A sodium layered manganese oxides as 3 V cathode materials for secondary lithium batteries

A sodium layered manganese oxides as 3 V cathode materials for secondary lithium batteries

Electrochimica Acta 52 (2006) 504–510 A sodium layered manganese oxides as 3 V cathode materials for secondary lithium batteries S. Bach a,∗ , J.P. P...

404KB Sizes 0 Downloads 21 Views

Electrochimica Acta 52 (2006) 504–510

A sodium layered manganese oxides as 3 V cathode materials for secondary lithium batteries S. Bach a,∗ , J.P. Pereira-Ramos a,b , P. Willmann b a

Laboratoire d’Electrochimie, Catalyse et Synth`ese Organique, CNRS, UMR 7582, 2, rue Henri-Dunant, 94320 Thiais, France b CNES, 18, Avenue Edouard Belin, 31401-Toulouse, Cedex 9, France Received 28 February 2006; received in revised form 16 May 2006; accepted 16 May 2006 Available online 5 July 2006

Abstract The synthesis of a new anhydrous sodium manganese oxide ␣-Na0.66 MnO2.13 obtained via a sol–gel process in organic medium is reported. The ˚ c = 11.09 A) ˚ allows to get a high and partial and limited removal of sodium ions from the layered host lattice (hexagonal symmetry; a = 2.84 A, stable specific capacity of 180 mAh g−1 at C/20 in the cycling limits 4.3/2 V with a mean working voltage of 3 V without the emergence of a spinel phase. By introducing acetylene black in solution during the sol–gel reaction, a composite material containing 8 wt.% AB has been obtained. The rate capability is shown to be significantly improved leading to an increase of the available specific capacity with for instance 200 and 90 mAh g−1 at C/20 and C rate. This effect is ascribed to a better electronic contact between particles and/or the modification of the oxide surface which makes the intercalation process more homogeneous and more efficient. © 2006 Elsevier Ltd. All rights reserved. Keywords: Manganese oxides; Sol–gel process; Lithium batteries

1. Introduction Manganese oxides are preferred as one of the promising electrodes in high energy density lithium batteries because of low cost, low toxicity and safer performance [1–3]. However, some problems are still standing in the utilisation of lithium manganese oxides as cathode in the practical Li-ion batteries. In the case of the LiMn2 O4 spinel, capacity fading upon cycling is always present due particularly to Jahn-Teller lattice distortion [4], slow dissolution of LiMn2 O4 into the electrolyte at high voltages [5] and above 50 ◦ C [6]. Another kind of lithium manganese oxide, LiMnO2 with a layered structure presents a significant drawback due to its crystallographic transformation to a spinel structure during extended cycling [7,8]. To overcome all these problems, many works have been performed on the synthesis and characterization of other kinds of electrochemically active manganese oxides based on the Na–Mn–O system [9–18]. Depending on the sodium content, tunnel structures (Na/Mn < 0.45) [10,15] or layered phases

Corresponding author. Tel.: +33 1 49 78 11 63. E-mail address: [email protected] (S. Bach).

0013-4686/$ – see front matter © 2006 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2006.05.046

(0.45 ≤ Na/Mn ≤ 0.7) are obtained [9,11,13–17]. Little is known on the electrochemical behaviour of lamellar Na-containing phases as cathodic materials due to the fact that many authors only used this group of lamellar phases as precursor for the synthesis of lithiated layered manganese oxides via an ionexchange procedure [16,18] while data found in [14,17] report low capacities in the range 100–150 mAh g−1 with a poor cycling behaviour. Previous results obtained in our group have shown the benefit of using the sol–gel method to get the hydrated lamellar sodium oxide Na0.45 MnO2.14 ·0.76H2 O by a sol–gel process involving the reduction of sodium permanganate by selected reductor organic agents such as for instance the fumaric acid [9,19,20]. The use of Na ions as pillaring species between MnO2 layers has been proven to limit the magnitude of the structural response of the host lattice during the Li insertion extraction reaction. This has resulted in a stabilization of the specific capacity around 130 mAh g−1 on extended cycling at C/6 rate in the potential range 4.2/2.0 V [9]. However, the rate capability was found to be rather low due to the presence of interlayer water. In this paper, we report an extension of the sol–gel method [19,20] applied here in an organic medium to prepare an anhydrous Na-lamellar manganese oxide with the ␣-Na0.66 MnO2.13

