analysis surprisingly shows that the substituting aliovalent vanadium ions occupy the Fe site while some of the Fe resides at the Li site, probably as sarcopside, which leads to an increase in the unit cell volume. Such substitution reduces the miscibility gap at room temperature and also significantly lowers the solid solution formation temperature in the two-phase region. The effect of the phase diagram modification results in improved kinetics, leading to better rate performance. Such substitution, however, significantly lowers the LiFePO4 capacity at moderate current densities. KEYWORDS: olivine, site substitution, Li electrochemistry, sarcopside

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INTRODUCTION Since the discovery of LiFePO4 as a cathode material,1 many studies have been devoted to the optimization of its electrochemical performance through carbon coating,2 substitution,3 and nanosizing of the particles. At present, LiFePO4 is capable of cycling at remarkably high rates despite its inherently low ionic and electronic conductivities. Recent focus on the material has shifted toward understanding the reaction mechanism of the LiFePO4; is it a two-phase or single-phase mechanism, and what controls the Li-ion intercalation− deintercalation processes? The possibility and effect of aliovalent substitution has attracted much interest following the earlier report by Chung et al.3 that an improvement in electronic conductivity by about 8 orders of magnitude was observed when aliovalent cations below 1 at % were doped at the Li (M1) site. This dramatic increase in conductivity was later attributed to the formation of a conductive coating of Fe2P4 and/or carbon5 on the surface of the LiFePO4 during synthesis at temperatures above around 650 °C, rather than to an increase of the conductivity of the LiFePO4 structure itself. Following the earlier debate on the possibility and the effect of aliovalent substitution at the M1 site, Chiang’s6,7 and Nazar’s8 groups have recently revisited this area. Using combined X-ray and neutron diffraction data, Wagemaker et al.8 showed the incorporation of a low concentration of aliovalent ions (less than 3%) at the M1 site, which was balanced by vacancies at the same site. Meethong et al.,6 on the other hand, have recently claimed the possibility of higher level aliovalent cation substitution at the © XXXX American Chemical Society

M1 site. In their work, in which the synthesis was done at 700 °C, the formation of a second NASICON phase in the Zr substituted samples was observed. The X-ray diffraction lines of the Zr-NASICON phase were detected even at a very low Zr concentration, 1.5 mol % of Zr. These intensities of the NASICON peaks were observed to increase with the increase in the Zr content. In addition, the above work did not show any crystallographic evidence of the decrease in the miscibility gap apart from relying on the PITT data7 to determine the solubility limit. In such instances, it might not be possible to isolate the role of the second phase and that of the substituted one in the electrochemical behavior. Many researchers have also investigated substitution at the Fe (M2) site.9−17 We have previously reported the possibility and role of aliovalent substitution at the Fe site on the crystallographic and electrochemical properties of LiFePO4. From our results, up to 10 mol % of aliovalent ions could be substituted and vacancies were found to be located at the M2 site. The electrochemical performance is improved upon aliovalent substitution,13 demonstrating the role of proper stoichiometric control and reduced synthesis temperature in enhancing the aliovalent solubility. Recently, Harrison et al.17 confirmed the temperature dependence of the aliovalent substitution with up to 20% substitution of V at the Fe site using microwave synthesis. Most of the research on vanadium substitution has Received: April 20, 2013 Revised: June 9, 2013

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focused on the Fe site, with only one paper on the Li site.18 In the latter report, single-phase olivine phosphate could only be obtained for a vanadium concentration below 2%. In this work, we show unambiguous proof of high levels of aliovalent substitution at the Li site through crystallographic and electrochemical evidence. This sheds a new light on the fundamental understanding of substitution in olivine materials: the possibility of substitution, the site occupancy of the substituent, the impact of such substitution on the phase diagram, and its effect on the electrode kinetics.

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applied after one complete charge and subsequent discharge of the electrodes to 2 V. A 5 mV step was then applied from 2 to 4.3 V; the current at each step was measured until the current dropped to less than 5 μA/cm2. The partially delithiated samples for ex situ XRD were initially electrochemically delithiated at a current density of 0.1 mA/ cm2. These samples were then disassembled in a He-filled glovebox and washed several times with DMC to remove all the electrolyte before being packed in a Kapton capillary tube for room-temperature diffraction or in a quartz capillary tube for in situ high-temperature diffraction.

