Construction of waffle-like NS-ZIF@V₂CT_x heterostructures for high-performance potassium-ion batteries†

Yue Qin ^a, Weifang Zhao *^b, Ting Wang ^a, Wenlong Liu ^a, Tengfei Zhou ^c, Xiaole Han ^a, Yi Liu ^a, Juncheng Hu ^a and Qingqing Jiang *^a
aSouth-Central Minzu University, Wuhan 430074, China. E-mail: 2014050@mail.scuec.edu.cn
^bGanfeng LiEnergy Technology Co., Ltd, Xinyu 338000, Jiangxi, China
^cAnhui University, Hefei 230601, China

First published on 25th June 2025

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Qingqing Jiang

Dr Qingqing Jiang's research is focused on potassium-ion batteries and interfacial problems encountered in energy storage systems. She received her BSc degree from China University of Petroleum in 2008. In 2014, she obtained her PhD degree from Dalian Institute of Chemical Physics, Chinese Academy of Sciences (CAS) (Advisor: Professor Can Li). Dr Jiang began her professional research career at South-Central Minzu University in 2014. She is currently an associate professor at the School of Chemistry and Materials Science in South-Central Minzu University.

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1. Introduction

Potassium-ion batteries (PIBs) are considered effective alternatives to lithium-ion batteries due to the abundant potassium resources, relatively low redox potential, small Stokes radius, inexpensive negative collector fluid,^1–7 and technical properties. Furthermore, the potassium-ion battery shows the same “rocking chair” type charge and discharge principle as the well-established lithium-ion batteries (LIBs) and extensively studied sodium-ion batteries (SIBs).^8–10 However, the practical application of PIBs is impeded by the larger ionic radius of the K-ion, which results in rapid capacity fading and limited reaction kinetics in electrodes. Therefore, the development of new electrode materials with high energy density, high rate performance and excellent stability has become an important challenge in the research and development of potassium ion batteries.

MXenes are an emerging group of transition metal carbides, nitrides, or carbonitrides with a two-dimensional layered structure. They have the molecular formula Mn₊₁AX_nT_x, where M is a transition metal, A is a group III/IV element of the periodic table, X is carbon or nitrogen, and T is a surface termination group (e.g., –O, –OH or –F).^11–13 MXene materials exhibit good hydrophilicity, excellent electrical conductivity, good mechanical stability and high surface area, and the abundant functional groups on their surfaces can serve as nucleation bases to anchor other active materials.^14–18 Among the common MXene materials, V₂CT_x MXene demonstrates superior mechanical strength and toughness, abundant surface groups and low potassium ion diffusion energy barriers compared to other MXene materials.^19–21 However, despite the high electrical conductivity of the MXene materials, they suffer from the self-stacking phenomenon after multiple charge/discharge cycles, leading to fast capacity decay and performance degradation.^22,23

Metal–organic frameworks (MOFs) are among the most attractive materials for energy storage and conversion due to their large surface area, tunable pore size and numerous functional ligands.^24,25 Meanwhile, the metal ions in MOF can act as redox active sites in electrochemical processes. Therefore, MOFs can also be used as electrode materials for PIBs.²⁶ It is noteworthy that pure MOFs suffer from poor electrical conductivity and agglomeration when used as electrode materials. Thus, researchers construct the MXene/MOF heterostructure by utilizing the abundant surface functional groups of MXenes to adsorb metal ions and subsequently grow MOF materials, which can realize the dispersion of MOF materials, while leveraging the high conductivity and flexible mechanical properties of MXenes.^27–29 In general, MOF/MXenes heterostructures exhibit the following characteristics: (i) the adsorption and insertion of porous MOF nanoparticles into the interlayers of MXenes can create porous nanostructures with high surface area, providing more active sites;^30–33 (ii) the composite structure can synergistically inhibit the agglomeration and self-stacking behaviour of MOFs and MXenes, and promote the charge transfer and mass transport;^34,35 (iii) MXenes can be an excellent substrate for supporting the MOF structure and improve its electrochemical performance, and the excellent conductivity of MXenes helps to overcome the low electronic conductivity problem of MOFs.^36–38

