Rational core–shell design of novel Ni–Co bimetallic MOF@PBA and Al_xV₂O₅@C as positive and negative electrode materials for high-performance aqueous Zn-ion hybrid supercapacitors

Abstract

Aqueous Zn-ion hybrid supercapacitors (ZIHSCs) represent a viable energy-storage technology for power-oriented applications, combining long cycle life with intrinsic safety, low cost, and environmental compatibility. However, their practical development is often hindered by kinetic mismatch between capacitive and battery-type electrodes in conventional asymmetric configurations. In this work, we demonstrate the feasibility of a core–shell electrode design strategy that enables improved kinetic compatibility by integrating complementary Faradaic and capacitive charge-storage mechanisms within each electrode of a ZIHSC. The positive electrode (NCBMOF@PBA) is designed with a new pseudocapacitive nickel, cobalt bimetallic metal–organic framework (NCBMOF) core and a Faradaic Prussian blue analogue (PBA) shell, combining fast redox reactions with stable Zn²⁺-storage behaviour. The negative electrode employs a Faradaic Al-doped vanadate (Al_xV₂O₅/AlVO) core encapsulated by a capacitive carbon shell (AlVO@C), which enhances electronic conductivity, buffers volume changes and facilitates charge-transfer processes. This complementary core–shell configuration serves as a model system to mitigate electrode-level kinetic imbalance and improve active-material utilization. An aqueous ZIHSC assembled with NCBMOF@PBA//AlVO@C electrodes and a 1 M ZnSO₄-silica gel electrolyte exhibits stable charge-storage behaviour within a 0–1.8 V operating window. Representative electrochemical performance includes a capacitance of 450 F g⁻¹ at 1 A g⁻¹, an energy density of 203 Wh kg⁻¹ at 900 W kg⁻¹, and a power density of 9000 W kg⁻¹ at 158 Wh kg⁻¹. These results validate the proposed core–shell design concept and highlight its potential as a generalizable electrode-engineering approach for advancing aqueous ZIHSCs.

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1. Introduction

The growing environmental crisis associated with fossil-fuel consumption has accelerated the adoption of renewable energy sources such as solar and wind.^1,2 Despite their sustainability advantages, the intermittent and unpredictable nature of these resources, governed by diurnal cycles and weather fluctuations, necessitates the development of efficient energy-storage systems to maintain grid reliability and continuous power delivery.³ Although lithium-ion and lead–acid batteries are currently employed for renewable energy storage, their widespread implementation is restricted by high cost and safety concerns in lithium-ion systems, along with the limited energy and power densities, poor cycle life, and environmental issues associated with lead–acid batteries.⁴ In this scenario, supercapacitors have emerged as attractive energy-storage devices due to their high power capability, fast charge–discharge kinetics, and exceptional cycling stability.^5,6 Among them, aqueous zinc-ion hybrid supercapacitors (ZIHSCs) are particularly promising, as they synergistically combine the high energy storage of Faradaic battery-type electrodes with the rapid power delivery of capacitive materials.⁷ Furthermore, the utilization of earth-abundant, low-cost zinc and environmentally benign aqueous electrolytes imparts inherent safety, sustainability, and economic viability to ZIHSCs, highlighting their strong potential for large-scale renewable energy storage applications.

In aqueous ZIHSCs, the Faradaic charge-storage process is predominantly governed by metallic Zn-based negative electrodes, which are widely adopted owing to their high theoretical energy density, low redox potential, suppressed hydrogen evolution during Zn plating/stripping, and advantages of low cost, non-toxicity, and environmental benignity.⁸ Despite these merits, the practical application of metallic Zn electrodes is severely constrained by several intrinsic challenges arising from repeated Zn deposition and dissolution. These include non-uniform Zn plating leading to dendritic growth, surface passivation caused by by-product formation, electrode shape deformation, and zincate crossover, all of which collectively deteriorate electrochemical reversibility, safety, and cycling stability.⁸ To address these issues, extensive efforts have been devoted to engineering Zn metal electrodes through morphology regulation, surface coatings, rational current collector and separator design, and electrolyte optimization.⁹ Nevertheless, achieving long-term stable and highly reversible Zn metal electrodes remains a formidable challenge due to the fundamental instability of the stripping/plating mechanism in aqueous media. In this context, intercalation-type Faradaic or pseudo-Faradaic electrode materials have emerged as promising alternatives to metallic Zn electrodes, as they enable charge storage via reversible ion insertion/extraction or surface redox reactions, thereby avoiding dendrite formation and offering improved electrochemical kinetics, structural stability, and cycling durability. In addition to Faradaic electrodes, engineered carbon-based non-Faradaic materials that store charge through reversible ion adsorption/desorption at the electrode–electrolyte interface can also be employed as either the negative or positive electrode in aqueous ZIHSCs.⁹ Such capacitive electrodes offer rapid kinetics, excellent cycling stability, and structural robustness, making them attractive for high-power applications. On the positive electrode side, both Faradaic and pseudo-Faradaic redox-active materials are commonly explored to enhance energy density while maintaining fast charge–discharge characteristics.⁷ However, the rational design of intercalation-type Faradaic materials for ZIHSCs requires careful consideration of the unique physicochemical characteristics of the Zn²⁺ ion.¹⁰ As a multivalent charge carrier, Zn²⁺ experiences strong electrostatic interactions with the anionic framework of the host lattice, which significantly hinders solid-state diffusion and slows reaction kinetics. Moreover, the strong solvation of Zn²⁺ in aqueous electrolytes results in a large hydration shell, leading to a high de-solvation energy barrier prior to insertion into the host structure. To circumvent this penalty, solvent co-insertion often occurs during the charge-storage process, which can induce lattice distortion, structural degradation, and eventual capacity fading. Consequently, developing stable and efficient Zn²⁺-hosting intercalation materials remains intrinsically challenging, necessitating precise control over the crystal structure, interlayer spacing, electronic conductivity, and electrolyte–electrode interactions to enable reversible Zn²⁺ storage with long-term stability.

Conventional ZIHSC configurations typically employ an asymmetric electrode design, in which one electrode exhibits capacitive or pseudocapacitive behaviour while the counter electrode operates via a battery-type Faradaic mechanism, or vice versa.⁷ Although this strategy enables partial decoupling of energy and power contributions, it often results in kinetic mismatch between the two electrodes, leading to sub-optimal utilization of active materials and unbalanced electrochemical performance. In this work, we address this limitation by engineering both the positive and negative electrodes with rationally designed core–shell architectures that simultaneously integrate Faradaic and capacitive charge-storage mechanisms. For the positive electrode, a new pseudocapacitive nickel and cobalt bimetallic metal–organic framework (NCBMOF) is employed as the core to ensure fast redox kinetics, while a Prussian blue analogue (PBA) shell provides Faradaic Zn²⁺ storage with higher capacity and structural stability. Conversely, the negative electrode is designed with a Faradaic aluminium-doped vanadate (Al_xV₂O₅·nH₂O, AlVO) core to deliver reversible Zn²⁺ intercalation, encapsulated by a purely capacitive carbon shell that enhances electronic conductivity, buffers volume changes, and accelerates charge transfer. This complementary core–shell electrode design enables synergistic coupling of energy-dominant Faradaic processes with power-dominant capacitive contributions within each electrode, thereby minimizing kinetic imbalance and improving charge utilization. As a result, the two electrodes exhibit well-matched electrochemical behaviour, enabling a balanced ZIHSC system with enhanced energy density, power capability, and cycling stability. Consequently, the device outperforms previously reported systems, as summarized in Table S1 (SI).