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RESULTS AND DISCUSSION The Li site substitution was achieved by controlling the stoichiometric composition of the precursors so as to hypothetically allow V3+ substitution and vacancies at the Li+ site as in the formula Li1−3y[VyFe]PO4. The as-synthesized substituted compounds show single olivine phases similar to the unsubstituted LiFePO4, with no observable impurity for up to 10 mol % substitution. The XRD patterns are shown in Figure 1. The as-synthesized samples with various vanadium

EXPERIMENTAL SECTION

The LiFePO4 samples were synthesized by solid-state reaction of Li2CO3, FeC2O4·2H2O, NH4H2PO4, and NH4VO3. These reactants were mixed together in amounts to give the transition-metal-rich composition Li1−3y[VyFe]PO4, where V + Fe > 1 > Li, assuming Li vacancies for charge compensation, and for y = 0, 0.025, 0.05, 0.10, and 0.20. After adding 5 wt % carbon black, the mixture was planetary ballmilled in acetone for 12 h. The acetone was evaporated and dried. The resulting precursors were preheated at 350 °C for 8 h before being sintered at 550 °C for 10 h in an 8.5% H2/He atmosphere. This low temperature was chosen so that only the effects of substitution were studied, as at higher temperatures, second phases such as Li3V2(PO4)3 and the conductive Fe2P are formed, which impact the electrochemistry.13 The phase composition and the crystal structure of the synthesized samples were determined by powder X-ray diffraction (XRD) using a Scintag XDS2000 θ-θ diffractometer equipped with a Ge(Li) solidstate detector and a Cu Kα sealed tube (λ = 1.54178 Å). The data were collected in the range of 2θ = 10−80° with a step size of 0.02° while spinning the sample to minimize preferred orientation. Highresolution synchrotron powder XRD data were collected using beamline 11-BM at the Advanced Photon Source (APS), Argonne National Laboratory, with an average wavelength of 0.413612 Å. In situ high-temperature X-ray diffraction analysis was performed at the National Synchrotron Light Source, beamline X7B, wavelength 0.3196 Å, with a heating rate of 5 °C min−1 under a helium gas flow. X-ray absorption data was collected at beamline X23A2, National Synchrotron Light Source, as described in our earlier work.13 Neutron powder diffraction (NPD) measurements were performed using the high-resolution neutron powder diffractometer (BT-1) at the NIST Center for Neutron Research, where measurements were taken with a neutron wavelength of 1.54030 Å using a Cu(311) monochromator. Time-of-flight (TOF) neutron diffraction data were collected on the POWGEN instrument in the Spallation Neutron Source (SNS) at Oak Ridge National Laboratory (ORNL). XRD and neutron powder Rietveld refinements were performed to determine the lattice parameters and site occupancy for the synthesized samples using the GSAS/EXPGUI package.19,20 The cathode materials for the electrochemical tests were prepared by mixing 80% active material, 10% carbon black (total amount of carbon in the electrode), and 10% polyvinylidene fluoride (PVDF) with 1-methyl-2-pyrrolidinone solvent. The slurry formed was then cast onto an Al foil current collector before drying. The dried electrodes, with an area of 1.2 cm2, containing 3−5 mg of active material were placed in 2325-type coin cells in a He-filled glovebox with pure lithium foil (Aldrich, thickness = 23 μm) as the counter and reference electrodes, and Celgard 2400 as the separator. LiPF6 (1 M) in a 1:1 volume ratio of ethylene carbonate (EC) and dimethyl carbonate (DMC) (LP30 from EM Industries) was used as the electrolyte. The cells were tested using a VMP2 multichannel potentiostat (Biologic). The galvanostatic charge and discharge experiments were performed at current densities of 0.05, 0.1, 0.25, 1, 2.5, and 5 mA/cm2 corresponding to 0.1, 0.2, 0.5, 1, 5 and 10 C rates (1 C corresponds to 160 mAh/g) over a 2.0−4.3 V voltage range. The cyclic voltammetry (CV) tests were done at different scan rates of 0.02, 0.05, 0.1, 0.5, 1, 5, and 10 mV/s in a voltage range of 2.2−4.3 V vs Li+ /Li. The potentiostatic intermittent test (PITT) technique was

Figure 1. High-resolution X-ray diffraction data for Li1−3y[VyFe]PO4, y = 0, 0.05, and 0.1 (λ = 0.413612 Å). The insets are magnifications of the 200, 101, and 210 reflection peaks.

concentrations in the structure show very small differences in the diffraction patterns; only a close assessment of the different diffraction peaks for the different vanadium concentrations reveals noticeable changes, as observed in Figure 1, inset. The different diffraction peaks shift in different directions with different magnitudes. For example, the 200 and the 201 diffraction peaks shift to the higher angles with an increase in vanadium content in olivine structures, but with varying degrees, and the 101 peak, on the other hand, shows a significant shift to the lower angles. The normalized intensities of the as-synthesized compounds also show slight variations; for example, the 200 and the 101 diffraction peaks show an observable decrease in normalized intensities with an increase in vanadium concentration, whereas the intensities of certain diffraction peaks, such as 210, are not affected much with the vanadium substitution level. These are strong indications of structural modification upon vanadium substitution. It is worth pointing out that, in Meethong’s work,6 formation of the NASICON phase in Zr substituted LiFePO4 was observed at very low Zr concentration; at 1.5 mol % of Zr, they could observe the formation of the NASICON phase with the B

