Construction of CuCo₂O₄@NiFe-LDH Core–Shell Heterostructure for High-Performance Hybrid Supercapacitors

Abstract

Transition metal oxides (TMOs) are considered to be highly promising materials for supercapacitor electrodes due to their low cost, multiple convertible valence states, and excellent electrochemical properties. However, inherent limitations, including restricted specific surface area and low electrical conductivity, have largely restricted their application in supercapacitors. In this paper, core–shell heterostructures of nickel–iron layered double hydroxide (NiFe-LDH) nanosheets uniformly grown on CuCo₂O₄ nanoneedles were synthesized by hydrothermal and calcination methods. It is found that the novel core–shell structure of CuCo₂O₄@NiFe-LDH improves the electrical conductivity of the electrode materials and optimizes the charge transport path. Under the synergistic effect of the two components and the core–shell heterostructure, the CuCo₂O₄@NiFe-LDH electrode achieves an ultra-high specific capacity of 323.4 mAh g⁻¹ at 1 A g⁻¹. And the capacity retention after 10,000 cycles at 10 A g⁻¹ is 90.66%. In addition, the assembled CuCo₂O₄@NiFe-LDH//RGO asymmetric supercapacitor device achieved a considerable energy density (68.7 Wh kg⁻¹ at 856.3 W kg⁻¹). It also has 89.36% capacity retention after 10,000 cycles at 10 A g⁻¹. These properties indicate the great potential application of CuCo₂O₄@NiFe-LDH in the field of high-performance supercapacitors.

1. Introduction

The intermittent nature of renewable energy sources (e.g., solar, wind) and the limitations of conventional energy storage systems in terms of energy/power tradeoffs are two of the core challenges facing the global energy sector. At the same time, a great deal of effort has been invested in developing new energy storage systems, such as lithium-ion batteries and other secondary batteries [1,2,3,4]. Compared with lithium-ion batteries and other secondary batteries, supercapacitors (SCs) have two energy storage mechanisms, namely, double-layer capacitance and pseudocapacitance. And they have fast charging and discharging capabilities, higher safety, and long cycle life [5,6]. These characteristics establish them as critical enablers for renewable energy grid integration and start–stop systems in electric vehicles. Nevertheless, the practical application of SCs remains constrained by their comparatively lower energy density relative to conventional electrochemical battery systems [7,8,9,10]. However, the conductivity, stability, and number of active sites of electrode materials are closely related to the composition and structure. Therefore, the design of electrode materials integrating high specific surface area, elevated energy density, and abundant redox-active sites is pivotal for practical implementation in energy storage systems [11,12,13]. Transition metal oxides (TMOs) have garnered significant academic attention by virtue of their rich redox chemistry coupled with multi-electron transfer capability, emerging as pivotal candidates in pseudocapacitive energy storage systems [14]. Among various TMOs, CuCo₂O₄ is one of the promising candidates due to its environmental friendliness, multiple convertible valence states, and excellent electrochemical properties [15,16]. Up to now, a variety of CuCo₂O₄ electrode structures have been prepared, and these structures have a significant impact on electrochemical performance. However, the actual capacity of CuCo₂O₄ is significantly lower than its theoretical value. This limitation hinders its practical application in SCs. The core–shell architecture synergistically combines the structural and functional advantages of different components, enhancing electrochemical performance through increased active sites and enlarged specific surface area. Combining it with other materials to build proper core–shell heterostructures can effectively boost its electrochemical performance [17]. Transition metal layered double hydroxides (LDHs) have drawn significant attention in the field of SCs because of their tunable chemical composition, anion exchangeability, and high redox activity [18,19,20]. In particular, NiFe-LDH is often used as an electrode material for SCs owing to its strong interlayer ion diffusion capability, environmental friendliness, and high efficiency [21,22,23].