S. Bach et al. / Electrochimica Acta 52 (2006) 504–510

composition. The electrochemical and structural properties of the material are discussed. In a second step, the synthesis method has been modified by adding acetylene black in solution during the sol–gel reaction in order to improve the electronic contact between particles and then the electrochemical behaviour of the cathode material ␣-Na0.66 MnO2.13 . 2. Experimental


apparatus. Galvanostatic discharge–charge tests between 2.0 and 4.3 V or 2.0 and 4.8 V versus Li/Li+ were carried out at C/2, C/5, C/10 and C/40 rate. Impedance spectroscopy experiements were carried out in the frequency range 104 –10−4 Hz using an EGG (PAR) Model 273A potentiostat coupled with a 1255 Schlumberger Frequency Response Analyzer. The excitation signal was 10 mV peak to peak. The equilibrium potential was considered to be reached when the drift in open-circuit voltage remained less than 1 mV during 5 h.

2.1. Chemical and structural analysis Manganese oxide is synthesized from the reduction of sodium permanganate by methanol acting both as a reducing organic agent and as the solvent. NaMnO4 .H2 O (Fluka® ) is dissolved in 60 mL of methanol. For the carbon composite material, ABMO, acetylene black (AB) is introduced in the desired ratio AB/manganese oxide = 9.25 wt.% in the methanol solution containing the sodium permanganate. In both cases, a gel is rapidly formed (20–30 min). Complete reduction of Mn(VII) is obtained when the purple coloration cannot be seen anymore in supernatant solution. The gel was dried in air at about 100 ◦ C for 12 h giving rise to a black amorphous powder called xerogel. Calcination is then performed up to 500 ◦ C. At this temperature, the elementary analysis for the un-doped and AB-doped compounds reveals the presence of 0.66 sodium ions per manganese with an experimental oxidation state of Mn, ZMn = 3.60. A convenient formula for the sodium manganese oxide is Na0.66 MnO2.13 , i.e., Na0.66 MnIV 0.6 MnIII 0.4 O2.13 . The average oxidation state, ZMn , of manganese in the sample was determined by the following procedure. The sample (100 mg) was dissolved in 50 cm3 of concentrated H2 SO4 , 50 cm3 of H2 O and in presence of an excess of ferrous(II) ammonium sulphate until complete dissolution. After cooling to 20 ◦ C, the excess of ferrous(II) ammonium sulphate is potentiometrically titrated with a potassium permanganate solution. At the same time a blank is run under identical conditions. XRD experiment was performed with a Brucker D8 diffractometer using Cu ˚ Step scan recordings were carried K␣ radiation (λ = 1.54178 A). out by using 0.02◦ 2θ steps of 10 s duration. 2.2. Electrochemical measurements The electrolyte used was 1 mol L−1 LiPF6 in a binary solution of ethylene carbonate (EC) and diethyl carbonate (DEC) solution (1:2, v/v). The working electrode consisted of a stainless steel grid (7 mm diam., 0.2 mm thickness) with a geometric area of 2 cm2 on which the cathode material was pressed (5 t/cm2 ). In the case of the pure material, the cathode was made of a mixture of active material (80 wt.%), graphite (7.5 wt.%), acetylene black (7.5 wt.%) and teflon as binder agent (5%). In the case of the composite material, the cathode is a mixture of active material (87.5 wt.%), graphite (7.5 wt.%) and teflon (5 wt.%), it means the acetylene black content is then very close to that used in a conventional cathode (8 wt.%). Under these conditions, the comparison of electrochemical properties between both materials is relevant. Electrochemical measurements were carried out in two-electrodes cells (Swagelok® type) using a Mac Pile

