To ensure low electrical resistance and high conductivity, minimizing energy loss, copper (Cu) foil of 99.9 % purity was used for this study. After cleaning with acetone and drying in an air drier, two pieces of foils placed in quartz crucibles were heated from room temperature up to 250 °C and 350 °C in a furnace and both samples were subjected to thermal oxidation in air for 2 h. The temperature was maintained at the set value during oxidation precisely and then the furnace was turned off allowing the samples to get cooled down to room temperature slowly. Later, the shape, pore size and distribution of the resultant nanostructured Cu-oxides (NCu_xO_y) layers were evaluated following removal of the copper substrate cautiously. For the readers' convenience, Cu foils oxidized at 250 °C and 350 °C are referred to as N250-Cu_xO_y and N350-Cu_xO_y, respectively hereafter.

The N250-Cu_xO_y and N350-Cu_xO_y samples were characterized using x-ray powder diffraction (XRD) to investigate the crystallographic data and especially identify the phases present. To account for various scattered and random orientations, scans were performed at different 2θ angles, ranging from 10° to 90° with a step size of 0.05° using CuK_α radiation (λ = 0.15406 nm) from an Empyrean XRD system (Panalytical, Netherlands). Morphological and compositional analysis of the NCu_xO_y were conducted using scanning electron microscopy (SEM) and energy dispersive X-ray spectroscopy (EDX) attached to the SEM. All the images were recorded using a high-quality JSM 7600F instrument (JEOL, Japan) at 5.0 kV at various magnifications.

Coin cell battery cases (CR2032, supplied by Lith Co., China) were used for assembly for the current study. The anode samples consisted of 18 mm diameter and 0.1 mm thick (~2 cm²) NCu_xO_y, while the cathodes were LiCoO₂-coated aluminum sheets. A 20 mm diameter battery-grade polyvinyl chloride (PVC) separator was placed between the electrodes. The PVC separator was used due to its excellent electrical insulating properties along with the chemical stability to prevent short circuits between electrodes as ions are conducted during the operational time of the LIBs. Both sides of the separator were soaked with a few drops of 1M LiPF₆ electrolyte (supplied by Ximen Tmax Battery Equipment Ltd., China), dissolved in a 3:7 volume ratio of ethylene carbonate (EC) and diethylene carbonate (DEC). The battery was then sealed air-tight by crimping at approximately 100 psi using a battery crimping machine (Metrology Lab, CUET).

XRD patterns of both the samples fabricated at different oxidation temperatures are presented in Fig. 1. Two distinct oxide phases, cubic Cu₂O (space group Pn–3m, JCPDS #05-0667) and monoclinic CuO (space group C2/c, JCPDS #45-0937), were identified for both samples. Cu₂O seems to be the dominant phase at a lower oxidation temperature of 250 °C, with characteristics peaks positioned at 2θ = 36.90°, 43.40°, 50.50°, 74.50°, and 90.00°. As the oxidation temperature increased to 350 °C, the CuO phase seems to be the dominant phase as evidenced from the increasing intensity of the CuO diffraction peaks, particularly the prominent peaks located at 2θ = 35.70°, 38.70°, and 66.60°. In addition, the presence of the low-intensity peaks can be attributed to the face-centered cubic Cu (space group Fm–3m, JCPDS #04-0836) due to partial exposure of the Cu substrate from non-uniform oxidation. The diffraction data were analyzed using Scherrer's formula,

D = 0.9λ/Bcos(θ),

where λ is the X-ray wavelength, B is the full width at half maximum (FWHM) of the peaks and 2θ is the angle between the incident and diffracted beams.

The average crystal size of Cu₂O and CuO of N250-Cu_xO_y is found to be 100 nm and 160 nm, respectively. In contrast, the average crystal size of Cu₂O and CuO of N350-Cu_xO_y is found to be 150 nm and 180 nm, respectively. It is noted that the average crystallite sizes of both Cu₂O and CuO increased with the increased oxidation temperature reflecting the improvement in crystalline quality and decreased micro-strain with reduced crystal imperfection due to higher solid-state diffusion [10–12]. Thus, a higher oxidation temperature leads to the formation of a larger amount of CuO as the dominant phase because of a higher degree of oxidation whereas a larger crystallite size results from the higher thermal energy for diffusion. The larger crystallite size of the CuO phase at both oxidation temperatures results because CuO is formed with a higher oxidation state of +2 compared to the lower oxidation state of Cu₂O with an oxidation state of +1. A higher oxidation state favors higher crystal size because of several factors including the increased phase stability and stronger interaction of copper ions with oxygen ions in the lattice promoting the growth of the larger crystallite size.

Figure 2 shows the SEM and EDX spectra of N250-Cu_xO_y and N350-Cu_xO_y. It is noted that N250-Cu_xO_y exhibits a surface resembling an ocean mesh, akin to a circular ripple, with forming wave crests and radial streaks, accompanied by a light sea spray mist above the surface as shown in Fig. 2a. In addition, the surface morphology of N350-Cu_xO_y entirely changed to a porous structure with catkin as shown in Fig. 2b. Usually, it is expected that higher oxidation temperatures should result in a denser structure. However, here a more porous surface morphology at a higher oxidation temperature can result from several factors like thermal stress and a faster oxidation rate. At higher oxidation temperatures the differential thermal expansion between the copper foil and the oxide layer can induce thermal stress leading to increased porous structure. The faster oxidation rate also can contribute to the more porous morphology differential oxidation rate across the thickness of the oxide layer.

