On the potential use of carbon-free mesoporous precursors of LiFePO4 for lithium-ion batteries electrode

Solid State Ionics 255 (2014) 30–38

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On the potential use of carbon-free mesoporous precursors of LiFePO4 for lithium-ion batteries electrode Rafael Trócoli a, Juliàn Morales a, Jesùs Santos Peña b,⁎ a Departamento de Química Inorgánica, Instituto Universitario de Investigación en Química Fina y Nanoquímica, Edificio Marie Curie, Campus de Rabanales, Universidad de Córdoba, 14071 Córdoba, Spain b Laboratoire de Physicochimie des Materiaux et Electrolytes pour L'energie (PCM2E), Bâtiment J, Campus de Grandmont, Université de Tours, 37200 Tours, France

a r t i c l e

i n f o

Article history: Received 6 May 2013 Received in revised form 11 November 2013 Accepted 14 November 2013 Available online 8 December 2013 Keywords: Mesoporous phosphates Lithium iron phosphate Mesopore order Mesoporous precursors Lithium ion electrodes

a b s t r a c t This paper describes the preparation, characterization and electrochemical study of various iron and lithium iron phosphates. The pristine iron phosphate samples obtained were amorphous materials with various textural properties, in terms of average pore size, pore order and specific surface area. Some solids were mesoporous, featuring high surface area, narrow pore size distribution and high pore volume. Lithium iron phosphates were obtained by chemical lithiation of these iron phosphates, using LiI dissolved in acetonitrile. These materials were amorphous and electrochemically inactive against lithium. Following heating at 550 °C, the materials crystallized and showed disparate electrochemical performance in lithium half-cells. This can be ascribed not only to the presence in the samples of impurities of different nature (and inherent to the synthesis method), but also to different textural and morphological properties. In fact, the best electrochemical performance was that of crystallized LiFePO4 obtained from the mesoporous iron phosphate and consisting of rounded nanoparticles with high surface area and adequate porosity. As a result, a better electrolyte impregnation in the porous structures is highly expected. Furthermore, enhancement of the lithium ion diffusion inside the network is observed. This justifies the potential use of carbon-free mesoporous precursors of LiFePO4. However, introduction of a conductive carbon coating onto the phosphate particles and/or removal of the impurities prove necessary for good capacity retention upon cycling. © 2013 Elsevier B.V. All rights reserved.

1. Introduction Mesoporous materials were initially based on silica derivatives [1,2] and later expanded with oxides of metals including some in the transition groups [3–6]. These materials have arisen the interest amongst researchers for application in various fields including catalysis and electronics. More recently, they have been used as precursors for different systems via the replica method [7]. This method has enabled the preparation of some interesting materials to be used in lithium ion batteries, where they demonstrate favorable electrochemical properties, as positive or negative electrodes. Their improved performance is due to their increased surface area and porosity (resulting in more efficient electrolyte/electrode contact), and also to their organized structure (facilitating transport properties within the electrochemical systems) [5,8,9]. Mesoporous iron phosphates, whose formulae are close to FePO 4·2H2 O and whose surface areas fall in a typical range for oxide materials, have been successfully prepared by some research groups [10–15] and used as catalysts [12,13] or electrochemical energy storage [11,14,15]. Regarding the latter, Santos-Peña et al. ⁎ Corresponding author. E-mail address: [email protected] (J. Santos Peña). 0167-2738/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.ssi.2013.11.038

[11] were the first to demonstrate the suitability of this compound for lithium cells. Later, other authors [15] have improved the synthetic method for the material and obtained an increased surface area and electroactivity as a result. Porosity and surface area are seemingly keys to understanding the good electrochemical properties of this material. Our group has recently synthesized LiFePO4 nanoparticles in a simple way [16]. Phospholivine is a very promising material for positive electrodes in lithium ion batteries [17–20] and is already available commercially. However, this material shows its limitations in terms of electronic conductivity and lithium ion diffusion inside its structure [17,18]. The latter is particularly apparent in LiFePO4 based cells subjected to fast current regimes; however, a number of authors have shown that decreasing their particle size to nanometric level can help improve electrochemical performance [16,19,20]. Also, the combined use of carbon coating (N 3.8% in weight ratio referred to LiFePO4) with a porous structure is sufficient to greatly improve the kinetics of the electrochemical reaction of the phosphate [21,25]. Gaberseck et al. [22] pointed out several advantages of these porous materials referred to nanoparticles, which include the presence of a mixed-conducting transport system and a better electrical contact between the porous particles. The aim of this work was to combine the favorable effects of using the mesoporous materials with the suitable preparation of LiFePO4

R. Trócoli et al. / Solid State Ionics 255 (2014) 30–38

from amorphous iron phosphate as indicated by Prosini et al. [26] Therefore we prepared mesoporous iron phosphate, which was then chemically lithiated and crystallized in vacuo to obtain an electroactive material. For comparison purposes, we also prepared a series of samples from amorphous iron phosphate that demonstrated poorer electrochemical properties than those obtained from pristine mesoporous materials.