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Table 1. Lattice Parameters of “Li1−3y[VyFe]PO4”, y = 0, 0.05, and 0.1 y in Li1−3y[VyFe]PO4

a (Å)

b (Å)

c (Å)

V (Å3)

0 0.05 0.1

10.3210(1) 10.3197(1) 10.3151(1)

6.0041(1) 5.9984(1) 5.9873(0)

4.6933(1) 4.7019(0) 4.7146(1)

290.847 291.058 291.203

an overall decrease in the unit cell volume. There is, therefore, a significant difference between the Li site and Fe site substitution; substitution on the Fe site leads to a unit cell volume decrease, whereas the attempted Li-site substitution by vanadium leads to a unit cell volume increase. These anisotropic lattice parameter changes are consistent with the changes observed in the diffraction peak position already illustrated in Figure 1 for the different vanadium concentrations. After establishing the possibility of vanadium substitution in the olivine lattice, the next issue to address is the site occupancy. To determine the site occupancy of the substituting ion and the vacancies in the substituted olivine structure, we carried out both a high-resolution XRD and an NPD data refinement. The XRD data refinement reveals a nearly full occupancy of the M2 site. However, due to the almost identical X-ray scattering powers of iron and vanadium, no information on the detailed site occupancy between the two could be obtained; the full occupancy of the M2 site could suggest either full occupancy by the Fe or the occupancy by both Fe and vanadium. The results further revealed that, close to the targeted composition of the transition metal, Fe or vanadium or both reside at the M1 site. The lithium and vacancies were mostly on the M1 site, while the concentration could not be precisely determined from the XRD refinement due to the low scattering of X-rays by lighter atoms. The XRD data thus gives the formula as Li1−3y[MyM]PO4, where M = Fe + V. The site occupancies of the various atoms were further evaluated by neutron powder diffraction. In contrast to XRD, neutron diffraction has sensitivity to light atoms, such as Li, and also provides good contrast between Fe and V since Fe has a positive scattering, whereas the vanadium has a small negative scattering factor. A combination of XRD and NPD techniques allows an accurate determination of the lithium, vanadium, and vacancies site occupancies. In the NPD data refinement, shown in Figure 3, the unit cell parameters of the substituted compound were constrained to the values obtained from the high-resolution X-ray diffraction data. The site occupancy refinement from the neutron diffraction data revealed about

intensity of this phase increasing with the increase of Zr added. This is in good agreement with our previous study,13 where we showed that, at temperatures ≥ 650 °C, some of the vanadium substituted in the olivine structure is ejected, forming the NASICON phase. When the precursor concentration of vanadium was increased to 0.2 mol, we observed the formation of pyrophosphates as the second phase, contrary to the Fe site substitution where high vanadium contents lead to the formation of the NASICON phase. The Rietveld refinement results further support vanadium solubility into the olivine structure; the lattice parameters shown in Table 1 follow Vegard’s law, as shown in Figure 2.

Figure 2. Lattice constants and unit cell volume comparison of Li1−3y[VyFe]PO4 and LiFe1−3y/2VyPO4 as a function of vanadium concentration.

The a and b lattice parameters show a small, but consistent, decrease with the increase in the vanadium content in the olivine structure, whereas the c lattice parameter shows a significant increase as compared to the change in the a and b lattice constants. The consequence of the combined effect of the anisotropic lattice parameter changes leads to an overall linear increase in the unit cell volume as a result of the larger change in the c lattice constant, which overcompensates for the small changes in the a and b parameters. In contrast, substitution on the Fe site, on the other hand, shows a significant decrease in both the a and the b lattice parameters with an almost negligible increase in the c parameter, leading to

Figure 3. Neutron diffraction data refinement for the 0.1 mol vanadium-substituted LiFePO4 (Rwp = 0.037, Rp = 0.0314, and χ2 = 1.18). C

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transition-metal ions and thereby lowering the magnetic transition and Curie−Weiss temperatures. It should be noted that the magnetic behavior of these vanadium-substituted compounds differs markedly from that of an Fe-excess single crystal, where a sharp upturn of the magnetic susceptibility was observed right at the Neel temperature.21 It supports the idea of the iron being distributed between the Fe and Li sites, resulting in a more disordered structure and the absence of a ferrimagnetic transition. X-ray absorption near-edge structure was used to determine the oxidation state and the local environment of vanadium in the olivine structure. By comparison with the references of different oxidation states, the V K-edge XANES spectra for the series of vanadium-substituted LiFePO4 (vanadium concentration = 0.05, 0.075, and 0.1 mol) exhibits an absorption edge at a similar energy to that of Li3V2(PO4)3, suggesting that the V is approximately trivalent in these substituted LiFePO 4 materials (Figure 5). The Figure 5 inset shows a weak pre-