However, the electrochemical properties of NiFe-LDH are greatly limited in practical applications due to its limited voltage window and tendency to agglomerate. Some rational structural engineering of electrode materials has been shown to enhance charge storage performance by optimizing ion diffusion kinetics and surface redox, such as porous structure design, heterogeneous structure construction, dynamic interface engineering, and three-dimensional conductive network construction [24]. The core–shell heterostructures help to enhance the performance of individual materials by achieving larger potential windows and capacity values, thereby boosting their practical utility [25]. Also, the shells of the core–shell heterostructure possess more redox sites. In recent years, an increasing amount of research endeavors have been dedicated to the fabrication of core–shell electrode materials. Zhao’s team [26] prepared dandelion-like CuCo₂O₄ core–shell nanoflower structures with ultrathin NiMn-LDH shells by a two-step hydrothermal method. The findings demonstrated that the CuCo₂O₄@NiMn-LDH core–shell nanoflower obtained a specific capacitance of 2156.5 F g⁻¹ at 1 A g⁻¹ and 94.6% cycle retention after 2500 cycles at 5 A g⁻¹. Zhang’s team [27] proficiently fabricated CuCo₂O₄@MoNi-LDH core–shell assemblies via hydrothermal and calcination procedures. The findings demonstrated that the CuCo₂O₄@MoNi-LDH attained a specific capacitance (1286 F g⁻¹ at 1 A g⁻¹). Mishra’s team [28] engineered ZnO@Co₃O₄ core–shell heterostructures (ZC-CSH) through a single-step in situ growth solvothermal method. It exhibits an outstanding specific capacitance of 177 F g⁻¹ at 1.4 A g⁻¹. In addition, ZC-CSH manifests remarkable stability, maintaining a high level of 92.8% even after undergoing 10,000 cycles. Tamtam’s team [29] pioneered the core–shell Mn@Al-MOF structure through a multidimensional layered design. The results showed that this core–shell structure greatly improved the specific capacitance and cycling stability (>95% capacity retention after 10,000 cycles). Chavan’s team [30] has synergistically integrated CoMn-LDH flakes with heterostructure engineering. Excellent bifunctionality was achieved for supercapacitor (energy density: 88.6 Wh kg⁻¹) and electrocatalytic applications (OER overpotential: 298 mV@10 mA cm⁻²).

Inspired by the above considerations, in an effort to further enhance the practical specific capacity of CuCo₂O₄, an ultrathin NiFe-LDH nanosheet core–shell structure was successfully grown on CuCo₂O₄ by the hydrothermal method in this study. In this core–shell configuration, CuCo₂O₄ offers a conductive framework for NiFe-LDH two-dimensional nanosheets and promotes ion and electron transport across the interface. Simultaneously, the NiFe-LDH nanosheets serve as the “shell” to expand the surface area and facilitate the contact between the ions and the electrolyte. Furthermore, the distinctive core–shell strategy increases the conductivity of the electrode material and optimizes the charge-transfer pathway. The prepared CuCo₂O₄@NiFe-LDH core–shell structure exhibits an extraordinarily high specific capacity of 323.4 mAh g⁻¹ at a current density of 1 A g⁻¹. Impressively, after 10,000 consecutive charge–discharge cycles at a current density of 10 A g⁻¹, it still retains a capacity retention rate of 90.66%.

2. Experimental Process

2.1. Synthesis of CuCo₂O₄ Nanowire Arrays

A total of 2 mmol of Co (NO₃)₂·6H₂O, 1 mmol of Cu (NO₃)₂·3H₂O, and 5 mmol of urea were dissolved in 35 mL of deionized (DI) water. The mixture was stirred for 30 min to form a uniform solution. Subsequently, the obtained solution was transferred to a 50 mL autoclave, and a precleaned nickel foam (NF, 1 × 2 cm⁻¹) was placed at the center of the solution. The autoclave was kept at 120 °C for 12 h. After cooling naturally, the nickel foam was rinsed several times with ethanol (C₂H₅OH) and DI water, then dried at 60 °C for 12 h. Finally, the obtained samples were annealed at 350 °C for 2 h at a heating rate of 2 °C/min to obtain CuCo₂O₄.