The multidimensional morphological structure of MOFs provides versatility for complexation with MXenes. For example, Xiao's team designed a Ti₃C₂T_x/NiCo-MOF heterostructure for lithium-ion batteries, where the bimetallic NiCo-MOF could provide large numbers of exposed active sites, while the Ti₃C₂T_x lamellar material with increased interlayer spacing can further facilitate the Li-ion diffusion kinetics.³⁹ Zhu's team exploited the chemical interactions between MXene and the functional groups on the surface of the 2D sheet Ni-MOF to develop 2D MXene/MOF heterojunction nanosheets for lithium anode electrochemical performance studies through a liquid-phase self-assembly process.⁴⁰ Sun's team synthesised the Bi_0.17Sb_0.83-MOF@MXene with 3D porous nanostructures by a self-assembly method using MXene as a substrate, which still showed an excellent reversible capacity of 260 mAh g⁻¹ after 900 cycles at 200 mA g⁻¹.^41,42 Generally, the large particle size of MOFs compromises the interfacial bonding strength, and their substantial mass reduces the energy density of PIBs. Constructing a hybrid structure with two-dimensional MOFs and MXenes that could provide larger contact areas, potentially accelerating the charge transfer channels between interfaces, is thus in high demand.^43–45

In the present work, V₂CT_x MXene nanosheets were synthesised by selectively etching the metallic Al layer in V₂AlC using NaF and HCl as a green etchant. The obtained V₂CT_x nanosheets were exfoliated by the intercalator TMAOH to obtain lamellar V₂CT_x nanosheets. The negatively charged V₂CT_x surface could attach Zn²⁺ and Co²⁺ to its interlayer and surface, while the metal ions rapidly reacted with ligands (2-methylimidazole) to form ZIF@V₂CT_x. Subsequently, the ZIFs were transformed from a rhombic dodecahedron to ultrathin porous nanosheets by solvothermal reaction in the presence of Co²⁺, forming porous waffle-like NS-ZIF/V₂CT_x heterostructures. The specific capacity of the NS-ZIF@V₂CT_x composite was 278 mAh g⁻¹ after 200 cycles at 100 mA g⁻¹. This work provides a simple strategy for the design of MXenes and MOF composite structures.

2. Experimental section

2.1 Synthesis of V₂CT_x MXenes

1.0 g of NaF was slowly added to a PTFE liner containing 20 mL of 38% concentrated hydrochloric acid and magnetically stirred for 15 min at room temperature to completely dissolve the NaF. Then, 1.0 g of V₂AlC was slowly added to the solution and stirring was continued for 15 min to obtain a homogeneous mixture. The liner was then placed in a hydrothermal vessel and transferred to a blast furnace to react for 72 h at 90 °C. Subsequently, the product was washed several times with anhydrous ethanol and deionised water alternately. Then, the obtained V₂CT_x was transferred to the PTFE liner and reacted at 90 °C. The product was then washed several times by centrifugation alternately with anhydrous ethanol and deionised water until the pH of the supernatant was close to neutral. The obtained V₂CT_x was then transferred to a PTFE liner. Thereafter, 3 mL of TMAOH, 200 mg of ascorbic acid and 20 mL of deionised water were added to the liner and the reaction was carried out hydrothermally for 24 h at 140 °C. The product was washed by centrifugation alternately with anhydrous ethanol and deionised water several times. The product was washed under a N₂ environment at room temperature for 15 min and then sonicated in an ice bath under a N₂ environment for 1 h. Finally, lyophilization was performed to obtain the stripped oligolayer V₂CT_x MXene.

2.2 Synthesis of ZIF@V₂CT_x

The V₂CT_x prepared above was added into a beaker containing 80 mL of methanol and stirred for 20 minutes. Then, 200 mg of PVP was added into the solution and stirred for 10 minutes. Thereafter, 0.595 g of Zn(NO₃)₂·6H₂O and 0.291 g of Co(NO₃)₂·6H₂O were added and stirred for 30 min, after which the mixture was sonicated for 20 min. The solution was then subjected to a further 30 minute sonication, followed by a 20 minute sonication cycle. Subsequently, 1 g of dimethylimidazole was added to a beaker containing 40 mL of deionised water, resulting in the formation of solution B. Solution B was then gradually added to solution A under constant stirring. The reaction was allowed to proceed for a duration of 2 hours, after which the mixture was left to stand for a period of 12 hours. Finally, the solution was subjected to a centrifugal process, followed by drying, to yield the desired product, ZIF@V₂CT_x.

