Elucidating Mn²⁺/Mn³⁺ and Ni⁰/Ni²⁺ Redox Synergy in Hair-Derived Carbon-Supported Ag/Ni–MnO _x Supercapacitor

Abstract

Despite their critical importance, developing sustainable high-performance supercapacitor (SC) electrodes with long-term stability poses significant challenges. Herein, we report a novel ternary composite electrode in which Ag/Ni-doped manganese oxide (Ag/NiO _x @Mn _y O _z ) is supported on human hair-derived activated carbon (HHC). This composite is synthesized via a one-pot hydrothermal process followed by thermal annealing at 800 °C, a strategy that creates a conductive Ag/Ni bimetallic network and abundant oxygen vacancies in the NiO _x and Mn _y O _z phases. During operation, operando X-ray absorption spectroscopy (XAS) confirms reversible dual-ion redox transitions (Mn²⁺/Mn³⁺ and Ni⁰/Ni²⁺) in the cathode, highlighting the material’s enhanced redox activity. As a result, HHC-supported Ag/NiO _x @Mn _y O _z exhibits an exceptional specific capacitance (Cs) of 1770 F g^–1 at 5 mV s^–1 in three-electrode tests. When assembled into an asymmetric hybrid supercapacitor (AHSC), the device delivers a high energy density of 37.53 Wh kg^–1 and a power density of 2251.8 W kg^–1 at 3 A g^–1 while retaining ∼82% of its initial capacitance after 5000 charge–discharge cycles. These results confirm the effectiveness of our sustainable HHC-supported Ag/NiO _x @Mn _y O _z framework in addressing the enduring trade-off between energy density, power density, and cycling stability in next-generation SCs.

1. Introduction

Due to environmental concerns associated with fossil fuel consumption, the growing demand for advanced energy storage devices has intensified. , Across various industries, rechargeable batteries, supercapacitors (SCs), and other energy storage technologies, including consumer electronics and automotive applications, are widely utilized. − In particular, SCs are emerging as promising candidates owing to their high power density, rapid charge/discharge capabilities, and exceptional cycling stability. , SCs are classified into two primary types based on their mechanisms of charge storage: electric double-layer capacitors (EDLCs) and pseudocapacitors (PCs). EDLCs store charge through ion accumulation at the electrode–electrolyte interface. In contrast, PCs rely on fast and reversible redox reactions within active materials. ,

As pseudocapacitive materials, transition metal organic frameworks, double-layered hydroxides, and transition metal oxides (TMOs) are widely explored. − Multiple oxidation states of transition metals enable rapid redox reactions for charge storage. Of the various TMOs (RuO₂, MnO₂, Fe₃O₄, WO₃, V₂O₅, MoO₃), manganese­(II) oxide (MNO) and manganese­(IV) oxide (MnO₂) are the most extensively studied metal oxides for SC electrodes due to their favorable properties. Manganese exists in several stable oxidation states, such as MnO, MnO₂, MnO₃, and MnO₄, each exhibiting distinct electrochemical behavior. Among these oxides, MnO theoretically offers the highest specific capacitance due to redox reactions between Mn²⁺ and Mn³⁺. Enhanced charge storage performance has also been ascribed to the intercalation and deintercalation of electrolyte ions, driven by the reversible redox transition between Mn²⁺ and Mn³⁺ within the electrode material. However, the practical application of unmodified MnO as a supercapacitor electrode is limited by its low electrical conductivity (10^–5–10^–6 S/cm) and poor structural stability, both of which severely hinder its overall performance. Likewise, MnO₂ offers fast and reversible redox kinetics, robust chemical and thermal stability, excellent charge–discharge behavior, a wide operating potential (∼1 V), and a high theoretical specific capacitance of 1370 F g¹. , However, pure MnO₂ has limited electronic conductivity, which can hinder its rate capability and the utilization of its full theoretical capacitance in practical devices. Strategies to improve the electrochemical performance of MnO₂ have therefore focused on electronic structure optimization, such as noble metal incorporation, defect engineering, and composite formation.

Noble metal nanoparticles are of great interest as conductive additives in nanocomposites with carbon materials, conducting polymers, and metal oxides/hydroxides due to their exceptional electrical conductivity. , Such nanoparticles facilitate efficient electron transport from redox-active sites to the current collector, thereby enhancing charge transfer and utilization of the active material. Among noble metals, silver (Ag) stands out for its superior conductivity and comparatively lower cost. , Nanoscale Ag particles can form conductive bridges and electrical contact points within an electrode, improving overall electrochemical performance. , The integration of Ag with carbon-based substrates or metal oxides has been extensively investigated, demonstrating significant enhancements in both SCs and lithium-ion batteries. In general, doping, or modifying MnO₂ with highly conductive components like noble metals, and creating structural defects such as oxygen vacancies, is an effective way to increase its electrical conductivity and expose more active sites for redox reactions. ,

Forming composite electrodes is another promising approach to overcome the limitations of a single metal oxide. By hybridizing two or more electroactive components, one can leverage synergistic effects that boost capacitive performance beyond what each component could achieve alone. , In this context, nickel oxide (NiO) is widely combined with other metal oxides. For example, NiO@MnO₂, NiO/ZnO, NiO@Co₃O₄, and NiO/NiCo₂O₄ composites have been reported to take advantage of multiple redox processes within an electrode. NiO is attractive because of its low toxicity, environmental benignity, and excellent chemical and thermal stability. However, NiO suffers from low intrinsic electrical conductivity that leads to high internal resistance and subpar rate capability; as a result, its practical capacitance often falls well below the theoretical maximum. , Integrating NiO, therefore, with additional redox-active oxides and combining both with conductive matrices, viz, carbon supports, is necessary to provide extra electron pathways and active sites, thus increasing overall pseudocapacitive behavior.

As sustainable alternatives to conventional carbons for energy storage, biomass-derived carbon materials have gained much attention, offering high electronic conductivity, tunable porosity, good stability, and low-cost production. As such, human hair, a widely available waste material that decomposes slowly, presents a promising precursor for porous carbon. Improper disposal of human hair waste can cause environmental problems, namely, harmful emissions from burning and airborne particulate hazards, but repurposing this waste into useful carbon materials addresses both waste management and energy storage needs. Chemically, human hair consists predominantly of keratin (>90 wt %), along with lipids (∼2 wt %), nucleic acids, and carbohydrates. Its elemental composition is roughly 50 wt % carbon along with substantial heteroatoms: ∼22 wt % oxygen, 16 wt % nitrogen, and 3–4 wt % sulfur. These heteroatoms, when retained in the carbon structure, can contribute faradaic pseudocapacitance by providing additional redox-active functional groups, improving charge storage capacity. , Utilizing human hair-derived carbon (HHC) for electrode applications not only mitigates environmental pollution but also yields a heteroatom-rich, high-surface-area conductive scaffold that can enhance electrochemical performance in SCs.

