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Solution synthesis of nanometric layered cobalt oxides for electrochemical applications
Electrochimica Acta 66 (2012) 306–312
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Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta
Solution synthesis of nanometric layered cobalt oxides for electrochemical applications Xavier Pétrissans a , Angélique Bétard a , Domitille Giaume a , Philippe Barboux a,∗ , Bruce Dunn b , Lorette Sicard c , Jean-Yves Piquemal c a
Chimie de la Matière Condensée de Paris, CNRS UMR 7574, Chimie-ParisTech, 11 rue Pierre et Marie Curie, 75005 Paris, France Department of Materials Science and Engineering, University of California Los Angeles, Los Angeles, CA 90095, USA c Laboratoire ITODYS, Université Paris-Diderot, Bâtiment Lavoisier, 15 rue Jean de Baïf, 75205 Paris Cedex 13, France b
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Article history: Received 12 December 2011 Received in revised form 24 January 2012 Accepted 28 January 2012 Available online 5 February 2012 Keywords: Nax CoO2 nanoparticles Colloidal oxide Ionic exchange Pseudocapacitor Cyclic voltammetry
a b s t r a c t Dispersed Na0.6 CoO2 ·yH2 O and Li0.5 CoO2 powders have been obtained at room temperature by rapid precipitation in aqueous solutions of LiOH or NaOH in the presence of a strong oxidizer. The precipitates are well crystallized and consist of nanoscale platelets with high specific area (above 100 m2 /g). The Li0.5 CoO2 phase is stable in aqueous electrolytes whereas the Na0.6 CoO2 ·yH2 O rapidly converts to CoOOH in neutral electrolytes or pure water. It also transforms to anhydrous Na0.6 CoO2 upon drying at moderate temperatures. Electrochemical studies show that at slow sweep rates the Na0.6 CoO2 ·yH2 O can store large amounts of charge in 10 M NaOH from a combination of both faradic and capacitive reactions. © 2012 Elsevier Ltd. All rights reserved.
1. Introduction The large scale application of batteries to electrical vehicles requires an improvement in their power density in order to reach an acceptable driving range. Moreover, the energy storage of these systems decreases considerably when the power demand increases [1] while safety problems arise due to heat dissipation and outof-equilibrium electrochemical side reactions [2]. To address the problem of peak power management, batteries can be coupled with complementary devices such as electrochemical capacitors, also known as supercapacitors, that provide high power but for short durations because of low energy density. These devices are based on the development of electrochemical double-layers that store charge at the solid-electrolyte interface and do not involve chemical reactions as in the case of batteries. Electrochemical capacitors possess a number of attractive features compared to batteries including higher power, shorter charging times and much longer cycle life. One method of increasing the energy density of electrochemical capacitors is to have reversible redox processes at the surface of high surface area oxide films or nanoparticles [3,4]. The resulting pseudocapacitance increases the level of energy storage by an order of magnitude or more, with the best properties exhibited by
∗ Corresponding author. Tel.: +33 1 53 73 79 25. E-mail address: [email protected] (P. Barboux). 0013-4686/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2012.01.104
hydrated RuO2 oxides [5,6]. The cost of RuO2 is prohibitive and for that reason, a number of less expensive candidates such as transition metal oxides or hydroxides have been investigated (MnO2 [7–9], Ni(OH)2 [10], Co3 O4 [11,12], CoOOH [13], FeOOH [14], CuO [15], V2 O5 [16–18], TiO2 [19,20], Zn2 CoO4 [21], LiCoO2 [22], etc.). Manganese based oxides have probably attracted the largest interest due to their low cost, and low toxicity. Lamellar MnO2 structures such as birnessite have been the subject of many studies due to their ability to reversibly intercalate monovalent cations in aqueous electrolytes [8]. To achieve significant capacity (on the order of hundreds of F/g or C/g in capacity or charge storage respectively), it is necessary to optimize the synthesis, chemistry and microstructure so that high surface area can be combined with favorable diffusion pathways for both ions and electrons to enable redox reactions to occur at the interface [7,19,23].More generally lamellar structures can achieve good ionic mobility and good exchange properties. Lamellar MnO2 structures such as birnessite have been the subject of many studies due to their high ability to reversibly intercalate monovalent cations in aqueous electrolytes [8]. Layered hydroxides also demonstrated large specific capacitance and good cycle reversibility [24,25]. In this paper, we particularly focus on layered cobalt oxides, especially on the Na0.6 CoO2 phase [26] and on its hydrated polymorphs [27]. The good electronic conductivity of this material associated to its fast ion exchange properties make it a good candidate as a model material to study the effect of ionic mobility on the distribution of pseudocapacity between double
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layer effects and faradic insertion mechanisms as a function of the cycling rate and the nature of the ions in the interlayers. The Ax CoO2 oxides exhibit lamellar structures with sheets of edge-sharing CoO6 octahedra separated by layers of alkali ions. Depending on the alkali ion and stoichiometry of the compound, the stacking of these sheets may differ and lead to octahedrically coordinated alkali ions as in LiCoO2 or to prismatic coordination around the alkali ion as in Na0.6 CoO2 [28,29]. The various Ax CoO2 phases can be obtained by conventional reaction at high temperature between cobalt oxide (Co3 O4 ) and alkali carbonates [27,30,31,34]. LiCoO2 , which has been studied extensively because of its application in lithium-ion batteries [30,32], can be readily synthesized at low temperature by various routes. Cobalt oxyhydroxide, CoOOH, can be reacted with lithium hydroxide solutions [33,34], Co(OH)2 can be refluxed in LiOH without addition of an oxidizer to transform Co(II) into Co(III) [35], or homogeneous lithium cobalt citrate precursors can be decomposed at 400 ◦ C [36]. These approaches are not successful for Nax CoO2 , or lead to multiphase products [32,37]. The formation of the sodium phase is much more difficult than that of LiCoO2 due to the different ionic radii and electronegativity of Li and Na. The electronic distribution is rather different in both phases with stronger hybridization of O and Co orbitals in the case of the sodium phase [38]. As a result, the sodium phase can become metallic whereas the lithium phases are semiconductors [39]. For this reason, it could be proposed as a conductive additive to the Nickel electrode in Ni–MH batteries [40,41].In this work, we present the direct synthesis of nanoparticles of lamellar cobalt oxides and discuss their ability to store energy by capacitive or faradic effects at different rates. The method described here for the synthesis of Mx CoO2 (M = Li, Na) is a low temperature route based on using an alkaline oxidizing medium. The Pourbaix diagram of the cobalt oxides, recently revised by Chivot et al. [42], shows that the MCoO2 form (M = H, Na, etc.) is stable only at high potential and high pH. This method has already been used to synthesize various transition metal oxide compounds including LiCoO2 , Li0.6 NiO2 or Nax MnO2 [43,44]. To the best of our knowledge, the direct precipitation of Na0.6 CoO2 has not been reported in the literature. In addition, many high values of capacitance have been reported on different cobalt oxides in the literature [11,13], however, most of these results have been obtained on thin films or mesoporous structures supported on various substrates, which do not allow to separate the structural and microstructural contributions. It is thus necessary to develop a method to assess the intrinsic electrochemical properties of the materials. The deposition of particles onto glassy carbon electrodes developed in this work allows one to determine the intrinsic properties of the material because the measurement is taken without the presence of either a conductive additive or binder. 2. Experimental 2.1. Synthesis of electrochemically active materials 2.1.1. Reagents All chemical reagents were 99.9% grade, commercially available at Aldrich chemicals and used as received. Glassy carbon electrodes were purchased from Alfa-Aesar. 2.1.2. Material synthesis via alkaline oxidation In a typical synthesis, 22 g of solid NaOH (respectively 6 g LiOH) are dissolved in 50 mL of deionized water in a 250 mL three-neckflask under O2 flow. A 10 M NaOH solution is obtained (respectively 3 M LiOH which is near the saturation limit). The flask temperature is kept around 25 ◦ C, and 0.7 mL of Br2 (14 mmol) are added to the solution which turns yellow. Immediately afterwards, 5 mL of a red