94% occupancy of the M2 site. Since the XRD data had revealed a full occupancy of the M2 site, the neutron refinement result, therefore, suggests that the vanadium, with almost zero scattering factor, occupies the Fe site. Thus, these metal-rich olivines are best represented by the formula Li1−3yFey[Fe1−yVy]PO4, where most, if not all, of the higher valent vanadium is on the Fe site. This formulation is used in the remainder of the paper. In most of the previous publications,3,6,8 the substituting ions have been proposed to occupy the M1 site; the use of vanadium as a substituting ion allows for accurate refinement of its site occupancy due to the good contrast between V and either Fe or Li in the neutron diffraction. The thermal evolution of the molar magnetic susceptibility of the pure LiFePO4 and 0.1 mol vanadium-substituted LiFePO4 were studied in the temperature range of 2−350 K with a 1000 Oe magnetic field. Both the substituted and the unsubstituted samples exhibit paramagnetic behavior in the high-temperature region and antiferromagnetic ordering at temperatures below 50 K. The Neel temperature TN is observed to decrease with substitution, from 49.2 K for LiFePO4 to 47.6 K for the 0.1 mol vanadium-substituted material. In addition, the substituted compounds show a more pronounced increase in the molar susceptibility values at lower temperatures below the Neel temperature as compared to the unsubstituted phase (Figure 4). The Curie−Weiss law was used to fit the paramagnetic

Figure 5. Normalized vanadium K-edge XANES for vanadiumsubstituted LiFePO4 samples together with standards, VCl2, Li3V2(PO4)3, VO4, and V2O5, corresponding to vanadium in +2, +3, +4, and +5 valence states. The inset shows the magnified pre-edge region for the 0.05, 0.075, and 0.1 mol of vanadium-substituted LiFePO4.

edge absorption, characterized by a triplet peak in all three samples. Such a feature arises from 1s−3d transitions, which is a formally electric dipole-forbidden transition and gains its intensity through electric quadrupole coupling and p−d orbital hybridization in the distorted MO6 octahedra. Crystal field splitting in a quadrupole transition has been attributed to the formation of multiplet pre-edge peaks, as well as a possible contribution from the dipole transition. The relatively low intensities in the V-LFP series corroborate that vanadium resides at the octahedral site (M1 or M2) in the substituted compounds. This pre-edge feature is observed to decrease in intensity upon the increase of vanadium concentration in the olivine structure, suggesting that the distorted octahedron

Figure 4. Temperature dependence of the magnetic susceptibility of LiFePO4 and 0.1 mol vanadium-substituted LiFePO4. The inset shows inverse molar susceptibility corrected for the temperature independent contribution and their fit to the Curie−Weiss law.

region, 150−350 K, to determine the Curie constant and the Curie−Weiss temperature (results summarized in Table 2). Both the Curie constant and the absolute value of the Curie− Weiss temperature decrease with substitution. These changes in the magnetic properties support the incorporation of vanadium ions in the olivine structure, even more specifically, onto the Fe site, which weakens the magnetic interactions between the

Table 2. Magnetic Parameters of Pristine and 0.1 mol Vanadium-Substituted LiFePO4a y in Li1−3y[VyFe]PO4

TN, K

χ0, 10−4 emu/mol

θ, K

C, emu K/mol

0 0.1

49.2 47.6

1.42 1.45

−85.6 −75.4

3.24 3.21

μeff

exp

, μB

5.09 4.83

μeff

theor

, μB

5.09 4.92

TN is determined as an inflection point of the M(T) dependence. μeffexp is determined using μ = [8C/(1 + y)]1/2. In calculations of μefftheor = [(yμV2 + μFe2)/(1 + y)]1/2, the magnetic moment of Fe2+ is assumed to be μFe = 5.09 μB, as in LiFePO4, since it is difficult to account for its orbital contribution. The magnetic moment of V3+ is assumed as spin-only μV = 2.82 μB. a