Despite progress in developing nanocomposites and using biobased carbons, achieving a holistic solution that combines high capacitance, excellent rate performance, and long-term cycling stability in a sustainable electrode is challenging. Many MnO₂- or NiO-based composites still face trade-offs. For instance, improved energy density often comes at the expense of cycle life; incorporating noble metals can increase cost or complexity if not executed judiciously. Moreover, the full potential of waste-derived carbon supports in such multicomponent systems is yet to be realized. Notably, an electrode design that combines the good pseudocapacitance properties of MnO₂, the complementary redox activity of NiO, and the superior conductivity of Ag is crucial. Incorporating these materials within a sustainable framework aims to maximize surface area and electrochemical response.

This work investigates a novel ternary composite electrode that synergistically integrates Mn and Ni oxides with Ag nanoparticles on HHC support. The novelty of our approach lies in two key aspects: (1) Waste-derived HHC is used as a high-surface area, heteroatom-doped support for the metal oxides. (2) In-situ formation of Ag/Ni bimetallic framework and oxygen-vacancy-rich oxide phases, achieved through thermal tuning, promotes dual redox reactions at the electrode. This sustainable design simultaneously improves the electrode’s surface area, electronic conductivity, and electrochemical stability. The synergistic effect of the HHC support and high-temperature annealing yields a robust composite architecture with optimized charge transport pathways, emphasizing its capability for future applications. Our findings demonstrate significant advancement in electrode design for SCs, highlighting the effectiveness of combining bioderived carbon supports with bimetallic oxide networks to achieve high energy and power densities concurrently with remarkable cycling stability.

2. Experimental Section

2.1. Chemicals/Materials

All chemicals utilized were used as purchased without further refining. The chemicals/reagents include manganese chloride (MnCl₂, ≥98%, grade AR, Kemaus), silver nitrate (AgNO₃, 99.9%, POCH), nickel nitrate hexahydrate (Ni­(NO₃)₃·6H₂O, 98%, Loba Chemie PVT LTD), sodium hydroxide (NaOH, 98%, Loba Chemie PVT Ltd.), urea (Co­(NH₂)₂, grade AR, QRec), ethanol (95% C₂H₅OH), potassium hydroxide (KOH, ≥99.0%, Kemaus), N-methyl-2-pyrrolidone (NMP, QReC), carbon black (Super P, TIMCAL), human hair (middle-aged African male), and nickel foam (NF) (thickness = 0.6 mm).

2.2. Synthesis of Human Hair-Based Activated Carbon (HHC)

In Figure , the step-by-step procedure for the synthesis of HHC is depicted. Human hair was collected from a middle-aged African and meticulously cleaned using deionized (DI) water and acetone to eliminate dust and oil. The hair was then dried in an oven at 80 °C for 24 h. After drying, the hair was placed in a quartz boat and subjected to pyrolysis within a tubular furnace under an inert (N₂) atmosphere, following a two-stage process. Initially, the furnace was purged with nitrogen for 15 min to eliminate residual oxygen. In the first stage of pyrolysis, the temperature was elevated to 220 °C at a rate of 5 °C per min and maintained at that temperature for 30 min to enable degasification and remove moisture and volatile substances.

1.

Schematic illustration for the synthesis of HHC and C–Ag/NiO _x @Mn _y O _z composites.

The second stage involved heating the material further to 900 °C at the same heating rate and maintaining it for 1 h to ensure complete carbonization. The resulting carbonized hair was then subjected to chemical activation by mixing it with KOH pellets in a weight ratio of 1:4. This KOH-infused carbon was reheated at 900 °C under identical conditions to bolster its porosity. Following activation, the sample was neutralized with 1 M HCl, thoroughly washed with DI water, and filtered. Finally, the material was dried at 80 °C for 24 h.

2.3. Synthesis of HHC-Supported Ag/NiO _x @Mn _y O _z Composites

First, equimolar solutions of 0.2 M MnCl₂, Ni­(NO₃)₂·6H₂O, and AgNO₃ were combined with a sonicated suspension of 200 mg of HHC in 30 mL of DI water. Then, the mixture was stirred continuously at 230 rpm using a magnetic stirrer for 30 min. After that, solutions of 0.4 M CO­(NH₂)₂ and 5 M NaOH were added to the mixture and stirred for an additional 30 min. The resultant colloidal suspension was transferred into a PTFE-lined autoclave and heated in an oil bath at 180 °C for 6 h. The precipitate was collected, washed thoroughly with water and ethanol, and dried in an oven at 80 °C for 24 h. The dried material was then divided into three portions, two of which were subjected to different annealing temperatures of 400 and 800 °C for 1 h under identical annealing conditions described in Section .

The samples were labeled as C–Ag/NiO _x @Mn _y O _z , C–Ag/NiO _x @Mn _y O _z @400 °C, and C–Ag/NiO _x @Mn _y O _z @800 °C. For comparison, a control sample (Ag/NiO _x @Mn _y O _z ) was prepared following the same procedure but without the inclusion of HHC and annealing.

2.4. Sample Characterization and Electrochemical Investigation

Detailed descriptions of the material characterization techniques and the methodologies employed to assess the electrochemical performance of the synthesized materials are provided in the Supporting Information.

2.5. Assemblage of C–Ag/NiO _x @Mn _y O _z @800 °C/HHC AHSC

AHSC was fabricated using a C–Ag/NiO _x @Mn _y O _z @800 °C-coated nickel foam (NF) as the cathode, with a mass loading of ∼1.5–2.0 mg, and an HHC-coated NF as the anode, with a mass loading of ∼2.0–2.5 mg. A Whatman glass microfiber, presoaked in 2 M aqueous KOH electrolyte solution for 30 min, was employed as the separator. The assembled components, including the electrodes and separator, were then integrated into a stainless steel CR2025 lithium coin cell casing.