2.2. Synthesis of CuCo₂O₄@NiFe-LDH

CuCo₂O₄@NiFe-LDH was synthesized via the hydrothermal method. A total of 6 mmol Fe (NO₃)₂·9H₂O, 18 mmol Ni (NO₃)₂·6H₂O, 2 mmol Na₃C₆H₅O·2H₂O, and 40 mmol urea were dissolved within 30 mL of DI water with magnetic stirring. Then, the above yellow solution was poured into a 50 mL hydrothermal autoclave. At the same time, the prepared CuCo₂O₄ was suspended at the center of the solution, and the autoclave was kept at 150 °C for 8 h. After cooling naturally, it was rinsed repeatedly with C₂H₅OH and DI water and dried under vacuum for 10 h to obtain CuCo₂O₄@NiFe-LDH. For more information on material preparation processes and testing of electrochemical materials, see the Supporting Information.

3. Results and Discussion

The synthesis process of CuCo₂O₄@NiFe-LDH is shown in Figure 1. CuCo₂O₄ microspheres were initially synthesized on NF substrates via a combined hydrothermal calcination approach. Then, NiFe-LDH nanosheets were grown on CuCo₂O₄ by a secondary hydrothermal process, and the unique morphology of CuCo₂O₄@NiFe-LDH was obtained by controlling the hydrothermal reaction time. Among these, CuCo₂O₄ provides conductive support for the core–shell structure, and NiFe-LDH nanosheets serve as the “shell”, boosting the surface area and making it easier for ions to come into contact with the electrolyte.

In Figure 2, the phase composition of CuCo₂O₄, NiFe-LDH, and their composite was investigated by X-ray diffraction (XRD) measurement. Among these, characteristic diffraction peaks observed at 44.4°, 51.7°, and 76.2° were attributed to the NF substrate (JCPDS 04-0850) [31]. For the spinel-type CuCo₂O₄ phase, six diffraction peaks were identified at 2θ values of 18.78° (111), 31.12° (220), 36.6° (311), 44.5° (400), 58.9° (511), and 64.8° (440), which correspond well to the standard cubic spinel structure (JCPDS 78-2177). It can be seen that the generated CuCo₂O₄ does not contain impurity peaks of copper oxide and cobalt oxide, indicating that the purity of the sample is high. Figure 2 shows the XRD characterization of NiFe-LDH. The diffraction peaks at 2θ of 11.3, 22.8, 33.9, 38.6, 46.2, 59.8, and 61.1° match well with the standard cards of NiFe-LDH, corresponding to the (003), (006), (012), (015), (018), (110), and (113) crystal faces [32]. Notably, the XRD pattern of CuCo₂O₄@NiFe-LDH exhibits well-preserved diffraction peaks from both the spinel CuCo₂O₄ and NiFe-LDH components. The absence of impurity-related peaks further verified the successful construction of the CuCo₂O₄@NiFe-LDH core–shell heterostructure.

The surface morphologies of CuCo₂O₄, NiFe-LDH, and their core–shell composite CuCo₂O₄@NiFe-LDH were investigated through scanning electron microscopy (SEM). Figure 3a,d presents SEM images of CuCo₂O₄, where smooth nanoneedles (2 μm in length, 30 nm in diameter) are uniformly distributed on nickel foam, forming a sea urchin-like morphology. The anisotropy of CuCo₂O₄ nanoneedles is very strong, enables their effective deployment as rapid transport channels for both electrons and ions [33]. In addition, there is enough space between the nanoneedles to promote the continuous growth of NiFe-LDH nanosheets on it. Figure 3b,e shows SEM images of NiFe-LDH, where 20–30 nm thick nanosheets are vertically aligned and uniformly distributed on the nickel foam, forming an interwoven honeycomb-like architecture. This interweaving has the ability to enlarge the contact area that the nanosheets have with the electrolyte. Consequently, it enhances the electrolyte transfer rate and reduces volume expansion. Modulation of the morphology of NiFe-LDH (Figure S1) loaded on sea urchin-like CuCo₂O₄ by controlling the hydrothermal time. The NiFe-LDH nanosheets obtained at the hydrothermal time of 8 h were uniformly covered with CuCo₂O₄ nanoneedles so that the stability of the core–shell nanostructures was well maintained (Figure 3c,f). The interwoven nanosheets expose abundant active sites that facilitate the diffusion of electrolyte ions, thus improving the reaction kinetics. Moreover, they can act as a protective layer to mitigate the volume change of electrode materials during the cycling process. In Figure 3g, the EDS mapping of CuCo₂O₄@NiFe-LDH was presented; the elements of Cu, Co, O, Ni, and Fe in the core–shell structure are uniformly distributed, which proves the successful synthesis of CuCo₂O₄@NiFe-LDH.