2.3 Synthesis of NS-ZIF@V₂CT_x

The above prepared ZIF@V₂CT_x, 0.291 g Co(NO₃)₂·6H₂O and 30 mL methanol were added to a PTFE liner and stirred for 5 min. Then, the mixture was hydrothermalised at 120 °C for 4 h, and centrifuged and dried to obtain NS-ZIF@V₂CT_x. The subsequent annealing treatment was carried out in an Ar gas environment at 500 °C, and NS ZIF@V₂CT_x was annealed for 2 h at an increasing temperature rate of 5 °C min⁻¹.

3. Results and discussion

Fig. 1 shows the schematic diagram for the synthesis of NS-ZIF@V₂CT_x. Firstly, V₂CT_x MXene nanosheets were synthesised by selectively etching the metallic Al layer in V₂AlC using NaF and HCl as etching agents.⁴⁶ The as-obtained V₂CT_x nanosheets were further exfoliated by the intercalating agent (TMAOH) to obtain fewer layered V₂CT_x nanosheets with a smooth surface. The few-layered V₂CT_x nanosheets have richer interlayer functional groups, more reactive active sites, larger layer spacing and larger specific surface area than the multilayers, all of which can promote the subsequent rapid growth of ZIFs on the V₂CT_x surface and in the interlayers. The negatively charged functional groups of the few-layered V₂CT_x nanosheets attract the positively charged Co²⁺ and Zn²⁺ to aggregate to the surface by electrostatic interaction.

These self-assembled ZIF further increase the interlayer spacing of V₂CT_x, and organise the stacking of the layers of V₂CT_x, forming a “column effect” that improved the ion diffusion efficiency and structural stability.⁴⁷ Subsequently, the ZIF nanoparticles with rhombic dodecahedron were transformed into ultrathin porous NS-ZIF nanosheets via solvothermal reaction in the presence of Co²⁺ because methanol can affect the coordination mode and induce a mild phase transition of the ZIF nanoparticles.^48,49

From Fig. 2a, it can be seen that V₂CT_x has a typical multilayer accordion-like structure, and the interlayer van der Waals forces and hydrogen bonding of V₂CT_x were destroyed after intercalation treatment with TMAOH, resulting in fewer layered V₂CT_x flakes (Fig. S1a†) with a larger interlayer spacing of 1.02 nm.⁵⁰Fig. 2b shows that a large number of self-assembled ZIFs are attached into the V₂CT_x framework, demonstrating the successful synthesis of ZIF@V₂CT_x. The particle size of ZIFs are in the range of 60–80 nm, which increase the interlayer spacing of V₂CT_x and maintain the structural stability of V₂CT_x. Fig. 2d and S4† further confirm the uniform distribution of the ZIFs nanoparticles. As exhibited in Fig. 2c and S1b,† the waffle-like NS-ZIF/V₂CT_x heterostructures have been successfully constructed via solvothermal reaction in the presence of Co²⁺. The ZIFs have been transformed from a rhombic dodecahedron into ultrathin porous ZIF nanosheets (Fig. S1d†), and these nanosheets are wrapped around the surface of V₂CT_x and the interlayer space (Fig. S1c†) to form a continuous porous ZIF network. Such a waffle-like NS-ZIF@V₂CT_x framework offers abundant active sites for K-ion storage and rapid diffusion channels for the electrolyte. As demonstrated in Fig. 2e, the surface and edges of V₂CT_x were coated uniformly by these ultrathin porous ZIF nanosheets. The high-resolution TEM image in Fig. 2f shows that the crystallite spacings of 0.256 nm and 0.246 nm correspond to the (100) and (110) crystallites of V₂CT_x, respectively. However, no clear lattice stripes of NS-ZIF could be observed. This may be due to the ZIF particles having already been transformed into amorphous nanosheets. The SAED patterns of ZIF@V₂CT_x, NS-ZIF@V₂CT_x and ANS-ZIF@V₂CT_x are provided in Fig. S2.† For ZIF@V₂CT_x, the ring patterns correspond to the (110) crystal plane of V₂CT_x and the (201), (132) crystal planes of the ZIF nanoparticles.^51,52 For NS-ZIF@V₂CT_x, there is no obvious diffraction ring owing to the transformation of the ZIF nanoparticles into amorphous nanosheets, which is consistent with the XRD results. For ANS-ZIF@V₂CT_x, the obvious bright diffraction rings could be ascribed to the (220), (111) and (311) crystal planes of the metallic Co phase.⁵³ The energy spectral analysis (EDX) and elemental mapping analysis shown in Fig. 2g and S3† indicate that the elements C, O, V, Co and Zn are uniformly distributed in the lamellar structure, which further confirms the coexistence of ZIF in the V₂CT_x MXene lamellar matrix.