3. Results and discussion For the un-doped compound, X-ray diffraction patterns performed on samples heat-treated at different temperatures are shown in Fig. 1. The starting material (Fig. 1) is amorphous up to 200 ◦ C. From 350 ◦ C, ill defined and low intensity peaks are observed at 2-θ values = 16.08◦ , 32.27◦ , 36.59◦ , 44.02◦ . At 500 ◦ C (Fig. 1d and e), all these peaks are well defined with an important increase in the intensity of the 0 0 2, 0 0 4 and 1 0 3 lines, a splitting of the broad peak at 36.59◦ and the emergence of the 1 0 2 and 1 1 4 lines. These diffraction patterns can be clearly indexed on the basis of an hexagonal structure (P63 /mmc) ˚ c = 11.09 A ˚ close to with the following parameters: a = 2.84 A, that reported by Parant et al. [15] for the ␣-Na0.7 MnO2+δ phase synthesized from solid state reactants at 600 ◦ C and in good agreement with the measured oxidation state of manganese (ZMn = 3.60). The structure of the ␣-Na0.7 MnO2+δ is characterized by two layers of edge sharing (MnO6 ) octahedra in the unit cell with a single layer of sodium ions localized between these sheets in trigonal prismatic sites (P) which correspond to a P2 structure, the number by convention giving the number of sheets contained in the unit cell (Fig. 2). Moreover, the XRD pattern obtained at 500 ◦ C (Fig. 1), shows that this sol–gel compound exhibits an important 0 0 l preferred orientation. The SEM micrographs indicates that particles of 1–3 ␮m wide (Fig. 3a) are obtained after heat-treatment at 500 ◦ C. These particles consist of thin platelets ≈0.3/0.5 ␮m long and 0.1 ␮m wide (Fig. 3b). The influence of the current density on the discharge–charge curves of ␣-Na0.66 MnO2.13 in 1 mol L−1 LiPF6 in EC/DEC between 4.3 and 2.0 V is reported Fig. 4. One main process for Li insertion and extraction is evidenced. The maximum Li uptake reached at C/20 corresponds to 0.55 Li per mole of ␣-

Fig. 1. Evolution of the X-ray diffraction patterns (Cu K␣) of the sodium manganese oxide as a function of temperature of the heat-treatment.


S. Bach et al. / Electrochimica Acta 52 (2006) 504–510

Fig. 4. Influence of the current density on ␣-Na0.66 MnO2.13 electrode in a potential window 4.3–2 V in 1 mol L−1 LiPF6 in EC / DEC: (a) C/40; (b) C/20; (c) C/10; (d) C/5; (e) C/2.

Fig. 2. Crystal structure of P2 ␣-Na0.66 MnO2.13 . Mn is located in octahedral sites and Na in prismatic sites.

Fig. 3. Scanning electron micrographs of ␣-Na0.66 MnO2.13 (a) and magnified (b).

Na0.66 MnO2.13 which is in good accord with the mean oxidation state of Mn (ZMn = 3.60). Thus a maximal specific capacity of 150 mAh g−1 is available (Fig. 4a) in the first discharge for ␣Na0.66 MnO2.13 . Based on available Mn4+ , an utilization of 92% is then found which is higher than the value of 20% reported for the ␣-Na0.7 MnO2.25 [17]. In the case of the latter compound, the chemical formula of the compound is not really known since no determination of the manganese oxidation state have been establish by the authors. The average potential of the ␣Na0.66 MnO2.13 /Li cell around 2.97 V, makes it attractive as a cathode in 3 V lithium batteries. A notable decrease of the discharge capacity is only observed for the highest C rate (i.e., C/2 rate). Whatever the C rate, a slight capacity loss is observed during the charge process, indicating that ca. 0.025–0.05 Li ions are systematically trapped inside the host lattice. Of interest is the low polarization observed (<200 mV) between the reduction and oxidation processes. Fig. 5 shows the initial galvanostatic discharge to 2.0 V followed by a first oxidation performed up to 4.8 V. A larger Faradaic yield of 0.82 F mol−1 is then recovered in the charge process due to an additional step at ≈4.55 V/4.6 V corresponding to the extraction of ≈0.40 Na+ ion from the interlayer space. From the second discharge curve, the electrochemical behaviour of the material is found to be significantly changed with a new discharge–charge profile and an improvement of the specific capacity (190 mAh g−1 ) corresponding to a maximum Li uptake of x = 0.70. Fig. 6 shows the XRD patterns for lithiated electrodes ␣Lix Na0.66 MnO2.13 with x = 0, 0.25 and 0.55 during the first discharge. Three of the four main lines of the pristine oxide