D

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The lattice mismatch results were further investigated by a Rietveld refinement of the partially delithiated phases. These refinement results, shown in Table 3, indicate a systematic decrease in the a and b lattice parameters and of the unit cell volume with the increase in the vanadium concentration in the olivine structure. The calculated volume changes between the Li-poor and the Li-rich phases coexisting in the partially delithiated samples show a systematic decrease from 6.0%, to 5.3%, to 4.5%, corresponding to LiFePO4, Li0.925Fe0.025[V0.025Fe0.975]PO4, and Li0.85Fe0.05[V0.05Fe0.95]PO4, respectively. The crystallographic changes of the single phase Li0.7V0.1FePO4 were further studied for Li solubility upon electrochemical delithiation to different states of charge. The partially delithiated samples lead to a much smaller lattice mismatch between the Li-rich and the Li-poor phases such that the next neighboring peaks (for example, the 200 peaks) belonging to the heterosite and triphylite phases could not be resolved by the XRD instrument, as shown in Figure 7. Considering the

becomes more symmetric in geometry with the increase in substitution level (Figure 5, inset) We further studied the valence state of the vanadium in electrodes during the charge process. It is observed that the vanadium edge feature remains unchanged in shape and position at different states of charge, which suggests that the vanadium does not participate in the electrochemical process. The XANES of the 0.1 V substituted sample charged to 4.3 V is plotted in Figure 5 along with pristine samples to illustrate this comparison. The effect of lithium nonstoichiometry upon delithiation was further examined by XRD on partially delithiated samples to determine the impact of substitution on the miscibility gap. The diffraction profiles from synchrotron-based XRD are shown in Figure 6 for both pure and vanadium-substituted LFP

Figure 6. X-ray diffraction patterns of partially delithiated Li1−3y−xFey[Fe1−yVy]PO4, y = 0, 0.025, 0.05, and 0.1 and x = 0.5, 0.45, 0.35, and 0.3, respectively (λ = 0.3196 Å). Figure 7. Ex situ synchrotron XRD patterns of Li(0.7−x)Fe0.1[Fe0.9V0.1]PO4 at different states of charge: x = 0.0, 0.25, 0.3, and 0.46. The inset shows the magnified section of the 200 reflections of the triphylite and heterosite phases (λ = 0.3196 Å).

electrodes at the 50% state of charge (Li1−3y−xFey[Fe1−yVy]PO4, for y = 0, 0.025, 0.05, 0.1 and x = 0.5, 0.45, 0.35, 0.3, respectively). At the given states of charge, the samples in the relaxed state at room temperature unquestionably show the coexistence of the two phases, triphylite and heterosite, without the existence of any other crystalline phase. Taking the 200 peaks, for example (inset in Figure 6), as the concentration of the vanadium is increased in the olivine structure, the mismatch between the heterosite and the triphylite 200 diffraction peaks decreases. For example, at 0.1 mol of vanadium substitution, the two diffraction peaks appear as one broad single peak. This clearly demonstrates that the miscibility gap is a function of the aliovalent substitution level. The miscibility gap shrinks systematically with the increase in vanadium solubility in the olivine structure, reaching a concentration where the two peaks cannot be readily resolved and they appear as one broad single peak. The mismatch in these samples is much smaller as compared to the substitution targeting the Fe site.

200 reflection peaks of the two phases of the partially delithiated sample, the two peaks appear as a single broad diffraction peak observed at around 2θ = 3.5° at all states of charge, as shown in Figure 7, inset. The unresolved peaks can be mistaken to be a single olivine phase. However, comparison of these peaks with those peaks of the same partially delithiated sample that are not affected by the contribution from the two phases, heterosite and triphylite, or with the similar diffraction peaks of the pristine material, shows that the broadness is a result of overlapping unresolved peaks. Because of these unresolved peaks of Li-rich and Li-poor phases, it may not be feasible to refine these phases of the partially delithiated Li0.7−xV0.1FePO4. This observation is strong evidence that controlling the stoichiometric composition of the starting

Table 3. Lattice Parameters of Partially Delithiated Li1−3y−xFey[Fe1−yVy]PO4 sample LiFePO4, 50% charged Li0.925Fe0.025[V0.025Fe0.975]PO4, 45% charged Li0.85Fe0.05[V0.05Fe0.95]PO4, 35% charged

T H T H T H

a (Å)

b (Å)

c (Å)

V (Å3)

Rwp

10.309(1) 9.835(1) 10.291(1) 9.862(1) 10.273(1) 9.906(2)

5.999(1) 5.805(1) 5.987(1) 5.818(1) 5.978(1) 5.840(1)

4.697 (1) 4.783(1) 4.702(1) 4.777(1) 4.711(1) 4.775(1)

