Article pubs.acs.org/JPCC
Mesoporous Cobalt Oxalate Nanostructures as High-Performance Anode Materials for Lithium-Ion Batteries: Ex Situ Electrochemical Mechanistic Study Wei An Ang,†,‡ Yan Ling Cheah,‡ Chui Ling Wong,‡ Raghavan Prasanth,†,‡ Huey Hoon Hng,† and Srinivasan Madhavi*,†,‡,§ †
School of Materials Science and Engineering, Nanyang Technological University, 50 Nanyang Avenue, Singapore 639798 Energy Research Institute @ NTU (ERI@N), Nanyang Technological University, 50 Nanyang Drive, Singapore 637553 § TUM-CREATE Center for Electromobility, Nanyang Technological University, 62 Nanyang Drive, Singapore 637459 J. Phys. Chem. C 2013.117:16316-16325. Downloaded from pubs.acs.org by UNIV OF KANSAS on 01/22/19. For personal use only.
‡
S Supporting Information *
ABSTRACT: Two distinct mesoporous nanostructures, that is, rod and sheet cobalt oxalate (CoC2O4), have been synthesized via facile chimie douce precipitation technique. The selective interaction between solvent type and crystallographic planes of the metal ion is the key factor in morphological variations. The morphology and microstructure are studied by high-resolution transmission electron microscopy. Structural characterization of the materials has been carried out by X-ray diffraction and confirmed phase pure CoC2O4·2H2O formation. The critical dehydration process of CoC2O4·2H2O led to anhydrous CoC2O4, and its thermal properties are investigated by thermogravimetric analysis. Electrochemical properties of anhydrous CoC2O4 in half-cells are studied by cyclic voltammetry, galvanostatic charge−discharge cycling, and electrochemical impedance spectroscopy. The studies showed that initial discharge capacity of anhydrous CoC2O4 nanorods and sheets is 1599 and 1518 mA h g−1, respectively, at 1C-rate. Anhydrous CoC2O4 nanostructures fabricated by this chimie douce process achieved higher reversible capacity, more stable cycling, and better rate capabilities than reported. The electrochemical performances of anhydrous CoC2O4 nanostructures are found to be significantly influenced by morphology and porosity. In addition, the interfacial electrochemical mechanism related to the transitional metal oxidation states, phase structural changes, and distribution during cycling are validated.
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INTRODUCTION As one of the critical electrochemical energy storage devices, rechargeable lithium-ion batteries (LIBs) have drawn worldwide interest for their varied applications such as portable consumer electronics, implantable medical devices, electric vehicles, and so on.1,2 Much attention has been devoted in recent years to improve the performances of LIBs, which are one of the prospective energy-storage systems for renewable energies and urban transportation.3,4 Today, the key to commercial success of hybrid electric vehicles lies in highperformance battery technology. These vehicles are promising because they combine the extended driving range and rapid refueling capacity of conventional vehicles with increased fuel economy and environmental friendliness. The relatively high specific energy and power characteristics of LIBs make them an attractive candidate to nickel metal-hydride batteries used in hybrid vehicles currently in the market. Most commercial LIBs use carbonaceous materials (i.e., graphite, theoretical capacity of 372 mA h g−1) as anode. However, emerging applications demand high-performance LIBs in terms of capacity, energy density, power, and safety. Hence, intensive research on alternative electrode materials that are efficient and inexpensive © 2013 American Chemical Society
is ongoing. Numerous materials have been widely investigated as anode materials in LIBs.5−7 Transition-metal oxides, such as cobalt oxides (CoO, Co3O4), are attractive due to high lithium storage capacities, which are about three times higher than those of graphite.5,8 However, the large initial irreversible capacity and poor capacity retention during charge−discharge cycling restrict their practical applications in LIBs.9,10 Thus, efforts are underway to look for alternative efficient anode materials for high energy and power density LIBs. Recently, transition-metal oxalates11−13 and mixed transitionmetal oxalates14,15 are being studied as potential anode materials for LIBs. Aragon et al. reported iron oxalate11 and cobalt oxalate12 (CoOx) nanoribbons, both synthesized by reverse micelles method exhibiting reversible capacity of ∼700 and ∼900 mA h g−1, respectively, at 2C-rate. However, considerable capacity fading was observed after 40 cycles. On the contrary, a simple, surfactant-free, and economical chimie douce technique for preparing mesoporous iron oxalates Received: April 24, 2013 Revised: July 4, 2013 Published: August 5, 2013 16316