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1 M Co(NO3 )2 solution is added dropwise over 10 min. A black precipitate forms immediately which is recovered after filtration using a sintered glass frit under vacuum. The precipitate is dried at 120 ◦ C for 24 h, producing a black hygroscopic powder. Subsequent washing with pure water to remove the adsorbed NaOH solution could be performed but would affect the nature of the resulting phase as will be discussed in the following section. 2.2. Characterizations 2.2.1. Structure X-ray diffraction (XRD) was carried out on the various synthesized powders using the X’Pert PRO, PANanalytical diffractometer ˚ equipped with an X’Celerator detector and (Co (K␣ Co ) = 1.78901 A) used in the Bragg-Brentano geometry. 2.2.2. Composition analysis Elemental analysis was performed by ICP–AES (Varian Vista Axial and ThermoFisher iCAP 6000). The cobalt mean oxidation state was determined by iodometric titration. 20 mg of powder was dissolved in 20 mL of a 6 M HCl solution containing excess potassium iodide (500 mg). The I2 resulting from the oxidation of iodide by Co3+ and Co4+ was back-titrated with a sodium thiosulphate solution (10−2 M) using starch as the indicator. 2.2.3. Particle size and morphology Particle size and morphology were determined for the synthesized powders using a Transmission Electron Microscope (Jeol-JEM-100CXII operating at 100 kV). In these experiments, the powder was dispersed in ethanol and a drop placed on a 200mesh carbon-coated copper grid. Scanning Electron Microscopy measurements using a Hitachi S2500 IMIX PGT microscope were also performed on powder pasted onto a carbon tape. Surface areas were measured using a gas adsorption analyzer (BEL Japan, Belsorp-max). The Brunauer, Emmett, Teller (BET) analysis was used to characterize the powder samples which were degassed 120 ◦ C under vacuum during 10 h prior to measurement. 2.3. Electrochemical measurements 2.3.1. Glassy carbon electrodes Electrodes were prepared following the method described by Wang et al. [20] by depositing the powder onto a glassy carbon substrate. This method allows direct contact between each oxide grain and the current collector and enables one to determine electrochemical properties of materials without the use of carbon additives or binder. First, a glassy carbon substrate (1 cm × 2 cm) is treated for 15 min in an oxygen plasma to become hydrophilic. A small amount of colloidal dispersion (7.0 L) of the material in the electrolyte to be used (vide infra) is then directly deposited onto the substrate and allowed to dry. To bind the particles to the substrate, it is necessary to anneal the samples at 300 ◦ C in air for one hour except for the hydrated Nax CoO2 ·yH2 O composition which can be annealed only at 220 ◦ C for 40 min. Still, some unbound particles are lost when the electrodes are dipped in the electrolyte. The quantitative loss is determined by chemical analysis of the electrolyte solution after acidification and ICP titration. Finally, the amount of oxide on the electrode is calculated from the concentration of the colloidal dispersion and the volume deposited, corrected with the cobalt concentration found in the electrolyte after the study (about 10–20%). During measurement, Teflon tape covers the other side of the glassy carbon substrate to avoid any leakage current. 2.3.2. Electrochemical studies Cyclic voltammetry (C-V) measurements were recorded using an Autolab PGSTAT 30. A three-electrode geometry was used with a
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large graphite sheet serving as the counter-electrode and a Ag/AgCl electrode as the reference for the glassy carbon experiments. In the C-V experiments, the sweep rate was varied between 0.5 and 100 mV/s and the total amount of charge storage was determined by integrating and averaging both the oxidation and reduction currents. The contribution of the glassy carbon is very low (5 C/g) and has been subtracted directly from the voltammograms. 3. Results 3.1. Synthesis 3.1.1. Crystalline structures The powders synthesized under various conditions were analyzed by XRD and are shown in Fig. 1. Just after synthesis in NaOH, without washing or drying, the material is a wet black paste which appears to be amorphous from XRD (Fig. 1a). Subsequent drying at 120 ◦ C allows the material to crystallize as a single phase corresponding to a Nax CoO2 structure (0.5 < x < 1) (Fig. 1b). All the diffraction peaks can be attributed to the rhombohedral Na0.6 CoO2 structure (JCPDS 00-071-1281) which is composed of CoO6 octahedra linked by edges and separated by an interlayer space in which sodium ions are in a trigonal prismatic environment [28]. The c-axis parameter is calculated to be 16.7 ± 0.3 A˚ based on the XRD results, which is slightly larger than the 16.53 A˚ reported for Na0.6 CoO2 [28]. The broadening of the X-ray peaks indicates that the crystalline domains are quite small. From the Debye-Sherrer formula, we estimate a crystallite size of ∼10 nm. It is interesting to note that some of the peaks are significantly broader than others, especially the (0 0 3) plane around 20◦ . In addition to a size effect, the shape of this (0 0 3) line may be influenced by other contributions, such as a cation deficiency in the interlayer space as discussed by Butel et al. [45]. The presence of cation vacancies can enhance repulsion between the oxide layers, leading to an increase in lattice strains, which is detected by the (0 0 l) line shape. The presence of both sodium ions and protons in the interlayer space can also generate extended structural defects that are caused by segregation of the Nax CoO2 -rich and HCoO2 -rich zones. Abundant washing with water leads to a diffraction pattern (Fig. 1d) which is attributed to a pure -HCoO2 phase (JCPDS 01072-2280). The c-axis cell parameter from the X-ray diagram gives 13.2 ± 0.1 A˚ while it is reported to be 13.13 A˚ for HCoO2 [45]. The