Fig. 3a shows the XRD spectra of V₂CT_x, ZIF@V₂CT_x and NS-ZIF@V₂CT_x. For V₂CT_x, the diffraction angle is dominated by the characteristic (002) peak at 9.2°, while the strong peaks at 13.5° and 41.2° essentially vanished, proving the complete etching Al from the MAX phase.⁵⁴ For ZIF@V₂CT_x, the diffraction angle at 7.4° corresponds to the (002) plan of V₂CT_x and the (011) facet of ZIF. Meanwhile, the peaks at 10.5°, 12.7° and 16.5° correspond to the (002), (112) and (222) planes of the ZIF nanoparticles, respectively, indicating that the two are successfully combined.^55,56 It should be noted that compared with V₂CT_x, the (002) diffraction peak of ZIF@V₂CT_x is shifted to a lower angle, demonstrating that ZIF grew in the interlayer of V₂CT_x, which drastically increased the spacing of the layers of V₂CT_x. For NS-ZIF@V₂CT_x, the (002) peak of V₂CT_x becomes inconspicuous, while the characteristic peaks of the ZIF nanoparticles almost disappear. No obvious information appeared in the low angle diffraction region. This indirectly confirms that the crystalline ZIF particles have been transformed into a nanoflake structure and V₂CT_x was wrapped with ZIF nanosheets.

From the Raman spectra in Fig. 3b, it can be seen that there are two broad peaks at 1359 cm⁻¹ and 1583 cm⁻¹ for V₂CT_x, ZIF@V₂CT_x and NS-ZIF@V₂CT_x, corresponding to disordered carbon and graphitic carbon, while the intensity ratios (I_D/I_G) are 0.71, 0.91 and 0.92, respectively. This indicates that NS-ZIF@V₂CT_x has a higher percentage of disordered carbon, which could provide a large number of active sites for the storage of K.⁵⁷ From the Fourier transform infrared spectra in Fig. 3c, it can be seen that the peaks at 3440 cm⁻¹, 1627 cm⁻¹, and 671 cm⁻¹ correspond to the –OH, CO, and V–C stretching vibrational modes of V₂CT_x, respectively.⁵³ For ZIF@V₂CT_x and NS-ZIF@V₂CT_x, the peak at 748 cm⁻¹ corresponds to the imidazole ring out-of-plane bending vibrational mode, while the peaks in the range of 900–1500 cm⁻¹ correspond to the imidazole ring out-of-plane bending vibrational mode.^58,59 The peaks at 1388 cm⁻¹, 1627 cm⁻¹, and 3440 cm⁻¹ correspond to C–H, CO, and –OH, respectively.

The pore size distribution was analyzed by using BJH models. As displayed in Fig. S5a,† the adsorption–desorption curve of V₂CT_x exhibits the type IV behavior with H3-type hysteresis loops, suggesting a mesoporous structure.⁶⁰ The pore size is concentrated between 10 nm to 40 nm with the average pore size of 20.4 nm and a mean cumulative pore size of 0.06 cm³ g⁻¹. As illustrated in Fig. S5b,† the adsorption–desorption curve of ZIF@V₂CT_x shows type I behavior with a distinct H3-type hysteresis line.⁶¹ ZIF@V₂CT_x presents a large specific surface area of 490.6 m² g⁻¹, while the pore size is concentrated at approximately 4 nm and a cumulative pore volume of 0.3 cm³ g⁻¹. The specific surface areas of V₂CT_x and ZIF@V₂CT_x were similar to those generally reported in the literature.^47,60,62 After combining V₂CT_x with the ZIF nanoparticles, the specific surface area underwent a substantial increase because the porous ZIF material inherently possesses a high specific surface area, while the ZIF particles enlarge the interlayer spacing of V₂CT_x, creating additional surface area. In contrast, the adsorption and desorption of NS-ZIF@V₂CT_x showed typical type IV behavior with a distinct H4-type hysteresis line with an average pore size of 11.4 nm and a mean cumulative pore size of 0.58 cm³ g⁻¹ (Fig. 3d). The adsorption and desorption curves of NS-ZIF@V₂CT_x showed significant hysteresis lines in the relative pressure hysteresis line (P/P₀) in the range of 0.45–0.98, indicating the presence of a mesoporous structure. These ZIF nanoparticles were able to have a large and homogeneous pore structure after transforming into nanosheets on the V₂CT_x substrate, which would be beneficial for the rapid transport of bulk K⁺. According to thermogravimetric analysis, the porous NS-ZIF@V₂CT_x can be stable above 250 °C (Fig. S6†).