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Perkin-Elmer Spectrum GX) spectra were obtained for the samples in KBr pellets at a resolution of 1 cm−1. Electrochemical Measurements. The anode was prepared by mixing 60 wt % A-CoOx, 30 wt % conductive additive (Super P Li carbon, Timcal), and 10 wt % PVDF binder (Kynar 2801) in N-methyl-2-pyrrolidone (NMP) as solvent for the binder to form the slurry. The resulting homogeneous slurry was casted onto etched copper foil by doctor-blade technique, followed by drying at 80 °C for 24 h under vacuum. Before being used, the film electrode was punched into circular disks of 16 mm diameter prior for assembly. Test cells, Li/CoC2O4, were fabricated in coin cell configuration (CR2016), and electrochemical tests were conducted using a battery test system (XWJ Neware) between 0 and 3 V at different C rates (1−10 C rates, 1 C = 1 Li h−1 mol−1, that is, ∼0.2 mA cm−2). 1 M LiPF6 in ethylene carbonate (EC)/diethyl carbonate (DEC) (1:1 wt/wt, Danvec) was used as electrolyte and Celgard 2400 membrane was used as the separator. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) measurements were carried out on a Solartron (1470E and SI 1255B impedance/gain-phase analyzer coupled to a potentiostat) electrochemical workstation. CV was studied at 0.1 mV s−1 in the voltage range of 0 to 3 V, and EIS was studied by applying an A.C. perturbation of 10 mV over frequency range of 100 kHz to 5 mHz at open circuit potential. The Nyquist plots were analyzed using Zplot and Zview programs (Version 2.2, Scribner Associates). All of the test cells were fabricated in Ar-filled glovebox (H2O, O2 < 1 ppm, Mbraun, Unilab) using Li metal foil (∼0.59 mm thick, Hohsen Corporation) as counter and reference electrode. The cells were then aged for 24 h at room temperature prior to measurements. For postcycled analysis, the cycled electrodes were carefully removed from the coin cell inside an Ar-filled glovebox and washed with DEC solvent for 30 min. After washing, the electrodes were left to dry in the glovebox at room temperature (until all of the DEC solvent evaporated) prior to further analysis.
(cocoons and rods) showed improved electrochemical properties, and excellent cycling performance (even at high C-rates) is reported.13 To the best of our knowledge, there are only limited reports on the effect of morphologies on electrochemical performance and interfacial reactions mechanism of metal oxalate anodes in LIBs. In this work, the study of different nanostructures, rods and sheets of CoC2O4 synthesized by surfactant-free chimie douce precipitation technique, is carried out as a potential anode material in LIBs. The present study also reports a comprehensive mechanistic investigation of CoOx nanostructures, in which the interfacial mechanism related to the transitional metal oxidation states, phase structural changes, and distribution during cycling are discussed.
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EXPERIMENTAL SECTION Synthesis of CoOx Nanostructures. All of the chemicals (analytical grade) were used as received without further treatment. CoOx nanostructures were prepared as reported with the iron precursor being replaced by CoCl2·6H2O.13 The pink colored precipitate obtained is separated by centrifugation, washed with excess ethanol, and dried under vacuum at 60 °C for 12 h (hydrated CoOx, i.e., H-CoOx). To get anhydrous CoOx, that is, A-CoOx, H-CoOx nanostructures were carefully dried under vacuum at 250 °C for 2 h and subsequently kept in drybox (∼20% relative humidity). Materials Characterization. Thermal behavior of samples was studied by thermogravimetric analysis (TGA, TA Instruments Q500) at a heating rate of 5 °C min−1 from room temperature to 600 °C under air. Crystallographic characterization was carried out by powder X-ray diffraction (XRD, Shimadzu XRD-6000) operating at 40 kV and 40 mA using Cu−Kα1 radiation (λ = 0.15406 nm) with the copper target and nickel filter. The samples were scanned between 10 to 80° (2θ) at a scan rate of 2° min−1. Elemental compositions of samples were analyzed by energy-dispersive X-ray spectroscopy (EDX) using INCA mapping attached with the field-emission scanning