narrowing of the peak corresponding to the (0 0 3) line also indicates better periodicity along the (0 0 l) lines. These results are confirmed by the elemental analysis performed on the samples whose diffraction diagrams are given in Fig. 1b and d. The detailed composition of the powder after synthesis, drying and without washing (Fig. 1b), Na0.6±0.1 CoO2 ·(0.8 ± 0.5)H2 O was determined by linking the Na (12.17 wt%) and Co (44.68 wt%) contents given by ICP with the mean oxidation state of cobalt given by iodometric titration. If we consider that the product obtained after soaking in water (Fig. 1d) should be also hydrated, Nax Hy CoO2 ·zH2 O, the iodometric titration linked with the ICP analysis gives us a composition of Na0.12 H0.83 CoO2 ·(0.4)H2 O. As the (Na + H) content is close to 1 (0.95), we expect nearly all cobalt in this phase to be Co3+ . For simplicity, in this paper we refer to the Na0.6±0.1 CoO2 ·(0.8 ± 0.5)H2 O phase as Na0.6 CoO2 , and to the Na0.12 H0.83 CoO2 ·(0.42)H2 O one as HCoO2 . When the synthesis takes place using LiOH, the resulting material has the same rhombohedral Nax CoO2 structure but with a composition of Li0.5±0.1 CoO2 . The diffraction pattern (Fig. 1c) shows the similarity between phases. As discussed in the introduction, the Lix CoO2 phase is much more stable than Nax CoO2 and it does not transform upon soaking in water. The oxidizer used in the LiOH synthesis is bromine (Br2 ) whose electrochemical potential is independent of pH. In the case of Lix CoO2 , hydrogen peroxide could also be used although its oxidizing potential is weaker at high pH. With this mild oxidizer, we have obtained a Lix CoO2 composition with x ≈ 1. 3.1.2. Particle size and morphology Electron micrographs of the amorphous paste deposited onto the microscope grids without washing show that it is composed of very small crystallites of approximately 30 nm long and 6 nm thick, embedded in an amorphous coating (Fig. 2a). Those crystallites appear to be lamellar, with an interlayer space of 6.6 A˚ as determined with Digital Micrograph software. This is consistent with the hydrated Nax CoO2 phase [45]. The samples dried before dispersion onto the microscope grids give an interlayer spacing of 5.51 A˚ in good agreement with the Na0.6 CoO2 composition [45]. The specific surface area of the unwashed precipitates measured by the BET method is 125 ± 30 m2 /g. This value is equivalent to the surface area calculated for spherical 10 nm non-porous particles of Na0.6 CoO2 (density = 4.55 g/cm). The Li0.5 CoO2 samples exhibit a somewhat greater surface area of 160 ± 20 m2 /g. Thus, the oxidizing alkaline conditions described in this work enable one to obtain high surface area crystalline nanoparticles of Na0.6 CoO2 ·yH2 O which cannot be obtained through high temperature routes. These nanoparticles are highly reactive and may transform upon washing or drying sequences. The large distribution of surface area (±25%) is related to different washing sequences or drying processes (between 80 ◦ C and 120 ◦ C) and can be associated with a variation in the packing of the crystallites at a macroscopic scale, as seen in Fig. 2c and d. This microstructure is favored by the particle morphology in platelets: aggregates of approximately 100 nm form, which themselves pack together into larger objects. 3.2. Stability of Na0.6 CoO2 in electrolyte solutions
Fig. 1. X-ray diffractograms of precipitates obtained with NaOH (a) just after synthesis; (b) after drying at 120 ◦ C during 24 h; (c) with LiOH; and (d) with NaOH after washing with deionised water. Reference diagrams of the -Na0.6 CoO2 phase (JCPDS 00-071-1281), the LiCoO2 phase (JCPDS 04-006-3441) and the HCoO2 phase (JCPDS 01-073-1213) are also shown.
An important issue with the Na0.6 CoO2 is whether immersion in typical electrolytes causes a phase transition to occur. The Na0.6 CoO2 powder was soaked in neutral water as well as in standard electrolyte solutions such as sulphates and hydroxides. In the discussion below, the X-ray patterns shown in Fig. 3 are compared to that of Na0.6 CoO2 (Fig. 3a). Upon soaking in water, the HCoO2 phase is obtained, similar to the effect of washing as indicated in Fig. 1d. With Na2 SO4 , as in pure water, there is a
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Fig. 2. Transmission electron microscopy images of (a) Na0.6 CoO2 ·yH2 O particles just after synthesis; (b) and (c) of Na0.6 CoO2 after drying at 120 ◦ C during 24 h; (d) scanning electron microscopy images of Na0.6 CoO2 powders after synthesis.
complete transformation to Hx CoO2 which indicates that the ion exchange of Na+ for H+ is strongly selective for the proton (Fig. 3c). These results have a number of similarities to those reported by Butel et al. who observed a structure collapse of Na0.6 CoO2 at low pH after Na+ /H3 O+ exchange; once all the alkaline cations were removed from the interlayer space, the -HCoO2 phase was obtained [45]. Soaking in 1 M NaOH solution also leads to the formation of the Hx CoO2 phase (not shown), however, in concentrated NaOH (10 M), the Na0.6 CoO2 phase is unchanged (Fig. 3d). For this reason, all the electrochemical experiments were carried out in 10 M NaOH. With another electrolyte, Li2 SO4 (neutral pH), the material also undergoes ion exchange, probably to a composition of Lix Hy CoO2 (Fig. 3b). The LiCoO2 phase is far more stable than the Nax CoO2 [34] and remains unchanged in water.
3.3. Electrochemical properties
Fig. 3. X-ray diffractograms of the Na0.6 CoO2 phase (a) after synthesis; and after soaking in (b) Li2 SO4 (0.5 M); (c) Na2 SO4 (0.5 M) and (d) NaOH (10 M). Reference diagrams of the Na0.6 CoO2 phase (JCPDS 00-071-1281), the LiCoO2 phase (JCPDS 04-006-3441) and the HCoO2 phase (JCPDS 01-073-1213) are present for comparison. * refers to Na2 CO3 impurities that appear after soaking in the NaOH 10 M solution.
Studies on glassy carbon demonstrate the intrinsic properties of the material, as pure oxide is deposited onto glassy carbon, without any additive. Cyclic voltammetry (CV) was used to characterize the electrochemical properties of the different samples. Fig. 4 shows cyclic voltammograms taken on the third cycle for samples deposited onto glassy carbon at sweep rates ranging from 1 mV/s to 100 mV/s. The samples include the as-synthesized Na0.6 CoO2 ·yH2 O powder without washing and drying (Fig. 4a), the same material
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Fig. 4. Cyclic voltammograms at different sweep rate (1, 5, 10, 50 and 100 mV/s) for (a) Na0.6 CoO2 ·yH2 O in NaOH 10 M; (b) Na0.6 CoO2 performed in NaOH 10 M; (c) Li0.5 CoO2 performed in LiOH 3 M; and (d) HCoO2 performed in NaOH 10 M. The specific current is given per gram of oxide.
after heating at 120 ◦ C to yield Na0.6 CoO2 (Fig. 4b), and after ion exchange in water to yield HCoO2 (Fig. 4d). For all these materials, the electrochemical studies were performed in 10 M NaOH as electrolyte. In addition, we characterized a LiCoO2 powder (Fig. 4c); the electrolyte used for that material was 3 M LiOH. To provide an appropriate comparison among the different samples, the currents are normalized to the mass of the oxide. In establishing the voltage window for the CV experiments, the positive limit was taken as the onset of water oxidation whereas the lower limit was taken at an appropriate potential that gave good cycle-to-cycle reproducibility. In characterizing these samples over tens of cycles, we noticed a slow decrease in current which is mainly associated with detachment of particles from the glassy carbon substrate. This problem was particularly evident with Na0.6 CoO2 ·yH2 O. Thus, the results for different sweep rates had to be collected with different samples of this material and this explains the error bars given from Fig. 5a for the Na0.6 CoO2 ·yH2 O composition. For all other materials the curves could be obtained with a single sample. The curves observed for Nax CoO2 (Fig. 4b) and HCoO2 (Fig. 4d) samples do not show oxidation or reduction peaks, features which are characteristic of an electrochemical reaction. This current–voltage response is the signature of a pure capacitance. In contrast, the curves obtained for the Na0.6 CoO2 ·yH2 O sample (Fig. 4a) or the Li0.5 CoO2 phase (Fig. 4c), exhibit broad oxidation and reduction peaks in addition to the large capacitive effect. At higher sweep rates, the oxidation and reduction peaks are less evident. Overall, the cyclic voltammetry measurements can be considered as being a large capacitive response in combination with faradic reactions.
