Microstructural Engineering of Conductive MOF Nanocrystals for Multifunctional Heterostructures

Abstract

Conductive metal–organic frameworks (cMOFs) have emerged as transformative materials, significantly expanding the functional landscape of conventional MOFs by integrating electrical conductivity. Despite this, harnessing their full potential in heterostructured architectures is constrained by challenges in finely tuning the critical structural parameters, notably the packing density and crystallographic orientation. This paper introduces an innovative approach to engineering the microstructure of cMOFs as conformal shell‐enveloping spherical nanoparticles. The compactness of the cMOF layer is modulated by the crystallization kinetics, while a deliberate seeding process directed the alignment of cMOF nanocrystallites. These methodologies collectively facilitated the formation of densely packed, highly ordered nanocrystallite assemblies in a core‐shell configuration, enhancing mechanical robustness, selective permeability, and electron transport. The resulting cMOF assembly markedly augmented the stability and electrochemical performance of the sulfur core, showing significant promise as a next‐generation cathode material for lithium‐sulfur batteries.

A coating strategy based on conductive metal–organic frameworks (cMOFs) is presented to construct advanced heterostructures. By controlling crystallization and employing pre‐deposition seeding, precise nanocrystal assembly is achieved, maximizing selective mass transport, catalytic functionality, and electrical conductivity. This approach, as demonstrated in lithium–sulfur cathodes, offers broad applicability across advanced technologies that demand multifunctional materials.

1. Introduction

The advent of conductive metal–organic frameworks (cMOFs) is a seminal breakthrough, seamlessly integrating electrical conductivity into traditionally insulating, porous crystalline architectures.^{[ 1} , ² , ^{3 ]} Unlike conventional strategies that enhance conductivity via conductive additives or MOF carbonization, cMOFs leverage a conjugated network generated by partially oxidized organic ligands.^{[ 4} , ⁵ , ⁶ , ^{7 ]} This intrinsic design establishes efficient charge transport pathways and preserves the inherent advantages of periodic pore architectures, expanding MOF applications beyond molecular storage and separation to encompass electrically driven functionalities. Substantial progress has been made recently by precisely tailoring electronic structures through the controlled manipulation of metal nodes and ligand chemistry, combined with advances in synthetic techniques, such as interfacial assembly, liquid‐phase exfoliation, and defect engineering.^{[ 8} , ⁹ , ¹⁰ , ^{11 ]} These targeted strategies have significantly enhanced electrical conductivity, finely tuned bandgaps, and improved multifunctionality, leading to breakthroughs in catalytic efficiency, molecular selectivity, and optoelectronic capabilities, advancing their applicability in real‐world technologies.^{[ 12} , ¹³ , ¹⁴ , ^{15 ]}

Considering these compelling advantages, the strategic assembly of cMOFs into heterostructures is a promising avenue for engineering advanced functional nanomaterials.^{[ 16} , ¹⁷ , ¹⁸ , ¹⁹ , ^{20 ]} Various composites incorporating cMOFs with materials such as carbon nanotubes (CNT), cellulose, and reduced graphene oxide (rGO) have been studied extensively to improve the conductivity and structural robustness, thus significantly broadening their functional applicability.^{[ 21} , ²² , ^{23 ]} In particular, encapsulation with cMOFs has a dual benefit; it integrates the structural merits of the framework with the core material, and it exploits the precise regulation of material transport through periodic pore architectures, ensuring controlled permeability.^{[ 24} , ²⁵ , ^{26 ]} Moreover, the conductivity of the cMOF layer activates robust electrical pathways, optimizing the functional performance of the encapsulated core.^{[ 27} , ²⁸ , ^{29 ]} On the other hand, previous studies on cMOF‐based heterostructures have predominantly emphasized compositional integration and associated performance improvements, with limited attention directed toward the systematic control and chemical understanding of a nanocrystal assembly at the microstructural level. Thus, a comprehensive chemical understanding and systematic control of spatial arrangements, particularly the packing density and crystal alignment within cMOF nanocrystal assemblies, is essential for fully harnessing their functional potential.