Fig. 5. First two discharge–charge profile of ␣-Na0.66 MnO2.13 in a potential window 4.8–2 V in 1 mol L−1 LiPF6 in EC/DEC.

S. Bach et al. / Electrochimica Acta 52 (2006) 504–510


Fig. 8. Evolution of the specific capacity as a function of the number of cycles for ␣-Na0.66 MnO2.13 at () C/20 (4.3–2 V), () C/5 (4.3–2 V), (䊉) C/2 (4.3–2 V) and () C/20 (4.8–2 V) discharge–charge rates.

(0 0 2, 1 0 1, 1 1 0) still exist with a loss of intensity when Li enters the host lattice (x = 0.25) but only the 0 0 2 line remains with a noticeable intensity for higher Li content (Fig. 6). It can be seen that this evolution is accompanied by a shift of the 0 0 2 lines towards larger 2 theta values. This evolution indi˚ in cates the c parameter only decreases by 2% from 11.09 A ˚ ␣-Na0.66 MnO2.13 to reach 10.79 A in ␣-Li0.55 Na0.31 MnO2.13 as lithium insertion proceeds (Fig. 7). This decrease of the c parameter is two up three to fold lower than that observed in the case of Li insertion in hydrated lamellar manganese oxides such as Na0.45 MnO2.14 ·0.76H2 O [9] or sol–gel birnessite [20]. After a discharge–charge cycle up to 4.3 V has been carried out (Fig. 6d), the initial values of the a and c parameter are restored. When a cut-off voltage of 4.8 V is used leading to the further extraction of sodium ions, the (0 0 2) peak shifts to higher 2-θ values, ˚ in ␣showing an expansion of the interlayer spacing of 11.2 A Na0.31 MnO2.13 (Figs. 6e, f and 7). Moreover, the removal of Na ions induces a disordering process with the decrease of 0 0 2, 1 0 1 lines and the disappearance of 1 1 0 line. These results suggest the layered structure is maintained all along the first reduction–oxidation process in the 4.3/2.0 V and in the 4.8/2.0 V range. The evolution of the hexagonal c parameter as a function

of the content of alkali cations (Na+ and Li+ ) is summarized in Fig. 7. The linear increase observed for the c parameter ˚ for ␣-Li0.55 Na0.66 MnO2.13 to reach 11.22 A ˚ for from 10.79 A ␣-Na0.31 MnO2.13 can be explained by an enhancement of electrostratic repulsive force between the negatively charged MnO2 sheets as the amount of interlayer cations decreases. Available data on the evolution of the c parameter as a function of the alkali content in layered oxides are limited to a very narrow Li composition range, (x = 0.2 Li) for the ␣-Na0.7 MnO2+δ compound [17,21–23]. The cycling performance of the sol–gel compound ␣Na0.66 MnO2.13 has been evaluated in the two potential limits 4.3/2 and 4.8/2 V. The discharge capacity obtained as a function of the number of cycles for different discharge–charge rates (C/2, C/5 and C/20) is reported in Fig. 8. During the first ten cycles the specific capacity slightly decreases to reach 155 mAh g−1 (96% of the first discharge) at C/20. A slight increase of the specific capacity is then observed due to a progressive extraction of sodium ions which occurs from the 10th cycle at C/20. The same phenomenom takes place for the 20th cycle at C/5. For example, Fig. 9 shows a comparison between the first, the 10th and the 50th cycle for ␣-Na0.66 MnO2.13 at C/20 discharge–charge rate. It is interesting to note that there is a difference in the shape of the discharge–charge profile with an evolution of the discharge curve towards a sloped voltage profile well known to characterize layered phases. In addition, an increase of the working potential can be reported from 2.970 V from the first cycle to 3.115 V in the 10th discharge to reach 3.135 V at the 50th cycle. From the 30th cycle, stable specific capacities are observed at

Fig. 7. Evolution of the hexagonal c parameter in the sol–gel ␣-Lix Nay MnO2.13 as a function of alkali content (Na+ and Li+ ).