The chemical states of CuCo₂O₄@NiFe-LDH were systematically probed by X-ray Photoelectron Spectroscopy (XPS). The survey spectrum (Figure 4a) confirms the coexistence of Cu, Co, O, Ni, and Fe. High-resolution Cu 2p analysis (Figure 4b) shows the Cu²⁺ 2p_3/2 (933.15 eV) and 2p_1/2 (953.49 eV) spin–orbit doublet ($\mathsf{\Delta }=20.34 \mathrm{eV}$) with satellite peaks at 942.3/962.2 eV, confirming the presence of Cu 2p_3/2 and Cu 2p_1/2. This multivalent property of copper facilitates the promotion of reversible redox reactions and electronic conductivity in the spinel structure. As shown in Figure 4c, the peaks situated within the ranges of 793–798 eV and 778–785 eV are assigned to Co 2p_1/2 and Co 2p_3/2, respectively. Moreover, the peaks positioned at 785.67 eV and 802.36 eV are identified as satellite peaks (labelled “Sat.”). Additionally, the peaks at 782.87 eV (Co 2p_3/2) and 796.4 eV (Co 2p_1/2) can be attributed to the presence of Co²⁺. On the other hand, the peaks at 778.5 eV (Co 2p_3/2) and 793.5 eV (Co 2p_1/2) strongly suggest the existence of Co³⁺, thereby indicating the concurrent presence of both Co³⁺ and Co²⁺ in the sample [34]. This mixed-valence property is crucial as it facilitates electron hopping within the spinel lattice and at the interface with the LDH layer. This electronic conductivity greatly facilitates charge transfer during charge/discharge cycling. In addition, the core-level spectrum of O 1s shows three main peaks (Figure 4d), with peaks at 529.9, 531.2, and 532.7 eV attributed to metal bonds (Cu/Co-O), oxygen vacancies, and adsorbed H₂O molecules [35]. The Ni 2p_2/3 (855.6 eV) and 2p_1/2 (873.2 eV) peaks with a splitting energy of 17.6 eV confirm the Ni²⁺/Ni³⁺ redox pair, supported by satellite peaks at 861.5 eV and 879.5 eV [36]. Furthermore, the XPS spectra of Fe 2p feature the peak at a binding energy of 712.8 eV, corresponding to Fe 2p_3/2, and the peak at 726.2 eV, corresponding to Fe 2p_1/2 (Figure 4f). These results collectively demonstrate that Ni and Fe in the LDH structure exist as Ni²⁺/Ni³⁺ and Fe³⁺, respectively, consistent with previous studies [37].

The microstructure of CuCo₂O₄@NiFe-LDH was further characterized by transmission electron microscopy (TEM). The TEM image of CuCo₂O₄@NiFe-LDH is shown in Figure 5a, and a clear hierarchical core–shell structure can be seen. Individual CuCo₂O₄ nanoneedles exhibit a width of ~60 nm, while the NiFe-LDH nanosheets possess lengths ranging from 50 to 80 nm. The interwoven structure of NiFe-LDH nanosheets broadens the contact area with the electrolyte, and the anisotropy of CuCo₂O₄ nanoneedles creates an ordered reaction path for rapid electron/ion transport. This cooperative structure advantage significantly improves the charge conduction efficiency of the material. In Figure 5b, the 0.25 nm crystal plane spacing corresponds to the (012) crystal plane of NiFe-LDH, while 0.28 nm corresponds to the (311) crystal plane of CuCo₂O₄. This further demonstrates the accomplished synthesis of CuCo₂O₄@NiFe-LDH. The polycrystalline nature of CuCo₂O₄@NiFe-LDH is clearly demonstrated by the selected electron diffraction map presented in Figure 5c. The elemental mapping test results in Figure 5d show that Cu, Co, O, Ni, and Fe are uniformly distributed in CuCo₂O₄@NiFe-LDH. This finding is in accordance with the results obtained from XRD and XPS analyses.