290.505 273.048 289.704 274.067 289.287 276.278

0.023

E

0.038 0.028

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precursors leads to the substitution of the transition metal at the M1 site, which affects the electrochemical performance of the LiFePO4. In addition to the XRD crystallographic information, PITT was used to study the delithiation mechanism to further understand the effect of substitution on the reaction pathway and the reaction kinetics. This technique has been used previously to determine the lithium nonstoichiometric compositions and the phase transformation of LiFePO4 on doped or nano-LiFePO4.6,7,22 The PITT curves for the substituted samples differ markedly from that of the unsubstituted one. First, a current maximum is observed in LiFePO4 but is absent in the substituted samples; a current relaxation maximum is indicative of a first-order phase transformation. Second, the current relaxation time at the “plateau region” for pure LiFePO4 is the longest, whereas those of the substituted materials are much shorter and decrease with the increase of the vanadium content. Such current relaxation times can be qualitatively used to understand the reaction kinetics in this class of materials. On the basis of the current− time response, the two-phase behavior is observed to decrease with the increase in the vanadium concentration (Figure 8). The diffusion-like behavior and continuous variation in the

voltage curve observed in the 10% vanadium material are indicative of a single-phase solid solution. Our XRD results, however, show that these partially delithiated samples demonstrate the coexistence of two phases, but with a narrow lattice mismatch between them. The current response in PITT is, therefore, very closely related to the mismatch of the heterosite and triphylite phases. To understand the effect of vanadium substitution on the LiFePO4 phase diagram, we performed an in situ hightemperature synchrotron XRD. The thermal behavior of partially electrochemically delithiated samples of y = 0.0, 0.025, and 0.05 at x = 0.5, 0.4, 0.35, and 0.3, respectively, were investigated. From our XRD data in Figure 9a, no matter the content of vanadium in the structure, at room temperature, these partially delithiated samples existed as a mixture of two phases, triphylite and the heterosite. The mismatch between the two phases at room temperature, however, decreases with the increase in vanadium solubility into the structure, as already discussed. The diffraction peaks between 3.5−4°, 6−6.5°, and 7−7.7°, are characteristic of the two-phase nature. Li0.5FePO4 was used as a baseline case. At temperatures below 150 °C, no observable transformation is found in the diffraction pattern, as shown in the inset in Figure 9a; the observable transformation starts at a considerably high temperature above 200 °C. A significant solid solution phase is observed at around 250 °C, where a single unresolved peak is observed, as evident in the 200 reflection. The complete transition to the single solid solution phase is not complete till temperatures above 300 °C. Introducing the vanadium into the olivine structure tends to change the solid solution formation temperature. When as little as 0.025 mol of vanadium is introduced into the olivine structure, the transformation temperature is slightly lowered, and at 150 °C, a small change can be observed in the 200 peak, as shown by the magnified 200 peaks (Figure 9b, inset). Compared to LiFePO4, a considerable solid solution formation is observed at 200 °C, and on further heating the sample, a complete single-phase solid solution is formed at a temperature much below 300 °C. When the concentration of the vanadium in the LiFePO4 structure is further increased to 0.05 mol (Figure 9c), the solid solution formation temperature is lowered and is observed at temperatures below 100 °C and a considerable solid solution is observed at 150 °C. At temperatures below 250 °C, a complete solid solution is formed, which is persistent up to temperatures around 500 °C. This shows that, upon increasing the vanadium concentration in the olivine structure, the solid solution formation temperature is systematically lowered, allowing the system to access the single-phase solid solution at fairly low temperatures as compared to the unsubstituted LiFePO4. The electrochemical properties based on galvanostatic charge discharge for the different vanadium substitution levels are shown in Figure 10. For all the curves between 2.0 and 4.3 V, we only observe a single plateau at around 3.45 V, which is characteristic of the LiFePO4. There is no additional redox plateau from the vanadium substitution. At low current densities, there is no significant difference between the charge and the discharge polarization in the vanadium-substituted samples and unsubstituted LiFePO4. Also, at the lowest current densities, we observe a decrease in the electrochemical capacity with the increase of the vanadium content in the olivine structure; this is consistent with the lower lithium content caused by the occupancy of a fraction of the Li sites by transition metals, and the lithium vacancies needed for the

Figure 8. PITT graphs of the 0.0, 0 5, and 10 mol % vanadiumsubstituted LiFePO4; the blue curve shows a voltage step of 5 mV and the corresponding current relaxation time in red. F

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Figure 10. Galvanostatic charge−discharge curves of Li1−3yFey[Fe1−yVy]PO4 discharged at different rates: (a) 0.1 C and (b) 10 C. The inset shows the OCV of the different samples.