The sweep rate dependence on charge storage (reported in C/g) for the different samples is shown in Fig. 5. In all cases, the amount of charge storage increases at slow sweep rates because of contributions from diffusion-related processes such as intercalation [20]. Interestingly, for Na0.6 CoO2 ·yH2 O, the charge stored at slow sweep rates can be as high as 900 ± 250 C/g of oxide, similar to that obtained for ruthenium oxide [5,6]. This value is comparable to the charge calculated, ∼920 C/g, for one-electron redox reactions for Co3+ /Co4+ with this material. About half of the measured charge storage can be attributed to the extraction of 0.5 sodium in the Na0.6 CoO2 phase (∼460 C/g). Another contribution, the electrical double layer capacitance at the electrode/electrolyte interface, can be ascribed to the density of surface ionisable groups that can reach a limit value of 50 C/cm2 . Based on the specific surface area of 125 m2 /g, this process yields an estimated capacitive charge of 65 C/g. The remaining amount of charge storage, in the range of 400 C/g, is attributed to a pseudocapacitance contribution. Electrochemical energy storage studies with transition metal oxides have shown that pseudocapacitive charge storage becomes more prominent with nanoscale particles [20,46]. Moreover, in the recent V2 O5 work, pseudocapacitive processes were identified as the charge storage mechanism which enabled the specific capacity to exceed that associated with traditional intercalation [46]. The other materials that were investigated, Na0.6 CoO2 , Li0.5 CoO2 and HCoO2 exhibit charge storage values in the range of 200–300 C/g at slow sweep rate which are consistent with previous work [47,48]. At higher sweep rates (Fig. 5a), the stored charge decreases to values below 200 C/g. At the highest sweep rates of 100 mV/s, the Na0.6 CoO2 ·yH2 O and Li0.5 CoO2 phases still store up to 160
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the double-layer capacitance. Only the Na0.6 CoO2 ·yH2 O material exhibits a significantly greater level of charge storage (200 ± 70 C/g) which, as discussed above, indicates that the high level of Na+ mobility in these materials enables faradic processes to occur. 4. Discussion We have prepared layered alkaline cobalt oxides by a synthetic route based on room-temperature precipitation in an oxidizing medium. The method yields small crystallites with high specific surface area. The LiCoO2 nanoparticles are readily obtained and they are stable in neutral aqueous solutions whereas the sodium phase can only be obtained in a hydrated form, Na0.6 CoO2 ·yH2 O, with a strong oxidizer. These nanoparticles of Nax CoO2 ·yH2 O are highly reactive. They easily dehydrate and show fast ion exchange properties. Sodium can be readily exchanged by protons in neutral solutions as well as by lithium in the presence of lithium salts. However, the sodium form remains stable upon soaking in concentrated NaOH. Samples deposited onto glassy carbon demonstrate the intrinsic properties of the materials. At low sweep rate it appears that the hydrated form Na0.6 CoO2 ·yH2 O can electrochemically exchange sodium and shows a charge storage property as high as 900 ± 250 C/g of oxide. This is a very promising result but there are some difficulties to overcome:
Fig. 5. (a) Total charge stored in the materials during the cyclic voltammetry at different sweep rates (0.5, 1, 2.5, 5, 10, 50 and 100 mV/s); and (b) charge stored as a function of the inverse of the square root of the sweep rate.
and 100 C/g, respectively. Since these values exceed charge storage associated with double-layer processes, it is evident that the cations in these materials are sufficiently mobile to enable faradic processes to contribute to charge storage. Moreover, at 20 mV/s, the storage capacity of Na0.6 CoO2 ·yH2 O is around 330 C/g (440 F/g if divided by the voltage step). This value is very acceptable when compared to other oxides (comparison with other works need to translate charges into capacitances), and could be very interesting in applications. To discriminate between the role of diffusion into the material and charge storage on the surface, we used the model introduced by Trasatti et al. [49]. In characterizing the charge storage kinetics for RuO2 , they assumed that the total charge accumulated, q∗T , is the sum of the charge contributed by the “outer” (more accessible) surface q∗S and the charge associated with the “inner” (less accessible) one q∗I . q∗T = q∗S + q∗I Since the inner charge q∗I is diffusion limited, it is assumed to have a square root dependence (i.e., semi-infinite diffusion) whereas the outer charge should not depend on the sweep rate. Thus, one should observe the following relationship: k q∗T (v) = q∗∞ + √
v
where q∗∞ is the charge that can be stored instantly, k a constant and v the sweep rate. Fig. 5b shows the results of this analysis for the various materials investigated. As expected, the extrapolated values for Na0.6 CoO2 (60 ± 10 C/g) and HCoO2 generally reflect
- First, the storage capacity rapidly decreases with the sweep rate. This is probably associated with the slow ionic diffusion of sodium at the solid- electrolyte interface. Still, at 20 mV/s, the capacity is around 330 C/g (440 F/g if divided by the voltage step). This value is very good when compared to other cobalt oxides in the literature [11,13,36]. For instance, a specific capacitance of amorphous cobalt oxide films in KOH 1 M of 165 F/g of oxide at the lower sweep rate of 10 mV/s was reported by Kandalkar et al. [50], and of more than 290 F/g in the case of CoOx xerogels (sweep rate 5 mV/s) [51]. More recently, Wei et al. have reported a high value of 600 F/g (420 C/g) on cobalt oxide aerogels, in NaOH 1 M at 25 mV/s but it rapidly decreases when the surface area decreases (239 F/g for a specific surface of 123 m2 /g) [52]. However, comparison with these systems found in the literature remains difficult because of the numerous geometries which do not allow to separate the intrinsic properties from the microstructural effects. - Second, the capacity decreases by dehydration of this product, as seen on Fig. 5a. Observation of the C-V (Fig. 4a and b) shows that this decrease is correlated with the extinction of the redox peaks in the dehydrated Na0.6 CoO2 . This is a consequence of the interlayer space decrease that limits the interlayer ionic diffusion. Na0.6 CoO2 should then be kept hydrated in order to present high capacitances. This is an issue for device application, as fabrication of scalable systems using this hydrated product, such as composite membranes, is not solved yet and needs further work. - Third, the potential window used in this study is quite narrow (0.7 V), and is limited by the electrolysis of water at such a high pH. This operating window is comparable to the ones found in the literature for other oxides in aqueous electrolytes [50,52]. For this reason, it could be used in asymmetric supercapacitor associated with carbon, as has already been demonstrated for Co(OH)2 nanoflakes in KOH electrolyte solutions [53]. However, as an electrolyte of 10 M NaOH might not be practical, efforts to work with organic electrolytes containing Na+ ions are currently in progress. 5. Conclusions Mx CoO2 (M = H, Li, Na) nanoparticles have been successfully synthesized with a low temperature route. The most interesting charge storage properties were those obtained with the hydrous
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sodium cobalt composition. At slow sweep rates, this material exhibits high capacity, up to 900 C/g, that arises from both ion insertion and pseudocapacitive processes. Although the storage capacity decreases with increasing sweep rates, to reach 160 C/g at 100 mV/s, the hydrated Na0.6 CoO2 ·yH2 O phase appears to be a very promising material for energy storage applications. Further works are in progress to understand the fundamental charge storage of this lamellar material, to stabilize this form and to prepare thicker composite electrodes containing both polymer binder and carbon for practical device structures. Acknowledgments This study is supported by the DGA. The authors wish to thank Patricia Beaunier (University Pierre et Marie Curie), Frédéric Prima (Chimie-ParisTech) and Frédéric Herbst (University Paris DiderotITODYS) for their help in electron microscopy and Diana Dragoe (CNRS-ICMPE -Thiais) for the ICP–AES experiments. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20]
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Electrochimica Acta journal homepage: www.elsevier.com/locate/electacta
Solution synthesis of nanometric layered cobalt oxides for electrochemical applications Xavier Pétrissans a , Angélique Bétard a , Domitille Giaume a , Philippe Barboux a,∗ , Bruce Dunn b , Lorette Sicard c , Jean-Yves Piquemal c a
Chimie de la Matière Condensée de Paris, CNRS UMR 7574, Chimie-ParisTech, 11 rue Pierre et Marie Curie, 75005 Paris, France Department of Materials Science and Engineering, University of California Los Angeles, Los Angeles, CA 90095, USA c Laboratoire ITODYS, Université Paris-Diderot, Bâtiment Lavoisier, 15 rue Jean de Baïf, 75205 Paris Cedex 13, France b
a r t i c l e
i n f o
Article history: Received 12 December 2011 Received in revised form 24 January 2012 Accepted 28 January 2012 Available online 5 February 2012 Keywords: Nax CoO2 nanoparticles Colloidal oxide Ionic exchange Pseudocapacitor Cyclic voltammetry
a b s t r a c t Dispersed Na0.6 CoO2 ·yH2 O and Li0.5 CoO2 powders have been obtained at room temperature by rapid precipitation in aqueous solutions of LiOH or NaOH in the presence of a strong oxidizer. The precipitates are well crystallized and consist of nanoscale platelets with high specific area (above 100 m2 /g). The Li0.5 CoO2 phase is stable in aqueous electrolytes whereas the Na0.6 CoO2 ·yH2 O rapidly converts to CoOOH in neutral electrolytes or pure water. It also transforms to anhydrous Na0.6 CoO2 upon drying at moderate temperatures. Electrochemical studies show that at slow sweep rates the Na0.6 CoO2 ·yH2 O can store large amounts of charge in 10 M NaOH from a combination of both faradic and capacitive reactions. © 2012 Elsevier Ltd. All rights reserved.