2. Experimental methods

2.1. Materials

Nickel(II) nitrate hexahydrate (Ni(NO₃)₂·6H₂O, 98%, Loba Chemie Pvt. Ltd), cobalt(II) chloride hexahydrate (CoCl₂·6H₂O, 98.0%, Merck KGaA), benzene-1,3,5-tricarboxylic acid (BTC, 95%, Sigma-Aldrich), cetyltrimethylammonium bromide (CTAB, 98%, SRL Pvt. Ltd), potassium ferricyanide (K₃[Fe(CN)₆], 99%, SRL Pvt. Ltd), vanadium pentoxide (V₂O₅, Merck KGaA), aluminium sulfate hexadecahydrate (Al₂(SO₄)₃·6H₂O, Merck KGaA), zinc sulfate heptahydrate (ZnSO₄·7H₂O, Merck KGaA), polyvinyl alcohol (PVA, Merck KGaA), hydrogen peroxide (H₂O₂ (30%), Merck KGaA), N,N-dimethylformamide (DMF, 99%, Merck KGaA) and ethanol (EtOH, 99.9%, CSS Reagent) were used without any further purification. Deionized water (DI H₂O) was used throughout the experiments.

2.2. Synthesis of active materials

2.2.1 Synthesis of nickel–obalt bimetallic MOFs (NCBMOF-XY, where X : Y = Ni : Co; XY = 11, 12 & 21). NCBMOF samples were synthesized by means of a solvothermal method (Scheme 1). For the synthesis of NCBMOF-11, a mixture of Ni(NO₃)₂·6H₂O (2 mmol), CoCl₂·6H₂O (2 mmol), BTC ligand (4 mmol) and CTAB (0.02 g) was added in a solvent mixture containing DMF, EtOH and H₂O in a 10 : 3 : 2 (v/v) ratio followed by constant stirring at room temperature until a homogeneous solution was formed. The resulting solution was then transferred into a 25 mL Teflon-lined autoclave and kept at 140 °C for 16 h. After cooling to room temperature, the resulting crystalline products were filtered and washed with EtOH several times. Finally, the products were dried at 60 °C. In a similar process, NCBMOF-12 and NCBMOF-21 were prepared with their respective Ni : Co molar ratios. The total metal to ligand molar ratio was maintained as 1 : 1.

Scheme 1 Schematic representation of the stepwise synthesis of NCBMOF@PBA and AlVO@C core–shell composite materials.

2.2.2 Synthesis of NCBMOF@PBA-XY core–shell composites. 0.2 g of each of the three pre-obtained NCBMOF samples was homogeneously dispersed in 2.5 mL H₂O (Solution A). Then a 2.5 mL aqueous solution of 1 M K₃[Fe(CN)₆] (Solution B) was quickly injected in Solution A and stirred for 2 h at room temperature. The resulting products were centrifuged and washed repeatedly with H₂O and EtOH to remove unreacted species, followed by drying in an oven at 60 °C overnight. The final products were designated as NCBMOF@PBA-11, NCBMOF@PBA-12 and NCBMOF@PBA-21 accordingly. Scheme 1 depicts the synthetic route for NCBMOF@PBAs (NCBMOF@PBA-XY core–shell composites).

2.2.3 Synthesis of AlVO and AlVO@C. To synthesize AlVO, 2.1 g of V₂O₅ was dissolved in 400 mL of deionized water with 50 mL of H₂O₂ (30%). Separately, 1 g of Al₂(SO₄)₃·16H₂O was dissolved in 40 mL of deionized water. Both solutions were then combined in a 1000 mL beaker and left undisturbed at room temperature for 15 days.¹¹ A dark brown precipitate formed, which was collected by centrifugation and washed three times with water. The washed precipitate was then dried overnight in a vacuum oven at 60 °C. Carbon-coated AlVO was synthesized by thoroughly grinding AlVO with polyvinyl alcohol (PVA) in a 1 : 1 mass ratio, followed by calcination at 650 °C for 3 h under Ar flow (50 mL min⁻¹) at a heating rate of 5 °C min⁻¹ (see Scheme 1). The resulting dark black carbon-coated AlVO powder was ground and subsequently used for characterization.

2.3. Physicochemical characterization methods

Suitable crystals of NCBMOF-11 were identified under a polarising optical microscope for single-crystal X-ray diffraction (SCXRD) studies. The data were collected at room temperature using a Rigaku Oxford Diffraction, XtaLAB Synergy-S diffractometer. The X-ray was generated at 50 kV and 0.8 mA by using MoKα (λ = 0.71073 Å) radiation (PhotonJet Mo, Ceramic Fine Focus). The cell refinement and data reduction for the SCXRD studies were carried out with CrysAlis^pro.¹² The Olex2 program package, which includes programs like SHELX, PLATON etc., was employed to solve the structure and structural refinement.^13–15 The position of the hydrogen atoms can be assigned from the difference Fourier maps, and during the structural refinement, the hydrogens were fixed at geometrically ideal positions and refined using the riding mode. The final refinement was carried out by including the atomic positions of all the atoms and keeping anisotropic thermal parameters for all the non-hydrogen atoms, and isotropic thermal parameters for all the hydrogen atoms. The details of cell parameters and refinement parameters are listed in Table S2 (SI). The powder X-ray diffraction (PXRD) analysis was performed on a Bruker D8 advance X-ray diffractometer (XRD) using Cu Kα radiation (λ = 1.5406 Å) at a scan speed of 0.5° min⁻¹ and step size of 0.02°. Morphological analysis was performed using a field emission scanning electron microscope (FESEM, Zeiss-Sigma) attached with an energy-dispersive X-ray (EDX) spectrometer and high annular angle-field scanning transmission electron microscope (HAADF-STEM, Thermo Scientific, Talos F200X G2). N₂ adsorption–desorption studies were carried out using a surface area and pore size analyzer (Quantachrome Autosorb iQ2 analyzer) in a liquid N₂ atmosphere (77.35 K) and the degassing temperature was 100 °C for 4 h. Fourier transform infrared (FTIR) spectra were obtained using an IR spectrophotometer (Shimadzu Corp., 01172) using the attenuated total reflectance (ATR) mode in the spectral range 400–4000 cm⁻¹. Raman spectra were obtained on a Raman spectrometer (Renishaw, Invia II) having a laser wavelength of 532 nm. Surface compositional study was carried out using X-ray photoelectron spectroscopy (XPS) (PHI, VersaProbe 4). Elemental analysis was carried out using inductively coupled plasma-optical emission spectroscopy (ICP-OES, Thermo Scientific, iCAP 7400 Duo).