2. Experimental 2.1. Materials Fe(III) nitrate nonahydrate (≥98%), Na2HPO4 (≥99%), HF (puriss; ≥ 48% in water), hexadecyltrimethylammonium bromide (CTAB) (≥ 98%), tetrapropyl hydroxide (TPAOH) (1 M in water), anhydrous lithium acetate (99.95%), lithium iodide (99%) and acetonitrile (99.8%) were purchased from Aldrich. The electrolyte used was Merck battery electrolyte LP 40, which consists of 1 M LiPF6 in ethylene carbonate (EC) and dimethyl carbonate (DMC) in a 1:1 w/w ratio. Amorphous FePO4·2H2O (a-FP) was prepared by mixing solutions containing iron(III) nitrate and Na2HPO4 in stoichiometric amounts. The precipitate thus obtained was allowed to settle, centrifuged and washed thoroughly with water. Details on the preparation of mesoporous iron phosphate can be found elsewhere [11]. Briefly, fresh amorphous iron phosphate was dissolved in HF. Then, a solution of CTAB was added, the mixture stirred for 30 min and heated at 60 °C for 2.5 h under continuous stirring. After cooling, the mesoporous phosphate was precipitated by adding TPAOH solution to raise the pH. A white solid (t-FP from templated-iron phosphate) was recovered by centrifugation, and thoroughly washed with water and acetone. In order to extract the template [10], the mesoporous material was immersed in a solution containing 0.05 M lithium acetate in ethanol and stirred at r.t. for 18 h. The solid was recovered by filtration and washed with ethanol first and then water. The yellow solid thus obtained was the mesoporous iron phosphate (m-FP) used in this work. An amount of 0.7 g of iron phosphate (a-FP or m-FP) was immersed in a 1 M LiI solution in acetonitrile, the system immediately changes to an orange (a-FP) or red color (m-FP). After stirring at r.t. for 18 h the solids were recovered by filtration and washed with acetonitrile first and then water. The samples thus obtained were named a-LFP (obtained from a-FP) and m-LFP (obtained from m-FP). The phosphates were crystallized by heating 0.4 g of a-LFP or m-LFP at 550 °C for 3 h under dynamic vacuum. The resulting samples were named a-LFP550 and m-LFP550, respectively. For clarity, Table 1 collects the preparation methods of all these samples.

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2.2. Methods X-ray diffraction (XRD) patterns were recorded on a Siemens D5000 X-ray diffractometer, using CuKα radiation and a graphite monochromator. Textural properties were established from nitrogen adsorption/ desorption isotherms on a MicromeriticASAP 2020 instrument using N2 gas as adsorbate. Pore size distribution was calculated by applying the Barret et al. method to the desorption branch of the isotherms. Transmission electron microscopy (TEM) images were obtained on a Jeol 2010 microscope operating at 200 keV. Thermogravimetric (TG) analysis were carried out on a Setsys Evolution 16/18 instrument from SETARAM using a N2 stream at 20 ml·min− 1at a heating rate of 5 °C·min−1 up to 800 °C. FTIR spectra were recorded on a FossNIRSystems 6500 System Instrument. Electrochemical measurements were made with a two-electrode cell. Swagelock cells were built with lithium as counter and referenceelectrode, and amorphous or crystalline “LiFePO4” samples as working electrode. “LiFePO4” electrode mixtures were made from active material, carbon black and Teflon in a 85:10:5 weight proportion, and subsequently pressed on stainless steel. As adsorbed water on the compounds could react with lithium in the half-cells, prior to entering the glove box, the electrodes are submitted to heating at 90 °C under vacuum in an oven, overnight. Potentiostatic cycling was done at 10 mV·10 s− 1. Galvanostatic cycling was carried out under a C/7 regime (1 mol of lithium per mole of LiFePO4 is deinserted/inserted in 7 h). Both electrochemical measurements were controlled via a MacPile II potentiostat−galvanostat. For electrochemical impedance spectroscopy, a TOYO™ three electrode cells was used, with lithium foil as counter electrode and lithium ribbon as reference electrode. Spectra were recorded with an EGG PAR 273 A potentiostat coupled to a 1255 Schlumberger Frequency Response Analyzer. Measurements were made over the range 84 kHz–3.4 mHz. The excitation signal was 10 mV peak to peak. The equilibrium potential was assumed to be reached when the drift in open-circuit voltage fell below 1 mV·h− 1. Impedance diagrams were fitted with the software Zplot®, from Scribner Associates.