1. Introduction The large scale application of batteries to electrical vehicles requires an improvement in their power density in order to reach an acceptable driving range. Moreover, the energy storage of these systems decreases considerably when the power demand increases [1] while safety problems arise due to heat dissipation and outof-equilibrium electrochemical side reactions [2]. To address the problem of peak power management, batteries can be coupled with complementary devices such as electrochemical capacitors, also known as supercapacitors, that provide high power but for short durations because of low energy density. These devices are based on the development of electrochemical double-layers that store charge at the solid-electrolyte interface and do not involve chemical reactions as in the case of batteries. Electrochemical capacitors possess a number of attractive features compared to batteries including higher power, shorter charging times and much longer cycle life. One method of increasing the energy density of electrochemical capacitors is to have reversible redox processes at the surface of high surface area oxide films or nanoparticles [3,4]. The resulting pseudocapacitance increases the level of energy storage by an order of magnitude or more, with the best properties exhibited by
∗ Corresponding author. Tel.: +33 1 53 73 79 25. E-mail address: [email protected] (P. Barboux). 0013-4686/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2012.01.104
hydrated RuO2 oxides [5,6]. The cost of RuO2 is prohibitive and for that reason, a number of less expensive candidates such as transition metal oxides or hydroxides have been investigated (MnO2 [7–9], Ni(OH)2 [10], Co3 O4 [11,12], CoOOH [13], FeOOH [14], CuO [15], V2 O5 [16–18], TiO2 [19,20], Zn2 CoO4 [21], LiCoO2 [22], etc.). Manganese based oxides have probably attracted the largest interest due to their low cost, and low toxicity. Lamellar MnO2 structures such as birnessite have been the subject of many studies due to their ability to reversibly intercalate monovalent cations in aqueous electrolytes [8]. To achieve significant capacity (on the order of hundreds of F/g or C/g in capacity or charge storage respectively), it is necessary to optimize the synthesis, chemistry and microstructure so that high surface area can be combined with favorable diffusion pathways for both ions and electrons to enable redox reactions to occur at the interface [7,19,23].More generally lamellar structures can achieve good ionic mobility and good exchange properties. Lamellar MnO2 structures such as birnessite have been the subject of many studies due to their high ability to reversibly intercalate monovalent cations in aqueous electrolytes [8]. Layered hydroxides also demonstrated large specific capacitance and good cycle reversibility [24,25]. In this paper, we particularly focus on layered cobalt oxides, especially on the Na0.6 CoO2 phase [26] and on its hydrated polymorphs [27]. The good electronic conductivity of this material associated to its fast ion exchange properties make it a good candidate as a model material to study the effect of ionic mobility on the distribution of pseudocapacity between double
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layer effects and faradic insertion mechanisms as a function of the cycling rate and the nature of the ions in the interlayers. The Ax CoO2 oxides exhibit lamellar structures with sheets of edge-sharing CoO6 octahedra separated by layers of alkali ions. Depending on the alkali ion and stoichiometry of the compound, the stacking of these sheets may differ and lead to octahedrically coordinated alkali ions as in LiCoO2 or to prismatic coordination around the alkali ion as in Na0.6 CoO2 [28,29]. The various Ax CoO2 phases can be obtained by conventional reaction at high temperature between cobalt oxide (Co3 O4 ) and alkali carbonates [27,30,31,34]. LiCoO2 , which has been studied extensively because of its application in lithium-ion batteries [30,32], can be readily synthesized at low temperature by various routes. Cobalt oxyhydroxide, CoOOH, can be reacted with lithium hydroxide solutions [33,34], Co(OH)2 can be refluxed in LiOH without addition of an oxidizer to transform Co(II) into Co(III) [35], or homogeneous lithium cobalt citrate precursors can be decomposed at 400 ◦ C [36]. These approaches are not successful for Nax CoO2 , or lead to multiphase products [32,37]. The formation of the sodium phase is much more difficult than that of LiCoO2 due to the different ionic radii and electronegativity of Li and Na. The electronic distribution is rather different in both phases with stronger hybridization of O and Co orbitals in the case of the sodium phase [38]. As a result, the sodium phase can become metallic whereas the lithium phases are semiconductors [39]. For this reason, it could be proposed as a conductive additive to the Nickel electrode in Ni–MH batteries [40,41].In this work, we present the direct synthesis of nanoparticles of lamellar cobalt oxides and discuss their ability to store energy by capacitive or faradic effects at different rates. The method described here for the synthesis of Mx CoO2 (M = Li, Na) is a low temperature route based on using an alkaline oxidizing medium. The Pourbaix diagram of the cobalt oxides, recently revised by Chivot et al. [42], shows that the MCoO2 form (M = H, Na, etc.) is stable only at high potential and high pH. This method has already been used to synthesize various transition metal oxide compounds including LiCoO2 , Li0.6 NiO2 or Nax MnO2 [43,44]. To the best of our knowledge, the direct precipitation of Na0.6 CoO2 has not been reported in the literature. In addition, many high values of capacitance have been reported on different cobalt oxides in the literature [11,13], however, most of these results have been obtained on thin films or mesoporous structures supported on various substrates, which do not allow to separate the structural and microstructural contributions. It is thus necessary to develop a method to assess the intrinsic electrochemical properties of the materials. The deposition of particles onto glassy carbon electrodes developed in this work allows one to determine the intrinsic properties of the material because the measurement is taken without the presence of either a conductive additive or binder. 2. Experimental 2.1. Synthesis of electrochemically active materials 2.1.1. Reagents All chemical reagents were 99.9% grade, commercially available at Aldrich chemicals and used as received. Glassy carbon electrodes were purchased from Alfa-Aesar. 2.1.2. Material synthesis via alkaline oxidation In a typical synthesis, 22 g of solid NaOH (respectively 6 g LiOH) are dissolved in 50 mL of deionized water in a 250 mL three-neckflask under O2 flow. A 10 M NaOH solution is obtained (respectively 3 M LiOH which is near the saturation limit). The flask temperature is kept around 25 ◦ C, and 0.7 mL of Br2 (14 mmol) are added to the solution which turns yellow. Immediately afterwards, 5 mL of a red