2.4. Electrochemical characterization methods

Electrochemical testing of the active materials was performed in a three-electrode configuration using cyclic voltammetry (CV), electrochemical impedance spectroscopy (EIS), and galvanostatic charge/discharge (GCD) in 1 M ZnSO₄ (aq.) electrolyte solution. Stainless steel (SS304) foil was used as a low-cost current collector; however, its susceptibility to chemical and electrochemical corrosion necessitated surface modification. A thin single-wall carbon nanotube (SWCNT) coating was applied to SS304, enabling it to function as a stable and conductive current collector.¹⁶ Working electrodes were prepared by brush-coating a slurry containing 80 wt% active material, 10 wt% carbon black, and 10 wt% polyvinylidene fluoride (PVDF) in dimethylformamide (DMF) solvent onto the SWCNT-coated SS304 substrate (1 × 1 cm², ∼5 mg cm⁻² mass loading). The coated electrodes were dried at 60 °C and subsequently pressed at 15 tons. Platinum foil (2 × 2 cm²) and Ag/AgCl served as the counter and reference electrodes, respectively. Full cells (electrode area: 5 × 5 cm²) were assembled using a 1 M ZnSO₄-silica gel electrolyte prepared by dispersing 6 wt% fumed silica (CAB-O-SIL® M-5) in the 1 M ZnSO₄ solution. A piece of polyvinyl chloride (PVC) mosquito netting was used as the separator. The full cells were positive-electrode-limited, maintaining a 1 : 2 mass ratio between the positive (∼5 mg cm⁻²) and negative (∼10 mg cm⁻²) electrodes. All measurements were carried out at room temperature using a multichannel AUTOLAB M204 potentiostat/galvanostat. Unless otherwise stated, all experiments were carried out at ambient temperature (∼25 °C).

3. Results and discussion

3.1. Physicochemical characterization studies of active materials

Three new bimetallic NCBMOFs were stabilized in three different compositions with a general formula of [Ni_xCo_y(BTC³⁻)₂(HCOO⁻)₆(DMF)₆] [where x = 3 and y = 3 for NCBMOF-11, x = 4 and y = 2 for NCBMOF-21, and x = 2 and y = 4 for NCBMOF-12]. As all three have identical structures, here we are describing the structure of NCBMOF-11. As suggested by SCXRD, NCBMOF-11 crystallizes in the trigonal P-3 (no. 147) space group. It contains 14 non-hydrogen atoms in its asymmetric unit (Fig. S1), which originates from one metal atom (Co^II/Ni^II), a third of a BTC³⁻ linker, one formate anion, and one DMF molecule. The metal atom is connected to six oxygen atoms to form a distorted octahedral geometry. Three out of six oxygen atoms come from three formate ions, and the remaining three come from two BTC carboxylates and one DMF solvent molecule (Fig. 1, S1 and S2). Here, the formate and BTC molecules act as connecting linkers, whereas the DMF acts as a terminal. In this case, the formate ion, generated in situ inside the autoclave at high temperatures and pressures, acts as a coligand. It is known that DMF solvent, in the presence of water/acid/base, and at high temperatures, can hydrolyse to form formate ions.¹⁷ The M–O bond distance varies in the range from 2.027 to 2.123 Å (Table S3). The M–O–M bond angles vary from 114.36° to 114.4° (Table S4).

Fig. 1 (a) Packing of 2D layers of NCBMOF-11 along the crystallographic b-axis. (b) PXRD plots of NCBMOF-11, NCBMOF-12 and NCBMOF-21. (c) Rietveld refinement of NCBMOF-12. (d) PXRD plots of NCBMOF@PBA-11, NCBMOF@PBA-12 and NCBMOF@PBA-21. Rietveld refinement of (e) NCBMOF@PBA-21 and (f) AlVO. The inset of (f) depicts the representative crystal structure of AlVO.

Six such octahedral metal atoms, bridged via six formate ions and six BTC linkers, form a hexagonal cluster-like arrangement. Here, the formate ion connects two adjacent metal atoms present within the hexagonal cluster via µ₃-η¹:η² mode of bridging, whereas the six carboxylates from six different BTC linkers connect one hexagonal cluster to six other nearest hexagonal metal cluster units via µ₂-η¹:η¹ mode of bridging from carboxylates. Each BTC linker connects three such hexagonal units via its three carboxylates to form a two-dimensional layer. The layers are arranged one above the other through AAA… packing. Several weak non-covalent interactions (such as π⋯π, C–H⋯π, and hydrogen bonding) are identified between the layers shown in Fig. 1a. Between the two layers, the bonded DMF molecules are present in a crisscross fashion. No substantial porous channels are found in the structure. However, along the abc axes, small porous channels of 7.80 and 9.99 Å can be observed.

Our literature search reveals that a few MOFs have been reported earlier with the similar cell parameters and structure (see Table S5). However, to the best of our knowledge, to date, there are no reports of bimetallic MOFs (Ni/Co) with the same cell parameters. A pure Ni-based MOF [Ni₆(BTC)₂(DMF)₆(HCOO)₆] has been reported earlier by Stock et al.¹⁸ Here, we have been able to stabilize the same crystal system with both Ni(II) and Co(II) metal centres in three different compositions. We present the crystal lattice parameters for these three bimetallic MOFs in Table S2. For NCBMOF-12 and NCBMOF-21, we could not find a perfect crystal to collect SCXRD data. However, the PXRD patterns of all three samples match with the simulated pattern obtained from the single crystal data of NCBMOF-11 (Fig. 1b). Therefore, for these MOFs, the cell parameters were extracted from the Le Bail refinement of the PXRD (Fig. 1c and S3). With the increment of Co percentage in the MOFs, a slight increase in the a and b cell axes values was observed. Elemental compositions (CHNS) of the NCBMOFs (Fig. S4) are in agreement with the empirical formula obtained from the SCXRD data. The Ni : Co molar ratios in the pristine NCBMOFs were determined by ICP-OES elemental analysis (Table S6). The Ni : Co molar ratios in NCBMOF-11, NCBMOF-12, and NCBMOF-21 were found to be 0.96, 0.54, and 2.03, respectively, which are in close agreement with the nominal values.

Upon PBA growth on the NCBMOFs, the PXRD patterns of NCBMOF@PBAs (Fig. 1d) differ markedly from those of the NCBMOFs (Fig. 1b), with varying degrees of crystallinity. The single phase of the NCBMOF@PBA-21 was determined by the refinement of the PXRD data (Fig. 1e and Table S7). The SEM (Fig. 2b), TEM (Fig. 2e and S16) and EDX (Fig. 2g–n and S11) analyses indicate that PBA shells have been formed on top of the NCBMOFs leading to a core–shell structure of NCBMOF@PBAs. The surface morphology and elemental composition of NCBMOF11, NCBMOF12, NCBMOF21, and NCBMOF@PBA-12 were analyzed using FESEM-EDX and HAADF-STEM EDX data (Fig. S6–S11) suggesting their uniformity and purity.

Fig. 2 (a)–(c) FESEM images and (d)–(f) TEM images of NCBMOF-12, NCBMOF@PBA-12 and AlVO@C, respectively; (g)–(n) SEM-EDX mapping of NCBMOF@PBA-12.

FTIR (Fig. 3a) and Raman (Fig. S17) analysis of the pristine MOFs suggest the formation of metal–BTC linkage whereas the FTIR spectra of the NCBMOF@PBAs (Fig. 3b) indicate the disappearance of the characteristic bands of NCBMOFs and appearance of a sharp absorption band corresponding to –C≡N group around 2094 cm⁻¹.^19,20 It supports the core–shell structure of the positive active materials.