3. Results and Discussion

3.1. X-ray Diffraction (XRD) Analysis

The synthesized samples were systematically analyzed via XRD, as illustrated in Figure a,b. In Figure a, two broad diffraction peaks at ∼2θ = 25.7 and 43.1° are displayed, corresponding to the (002) and (100) planes of the turbostratic carbon structure. These peaks confirm the presence of crystallites, as indicated by the graphite-like (002) and (100) reflections observed in the HHC spectrum. The peak at 2θ = 43.1° specifically denotes a higher degree of interlayer condensation within HHC, a structural attribute expected to significantly augment its electrical conductivity. Additionally, the peak observed at 2θ = 10.5° suggests a transition of the amorphous phase in HHC toward graphitic structures. However, the relatively low intensity of the 2θ = 43.1° peak suggests that only a minimal amount of graphitic material is present within the overall structure. ,

2.

XRD patterns and XPS spectra of synthesized materials: (a) XRD patterns of human HHC, (b) XRD patterns of C–Ag/NiO _x @Mn _y O _z composite materials at different conditions, (c) XPS wide scan, (d) Mn 2p XPS spectra of C–Ag/NiO _x @Mn _y O _z @800 °C, (e) Ni 2p XPS spectra of C–Ag/NiO _x @Mn _y O _z @800 °C, (f) Ag 3d XPS spectra of C–Ag/NiO _x @Mn _y O _z @800 °C, (g) C 1s XPS spectra of C–Ag/NiO _x @Mn _y O _z @800 °C, and (h) O 1s XPS spectra of C–Ag/NiO _x @Mn _y O _z @800 °C.

In Figure b, the presence of metallic silver (Ag⁰) is confirmed by the diffraction peaks observed at 2θ angles of 38.1, 64.1, 77.1, and 81.4°, corresponding to the (111), (220), (311), and (222) planes, respectively. These peaks are in close agreement with the reference data provided by JCPDS card No. 04–0783. Further, as the annealing temperature increases up to 800 °C, distinct differences in the Mn oxide phases are revealed. The diffraction peaks of Mn _y O _z in the composites annealed at 400 °C, occurring at 2θ = 20.5, 29.4, and 44.0°, are indicative of the tetragonal phase of Mn₃O₄ as referenced by JCPDS Card No. 00-024-0734. This observation suggests the coexistence of Mn₂O₃ and MnO within these samples. However, at 800 °C, a pronounced molecular transformation from Mn₂O₃ to MnO occurs, which is attributed to thermal reduction. The transition from Mn₂O₃ to MnO is expected to improve electrochemical performance by amplifying electrical conductivity and oxygen vacancy density, resulting in quicker charge transfer and ion diffusion. The lower oxidation state of MnO compared to Mn₂O₃ is capable of promoting more reversible redox processes, increasing specific capacitance and cycle stability. This phase shift also improves structural flexibility, allowing the electrode to tolerate several charge–discharge cycles with minimum degradation. The thermal treatment process also promotes the formation of metallic Ni, as indicated by the appearance of a diffraction peak, corresponding to the (200) plane of Ni. This peak is exclusively observed in the C–Ag/NiO _x @Mn _y O _z @800 °C sample, confirming the influence of high-temperature annealing on structural modification and phase evolution.

3.2. X-ray Photoelectron Spectroscopy (XPS) Analysis

To gain a comprehensive understanding of the electronic structure and elemental composition of the composites, XPS analyses were conducted. As depicted in Figure c, the XPS survey spectrum reveals distinct signals corresponding to Mn, Ag, Ni, C, and O, affirming their successful incorporation into the composite matrix. In Figure d, the high-resolution XPS spectrum of Mn 2p for the C–Ag/NiO _x @Mn _y O _z @800 °C sample was deconvoluted into three peaks located at 641.4, 643.1, and 646.7 eV, which are ascribed to Mn²⁺, Mn³⁺, and Mn⁴⁺ species, respectively. The high peak intensity of the Mn²⁺ species demonstrates the prevalence of MnO within the sample, aligning well with the XRD findings.

In Figure e, the high-resolution XPS spectrum of Ni 2p is illustrated, which comprises two spin–orbit doublets at binding energies of 855.7 and 874.2 eV, with a spin energy separation of 18.5 eV, corresponding to the Ni 2p_3/2 and Ni 2p_1/2 components. Satellite peaks detected at 861.3 and 879.7 eV are credited to the shakeup phenomena of Ni 2p_3/2 and Ni 2p_1/2. A meticulous peak fitting analysis of the Ni 2p_3/2 component reveals two peaks at 855.7 and 857.9 eV, consistent with the binding energy values associated with Ni²⁺ and Ni³⁺ species. The presence of atomic Ni at 853.3 eV on the surface of the C–Ag/NiO _x @Mn _y O _z @800 °C sample results from the thermal reduction of NiO at elevated annealing temperature. In Figure S1, the comparative XPS spectra of Mn 2p of C–Ag/NiO _x @Mn _y O _z @400 °C and C–Ag/NiO _x @Mn _y O _z @800 °C samples are seen to exhibit different patterns, signifying the impact of thermal tuning derived from high-temperature annealing.

In Figure f, the XPS spectrum of Ag 3d is presented, showing two prominent peaks at 368.2 and 374.2 eV, attributed to Ag 3d_5/2 and Ag 3d_3/2 of metallic silver (Ag⁰). This observation highlights the dominance of metallic silver within the sample. In Figure g, the C 1s spectrum is shown, providing additional chemical information, with peaks observed at 289.0, 287.6, 286.4, and 284.9 eV, which are ascribed to O–C–O, C = O, C–O, and C–C bonds, respectively. In Figure h, the O 1s spectrum reveals peaks at 532.9, 531.4, and 529.9 eV, which conform to adsorbed oxygen, oxygen associated with defect sites (O₂-containing species), and metal–oxygen bonds (O–Mn or O–Ni). , The combined results from XRD and XPS analyses verify the successful integration of metallic silver (Ag) into the HHC-supported NiO _x and Mn _y O _z nanostructure.

3.3. Morphology Characterization

In Figure , a comprehensive analysis of the surface morphology and compositional attributes of the synthesized materials, utilizing FE-SEM, EDS, and HR-TEM techniques, is presented. In Figure a,b, the FE-SEM micrographs of HHC are depicted, exhibiting a porous nanostructure essential for enhancing surface area and promoting electrochemical performance. In Figure c, EDS elemental mapping is illustrated, confirming the consistent distribution of Ag, Mn, Ni, O, and C within the composite, affirming their successful incorporation. In Figure S2, the EDS spectra and percentage weight composition of all the constituent elements are given. Figure S3 presents the relative concentrations of constituent metals (Ag, Mn, and Ni) in the samples, as determined by inductively coupled plasma (ICP) analysis. The results indicate a high degree of compositional homogeneity across the samples.