This study addressed this challenge using a systematic approach for fine‐tuning the microstructure of cMOF layers within core‐shell constructs. The investigation centered on elucidating the formation mechanism and microstructural evolution of a Ni₃(HITP)₂ shell (where HITP = 2,3,6,7,10,11–hexaiminotriphenylene) as it encapsulates the nanoparticle cores. The Ni₃(HITP)₂ shell was assembled via nanocrystal aggregation on a nanoparticle substrate. The crystal size and packing density of the Ni₃(HITP)₂ layer were adjusted precisely by modulating crystallization kinetics. Furthermore, implementing a seeding and secondary growth protocol directs crystal alignment, culminating in the formation of ordered porous channels that enhance conductivity and functionality. This meticulously engineered Ni₃(HITP)₂ shell imparts atomically dispersed catalytic activity and finely tuned selective permeability while concurrently establishing continuous conductive pathways. The encapsulation of insulating sulfur nanoparticles with the Ni₃(HITP)₂ shell was performed to validate the proposed approach, achieving marked improvements in stability and performance even with a minimal shell content. These findings underscore the critical role of cMOF microstructural control in unlocking the full functional potential of heterostructured systems.

2. Results and Discussion

2.1. SNP@cMOF Core–Shell Structures

Figure 1a presents the concept of coating spherical nanoparticles (SNPs) with a conducting MOF (cMOF) layer. Uniform sulfur SNPs were chosen as the model core. They were synthesized by the disproportionation of sodium thiosulfate in the presence of polyvinylpyrrolidone (PVP) and HITP, yielding particles with a consistent size and morphology. Subsequently, the Ni₃(HITP)₂ shell was constructed via a coordination reaction between Ni²⁺ and HITP ligands in the presence of ammonia. In this system, the amino groups in ammonia facilitate the deprotonation of the HITP ligand, promoting cMOF crystallization on the SNP surface. The structural characteristics of the resulting core–shell assembly, designated as SNP@cMOF, were elucidated using a suite of analytical techniques.

Figure 1.

a) Schematic representation of the encapsulation strategy of cMOF on the SNP. b) SEM image of the SNP, and c) SNP@cMOF (inset: particle size distribution). d) TEM image of SNP@cMOF (inset: magnified view of (100) plane of the Ni₃(HITP)₂ layer). e) XRD patterns for SNP and SNP@cMOF (inset: expanded view of the Ni₃(HITP)₂ XRD peak). f) Raman spectrum for SNP and SNP@cMOF. g) Micropore size distribution curves for SNP@cMOF were calculated using the HK method (inset: enlarged pore size distribution calculated using the HK and BJH methods).

Scanning electron microscopy (SEM) (Figure 1b) confirmed the formation of uniform SNPs with smooth surfaces, exhibiting a mean diameter of 409.3 nm. After the cMOF coating, the particle surface became rougher, and the mean size increased to 457.6 nm, corresponding to an estimated shell thickness of ≈24.2 nm (Figure 1c). Transmission electron microscopy (TEM) corroborated these findings, showing that the cMOF shell is composed of aggregated nanocrystallites (Figure 1d). The inset in Figure 1d shows the (100) plane of Ni₃(HITP)₂, with an interplanar spacing of 1.85 nm. Complementary scanning transmission electron microscopy (STEM) coupled with energy‐dispersive X‐ray spectroscopy (EDS) line scans across the SNP@cMOF particles confirmed a uniform elemental distribution of Ni throughout the shell, supporting its conformal nature and consistent thickness in agreement with SEM observations (Figure S1, Supporting Information).

X‐ray diffraction (XRD) of the cMOF‐coated particles revealed new characteristic peaks at 4.7° and 9.5° 2θ, which were indexed to the (100) and (200) planes of Ni₃(HITP)₂, respectively, confirming the successful formation of the cMOF layer (Figure 1e).^{[ 30 ]} Raman spectroscopy further supported the presence of the cMOF coating. Initially, the D and G peaks in the SNP were attributed to the HITP layer embedded within the PVP matrix (Figure 1f). Upon cMOF incorporation, the Raman spectrum displayed prominent Ni─N and C═C peaks alongside the enhanced D and G bands, affirming the coordination of Ni with the HITP ligand.^{[ 31} , ^{32 ]} These combined XRD and Raman results closely align with independently synthesized pure Ni₃(HITP)₂, confirming the successful construction of a Ni₃(HITP)₂‐based heterostructure (Figure S2, Supporting Information). Thermogravimetric analysis (TGA) suggested that the cMOF shell contained 3.6 wt.% of the SNP@cMOF structure (Figure S3, Supporting Information). Moreover, the sulfur volatilization‐induced weight loss in SNP@cMOF shifted to higher temperatures relative to bare SNP because of the stronger interactions between the sulfur core and the cMOF shell.^{[ 33 ]}