Fig. 3 FTIR plots of (a) NCBMOFs and (b) NCBMOF@PBAs; (c) Raman plots of AlVO@C; deconvoluted XPS spectra of (d) Co 2p, (e) Ni 2p, and (f) O 1s of NCBMOF-12 and NCBMOF@PBA-12; deconvoluted XPS spectra of (g) Fe 2p of NCBMOF@PBA-12, (h) V 2p and (i) O 1s of AlVO@C; (j) comparative N₂ adsorption–desorption plots for NCBMOF-12 and NCBMOF@PBA-12.

The textural properties and porosity of NCBMOF-12 and NCBMOF@PBA-12 have been analyzed by means of N₂ adsorption/desorption isotherms (Fig. 3j and S18). NCBMOF-12 shows a type I adsorption/desorption isotherm (Fig. 3j) whereas NCBMOF@PBA-12 exhibits a combined I/IV type adsorption/desorption isotherm with strong N₂ adsorption at low relative pressure values and slightly steep adsorption at higher relative pressure values (0.8–1.0).²¹ It indicates the porous nature of both the materials. The Brunauer–Emmett–Teller (BET) model suggests an increase of almost six times in specific surface area upon PBA growth (521.1 m² g⁻¹) on bare NCBMOF-12 (89.7 m² g⁻¹) arising from the heterostructured core–shell formation. Such a high increase in the specific surface area for NCBMOF@PBA-12 is likely to provide a larger electrode/electrolyte interface for the electrochemical processes.

XPS analysis has been carried out on NCBMOF-12 as well as NCBMOF@PBA-12 to assess the oxidation state of the elements involved in these materials. The C 1s core level spectrum of the pristine MOF is deconvoluted to three components (Fig. S19) centred around 284.8 eV, 285.9 eV and 287.5 eV which correspond to aromatic carbon, –C–H and carboxylate carbon of the ligands, respectively. The O 1s region (Fig. 3f) also consists of three components with maxima around 529.6 eV, 532.6 eV and 534.6 eV which are assigned to M–O, M–O–M and surface oxygen species, respectively. On the other hand, the nickel and cobalt spectra (Fig. 3d and e) indicate the presence of divalent metal ions as expected from the empirical formula of the pristine MOF.^22,23

In case of NCBMOF@PBA-12, the C 1s spectrum (Fig. S19) also consisted of three species, adventitious carbon (284.8 eV), –CN (285.6 eV) and surface carbonates (287.2 eV). The O 1s spectrum (Fig. 3f) is resolved into two components arising from crystal water molecules (530.6 eV) and surface oxygen species (534.4 eV). Upon PBA growth on the pristine MOFs, the oxidation state of Ni remains unchanged, whereas Co changes from a divalent state to mixed divalent and trivalent states (Fig. 3d and e). Presence of +2 and +3 oxidation states of Fe in the Fe 2p deconvoluted spectrum (Fig. 3g) confirms the formation of PBA core–shell of the pristine MOFs.^24,25

The PXRD pattern of AlVO is refined based on the structure of hydrated magnesium decavanadate (Fig. 1f and Table S7) in which decavanadate anions are bound together with the hydrated metal cations.²⁶ AlVO is mixed with PVA and calcined in an Ar atmosphere to get the negative active material AlVO@C. The surface morphology and elemental composition of AlVO and AlVO@C were analyzed using FESEM-EDX and HAADF-STEM EDX data (Fig. S12–S16) suggesting their uniformity and purity. Due to carbonization of PVA on top of AlVO, the PXRD pattern (Fig. S5) of AlVO@C exhibits amorphous characteristics. The SEM (Fig. 2c), TEM (Fig. 2f and S16), and EDX (Fig. S13 and S15) analyses of the negative active material indicate the core–shell structure of AlVO@C (Fig. S13 and S15). The Raman spectrum of AlVO@C (Fig. 3c) shows D (1368 cm⁻¹) and G (1603 cm⁻¹) bands corresponding to lattice defects and graphitic carbon, respectively.²⁷ The XPS analysis of the V 2p spectrum (Fig. 3h) shows two peaks at 517.5 eV (2p_3/2) and 524.9 eV (2p_1/2) which indicates the presence of V⁵⁺ in the negative active material which can be electrochemically reduced to V⁴⁺ and V³⁺.²⁸ A significantly lower amount of V⁴⁺ is also detected from the deconvolution of the spectrum. The O 1s core level spectrum (Fig. 3i) is deconvoluted to the components corresponding to lattice oxygen (530 eV), intercalated H₂O molecules (532.5 eV) and surface oxygen species (534.6 eV).²⁸ All these XPS peaks are calibrated against the C 1s spectrum (Fig. S20).

3.2. Electrochemical characterization studies of active materials

3.2.1. Evaluation of charge storage mechanisms for core, shell and core–shell structures. In electrochemical energy-storage systems, charge storage typically arises from a combination of diffusion-controlled Faradaic reactions and surface-controlled capacitive or pseudo-Faradaic processes. To quantitatively distinguish these contributions, cyclic voltammetry (CV) at varying scan rates is employed, enabling analysis of how current responds to changes in sweep speed. The relationship between peak current (i) and scan rate (ν) follows the power-law expression.²⁹

i = a × ν^b (1)

where the exponent b provides mechanistic insight: values near 0.5 indicate diffusion-limited Faradaic behaviour, while values approaching 1 signify surface-dominated capacitive or pseudo-Faradaic processes. This approach allows clear identification of whether bulk ion diffusion or surface redox kinetics govern charge storage in the material. The charge-storage mechanisms of NCBMOFs, PBAs (synthetic route described in Section S1), NCBMOF@PBAs, AlVO, and AlVO@C were systematically investigated through CV at varying scan rates (5–25 mV s⁻¹) by using eqn (1) (Fig. 4 and S21–S24). All samples exhibit prominent redox features, confirming mixed contributions from both mechanisms. The MOF samples (NCBMOF-11, NCBMOF-12, and NCBMOF-21) exhibit a broad and merged redox peak pair, with the anodic feature appearing around 0.8–1.0 V and the corresponding cathodic peak at approximately 0.6–0.8 V (Fig. 4a and S21). These wide peaks arise from the overlapping redox contributions of both Ni²⁺/Ni³⁺ (minor) and Co²⁺/Co³⁺ (major) centres, reflecting their simultaneous and highly coupled electrochemical activity within the MOF framework.³⁰ From the mechanistic side, MOFs show higher b values (0.63–0.86), indicating that their charge storage is dominated by surface-driven pseudo-Faradaic reactions associated with the redox-active Ni/Co sites; these higher b values highlight the rapid kinetics and accessibility of surface redox centres. In contrast, the PBAs (PBA-11, 12, and 21) exhibit lower b values (0.21–0.47), much closer to the theoretical diffusion-controlled limit (b = 0.5), demonstrating that their Co/Fe redox activity involves deeper ion insertion and diffusion-regulated Faradaic reactions (Fig. S22). For the PBA samples, the Ni sites remain electrochemically inactive within the 0–1.2 V window, while the Co²⁺/Co³⁺ and Fe²⁺/Fe³⁺ redox couples dominate the charge-storage process.³¹ These coupled redox transitions produce anodic peaks in the 0.7–0.9 V range and corresponding cathodic peaks around 0.5–0.6 V. In the NCBMOF@PBA core–shell composites, the anodic and cathodic features appear as broad, overlapped peaks centred around 0.80–0.85 V and 0.65–0.70 V, respectively (Fig. 4b, S23(a and b)). Mainly Co redox centres from the MOF core and the Co/Fe centres from the PBA shell contribute to the overall response; however, the electrochemical activity is predominantly governed by the Co and Fe sites in the PBA shell, which dictate the composite's redox behaviour. Importantly, the core–shell NCBMOF@PBA composites present intermediate b values (0.53–0.68) between MOFs and PBAs, providing direct evidence of the synergistic interplay between the pseudo-Faradaic MOF core and the Faradaic PBA shell (Fig. S23). These intermediate b values signify a co-regulated mechanism, where the PBA shell promotes diffusion-governed Faradaic charge storage while the MOF core contributes to fast surface-controlled pseudo-Faradaic processes, collectively enhancing capacity and moderating the strongly overlapped redox responses of the individual components.