3.

FE-SEM Micrographs and EDS mapping of synthesized materials: (a, b) FE-SEM micrographs of HHC, (c) EDS mapping of Ag/NiO _x @Mn _y O _z @800 °C, (d) FE-SEM micrograph of Ag/NiO _x @Mn _y O _z , (e) FE-SEM micrograph of C–Ag/NiO _x @Mn _y O _z , (f, g) FE-SEM micrographs of C–Ag/NiO _x @Mn _y O _z @400 °C, (h, i) FE-SEM micrographs of C–Ag/NiO _x @Mn _y O _z @800 °C, (j, k) HR-TEM micrograph of C–Ag/NiO _x @Mn _y O _z @400 °C, and (l) HR-TEM micrographs of C–Ag/NiO _x @Mn _y O _z @800 °C.

In addition, In Figure d shows the compact structure of the Ag/NiO _x @Mn _y O _z composite synthesized without HHC. The absence of HHC results in a more aggregated and less crystalline morphology.

Under identical conditions, the incorporation of HHC with Ag/NiO _x @Mn _y O _z produces a more loosely packed and crystalline architecture (Figure e), as evident in the more pronounced XRD peaks associated with the C–Ag/NiO _x @Mn _y O _z sample. This enhancement in structural characteristics emphasizes the critical role of HHC as an effective support matrix for the composite. In addition, In Figure f,g, the formation of hexagonal NiO crystals attached to the Mn₃O₄ nanosheets, when annealed at 400 °C, is observed. Notably, Figure h,i reveal substantial impregnation of Ag and Ni on MnO as the partial reduction of Mn₃O₄ to MnO occurs at 800 °C. Furthermore, HR-TEM analysis offers a detailed visualization of the crystalline nanostructures. In Figure j,k, the overlapping hexagonal NiO molecules formed at 400 °C are highlighted. In Figure l, the uniform distribution of Ag nanospheres and Ni nanoparticles across the MnO surface is demonstrated.

3.4. Surface Area and Particle Size Measurements

In Figure a, a characteristic type I Brunauer–Emmett–Teller (BET) isotherm, corresponding to the HHC sample, is illustrated. This isotherm portrays a remarkable specific surface area (SSA) of 2605.93 m² g^–1 with an average pore size of 2.18 nm. The negligible hysteresis observed in the isotherm indicates a predominantly closed-ended pore structure with similar adsorption and desorption mechanisms. In Figure b, a comparative BET analysis conducted on the composites, with and without HHC support, is presented. Both samples exhibit type IV isotherms, which are indicative of mesoporous materials. Significantly, the composite incorporating HHC exhibits a substantially higher SSA (91.4 m² g^–1): approximately double that of the sample synthesized without HHC (40.6 m² g^–1). This outcome highlights the critical function of HHC as a conductive supporting material, capable of boosting the specific surface area and electrochemical properties of the composites. The corresponding pore diameters measure 7.08 and 11.09 nm for the samples, with and without HHC, respectively. These results further demonstrate the beneficial impact of HHC inclusion.

4.

BET and particle size analyses (PSA) of the composites: (a) BET adsorption–desorption isotherm for HHC, (b) Combined BET adsorption–desorption isotherms for Ag/NiO _x @Mn _y O _z and C–Ag/NiO _x @Mn _y O _z composites, (c) PSA of Ag/NiO _x @Mn _y O _z , (d) PSA of C–Ag/NiO _x @Mn _y O _z , (e) PSA of C–Ag/NiO _x @Mn _y O _z @400 °C, and (f) PSA of C–Ag/NiO _x @Mn _y O _z @800 °C.

In Figure c–g, an analysis of particle size is shown. Average particle sizes of 1.26, 2.08, 2.35, and 12.59 μm for Ag/NiO _x @Mn _y O _z , C–Ag/NiO _x @Mn _y O _z @, C–Ag/NiO _x @Mn _y O _z @400 °C, and C–Ag/NiO _x @Mn _y O _z @800 °C are displayed. These findings are consistent with the FE-SEM analysis results, affirming that the agglomeration of Ag and Ni nanoparticles becomes increasingly prominent at elevated annealing temperatures. The trend of increasing particle size, as observed in Figure c–f, reveals a clear correlation between annealing temperature and particle growth. This outcome indicates that thermal treatment facilitates the agglomeration and coalescence of metallic nanoparticles, particularly Ag and Ni. Such particle growth at elevated temperatures can significantly enhance electrochemical performance by reducing interfacial resistance and improving electron transport pathways. Larger particles typically possess a reduced number of grain boundaries, which act as barriers to charge mobility. Consequently, the observed increase in particle size at elevated temperature thermal treatment can contribute to a more continuous and conductive framework within the electrode material with improved crystallinity and structural stability, crucial for long-term cycling and capacitive retention.

In Figure h, a broad particle size distribution is observed. This trend can be attributed to a combination of rapid particle growth dynamics and concurrent phase transformations within the composite’s constituents. At such a high annealing temperature, a diverse growth rate of particle sizes is expected, especially in a ternary composite, as some particles grow more extensively than others due to local energy and compositional differences. Moreover, XRD and subsequent ex-situ XAS analyses confirm that at 800 °C, the formation of a new molecular phase, MnO, becomes evident alongside the reduction of NiO _x to atomic Ni. These transformations involve substantial structural rearrangements that contribute to nonuniform growth rates and morphological heterogeneity. The emergence of MnO and metallic Ni not only introduces new nucleation sites but also alters the thermodynamic and kinetic growth behavior of the existing phases, resulting in varied crystallite sizes.

3.5. Electrochemical Measurements in Half Cell

In Figure a–e, for all samples, the comparative cyclic voltammetry (CV) curves are delineated. The CV curves were measured within a specified voltage window of 1.0 V for HHC and 0.6 V for the composite materials, respectively. In Figure S4, a three-electrode configuration used for electrochemical analysis is depicted. As shown in Figure a, the CV curves for HHC display nearly rectangular profiles, characteristic of EDLC typically observed in carbon-based materials. In contrast, in Figure b–e, the CV curves of the composite samples reveal the presence of two prominent oxidation and reduction peaks, particularly at lower scan rates. This observation indicates the occurrence of at least two simultaneous redox reactions within the system, improving overall capacitance.