Pore size analysis conducted using the Horvath–Kawazoe (HK) and Barrett–Joyner–Halenda (BJH) methods revealed nanopores with a mean diameter of 1.3 nm (Figure 1g), which is consistent with the literature on Ni₃(HITP)₂.^{[ 34 ]} The introduction of the cMOF shell significantly enhanced the Brunauer–Emmett–Teller (BET) specific surface area despite the much lower Ni₃(HITP)₂ content than the sulfur core, increasing the BET surface area from 5 to 21 m² g⁻¹ (Figure S4, Supporting Information). Collectively, these comprehensive investigations confirmed that a uniform and functional cMOF coating formed on the surface of the spherical nanoparticles. The encapsulation of Ni₃(HITP)₂ onto polystyrene cores modified by PVP–HITP, verified through SEM and Raman spectroscopy, underscores the versatility of our nanocrystal assembly strategy (Figure S5, Supporting Information). This result demonstrates its potential compatibility with diverse core materials and highlights the broader applicability for designing functional heterostructures.

2.2. Tunable Physical Properties of cMOF Shell

The spatial organization of cMOF nanocrystallites was refined systematically to impart programmable control over the physical properties of the shell structures. The high donicity of DMSO promotes strong coordination with metal ions through its incorporation into the aqueous reaction medium, decelerating the crystallization kinetics.^{[ 35} , ^{36 ]} This deliberate slowdown stabilizes intermediate species and suppresses rapid crystal growth. F and M represent the DMSO volume fraction in the cosolvent and the precursor concentration, respectively (further details are provided in the Experimental Section). For example, increasing the DMSO ratio from F0 to F0.2 at a fixed precursor concentration (M2.5) reduces the surface roughness and yields a more uniform cMOF shell (Figure 2a). In contrast, the surface reverts to that of the original SNPs at a higher DMSO ratio of F0.33, indicating incomplete crystal formation because of the overly retarded kinetics. TEM of the F0.2–M2.5 sample revealed a more compact and uniformly assembled crystallite layer, with a densely packed arrangement that minimizes interparticle voids compared to F0–M2.5 (Figure 2b). Under the F0.33–M2.5 condition, however, only minimal amorphous deposits were observed, which is in agreement with the SEM observations. The XRD patterns of F0–M2.5 and F0.2–M2.5 confirm the formation of cMOF nanocrystals, while the indistinct cMOF peaks in F0.33–M2.5 indicate incomplete crystallite formation, which is consistent with morphological analysis (Figure S6, Supporting Information).

Figure 2.

a) SEM images and b) TEM images of SNP@cMOF synthesized using various volumetric fractions of cosolvents and precursor concentrations. c) Mapping of the physical properties of the cMOF shell based on the volumetric fraction of the cosolvents and precursor concentrations. d) Schematic diagram of the spatial configuration of cMOF nanocrystallites under mono‐solvent synthesis versus cosolvent synthesis. e) SEM images of SNP@cMOF after soaking in a toluene solution. f) Fraction of ruptured particles of various SNP@cMOF after soaking in a toluene solution. g) Schematic diagram of the structural robustness of assembly cMOF nanocrystallites according to packing patterns.

The combined effects of the DMSO ratio and precursor concentration were assessed by varying the DMSO content at F0, 0.17, 0.2, and 0.25 and the precursor concentration at M1.0, 1.5, 2.0, and 2.5. At a constant precursor concentration, an increased DMSO ratio generally decreases the shell thickness and its overall content (Figure S7 and Table S1, Supporting Information). In particular, the decrease in thickness is more pronounced than the decrease in content, resulting in an overall increase in shell density, as shown in Figure 2c. Elevated DMSO ratios (F0.17, F0.2, and F0.25) promote further shell densification, enhancing its structural compactness. The DMSO–water cosolvent slowed the crystallization kinetics sufficiently to activate a greater number of nucleation sites compared to a purely aqueous reaction, as shown in Figure 2d. These findings underscore the pivotal role of the crystallization kinetics in achieving an organized, spatially confined assembly of nanocrystallites on the core surface.