Fig. 9. First (—), 10th (– – –) and 50th (- - -) discharge–charge cycles for ␣Na0.66 MnO2.13 (cycling limits: 4.3–2 V in 1 mol L−1 LiPF6 in EC/DEC).

Fig. 6. XRD patterns of (a) ␣-Na0.66 MnO2.13 , (b) ␣-Li0.25 Na0.66 MnO2.13 , (c) ␣-Li0.55 Na0.66 MnO2.13 , (d) after one oxidation process at 4.3 V: ␣Li0.04 Na0.66 MnO2.13 , (e) after one oxidation process at 4.8 V: ␣-Na0.31 MnO2.13 and (f) zoom of the XRD patterns between 10 < 2␣ < 25 (* graphite peak).


S. Bach et al. / Electrochimica Acta 52 (2006) 504–510

Fig. 10. The 60th discharge–charge curves (C/20 rates) for ␣-Na0.66 MnO2.13 electrode in a potential window 4.8–2 V in 1 mol L−1 LiPF6 in EC/DEC at C/20 (4.8–2 V).

Fig. 12. Comparison of XRD patterns of (a) ␣-Na0.66 MnO2.13 and (b) AB/␣Na0.66 MnO2.13 .

C/20 and C/2. After 60 cycles, attractive specific capacities of 180, 150 and 55 mAh g−1 are obtained at C/20, C/5 and C/2 rate, respectively. Therefore this compound can be considered as among the best manganese oxides for rechargeable lithium batteries. The best results reported by Strobel et al. [14] for anhydrous sodium phyllomanganates do not exceed 100 mAh g−1 after the third cycle in the wide voltage range 3.8–1.2 V. In other respects, attractive properties have been recently reported with the 2D Li0.54 Na0.02 Mn0.86 O2 with a specific capacity of 160 mAh g−1 achieved in the voltage range 4.3–2.5 V at C/15 rate after only 20 cycles [16]. Surprisingly, the discharge capacity of the ␣-Na0.66 MnO2.13 strongly decreases from the 20th cycle when the enlarged voltage range (4.8–2.0 V) is used. The material delivers a discharge capacity of 130 mAh g−1 after the 60th cycle which corresponds to a loss of 40%. To understand this capacity fading, the 60th discharge–charge cycle for ␣-Na0.66 MnO2.13 electrode is reported in Fig. 10. The emergence of two voltage plateaus located at 2.9 and 4.1 V already observed in other Li–Mn–O compounds suggests the layered structure progressively converts into a spinel structure. This structural transition is clearly confirmed by XRD experiments performed after 40 and 60 cycles (Fig. 11). Fig. 11a clearly indicates this compound is a mix˚ ture of two phases. The diffraction lines, (0 0 2) : d = 5.63 A, ˚ (1 0 3): d = 2.04 A ˚ and (1 1 0): d = 1.42 A, ˚ (1 0 1) : d = 2.46 A, can be indexed on the basis of the pristine hexagonal lattice of the ␣-Na0.66 MnO2.13 (P63 /mmc), with the following param˚ c = 11.12 A, ˚ while the other diffraction peak eters: a = 2.85 A, (1 1 1) located at 2 theta values = 18.67◦ , can be ascribed to a ˚ with the main lines 1 1 1, 3 1 1 and spinel phase (a = 8.21 A)