The electrochemical performance of CuCo₂O₄@NiFe-LDH was evaluated using a three-electrode system in a 2 mol L⁻¹ KOH electrolyte. The working electrode was CuCo₂O₄@NiFe-LDH, with a Pt plate as the counter electrode and a saturated calomel electrode as the reference electrode. Figure 6a presents the cyclic voltammetry (CV) comparison curves of the three electrodes at 10 mV s⁻¹. It reveals that, CuCo₂O₄@NiFe-LDH has a larger CV integration area and redox peaks compared to CuCo₂O₄ and NiFe-LDH at the same current density. This outcome is attributable to the distinctive core–shell architecture of the CuCo₂O₄@NiFe-LDH electrode material. This structure maximizes the surface area-to-volume ratio, creating more exposed surfaces for ion interaction and minimizing the ion migration length. At the same time, both CuCo₂O₄@NiFe-LDH and CuCo₂O₄ have a pair of distinct redox peaks. This results from the redox reactions of copper, cobalt, nickel, and iron ions, suggesting that all of them exhibit excellent pseudocapacitive responses. Moreover, the redox peak peaks of CuCo₂O₄@NiFe-LDH and CuCo₂O₄ are not much different from each other. This indicates that both materials have excellent reversibility in redox reactions. The galvanostatic charge–discharge (GCD) curves of CuCo₂O₄, NiFe-LDH, and CuCo₂O₄@NiFe-LDH at 1 A g⁻¹ are shown in Figure 6b. It can be seen that CuCo₂O₄@NiFe-LDH has a longer discharge time than other electrode materials at the same current density. This observation suggests that the core–shell structure effectively increases the availability of active sites in the electrode material. In short, advanced structural design bridges the gap between materials science and electrochemistry. The electrochemical performance of the device is improved. These findings are consistent with previous CV analyses. To quantitatively analyze the impedance behavior, an equivalent circuit model (Figure S2) was employed. The electrochemical impedance comparison of the three electrodes is shown in Figure 6c. The Nyquist plot reveals that CuCo₂O₄@NiFe-LDH exhibits the lowest internal resistance ($Rs=0.703 \mathsf{\Omega }$) in the high-frequency region. Meanwhile, the charge transfer resistance (R_ct) of CuCo₂O₄@NiFe-LDH (0.12 Ω) is smaller than that of CuCo₂O₄ (0.18 Ω) and NiFe-LDH (0.27 Ω). This result indicates that the introduction of NiFe-LDH effectively lowers interfacial resistance and improves the charge transfer rate. Furthermore, the steeper slope of CuCo₂O₄@NiFe-LDH in the low-frequency region suggests enhanced ion diffusion dynamics within the electrode. The composite material demonstrates enhanced ionic diffusion capability, as evidenced by its steeper low-frequency slope in impedance spectra. As shown in Figure 6d, the CV curves of CuCo₂O₄@NiFe-LDH retain distinct redox peaks at varying scan rates. A slight shift of the peak can be observed due to progressive polarization and internal diffusion resistance, but the overall shape of the curve is essentially unchanged. According to the CV curve, the following redox reactions occur at this electrode [38]:

$${\mathrm{CuCo}}_2{\mathrm{O}}_4{+\mathrm{OH}}^{-}{+\mathrm{H}}_2\mathrm{O} \leftrightarrow { \mathrm{CuOOH}+2\mathrm{CoOOH}+\mathrm{e}}^{-}$$

