In situ built nanoconfined Nb₂O₅ particles in a 3D interconnected Nb₂C MXene@rGO conductive framework for high-performance potassium-ion batteries

Abstract

Exploring novel anode materials with excellent electrochemical performance is of great significance for the development of potassium-ion batteries (KIBs). Here, a 3D interconnected Nb₂C/rGO conductive framework with in situ generated Nb₂O₅ nanoparticles (Nb₂O₅/Nb₂C/rGO) is successfully constructed by a simple one-step hydrothermal method and subsequent freeze-drying and annealing treatments. The unique structure formed by the intimate contact of the three components has a 3D conductive network, abundant pores and a large specific surface area, which can not only inhibit the self-restacking of Nb₂C nanosheets and the agglomeration of Nb₂O₅ nanoparticles and alleviate the volume change during the charge–discharge process, but also expose more active sites and provide unimpeded channels for the diffusion of K⁺ and infiltration of the electrolyte. Meanwhile, Nb₂O₅ nanoparticles produced by in situ oxidation of surface Nb₂C and the residual subsurface Nb₂C with a low potassium ion diffusion barrier and a high conductivity can shorten the diffusion distance and promote the diffusion kinetics of electrons/ions. Benefiting from the elaborately designed structure and synergistic effects of three different components, as an anode for KIBs, the resulting Nb₂O₅/Nb₂C/rGO exhibits a superior specific capacity of 410.6 mA h g⁻¹ after 100 cycles at 0.1 A g⁻¹, an exceptional rate performance of 159.0 mA h g⁻¹ at 5 A g⁻¹, a capacity retention of 88.8% and a coulombic efficiency over 99.8% after 1000 cycles at 2.0 A g⁻¹. Moreover, Nb₂O₅/Nb₂C/rGO also shows a good potassium storage performance in a KIB full-cell. Furthermore, the combined potassium storage mechanism of K⁺ intercalation/deintercalation is revealed by CV and in/ex situ analyses. This work can provide more meaningful guidance for the rational design and construction of anode materials for high-performance KIBs.

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1. Introduction

Potassium-ion batteries (KIBs) are promising candidates for replacing lithium-ion batteries (LIBs) as large-scale energy storage devices due to their low cost and K/K⁺ potential (−2.93 V), abundant potassium resources, fast ionic conductivity in the electrolyte and similar “rocking chair” mechanism.^1,2 Unfortunately, the widespread application of KIBs is still hindered owing to the large radius of K⁺ (1.38 Å), which will result in sluggish kinetics, low discharge capacity and large volume changes of the electrode material during repeated charge–discharge processes.³ Therefore, it remains a great challenge to develop suitable electrode materials with improved diffusion kinetics, a high specific discharge capacity and the capability of accommodating large volume changes.

Undoubtedly, the intrinsic phase structure of electrode materials has an important influence on the electrochemical performance.^3–6 Transition metal oxides (TMOs) are the research focus because of their advantages such as low cost, abundant resources, high theoretical specific capacity, safety and adjustable interlayer spacing.^7,8 In particular, Nb₂O₅ is an ideal electrode material with merits of a relatively small volume change, good electrochemical stability, nontoxicity and environmental friendliness.^9,10 In fact, Nb₂O₅ has been extensively studied in the field of ion batteries, showing considerable electrochemical performance.¹¹ Nevertheless, the intrinsic low electrical conductivity (∼3 × 10⁻⁶ S cm⁻¹) of Nb₂O₅ seriously limits its specific capacity and high rate capability.¹²

To address this problem, various strategies have been devised, such as controlling the morphology and structure, hybridization with highly conductive materials, heteroatom doping, etc.^13–15 For instance, Zhang et al. synthesized a Nb₂O₅/C@N-coated carbon nanocomposite through a combination of hydrothermal and multi-step annealing treatments.¹⁶ As an anode for sodium ion batteries (SIBs) and KIBs, the obtained material showed relatively high specific capacities of 201 and 175 mA h g⁻¹ after 100 cycles at a current density of 0.1 A g⁻¹, respectively. Li et al. designed an urchin-like Nb₂O₅ nanocomposite encapsulated in nitrogen-doped carbon sheaths (Nb₂O₅/C), showing an enhanced rate capability and cycling stability as an anode of KIBs.¹⁷ Therefore, nanostructured Nb₂O₅-based composites with a short transport distance of electrons/ions, high electronic conductivity and excellent electrolyte accessibility are powerful guarantees for obtaining superior electrochemical performance.

Transition metal carbides (MXenes), another kind of typical layered material, have shown great competitiveness in efficient energy storage due to their unique 2D structure, high conductivity, large interlayer spacing, low diffusion barrier of ions (Li⁺, Na⁺, K⁺ and Mg²⁺) and so on.¹⁸ However, the large number of surface functional groups and self-restacking properties of MXenes significantly affect their electrochemical performance, leading to a limited potassium storage capacity (50–150 mA h g⁻¹).¹⁹

In general, the electrochemical performance of MXenes can be effectively improved by changing the surface state of MXenes or constructing MXene-based nanocomposites.^20,21 For instance, Gogotsi and co-workers used TiO₂ and Ti₃C₂ as raw materials to synthesize TiO₂/Ti₃C₂ nanocomposites through van der Waals interactions, showing improved lithium storage performance.²² The tightly coupled Ti₃C₂/NiCoP nanocomposite reported by Yin et al. has a unique structure with high conductivity and a large specific surface area, thus exhibiting increased sodium storage performance.²³ Zhang et al. fabricated Ti₃C₂/MoS₂ with a low concentration of surface functional groups by hydrothermal and annealing methods, possessing a relatively high specific capacity as an anode of KIBs.²⁴ Besides, other nanomaterials, such as MnO₂, Sn, Bi, Si, SnO₂, etc., have also been introduced into MXenes to form the corresponding nanocomposites to improve the electrochemical performance.^18,25 Nevertheless, the practical application of the above-mentioned methods is limited by the complex and cumbersome steps. On the other hand, the specific capacity and cycling stability of the obtained MXene-based nanocomposites still need to be further improved to meet the practical requirements. Therefore, developing simple and effective methods for the preparation of well-structured MXene-based nanocomposites with good electrochemical properties is of great significance.