3. Results and discussion 3.1. Physicochemical characterization of the amorphous and mesoporous phosphate based materials The XRD patterns (results not shown) confirmed the amorphous nature of a-FP. Fig. 1 shows the XRD patterns at a small angle for the different mesoporous materials prepared in this work. The as synthesized

Table 1 Preparation methods of the phosphate materials studied in this work. Sample

Method

Heat Main formula treatment

a-FP

Spontaneous precipitation of Fe3+ with PO3− 4 Reaction between LiI in acetonitrile + a-FP Heating of a-LFP under vacuum Treatment of a solution of a-FP in aqueous HF with CTAB Extraction of t-FP with LiAc in ethanol Reaction between LiI in acetonitrile + m-FP Heating of m-LFP under vacuum

NO

FePO4·2.8H2O

NO

LiFePO4·xH2O (x b 2.8)

3 h, 550 °C NO

LiFePO4

NO

FePO4·(CTAB)0.0032·1.8H2O

NO

LiFePO4·yH2O (y b 1.8)

3 h, 550 °C

LiFePO4

a-LFP a-LFP550 t-FP

m-FP m-LFP m-LFP550

FePO4·(CTAB)0.25·1.7H2O

Fig. 1. Small-angle XRD patterns for various iron phosphate based mesoporous materials.

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mesoporous phosphate, t-FP, exhibited an ordered pore distribution, with peaks corresponding to the [100], [110] and [200] planes at 2.1°, 3.6° and 4.1° (2Θ). Similar values have been reported for mesoporous samples obtained with SDS [10,12,13] or CTAB [11] templates. Sample m-FP, which was obtained by removing CTAB from sample t-FP, also exhibited ordered mesoporosity (Fig. 1). However, removing the template caused slight broadening in its peaks and increased the intensity of [h00]. This has often been observed in silica matrices and associated to a high organization in the pores by effect of the template being chemically or thermally removed. A slight shift in the peaks to higher angles was also observed. Cell parameter a0 was calculated from Fig. 1, using the expression [10]a0 = 2d100/√3. a0 was 4.1 nm for m-FP. Lithiation of this sample resulted in the cell contracting from 4.1 to 3.88 nm. This contraction was accompanied by an increased pores disorder, suggested by the broadening in the peak for the [100] plane. The four phosphates (a-FP, a-LFP, m-FP and m-LFP) were subjected to nitrogen adsorption/desorption measurements. The isotherms for the former two samples, Fig. 2a exhibited a small hysteresis loop at high relative pressures suggestive of a limited mesoporous system. The isotherms for the latter two samples shown in Fig. 2b, were markedly different in shape. Thus, they were type IV in the IUPAC classification and exhibited a pronounced hysteresis loop which was assigned to a well-defined mesoporous system. Table 2 lists the specific surface areas, derived from the figures. Sample a-FP sample had a relatively high surface area, 91 m2·g− 1, exceeding that of a nanomaterial prepared from solutions containing Möhr salt and NH4H2PO4, and oxidized with hydrogen peroxide [26]. Under the assumption that the solid consisted of spherical particles, a diameter of 30 nm was

Fig. 2. Nitrogen sorption isotherms for (a) amorphous and (b) mesoporous samples synthesized, at 77 K. The inset shows the corresponding BJH pore size distribution curve as calculated from the adsorption branch of the mesoporous samples.

Table 2 Textural properties of the phosphate materials prepared in this work. Sample

SBET (m2·g−1)

Vpore (cm3·g−1)

Vmicropore (cm3·g−1)

bdporeN (nm)

a-FP a-LFP a-LFP550 m-FP m-LFP m-LFP550

91 43 1.5 395 136 13

0.589 0.336 0.011 0.373 0.261 0.142

0.007 0.007 0.0003 0.090 0.036 0.005

26.2 31.9 33 4.7 8.7 47

calculated from SBET and the phosphate density. Lithiation of sample a-LFP decreased its surface area (average diameter, 40 nm). Our results depart from those of Prosini et al. [26], who found a decrease in particle size upon lithiation (from 109–136 to 86–107 nm). Samples m-FP and m-LFP exhibited a significant increase in SBET values, which were 3.5–4.5 times higher (Table 2). Also, reaction with LiI halved their surface area value. The pore size distribution of both samples is shown as an inset in Fig. 2b. Sample m-FP gave a rather symmetric peak centered around 3.7 nm, somewhat higher than that found in mesoporous iron phosphate nanotubes [12]. A similar peak in the same range was also observed for m-LFP; the profile, however, was somewhat more complex owing to the presence of a weak, broad peak centered at 8.0 nm. The mesopore diameter was estimated from [27,28] h
i1=2 dpore ¼ cd100 rVp = 1 þ rVp ½1 where Vp is the mesopore volume, ρ the density of pore walls, c a geometric parameter equal to 1.213 for cylindrical pores, and d100 the spacing corresponding to the [100] plane. If we assume pore to be cylindrical and pore walls to be built from FePO4·2H2O (average density ≈ 2.79 g·cm−3), then dpore is 3.55 nm for m-FP. This value is close to that derived from the isotherms. However, the corresponding value for m-LFP is 3.27 nm or 3.05 nm depending on whether pore walls consist of LiFePO4 or FePO4·2H2O—could not obtain the value for LiFePO4·2H2O—, respectively. These values are somewhat lower than those obtained from adsorption measurements. In any case, the pore system observed by XRD is more ill-defined than that of sample m-FP, which is consistent with a more heterogenous pore size distribution as suggested by the curve profile (see inset of Fig. 2b). TEM images of the m-LFP sample confirmed its porous structure. Fig. 3a reveals that the sample is made up of rounded-shape nanoparticles of heterogeneous size that keep an excellent connection. A closer examination of these particles (Fig. 3b and inset) reveals a distribution of mesopores, quite similar to that found in a mesoporous iron phosphate synthesized by Guo et al. [10] and Ding et al. [12]. In order to facilitate their discussion, FTIR spectra were split into two regions: 4000–1200 cm−1 which contained the peaks for the organic components of some samples, and 1300–250 cm−1 which contained information about the nature of the phosphate units. Fig. 4a shows the FTIR spectra for the mesoporous materials synthesized in this work (t-FP, m-FP and m-LFP). The presence of CTAB caused several additional peaks to appear in the spectra for t-FP, mainly two strong peaks assigned to asymmetric (2924 cm−1) and symmetric (2852 cm−1) stretching vibrations of C\CH2 bonds in methylene chains. Deformation vibrations of \CH2\ and \CH3 groups provided sharp bands in the region 1450–1500 cm−1. Furthermore, a weak band at 3020 cm−1 was observed which was assigned to C\CH3 asymmetric stretching and N\CH3 symmetric stretching vibrations in the solid surfactant. Compositional water was responsible for a broad band in the 3700–3050 cm−1 region and a peak at 1633 cm−1 corresponding to O\H stretching and deformation vibrations, respectively. An additional unexpected peak at 1744 cm−1 was observed. Treatment with lithium acetate in ethanol resulted in incomplete removal of the surfactant since the two peaks around