Fig. 4 Cyclic voltammograms of positive and negative electrodes at variable scan rates: (a) NCBMOF-12, (b) NCBMOF@PBA-12, (c) AlVO, and (d) AlVO@C.

CV analysis of the negative electrodes reveals two redox couples corresponding to the V⁵⁺/V⁴⁺ and V⁴⁺/V³⁺ transitions in both AlVO and AlVO@C. In the anodic sweep, the V³⁺ → V⁴⁺ and V⁴⁺ → V⁵⁺ peaks appear at approximately −0.2 V and 0.5 V, respectively, while the reverse V⁵⁺ → V⁴⁺ and V⁴⁺ → V³⁺ peaks emerge near −0.2 V and −0.5 V, respectively, in the cathodic sweep. Although both electrodes exhibit the same fundamental redox features, their kinetic behaviour differs significantly. Pristine AlVO shows pronounced peak shifting at higher scan rates, indicating sluggish charge-transfer kinetics and substantial mass-transfer limitations. In contrast, the core–shell AlVO@C displays minimal peak movement even at fast scan rates, reflecting more stable redox kinetics and reduced polarization. This discrepancy arises from the stabilizing effect of the carbon shell, which acts as a protective barrier that suppresses vanadium dissolution and preserves the structural integrity of the layered AlVO framework. In pristine AlVO, dissolution- and degradation-induced surface layers form readily, creating passivating barriers that impede Zn²⁺ transport and elevate mass-transfer resistance. The absence of such barriers in AlVO@C enables more efficient ion movement and maintains rapid redox responsiveness. Overall, the CV behaviour clearly highlights the superior rate capability and electrochemical durability of the core–shell AlVO@C electrode compared to its uncoated counterpart. From the charge storage mechanistic perspective, AlVO exhibits b values of 0.59–0.66, highlighting its predominantly diffusion-controlled Faradaic behaviour associated with reversible ion insertion. Upon carbon coating, AlVO@C displays increased b values (0.73–0.74), indicating a shift toward surface-controlled capacitive and pseudo-Faradaic contributions. This enhancement is attributed to the conductive carbon shell, which increases surface area, reduces charge-transfer resistance, and enables faster ion/electron transport, thereby significantly boosting rate capability.

However, the evolution of b values across the materials clearly demonstrates how rational core–shell engineering in both NCBMOF@PBA and AlVO@C finely tunes the balance between diffusion-dominated Faradaic and surface-dominated pseudo-Faradaic processes, resulting in optimized charge-storage performance crucial for high-efficiency hybrid supercapacitor systems.

3.2.2. Electrochemical performance testing of positive electrode materials. The GCD profiles of the NCBMOFs (Fig. 5a and S25), PBAs (Fig. S26), and NCBMOF@PBA composites (Fig. 5b and S27) reveal distinct charge-storage behaviours across the different materials. The inset of each GCD profile reports the corresponding Coulombic efficiency (CE) alongside the specific capacity (mAh g⁻¹), providing a clear assessment of charge–discharge reversibility. The MOFs (NCBMOF-11, NCBMOF-12, and NCBMOF-21) exhibit characteristically sloping charge–discharge curves, a hallmark of pseudocapacitive behaviour where surface redox reactions dominate over discrete voltage plateaus. Among the MOFs, NCBMOF-12 delivers the highest capacitance, ranging from 84 to 37 F g⁻¹ as the current density increases from 100 to 500 mA g⁻¹, whereas NCBMOF-11 (70–26 F g⁻¹) and NCBMOF-21 (25–8 F g⁻¹) show comparatively lower performance. Notably, NCBMOF-21 demonstrates the poorest electrochemical response, which correlates with its higher Ni content relative to Co. This compositional effect is chemically meaningful: oxidation of Ni²⁺ to Ni³⁺ in the positive electrode is energetically demanding due to the inherent stability of the Ni²⁺ (d⁸) electronic configuration, which resists further oxidation. In contrast, Co²⁺ (d⁷) is less stable and can be more readily oxidised to Co³⁺, enabling more favourable and reversible redox processes. Thus, a higher proportion of Co in the MOF framework enhances redox activity and boosts capacitive performance, while Ni-rich systems exhibit suppressed electrochemical response due to the less accessible Ni²⁺/Ni³⁺ redox couple. However, the MOF electrodes show capacitance losses of 63, 56, and 68% for NCBMOF-11, NCBMOF-12, and NCBMOF-21, respectively, as the current density increases from 100 to 500 mA g⁻¹. When the current density is returned to 100 mA g⁻¹ after 25 rate-capability steps, the electrodes recover 80, 87, and 76% of their initial capacitance, respectively, indicating partial reversibility but overall modest rate performance (Fig. S25c). For the same electronic-structure reasons discussed earlier, the Ni centres in the PBA systems remain electrochemically inactive within the positive-electrode potential window, leaving Co and Fe as the primary redox-active species. The charge–discharge profiles of the PBA systems display well-defined plateau regions, confirming their diffusion-controlled Faradaic charge-storage behaviour. The GCD profiles of PBA-11, PBA-12, and PBA-21 (Fig. S26a–c) and their rate-capability data (Fig. S26d) further support this behaviour. As expected, PBA-12 delivers the highest capacitance, ranging from 210 to 135 F g⁻¹ as the current density increases from 100 to 500 mA g⁻¹, outperforming PBA-11 (186–87 F g⁻¹) and PBA-21 (117–48 F g⁻¹). The rate-capability analysis shows capacitance losses of 53%, 36%, and 69% for PBA-11, PBA-12, and PBA-21, respectively, across the 100–500 mA g⁻¹ range, with subsequent capacitance recovery of 90%, 98%, and 87%, respectively, when the current density is returned to 100 mA g⁻¹ after 25 rate steps. The incorporation of the PBA shell introduces strong Faradaic charge-storage contributions, thereby significantly enhancing the capacitance of the NCBMOF@PBA composites. The charge–discharge profiles of the NCBMOF@PBA composites exhibit a hybrid “slanted-plateau” shape, reflecting a well-balanced combination of MOF-like pseudocapacitive behaviour and PBA-like diffusion-controlled Faradaic processes. Among the three composite core–shell materials, NCBMOF@PBA-12 delivers the highest performance, exhibiting 355–225 F g⁻¹ across 100–500 mA g⁻¹ (Fig. 5b), clearly surpassing NCBMOF@PBA-11 (188–97 F g⁻¹; Fig. S27a) and NCBMOF@PBA-21 (109–35 F g⁻¹; Fig. S27b). The rate-capability results (Fig. S27c) show capacitance losses of 48%, 35%, and 68% for NCBMOF@PBA-11, NCBMOF@PBA-12, and NCBMOF@PBA-21, respectively, over the 100–500 mA g⁻¹ range, with corresponding capacitance recoveries of 93%, 98%, and 84%, respectively, when the current density is returned to 100 mA g⁻¹ after 25 steps. In brief, integrating the PBA shell onto the MOF core markedly enhances both the capacitance and rate capability of the composites. This improvement arises from the synergistic coupling of the MOF's pseudocapacitive behaviour with the strong Faradaic activity of the PBA shell. Among all variants, the Ni : Co = 1 : 2 composition delivers the most optimized performance, as highlighted in Fig. 5c. Analysis of the CE values from the GCD profiles provides important insights into the reversibility of the charge–discharge processes. The results indicate that MOF electrodes exhibit significantly lower reversibility compared to PBAs and MOF@PBA composites, reflecting their limited structural stability and slower redox kinetics. Furthermore, among the different Ni/Co ratios studied in the core, shell, and core–shell systems, the Ni : Co = 2 : 1 variant consistently shows the lowest CE values, highlighting that a higher Ni content adversely affects both redox reversibility and overall electrochemical efficiency. The ex situ XPS analysis of NCBMOF@PBA-12 was performed at both charged and discharged states (see Fig. S28). The results indicate that the oxidation state of Ni remains unchanged, confirming the exclusive presence of Ni²⁺ without any detectable redox activity. In contrast, Co²⁺ and Fe²⁺ are oxidized to Co³⁺ and Fe³⁺ during the charged state, while a higher proportion of Co²⁺ and Fe²⁺ is restored in the discharged state. This behaviour demonstrates the electrochemical inertness of the Ni centre, which explains why a higher Ni content adversely affects both redox reversibility and the overall electrochemical efficiency.