Recent studies have shown that MXenes can be oxidized in situ to form composites of TMO nanoparticles and MXenes nanosheets due to the presence of thermodynamically metastable transition metal atoms on the surface.^26,27 For example, Gogotsi et al. once prepared Nb₂O₅@Nb₄C₃T_x and TiO₂@Ti₃C₂T_x nanocomposites by simple one-step oxidation of the corresponding MXenes in a CO₂ atmosphere, exhibiting a rather good cycle performance in Li-ion capacitors.²⁶ Yuan's group constructed an accordion-like TiO₂/Ti₃C₂T_x composite by the hydrothermal treatment of Ti₃C₂T_x MXenes, exhibiting high specific capacities of 267 and 101 mA h g⁻¹ at a current density of 0.2 A g⁻¹ as anodes for LIBs and SIBs, respectively.²⁷ Therefore, the in situ oxidation of MXenes is an effective strategy to obtain TMO/MXene nanocomposites, which not only greatly suppresses the self-restacking of MXenes, but also fully utilizes the properties of TMO nanoparticles.

On the other hand, the introduction of 3D structures with abundant pores and a large specific surface area into MXenes, TMOs and their composites can further improve their specific capacity and diffusion kinetics.^28,29 For example, Tong et al. once designed 3D porous Nb₂O₅/rGO composites, showing reversible capacities of 225 and 160 mA h g⁻¹ as anodes of SIBs and KIBs at a current density of 0.2 A g⁻¹.²⁹ The sandwich-structured TiO₂/Ti₃C₂/rGO prepared by Zhang et al. shows high energy and power densities as an anode for Li-ion capacitors.²⁸ In addition, other 3D porous structures, such as Ti₃C₂ hollow spheres and Ti₃C₂/rGO aerogels, have also been reported.³⁰ These results indicate that well-designed 3D TMO/MXene/rGO nanocomposites with high conductivity, a small crystal size and an abundant porous structure can overcome the limitations of TMOs and MXenes to obtain high specific capacity and excellent cycling stability.

To date, Ti₃C₂-derived TiO₂/Ti₃C₂ nanocomposites have been intensively studied in different fields.^31–34 Compared with Ti₃C₂, Nb₂C has a relatively higher practical capacity (170 mA h g⁻¹vs. 100.0 mA h g⁻¹ for Ti₃C₂) and a much lower diffusion barrier of K⁺ (0.004 eV vs. 0.103 eV for Ti₃C₂).³⁵ Therefore, it is reasonable to expect that a better electrochemical performance can be obtained by using Nb₂C as the matrix material. For example, Nb₂C/rGO hybrid aerogels were synthesized by a low-temperature hydrothermal self-assembly method in our previous work, achieving an excellent potassium storage performance.³⁶ However, the presence of thermodynamically metastable transition metal atoms on the surface of Nb₂C imposes extremely strict requirements on the practical use and storage conditions of Nb₂C/rGO. This problem could be effectively solved by converting Nb₂C/rGO into Nb₂O₅/Nb₂C/rGO with good environmental stability through in situ oxidation. Meanwhile, the in situ generated Nb₂O₅ nanoparticles with a small volume effect can also expose more active sites and shorten the diffusion distance of ions/electrons. Undoubtedly, it can be expected that 3D Nb₂O₅/Nb₂C/rGO hybrid aerogels with the above advantages will have a potential application in KIBs. Nevertheless, until now there is no report on Nb₂O₅/Nb₂C/rGO as an anode for KIBs and the related energy storage mechanisms.

Based on the above considerations, in this work, the 3D interconnected Nb₂O₅/Nb₂C/rGO conductive framework is successfully synthesized by a simple one-step hydrothermal method and subsequent freeze-drying and annealing treatments. The morphology, microstructure and chemical composition of the 3D porous aerogel, Nb₂O₅ nanoparticles formed by in situ oxidation of Nb₂C and residual Nb₂C nanosheets are characterized in detail. The electrochemical properties and potassium ion diffusion behaviors of the prepared samples are investigated thoroughly, and the structure–property relationship is disclosed. Moreover, the synergistic potassium storage mechanism of K⁺ intercalation/deintercalation is also revealed explicitly through CV and in/ex situ analyses.

2. Experimental section

2.1 Materials

Tetramethylammonium hydroxide ((CH₃)₄NOH, TMAOH, AR), ethanol (C₂H₅OH, AR) and hydrofluoric acid (HF, 49 wt%) were obtained from Sinopharm Chemical Reagent Co., Ltd. Nb₂AlC was purchased from Nanjing Mission New Materials Co., Ltd.

2.2 Synthesis of the Nb₂O₅/Nb₂C/rGO hybrid aerogel

Firstly, Nb₂CT_x and rGO suspensions were obtained according to our previous report.³⁶ Subsequently, the mixed suspensions of Nb₂CT_x (50.0 mL, 5.6 mg mL⁻¹), rGO (9.5 mL, 15.2 mg mL⁻¹) and deionized water (10.0 mL) were stirred for 1 h, sonicated for 1 h, and then transferred to a 100 mL reactor and hydrothermally treated at 180 °C for 12 h. After cooling to room temperature, the obtained Nb₂O₅/Nb₂CT_x/rGO hybrid hydrogel was washed three times with deionized water and then freeze-dried to obtain the Nb₂O₅/Nb₂CT_x/rGO hybrid aerogel. Finally, the Nb₂O₅/Nb₂CT_x/rGO aerogel was annealed at 600 °C for 3 h to obtain the Nb₂O₅/Nb₂C/rGO hybrid aerogel with high crystallinity and a low concentration of surface functional groups.

For comparison, the Nb₂O₅/Nb₂CT_x/rGO aerogel and Nb₂CT_x nanosheets were respectively soaked in deionized water containing H₂O₂ (10 mL) for 24 h, resulting in a fully oxidized Nb₂O₅/rGO hybrid aerogel and Nb₂O₅ nanoparticles. The other steps were consistent with those for the preparation of the Nb₂O₅/Nb₂C/rGO hybrid aerogel.