X-ray photoelectron spectroscopy (XPS) was employed to examine the surface chemical state of V₂CT_x and NS-ZIF@V₂CT_x. The presence of Zn, Co, V, C and O elements was confirmed in the XPS spectra of V₂CT_x and NS-ZIF@V₂CT_x. The presence of F was derived from the etchants (NaF). In addition, Al peaks were not detected in the XPS survey, which also indicated that the Al layers have been essentially removed. Fig. 4b shows that the C 1s region spectrum of NS-ZIF@V₂CT_x consists of C–C, C–O and CO species at 248.8 eV, 285.8 eV and 288.5 eV, respectively, which are mainly derived from the synthesised ZIF organic ligands and V₂CT_x. The strong C–V bond at 282.7 eV could not be observed in NS-ZIF@V₂CT_x, owing to the surface of V₂CT_x having been covered by NS-ZIF.

In the O 1s spectrum, there are three peaks at 530.4 eV, 531.8 eV and 533.6 eV, corresponding to C–V–O, C–V–OH and absorbed H₂O, respectively. The C–V–O and C–V–OH peaks originate from the organic ligands and oxygen-containing functional groups on the surface of V₂CT_x. Additionally, the C–V–O peaks could be attributed to the surface oxidation of V₂CT_x during the synthesis of NS-ZIF@V₂CT_x, which was also confirmed by V 2p spectra.^55,63,64 In the V 2p spectrum, the binding energy peaks of 516.8 eV and 517.2 eV of V 2p_3/2 belong to V⁵⁺ and V⁴⁺, respectively, and the binding energy peaks of 523.6 eV and 525 eV of V 2p_1/2 correspond to V⁵⁺ and V⁴⁺, respectively. The presence of the oxidative valence is probably related to the surface oxidation of V₂CT_x during the synthesis of NS-ZIF@V₂CT_x. In general, the trend of the vanadium species is consistent with the C 1s spectrum. For the Zn 2p spectrum (Fig. 4e), the two peaks located at 1022.2 eV and 1045.0 eV could be ascribed to Zn²⁺ in ZIF-8.⁶⁵ Regarding the Co 2p spectrum (Fig. 4f), the curve can be divided into Co 2p_3/2, Co 2p_1/2, and two satellite peaks. The peaks at 782.4 and 797.2 eV reflect the presence of Co²⁺, while the peaks at 780.0 and 795.1 eV are ascribed to Co³⁺ in ZIF-67.^{45,59,66–68}

The potassium storage performance was tested with pure V₂CT_x, ZIF@V₂CT_x and NS-ZIF@V₂CT_x as the anode materials (K-foil as the counter electrode) in a CR2032 coin half-cell. Fig. 5a shows the cyclic voltammetry curve of the NS-ZIF@V₂CT_x electrode at a sweep rate of 0.1 mV s⁻¹.⁶⁹ Two distinct cathodic peaks appear at around 0.30 V and 0.80 V during the first round of cathodic scanning. The peak at 0.30 V disappears in the subsequent scans, which could be attributed to the occurrence of the electrolyte decomposition and the formation of the SEI film.⁷⁰ The cathodic peak at 0.80 V indicates the embedding behaviour between K⁺ and the NS-ZIF@V₂CT_x framework. The K⁺ could be stored in the structural pores of the ZIF nanosheets, as well as in the interlayer spaces of V₂CT_x. There are two types of K-ion battery storage mechanisms in the MOF-based electrode materials: conversion and insertion. In the insertion-type mechanism, the host–guest interaction between K⁺ and the amine group/N atoms in the imidazole or carboxylic acid ligands allows K⁺ to be stored in the pores. During the insertion process, the multidimensional network and coordination environment of the MOFs would not be affected because the central metal ions do not participate in the reaction during the charging and discharging process. Thereby, the intercalation/conversion reaction may be illustrated as follows,^61,63,71,72