(1)

$${\mathrm{CuOOH}+\mathrm{OH}}^{-} \leftrightarrow { \mathrm{CuO}}_2{+\mathrm{H}}_2{\mathrm{O}+\mathrm{e}}^{-}$$

(2)

$${\mathrm{CoOOH}+\mathrm{OH}}^{-} \leftrightarrow { \mathrm{CoO}}_2{+\mathrm{H}}_2{\mathrm{O}+\mathrm{e}}^{-}$$

(3)

$$\mathrm{Ni}{\left(\mathrm{OH}\right)}_2{+\mathrm{OH}}^{-} \leftrightarrow { \mathrm{NiOOH}+\mathrm{H}}_2{\mathrm{O}+\mathrm{e}}^{-}$$

(4)

$$\mathrm{Fe}{\left(\mathrm{OH}\right)}_{2 }{+\mathrm{OH}}^{-} \leftrightarrow { \mathrm{FeOOH}+\mathrm{H}}_2{\mathrm{O}+\mathrm{e}}^{- }$$

(5)

The GCD curves of CuCo₂O₄@NiFe-LDH are presented in Figure 6e. The specific capacities are 323.4, 301.1, 293.3, 280.6, 271.3, and 260.6 mAh g⁻¹ at 1, 2, 3, 5, 8, and 10 A g⁻¹, respectively. The retention rate of 10 A g⁻¹ is 80.7% compared with 1 A g⁻¹, which provides excellent rate capability performance. The outstanding rate capability performance and high specific capacity may be attributed to the good synergy between CuCo₂O₄ and NiFe-LDH. This heterostructure design simultaneously boosts the specific capacity of CuCo₂O₄ and suppresses NiFe-LDH agglomeration. At the same time, it can be observed that each curve has an obvious charge and discharge platform. It can be proved that the redox reaction of the electrode material is highly reversible. Figure 6f presents the long-term cycling performance of the CuCo₂O₄@NiFe-LDH electrode at a high current density of 10 A g⁻¹. Remarkably, the electrode retains 90.66% of its initial capacity after 10,000 cycles, demonstrating exceptional cycling stability and structural robustness.

To further analyze the charge storage behavior of CuCo₂O₄@NiFe-LDH, CV tests were performed at low scan rates of 0.4, 0.6, 0.8, 1, and 1.2 mV s⁻¹. The value of b was obtained by fitting the peak data and the scan rate from Figure 7a. To further understand the proportion of contribution of each of the diffusion control process and the capacity control process, the formula is calculated as follows [39,40,41]:

$${\mathrm{i}=\mathrm{a} \times \mathrm{v}}^{\mathrm{b}}$$

(6)

$${\mathrm{i}\left(\mathrm{v}\right)=\mathrm{k}}_1{ \times \mathrm{v}+\mathrm{k}}_2{ \times \mathrm{v}}^{1/2}$$

(7)

Calculations reveal b-values of 0.72 and 0.75 (both within the range of 0.5–1), indicating that the CuCo₂O₄@NiFe-LDH electrode exhibits a hybrid energy storage mechanism combining battery-type and capacitive behavior. To elucidate the contribution of pseudocapacitance, the capacitive and ionic diffusion processes were quantitatively analyzed using Equations (6) and (7). As can be seen from Figure 7c, the contribution of surface control is 65%, 70.4%, 72.2%, 73.5%, and 76.59% at 0.4, 0.6, 0.8, 1, and 1.2 mV s⁻¹, respectively. This may be related to the core–shell structure, where the larger specific surface area of CuCo₂O₄@NiFe-LDH provides shorter ion diffusion channels for ion diffusion. In addition, the percentage contribution of the surface capacity rises as the scan rate increases [42]. This can be attributed to the shortened diffusion time of electrolyte ions from the electrolyte to the lattice at high scan rates.