capacity drop in the highly substituted vanadium electrodes, as the current is increased. The kinetics of reaction are much improved by vanadium substitution, but the overall capacity is reduced, as clearly shown in Figure 10a. The reaction kinetics dependent on the vanadium concentration in the olivine structure were further evaluated using cyclic voltammetry (CV). Figure 11 shows the CV curves for the vanadium-substituted LiFePO4 over a range of scan rates. The CV profiles were normalized with the active mass of each electrode for comparison of the effect of substitution on the electrochemical properties. The CV curves show only one distinct anodic and cathodic peak at all scan rates within the 2.2−4.3 V range. At slower scan rates, 0.02 or 0.05 mV/s, the peak current is observed to decrease with the increase in the vanadium substitution level. The peak positions show only a very small shift, if any, at such slow scan rates, as shown in Figure 11. This suggests good electrochemical performance of LiFePO4 at such a slow scan rate. As the scan rate increases, first, the current peak intensities, which are smaller for the 5 and 10 mol % of vanadium-substituted samples, significantly grow relative to the low vanadium substitution with the increase in the scan rate. For example, at 0.1 C, the y = 0 and 0.05 have similar current densities and, at an even much higher scan rate, 1 or 5 mV/s, the peak current of the y = 0.05 is higher than y = 0. The 0.1 mol % vanadium-substituted sample has the smallest current peak; however, at a higher scan rate, it is observed to significantly increase relative to the unsubstituted and 0.05 mol % vanadium-substituted one. This shows that the kinetics of the higher vanadium-substituted LiFePO4 is superior to the low-level substituted and unsubstituted LiFePO4. Second, the peak separation between the anodic and the cathodic peaks is generally similar at slow scan rates, but as the scan rate increases, the 10 mol % vanadium-substituted sample has the least peak separation, followed by the 5%, whereas the

Figure 9. In situ XRD patterns of (a) LiFePO4, (a) Li0.925−xFe0.025[Fe0.975V0.025]PO4, x = 0.4, and (c) Li0.85−xV0.05FePO4Li0.85−xFe0.05[Fe0.95V0.05]PO4, x = 0.35, collected under a He atmosphere; the rectangular boxes show peaks affected by solid solution formation. The insets show the 200 peaks at 25, 100, and 150 °C (λ = 0.3196 Å).

transition-metal charge compensation. On increasing the discharge current densities from low rates, 0.1 C, to high rates, 10 C, an increase in polarization is observed as expected. At a 10 C current density, the unsubstituted LiFePO4 has the highest polarization and the polarization decreases with the increase of the vanadium content. There is thus a much less G

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Figure 11. Normalized CV curves for different concentrations of vanadium in Li1−3yFey[Fe1−yVy]PO4 at 0.02, 0.1, 1, and 5 mV/s scan rates.

unsubstituted LiFePO4 has the highest peak separation. These peak potential shifts are attributed to polarization as a result of the kinetic limitations. Our results are consistent with the lower polarization observed in galvanostatic charge discharge curves at higher current densities for increasing the vanadium concentration in the structure. At a 5 mV/s scan rate, the unsubstituted LiFePO4 shows the least electrochemical activity as compared to other vanadium-substituted counterparts at the same scan rate due to the shift of its redox peaks and very reduced peak intensities. The rate performance of the vanadium-substituted LiFePO4 was also evaluated. The cells were charged at the same current densities to 4.3 V and then discharged at different current densities from 0.2 to 10 C; the results are shown in Figure 12. In the comparative analysis of the systematic variation of vanadium concentration level in the structure, we observe that, at the lower rates of 0.1 and 0.2 C, the unsubstituted LiFePO4

shows the highest electrochemical capacity. The substituted samples exhibit inferior capacities at such low rates, as expected because of their lower lithium content; this was also noted by Meethong et al.7 for Mg and Zr substitution. At higher current densities, the unsubstituted LiFePO4 loses much more of its capacity to such an extent that, at 1 C, the capacities of the 2.5% and 5% vanadium-substituted samples outperform the LiFePO4. The capacity drop as the rate increases is much lower in the vanadium-substituted samples, similarly for Mg and Zr substitution.7 The higher the vanadium content in the olivine structure, the better the rate performance. If the capacities of the various systems are normalized with the achievable capacity at 0.1 C for each system, then, at higher current densities, 10 C, the LiFePO4, 2.5, 5, and 10 mol % vanadium-substituted compounds have a capacity relative to the 0.1 C value of around, 40, 55, 60, and 70%, respectively. This shows that the capacity drop decreases with the increase in the vanadium content. The faster kinetics upon vanadium substitution is consistent with the CV curve results. The improved rates on substitution can be attributed to the reduced lattice mismatch between the Li-rich and the Li-poor phases, which decreases with the increase of vanadium content in the structure and also the ease of accessibility of the solid solution at lower temperatures as compared to LiFePO4, as already discussed. It is clear from the high rates achieved in these metal-rich materials that the iron atoms at the lithium sites are not blocking the lithium ions from fast diffusion. This implies that most of the diffusion pathways, tunnels, must remain free of Fe atoms, and that these iron atoms are probably clustered in the structure, forming nanodomains of an iron-rich phase, such as sarcopside, Fe3(PO4)2. The high rates in these materials may be associated with these nanodomains introducing disorder into the structure, thereby reducing the overpotential required for single-phase formation, or the nucleation energy to form the second phase. Studies are underway to determine the exact nature of this phase.23 It is likely that, whenever the substituent ion does not form a sarcopside phase, then that ion goes to the Fe site, displacing Fe to the sarcopside phase.