The compact nanocrystallite arrangement enhances the structural robustness of the cMOF assembly.^{[ 37} , ^{38 ]} The integrity was evaluated using toluene to dissolve the sulfur core selectively while preserving the crystallinity of the cMOF (Figure S8, Supporting Information). The dissolved sulfur exited through the weak points in the shell, rupturing the core‐shell structure. In particular, despite the lower shell content, F0.2–M2.5 achieved a higher density of 0.82 g cm⁻³, which was attributed to a more pronounced thickness reduction effect, and consequently exhibits significantly fewer ruptured particles than F0–M2.5, which had a density of 0.70 g cm⁻³ (Figure 2e). Moreover, a systematic study suggested that cMOF shells synthesized with DMSO fractions of F0.17, F0.2, and F0.25 showed a markedly lower proportion of ruptured particles than those prepared without DMSO (F0; Figure 2f; Figure S9, Supporting Information). These results highlight the importance of configurational compactness in strengthening the structural robustness of the cMOF nanocrystallite assembly (Figure 2g). Overall, the kinetics‐based strategy provides a straightforward approach for precise control over the physical properties of the cMOF shell.

2.3. Oriented Growth of cMOF Assemblies

Most members of the cMOF family, including Ni₃(HITP)₂, have anisotropic crystal lattice structures.^{[ 39 ]} Aligning individual cMOF nanocrystallites is essential for establishing a collective network with oriented pore channels, facilitating directional diffusion.^{[ 40} , ^{41 ]} Moreover, the relaxation of grain boundaries between nanocrystallites enhances the overall electrical conductivity.^{[ 42 ]} This study aimed to direct the oriented growth of cMOF nanocrystallites on SNP to exploit the functionality of the cMOF shell.

Figure 3a compares two cMOF nano‐assemblies: one obtained via random growth (top) and the other achieved through oriented growth on a pre‐deposited metal–ligand seed (bottom). The pre‐deposited seed acts as a nucleation site, directing the assembly of cMOF along specific crystallographic planes to yield an ordered, aligned structure.^{[ 43} , ^{44 ]} Accordingly, a pre‐deposition seeding step was implemented using an ammonia‐free aqueous solution of Ni²⁺ and HITP before cMOF shell growth.

Figure 3.

a) Schematic diagram of the direct growth and oriented growth of the cMOF shell. b) SEM images of SNP‐3 and SNP‐12. c) EDS image of SNP‐12. d) Raman spectra and e) coordination numbers of SNP‐# with varying seeding times. f) Estimation of the oriented angle of cMOF nanocrystallites using TEM images of SNP‐3– and SNP‐12@cMOF. g) Distribution of the oriented angles of cMOF nanocrystallites in SNP‐#@cMOF as a function of the seeding time. h) Electrical conductivity of SNP‐#@cMOF paper at different seeding times.

The evolution of the Ni–HITP seed layer on SNP was monitored over time, designating “SNP‐#” as the sample subjected to a seeding step for # hours. SEM (Figure 3b; Figure S10, Supporting Information) showed no distinct crystalline features on the SNP surface immediately after pre‐deposition. On the other hand, the progressively smoother surface with extended seeding times indicated a more uniform seed distribution, providing an optimal substrate for oriented growth. EDS of SNP‐12 confirmed the homogeneous distribution of Ni, indicative of consistent Ni–HITP pre‐deposition (Figure 3c). Raman spectroscopy further corroborated the gradual formation of the seed layer, as evidenced by the increasing intensity of C═C and Ni─N vibrational peaks relative to sulfur with prolonged seeding (Figure 3d; Figure S11, Supporting Information).