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1 M Co(NO3 )2 solution is added dropwise over 10 min. A black precipitate forms immediately which is recovered after filtration using a sintered glass frit under vacuum. The precipitate is dried at 120 ◦ C for 24 h, producing a black hygroscopic powder. Subsequent washing with pure water to remove the adsorbed NaOH solution could be performed but would affect the nature of the resulting phase as will be discussed in the following section. 2.2. Characterizations 2.2.1. Structure X-ray diffraction (XRD) was carried out on the various synthesized powders using the X’Pert PRO, PANanalytical diffractometer ˚ equipped with an X’Celerator detector and (Co (K␣ Co ) = 1.78901 A) used in the Bragg-Brentano geometry. 2.2.2. Composition analysis Elemental analysis was performed by ICP–AES (Varian Vista Axial and ThermoFisher iCAP 6000). The cobalt mean oxidation state was determined by iodometric titration. 20 mg of powder was dissolved in 20 mL of a 6 M HCl solution containing excess potassium iodide (500 mg). The I2 resulting from the oxidation of iodide by Co3+ and Co4+ was back-titrated with a sodium thiosulphate solution (10−2 M) using starch as the indicator. 2.2.3. Particle size and morphology Particle size and morphology were determined for the synthesized powders using a Transmission Electron Microscope (Jeol-JEM-100CXII operating at 100 kV). In these experiments, the powder was dispersed in ethanol and a drop placed on a 200mesh carbon-coated copper grid. Scanning Electron Microscopy measurements using a Hitachi S2500 IMIX PGT microscope were also performed on powder pasted onto a carbon tape. Surface areas were measured using a gas adsorption analyzer (BEL Japan, Belsorp-max). The Brunauer, Emmett, Teller (BET) analysis was used to characterize the powder samples which were degassed 120 ◦ C under vacuum during 10 h prior to measurement. 2.3. Electrochemical measurements 2.3.1. Glassy carbon electrodes Electrodes were prepared following the method described by Wang et al. [20] by depositing the powder onto a glassy carbon substrate. This method allows direct contact between each oxide grain and the current collector and enables one to determine electrochemical properties of materials without the use of carbon additives or binder. First, a glassy carbon substrate (1 cm × 2 cm) is treated for 15 min in an oxygen plasma to become hydrophilic. A small amount of colloidal dispersion (7.0 L) of the material in the electrolyte to be used (vide infra) is then directly deposited onto the substrate and allowed to dry. To bind the particles to the substrate, it is necessary to anneal the samples at 300 ◦ C in air for one hour except for the hydrated Nax CoO2 ·yH2 O composition which can be annealed only at 220 ◦ C for 40 min. Still, some unbound particles are lost when the electrodes are dipped in the electrolyte. The quantitative loss is determined by chemical analysis of the electrolyte solution after acidification and ICP titration. Finally, the amount of oxide on the electrode is calculated from the concentration of the colloidal dispersion and the volume deposited, corrected with the cobalt concentration found in the electrolyte after the study (about 10–20%). During measurement, Teflon tape covers the other side of the glassy carbon substrate to avoid any leakage current. 2.3.2. Electrochemical studies Cyclic voltammetry (C-V) measurements were recorded using an Autolab PGSTAT 30. A three-electrode geometry was used with a
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large graphite sheet serving as the counter-electrode and a Ag/AgCl electrode as the reference for the glassy carbon experiments. In the C-V experiments, the sweep rate was varied between 0.5 and 100 mV/s and the total amount of charge storage was determined by integrating and averaging both the oxidation and reduction currents. The contribution of the glassy carbon is very low (5 C/g) and has been subtracted directly from the voltammograms. 3. Results 3.1. Synthesis 3.1.1. Crystalline structures The powders synthesized under various conditions were analyzed by XRD and are shown in Fig. 1. Just after synthesis in NaOH, without washing or drying, the material is a wet black paste which appears to be amorphous from XRD (Fig. 1a). Subsequent drying at 120 ◦ C allows the material to crystallize as a single phase corresponding to a Nax CoO2 structure (0.5 < x < 1) (Fig. 1b). All the diffraction peaks can be attributed to the rhombohedral Na0.6 CoO2 structure (JCPDS 00-071-1281) which is composed of CoO6 octahedra linked by edges and separated by an interlayer space in which sodium ions are in a trigonal prismatic environment [28]. The c-axis parameter is calculated to be 16.7 ± 0.3 A˚ based on the XRD results, which is slightly larger than the 16.53 A˚ reported for Na0.6 CoO2 [28]. The broadening of the X-ray peaks indicates that the crystalline domains are quite small. From the Debye-Sherrer formula, we estimate a crystallite size of ∼10 nm. It is interesting to note that some of the peaks are significantly broader than others, especially the (0 0 3) plane around 20◦ . In addition to a size effect, the shape of this (0 0 3) line may be influenced by other contributions, such as a cation deficiency in the interlayer space as discussed by Butel et al. [45]. The presence of cation vacancies can enhance repulsion between the oxide layers, leading to an increase in lattice strains, which is detected by the (0 0 l) line shape. The presence of both sodium ions and protons in the interlayer space can also generate extended structural defects that are caused by segregation of the Nax CoO2 -rich and HCoO2 -rich zones. Abundant washing with water leads to a diffraction pattern (Fig. 1d) which is attributed to a pure -HCoO2 phase (JCPDS 01072-2280). The c-axis cell parameter from the X-ray diagram gives 13.2 ± 0.1 A˚ while it is reported to be 13.13 A˚ for HCoO2 [45]. The
narrowing of the peak corresponding to the (0 0 3) line also indicates better periodicity along the (0 0 l) lines. These results are confirmed by the elemental analysis performed on the samples whose diffraction diagrams are given in Fig. 1b and d. The detailed composition of the powder after synthesis, drying and without washing (Fig. 1b), Na0.6±0.1 CoO2 ·(0.8 ± 0.5)H2 O was determined by linking the Na (12.17 wt%) and Co (44.68 wt%) contents given by ICP with the mean oxidation state of cobalt given by iodometric titration. If we consider that the product obtained after soaking in water (Fig. 1d) should be also hydrated, Nax Hy CoO2 ·zH2 O, the iodometric titration linked with the ICP analysis gives us a composition of Na0.12 H0.83 CoO2 ·(0.4)H2 O. As the (Na + H) content is close to 1 (0.95), we expect nearly all cobalt in this phase to be Co3+ . For simplicity, in this paper we refer to the Na0.6±0.1 CoO2 ·(0.8 ± 0.5)H2 O phase as Na0.6 CoO2 , and to the Na0.12 H0.83 CoO2 ·(0.42)H2 O one as HCoO2 . When the synthesis takes place using LiOH, the resulting material has the same rhombohedral Nax CoO2 structure but with a composition of Li0.5±0.1 CoO2 . The diffraction pattern (Fig. 1c) shows the similarity between phases. As discussed in the introduction, the Lix CoO2 phase is much more stable than Nax CoO2 and it does not transform upon soaking in water. The oxidizer used in the LiOH synthesis is bromine (Br2 ) whose electrochemical potential is independent of pH. In the case of Lix CoO2 , hydrogen peroxide could also be used although its oxidizing potential is weaker at high pH. With this mild oxidizer, we have obtained a Lix CoO2 composition with x ≈ 1. 