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Fig. 3. Selected TEM images of the m-LFP sample at different magnifications: (a) ×92,000, (b) ×340,000. Inset corresponds to a magnification of the area (b).

2900 cm−1 continued to be present, albeit with a lower intensity. The treatment also removed the impurity giving the peak at 1744 cm−1. By contrast, the exchange reaction introduced two new peaks at 1550 and 1418 cm−1 associated to COO− from acetate: [11,12,29] The remaining peaks were due to compositional water. Finally, chemical lithiation with LiI in acetonitrile caused peak shifts and broadening, and removed completely the surfactant. Fig. 4b shows the infrared spectra in the 1300–205 cm−1 region for samples a-FP, a-LFP, m-FP and m-LFP, in addition to Li3PO4 and LiFePO4. As shown, the vibrational spectrum was dominated by the fundamental vibrations of phosphate polyanions. Julien et al. [30] previously characterized FePO4 · 2H2O, chemically related with a-FP and m-FP, by infrared spectroscopy. These authors identified the symmetric stretching mode at ν1 = 990 cm−1 (P\O bond), the asymmetric doublet ν2 = 447–485 cm−1 (P\O bond) and two triplets associated to asymmetric O\P\O vibrations (ν3 = 1000–1085 cm−1 and ν4 = 570–640 cm− 1). The strongest peak for a-FP, located at 1051 cm−1, corresponded to ν3 vibration; the triplet, however, was ill-defined. Besides, the symmetric mode was shifted to a lower wavenumber (925 cm−1). The other signals were very broad and exact discrimination of doublet or triplet positions was very difficult—peaks at 545 and 594 cm− 1 (ν4) and 344 and 425 cm− 1 (ν2) were clearly seen, however. A similar spectrum was obtained for m-FP except for the lack of a ν1 band. However, bands at 561, 606 and 632 cm−1, associated to ν4, were well-defined. The lack or extreme weakness of ν1 suggests that the phosphate group was more weakly associated to the iron atoms in the mesoporous form owing to its special architecture [13,30]. On the other hand, the amorphous nature of a-FP resulted in poorer definition of the triplets at lower frequencies in a-FP compared with m-FP. Finally, the spectra for m-FP were similar to those obtained for mesoporous iron phosphates by other authors [13,30].

Fig. 4. FTIR spectra for the mesoporous samples prepared in this work; the low wavenumber spectra for LiFePO4 and Li3PO4 are included for comparison.

The FTIR spectra for chemically lithiated a-LFP and m-LFP were expected to be similar to those previously reported for LixFePO4·2H2O or LiFePO4 [30]. The spectra for the latter are included in Fig. 4b for comparison. As shown, the infrared fingerprints for phospholivine are not equivalent to the spectra exhibited by the phosphates involved in this work. Nevertheless, the main differences between the lithium free and lithium intercalated iron phosphates resemble those observed by Julien al. [30] in their study of dihydrated iron phosphate electrochemically intercalated with lithium. Thus, the stretching vibrations for this material were significantly shifted to higher frequencies, and new bands appeared below 500 cm−1, upon lithiation. Both effects were associated to the reduction of Fe(III) to Fe(II), and presence of lithium in the channels defined by the [FePO4] networks [30]. Unfortunately, the amorphous character of a-LFP and m-LFP resulted in the absence of new bands below 500 cm−1, so no information about the location of lithium could be obtained from their spectra. The spectrum for a-LFP exhibited vibrational shifts in the low frequency asymmetric vibrations ranging from 35 cm−1 for ν2 to 6 cm−1 for ν4. The shifts for m-LFP were rare and only a few reciprocal centimeters. One other effect of lithiating m-LFP was improved discrimination of the ν3 triplet. This indicates that inserting lithium boosts interactions between the iron and phosphate entities [13,30]. Fig. 5 shows the thermograms under N2 for the four phosphate samples. The curves for a-FP and m-FP were quite similar to those obtained by our group in earlier work [11]. The a-FP profile exhibited a continuous weight loss from r.t. to 400 °C associated to the loss of compositional and adsorbed water, followed by surface dehydroxylation. Crystallization of the α form of the anhydrous phosphate has