electron microscopy (FE-SEM, JEOL 7600F) with an accelerating voltage of 15 kV. Prior to recording FE-SEM, the samples were placed on carbon tape and sputter coated with platinum (Pt) for 40 s. Morphology and microstructure of samples were observed by high-resolution transmission electron microscopy (HR-TEM, JEOL 2100F) at an accelerating voltage of 200 kV. The samples were prepared by dropping the dispersion of A-CoOx nanostructures in ethanol onto a carboncoated copper grid and dried at room temperature. Specific surface area was determined by Brunauer−Emmett−Teller (BET, Micromeritics ASAP-2020) test. Surface analysis of the studied samples was performed using X-ray photoelectron spectroscopy (XPS, Kratos AXIS Ultra DLD). XPS measurements were carried out using a focused monochromatized Al Kα radiation (hv = 1486.6 eV). The spectrometer was calibrated using the photoemission line Ag 3d5/2 (binding energy (BE) 368.26 eV). For the Ag 3d5/2 line, the full width at halfmaximum (fwhm) was 0.61 eV under the recording conditions. The analyzed area of the samples was 300 × 700 μm2. All of the measurements were performed under charge neutralization. The pressure in the analysis chamber was ca. 5 × 10−8 Pa. The BE scale was calibrated from the carbon contamination using the C1s peak at 284.5 eV. The spectra were analyzed using a peak synthesis program (CasaXPS, version 2.3.15), fitting with distribution of Gaussian (80%), Lorentzian (20%), and Shirley background. Fourier transform infrared spectroscopy (FT-IR,
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RESULTS AND DISCUSSION Thermal Analysis. Thermal properties of CoOx are studied by TGA. The as-synthesized H-CoOx nanostructures (rods and sheets, that is, H-CoOx-NR and H-CoOx-NS) contained two water molecules of crystallization. A gradual weight loss (∼20 wt %) is observed for H-CoOx nanostructures from 100 to 200 °C (Figure 1), which is in good agreement with the theoretical weight loss according to eq 1 (water loss). At 290 °C, a drastic weight loss of ∼55.4 wt % (theoretical weight loss: 56.1 wt %
Figure 1. Thermogravimetric analysis (TGA) curves (5 °C min−1 from room temperature to 600 °C under air). 16317
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Å, b = 5.03 Å, and c = 5.40 Å and β = 113.62°.17 The change in crystal structure after drying of H-CoOx is due to the removal of crystallized water molecules.13 No impurity phase was detected for H-CoOx and A-CoOx nanostructures. It is also observed that all samples exhibited broad peak widths, indicating small crystallites size, which is further confirmed by HR-TEM. A representative EDX spectrum of A-CoOx-NR indicating the material is phase pure and is well supported by the XRD data. The detected platinum (Pt) is sputter-coated onto the samples prior to EDX and no other contaminants are being detected (see Supporting Information, Figure S1). TEM/HR-TEM micrographs revealing the fine morphological details of A-CoOx nanostructures are shown in Figure 3. Two different morphologies are obtained depending on the solvent medium in which the synthesis takes place. CoOx prepared in N,N-diethylacetamide (DEAc) exhibited nanosized rods (Figure 3a) with 20−45 nm in width and 0.2 to 0.3 μm in length. When N-methyl-2-pyrrolidione (NMP) is used as the solvent, rectangular nanosheets are formed (Figure 3b). It is seen that both A-CoOx-NR and NS samples agglomerate slightly, and a magnified SEI micrograph (Figure 3a,b insets) provided a clearer morphological differentiation. It is clearly observed that A-CoOx-NR shows interior porous structure, which may be generated due to release of crystallized water molecules during dehydration process. There are no considerable pores observed in A-CoOx-NS nanostructure. From HRTEM micrographs (Figure 3c,d insets), both A-CoOx-NR and NS exhibited observable lattice fringes, indicating crystalline nature with an average of 0.2213 and 0.2472 nm interplanar spacing indexed to (12−1) and (002) planes, respectively, which is in good accordance with XRD observations. BET specific surface area for A-CoOx-NR and NS is 82 and 67 m2 g−1, and average pore diameter is 186 and 80 Å, respectively, as calculated by the analysis program (see Supporting Information, Figure S2). A-CoOx nanostructures exhibit mesoporous nature. BET measurements are consistent with the TEM micrographs because much larger pores are visible in A-CoOxNR than in NS. There are no significant differences in the size and shape of the samples