3.1.2. Particle size and morphology Electron micrographs of the amorphous paste deposited onto the microscope grids without washing show that it is composed of very small crystallites of approximately 30 nm long and 6 nm thick, embedded in an amorphous coating (Fig. 2a). Those crystallites appear to be lamellar, with an interlayer space of 6.6 A˚ as determined with Digital Micrograph software. This is consistent with the hydrated Nax CoO2 phase [45]. The samples dried before dispersion onto the microscope grids give an interlayer spacing of 5.51 A˚ in good agreement with the Na0.6 CoO2 composition [45]. The specific surface area of the unwashed precipitates measured by the BET method is 125 ± 30 m2 /g. This value is equivalent to the surface area calculated for spherical 10 nm non-porous particles of Na0.6 CoO2 (density = 4.55 g/cm). The Li0.5 CoO2 samples exhibit a somewhat greater surface area of 160 ± 20 m2 /g. Thus, the oxidizing alkaline conditions described in this work enable one to obtain high surface area crystalline nanoparticles of Na0.6 CoO2 ·yH2 O which cannot be obtained through high temperature routes. These nanoparticles are highly reactive and may transform upon washing or drying sequences. The large distribution of surface area (±25%) is related to different washing sequences or drying processes (between 80 ◦ C and 120 ◦ C) and can be associated with a variation in the packing of the crystallites at a macroscopic scale, as seen in Fig. 2c and d. This microstructure is favored by the particle morphology in platelets: aggregates of approximately 100 nm form, which themselves pack together into larger objects. 3.2. Stability of Na0.6 CoO2 in electrolyte solutions
Fig. 1. X-ray diffractograms of precipitates obtained with NaOH (a) just after synthesis; (b) after drying at 120 ◦ C during 24 h; (c) with LiOH; and (d) with NaOH after washing with deionised water. Reference diagrams of the -Na0.6 CoO2 phase (JCPDS 00-071-1281), the LiCoO2 phase (JCPDS 04-006-3441) and the HCoO2 phase (JCPDS 01-073-1213) are also shown.
An important issue with the Na0.6 CoO2 is whether immersion in typical electrolytes causes a phase transition to occur. The Na0.6 CoO2 powder was soaked in neutral water as well as in standard electrolyte solutions such as sulphates and hydroxides. In the discussion below, the X-ray patterns shown in Fig. 3 are compared to that of Na0.6 CoO2 (Fig. 3a). Upon soaking in water, the HCoO2 phase is obtained, similar to the effect of washing as indicated in Fig. 1d. With Na2 SO4 , as in pure water, there is a
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Fig. 2. Transmission electron microscopy images of (a) Na0.6 CoO2 ·yH2 O particles just after synthesis; (b) and (c) of Na0.6 CoO2 after drying at 120 ◦ C during 24 h; (d) scanning electron microscopy images of Na0.6 CoO2 powders after synthesis.
complete transformation to Hx CoO2 which indicates that the ion exchange of Na+ for H+ is strongly selective for the proton (Fig. 3c). These results have a number of similarities to those reported by Butel et al. who observed a structure collapse of Na0.6 CoO2 at low pH after Na+ /H3 O+ exchange; once all the alkaline cations were removed from the interlayer space, the -HCoO2 phase was obtained [45]. Soaking in 1 M NaOH solution also leads to the formation of the Hx CoO2 phase (not shown), however, in concentrated NaOH (10 M), the Na0.6 CoO2 phase is unchanged (Fig. 3d). For this reason, all the electrochemical experiments were carried out in 10 M NaOH. With another electrolyte, Li2 SO4 (neutral pH), the material also undergoes ion exchange, probably to a composition of Lix Hy CoO2 (Fig. 3b). The LiCoO2 phase is far more stable than the Nax CoO2 [34] and remains unchanged in water.
3.3. Electrochemical properties
Fig. 3. X-ray diffractograms of the Na0.6 CoO2 phase (a) after synthesis; and after soaking in (b) Li2 SO4 (0.5 M); (c) Na2 SO4 (0.5 M) and (d) NaOH (10 M). Reference diagrams of the Na0.6 CoO2 phase (JCPDS 00-071-1281), the LiCoO2 phase (JCPDS 04-006-3441) and the HCoO2 phase (JCPDS 01-073-1213) are present for comparison. * refers to Na2 CO3 impurities that appear after soaking in the NaOH 10 M solution.
Studies on glassy carbon demonstrate the intrinsic properties of the material, as pure oxide is deposited onto glassy carbon, without any additive. Cyclic voltammetry (CV) was used to characterize the electrochemical properties of the different samples. Fig. 4 shows cyclic voltammograms taken on the third cycle for samples deposited onto glassy carbon at sweep rates ranging from 1 mV/s to 100 mV/s. The samples include the as-synthesized Na0.6 CoO2 ·yH2 O powder without washing and drying (Fig. 4a), the same material
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Fig. 4. Cyclic voltammograms at different sweep rate (1, 5, 10, 50 and 100 mV/s) for (a) Na0.6 CoO2 ·yH2 O in NaOH 10 M; (b) Na0.6 CoO2 performed in NaOH 10 M; (c) Li0.5 CoO2 performed in LiOH 3 M; and (d) HCoO2 performed in NaOH 10 M. The specific current is given per gram of oxide.
after heating at 120 ◦ C to yield Na0.6 CoO2 (Fig. 4b), and after ion exchange in water to yield HCoO2 (Fig. 4d). For all these materials, the electrochemical studies were performed in 10 M NaOH as electrolyte. In addition, we characterized a LiCoO2 powder (Fig. 4c); the electrolyte used for that material was 3 M LiOH. To provide an appropriate comparison among the different samples, the currents are normalized to the mass of the oxide. In establishing the voltage window for the CV experiments, the positive limit was taken as the onset of water oxidation whereas the lower limit was taken at an appropriate potential that gave good cycle-to-cycle reproducibility. In characterizing these samples over tens of cycles, we noticed a slow decrease in current which is mainly associated with detachment of particles from the glassy carbon substrate. This problem was particularly evident with Na0.6 CoO2 ·yH2 O. Thus, the results for different sweep rates had to be collected with different samples of this material and this explains the error bars given from Fig. 5a for the Na0.6 CoO2 ·yH2 O composition. For all other materials the curves could be obtained with a single sample. The curves observed for Nax CoO2 (Fig. 4b) and HCoO2 (Fig. 4d) samples do not show oxidation or reduction peaks, features which are characteristic of an electrochemical reaction. This current–voltage response is the signature of a pure capacitance. In contrast, the curves obtained for the Na0.6 CoO2 ·yH2 O sample (Fig. 4a) or the Li0.5 CoO2 phase (Fig. 4c), exhibit broad oxidation and reduction peaks in addition to the large capacitive effect. At higher sweep rates, the oxidation and reduction peaks are less evident. Overall, the cyclic voltammetry measurements can be considered as being a large capacitive response in combination with faradic reactions.