5.

Three-electrode’s CV and GCD features of the electrode: (a) CV of the HHC, (b) CV of Ag/NiO _x @Mn _y O _z , (c) CV of C–Ag/NiO _x @Mn _y O _z , (d) CV of C–Ag/NiO _x @Mn _y O _z @400 °C, (e) CV of C–Ag/NiO _x @Mn _y O _z @800 °C, (f) Specific capacitance of the electrodes at various scan rate, (g) GCD plot of HHC, (h) GCD curve of Ag/NiO _x @Mn _y O _z , (i) GCD curve of C–Ag/NiO _x @Mn _y O _z , (j) GCD curve of C–Ag/NiO _x @Mn _y O _z @400 °C, (k) GCD curve of C–Ag/NiO _x @Mn _y O _z @800 °C, and (l) Specific capacitance of the electrodes at various specific currents.

In Figure f, HHC exhibits a Cs of 418 F g^–1 at a scan rate of 5 mV s^–1. While all composites demonstrated impressive specific capacitances, the highest Cs of 1770 F g^–1 at 5 mV s^–1 was achieved by the Ag/NiO _x @Mn _y O _z @800 °C sample. This exceptional performance can be attributed to the presence of impregnated, free-standing redox-active Ni atoms on the surface of HHC-supported porous MnO molecules, which facilitate enhanced charge storage. In Figures S5 and S6, an average b-value of 0.54 is denoted, signifying a combination of capacitive and diffusion-controlled processes. Moreover, Figure S7 illustrates the proportion of current signal contributed by capacitive and diffusion processes at varying sweep rates. In Table S1, a comparison of HHC with other bioderived electrode materials is provided.

Galvanostatic charge–discharge (GCD) analysis of HHC, conducted at a current density range of 1–10 A g^–1, reveals symmetrical and triangular profiles with a nearly steady specific capacitance of 271 F g^–1 at 1 A g^–1. During charge–discharge processes, the GCD curves of the composite materials display distinct two-stage profiles, indicating dual redox-active transitions associated with Mn²⁺/Mn³⁺ and Ni⁰/Ni²⁺ within the electrode (Figure g–k). In Figure l, the specific capacitance values for each composite sample are summarized. In Table , a comparative analysis of the electrochemical performance of this study against recent literature is provided.

1. Comparative Electrochemical Outputs of This Study with Recently Reported Literature.

material	synthesis method	electrolyte	specific capacitance Cs(F g –1 )	energy density(Wh kg–1)/Cs retention	refs
MnO2/CNTs	electrodeposition	0.5 M Na2SO4	3310	–/65.4%
Na+-MnO2	solvothermal	1 M KOH	655	27/95%
Cu-MnO x	electrodeposition	1 M Na2SO4	92.3	25.13/44.9%
MnO2@Ti3C2T x	solution immersion	1 M KOH	324.1	30.8/95.98%
MnO2/NiCo	hydrothermal	1 M KOH	2572.7	–/78.9
MnO2/MnS	hydrothermal		977.6	122.23/91.5
Al-doped δ-MnO2	hydrothermal	6 M KOH	174	18.5/90%
MnO2/CNTs	chemical reduction		617.7	111/71.4%
PCDWT/MnO2/PANI	pyrolysis	trifluoroacetic acid (TFA)	369.6	36.7/87.7%
C–Ag/NiO x @Mn y O z @800 °C	hydrothermal	2 M KOH	1770	37.53/82%	this study

In Figure a, a schematic representation of the redox reactions occurring at the C–Ag/NiO _x @Mn _y O _z @800 °C/KOH interface is presented. The illustration highlights the dual role of Mn and Ni species in promoting OH^– reduction when a potential difference is applied across the cell terminals. In Figure d, the electrochemical impedance spectroscopy (EIS) results are presented through Nyquist plots. The C–Ag/NiO _x @Mn _y O _z @800 °C sample exhibits the lowest equivalent series resistance (R₁) of ∼0.48 Ω and a charge transfer resistance (R₂) of 1.41 Ω, which is significantly lower than that of the sample annealed at 400 °C (Table S2). This observation is consistent with the fact that the infused metallic Ni derived at 800 °C acts as a superior electrical conductor compared to the NiO present in the composite material annealed at 400 °C.

6.

Reaction mechanisms, diffusion coefficient, charge transfer coefficient, and reaction rate constant evaluation plots: (a) Simultaneous redox reactions occurring at the cathode, (b) Relationship between peak current (A) and υ^0.5, (c) Relationship between peak potentials (E _p) and ln (υ), (d) EIS analysis of the composites, and (e) EIS-fitted equivalent circuit.

Further, both C–Ag/NiO _x @Mn _y O _z @400 °C and C–Ag/NiO _x @Mn _y O _z @800 °C samples exhibit nearly identical spectral profiles, characterized by low electrochemical impedance across R₁, R₂, and double-layer resistance (R₃), as detailed in Table S2. These findings indicate that increasing the annealing temperature from 400 to 800 °C does not negatively affect the material’s conductivity. Instead, the higher annealing temperature amplifies both the material’s electrical conductivity and charge transport capabilities by generating oxygen vacancies through thermal reduction. In Figure e, the fitted EIS equivalent circuit is illustrated, providing a visual representation of the circuit elements that impede current flow. The excellent impedance feature of the C–Ag/NiO _x @Mn _y O _z @800 °C sample highlights its potential for efficient charge transfer, thereby improving the overall electrochemical performance of the composite material.