2.2 Characterization

The structure, surface morphology and composition of the as-prepared samples were identified using an X-ray diffractometer (XRD, Cu Kα, λ = 1.5418 Å, Philips X′Pert), a scanning electron microscope (SEM, JEOL JEM-6300F), a transmission electron microscope (TEM, JEOL JEM-2100F) and an X-ray photoelectron spectrometer (XPS, ESCALAB250Xi). The specific surface area and pore-size distribution were measured using a Micromeritics ASAP 2020 surface area analyzer. TG analyses (Netzsch STA449C) were performed to determine the content of rGO in the hybrid aerogel. Raman spectra were collected on an automatic laser confocal Raman microspectrometer (LabRAM Aramis).

2.3 Electrochemical measurements

The working electrode was prepared by mixing the active material (70 wt%), Super P (15 wt%) and polyvinylidene difluoride (PVDF, 15 wt%) in N-methyl-2-pyrrolidone (NMP) and stirring for about 24 h to obtain a uniform dispersion slurry. The mixed slurry was coated on copper foil and dried at 100 °C for 10 h. Subsequently, the obtained working electrode (0.7–1.0 mg cm⁻²) was assembled into a 2032 type coin cell, in which potassium foil was used as the counter electrode, a glass microfiber filter (Whatman, GF/D) was utilized as the separator and 0.8 M KPF₆ EC/DEC (1 : 1 by volume) was employed as the electrolyte (0.02 mL). The assembly of the KIB full cell is consistent with the above-mentioned half cells except that Prussian blue (PB, K_xFe[Fe(CN)₆]) was used as the cathode material. The cathode was prepared by coating a slurry mixture of Prussian blue (70 wt%), carbon black (15 wt%) and PVDF (15 wt%) in NMP onto aluminium foil. The designed amount of the cathode material was 1.1 times that of the anode material. Moreover, in order to better reveal the characteristics of the full cell, the anode and cathode electrodes were first discharged and charged for 3 cycles to form a stable SEI.

Galvanostatic charge–discharge (GCD) and galvanostatic intermittent titration technique (GITT) tests were performed on the LAND test instrument (CT2001A, Wuhan) within the voltage range of 0.01–3.0 V. Electrochemical impedance spectra (EIS, 0.01–100 000 Hz, 5 mV) and cyclic voltammetry (CV) curves were collected at the electrochemical workstation (CHI 660E, Chenhua).

3. Experimental section

3.1 Synthesis and characterization

Fig. 1 shows a schematic illustration for the preparation of the Nb₂O₅/Nb₂C/rGO aerogel. The suspension of Nb₂CT_x nanosheets was obtained by etching Nb₂AlC powder in HF, followed by exfoliation through intercalation and sonication.³⁶ To assemble Nb₂CT_x nanosheets into a 3D porous structure, GO nanosheets with a strong gelling ability were selected as the gelling agent due to the weak self-gelling ability of Nb₂CT_x. Benefiting from the similar hydrophilicity and negative charge of Nb₂CT_x and GO nanosheets, a stable homogeneously mixed suspension of Nb₂CT_x and GO nanosheets can be easily formed in water. During the subsequent hydrothermal reaction, the hydrophilic GO was reduced to rGO with a more hydrophobic surface, and the resulting rGO nanosheets formed a 3D interconnected framework structure. Meanwhile, Nb₂CT_x nanosheets were self-assembled onto the surface of rGO nanosheets in a face-to-face form via interfacial interaction, thus generating the Nb₂O₅/Nb₂CT_x/rGO hybrid hydrogel, in which Nb₂O₅ nanoparticles came from the in situ surface oxidation of Nb₂CT_x nanosheets at high temperature. Finally, the Nb₂O₅/Nb₂C/rGO hybrid aerogel was obtained through freeze-drying and annealing the Nb₂O₅/Nb₂CT_x/rGO hybrid hydrogel.

Fig. 1 A schematic illustration for the preparation of the Nb₂O₅/Nb₂C/rGO hybrid aerogel.

SEM and (HR)TEM were applied to observe the morphology of the prepared samples. As shown in Fig. S1,† the raw material Nb₂AlC has a densely-packed bulk structure, which can be transformed into multilayered Nb₂CT_x (m-Nb₂CT_x) with a typical accordion-like structure after HF etching. Fig. 2a depicts the TEM images of single- and few-layered Nb₂CT_x nanosheets obtained by exfoliating m-Nb₂CT_x with TMAOH intercalation and sonication. It can be observed that they have a smooth surface. Besides, the translucent state of Nb₂CT_x nanosheets indicates that they have an ultrathin layer thickness. By contrast, GO nanosheets not only have an ultrathin thickness, but also a wrinkled surface with large lateral dimensions (Fig. 2b). The large GO nanosheets can easily wrap small Nb₂CT_x nanosheets under the interfacial interaction.³⁷ It is verified by the Tyndall effect that Nb₂CT_x and GO nanosheets with abundant surface functional groups can be stably dispersed in water (the inset of Fig. 2c), which is the prerequisite for the preparation of electrode materials.³⁸Fig. 2c demonstrates the ultra-light property of the obtained Nb₂O₅/Nb₂C/rGO aerogel.

Fig. 2 TEM images of (a) Nb₂CT_x and (b) GO nanosheets. (c) Ultralight Nb₂O₅/Nb₂C/rGO aerogel can be placed on the top of a lotus leaf (inset shows the Tyndall effect of Nb₂CT_x and GO suspensions). (d–f) SEM images, (g and h) TEM images and (i) HRTEM image of Nb₂O₅/Nb₂C/rGO. (j) SEM image of Nb₂O₅/Nb₂C/rGO and the corresponding elemental mapping images of C, Nb and O.

As detailed in Fig. 2d, Nb₂O₅/Nb₂C/rGO has a randomly distributed porous structure. The pores are interconnected, with a size of 10 to 20 μm, which can not only promote the diffusion of ions/electrons, but also accommodate volume changes. From the magnified SEM image of the Nb₂O₅/Nb₂C/rGO aerogel (Fig. 2e and f), it can be seen that the aerogel is composed of wrinkled nanosheets and the nanosheets are interconnected, being beneficial for the transport of electrons.

As shown in Fig. S2 (ESI†), the fully oxidized Nb₂O₅/rGO sample exhibits a similar aerogel structure to Nb₂O₅/Nb₂C/rGO. However, a large amount of Nb₂O₅ nanoparticles almost completely cover rGO, and even lead to the blockage of the porous structure, being not conducive to the transport of electrolyte ions. On the other hand, in the Nb₂O₅ sample without rGO (Fig. S3, ESI†), severe agglomeration of Nb₂O₅ nanoparticles is observed. This further indicates that the addition of graphene effectively prevents the self-restacking of Nb₂C nanosheets and the agglomeration of Nb₂O₅ nanoparticles.