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a-LFP. This step, which was also observed by Prosini et al. [26] was not identified here. In addition, these authors found crystallization to yield LiFePO4 between 496 and 567 °C. Besides, acetate was detected by FTIR in m-LFP, so partial pyrolysis of this organic compound also occurred (between 200 and 600 °C) [11]. An amount of carbon (smaller than in m-FP) reduced Fe(III) to Fe(0) in the same temperature region. Finally, a significant weight loss, observed for the m-LFP over 700 °C, can be explained by a decomposition of the LiFePO4 structure, catalyzed by the presence of iron particles, releasing P2O5. 3.2. Crystalline phosphate based materials

Fig. 5. Thermogravimetric curves under N2 for different iron phosphates synthesized in this work.

been reported to occur close to 600 °C [26,31]. The system contained ca. 2.8 mol water per mol FePO4. The m-FP curve exhibited a smaller weight loss below 200 °C than did a-FP, as a consequence of dehydrating action of the ethanolic solution on the phosphate. The rest of the profile was influenced by the presence of small amounts of CTAB and acetate (both detected by FTIR), which introduced two new steps in the 200–600 °C region. Thus, desorption and decomposition of the template/acetate occurred in the 200–350 °C range and was followed by removal of residual organic compounds and surface dehydroxylation between 350 °C and 600 °C. Because the tests were carried out under nitrogen, some carbon was formed during heating and Fe(III) reduced to Fe(0), as discussed below. This reaction decreased the overall system weight since carbon was removed from the sample as CO. The formula of m-FP, obtained from the distinct weight loss steps under the assumption of a negligible presence of acetate ions, was FePO4·(CTAB)0.032·1.8H2O. Furthermore, Fig. 5 shows the thermograms for the lithiated products. The total weight loss for both a-LFP and m-LFP was lower than for the non-lithiated pristine compounds. Removal of water during contact with the acetonitrile solution of LiI accounts for this feature. The thermograms for the a-LFP sample are different to those previously obtained by Prosini et al. [26]. As a matter of fact, it lost 20% weight up to 600 °C, whereas their amorphous sample lost less than 10%. This is consistent with low water content in the reported material compared to a-LFP (2% versus 10%). Removal of iodine (a by-product of the lithiation reaction trapped between the phosphate particles) gave a well-defined peak in the 200–300 °C region for a-LFP and m-LFP. Additionally, a weight loss between 300 and 575 °C was detected for

The a-LFP and m-LFP systems were heated under dynamic vacuum at 550 °C for 3 h. Elemental analysis using a CHNS analyzer confirmed the lack of carbon in the resulting materials. The XRD patterns for both powders, a-LFP550 and m-LFP550, (Fig. 6) warrant several comments. Firstly, both systems crystallized at 550 °C, which made it unnecessary to raise the heating temperature in order to obtain ordered materials [26]. Secondly, both patterns contained peaks of Li3PO4, with higher intensity in the m-LFP550 system. The presence of this phosphate cannot be explained. The FTIR spectra discussed in the previous section suggested that this phase was absent from a-LFP and m-LFP. In order to obtain Li3PO4 free products, we reacted a-FP and m-FP with a more dilute LiI solution (0.1 M). However, Li3PO4 was always present in the crystallized material. A possible explanation for the occurrence of this salt is that the necessary heating for crystallization induces the partial dissociation of the lithium iron phosphate. Nevertheless, additional peaks in the 25–30° region were observed for a-LFP550. Unfortunately, up to now, its nature is unknown although it probably consists of iron oxide­derivatives in appreciable amounts. Finally, the occurrence of a strong peak located close to 45° for m-LFP550 and easily ascribed to metallic iron, is noteworthy. At this point, we should note that the thermograms for the m-LFP system revealed the presence of a carbothermal process below 600 °C involving residual carbon resulting from incomplete pyrolysis of CTAB and acetates. This residual carbon can reduce the potential iron oxides formed during the heating (similarly to a-LFP550) and results in the conspicuous presence of metallic iron in sample m-LFP550. Conversely, the iron oxide impurities vanish from the XRD pattern. An estimation of the ratio of the phases present in a-LFP550 and m-LFP550 products was obtained by weighting their relative XRD peak areas. The compositional results are 9.4% Fe2O3, 19.7% Li3PO4, 70.9% LiFePO4 (for a-LFP550) and 4% Fe, 26.6% Li3PO4, 69.4% LiFePO4 (for m-LFP550). Therefore for the same amount of phospholivine, the pristine mesoporous phosphate provides more Li3PO4 than the amorphous sample. The nitrogen adsorption/desorption isotherms for a-LFP550 and m-LFP550 are shown in Fig. 7. As expected, the thermal treatment provokes collapsing of the pores and, as a result, the m-LFP system

Fig. 6. XRD patterns for samples a-LFP550 and m-LFP550. (■) LiFePO4, (*) Li3PO4.

