electrode in Na-ion batteries is also advantageous to the cost reduction, while copper foil is necessary in Li-ion batteries. [30,31] Studies on room-temperature Na-ion batteries have started since 1970s and electrochemical properties of Na//Na x CoO 2 and Na//TiS 2 cells were reported in 1980 [32,33] when that of Li//LiCoO 2 was also first reported. [34] Then, Li-ion batteries was commercialized in 1991. In principle, Li-ion batteries consist of a Li-containing material such as LiCoO 2 as a positive electrode material and a Li-insertion material such as carbon as a negative one. On charge process, Li + ions move from the posi tive to negative electrode through electrolyte solution with simultaneous movement of electrons through an external circuit. A discharge process proceeds in the opposite direction. Li + ions and electrons come back to the positive electrode on the discharge. For ease in handling, Licontaining materials are utilized for the positive and non-Li ones for the negative electrode. Li-ion batteries have attracted much attention as high-voltage rechargeable batteries. On the other hand, Na-ion batteries essentially consist of the same technology with Li-ion batteries, except charge carriers. Na-containing materials are used for the positive and non-Na ones for the negative electrode. Na//Na x CoO 2 , however, shows working voltage ≈1 V lower than Li//LiCoO 2 , resulting in lower energy density. [35] As a result, Na-ion batteries have never been commercialized so far. [18] Indeed, Allied Corp. in USA, Showa Denko K. K., and Hitachi, Ltd. in Japan had collaborative work on Na-ion batteries and filed patents of Na-Pb alloy//γ-Na x CoO 2 cells [36][37][38][39] exhibiting good cycle stability but they had never commercialized the Na-ion batteries. The Na-Pb alloy//γ-Na x CoO 2 cells needed presodiation process for the Pb negative electrode. Therefore, primary drawback of Na-ion batteries was no candidates as practical negative electrode materials without Na predoping. Now, nongraphitizable carbon, so called hard carbon, is known to deliver large reversible capacity with good capacity retention. [40,41] Thus, secondary issue is low working potential of positive electrode materials. Open circuit potential of αand γ-Na x CoO 2 in Na cells is lower than that of α-NaFeO 2 type LiCoO 2 in the Li cell at the end of discharge. Indeed, standard redox potential of Na metal is lower than that of Li metal by ≈0.3 V. [42,43] The difference is, however, much smaller than that between Na x CoO 2 and LiCoO 2 at the end of discharge (ΔV ≈ 1.5 V), [14] which is probably due to larger ionic size and lower Lewis acidity of Na + in comparison to Li + as already discussed by Goodenough and Mizushima et al. in 1980. [35] Since our group demonstrated hard carbon//NaNi 1/2 Mn 1/2 O 2 full cells exhibiting acceptable cycle stability in 2009 [44] and Sodium 3d transition metal oxides for Na-ion batteries have attracted attention of battery researchers because of their new chemistries and abundant material resources in the eart...
Solid solution samples of P2-type Na 2/3 Ni 1/3-x Mn 2/3 Cu x O 2 (0 ≤ x ≤ 1/9) are successfully synthesized and their electrochemical performance are examined in non-aqueous Na cells. Non-substituted Na 2/3 Ni 1/3 Mn 2/3 O 2 delivers the highest reversible capacity compared to those of the solid solution samples; however the capacity rapidly decays during cycles. Partial substitution of Cu for Ni in Na 2/3 Ni 1/3 Mn 2/3 O 2 effectively improves its cyclability and rate capability. In particular, a Na//Na 2/3 Ni 1/4 Mn 2/3 Cu 1/12 O 2 cell delivers reversible capacity of 138 mAh g −1 and the average discharge voltage reaches 3.54 V on initial discharge with good cycle stability. Operando XRD reveals that the original P2 phase transforms to P2-O2 intergrowth phase having stacking faults during sodium extraction, as is supported by XRD data and simulation with DIFFaX software. Although non-substituted Na 2/3 Ni 1/3 Mn 2/3 O 2 shows P2-O2 phase transition as a two-phasic reaction, P2-type Cu-substituted material does not transform to O2-type one and P2type layers with wider interlayer distance than that of the O2-type Ni 1/3 Mn 2/3 O 2 are randomly and partly retained in the desodiated phase because Cu is distributed in the initial phase at random and Na in interslab space surrounding the electrochemically inactive Cu is not extracted from the structure. The suppressed volume change during charge/discharge results in the excellent electrode performance compared to non-substituted Na 2/3 Ni 1/3 Mn 2/3 O 2 .
P2-Na2/3Ni1/3Mn2/3O2 (P2-NiMn) is one of the promising positive electrode materials for high-energy Na-ion batteries because of large reversible capacity and high working voltage by charging up to 4.5 V versus Na+/Na. However, the capacity rapidly decays during charge/discharge cycles, which is caused by the large volume shrinkage of ca. 23% by sodium deintercalation and following electric isolation of P2-NiMn particles in the composite electrode. Serious electrolyte decomposition at the higher voltage region than 4.1 V also brings deterioration of the particle surface and capacity decay during cycles. To solve these drawbacks, we apply water-soluble sodium poly-γ-glutamate (PGluNa) as an efficient binder to P2-NiMn instead of conventional poly(vinylidene difluoride) (PVdF) and examined the electrode performances of P2-NiMn composite electrode with PGluNa binder for the first time. The PGluNa electrode shows Coulombic efficiency of 95% at the first cycle and capacity retention of 89% after 50 cycles, whereas the PVdF electrode exhibits only 80 and 71%, respectively. The alternating current impedance measurements reveal that the PGluNa electrode shows a much lower resistance during the cycles compared with the PVdF one. From the surface analysis and peeling test of the electrodes, the PGluNa binder was found to cover the surface of the P2-NiMn particles and suppresses the electrolyte decomposition and surface degradation. The PGluNa binder further enhance the mechanical strength of the electrodes and suppresses the electrical isolation of the P2-NiMn particles during sodium extraction/insertion. The efficient binder with noticeable adhesion strength and surface coverage of active materials and carbon has paved a new way to enhance the electrochemical performances of high-voltage positive electrode materials for Na-ion batteries.
O3 type NaNi1/2Mn1/2O2 materials with Mg and Ti co-substitution demonstrate better capacity capability with an initial discharge capacity of 200 mA h g−1 in non-aqueous Na cells without any capacity loss due to substitution.
Although O3‐NaFe1/2Mn1/2O2 delivers a large capacity of over 150 mAh g−1 in an aprotic Na cell, its moist‐air stability and cycle stability are unsatisfactory for practical use. Slightly Na‐deficient O3‐Na5/6Fe1/2Mn1/2O2 (O3‐Na5/6FeMn) and O3‐Na5/6Fe1/3Mn1/2Me1/6O2 (Me = Mg or Cu, O3‐FeMnMe) are newly synthesized. The Cu and Mg doping provides higher moist‐air stability. O3‐Na5/6FeMn, O3‐FeMnCu, and O3‐FeMnMg deliver first discharge capacities of 193, 176, and 196 mAh g−1, respectively. Despite partial replacement of Fe with redox inactive Mg, oxide ions in O3‐FeMnMg participate in the redox reaction more apparently than O3‐Na5/6FeMn. X‐ray diffraction studies unveil the formation of a P‐O intergrowth phase during charging up to >4.0 V.
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