observed after dehydration. In this chimie douce precipitation method, the metal ion and oxalic acid do not react immediately due to large differences in ionization ability in organic solvent medium.18 Previous reports show that differences in solubility, nucleation and growth of the products, reaction kinetics, solution properties, and stability of particles in different solvent systems are critical factors for forming various shapes.19 In other words, the selective interaction between solvents and crystallographic planes of CoOx is the key factor in morphological variations. Referring to a previous report,13 the pairing between the type of metal ion (using cobalt instead of iron precursor) and solvent types (keeping to the same two organic solvents, i.e., DEAc and NMP) has a great influence on the morphologies being achieved, as evidenced in this work. Hence, morphological variations can be attained by controlling the ion exchange reaction with the different combination of metal ions and organic solvents. Cyclic Voltammetry. CV measurements were performed to elucidate the electrochemical process of A-CoOx-NR and NS electrodes and its CV cycles are shown in Figure 4a,b, respectively. All A-CoOx electrodes depicted similar CV profiles. For A-CoOx-NR, two cathodic current peaks at 1.18 and 0.68 V were observed in the first cycle. The high intensity peak located at 1.18 V in the first cycle shifted to higher
according to eq 2) is observed for H-CoOx nanostructures, which is ascribed to oxalate decomposition. After dehydration, there is no considerable weight loss observed up to 300 °C for A-CoOx nanostructures (rods and sheets, that is, A-CoOx-NR and A-CoOx-NS). A weight loss of ∼45.5 wt % (theoretical weight loss: 45.4 wt % according to eq 3) is observed between 300 and 320 °C for A-CoOx nanostructures. No further significant weight loss is observed above 320 °C, indicating complete decomposition of oxalate for H-CoOx and A-CoOx nanostructures. The results indicate that at 250 °C H-CoOx converts to A-CoOx by the removal of crystallized water molecules and without breakage of the oxalate anion, as shown in the following equations: Dehydration: 100 to 200 °C CoC2O4 · 2H 2O → CoC2O4 + 2H 2O↑
(1)
Decomposition: >290 °C 3CoC2O4 · 2H 2O → Co3O4 + 2CO2 ↑ + 4CO ↑ + 6H 2O↑
3CoC2O4 → Co3O4 + 2CO2 ↑ + 4CO↑
(2) (3)
Structural and Morphological Characterization. HCoOx (chemical formula: CoC2O4.2H2O) crystallizes in two allotropic forms, α monoclinic (space group: C2/c) and β orthorhombic (space group: Cccm).16 Figure 2a,b shows the
Figure 2. XRD patterns for (a) H-CoOx and (b) A-CoOx samples.
XRD patterns of both H-CoOx and A-CoOx nanostructures, respectively. All of the diffraction peaks for H-CoOx samples are readily indexed to the orthorhombic β-phase (ICDD PDF no. 00-025-0250) with a = 11.88 Å, b = 5.42 Å, and c = 15.62 Å. In contrast, the diffractograms are entirely different for ACoOx. For A-CoOx nanostructures, the pattern is indexed in a monoclinic lattice (ICDD PDF no. 00-037-0719) with a = 5.74 16318
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Figure 3. TEM/HR-TEM micrographs: (a,c) A-CoOx-NR and (b,d) A-CoOx-NS. Insets: (a) and (b) showing magnified FE-SEM micrographs. (c,d) Lattice fringes.
Li. From Figure 5a, A-CoOx-NR electrodes delivered an initial discharge capacity of 1599 mA h g−1 and charge capacity of 994 mA h g−1 with a Coulombic efficiency of 62%. From the 50th discharge profile, only one voltage plateau between 0.8 and 0.6 V corresponding to the reduction of Co2+ to Co0 and a discharge capacity of 1203 mA h g−1 are observed. Up to the 100th cycle, the discharge and charge capacity remained ∼959 and 947 mA h g−1, respectively, which is much higher than commercial graphite (372 mA h g−1) and those recently reported.12 As compared with the 50th discharge profile, ACoOx-NR also exhibits a similar voltage plateau at the 100th discharge cycle. A similar trend of electrochemical reactions of A-CoOx-NS with Li is depicted in Figure 5b. However, ACoOx-NS showed a slightly lower discharge and charge capacity of about 1518 and 985 mA h g−1, respectively, for the first cycle. This probably originates from morphological differences between the samples, and a larger surface area of ACoOx-NR can provide more sites for lithium ion conversion. The discharge capacities at the 50th and 100th cycles for ACoOx-NS are 833 and 741 mA h g−1, respectively. This electrochemical behavior of A-CoOx nanostructures is consistent with the CV studies.