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3.3. Electrochemical properties of the amorphous materials in lithium half-cells Firstly, we investigated the electroactivity of non-crystalline lithium iron phosphates (a-LFP and m-LFP). To our knowledge, this is the first time amorphous “LiFePO4” and “LiFePO4 · 2H2O” have been tested as electrodes in lithium half-cells with an organic electrolyte. Fig. 9 shows the potentiostatic curves obtained between 3 and 4 V versus the Li/Li+ redox couple. During charge, the simplified reaction expected to take place in a LiFePO4 electrode is a two-phase process described by [16–26] II

þ

III

–

LiFe PO4 →Fe PO4 þ Li þ e :

ð1Þ

No similar scheme has been reported for LiFePO4·2H2O. However, the reverse reaction has been found to take place at 3 V and confirmed to follow the scheme [11,14,15,32–34] Fig. 7. Nitrogen sorption isotherms for samples a-LFP550 and m-LFP550. III

þ

–

II

Fe PO4 :2H2 O þ Li þ e →LiFe PO4 :2H2 O: does not retain its mesoporosity. However, some textural properties like pore volume or specific area surface, which are both beneficial properties for electrochemical purposes, are significantly bigger for m-LFP550 than for a-LFP550 (Table 2). Finally, the examination of both materials by means of TEM microscopy (Fig. 8) confirms their morphological differences. Whereas m-LFP550 is constituted by round shaped nanoparticles (average radii smaller than 100 nm) leaving pores between them, the a-LFP550 particles are several times bigger and have well defined edges typical of other LiFePO4 materials [16]. These observations agree with the textural parameters collected in Table 2.

ð2Þ

However, based on the curves in Fig. 9, none of these phosphates is electrochemically active versus lithium; in fact, no oxidation or reduction bands were observed during charge and discharge processes, respectively. The broad, weak peak in the first charge of m-LFP is more consistent with mitigated, irreversible oxidation of the stainless steel grid. These results unambiguously indicate that lithium ions require a well-defined pathway for deinserting/inserting in the LiFePO4 structure [35,36] (note that the iron phosphate units in a-LFP and m-LFP are disordered in the long range); therefore, the LiFeIIPO4·2H2O phase must be created in situ in the electrode in order to be electrochemically active. 3.4. Electrochemical properties of the crystalline materials in lithium half-cells.

Fig. 8. TEM image of the (a) m-LFP550 and (b) a-LFP550 samples.

Unlike the amorphous materials, the crystalline systems a-LFP550 and m-LFP550 exhibited electroactivity (Fig. 9) in the lithium half-cells, albeit to a different extent. Thus, the potentiostatic tests on m-LFP550 based cell revealed a strong oxidation peak at ca. 3.55 V and mainly associated to reaction (1), although the oxidation of iron present in this sample is also possible in this voltage range [37]. During the reduction process, a cathodic wave centered at 3.35 V was observed. The difference between the centers of the two peaks (ΔEp) was 200 mV, which is consistent with a LiFePO4-based lithium ion electrode [16–26]. However, peak profiles were strongly retained upon cycling. Furthermore, the electrochemical behavior of a-LFP550 was not what we had expected. The first cycle acts as a conditioning step. Thus, its I/V profile is rather different to those observed in the second and subsequent cycles. It is consistent with the observed lower surface area for this sample, which implies that complete electrolyte impregnation of the electrode requires more time than for the m-LFP550 sample. The first charge gives a very broad, incomplete peak at 4.0 V seemingly centered at 3.65 V, whereas the first discharge provides a broad reduction peak that is well defined and centered at 3.3 V. The difference between the oxidation and reduction potentials (ΔEp = 350 mV) indicates that reaction (1) is more kinetically hindered [18,21–26] than it is for m-LFP550, which can be easily associated to the presence of insulating impurities (FexOy, Li3PO4) together with the lack of conductive iron in a-LFP550. Note that ΔEp was 75% higher than for m-LFP550. In addition, the gravimetric peak intensities were smaller and, even upon cycling, only amounted, at most, to half of that of m-LFP550. After the first and conditioning cycle, the curve profiles were similar and ΔEp was reduced to 280 mV; overall, however, a-LFP550 was much less electroactive than m-LFP550. Selected galvanostatic charge/discharge curves for the LiFePO4-based electrodes are shown in Fig. 10. In agreement with the potentiostatic results, the first charge curves exhibited a voltage pseudoplateau located

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Fig. 9. Potentiostatic tests on Li/LiPF6(EC,DMC) half-cells containing the phosphates prepared in this work as positive electrodes. Voltage steps: 10 mV·10 min−1.