Fig. 5 Performance testing of positive electrodes: GCD profiles for (a) NCBMOF-12 and (b) NCBMOF@PBA-12; comparative studies between NCBMOF-12 and NCBMOF@PBA-12: (c) rate capability, (d) Nyquist plots from EIS, (e) variation of D as a function of electrode potential at the second cycle, and (f) cycling performance at 500 mA g⁻¹ within 0–1.2 V vs. Ag/AgCl window.

Electrochemical impedance spectroscopy (EIS) was used to analyse the resistive and Zn²⁺ diffusion characteristics of NCBMOF-12 and NCBMOF@PBA-12, with the corresponding Nyquist plots and diffusion coefficients (D) shown in Fig. 5d and e. The inset in Fig. 5d illustrates the equivalent circuit comprising solution resistance (R_S), charge-transfer resistance (R_CT), a constant phase element (CPE), and the Warburg diffusion element (W). While both materials display nearly identical R_S values (∼6 Ω), NCBMOF@PBA-12 exhibits a markedly lower R_CT (∼2 Ω) with respect to NCBMOF-12 (∼10 Ω), confirming improved charge-transfer kinetics and enhanced electrochemical activity in the core–shell system. This reduction in R_CT directly supports the higher capacitance observed in GCD measurements. The D values of Zn²⁺ ions in the solid matrix were extracted from the low-frequency Warburg region using the formulations provided in Section S2, and the D-voltage profiles clearly show that NCBMOF@PBA-12 possesses consistently higher diffusion coefficients than NCBMOF-12 across different states of charge/discharge. The characteristic variations in D with voltage reflect structural and ion-migration changes occurring within the Zn²⁺ host framework during charge and discharge.

Fig. 5f presents the long-term cycling performance of NCBMOF-12 and NCBMOF@PBA-12 at 500 mA g⁻¹ within the 0–1.2 V window. The core–shell NCBMOF@PBA-12 shows good cycling stability, retaining ∼80% of its initial capacitance with ∼98% CE after 2000 cycles, whereas NCBMOF-12 suffers severe degradation, retaining only ∼6% with ∼90% CE. The Nyquist plots recorded before and after cycling (Fig. S29 and Table S8) further clarify this behaviour: NCBMOF-12 exhibits a drastic increase in R_CT (∼6 times), consistent with its near-complete capacitance loss, while NCBMOF@PBA-12 shows almost unchanged R_CT, reflecting its preserved electrochemical activity. The poor durability of NCBMOF-12 arises from the intrinsic structural instability of MOF frameworks under repeated redox cycling, which leads to gradual collapse and loss of active sites. In contrast, the robust PBA shell effectively stabilizes the MOF core, preventing structural degradation and retaining interfacial integrity. As a result, the core–shell architecture ensures long-term mechanical and electrochemical stability, enabling superior cycling performance.

3.2.3. Electrochemical performance testing of negative electrode materials. Fig. 6a and b presents the GCD profiles of AlVO and AlVO@C, with the corresponding capacity (mAg g⁻¹), capacitance (F g⁻¹) and CE (%) values at different current densities shown in the insets. Pristine AlVO delivers a remarkably high capacity of 200–103 mAh g⁻¹ (i.e., 720–371 F g⁻¹) within the wide potential window of 0.4 to −0.6 V when the current density increases from 100 to 500 mA g⁻¹, consistent with the two-electron redox process involving the V⁵⁺/V⁴⁺ and V⁴⁺/V³⁺ couples confirmed by CV. The ex situ XPS analysis of the V 2p spectra reveals that, in the charged state of AlVO, mixed valence states of V³⁺, V⁴⁺, and V⁵⁺ coexist, with V⁴⁺ being the dominant species. In contrast, during the discharged state, V⁵⁺ becomes the predominant component, accompanied by a relative decrease in V⁴⁺, while the V³⁺ state is no longer observed. These systematic changes in vanadium oxidation states (see Fig. S30) clearly elucidate the redox mechanism of the vanadate based negative electrode during electrochemical cycling. Despite its high intrinsic capacity, AlVO suffers from poor rate capability, showing a ∼48% capacitance loss from 100 to 500 mA g⁻¹ and recovering only ∼63% of its initial capacitance upon returning to 100 mA g⁻¹ (Fig. 6c). This degradation is attributed to vanadium dissolution and the collapse of the layered framework during repeated Zn²⁺ insertion and extraction.³² In contrast, the carbon-coated AlVO (AlVO@C) exhibits lower capacity/capacitance (82–64 mAh g⁻¹ & 296–230 F g⁻¹) but significantly improved rate performance and reversibility, losing only ∼22% capacitance across the same current range and retaining about 97% of its initial capacitance after 25 current-density steps (Fig. 6c). The carbon shell stabilizes the layered structure, suppresses V-dissolution, and enhances electronic and ionic pathways, which collectively account for the superior rate capability of AlVO@C. Although the carbon coating compromises some capacity, this trade-off is practically acceptable because most positive electrodes in aqueous Zn-ion systems do not reach capacities above 150 mAh g⁻¹, making the unutilized high capacity of pristine AlVO largely impractical in full-cell configurations.¹⁰ Therefore, the moderate capacity (∼80 mAh g⁻¹) combined with excellent rate retention and structural stability makes AlVO@C a far more viable and balanced negative electrode for high-performance aqueous Zn-ion hybrid supercapacitors.