4 4 0. After 60 cycles, we only observe a single well crystallized spinel phase with a lattice parameter close to that observed ˚ The structural change for the LiMn2 O4 cubic phase: a = 8.24 A. observed during intercalation/deintercalation may be due to an important removal of sodium ions at this high cut-off voltage limit. The loss of these pillaring species allows the progressive transformation of the initial layered phase into a spinel structure. We have then evaluated the influence of acetylene black on the structure and electrochemical properties of the cathodic material. Comparison of XRD patterns recorded for the un-doped and the carbon composite material, AB-␣-Na0.66 MnO2.13 , labeled ABMO, heat-treated at 500 ◦ C (Fig. 12) shows there is no change in the position and intensities of diffraction lines. The influence of the discharge rates on the Faradaic yield reported for the pristine material and ABMO (Fig. 13) clearly emphasizes the better electrochemical performance found for the ABMO compound especially at high rate. Comparison of the discharge–charge curves for the first (Fig. 14a) and 60th cycle (Fig. 14b) shows the presence of acetylene black in the composite material significantly increases the Faradaic yield and minimizes the polarization without any modification of the electrochemical fingerprint. Clearly, the AB-composite material provides the best results. Indeed, the Faradaic yield recovered with the ABMO is higher by 10% and by 40% at C/20 and C/2 respectively while the AB/manganese oxide ratio is very close in both cases (8 wt.% in ABMO, 7.5 wt.% in the pristine material). The cycling properties of the composite compound are reported in Fig. 15. After the 40th cycle, the material delivered stable discharge capacities of 180 and 200 mAh g−1 for ␣-Na0.66 MnO2.13 and AB-␣-Na0.66 MnO2.13 at C/20 (Fig. 15).

Fig. 11. XRD patterns of an electrode after 40th cycle and after 60th cycle (* graphite peak).

Fig. 13. Influence of the discharge rates on the Faradaic yield reported for ␣Na0.66 MnO2.13 (䊉) and the AB/␣-Na0.66 MnO2.13 ().

S. Bach et al. / Electrochimica Acta 52 (2006) 504–510


Fig. 16. ac impedance diagrams for ␣-Li0.3 Na0.66 MnO2.13 (䊉) and the AB/␣Li0.3 Na0.66 MnO2.13 ().

Fig. 14. Comparison of the first discharge–charge curves for ␣-Na0.66 MnO2.13 (—) and AB/␣-Na0.66 MnO2.13 (- - -) (a) and the 60th discharge–charge curves for ␣-Na0.66 MnO2.13 (—) and AB/␣-Na0.66 MnO2.13 (- - -) (b).

At C/2 rate, the improvement is more important with a specific capacity of 100 mAh g−1 which is twice that exhibited by the pristine ␣-Na0.66 MnO2.13 electrode. These results compare very well with data from reference [24] which reported a stable capacity of ≈200 mAh g−1 for the lamellar compound LiMn0.9 Co0.1 O2 . Values between 165 and 200 mAh g−1 up to 50 cycles are often reported for other compounds among Li[NiCoMn]O2 series [25–27]. We have performed ac impedance measurements for Li0.3 Na0.66 MnO2.13 and AB-Li0.3 Na0.66 MnO2.13 (Fig. 16). The diagrams present three regions: in the high frequency range (104 –10 Hz), only one semi-circle corresponding to the charge transfer appears, followed at medium frequency (10–10−1 Hz) by a straight line with a phase angle of 45◦ from the real axis corresponding to the Warburg impedance and a capacitive line at low frequency (f < 10−2 Hz). The characteristic frequency, f* , associated with the charge transfer semi-circle changes from 230 to 44 Hz while the charge transfer resistance (Rtc ) decreases

Fig. 15. Evolution of the specific capacity as a function of the number of cycles for ␣-Na0.66 MnO2.13 () at C/2 and C/20, and for AB/␣-Na0.66 MnO2.13 () at C/2 and C/20 [cycling limits 4.3–2 V].