The mass loading ratio of both electrodes, which forms the basis of electrochemical performance indices, was estimated at 0.39, using eq :

$$\frac{M_{+}}{M_{-}}=\frac{C_{\mathrm{m}-}}{C_{\mathrm{m}+}}\times \frac{\mathrm{\Delta }{V}_{-}}{\mathrm{\Delta }{V}_{+}}$$ 1

where C _m+ and C _m– represent the capacitance of the positive and negative electrodes, respectively, corresponding to their voltage windows ΔV ₊ and ΔV _–. By applying Laviron equations, good linear regressions were observed in the plots of both anodic and cathodic peak potentials against ln (υ), as shown in Figure b, yielding the following linear eqs and :

$$E_{\mathrm{pa}}(V)=0.0364\ln v+0.6907(R=0.97)$$ 2

$$E_{\mathrm{pc}}(V)=-0.0325\ln v-0.0273(R=0.97)$$ 3

By utilizing the gradient of either E _pc (V) or E _pa (V) in eqs and and assuming α = 0.5, the number of electrons transferred (n) is estimated at ∼2, indicating a two-electron reduction process at the cathode. The charge transfer coefficient (α) was subsequently evaluated at 0.56. The heterogeneous electron transfer rate constant (k _s) was determined using eq , as given by Reddy et al. In Table S3, for all scan rates, the values of the heterogeneous reaction rate constant (k _s) are summarized. Notably, k _s increased with increasing scan rate, confirming the rapid redox kinetics at higher sweep rates:

$$E_{\mathrm{pc}}(V)=E^{\mathrm{o}}-\frac{\mathit{RT}}{\alpha \mathit{nF}}\ln v$$ 4

$$E_{\mathrm{pa}}(V)=E^{\mathrm{o}}-\frac{\mathit{RT}}{(1-\alpha )\mathit{nF}}\ln v$$ 5

$$\log k_{\mathrm{s}}=\alpha \log (1-\alpha )+(1-\alpha )\log \alpha -\log \frac{\mathit{RT}}{\mathit{nFv}}-\alpha (1-\alpha )\frac{\mathit{nF}\mathrm{\Delta }\mathit{Ep}}{2.3\mathit{RT}}$$ 6

where E° is the formal potential, R is the universal gas constant, T is the absolute temperature, α is the electron transfer coefficient, n is the electron transfer number, and F is the Faraday constant. In addition, the diffusion coefficient (D), which quantifies the ion transport kinetics within the electrode interface, was determined using the Randles-Sevcik equation, as expressed in eq . By analyzing the linear relationship between the peak current (i _p) and the square root of the scan rate (υ), the diffusion coefficient was estimated to be 1.41 × 10^–4 cm² s^–1 (Figure c). This value highlights the efficient ion mobility within the electrode, reinforcing the material’s suitability for high-performance electrochemical applications:

$$i_p=2.69\times {10}^5{n}^{3/2}\mathit{AC}{(\mathit{Dv})}^{1/2}$$ 7

where n = number of transferred electrons, A = area of electrode (cm²), C = bulk concentration (mol cm^–3) D = diffusion coefficient (cm² s^–1), and v = scan rate (mV s^–1).

The suspected redox reactions occurring within the system can be represented as

$${\mathrm{Ni}}^0+{\mathrm{Mn}}^{2+}+x[{\mathrm{OH}}^{-}]\leftrightarrow {\mathrm{Ni}}^{2+}+{\mathrm{Mn}}^{3+}+{y\mathrm{H}}_2\mathrm{O}+z[e^{-}]$$ 8

where x, y, and z are positive integers.

3.6. Electrochemical Measurements of AHSC

In Figure a, a comparative cyclic voltammogram of HHC and C–Ag/NiO _x @Mn _y O _z @800 °C, at a scan rate of 5 mV s^–1, is presented. The substantially enhanced peak current achieved by the composite material, in comparison to the HHC anode, is highlighted. Notably, the current signals of both materials effectively complement one another, enabling a broader voltage window of 1.5 V. In Figure b, CV performed across a 1.5 V potential window at sweep rates ranging from 5 to 100 mV s^–1, demonstrates a consistent and stable pattern, indicating the robustness of AHSC. This stability is attributed to the synergistic interaction between the EDLC mechanism at the HHC anode and the faradaic redox reactions occurring at the C–Ag/NiO _x @Mn _y O _z @800 °C cathode. AHSC exhibits a maximum Cs of 150.7 F g^–1 at a scan rate of 5 mV s^–1, as shown in Figure S8.

7.

Electrochemical analyses of AHSC: (a) Combined CV curves of the HHC and C–Ag/NiO _x @Mn _y O _z @800 °C at 30 mV s^–1, (b) CV curves of AHSC, (c) GCD curve of AHSC, (d) AHSC assembling and illumination of 3.2 V LED bulb, (e) Capacitance retention after 5000 charge–discharge cycles, and (f) Specific capacitance, energy density, and power density values of AHSC.

In Figure c, GCD tests were conducted over a range of current densities from 3 to 15 A g^–1, within the same potential window. The GCD curves, particularly at lower current densities, exhibit nontriangular profiles, confirming the supercapacitive behavior of the AHSC device. Figure c exhibits an asymmetric feature due to different mechanisms of charge and discharge. In supercapacitors, discharging typically occurs faster than charging, reflecting an inherent asymmetry in their operation. In AHSC, during charging, OH^– and K⁺ migrate and adsorb onto the electrode surfaces to undergo redox reactions, a process slowed by ion diffusion and the porous structure of the electrodes. In contrast, discharging happens more rapidly as the electrostatic forces holding the ions weaken, allowing for fast ion desorption and electron flow. Factors like internal resistance and electrode polarization also accelerate voltage drop during discharge, further reducing its duration compared to charging. However, the electrochemical outputs of C–Ag/NiO _x @Mn _y O _z @800 °C/HHC AHSC are irrespective of the asymmetric feature of the charge–discharge curves.

Furthermore, in Figure d, the practical assembly of AHSC and its capability to successfully illuminate a 3.2 V LED bulb, is shown. In Figure S9, the practical applicability of the device is highlighted. The exceptional performance of this device is attributable to the availability of high volume of oxygen vacancy observed in C–Ag/NiO _x @Mn _y O _z @800 °C. Figure S10 shows the spectra an electron paramagnetic resonance (EPR) conducted to assess the presence of oxygen vacancies. Initial evidence of vacancy formation was detected in C–Ag/NiO _x @Mn _y O _z @400 °C. Using a comparable mass loading of approximately 94–98 mg across samples, the spectrum of the 800 °C-annealed sample was significantly distorted and out of range, suggesting the presence of a high concentration of unpaired electrons due to oxygen vacancies. Consequently, the mass loading for this sample was reduced to 1.5 mg to obtain a symmetrical and interpretable spectrum. The EPR signals for C–Ag/NiO _x @Mn _y O _z @800 °C and @400 °C appeared at g-values of 2.0 and 2.002, respectively–both characteristic of unpaired electrons resulting from oxygen vacancies.