Fig. 6 Performance testing of negative electrodes: GCD profiles for (a) AlVO and (b) AlVO@C; comparative studies between AlVO and AlVO@C: (c) rate capability, (d) Nyquist plots from EIS, (e) variation of D as a function of electrode potential at the second cycle, and (f) cycling performance at 500 mA g⁻¹ within 0.4 to −0.6 V vs. Ag/AgCl window.

EIS measurements for AlVO and AlVO@C (Fig. 6d) show that both electrodes possess similar R_S (∼6 Ω), indicating comparable electrolyte and contact contributions. However, a pronounced difference appears in the R_CT values: pristine AlVO exhibits a much higher R_CT (∼13 Ω) than AlVO@C (∼8 Ω). This elevated R_CT reflects sluggish interfacial charge transfer in AlVO, which is consistent with the formation of surface passivation layers arising from vanadium dissolution and structural degradation of the exposed layered framework. In contrast, the carbon shell in AlVO@C stabilizes the surface, suppresses vanadium dissolution, and provides improved electronic pathways, resulting in faster charge-transfer kinetics and demonstrating the clear advantage of the core–shell design. Although one might assume that the carbon shell could impede Zn²⁺ diffusion into the layered AlVO core, the EIS-derived Warburg diffusion profiles in Fig. 6e clearly show otherwise. Across the entire charge–discharge process, both AlVO and AlVO@C exhibit comparable diffusion coefficients, and no substantial decrease in D is observed for the carbon-coated sample at any of the states of charge and discharge. The similarity in the magnitude and trend of D values for both electrodes indicates that the carbon shell does not hinder Zn²⁺ transport; instead, it preserves the intrinsic diffusion characteristics of the AlVO framework. This behaviour further confirms that the improved rate capability of AlVO@C arises not from enhanced diffusion kinetics but from the suppression of surface degradation and the stabilization of interfacial charge transfer, allowing the intrinsic ion-transport properties of the layered structure to be fully utilized.

Fig. 6f compares the long-term cycling stability of AlVO and AlVO@C negative electrodes. The AlVO@C electrode delivers markedly superior durability, retaining ∼76% of its initial capacitance with ∼97% CE after 2000 cycles at 500 mA g⁻¹ within the 0.4 to −0.8 V window. In contrast, pristine AlVO suffers severe degradation, retaining only ∼5% of its initial capacitance and showing a reduced CE (∼85%) under identical cycling conditions. AlVO@C exhibits a modest performance increase during the first ∼50 cycles, which can be attributed to activation of the carbon shell and improved electrode–electrolyte interfacial wetting.³³ After activation, the capacity gradually declines but remains far more stable than that of pristine AlVO. The latter undergoes rapid deterioration, losing nearly 80% of its capacitance within the first 500 cycles, highlighting its inherent structural vulnerability. The EIS data before and after cycling (Fig. S31 and Table S9) further reinforce this observation. For AlVO@C, the increase in R_CT after 2000 cycles is minimal (∼1.5 Ω), indicating preserved interfacial integrity. In stark contrast, pristine AlVO shows a dramatic increase in R_CT, more than 5.5-fold, reflecting the formation of resistive surface layers and collapse of the layered structure. These changes are consistent with extensive vanadium dissolution and structural instability in the uncoated material. The carbon shell in AlVO@C effectively mitigates these degradation pathways by stabilizing the surface, buffering the volume change (see Fig. S32 for detail), and restricting vanadium-leaching, thereby enabling the robust long-term cycling performance characteristic of the core–shell architecture.

3.2.4. Electrochemical performance testing of the NCBMOF@PBA-12//AlVO@C ZIHSC full cell. Based on the electrochemical evaluation of all candidate electrodes, NCBMOF@PBA-12 and AlVO@C were selected as the positive and negative electrodes, respectively, for the ZIHSC full-cell assembly. A ZnSO₄-silica gel electrolyte was employed as a practical, user-friendly electrolyte in which the aqueous solution is immobilized within an extended 3D silica network. This configuration offers two key advantages: (i) confinement of water molecules through weak hydrogen bonding increases the water-splitting overpotential, thereby widening the electrochemical stability window³⁴ and (ii) the semi-solid electrolyte provides a maintenance-free design and improved handling safety. To understand the real-time electrochemical behaviour within the assembled full cell, the in-operando voltage profiles of the individual electrodes were recorded under slow rate constant-current (500 mA g⁻¹) charge–discharge within the 0–1.2 V regime. Fig. 7a shows the potential evolution of the NCBMOF@PBA-12 positive electrode and the AlVO@C negative electrode versus Ag/AgCl during full-cell operation. The positive electrode exhibits well-defined charge/discharge plateaus, whereas the negative electrode displays slanted, capacitive-like profiles. This contrast confirms the predominantly Faradaic reaction mechanism in NCBMOF@PBA-12 and the pseudo-Faradaic behaviour of AlVO@C, consistent with the individual electrode analyses discussed earlier. Fig. S33 and inset of Fig. 7b present the GCD profiles of the NCBMOF@PBA-12//AlVO@C ZIHSC full cell measured over a current-density range (1–10 A g⁻¹) within 0–1.2 V, 0–1.5 V and 0–1.8 V windows. The corresponding capacitance (F g⁻¹) and CE values are summarized in Fig. 7b. It is noteworthy that the specific capacitance of the full cell was calculated based on the total mass of both electrodes, including the active material, conductive carbon, and PVDF binder. The charge–discharge curves show limiting contribution from the sloping behaviour associated with the pseudo-Faradaic nature of the AlVO@C negative electrode. The CE values remain high at all current densities, with about 98%, 97%, and 96% for voltage windows of 0–1.2 V, 0–1.5 V, and 0–1.8 V, respectively. This indicates excellent reversibility, minimal extent of side reactions, and stable charge-transfer behaviour in the full cell. The energy and power densities of the hybrid supercapacitors were calculated using the full-cell capacitance (C_cell in F g⁻¹) and the operating voltage window (V_max/V_min), according to the following equations:¹⁰

Fig. 7 Performance testing of the NCBMOF@PBA-12//AlVO@C full cell: (a) GCD profile of the full cell at 500 mA g⁻¹ shown together with the corresponding in-operando voltage profiles of the individual positive and negative electrodes, (b) capacitance and cCoulombic efficiency data derived from GCD profiles at multiple current densities (1–10 A g⁻¹) (inset: GCD profiles of the full cell in the 0–1.5 V regime), (c) Ragone plot depicting energy and power rating, (d) Bode plot originated from the EIS study, (e) durability study at 5 A g⁻¹ current density in multiple voltage regimes, and (f) GCD profile of the 5.5 V/4 F prototype device at 6 mA (C-rate) current (inset: demonstration of real-life operation of the prototype device, showing its ability to power a laminated LED setup).