Zn–(2IM)2 + K+ + e− ↔ K+–(2IM)2–Zn	(1)

Co–(2IM)2 + K+ + e− ↔ K+–(2IM)2–Co	(2)

In addition, the peak position of the CV curve essentially corresponds to the voltage plateau of the first scan in the charge/discharge curve of the NS-ZIF@V₂CT_x electrode (Fig. 5b). During the first scan, the positive electrode undergoes an oxidation reaction and electrons, and K⁺ ions are released from the NS-ZIF@V₂CT_x electrode. This process is associated with the appearance of an oxidation peak with this electrode reaction in the voltage range of approximately 1.40 V. The subsequent cycles have similar curve shapes with good overlap, indicating that the NS-ZIF@V₂CT_x electrode presents superior cycling reversibility and stability.

Fig. 5c shows the cycling performance of V₂CT_x, ZIF@V₂CT_x and NS-ZIF@V₂CT_x at the current density of 100 mA g⁻¹. After 200 lap cycling tests, the specific discharge capacity of NS-ZIF@V₂CT_x is 278 mAh g⁻¹, while the capacities of pure V₂CT_x and ZIF@V₂CT_x are only 98 mAh g⁻¹ and 93 mAh g⁻¹, respectively. NS-ZIF@V₂CT_x presents much larger specific capacities than pure V₂CT_x and ZIF@V₂CT_x, which is mainly due to the insertion of ZIF particles into V₂CT_x, creating a ‘column effect’ to increase the interlayer spacing. The subsequent network of ZIF nanosheets thus provides more pores for K⁺ storage.

Fig. 5d depicts the rate performance for the NS-ZIF@V₂CT_x electrode. The reversible capacities are 560.4, 384.9, 245.2, 166.4 and 107.9 mAh g⁻¹ at current densities ranging from 0.1 to 2.0 A g⁻¹. It is worth noting that the reversible capacity of the NS-ZIF@V₂CT_x electrode showed a strong decreasing trend with increasing current density, but a reversible capacity of 166.4 mAh g⁻¹ can still be achieved at a current density of 1 A g⁻¹. In addition, when the current density is reduced to 0.1 A g⁻¹, the discharge specific capacity can be quickly recovered to 387 mAh g⁻¹ and remains stable, confirming the structural stability of NS-ZIF@V₂CT_x after rapid K⁺ insertion/extraction. The GDV curves at current densities from 0.1 to 2.0 A g⁻¹ also confirm the reliable rate performance of the NS-ZIF@V₂CT_x electrode (Fig. 5e). It also has a capacity retention rate of over 98%, indicating excellent long-cycle performance. Compared to most of the previously reported pure MOFs and its derivative electrodes, the NS-ZIF@V₂CT_x electrode has better electrochemical performance (Fig. 5h).^72–79

To improve the rate performance under high current density, NS-ZIF@V₂CT_x was further annealed under the protection of an Ar atmosphere to obtain ANS-ZIF@V₂CT_x. ANS-ZIF@V₂CT_x had a specific capacity of at least 300 mAh g⁻¹ in the 200-lap cycling test, with better cycling performance than NS-ZIF@V₂CT_x (Fig. S7a†). At current densities ranging from 0.1 to 2.0 A g⁻¹, the reversible capacities of the ANS-ZIF@V₂CT_x electrode were 336.9, 311.1, 277.3, 225.4 and 155.9 mAh g⁻¹ (Fig. 5f), indicating that the ANS-ZIF@V₂CT_x electrode has significantly higher reversible capacity and stability at high current densities than the NS-ZIF@V₂CT_x electrode. In addition, the long-term cycling stability of the ANS-ZIF@V₂CT_x electrode was evaluated at current densities of 1.0 A g⁻¹ and 2.0 A g⁻¹. After 500 cycles at a current density of 1.0 A g⁻¹, a high reversible specific capacity of 163 mAh g⁻¹ was achieved (Fig. S7b† and 6c), with a capacity retention rate of 80% (Fig. S11–S16†).