In Figure S11, the EIS spectra for both precycling and postcycling analyses through Nyquist plots are presented. Notably, before cycling, EIS analysis of the AHSC device shows an R₁ value of 0.48 Ω, with minor increases in R₂ and R₃, recorded at 3.66 and 11.48 Ω, respectively (Table S3). This result underscores the stability of the C–Ag/NiO _x @Mn _y O _z @800 °C sample and the excellent electrochemical interaction between the electrode and electrolyte as well as between the electrode material and the NF current collector. The minimal semicircle observed in the high-frequency region signifies low charge transfer resistance, corroborated by the recorded value of 3.66 Ω obtained via equivalent circuit fitting.

In the low-frequency region, the nearly vertical profile of the spectrum suggests enhanced ion diffusion capability within the electrode material. As expected, the postcycling EIS spectrum displayed increased impedance values for all three components, with R₁, R₂, and R₃ measured at 0.75, 3.9, and 25.6 Ω, respectively. Such an increase is likely due to irreversible redox reactions, the potential formation of a passivation layer, and electrolyte depletion, which can impede the redox activity of the active species at the cathode. This trend emphasizes the impact of long-term cycling on the electrode’s electrochemical performance. As shown in Figure e, cyclic stability tests over 5000 charge–discharge cycles demonstrated remarkable durability, with AHSC retaining 82% of its initial capacitance after 5000 cycles. This outstanding performance confirms AHSC’s robust long-term operational stability and reliability.

In Figure f, Cs, energy density, and power density values achieved at various current densities are summarized. It is significant that AHSC exhibited an impressive energy density of 37.53 Wh kg^–1 and a power density of 2251.8 W kg^–1 at 3 A g^–1. The electrochemical performance of the C–Ag/NiO _x @Mn _y O _z @800 °C/HHC composite surpassed several previously reported values for manganese oxide-based materials, indicating its superior performance and potential for practical applications.

3.7. Ex-situ and Operando XAS Analysis of C–Ag/NiO _x @Mn _y O _z @800 °C Electrode

In Figure S12a, the Ag L-edge XANES spectra provide critical insights into the electronic state and local structural environment of Ag within the C–Ag/NiO _x @Mn _y O _z @400 °C and C–Ag/NiO _x @Mn _y O _z @800 °C samples, compared to reference materials: Ag foil, Ag₂O, and AgO. The spectra of both samples display pre-edge and postedge features analogous to those of Ag foil, indicating the predominance of metallic silver (Ag⁰) within the composites. In both samples, however, a subtle shift of the absorption edge toward lower energy, relative to the Ag foil, was detected. In the local electronic environment of Ag atoms, this shift indicates minor perturbations, possibly induced by structural modifications in the Ag crystal lattice.

To further elucidate the oxidation state and phase composition of Ag, linear combination fitting (LCF) of the XANES spectra for the C–Ag/NiO _x @Mn _y O _z @400 °C and C–Ag/NiO _x @Mn _y O _z @800 °C samples was carried out over an energy range of 3331–3381 eV, using ATHENA software. The fitting procedure, conducted without floating the E₀ parameter and constrained to a sum of one, yielded high-quality fits with R-factors below 0.05. LCF analysis indicated that the C–Ag/NiO _x @Mn _y O _z @400 °C sample comprised ∼81% Ag⁰ (from Ag foil) and 19% Ag₂O. In contrast, the C–Ag/NiO _x @Mn _y O _z @800 °C sample consisted of 74% Ag⁰ and 26% AgO. This outcome demonstrates the subtle variation of Ag between the two samples attributable to different annealing conditions.

Initially, the hydrothermal synthesis yielded a composite predominantly featuring the Mn₃O₄ phase, as confirmed by XRD analysis. Subsequent thermal treatments induced controlled molecular transitions that were systematically examined. At 400 °C, the material transformed from Mn₃O₄ into Mn₂O₃, as revealed by the ex-situ XAS spectra at the Mn K-edge region. This intermediate phase is crucial for facilitating subsequent reduction and enhancing redox-active sites.

In Figure b, ex-situ Mn K-edge XANES analysis of the C–Ag/NiO _x @Mn _y O _z @800 °C sample reveals spectral features closely resembling those of MnO, with pre-edge and postedge characteristics as well as excitation energy (∼6554.3 eV), mirroring the standard MnO reference. This outcome signifies that at 800 °C, the desired MnO phase became predominant, with a quantified phase composition of 95.6% determined via linear combination fitting (LCF) of the XANES spectra for C–Ag/NiO _x @Mn _y O _z @800 °C. At this temperature, the MnO phase was also clearly detected in the XRD pattern, showing a distinct diffraction peak of MnO at a 2θ angle of 40.5°, which corresponds to the (200) plane (Figure b). This progressive thermal reduction process not only enables precise control over the redox-active species but also facilitates the formation of metallic Ni⁰, thus establishing the synergistic redox interplay between Mn²⁺/Mn³⁺ and Ni⁰/Ni²⁺. This synergy is essential in amplifying electrochemical performance, as it promotes improved charge storage capacity and cycling stability of the electrode material.

8.

Ex-situ XAS analysis and operando charge-storage mechanism assessment via synchrotron light: (a) Assemblage and operando XAS study of AHSC, (b) Normalized ex-situ Mn K-edge XANES spectra, (c) Mn K-edge XANES spectra recorded via operando XAS technique during 7 complete charging cycles, (d) Normalized ex-situ Ni K-edge XANES spectra, (e) Ni K-edge XANES spectra recorded via operando XAS technique during 7 complete charging cycles, (f) FT Mn K-edge EXAFS spectra of C–Ag/NiO _x @Mn _y O _z @800 °C, and (g) FT Ni K-edge EXAFS spectra of C–Ag/NiO _x @Mn _y O _z @800 °C.

These results deviate slightly from that of the C–Ag/NiO _x @Mn _y O _z @400 °C sample; the Mn K-edge spectrum mimics the patterns of the standard Mn₂O₃ reference (Figure S12b). Moreover, the Ni K-edge XANES spectra of the C–Ag/NiO _x @Mn _y O _z @800 °C sample exhibit remarkable similarity to the spectral patterns of the reference Ni foil, indicating the predominance of metallic Ni within the sample. This observation aligns well with the XRD findings, which display a prominent peak of Ni at a 2θ angle of 57.1°, corresponding to the (200) plane of metallic Ni. It is noted that the C–Ag/NiO _x @Mn _y O _z @400 °C sample exhibits an XANES spectral pattern identical to that of the NiO reference (Figure S12c).