From Fig. 2g it is once again proved that Nb₂O₅/Nb₂C/rGO has a rich porous structure. It is also observed that Nb₂C nanosheets adhered tightly to rGO nanosheets in a face-to-face form (Fig. 2h). Being relative to the initial Nb₂C nanosheets (Fig. 2a), the smaller Nb₂C nanosheets in Nb₂O₅/Nb₂C/rGO indicate that Nb₂C nanosheets undergo structural degradation and partial oxidation under high temperature conditions. As a result, it can be observed that Nb₂O₅ nanoparticles with an average diameter of 20–40 nm are tightly and homogeneously anchored on the surface of Nb₂C and/or rGO nanosheets (Fig. 2h), thus facilitating the electron transport of Nb₂O₅ nanoparticles through the interconnected highly conductive aerogel structure.

In the HRTEM image (Fig. 2i), the clear lattice fringes indicate the high crystallinity of Nb₂O₅/Nb₂C/rGO. The interlayer spacings of 10.4 and 3.9 Å match well with the (002) plane of Nb₂C and the (001) plane of orthorhombic Nb₂O₅, respectively.^29,35,39 Besides, the HRTEM image also shows good integration of Nb₂C nanosheets and Nb₂O₅ nanoparticles within the amorphous rGO network. Generally speaking, the diffusion barrier of the electrolyte ions in the (001) plane of Nb₂O₅ is much lower than that of other directions.³⁹ Therefore, the combination of nano-sized Nb₂O₅ with a preferentially exposed (001) plane is expected to provide a high energy storage performance. In addition, the elemental mapping images in Fig. 2j show the uniform distribution of C, Nb and O in the Nb₂O₅/Nb₂C/rGO aerogel, proving that Nb₂C and Nb₂O₅ are evenly adhered to the rGO nanosheets.

To investigate the crystalline structure of the as-prepared samples, XRD was employed. As displayed in Fig. S4,† Nb₂AlC has typical characteristic diffraction peaks (JCPDS No. 30-0033).³⁵ As expected, the diffraction peaks of Nb₂AlC undergo obvious changes after HF treatment. For example, the (002) diffraction peak of the resulting m-Nb₂CT_x shifts to a low 2 θ angle, which indicates the successful removal of Al layers from Nb₂AlC.³⁵

Fig. 3a presents XRD patterns of the as-prepared samples. The obtained Nb₂O₅/Nb₂CT_x/rGO aerogel has a distinct diffraction peak of rGO at 2 θ = 26.0°, indicating that GO is successfully reduced to rGO.^40,41 Nevertheless, no diffraction peaks of Nb₂O₅ and Nb₂CT_x are detected in the Nb₂O₅/Nb₂CT_x/rGO aerogel without annealing. It could be related to the amorphous nature of Nb₂O₅ and highly dispersed ultrathin Nb₂CT_x nanosheets.^42–44 After annealing at 600 °C, new peaks centered at 22.63°, 28.35°, 36.59°, 46.14°, 50.89° and 55.15° appear in the obtained Nb₂O₅/Nb₂C/rGO aerogel. These peaks correspond to the (001), (180), (181), (002), (380) and (182) crystal planes of Nb₂O₅, respectively, proving the formation of well-crystalline orthorhombic Nb₂O₅ (T-Nb₂O₅, JCPDS No. 30-0873).^15,29,42 In addition, Nb₂O₅ and Nb₂O₅/rGO also exhibit typical diffraction peaks of T-Nb₂O₅, suggesting that Nb₂CT_x in the raw materials of these two samples is successfully oxidized to Nb₂O₅.

Fig. 3 (a) XRD patterns of Nb₂O₅/Nb₂CT_x/rGO, Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅, (b) Raman spectra of GO, Nb₂O₅/Nb₂C/rGO and Nb₂O₅/rGO. (c) XPS survey spectra of Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅. Refined Nb 3d XPS spectra of Nb₂O₅/Nb₂C/rGO (d) before and (e) after Ar⁺ ion sputtering. Refined XPS spectra of (f) C 1s and (g) O 1s, and (h) nitrogen adsorption and desorption isotherms and (i) pore-size distribution curve of Nb₂O₅/Nb₂C/rGO.

To determine the structural characteristics of the prepared samples, Raman spectra were recorded. As shown in Fig. 3b, two broad peaks are observed at 1350 and 1590 cm⁻¹, which can be indexed to the D-band (disordered carbon) and the G-band (ordered crystalline graphite), respectively.^29,45 In general, the intensity ratio of the D- to G-band (I_D/I_G) can be used to measure the degree of disorder and defects in the prepared materials.³⁵ The increased I_D/I_G value from 0.95 for GO to 1.03 for Nb₂O₅/Nb₂C/rGO and 1.11 for Nb₂O₅/rGO indicates that GO has been successfully reduced in the hydrothermal process.⁴⁶ Compared with Nb₂O₅/rGO, the higher graphitization degree of Nb₂O₅/Nb₂C/rGO is beneficial for improving the electronic conductivity.³⁵ Furthermore, according to TGA (Fig. S5, ESI†), the content of rGO in the Nb₂O₅/Nb₂C/rGO sample is calculated to be 27.0 wt%.