Overall, the dense nano-network of the NS-ZIF@V₂CT_x electrode could not only increase the interlayer spacing of the V₂CT_x MXene, but also impede the occurrence of the self-stacking phenomenon of V₂CT_x, thus accelerating the diffusion and transport efficiency of electrons and ions. The waffle-like lightweight structures can increase the energy density of potassium-ion batteries. Furthermore, after calcination treatment of NS-ZIF@V₂CT_x, the resulting ANS-ZIF@V₂CT_x electrode with the carbon skeleton supported waffle-like morphology exhibits further enhanced performance under large current densities.

As demonstrated in Fig. 6, CV, EIS and GITT measurements were performed to investigate the reaction kinetics of the NS-ZIF@V₂CT_x electrode, as well as the K⁺ diffusion coefficient. With the increasing scan rate from 0.1 to 1 mV s⁻¹, the closed region of the CV curve became more prominent, and the anodic and cathodic peaks were shifted to the positive and negative directions, respectively. As shown in Fig. 6a, the anodic peaks (peak 1 and peak 2) shifted to higher voltages with increasing scan rate, and this change can be attributed to diffusion-controlled reactions and concentration polarisation. The peak current (i) varies with the scan rate (v), and generally satisfies the relation:

i = avb	(3)

where a and b are adjustable coefficients, and the electrochemical reaction is mainly controlled by diffusion when the b value is close to 0.5, and mainly by pseudocapacitance when the b value is close to 1.0. For the NS-ZIF@V₂CT_x electrode, the b values of the two anodic peaks were calculated to be 0.67 and 0.59, respectively, based on the linear relationship between the scan rate and peak current (Fig. 6b). The b values of the two anodic peaks were in the range of 0.5–1.0, indicating that the NS-ZIF@V₂CT_x electrode controlled by both diffusion-controlled and pseudo-capacitive behaviour. By means of the following relational formula,

i = k1v + k2v1/2	(4)

the pseudocapacitance contribution can be calculated for different scan rates, where k₁v represents the contribution of pseudocapacitance and k₂v^1/2 represents the contribution of diffusion. As shown in Fig. 6c, for the NS-ZIF@V₂CT_x electrode, the percentage of pseudocapacitance contribution is 21, 24, 31, 36, 41, and 45% as the scanning rate increases from 0.1, 0.2, 0.4, 0.6, 0.8 to 1.0 mV s⁻¹, respectively, and the pseudocapacitance contribution gradually increases with the scanning rate. In addition, the enlarged layer spacing of NS-ZIF@V₂CT_x facilitates the diffusion-controlled electrochemical behaviour.

According to the GITT curves of pure V₂CT_x, ZIF@V₂CT_x and NS-ZIF@V₂CT_x (Fig. S8†), the diffusion coefficient (D_K⁺) of the potassium ions can be calculated based on Fick's second law, as shown in Fig. 6d. The D_K⁺ values of the NS-ZIF@V₂CT_x electrode during discharge were 10⁻⁹ to 10⁻¹¹ cm² S⁻¹. Meanwhile, for the pure V₂CT_x electrode and ZIF@V₂CT_x, the D_K⁺ values were 10⁻¹⁰ to 10⁻¹³ cm² S⁻¹ and 10⁻¹⁰ to 10⁻¹² cm² S⁻¹, respectively (Fig. S9†). Because the adsorbed K⁺ repels the newly embedded K⁺, the D_K⁺ thus gradually decreases. The K⁺ diffusion coefficients of NS-ZIF@V₂CT_x during K⁺ detachment and embedding are at least one order of magnitude higher than those of pure V₂CT_x and ZIF@V₂CT_x, confirming the convenient potassium ion transport channels and fast diffusion kinetics. This could be attributed to the following three aspects: firstly, the waffle-like ZIF/V₂CT_x heterostructures can provide sufficient and fast channels and facilitate electron/ion transfer for K⁺, and ensure sufficient penetration of the electrolyte and electrode. Secondly, V₂CT_x has a large layer spacing in NS-ZIF@V₂CT_x, which provides multiple channels for K⁺ transfer. In addition, the high conductivity of V₂CT_x MXene promotes K-ion diffusion and efficient electron transfer, ensuring superior electrochemical performance.