potential (1.53 V) during the second cycle and disappeared after the third cycle. This may be attributed to the following reasons: (i) decomposition of electrolyte that results in the formation of an organic layer deposited on the surface of the particles20,21 and (ii) a conversion reaction of lithium-ion intercalation to form Li2C2O4 and metallic cobalt (Co0) particles as shown in eq 4. CoC2O4 + 2Li+ + 2e− → Co + Li 2C2O4
(4)
Meanwhile, an oxidation peak at about 1.21 V was recorded during the anodic process, corresponding to the reversible oxidation step of Co0 to Co2+ and decomposition of Li2C2O4, as seen in subsequent cycles. The CV agrees well with that previously reported for CoC2O4.12 Evaluation in Half-Cells. The electrochemical performance of A-CoOx with respect to Li+-ion conversion/deconversion was investigated in Li/CoC2O4 test cells by galvanostatic charge−discharge testing in the voltage window of 0 to 3 V. Figure 5 shows the voltage profile at a constant 1 C rate and the cycling behavior of the samples at various C rates. It can be seen that A-CoOx-NR electrode materials exhibited the initial discharge potential plateau at ∼1.2 V, followed by weak slope corresponding to the electrochemical reaction of CoC2O4 with 16319
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irreversible. The A-CoOx charge−discharge profile is observed to be significantly different from other conventional oxide electrodes, demonstrating that the conversion reaction may be due to the lithium oxalate formation during the discharge instead of lithium oxide.11,22 Figure 5c,d show the rate capabilities and galvanostatic cycling performance of A-CoOx in Li/CoC2O4 test cells cycled between 0 to 3 V under various C rates ranging from 1 to 10 C rate. In the case of A-CoOx-NR, it is observed that there is fast capacity fading up to seven cycles and starts to increase thereafter. For 1 C rate, a peak discharge capacity of 1240 mA h g−1 is observed at the 62nd cycle, and a gradual decrease in the capacity ensues thereafter. It is observed that after 55 cycles, the cells deliver stable reversible capacity of ∼903, 772, and 735 mA h g−1 at 2, 5 and 8 C rate, respectively (up to 200 cycles), which is 200 to 247% higher than theoretical capacity. As the current density increases, the discharge and charge capacities of A-CoOx-NR decreased evidently. At 10 C rate, a very stable reversible capacity of ∼310 mA h g−1 (up to 100 cycles) is observed after 10 cycles. A-CoOx-NS showed initial capacity fade up to nine cycles and started to increase thereafter, as shown in Figure 5d. With comparison to A-CoOx-NR, a lower peak discharge capacity of 977 mA h g−1 (at 22nd cycle) is observed for A-CoOx-NS at 1 C rate, and a similar gradual decrease in the capacity follows in subsequent cycles. After 100 cycles, the cells deliver a discharge capacity of ∼824, 727, 529, and 330 mA h g−1, respectively, at 2, 5, 8, and 10 C rates. The Coulombic efficiency of more than 98% was attained for ACoOx nanostructures at different C rates after 10 cycles. It is surprising to note that capacity increases in A-CoOx with cycling (more significant at low C rates) after an initial capacity fade. This phenomenon is not yet completely understood. During cycling, the capacity is observed to be higher than theoretical capacity, as shown in the proposed conversion mechanism (eq 4). The origin of this extra capacity is mainly
Figure 4. CV profile at 0−3 V at a scan rate of 0.1 mV s−1: (a) ACoOx-NR and (b) A-CoOx-NS.