However, the higher magnification image of N350-Cu_xO_y exhibits a more polished surface compared to N250-Cu_xO_y due to the more efficient oxidation process, improved atomic diffusion, and the formation of a more uniform and continuous oxide layer [13–15]. More smother surface morphology even with higher porosity evident in the presence of small size evenly distributed porosity for the oxide layer oxidized at higher oxidation temperature.

Figure 3 shows the change of cell current and voltage with respect to the time when N350-Cu_xO_y was used as the anode of the LIB. It shows the change of current and voltage with respect to time in Fig. 3a for 1–20 and in Fig. 3b for 21–40 charge-discharge cycles.

It is to be noted that the assembled coin cell with N250-Cu_xO_y as an anode of LIB was not electrochemically cycled due to a very small amount of oxide formation and loose connections of oxides with Cu foil. The above-discussed characterization also showed that the oxide formed at an oxidation temperature of 350 °C was of better quality with a higher degree of crystallinity. Hence, we discarded the lower-temperature oxidation sample for further analysis. During the charge-discharge cycles, copper oxides undergo electrochemical reactions where lithium ions intercalate the anode during charging and leave the anode during discharging. During the discharging cycle, lithium ion is oxidized at the anode and forms metallic Cu at anode and Li₂O at cathode. The plotted data represents galvanostatic charge-discharge cycles, where voltage pulses were applied, and the current response was measured. The sharp peaks in the current correspond to the rapid lithium-ion insertion/extraction, while the gradual decay indicates diffusion limitations and internal resistance effects. Some points are higher because of variations in the electrochemical kinetics, including differences in Li-ion diffusion, surface reactions, and resistance changes as the electrode material undergoes structural transformations. The progressive decrease in peak height suggests capacity is fading, which can result from material pulverization, loss of electrical contact, or side reactions.

Figure 4 shows the differential capacity (dQ/dV) vs. voltage for a LIB with copper oxides (N350-Cu_xO_y) as the anode and LiCoO₂ as the cathode. The two graphs correspond to electrochemical impedance spectroscopy (EIS) data, taken at two different states of the battery: Fig. 4a is for the fresh cell before cycling, showing the initial electrochemical behavior, and Fig. 4b is obtained after 40 charge-discharge cycles displaying changes due to cycling. It is to be noted that the wider, symmetric triangular shape indicates that the battery initially has a well-defined electrochemical response in Fig. 4a. It indicates a highly reversible electrochemical reaction means charge and discharge behavior is quite consistent. In addition, the multiple peaks suggest distinct redox reactions, likely corresponding to the lithiation and delithiation of CuO or Cu₂O and LiCoO₂. Also, the steep nature of the curves may indicate low internal resistance and high lithium-ion mobility in the fresh cell [16]. In contrast, after 40 charge-discharge cycles, the curves become more asymmetric and narrower, indicating increased resistance and capacity fading over cycling, as shown in Fig. 4b. A noticeable shift in the peaks suggests a change in reaction kinetics, likely due to SEI layer formation, electrode degradation, or lithium plating. In addition, the increased broadness in dQ/dV peaks indicates possible irreversible side reactions or electrode structural changes affecting charge storage [17]. However, the higher resistance (steeper voltage shifts) at low voltages suggests the development of passivation layers, electrode pulverization, or loss of active material [18, 19].

Figure 5 shows the change of specific capacity and coulombic efficiency during (a) 1–20 and (b) 21–40 charge-discharge cycles at 1C current rating. Figure 5a represents the specific capacity (mAh/g) vs cycle number for the first 20 charge-discharge cycles. The blue and orange lines likely correspond to the charge and discharge capacities, respectively, while the gray curve represents the Coulombic efficiency (%). The 1^st cycle charge and discharge capacity are observed to be 455.4 and 431.9 mAh/g, respectively, with a Coulombic efficiency of 94.8 %. However, the battery shows a gradual decrease in specific capacity, indicating some capacity fading and discharge capacity reduced to 352.2 mAh/g, which could be attributed to electrode degradation, SEI layer formation, or lithium loss [17]. The large fluctuation (spike) in Coulombic efficiency around cycles 10 to 12 suggests an experimental anomaly, electrolyte instability, or lithium dendrite formation, which could temporarily affect performance [20].

Figure 5b shows the cycling performance from cycle 21 to cycle 40, focusing on specific capacity and Coulombic efficiency. Here, the blue and orange lines stabilize, suggesting that the battery reaches a more steady-state cycling regime. The Coulombic efficiency improves and remains stable (~82–85 %), meaning that most of the lithium is being effectively cycled, with minimal losses. The charge and discharge capacity at the 21^st cycle is observed to be 440.2 mAh/g and 339 mAh/g, respectively. The discharge capacity was found to be 339 mAh/g at the end of the 40^th cycle. The capacity retention appears relatively stable after 40^th charge-discharge cycle, however, the gradual decline in Coulombic efficiency suggests ongoing degradation processes such as electrode pulverization, SEI layer thickening, or lithium trapping [20].

Hence, the data suggests that the N350-Cu_xO_y anode experiences initial instability during the first few cycles, likely due to SEI formation and structural changes. However, after approximately 20 cycles, the system stabilizes, leading to more consistent cycling performance. Despite this stabilization, long-term capacity fading is evident, which is a common challenge in the transition of metal oxide anodes due to their volume expansion and structural degradation.