Fig. 7c shows the Ragone plot of the NCBMOF@PBA-12//AlVO@C ZIHSC. The device delivers a moderate energy density of 203 Wh kg⁻¹ at a power density of 900 W kg⁻¹ (1 A g⁻¹), and exhibits a high power density of 9000 W kg⁻¹ at 158 Wh kg⁻¹ (10 A g⁻¹) in the 0–1.8 V regime. Such high power density reflects the rapid charge–discharge kinetics of the hybrid Zn-ion supercapacitor. Fig. 7d presents the Bode plot, where the characteristic breaking frequency (f) is identified at the phase angle of −45°. At this frequency, the capacitive component (–Z″) begins to dominate over the resistive component (Z′), marking the transition from resistive to capacitive behaviour. The inverse of this break frequency defines the response time (τ = 1/2πf) of the device. The NCBMOF@PBA-12//AlVO@C ZIHSC exhibits a response time of ∼4 ms, confirming its excellent rate capability and fast electrochemical dynamics. The long-term durability of the NCBMOF@PBA-12//AlVO@C ZIHSC was evaluated over 5000 charge–discharge cycles at a current density of 5 A g⁻¹ within 0–1.2 V, 0–1.5 V and 0–1.8 V windows (Fig. 7e). The capacitance retention is 87%, 80%, and 71% for voltage windows of 0–1.2 V, 0–1.5 V, and 0–1.8 V, respectively, with corresponding CE values of about 98%, 97%, and 95%, respectively. EIS measurements were carried out before and after cycling, and the corresponding Nyquist plots are shown in Fig. S34, with the R_S and R_CT values summarized in Table S10. The results show a small increase (∼0.4 Ω) in R_CT after cycling in the 0–1.2 V and 0–1.5 V windows, indicating stable interfacial properties and high durability. In contrast, the R_CT nearly doubles in the 0–1.8 V window, likely due to current collector corrosion at higher voltages. In all cases, R_S remains unchanged. After completing the durability tests, the full cell, operated within 0–1.8 V, was disassembled, and both the positive and negative electrodes were analysed using SEM and PXRD. The surface morphology of both electrodes remained largely unchanged before and after cycling, with no evidence of pulverization, cracking, or deformation (Fig. S35). These observations confirm that the electrode structures remain stable throughout the cycling process. The structural integrity of the MOF@PBA-12 positive and AlVO@C negative electrodes was further confirmed by PXRD analysis (Fig. S36). After cycling at 5 A g⁻¹ within 0–1.8 V, the electrodes retain the characteristic diffraction peaks of the pristine materials at the same positions, indicating preservation of the crystal structure. Additional peaks corresponding to the SS304 substrate and amorphous carbon are also observed, and the intensity of the active material peaks decreases, as expected due to electrode processing effects. Overall, the combined post-mortem SEM and PXRD analyses demonstrate that the electrodes retain their morphology and structural coherence, confirming the robustness of the full-cell materials under prolonged cycling. Overall, the electrochemical evaluations and post-mortem analyses indicate that the NCBMOF@PBA-12//AlVO@C full-cell system delivers promising performance with robust structural and configurational stability, highlighting its potential for future scale-up and practical applications. For a practical demonstration, a 5.5 V/4 F integrated prototype device was assembled using three 1.8 V full cells connected in series, with each unit comprising two cells in parallel. The charge–discharge characteristics of the integrated device are presented in Fig. 7f. The inset illustrates the successful operation of a LED setup powered by the 5.5 V prototype, highlighting its capability for real-world energy delivery.

4. Conclusions

This study establishes a rational core–shell design strategy to harmonize Faradaic and capacitive charge-storage processes in aqueous Zn-ion hybrid supercapacitors, addressing the long-standing issues of kinetic mismatch and structural instability in conventional asymmetric configurations. For positive electrodes, three new Ni–Co bimetallic MOFs (NCBMOFs) with varying Ni : Co ratios (2 : 1, 1 : 1, and 1 : 2) were synthesized. Due to restricted bulk redox activity, the pristine NCBMOFs exhibit surface-controlled pseudo-Faradaic charge-storage behaviour, with NCBMOF-12 showing the best performance. To enhance energy-storage capacity, Faradaic Prussian blue analogue (PBA) shells were grown on the pseudo-Faradaic MOF cores, forming NCBMOF@PBA core–shell hybrids. These core–shell materials significantly outperform their pristine counterparts, and NCBMOF@PBA-12 again delivers the best Zn²⁺ storage performance, exhibiting a capacitance of 355–225 F g⁻¹ (100–500 mA g⁻¹), good rate capability (∼65%), and stable cycling (∼80% retention after 2000 cycles at 500 mA g⁻¹). As negative electrode materials, aluminium-doped vanadate (AlVO) shows diffusion-controlled Faradaic Zn²⁺ insertion owing to its flexible structure and vanadium redox activity. However, poor electronic conductivity, high charge-transfer resistance, and large volume changes during Zn²⁺ intercalation lead to rapid capacity fading. These limitations are effectively mitigated by encapsulating AlVO with a capacitive carbon shell (AlVO@C). The AlVO@C electrode delivers a moderate yet practical capacity of ∼82–64 mAh g⁻¹ (296–230 F g⁻¹ at 100–500 mA g⁻¹), along with excellent rate retention (∼78%) and cycling stability (∼76% retention after 2000 cycles at 500 mA g⁻¹). Notably, the carbon shell does not hinder Zn²⁺ diffusion, as confirmed by nearly unchanged diffusion coefficients during cycling.

A hybrid full cell assembled with a pseudo-Faradaic@Faradaic//Faradaic@capacitive (NCBMOF@PBA-12//AlVO@C) configuration and a ZnSO₄-silica gel electrolyte operates stably within a 0–1.8 V window. The ZIHSC delivers a capacitance of 450 F g⁻¹ at 1 A g⁻¹ and a maximum energy density of 203 Wh kg⁻¹ at 900 W kg⁻¹, and retains 158 Wh kg⁻¹ at 9000 W k⁻¹ with a fast response time of ∼4 ms. The device retains ∼71% capacitance after 5000 cycles at 5 A g⁻¹ with high cCoulombic efficiency (∼95%), and post-mortem analyses confirm preserved electrode morphology and crystallinity. Overall, this study demonstrates that electrode-level core–shell engineering effectively synchronizes charge-storage kinetics, enhances material utilization, and improves durability in aqueous Zn-ion hybrid supercapacitors, providing a practical design framework for safe, low-cost, and scalable hybrid energy-storage systems.

Conflicts of interest

There are no conflicts to declare.

Data availability

CCDC 2517104 contains the supplementary crystallographic data for this paper.³⁵

The datasets generated during and/or analysed during the current study are not publicly available due to very large size but are available from the authors on reasonable request.

Supplementary information (SI): literature survey tables, synthesis protocols, experimental methodologies, and detailed crystallographic, spectroscopic, morphological, and electrochemical characterization data of the positive and negative active materials. See DOI: https://doi.org/10.1039/d5ta10451c.

Acknowledgements

Financial support from the Science and Engineering Research Board (SERB), India (File No. SRG/2019/000296, SRG/2023/000238 & SRG/2023/000521) and University Grants Commission (UGC), India (No. F.30-509/(BSR)) is gratefully acknowledged. NB thanks the Council of Scientific and Industrial Research (CSIR), India for a JRF (File No. 08/0155(22864)/2025-EMR-I). PKM, PP and PM acknowledge NIT Durgapur for providing a fellowship.

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Footnote

† These authors contributed equally to this work as co-first authors.

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