To reveal the surface and near-surface elemental compositions of Nb₂O₅, Nb₂O₅/rGO and Nb₂O₅/Nb₂C/rGO samples, XPS tests were carried out. As shown in Fig. 3c, the survey spectra prove the existence of C, O and Nb elements in all prepared samples. Compared to Nb₂O₅, the significant increase of the C element in Nb₂O₅/rGO and Nb₂O₅/Nb₂C/rGO indicates the successful introduction of rGO. The absence of the F element indicates that it is successfully removed during the synthetic process, which can greatly reduce side reactions and facilitate the transport of electrolyte ions.⁴⁷

As illustrated in Fig. 3d, the high-resolution Nb 3d XPS spectrum for the surface of the Nb₂O₅/Nb₂C/rGO sample shows two peaks at 210.0 eV (Nb 3d_3/2) and 207.2 eV (Nb 3d_5/2), consistent with the binding energies of Nb₂O₅.²⁹ This suggests that a substantial amount of Nb₂C on the surface of Nb₂C/rGO samples is converted into Nb₂O₅, being in agreement with the above XRD results. The surface layer of the Nb₂O₅/Nb₂C/rGO sample is removed by Ar⁺ sputtering for 300 s, and it is obvious that the high-resolution Nb 3d XPS spectrum for the near-surface of Nb₂O₅/Nb₂C/rGO shows a great difference. Specifically, the near-surface of Nb₂O₅/Nb₂C/rGO exhibits four peaks at 210.0, 207.2, 206.3 and 204.0 eV, respectively (Fig. 3e). The first two peaks are assigned to Nb₂O₅, while the last two peaks can be indexed to C–Nb bonds.^29,35 This indicates that the interior of the Nb₂O₅/Nb₂C/rGO sample is closely related to Nb₂C. That is, after hydrothermal and annealing treatments, the surface Nb₂C is oxidized to Nb₂O₅, while the pristine Nb₂C is still present in the interior bulk of the material. In contrast, Nb 3d spectra of Nb₂O₅/rGO and Nb₂O₅ (Fig. S6, ESI†) show that only Nb₂O₅ exists on the surface and near-surface, which further proves that Nb₂C is completely oxidized with the assistance of H₂O₂.

As shown in Fig. 3f, the C 1s spectrum of Nb₂O₅/Nb₂C/rGO can be deconvoluted into three peaks centered at 288.6 eV, 286.5 eV and 284.8 eV, which are assigned to O–C=O, C–O and C–C, respectively.³⁵ As displayed in Fig. 3g, the O 1s spectrum can only be divided into two peaks, namely the C–O peak (532.2 eV) and the Nb–O peak (530.8 eV).³⁵ These results suggest that –OH is successfully removed after the annealing process, which is in favor of the transport of electrolyte ions.⁴⁷

The specific surface area and porous nature of Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅ samples were examined using N₂ adsorption–desorption isotherms. As displayed in Fig. 3h, typical type I and IV isotherms indicate that there are abundant micropores and mesopores in Nb₂O₅/Nb₂C/rGO.^16,47 Obviously, Nb₂O₅/Nb₂C/rGO exhibits a larger specific surface area of 197.1 m² g⁻¹ compared with that of Nb₂O₅/rGO (127.6 m² g⁻¹) and Nb₂O₅ (70.8 m² g⁻¹) (Fig. S7a, ESI†). The largest specific surface area of Nb₂O₅/Nb₂C/rGO may be attributed to the 3D interconnected porous structure of the aerogel and the well-dispersed Nb₂O₅ nanoparticles on the surface of Nb₂C and/or rGO. The significant decrease in the specific surface area of the Nb₂O₅/rGO sample is related to the large number of aggregated Nb₂O₅ nanoparticles generated by the complete oxidation of Nb₂C, which severely blocked the porous structure of the resulting aerogel. The smallest specific surface area of Nb₂O₅ is caused by the serious agglomeration of Nb₂O₅ nanoparticles and the self-restacking of Nb₂C. As shown in Fig. 3i and Fig. S7b (ESI†), the Nb₂O₅/Nb₂C/rGO sample contains abundant mesopores that may originate from the 3D porous structure of the aerogel, while both Nb₂O₅/rGO and Nb₂O₅ lack this important feature. In addition, Nb₂O₅/Nb₂C/rGO also has the largest total pore-volume (0.33 cm³ g⁻¹vs. 0.17 cm³ g⁻¹ for Nb₂O₅/rGO and 0.16 cm³ g⁻¹ for Nb₂O₅). Benefiting from the abundant mesopores, large pore-volume and high specific surface area of Nb₂O₅/Nb₂C/rGO, it can ensure sufficient contact area between the active material and electrolyte and promote the migration of electrolyte ions, being conducive to the improvement of specific capacity and rate capability.

3.2 Potassium storage performance

To assess the advantages of the as-prepared samples, KIBs were assembled and tested under different conditions. Fig. 4a is a schematic illustration of the working principle of KIBs. K⁺ ions will be intercalated/adsorbed to the host material in the initial discharging process along with the gradual decrease of the voltage of KIBs. During the subsequent charging process, the intercalated/adsorbed K⁺ ions will be deintercalated/desorbed from the host material and move to the counter electrode simultaneously.¹

Fig. 4 (a) A schematic illustration for the working principle of the assembled KIBs, (b) CV curves of Nb₂O₅/Nb₂C/rGO at 0.1 mV s⁻¹, (c) GCD curves of Nb₂O₅/Nb₂C/rGO at the 1st, 10th, 30th and 50th cycles. The cycling performance at current densities of (d) 0.1 A g⁻¹ and (e) 0.5 A g⁻¹, and (f) rate performance (the numbers represent current densities with the unit of A g⁻¹) of Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅. (g) Long-term cycling stability of Nb₂O₅/Nb₂C/rGO at 2.0 A g⁻¹ and (h) a comparison of electrochemical performance between Nb₂O₅/Nb₂C/rGO and previously reported Nb₂O₅- and MXene-based anode materials.

Fig. 4b presents typical CV curves of Nb₂O₅/Nb₂C/rGO at a scan rate of 0.1 mV s⁻¹. A distinct cathodic peak appears at 0.50 V in the first cycle and disappears during the subsequent cycles, which can be attributed to the formation of a SEI film.^23,35 Another cathodic peak at 0.76 V is derived from the intercalation of K⁺ into Nb₂C (Nb₂C + xK⁺ + xe⁻ → K_xNb₂C). Moreover, the anodic peak at 0.58 V corresponds to the deintercalation of K⁺ from Nb₂O₅/Nb₂C/rGO during the charging process.²³ Encouragingly, after the first cycle, the almost completely overlapped CV curves suggest that the electrochemical reaction of Nb₂O₅/Nb₂C/rGO is highly reversible during potassiation/depotassiation processes.