We assembled a two-electrode system with CuCo₂O₄@NiFe-LDH as the positive electrode and RGO as the negative electrode. The CV curves of the RGO electrodes are shown in Figure 8a, and the quasi-rectangular shape can be clearly seen, indicating the ideal capacity of RGO. The GCD curves of RGO at 1–10 A g⁻¹ are almost symmetrical (Figure 8b). According to the GCD curves, the specific capacities of RGO are 56.4, 54.2, 52.6, 50.4, 49.0, and 47.1 mAh g⁻¹ at different current densities. Figure 8c shows the CV curves of RGO and CuCo₂O₄@NiFe-LDH. Among them, the voltage window of −1–0 V is contributed by RGO, and the potential window of 0–0.6 V is contributed by CuCo₂O₄@NiFe-LDH.

In order to further evaluate the practicality of CuCo₂O₄@NiFe-LDH, the CuCo₂O₄@NiFe-LDH//RGO aqueous ASC device was assembled (Figure 9a). As shown in Figure 9b, CV analysis across varying voltage windows revealed stable operation within 0–1.6 V. However, pronounced polarization was observed beyond this range. Consequently, 1.6 V was identified as the optimal operational limit. Remarkably, the CV profiles (Figure 9c) retained their quasi-rectangular shape even at high scan rates (up to 100 mV s⁻¹), reflecting rapid charge transfer kinetics and efficient ion diffusion. In addition, the enclosed area enclosed by the CV curve increases steadily with increasing sweep speed, indicating excellent rate performance and a good charge distribution on the electrodes. Figure 9d illustrates the GCD curves of the ASC device. The nonlinear symmetry demonstrated by the curves suggests that the electrode materials possess excellent reversibility. According to Equations (S2), (S5), and (S6), the specific capacity of the device declines from 78.5 mAh g⁻¹ to 66.51 mAh g⁻¹ as the current density rises from 1 to 10 A g⁻¹. Possessing a rate performance of 85.1% indicates that the device possesses a fast redox reaction and excellent kinetic equilibrium. As shown in Figure 9e, the Ragone plot depicts the relationship between energy density and power density. At a power density of 856.3 W kg⁻¹, the energy density is 68.7 Wh kg⁻¹. Even with the power density being boosted to 9213.6 W kg⁻¹, the energy density remains at 41.2 Wh kg⁻¹. The tested results far exceed those previously reported for the SCs [43,44,45,46]. The cycling performance of the ASC device at 10 A g⁻¹ is presented in Figure 9f. It can be observed that 89.36% cycling efficiency is achieved even at 10,000 cycles. The gradual weakening of the capacity may be due to the collapse of the interlocking nanosheets and the decrease in porosity after 10,000 cycles.

4. Conclusions

In this work, ultrathin NiFe-LDH nanosheets were vertically inserted on urchin-like CuCo₂O₄ nanoneedles by the hydrothermal method. Among them, the ultrathin NiFe-LDH nanosheets form a shell layer on CuCo₂O₄, providing abundant open spaces and active surface sites that facilitate rapid electrolyte penetration and ion diffusion. This core–shell structure not only significantly enhances the specific surface area of the material but also effectively mitigates the stress effects experienced by the electrode during charge/discharge processes. In three-electrode testing, the CuCo₂O₄@NiFe-LDH electrode exhibited a specific capacity of 323.4 mAh g⁻¹ at a current density of 1 A g⁻¹. At a current density of 10 A g⁻¹, the capacity was 260.6 mAh g⁻¹ with a retention of 80.6%, which showed good rate performance. The specific capacity still maintained 90.6% after 10,000 cycles (at a current density of 10 A g⁻¹), demonstrating excellent electrochemical performance. Kinetic analysis revealed that the energy storage mechanism involves a combination of capacitive and battery-type contributions. At 0.4 mV s⁻¹, surface-controlled processes accounted for 65% of the total capacity, with 35% attributed to diffusion-controlled processes. At 1.2 mV s⁻¹, the surface contribution increased to 76.5%, while the diffusion contribution decreased to 23.5%. The assembled CuCo₂O₄@NiFe-LDH//RGO aqueous ASC device exhibited a max energy density of 68.7 Wh kg⁻¹ at 856.3 W kg⁻¹. After 10,000 cycles at 10 A g⁻¹, its capacity retention was 89.36%.