Thermogravimetric analysis (TGA) was used to assess the Ni‐to‐HITP stoichiometric ratio within the seed layer (Figure 3e; Figure S12, Supporting Information). The X‐ray photoelectron spectroscopy (XPS) results were in agreement with the TGA findings (Figure S13, Supporting Information).^{[ 45 ]} The weaker O 1s signal in SNP‐3 results from limited adsorption of synthesis‐derived solvent water due to fewer available metal coordination sites in the initial Ni–HITP layer. Extended seeding increases Ni–HITP accumulation, enhancing the adsorption of water and hydroxyl groups, thus intensifying the O 1s peak.^{[ 46} , ^{47 ]} After three hours of pre‐deposition, the Ni–HITP seed layer exhibited a Ni‐to‐HITP ratio of 2.10, which increased to 3.45 after 12 h of seeding. Although extended seeding approached the theoretical Ni‐to‐HITP ratio of 4, it also induced pronounced particle aggregation (Figure S14, Supporting Information).

Figure S15 (Supporting Information) presents SEM images of SNP‐#@cMOF under F0.2–M2.0 conditions after introducing the cMOF shell. The orientation angles of the cMOF nanocrystallites were quantified using the (100) facet of Ni₃(HITP)₂ as a reference (Figure S16, Supporting Information). Representative TEM images (Figure 3f; Figure S17, Supporting Information) were analyzed by randomly selecting at least 50 nanocrystals from three independent images per condition. For each nanocrystal, the orientation angle was determined by measuring the angle between the direction of the (100)‐facet‐associated pores and a local tangent line at the nanoparticle core–shell interface, based on previously reported methodologies using lattice fringes as references.^{[ 48} , ^{49 ]} Statistical analysis (Figure 3g) showed that SNP@cMOF exhibits a randomly oriented distribution (average angle: 49.4° ± 17.3°), whereas SNP‐12@cMOF showed a significantly higher proportion of vertically aligned domains (average angle: 62.9° ± 18.8°). The intermediate seeding durations (SNP‐3@cMOF, SNP‐6@cMOF, and SNP‐9@cMOF) showed progressively increased vertical alignment, with average angles of 55.9° ± 23.1°, 60.3° ± 18.9°, and 59.5° ± 21.1°, respectively. Four‐point probe measurements of compressed SNP@cMOF papers (Figure 3h; Figure S18, Supporting Information) showed that the electrical conductivity increased systematically with improved nanocrystallite alignment. This measurement directly reflects the quality of the conductive network formed by the cMOF shell because the sulfur core is non‐conductive. The enhanced oriented growth correlated with improved electrical conductivity of the core–shell structures because of the more efficient electron transport across crystallite boundaries. In summary, the pre‐deposited metal–ligand seed layer on the SNP surface effectively directs the oriented growth of cMOF nanocrystallites, resulting in a highly ordered assembly that augments the structural integrity and electrical performance of the core–shell architectures.

2.4. Demonstration of cMOF Shell Functionality in Li─S Batteries

The core–shell nanoparticles were evaluated as advanced cathode materials for lithium–sulfur (Li─S) batteries to demonstrate the functionality of the cMOF shell encapsulating the sulfur core. Li─S batteries suffer from issues such as the shuttle effect, sluggish reaction kinetics, and cathode volume expansion, which impede sustainable high performance.^{[ 50 ]} We hypothesized that the cMOF shell could mitigate these challenges through its selective permeation, catalytic activity, and robust structural stability.

A beaker cell was assembled to monitor the confinement of polysulfides via changes in electrolyte color and assess selective permeation (Figure 4a). For SNP@cMOF (F0–M2.5), the loosely packed cMOF layer permits polysulfide leakage during discharge, as evidenced by the distinct color change.^{[ 51 ]} In contrast, the electrolyte in SNP‐9@cMOF (F0.2–M2.0) remains clear even after 16 h of discharge, showing that the densely packed and oriented Ni₃(HITP)₂ layer effectively blocks polysulfide leakage while preserving the Li⁺ transport pathways. UV–vis spectroscopy was performed on the electrolytes after discharge to validate these observations (Figures S19, Supporting Information). The UV–vis spectrum for SNP‐9@cMOF showed no significant peaks corresponding to dissolved polysulfide species, which contrasted with the pronounced absorption signals observed for loosely packed cMOF shells. These results support the visual evidence from the beaker cell tests, confirming the superior selective permeation performance of SNP‐9@cMOF.