Fourier transform (FT) EXAFS analysis of Mn K-edge spectra further supports the presence of MnO and metallic Ni within the C–Ag/NiO _x @Mn _y O _z @800 °C sample, with a recorded Mn–O bond length of 1.68 Å, consistent with the MnO reference (Figure f). Similarly, the FT EXAFS of Ni K-edge spectra highlight congruent patterns between the C–Ag/NiO _x @Mn _y O _z @800 °C sample and the Ni foil reference, both displaying a Ni–TM bond length of 2.17 Å.

To elucidate the charge-storage mechanism and redox activity of both Mn and Ni components within the cathode material of the AHSC, synchrotron operando XAS was conducted at Beamline 2.2 of the Synchrotron Light Research Institute (SLRI), Thailand (Figure S13). Measurements were conducted using a custom-designed cell optimized for transmission XAS mode (Figure a). To ensure the acquisition of high-quality XAS data, the cathode material loading was meticulously controlled at ∼ 3.0 mg cm^–2. The collected XAS data included XANES spectra, providing critical insights into the oxidation states and local electronic structure of the active materials. Data processing and interpretation were carried out using Athena–Demeter (version 0.9.26): a sophisticated software suite for XAS analysis.

The operando XAS technique was employed to investigate the structural evolution of Mn and Ni within the C–Ag/NiO _x @Mn _y O _z @800 °C samples under real-cell conditions. Mn K-edge XANES spectra reveal a noticeable increase in absorption edge energy upon charging, corresponding to the oxidation state evolution of Mn (Figure c). When compared to standard references of MnO, Mn₂O₃, and MnO₂, the absorption K-edge energy of C–Ag/NiO _x @Mn _y O _z @800 °C, during charging, overlaps with that of Mn₂O₃. This observation suggests a valence transition of Mn from ∼+2 to +3, indicating that the charge-storage redox mechanism is primarily governed by the valence evolution of Mn within the solid phase.

While the white-line region in the Mn K-edge spectrum does not directly provide information about Mn 3d orbitals, its intensity serves as a useful indicator of 3d orbital activity, particularly when Mn is coordinated with O atoms. During charging, the increased white-line intensity points to a higher number of unoccupied Mn 3d orbitals due to electron transfer to electrolyte species. This observation supports the gradual formation of an intercalated state where mobile ions from the electrolyte are incorporated into the C–Ag/NiO _x @Mn _y O _z @800 °C layers. In general, the pre-edge peak arises from both the quadrupole transition and the dipole transition associated with d–p hybridization. In the case of Ni foil, the pre-edge peak is primarily attributed to the direct dipole transition of Ni 1s → O 2p, resulting from strong hybridization, which is notably absent in the K-edge spectra of the reference NiO and C–Ag/NiO _x @Mn _y O _z @800 °C (Figure e). The distinct alignment of the white line between C–Ag/NiO _x @Mn _y O _z @800 °C and the standard NiO reference indicates the extensive accommodation of Ni 3d unoccupied orbitals by O atoms. This finding implies a gradual transition in Ni’s oxidation state from 0 to +2, resulting from a partial reduction reaction during the charging process.

3.8. Aging Effects on Electrochemical Properties of AHSC

To evaluate the aging effects on AHSC, postcycling electrochemical characterization was conducted after two months, including CV and GCD analyses. Results are presented in Figure S14a–S14d. Here, critical insights into the long-term electrochemical stability and performance retention of the device are provided, including Cs at different scan rates and current densities. The CV profile of AHSC is seen to retain its characteristic oxidation and reduction peaks, confirming the sustained redox activity of Mn²⁺/Mn³⁺ and Ni⁰/Ni²⁺ species within the electrode. As shown in Figure S14a, the CV curves maintain stable redox behavior across all scan rates, ranging from 5 to 100 mV s^–1, demonstrating the preserved electrochemical activity of the redox-active materials over time. The recorded Cs value of 51.2 F g^–1, at a scan rate of 5 mV s^–1, indicates the robustness of the electrode material (Figure S14b).

Likewise, in Figure S14c, the GCD curves retain their characteristic nontriangular profiles, further validating the supercapacitive behavior of AHSC. After two months, the device achieved a specific capacitance of 47 F g^–1 at a current density of 3 A g^–1 (Figure S14d). However, compared to the freshly prepared cell, the discharge time decreased across all current densities. This decline suggests a gradual decrease in energy storage capacity, which can be attributed to natural aging effects such as minor structural modifications within the electrode or electrolyte depletion over time. Overall, the electrochemical stability of AHSC remained well-preserved, reinforcing its potential for long-term applications.

4. Conclusions

In summary, we successfully developed a high-performance AHSC electrode by integrating a bimetallic Ag/Ni–MnO nanostructure with HHC support. The incorporation of HHC as a sustainable conductive scaffold significantly enhanced the composite’s surface area, electronic conductivity, and electrochemical stability. Comprehensive structural and spectroscopic analyses, applying FE-SEM, HR-TEM, XRD, XPS, and operando XAS revealed that the 800 °C thermal treatment induces a phase transformation of Mn₂O₃ to MnO and the partial reduction of NiO to metallic Ni as well as the generation of oxygen vacancies. These thermally driven changes give rise to a highly conductive framework and facilitate dual redox activity (Mn²⁺/Mn³⁺ and Ni⁰/Ni²⁺) within the electrode. As a result, the optimized Ag/NiO _x @Mn _y O _z @800 °C–HHC composite achieved exceptional Cs of 1770 F g^–1 at 5 mV s^–1 in three-electrode measurements. When assembled into an AHSC, using an HHC anode and the C–Ag/NiO _x @Mn _y O _z @800 °C cathode, the device delivered a Cs of ∼ 150.7 F g^–1 at 5 mV s^–1, an energy density of 37.53 Wh kg^–1 at a discharge current of 3 A g^–1, and a corresponding power density of 2251.8 W kg^–1. AHSC also demonstrated outstanding cycling stability, retaining about 82% of its initial capacitance after 5000 cycles. These findings provide a valuable pathway for designing next-generation energy storage devices with sustainable materials. In essence, they reinforce the potential of HHC-supported multicomponent composites for durable SCs.

The authors declare that all data supporting the findings of this study are available within the paper and its Supporting Information file.

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acsami.5c07064.

• Additional experimental details, comprising the descriptions of morphological characterization equipment, photos of the experimental setup, and additional experimental results (PDF)

The authors declare no competing financial interest.