from 62 to 36  for the pristine material and ABMO compounds. This indicates an increase in the exchange current density from 0.42 to 0.72 mA cm−2 for ABMO. The presence of acetylene black particles mixed during the sol–gel process leading to the formation of the solid oxide network is thought to ensure a better electronic contact between particles and/or a modification of the surface of the oxide. The low frequency response is characterized by a straight line with a phase angle of 45◦ from the real axis corresponding to the Warburg impedance from which the numerical values of the apparent chemical diffusion coefficient DLi are calculated using Eq. (1) when √ (1) ω  2DLi /L2 : DLi = VM (dE/dx)x /(F 2AS) where VM is the molar volume of the compound (= 52 cm3 mol−1 ), S is the apparent surface area of the electrode, i.e., 2 cm2 , and (dE/dx)x is the slope, at fixed x, of the equilibrium potential composition curve. L is the maximum length of the diffusion pathway (cm). The analysis of the Warburg impedance of the system plotted in the complex plane −Im Z = Aω−1/2 (3) or Re Z = Aω−1/2 (4) allows to get the Warburg prefactor A and then to calculate DLi . This region corresponds to a frequency range where the kinetics of the system is almost entirely limited by the rate of the chemical diffusional process in the host material under semi-infinite conditions. Fig. 17 illustrates the variation of the imaginary impedance part versus ω−1/2 , i.e., according to equation (1). Same values of the Warburg prefactor A have been found for the pristine material and its ABMO form, allowing to cal-

Fig. 17. Evolution of the imaginary impedance part, Zim as a function of ω−1/2 for ␣-Li0.3 Na0.66 MnO2.13 (䊉) and the AB/␣- Li0.3 Na0.66 MnO2.13 ().


S. Bach et al. / Electrochimica Acta 52 (2006) 504–510

culate a value of ≈5 × 10−10 cm2 s−1 for the apparent chemical diffusion coefficient of lithium ions. Such a value is consistent with the kinetics of Li transport reported in various MnO2 structures [28–30]. The same limiting frequency fl ≈ 5 × 10−3 Hz (≈DLi /L2 ) is found in both cases showing the kinetics for Li transport into ␣-Na0.66 MnO2.13 and AB-␣-Na0.66 MnO2.13 is the same. In order to investigate the influence of the dispersion state of electro-conductive additives on the properties of lithium-ion batteries, other works reported the possibility to add carbon during the synthesis of MnO2 . Hibino et al. [11,13,31,32] used acetylene black or ketjen black, in the sol–gel and the sonochemical reaction to get rapid discharge properties and an improvement of specific capacities. However, they obtained amorphous or disordered hydrated sodium manganese oxide phases with attractive initial capacities but poor cycle life quite different from that reported in this work. In addition, they used a large amount of carbon, near 50% by weight. Therefore, any comparison is difficult. 4. Conclusion A new lamellar sodium manganese oxide ␣-Na0.66 MnO2.13 has been synthesized via a sol–gel process performed in methanol. The structure of sol–gel ␣-Na0.66 MnO2.13 is char˚ acterized by the following hexagonal parameters a = 2.84 A, ˚ c = 11.09 A. In spite of a low mean oxidation stated for manganese (ZMn = 3.60), a high and stable specific capacity of 180 mAh g−1 between 4.3 and 2.0 V can be reached without the formation of the spinel phase. Such a result is explained by the extraction of a limited amount of sodium ions from the interlayer spacing while maintaining the layered structure upon cycling. The small change of the interlayer spacing (≈2%) as Li accommodation proceeds allows a high rechargeability and a good cycle life. A notable improvement of the maximum capacity combined with a decrease of the polarization is obtained by using a composite material containing 8 wt.% of acetylene black. The kinetics of charge transfer is improved by a factor two. Then high stable capacities are achieved with 200 mAh g−1 at C/20 after 60 cycles and 90 mAh g−1 at C/2 rate. Carbon particles introduced during the synthetic reaction in solution probably form a conductive network between MnO2 particles and carbone making the electrochemical process more efficient than in a conventional cathode material resulting from a mechanical mixture. However, the existence of carbon particles acting as microcurrent collectors on which the MnO2 oxide is formed cannot be discarded. Further experiments are needed to localize the carbon particles and we believe further improvement both in terms of capacity and rapid discharge performances can