As shown in Fig. 6f, the EIS curves have been calculated by applying the equivalent circuit model, including the internal resistance (R_s), SEI resistance (R_SEI) and charge transfer resistance (R_ct). The R_ct value of NS-ZIF@V₂CT_x shows a significant decline after 20 cycles of charging and discharging process, indicating that the stabilised SEI film on the surface of the electrode could accelerate electron transfer and improve the reaction kinetics. In addition, the decrease in R_ct could be related to the waffle-like network, which remains stable during charging and discharging, providing a fast transport channel for charge transfer. This is also supported by the SEM images of the NS-ZIF@V₂CT_x electrode after 200 charge/discharge cycles, demonstrating no obvious changes in the structure (Fig. S10a†). The increased R_ct resistance of ZIF@V₂CT_x may be related to the disruption of the structural integrity of the ZIF, which reduces the stability of the SEI membrane. V₂CT_x has the lowest R_ct, suggesting that the introduction of good conductivity of V₂CT_x can also effectively improve the electrochemical properties of the porous ZIF nanosheet materials.

The composition of the SEI film on the surface of the NS-ZIF@V₂CT_x electrode was further analysed by XPS and in situ electrochemical FT-IR spectra (Fig. 7). Fig. 7a shows that the uniform and stable SEI films have been generated on the NS-ZIF@V₂CT_x electrode during the charging and discharging process, with a thickness of about 18–20 nm, and the SEI maintains a stable structure throughout the voltage charging and discharging stages of the electrode (Fig. 7b). From the in situ electrochemical FT-IR spectra, the C–O vibrational intensity of the fatty ether at 1181 cm⁻¹ is low at the initial stage. As the voltage changes, the KFSI/DME electrolyte decomposes, while the decomposed electrolyte participates in the formation of a stable SEI film. The peak at 1181 cm⁻¹ corresponds to the transformation of DME into R–OK.^78,79 As shown in Fig. 7d, C 1s can be divided into four peaks, C–C (284.1 eV), C–O (284.8 eV) and CO (286.4 eV). The O 1s spectrum can be divided into R–OK (534.1 eV), C–O (533.1 eV), CO (532.2 eV) and –SO_xF (530.6 eV) (Fig. 7e). The peak at –SO_xF is due to the decomposition of the potassium salt KFSI, and the peak at R–OK is due to the decomposition of the ether-based solvent molecules.^80,81 The F 1s spectrum shows that the SEI membrane components include –SO_xF (687.9 eV) and KF (683.2 eV) (Fig. 7f). The major inorganic salt component in the SEI membranes is KF, which has low electronic conductivity and high surface energy, favouring the K⁺ diffusion kinetics.⁸² The stable and KF-rich organic–inorganic SEI layer is favourable for alleviating the volume expansion of the electrode.^83,84 However, due to the high KF content and large thickness of the SEI film (Fig. S10b†), the electrode easily hinders the reaction between the electrode and the electrolyte, leading to a significant attenuation of the electrode charge/discharge capacity at higher current densities and a poor rate performance.

4. Conclusions

In conclusion, the alternately stacked waffle-like NS-ZIF/V₂CT_x heterostructures have been fabricated via an innovative solvothermal reaction. NS-ZIF/V₂CT_x and its derivative demonstrate good potential in the field of potassium-ion batteries, which was mainly manifested in the toleration of volume expansion with more than 98% capacity retention, as well as the large potassium ion diffusion coefficient. Ex situ characterization and electrochemical kinetics tests have been performed to verify the superiority of the waffle-like nanostructure, which provide plentiful exposed active sites, enhanced electrical conductivity and superior structural stability. Especially, the waffle-like lightweight structures can increase the energy density of the potassium-ion batteries. This work provides a valuable reference for the diverse applications of MXene-based materials in alkali metal ion storage.

Data availability

The other data have been supplied in the ESI† files.

Author contributions

Yue Qin: original draft and data curation. Weifang Zhao, Ting Wang, and Wenlong Liu: data curation. Tengfei Zhou, Xiaole Han, Yi Liu, and Juncheng Hu: methodology and funding acquisition. Qingqing Jiang: writing, review, editing, investigation, and funding acquisition.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work was supported by the National Natural Science Foundation of China (22272207, 21806187, 32271538, and U23A2089).

Notes and references

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5ta03245h

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