The sweep rate dependence on charge storage (reported in C/g) for the different samples is shown in Fig. 5. In all cases, the amount of charge storage increases at slow sweep rates because of contributions from diffusion-related processes such as intercalation [20]. Interestingly, for Na0.6 CoO2 ·yH2 O, the charge stored at slow sweep rates can be as high as 900 ± 250 C/g of oxide, similar to that obtained for ruthenium oxide [5,6]. This value is comparable to the charge calculated, ∼920 C/g, for one-electron redox reactions for Co3+ /Co4+ with this material. About half of the measured charge storage can be attributed to the extraction of 0.5 sodium in the Na0.6 CoO2 phase (∼460 C/g). Another contribution, the electrical double layer capacitance at the electrode/electrolyte interface, can be ascribed to the density of surface ionisable groups that can reach a limit value of 50 C/cm2 . Based on the specific surface area of 125 m2 /g, this process yields an estimated capacitive charge of 65 C/g. The remaining amount of charge storage, in the range of 400 C/g, is attributed to a pseudocapacitance contribution. Electrochemical energy storage studies with transition metal oxides have shown that pseudocapacitive charge storage becomes more prominent with nanoscale particles [20,46]. Moreover, in the recent V2 O5 work, pseudocapacitive processes were identified as the charge storage mechanism which enabled the specific capacity to exceed that associated with traditional intercalation [46]. The other materials that were investigated, Na0.6 CoO2 , Li0.5 CoO2 and HCoO2 exhibit charge storage values in the range of 200–300 C/g at slow sweep rate which are consistent with previous work [47,48]. At higher sweep rates (Fig. 5a), the stored charge decreases to values below 200 C/g. At the highest sweep rates of 100 mV/s, the Na0.6 CoO2 ·yH2 O and Li0.5 CoO2 phases still store up to 160
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the double-layer capacitance. Only the Na0.6 CoO2 ·yH2 O material exhibits a significantly greater level of charge storage (200 ± 70 C/g) which, as discussed above, indicates that the high level of Na+ mobility in these materials enables faradic processes to occur. 4. Discussion We have prepared layered alkaline cobalt oxides by a synthetic route based on room-temperature precipitation in an oxidizing medium. The method yields small crystallites with high specific surface area. The LiCoO2 nanoparticles are readily obtained and they are stable in neutral aqueous solutions whereas the sodium phase can only be obtained in a hydrated form, Na0.6 CoO2 ·yH2 O, with a strong oxidizer. These nanoparticles of Nax CoO2 ·yH2 O are highly reactive. They easily dehydrate and show fast ion exchange properties. Sodium can be readily exchanged by protons in neutral solutions as well as by lithium in the presence of lithium salts. However, the sodium form remains stable upon soaking in concentrated NaOH. Samples deposited onto glassy carbon demonstrate the intrinsic properties of the materials. At low sweep rate it appears that the hydrated form Na0.6 CoO2 ·yH2 O can electrochemically exchange sodium and shows a charge storage property as high as 900 ± 250 C/g of oxide. This is a very promising result but there are some difficulties to overcome:
Fig. 5. (a) Total charge stored in the materials during the cyclic voltammetry at different sweep rates (0.5, 1, 2.5, 5, 10, 50 and 100 mV/s); and (b) charge stored as a function of the inverse of the square root of the sweep rate.
and 100 C/g, respectively. Since these values exceed charge storage associated with double-layer processes, it is evident that the cations in these materials are sufficiently mobile to enable faradic processes to contribute to charge storage. Moreover, at 20 mV/s, the storage capacity of Na0.6 CoO2 ·yH2 O is around 330 C/g (440 F/g if divided by the voltage step). This value is very acceptable when compared to other oxides (comparison with other works need to translate charges into capacitances), and could be very interesting in applications. To discriminate between the role of diffusion into the material and charge storage on the surface, we used the model introduced by Trasatti et al. [49]. In characterizing the charge storage kinetics for RuO2 , they assumed that the total charge accumulated, q∗T , is the sum of the charge contributed by the “outer” (more accessible) surface q∗S and the charge associated with the “inner” (less accessible) one q∗I . q∗T = q∗S + q∗I Since the inner charge q∗I is diffusion limited, it is assumed to have a square root dependence (i.e., semi-infinite diffusion) whereas the outer charge should not depend on the sweep rate. Thus, one should observe the following relationship: k q∗T (v) = q∗∞ + √
v
where q∗∞ is the charge that can be stored instantly, k a constant and v the sweep rate. Fig. 5b shows the results of this analysis for the various materials investigated. As expected, the extrapolated values for Na0.6 CoO2 (60 ± 10 C/g) and HCoO2 generally reflect
- First, the storage capacity rapidly decreases with the sweep rate. This is probably associated with the slow ionic diffusion of sodium at the solid- electrolyte interface. Still, at 20 mV/s, the capacity is around 330 C/g (440 F/g if divided by the voltage step). This value is very good when compared to other cobalt oxides in the literature [11,13,36]. For instance, a specific capacitance of amorphous cobalt oxide films in KOH 1 M of 165 F/g of oxide at the lower sweep rate of 10 mV/s was reported by Kandalkar et al. [50], and of more than 290 F/g in the case of CoOx xerogels (sweep rate 5 mV/s) [51]. More recently, Wei et al. have reported a high value of 600 F/g (420 C/g) on cobalt oxide aerogels, in NaOH 1 M at 25 mV/s but it rapidly decreases when the surface area decreases (239 F/g for a specific surface of 123 m2 /g) [52]. However, comparison with these systems found in the literature remains difficult because of the numerous geometries which do not allow to separate the intrinsic properties from the microstructural effects. - Second, the capacity decreases by dehydration of this product, as seen on Fig. 5a. Observation of the C-V (Fig. 4a and b) shows that this decrease is correlated with the extinction of the redox peaks in the dehydrated Na0.6 CoO2 . This is a consequence of the interlayer space decrease that limits the interlayer ionic diffusion. Na0.6 CoO2 should then be kept hydrated in order to present high capacitances. This is an issue for device application, as fabrication of scalable systems using this hydrated product, such as composite membranes, is not solved yet and needs further work. - Third, the potential window used in this study is quite narrow (0.7 V), and is limited by the electrolysis of water at such a high pH. This operating window is comparable to the ones found in the literature for other oxides in aqueous electrolytes [50,52]. For this reason, it could be used in asymmetric supercapacitor associated with carbon, as has already been demonstrated for Co(OH)2 nanoflakes in KOH electrolyte solutions [53]. However, as an electrolyte of 10 M NaOH might not be practical, efforts to work with organic electrolytes containing Na+ ions are currently in progress. 5. Conclusions Mx CoO2 (M = H, Li, Na) nanoparticles have been successfully synthesized with a low temperature route. The most interesting charge storage properties were those obtained with the hydrous
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sodium cobalt composition. At slow sweep rates, this material exhibits high capacity, up to 900 C/g, that arises from both ion insertion and pseudocapacitive processes. Although the storage capacity decreases with increasing sweep rates, to reach 160 C/g at 100 mV/s, the hydrated Na0.6 CoO2 ·yH2 O phase appears to be a very promising material for energy storage applications. Further works are in progress to understand the fundamental charge storage of this lamellar material, to stabilize this form and to prepare thicker composite electrodes containing both polymer binder and carbon for practical device structures. Acknowledgments This study is supported by the DGA. The authors wish to thank Patricia Beaunier (University Pierre et Marie Curie), Frédéric Prima (Chimie-ParisTech) and Frédéric Herbst (University Paris DiderotITODYS) for their help in electron microscopy and Diana Dragoe (CNRS-ICMPE -Thiais) for the ICP–AES experiments. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20]
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