Figure 12. Rate perfomance as a function of vanadium concentration in the olivine structure, Li1−3yFey[Fe1−yVy]PO4, y = 0, 0.025, 0.05, and 0.1 mol of vanadium. H

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(11) Hong, J.; Wang, X.-L.; Wang, Q.; Omenya, F.; Chernova, N. A.; Whittingham, M. S.; Graetz, J. J. Phys. Chem. C 2012, 116 (39), 20787−20793. (12) Sun, C. S.; Zhou, Z.; Xu, Z. G.; Wang, D. G.; Wei, J. P.; Bian, X. K.; Yan, J. J. Power Sources 2009, 193 (2), 841−845. (13) Omenya, F.; Chernova, N. A.; Upreti, S.; Zavalij, P. Y.; Nam, K.W.; Yang, X.-Q.; Whittingham, M. S. Chem. Mater. 2011, 23 (21), 4733−4740. (14) Zhang, L.-L.; Liang, G.; Ignatov, A.; Croft, M. C.; Xiong, X.-Q.; Hung, I. M.; Huang, Y.-H.; Hu, X.-L.; Zhang, W.-X.; Peng, Y.-L. J. Phys. Chem. C 2011, 115 (27), 13520−13527. (15) Ma, J.; Li, B. H.; Du, H. D.; Xu, C. J.; Kang, F. Y. J. Electrochem. Soc. 2011, 158 (1), A26−A32. (16) Hua, N.; Wang, C. Y.; Kang, X. Y.; Wumair, T.; Han, Y. J. Alloys Compd. 2010, 503 (1), 204−208. (17) Harrison, K. L.; Bridges, C. A.; Paranthaman, M. P.; Segre, C. U.; Katsoudas, J.; Maroni, V. A.; Idrobo, J. C.; Goodenough, J. B.; Manthiram, A. Chem. Mater. 2013, 25 (5), 768−781. (18) Chiang, C.-Y.; Su, H.-C.; Wu, P.-J.; Liu, H.-J.; Hu, C.-W.; Sharma, N.; Peterson, V. K.; Hsieh, H.-W.; Lin, Y.-F.; Chou, W.-C.; Lee, C.-H.; Lee, J.-F.; Shew, B.-Y. J. Phys. Chem. C 2012, 116 (46), 24424−24429. (19) Toby, B. H. J. Appl. Crystallogr. 2001, 34 (2), 210−213. (20) Larson, A. C.; VonDreele, R. B. Los Alamos Natl. Lab., [Rep.] LA (U.S.) 2000, 86, 748. (21) Chen, J.; Vacchio, M. J.; Wang, S.; Chernova, N.; Zavalij, P. Y.; Whittingham, M. S. Solid State Ionics 2008, 178 (31−32), 1676−1693. (22) Meethong, N.; Huang, H. Y. S.; Carter, W. C.; Chiang, Y. M. Electrochem. Solid-State Lett. 2007, 10 (5), A134−A138. (23) Khalifah, P.; Meng, S.; Chernova, N.; Whittingham, M. S., work in progress at NECCES.

CONCLUSION We have systematically and unambigously demonstrated crystallographic evidence for higher aliovalent substitution in LiFePO4 and its effect on the electrochemical properties. Unlike previous reports on aliovalent substitution at the Li site, our combined XRD and NPD analysis shows the distribution of aliovalent V to the Fe site with some Fe occupying the Li site. These Fe atoms on the Li site are clustered in microdomains of a sarcopside-like phase, leaving the Li tunnels free for fast Li-ion diffusion. The overal effect leads to an increased Li solubility range accompanied by a decrease in the unit cell volume change between the heterosite and the triphylite phases, which may facilitate the Li-ion mobility. Such substitution lowers the solid solution temperature as a function of the concentration of substituted transition-metal ions in the structure. The reaction kinetics are also improved with substitution; however, this comes at the expense of lowering the electrochemical capacity.

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AUTHOR INFORMATION

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Shirley Meng for the initial suggestion of iron clustering in LiFePO4. This research is supported as part of the Northeastern Center for Chemical Energy Storage, an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, Basic Energy Sciences, under Award Number DE-SC0001294. Use of the Advanced Photon Source at Argonne National Laboratory and the National Synchrotron Light Source at Brookhaven National Laboratory is supported by the U.S. Department of Energy, Office of Science, Basic Energy Sciences, under Contract Nos. DE-AC02-06CH11357 and DE-AC02-98CH10886, respectively. We also acknowledge support from Oak Ridge National Laboratory’s Spallation Neutron Source sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, U.S. Department of Energy, and the National Institute of Standards and Technology, U.S. Department of Commerce, in providing the neutron research facilities.

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