be still obtained through an optimization of the method we used here. References [1] J.M. Tarascon, W.R.Mc. Kinnon, F. Coowar, T.N. Bowmer, G. Amatucci, D. Guyomardand, J. Electrochem. Soc. 141 (1994) 1421. [2] Y. Xia, M. Yoshio, J. Electrochem. Soc. 143 (1996) 825. [3] T. Horiba, K. Hironaka, T. Matsumura, T. Kai, M. Koseki, Y. Muranaka, J. Power Sources 119–121 (2003) 893. [4] R.J. Gummow, A. DeKock, M.M. Thackeray, Solid State Ionics 69 (1994) 59. [5] A. Yamada, J. Solid State Chem. 122 (1996) 100. [6] G. Amatucci, A. Du Pasquier, A. Blyr, T. Zheng, J.M. Tarascon, Electrochim. Acta 45 (1999) 255. [7] G. Vittins, K. West, J. Electrochem. Soc. 144 (1997) 2587. [8] Y.I. Jang, B. Huang, Y.M. Chiang, D.R. Sadoway, Electrochem. Solid State Lett. 1 (1998) 13. [9] P. Le Goff, N. Baffier, S. Bach, J.P. Pereira-Ramos, R. Messina, Solid State Ionics 61 (1993) 309. [10] F. Hu, M.M. Doeff, J. Power Sources 129 (2004) 296. [11] H. Kawaoka, M. Hibino, H. Zhou, I. Honma, Electrochem. Solid State Lett. 8 (2005) A253. [12] M. Tsuda, H. Arai, Y. Sakurai, J. Power Sources 110 (2002) 52. [13] M. Hibino, H. Kawaoka, H. Zhou, I. Honma, Electrochim. Acta 49 (2004) 5209. [14] P. Strobel, C. Mouget, Mater. Res. Bull. 28 (1993) 93. [15] J.P. Parant, R. Olazcuaga, M. Devalette, C. Fouassier, P. Hagenmuller, J. Solid State Chem. 3 (1971) 1. [16] L. Bordet-Le Guennec, P. Deniard, P. Biensan, C. Siret, R. Brec, J. Mater. Chem. 10 (2000) 2201. [17] A. Mendiboure, C. Delmas, P. Hagenmuller, J. Solid State Chem. 57 (1985) 323. [18] J.M. Paulsen, J.R. Dahn, Solid State Ionics 126 (1999) 3. [19] S. Bach, M. Henry, J. Livage, J. Solid State Chem. 88 (1990) 325. [20] S. Bach, J.P. Pereira-Ramos, N. Baffier, R. Messina, Electrochim. Acta 36 (1991) 1595–1603. [21] C. Delmas, J.J. Braconnier, C. Fouassier, P. Hagenmuller, Solid State Ionics 3/4 (1981) 165. [22] J.J. Braconnier, C. Delmas, P. Hagenmuller, Mater. Res. Bull. 17 (1982) 993. [23] D.-C. Li, T. Muta, L.-Q. Zhang, M. Yoshio, H. Noguchi, J. Power Sources 132 (2004) 150. [24] P.G. Bruce, A.R. Armstrong, R.L. Gitzendanner, J. Mater. Chem. 9 (1999) 193. [25] D.C. Li, T. Muta, L.Q. Zhang, M. Yoshio, H. Noguchi, J. Power Sources 132 (2004) 150. [26] S.H. Na, H.S. Kim, S.I. Moon, Solid State Ionics 176 (2005) 313. [27] Z. Lu, D.D. Mac Neil, J.R. Dahn, Electrochem. Solid State Lett. 4 (2001) A200. [28] P.G. Dickens, G.F. Reynolds, Solid State Ionics 5 (1981) 331. [29] J. Vondrak, I. Jazubec, J. Power Sources 14 (1985) 141. [30] S. Franger, S. Bach, J. Farcy, J.-P. Pereira-Ramos, N. Baffier, J. Power Sources 109 (2002) 262. [31] H. Kawaoka, M. Hibino, H. Zhou, I. Honma, J. Power Sources 119–121 (2003) 924. [32] H. Kawaoka, M. Hibino, H. Zhou, I. Honma, J. Power Sources 125 (2004) 85.