Fig. 4c shows GCD profiles of Nb₂O₅/Nb₂C/rGO at a current density of 0.05 A g⁻¹. There is no obvious charge–discharge plateau throughout the potassiation/depotassiation process, which is consistent with previous reports.^15,16,35 A high specific discharge capacity of 473.0 mA h g⁻¹ is achieved in the first cycle with a corresponding coulombic efficiency (CE) of 51.3%. The large capacity loss in the first cycle is quite normal.⁴⁸ On the one hand, it is caused by the formation of a SEI, and on the other hand, it may be ascribed to the side reactions of K⁺ with the –O group and carbon defects on rGO nanosheets and the irreversible insertion/adsorption of K⁺ in the aerogel with a porous structure.⁴⁹ However, the CE of Nb₂O₅/Nb₂C/rGO rapidly reaches above 98.0% in the 10th cycle and remains stable during the subsequent cycles, implying the good reversibility of Nb₂O₅/Nb₂C/rGO.^50,51

Fig. 4d displays the specific discharge capacities and CEs of Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅ at 0.1 A g⁻¹. Notably, the initial CE of Nb₂O₅/Nb₂C/rGO (50.9%) is lower than those of Nb₂O₅ (55.2%) and Nb₂O₅/rGO (62.5%), which may be caused by the larger specific surface area of Nb₂O₅/Nb₂C/rGO. However, Nb₂O₅/Nb₂C/rGO renders a high reversible capacity of 410.6 mA h g⁻¹ after 100 cycles, far exceeding 287.1 mA h g⁻¹ of Nb₂O₅/rGO and 195.4 mA h g⁻¹ of Nb₂O₅ in this paper and 352.3 mA h·g⁻¹ of Nb₂C/rGO-2 in our previous work.³⁶ In addition, Nb₂O₅/Nb₂C/rGO also exhibits good cycling stability. As shown in Fig. 4e, the specific capacity of Nb₂O₅/Nb₂C/rGO remains at 304.8 mA h g⁻¹ after 200 cycles at a current density of 0.5 A g⁻¹, the corresponding capacity retention is 90.3% relative to the 5th cycle. By comparison, the capacities of Nb₂O₅/rGO and Nb₂O₅ decrease significantly to 200.3 and 80.4 mA h g⁻¹, and the corresponding capacity retentions are 74.7 and 62.7% relative to the 5th cycle, respectively. Overall, Nb₂O₅/Nb₂C/rGO has the best specific discharge capacity and cycling stability.

The rate capabilities of Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅ were also tested to evaluate the practical application potential. Fig. S8 (ESI†) reveals the GCD curves of Nb₂O₅/Nb₂C/rGO at different current densities. It can be seen that the shape of GCD curves does not notably change with the increase of current density, except for the gradual decrease of specific capacity, which further proves the reversibility of the electrochemical reaction. As shown in Fig. 4f, Nb₂O₅/Nb₂C/rGO presents a high reversible capacity of 410.2 mA h g⁻¹ at 0.1 A g⁻¹. As the current density is gradually increased to 0.2, 0.5, 1.0 and 2.0 A g⁻¹, Nb₂O₅/Nb₂C/rGO delivers good capacity retention with specific discharge capacities of 354.0, 318.7, 284.4 and 230.5 mA h g⁻¹, respectively. Even at a high current density of 5.0 A g⁻¹, a discharge capacity of 159.0 mA h g⁻¹ can still be obtained. More strikingly, Nb₂O₅/Nb₂C/rGO still provides a high specific capacity of 329.5 mA h g⁻¹ when the current density is restored to 0.1 A g⁻¹. However, for Nb₂O₅/rGO, it just displays the discharge capacities of 320.1, 265.0, 213.1, 184.0, 139.6 and 69.4 mA h g⁻¹ at 0.1, 0.2, 0.5, 1.0, 2.0 and 5.0 A g⁻¹. Even worse, Nb₂O₅ has almost lost its capacity when the current density is increased to 5 A g⁻¹. The discharge capacity of Nb₂O₅/rGO and Nb₂O₅ degrades sharply with the current density, which is mainly caused by the limited transfer ability of ions and/or electrons at a high current density.

Furthermore, Nb₂O₅/Nb₂C/rGO also exhibits promising long-term cycling stability, as illustrated in Fig. 4g. Distinctly, Nb₂O₅/Nb₂C/rGO realizes an ultra-stable capacity of 208.9 mA h g⁻¹ over 1000 cycles at 2 A g⁻¹. Not only that, it also has a high capacity retention of 88.8% and CE of over 99.5%. It is worth mentioning that the preeminent electrochemical performance of Nb₂O₅/Nb₂C/rGO as an anode of KIBs in this work exceeds those of all Nb₂O₅-based materials, as well as most anode materials in the literature (Fig. 4h and Table S1†).^{15,16,29,52–54}

3.3 Potassium storage mechanism

To gain a deep insight into the excellent K⁺ storage performance of Nb₂O₅/Nb₂C/rGO, a series of CV tests were carried out. Fig. 5a depicts CV curves at different scan rates (0.1–1.0 mV s⁻¹). According to the previous literature, the overall specific capacity usually comes from the contribution of capacitive- and diffusion-controlled behaviors. The capacitive-controlled behavior can be qualitatively analyzed by the value of b in the following equation:^15,16

Log (i) = blog (v) + log (a) (0.5 ≤ b ≤ 1) (1)

where i and v are current (A) and scan rate (mV s⁻¹), and a and b are two changeable parameters. The value of b can be achieved through the slope of log (i) versus log (v). Typically, the value of b varies between 0.5 for diffusion-controlled behavior and 1.0 for capacitive-controlled behavior. As shown in Fig. 5b, the b-values of potassiation and depotassiation peaks of Nb₂O₅/Nb₂C/rGO are 0.74 and 0.87, respectively, indicating that the electrochemical reaction combines capacitive and diffusion behaviors.

Fig. 5 (a) CV curves at various scan rates and (b) log (i) – log (v) plots for Nb₂O₅/Nb₂C/rGO, (c) contribution of diffusion and capacitive behaviors at various scan rates for Nb₂O₅/Nb₂C/rGO. (d) GCD curves and in situ XRD patterns of Nb₂O₅/Nb₂C/rGO. Ex situ XPS spectra of (e) Nb 3d and (f) K 2p for Nb₂O₅/Nb₂C/rGO at selected potentials.