The initial discharge capacity for A-CoOx is observed to be larger than the theoretical capacity of CoC2O4 (365 mA h g−1 based on eq 4), which may be due to the decomposition of nonaqueous electrolyte during discharge.21 It is clear that first discharge capacity is mostly due to faradaic contribution and is
Figure 5. Charge−discharge profile versus specific capacity (1 C rate) and galvanostatic cycling performance at various C rates: (a,c) A-CoOx-NR and (b,d) A-CoOx-NS. 16320
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relative current contribution of faradaic and nonfaradic phenomenon. It is clearly seen that the primary contribution is the faradaic component for A-CoOx nanostructures at different sweep rates. This indicates that at low sweep rate the reaction mechanism is dominated by Li conversion process. However, the nonfaradaic component for A-CoOx increases with sweep rates, as dictated by the decrease in faradaic component. From Figure 6a, A-CoOx-NR (fresh cell) presents larger faradaic contribution at low sweep rates. The faradaic current contribution is decreased from 89 to 71% with increased sweep rate (0.5 to 5 mV s−1). A-CoOx-NR has a better rate capability than A-CoOx-NS (Figure 6b), as observed from the greater decrease in faradic contribution. The difference is probably attributed to physical characteristics like morphology and pore size. In addition, A-CoOx-NR showed more stable cycling behavior than A-CoOx-NS, as seen from the minimal changes between the 50th and 100th cycle analyzed data shown in Figure 6a, which is in good agreement with the EIS results (as discussed in the following section). The voltammetric sweep study complements the behavioral trend from galvanostatic charge−discharge testing at various C rates. These results are also consistent with previous reports.13−15 To elucidate the electrochemical processes occurring in ACoOx electrodes during charge−discharge, we performed EIS on the cells cycled at 1 C rate (1st, 50th, and 100th cycle). An equivalent circuit is constructed to evaluate the formation of surface film and charge-transfer process.13 Nyquist plots of ACoOx electrodes are recorded for fresh cells (equilibrated overnight at room temperature) at open circuit potential as well as for the cycled cells at 1 C rate and shown in Figure 7. Typically in LIBs, the high-frequency semicircle (100 kHz to 10 Hz) is related to the charge-transfer processes, while the spike at middle-to-low frequencies describes lithium diffusion kinetics.26 The equivalent circuit (Figure 7d) of built-in series and parallel combinations of intrinsic resistance (Rs), surface and charge-transfer resistance (Rsf+ct), double-layer capacitance across the surface of the electrode (CPEsf+dl) (with angle of distortion α),13 Warburg impedance (Zw), and intercalation capacitance (CPEint)27 enables the description of the various processes occurring, and the values of the elements via fitting are presented in Table 1. Comparing the Rs values of the fresh cells of A-CoOx-NR (5.4 Ω) and A-CoOx-NS (2.9 Ω), as shown in Figure 7a, it is noted that there is negligible difference between the intrinsic resistances of these morphologies. The higher Rsf+ct value in ACoOx-NR indicates larger formation of passivation layer on the surface of the electrode. The higher intercalation capacitance CPEint of A-CoOx-NR (∼6.2 mF) as compared with A-CoOxNS (∼4.8 mF) indicates a higher number of lithium diffusion pathways. In general, larger CPEint values are three orders of magnitude higher than CPEsf+dl, indicating larger dependence on lithium diffusion kinetics, rather than double-layer kinetics, as shown in Table 1. There are negligible differences in the Rs values of the Li/A-CoOx-NR cells observed before and after cycling (Figure 7b). During the first charge−discharge cycle, the accumulation of the passivation layer is reduced slightly (compared with fresh cells), judging from the decrease in Rsf+ct values from 95 to 37 Ω. From the first electrochemical cycle, the initial charge−discharge propagates the movement of charge-carrying ions within the electrolyte and leads to an initial breakdown of the solid−electrolyte interface (SEI). Further cycling up to 50th cycle leads to possible increase in the formation of SEI, which could be due to continuous charge
due to the faradaic contribution that comes from the side reactions with the electrolyte or reversible formation of a gellike layer as previously reported.23,24 In addition, the porous nature of A-CoOx nanostructures can also benefit the electrochemical lithium storage processes as well. To evaluate cycling stability for A-CoOx nanostructures, we made comparisons after the region of increasing capacity with cycling, that is, between 100th and 200th cycles. It is observed that A-CoOx-NR obtained higher specific capacity values than NS at the end of 200th cycle for all C rates. Noticeably, the cycling stability of both A-CoOx-NR and NS is relatively comparable to the exception at 1 C rate, where A-CoOx-NR has a higher capacity fade of ∼17% than NS. In general, ACoOx-NR achieved a higher reversible capacity, better cyclability, and rate capabilities than A-CoOx-NS (especially at