at 3.5 V and 3.6 V for m-LFP550 and a-LFP550, respectively. The charge capacities obtained were also found to depend on the material, with values of 107 and 92 mAh·g−1 for m-LFP550 and a-LFP550, respectively, being lower compared to that reported for carbon-coated lithium iron phosphates [18–25]. However, these capacities reach 153 and 131 mAh·g−1 for m-LFP550 and a-LFP550 respectively when referred to the mass of LiFePO4 (70% of the materials). In addition, the first discharge curves for both cells, shown in Fig. 10, indicated a limited reversibility of reaction (1). The two profiles contain a wide pseudoplateau, close to 3.35 V, the a-LFP550 electrode showing a high polarization. The discharge capacities, 58 and 44 mAh·g−1 for m-LFP550 and a-LFP550 respectively, are also limited. Referred to the mass of LiFePO4, the

Fig. 10. Selected galvanostatic charge/discharge cycles of Li/LiPF6(EC,DMC) half-cells containing a-LFP550 (blue line) or m-LFP550 (red line) as positive electrodes, under a C/7 regime.

capacities shift to 82 and 63 mAh·g−1 for m-LFP550 and a-LFP550 respectively. After the first cycle, the pseudoplateau in the a-LFP550 electrode charge curve is developed at 3.5 V, which agrees with the conditioning undergone in the first charge (Fig. 9). By contrast, the successive charge profiles of the m-LFP550 electrode remain similar. However, after 25 cycles, and even as the m-LFP550 kept providing the highest values amongst the series of electrodes implied, the capacities provided are in the range of 34–51 mAh·g−1 (49–73 mAh·g−1 if referred to LiFePO4) which is below those found in carbon-coated systems [18–25]. Electrochemical impedance spectroscopy was applied on a-LFP550 and m-LFP550 electrodes at OCV in order to shed more light on the difference in their behaviors. The corresponding Nyquist plots, shown in Fig. 11, were found to fit well a Randles classic model containing an electrolyte resistance, Rel, a charge transfer resistance, Rct, a Warburg element, W0 and a double layer capacitance, Cdl, equivalent to a constant phase element. The characteristic frequencies of the processes associated to the Rct/Cdl element in the second part of the circuit are 68.34 Hz and 14.24 Hz for a-LFP550 and m-LFP550 respectively. Such values are similar and discard the occurrence of additional interfacial process. Since the specific surface area is bigger for m-LFP550 than for aLFP550, the impedance values are also bigger for the first system. For an adequate discussion, the electrical parameters are normalized by the theoretical surface area exposed to the electrolyte (despite the fact that inner layers of the electrode material are not exposed to the electrolyte). It is noteworthy that the electrolyte resistance is smaller (0.01 Ω·cm−2) for the m-LFP550 than for a-LFP550 (0.2 Ω·cm− 2) which proves better electrolyte impregnation in the porous and higher surface area structure than in the compact a-LFP550. However, the values found for Cdl were somewhat smaller for m-LFP550 (4.16 μF·cm−2 vs. 5.62 μF·cm−2 for a-LFP550). Please note that both composite electrodes contain carbon black, although surely not dispersed in a similar manner because of the different morphology of a-LFP550 and m-LFP550. The contribution of carbon black to the total capacitance can explain the unexpected smaller values of Cdl found for m-LFP550. Finally, Rct is somewhat smaller for m-LFP550 (0.85 Ω·cm−2 vs. 1.36 Ω·cm−2 for a-LFP550),

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increase electrode resistivity. As a result, the charge/discharge plateaux should be expanded. Indeed, the m-LFP550 based electrode always behaved more effectively than a-LFP550. We believe that this was a result of the combination of appropriate textural and morphological properties in LiFePO4 in addition to enhanced lithium diffusion into the network. In fact, the presumed favorable effect of iron is compromised during the first charge since metallic iron must be removed from the electrode in the first charge up to 4 V [37] although some iron could remain in the bulk material in a similar manner to copper in copper/LiFePO4 composites [43]. Therefore, after a few cycles, the better electrochemical response of m-LFP550 versus lithium compared to a-LFP550 must be a result of other factors such as porosity, surface area, smoother particles morphology and faster lithium diffusion (Table 2) [21–25]. 4. Conclusions Fig. 11. Nyquist plots recorded for a-LFP550 (blue squares) or m-LFP550 (red circles) as positive electrodes at OCV. Lines resulting from the fitting (blue and red for a-LFP550 and m-LFP550, respectively) are also included.