The contribution of diffusion and capacitive behaviors in the total specific capacity can also be quantified using the following equation:^15,16

i(v) = k₁v + k₂v^1/2 (2)

where k₁ and k₂ are two changeable parameters. As displayed in Fig. 5c, as the scan rate is increased, the diffusion behavior is suppressed, while the contribution of the capacitive behavior gradually increases. Specifically, the capacitive contribution is increased from 34.0% at 0.1 mV s⁻¹ to 57.7% at 0.6 mV s⁻¹, and finally reaches a maximum value of 64.9% at 1 mV s⁻¹ (Fig. S9, ESI†). The fast diffusion kinetics of capacitive behavior enables Nb₂O₅/Nb₂C/rGO to achieve an outstanding rate performance.

To further illustrate the structural evolution of Nb₂O₅/Nb₂C/rGO during potassiation/depotassiation, in situ XRD and ex situ XPS were employed at selected potentials. As shown in Fig. 5d, along with the proceeding of the potassiation process, the typical (001) diffraction peak of Nb₂O₅ gradually shifts to a low angle. Meanwhile, the intensity reduces greatly. This indicates that K⁺ is successfully intercalated into the Nb₂O₅ lattice, and the intercalation of K⁺ may lead to a decrease in crystallinity.^16,17 In contrast, the (001) diffraction peak returns to the original position during the subsequent charging process, corresponding to the reversible deintercalation of K⁺ from the Nb₂O₅ lattice.¹⁷ In addition, a new peak appears at 2 θ = 24.0° (KC₈) upon discharging to 0.01 V and disappears during the subsequent charging process, indicating that K⁺ can also be reversibly intercalated/deintercalated in rGO.⁴⁹

Combined CV (Fig. 4b), HRTEM (Fig. S10, ESI†) and in situ XRD results, it can be suggested that the following reactions occur in the potassiation/depotassiation process:

Nb₂C + xK⁺ + xe⁻ ↔ K_xNb₂C (3)

Nb₂O₅ + yK⁺ + ye⁻ ↔ K_yNb₂O₅ (4)

8C + zK⁺ + ze⁻ ↔ K_zC₈ (5)

Fig. 5e and Fig. S11 (ESI†) show the Nb 3d high-resolution XPS spectra of Nb₂O₅/Nb₂C/rGO at different states. The binding energies of Nb 3d_3/2 and Nb 3d_5/2 are shifted from 210.0 and 207.2 eV (fresh electrode, spectrum I in Fig. 5e) to lower binding energies of 209.7 and 206.9 eV (fully potassiated electrode, spectrum II), indicating the formation of Nb⁴⁺.²⁹ After charging to 3.0 V (spectrum III), the two peaks almost return to their original positions, revealing a reversible electrochemical redox reaction of Nb.^15,55

The K 2p high-resolution XPS spectra were also recorded at different states. As shown in Fig. 5f, compared with the fresh electrode (spectrum I), two peaks can be observed at 295.9 eV (K 2p_1/2) and 293.0 eV (K 2p_3/2) for the fully potassiated electrode (spectrum II), proving the intercalation of K⁺.^49,56 On the other hand, the intensity of K 2p peak is significantly weakened in the fully depotassiated electrode (spectrum III), once again demonstrating the reversible deintercalation of K⁺ during the cycling.⁵⁷ The remaining K 2p peaks may be attributed to different potassium salts in the SEI film.⁵⁸

3.4 Diffusion kinetics

To investigate the diffusion capability of K⁺ inside the electrode material, the K⁺ diffusion coefficients (D_K+s) of Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅ were obtained through GITT. As shown in Fig. 6a, the GITT curves were recorded at a current density of 0.1 A g⁻¹ with a galvanostatic pulse of 10 min and a relaxation of 30 min. D_K⁺ values of the prepared samples were calculated based on the GITT curves using the following equation:⁵⁹where τ is the pulse time (s); m, M, A and V_m are the mass (g), molar mass (g mol⁻¹), geometric area (cm²) and molar volume (cm³ mol⁻¹) of the electrode; and ΔE_s and ΔE_τ are the variations of the steady state voltage and total voltage (see the inset of Fig. 6a). The calculated dynamic (D_K⁺) values during potassiation and depotassiation are plotted in Fig. 6b. It is obvious that the (D_K⁺) value of Nb₂O₅/Nb₂C/rGO is larger than those of Nb₂O₅/rGO and Nb₂O₅ during the whole potassiation/depotassiation process. For example, the (D_K⁺) values of Nb₂O₅/Nb₂C/rGO are 1.1 × 10⁻¹¹, 3.5 × 10⁻¹² and 2.3 × 10⁻¹² cm² s⁻¹ at 1.0, 0.5 and 0.1 V, respectively, while the corresponding (D_K⁺) values of Nb₂O₅/rGO and Nb₂O₅ are 8.6 × 10⁻¹², 2.7 × 10⁻¹² and 1.5 × 10⁻¹², and 7.4 × 10⁻¹², 2.0 × 10⁻¹² and 4.1 × 10⁻¹³ cm² s⁻¹, respectively. These results well explain the superior rate capability of Nb₂O₅/Nb₂C/rGO, confirming the positive effect of rational structure design on promoting the diffusion of K⁺.

Fig. 6 (a) GITT curves (inset shows the single-step GITT titration of Nb₂O₅/Nb₂C/rGO) and (b) diffusion coefficients of K⁺ for Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅.

The improved diffusion kinetics of ions/electrons was further verified by EIS. As shown in Fig. S12,† the EIS spectra of Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅ have similar compositions. It can be fitted using an equivalent circuit including solution resistance (R_s), charge transfer resistance (R_ct) and Warburg impedance (W_s).⁴⁰ All samples exhibit similar R_s, but their R_ct values have a great difference. Compared with 875 Ω of Nb₂O₅/rGO and 1160 Ω of Nb₂O₅, Nb₂O₅/Nb₂C/rGO displays the smallest R_ct value of 720 Ω, suggesting the fastest charge transfer in Nb₂O₅/Nb₂C/rGO. This may be ascribed to the tight coupling of Nb₂C, Nb₂O₅ and rGO, and the high conductivity of the residual Nb₂C and rGO. In addition, the largest linear slope of Nb₂O₅/Nb₂C/rGO in the low frequency region implies the smallest W_s, i.e., the smallest diffusion resistance of K⁺ inside the Nb₂O₅/Nb₂C/rGO electrode, being in good agreement with the above GITT results.³⁵ Hence, the optimized electron/ion diffusion capability helps Nb₂O₅/Nb₂C/rGO achieve preeminent rate performance.

3.5 Structural stability

To intuitively verify the structural stability of Nb₂O₅/Nb₂C/rGO, Nb₂O₅/rGO and Nb₂O₅, the SEM images after 100 cycles at a current density of 0.1 A g⁻¹ were recorded. As illustrated in Fig. S13 (ESI†), the structural integrities of Nb₂O₅/rGO and Nb₂O₅ electrodes are severely damaged. For example, the porous structure of the Nb₂O₅/rGO electrode undergoes partial collapse and the agglomeration of Nb₂O₅ nanoparticles further blocks the pores; and the Nb₂O₅ electrode shows obvious cracks. However, compared with the initial morphology (Fig. 2d–f), the 3D porous structure of the Nb₂O₅/Nb₂C/rGO electrode as well as the Nb₂O₅ nanoparticles are basically preserved, highlighting the structural stability of Nb₂O₅/Nb₂C/rGO.

3.6 Electrochemical performance of the KIB full cell

Nb₂O₅/Nb₂C/rGO shows excellent electrochemical performance in the above-mentioned half-cells. To demonstrate the feasibility of Nb₂O₅/Nb₂C/rGO for practical applications, a full-cell was assembled by using Nb₂O₅/Nb₂C/rGO as the anode and Prussian blue (PB, K_xFe[Fe(CN)₆]) as the cathode. The GCD curves of the full cell at a current density of 0.1 A g⁻¹ are presented in Fig. S14b (ESI†), the charge and discharge capacities are 160.0 and 86.8 mA h g⁻¹ (based on the weight of the cathode), respectively. Moreover, the assembled full-cell also exhibits good cycling stability. As shown in Fig. S14c (ESI†), the specific capacity of Nb₂O₅/Nb₂C/rGO//PB remains at 75.1 mA h g⁻¹ after 100 cycles at a current density of 0.1 A g⁻¹, and the corresponding capacity retention is 86.5%. Interestingly, the fully charged full-cell can successfully light up 34 light-emitting diodes (LEDs), proving that the Nb₂O₅/Nb₂C/rGO anode has a good prospect for practical applications in energy storage devices.

The excellent electrochemical performance of the Nb₂O₅/Nb₂C/rGO aerogel can be attributed to the rationally designed unique aerogel structure and the synergistic effects of three different components (Fig. 7): (1) the 3D porous structure with a large specific surface area can ensure the full contact between the electrolyte and the active material, exposing more active sites. In addition, the Nb₂O₅ nanoparticles in the composite can further expose more active sites. These are beneficial for the improvement of specific capacity. (2) The 3D interconnected conductive network formed by the intimate contact of three components, the Nb₂O₅ nanoparticles generated by in situ oxidation, and the residual Nb₂C with a low potassium diffusion barrier and a high conductivity can shorten the diffusion distance and promote the diffusion kinetics of ions/electrons, which is conducive to obtaining outstanding rate performance. (3) The porous aerogel structure can accommodate the volume change during the charge–discharge process, and suppress the self-restacking of Nb₂C and agglomeration of Nb₂O₅ nanoparticles, thereby retaining the integrity and stability of the Nb₂O₅/Nb₂C/rGO structure.

Fig. 7 A schematic illustration of the enhanced energy-storage mechanism of Nb₂O₅/Nb₂C/rGO.

4. Conclusions

In summary, the Nb₂C/rGO conductive framework with in situ generated Nb₂O₅ nanoparticles (Nb₂O₅/Nb₂C/rGO) was successfully synthesized by a simple one-step high-temperature hydrothermal method and subsequent freeze-drying and annealing treatments, in which small-sized Nb₂C nanosheets are encapsulated by large-sized rGO nanosheets, and the Nb₂O₅ nanoparticles generated by in situ oxidation of Nb₂C are uniformly dispersed on the surface of Nb₂C and/or rGO, and the intimate contact of three components makes full use of the advantages of the 3D conductive network. The resulting 3D porous hybrid aerogel with abundant pores and a large specific surface area can not only inhibit the self-restacking of Nb₂C nanosheets and the agglomeration of Nb₂O₅ nanoparticles and alleviate the volume change during the charge–discharge process to improve the cycling stability, but also provide unimpeded channels for the diffusion of K⁺ and the infiltration of the electrolyte and expose more active sites to improve the specific capacity and rate performance. Meanwhile, Nb₂O₅ nanoparticles and residual Nb₂C nanosheets with a low potassium diffusion barrier and a high conductivity can shorten the diffusion distance and facilitate diffusion kinetics of electrons/ions. As a result, compared with Nb₂O₅/rGO, Nb₂O₅ and previous reports, the obtained Nb₂O₅/Nb₂C/rGO exhibits a much better potassium storage behavior as an anode of KIBs. Specifically, Nb₂O₅/Nb₂C/rGO presents a large specific capacity of 410.6 mA h g⁻¹ after 100 cycles at 0.1 A g⁻¹, and an attractive rate performance of 159.0 mA h g⁻¹ at 5 A g⁻¹. In addition, after 1000 consecutive high-rate cycles at 2.0 A g⁻¹, it still has a specific capacity of 208.9 mA h g⁻¹, with an ultralow capacity decay of 0.011% per cycle. More strikingly, Nb₂O₅/Nb₂C/rGO also shows a good potassium storage performance in a KIB full-cell. We believe that this facile synthesis method and the unique electrode structure can pave a new way to improve the performance of MXenes and accelerate the application of MXenes.

Author contributions

Cong Liu: conceptualization, methodology, software, investigation, data curation, visualization, and writing – original draft. Zhitang Fang: methodology, software, validation, and data curation. Weizhi Kou: methodology, validation, and data curation. Xiaoge Li: methodology, validation, and data curation. Jinhua Zhou: validation and investigation. Gang Yang: conceptualization, formal analysis, resources, and supervision. Luming Peng: conceptualization, formal analysis, resources, and supervision. Xuefeng Guo: conceptualization, formal analysis, resources, and supervision. Weiping Ding: conceptualization, formal analysis, resources, and supervision. Wenhua Hou: conceptualization, formal analysis, resources, writing – review & editing, supervision, project administration, and funding acquisition.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

We acknowledge the financial support of the National Natural Science Foundation of China (22379063) and the Modern Analysis Center of Nanjing University.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d3qi01775c

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