which is in line with the expected increased conductivity in m-LFP550 by virtue of iron presence. However such increase is limited since a bigger amount of Li3PO4 is found in this sample than in a-LFP550. The observed ΔEp in the potentiostatic and the galvanostatic tests can also reflect the presence of a kinetic factor resulting from lithium insertion [18,38]. The use of the Vaccaro's method [11,39] strongly supported the latter statement. The method analyzes the limit frequency, f*, that separates the Warburg domain (semi-infinite diffusion, defined by a 45° zone in the Nyquist plots) from the lithium ions accumulation (limited diffusion, defined by a Zim vs Zre variation approaching 90° for a non-porous material) (Fig. 11). This frequency can be used to estimate the lithium ion diffusion coefficient (D = L2 / 2πf*), assimilating L, the lithium diffusion path, to the average particle radius. The estimated values are 1.3 · 10− 15 cm2·s−1 and 8.8 · 10−15 cm2·s−1 for a-LFP550 and m-LFP550 respectively. These values are somewhat lower than those found in the phosphate based materials [18] and can explain to some extent their low capacities. Indeed, the coefficient for m-LFP550 is 7 times bigger than for a-LFP550, which agrees with the smaller polarization observed in the galvanostatic and the potentiostatic curves in the latter material and with its better capacity retention. In spite of some advantageous textural properties and particle size, other features account for the limited performances shown by the phosphates under study. Regarding the technological advances in the development of lithium iron phosphates, it seems that the most important feature is the lack of an adequate, conductive carbon coating onto the phosphate particles [18,21–26,40]. In fact, the performance of mesoporous lithium iron phosphates reported in the literature, none of which showing pores order, depends greatly of the carbon content [24]. For instance, at least a 3.8% weight ratio is required for a suitable electrochemical response [25]. Thus, further work on our samples should be focused on the production, after or during the formation of the mesoporous lithium iron phosphate, of a carbon film by filling the pores with a carbon precursor and by heating the whole system at mild temperatures. Nevertheless, firstly Delacourt et al. [41] and later, Dominko et al. [42] pointed out that carbon-free electrodes could also show adequate cycling behaviors, provided that a “carbon phase or other electronic conductor only point contacts each individual active particle”. On this basis, the electrochemical performance of the phosphates synthesized in this work should be improved by eliminating the insulating impurities (Li3PO4 and/or iron oxides) that reduce the contact points available to the carbon black present in the electrode mixture and

In this work, we suggest a novel mesoporous approach in the field of phospholivines electrode material preparation. It proves successful in obtaining LiFePO4 and could be extended to the synthesis of others systems such as LiMnPO4, Li3V2(PO4)3 or LiFexMn1 − xPO4 solid solutions. We have described a simple method for building ordered amorphous mesoporous LiFePO4, starting from ordered mesoporous FePO4·2H2O and a LiI solution in acetonitrile for simultaneous Li+ insertion and Fe3+ reduction. Heating the mixed phosphate for transformation into a crystalline compound at mild temperatures (550 °C) under vacuum causes partial decomposition and the formation of Li3PO4. Also, the solid contains some Fe originating from carbothermal reduction of hematite, another product intrinsic to the phosphates crystallization. This crystalline compound showed better electrochemical response in lithium half-cells than another crystallized product prepared from a disordered porous material. The improved performance seems to be bound to enhanced lithium ion diffusion inside the network, rather than to the presence of conductive iron nanoparticles. In addition, its porous structure and smooth surface could account for better electrolyte impregnation and suitable lithium ions transport. However, an efficient electronic transport upon cycling is not guaranteed in these systems and therefore, the crystallized material-based half-cells showed small capacities and noticeable capacity fading. In order to overcome these drawbacks, the introduction of a suitable carbon film surrounding the phosphate particles and/or the elimination of impurities hindering the contact between the percolating agent (carbon black) and the phosphate can be claimed. Finally, this work supports the idea that using carbon-free mesoporous precursors is a suitable alternative for the preparation of LiFePO4 lithium-ion electrodes. Acknowledgment This work was supported by the Ministerio de Ciencia e Innovación (MAT2010-16440) and Junta de Andalucía (Group FQM 175). The authors are also grateful to F. Gracia (Electronic Transmission Microscopy Service, SCAI) and J. García (NIR/MIR Spectroscopy Service, SCAI) at the University of Córdoba, for their help in collecting the TEM images and the infrared spectra respectively. JSP is indebted to C. Bihr who kindly corrected the English language of this manuscript. References [1] C.T. Krege, M.E. Leonowic, W.J. Roth, J.C. Vartuli, J.S. Beck, Nature 359 (1992) 710. [2] A. Sayan, P. Liu, Microporous Mesoporous Mater. 12 (1997) 149. [3] U. Sciela, D. Demuth, R. Leon, P. Petroff, G. Stucky, K. Unger, F. Schüth, J. Chem. Soc. Chem. Commun. (1994) 1387. [4] M. Antonietti, G.A. Ozin, Chem. Eur. J. 10 (2004) 28. [5] A.S. Arico, P.G. Bruce, B. Scrosati, J.M. Tarascon, W. van Schalkwijk, Nat. Mater. 4 (2005) 366. [6] L. Guo, H. Arafune, N. Teramae, Langmuir 29 (2013) 4404. [7] , see for example. K. Moller, T. Bein, Chem. Mater. 10 (1998) 2950. [8] H. Zhou, S. Zhu, M. Hibino, I. Honma, M. Ichihara, Adv. Mater. 15 (2003) 2107.

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