Figure 4.

a) Visualization of the electrolyte color change after discharge in the beaker cell. b) SEM images of the electrodes after discharge and charge with the corresponding particle size distributions. Histograms demonstrate reversible particle diameter changes after cycling. Vertical dashed lines indicate the original SNP‐9@cMOF particle diameter (458.1 nm) and the theoretical diameter after 80% volume expansion (558.1 nm). c) Cyclic voltammetry profiles of SNP‐#@cMOF electrodes according to the seeding time. d) Correlation between the C₁ peak current, orientation of crystal domains, and electrical conductivity. e) Correlation between the C₂ peak, coordination number, and binding energy. f) Potentiostatic discharge curves of SNP@cMOF electrodes. g) Voltage profiles of SNP@cMOF electrodes, and h) corresponding rate capability from 0.2 to 5 C. i) Extended cycling performance of SNP‐9@cMOF electrodes at 3 and 5 C. Here, 1 C corresponds to 1672 mA g⁻¹. Shaded regions in (h) and (i) represent standard deviation.

SEM images of the SNP‐9@cMOF (F0.2–M2.0) electrodes (Figure 4b; Figure S20a, Supporting Information) showed that the core–shell structure remained intact after the discharge and charge cycles. Furthermore, SNP‐9@cMOF exhibited 72% volume expansion upon discharge. Upon recharge, the particle diameter returned to an average of 461.5 nm, closely matching the initial average diameter of 458.1 nm (Figure S15c, Supporting Information), with only a 0.7% deviation, confirming the high sulfur utilization and structural reversibility within the cMOF shell.^{[ 52 ]} By contrast, core–shell structures with a loosely packed cMOF layer lose their sulfur core during cycling, resulting in hollow, shell‐only structures (Figure S20b, Supporting Information). This comparison highlights the importance of cMOF microstructures in maintaining selective permeation and mechanical resilience against volumetric changes.

Cyclic voltammetry (CV) was conducted as a function of the seeding time to elucidate the interplay between shell microstructure and catalytic performance (Figure 4c). The sulfur contents of electrochemically evaluated SNP and SNP‐derived heterostructures are summarized in Figure S21 (Supporting Information). The CV profiles consistently exhibited characteristic C₁ and C₂ peaks corresponding to the reduction of S₈ to Li₂S, along with an anodic peak (A) representing the reverse reaction.^{[ 53} , ^{54 ]} The C₁ peak reflects the initial conversion of elemental sulfur (S₈) into high‐order soluble polysulfides, a process critically influenced by the intrinsic conductivity of the electrode material.^{[ 55 ]} Accordingly, the observed increase in the C₁ peak current with extended seeding time (Figure 4d) strongly suggests enhanced electron transfer efficiency arising from the improved crystallite orientation within the cMOF shell.

First‐principles density functional theory (DFT) calculations of the binding energies between Ni₃(HITP)₂ and Li₂S₄ at various coordination states were performed to demonstrate the catalytic role underlying the observed electrochemical performance, particularly the prominent C₂ peak corresponding to polysulfide conversion into Li₂S (Figure 4e; Figure S22, Supporting Information). Effective polysulfide‐to‐Li₂S conversion critically depends on overcoming the significant kinetic barriers. Atomically dispersed catalytic sites, such as transition‐metal centers, lower these barriers significantly by forming stable intermediate interactions with polysulfide species.^{[ 56 ]} Theoretical analysis confirmed that unsaturated Ni─N sites (coordination number of 3) within the cMOF shell yield asymmetric electron distributions, strongly facilitating the formation of robust Ni─S and Li─N bonds with polysulfides.^{[ 57} , ^{58 ]} Consequently, this coordination state exhibited the highest binding energies, aligning closely with the experimentally observed maximum in the C₂ peak current for the SNP‐9@cMOF electrode. As the coordination saturation increased (coordination number approaching 4), the reduced binding energies correlated directly with the diminished catalytic activity and lower C₂ peak intensities. Potentiostatic discharge experiments further validated this correlation, showing the highest Li₂S deposition capacity for SNP‐9@cMOF and confirming enhanced catalytic polysulfide conversion (Figure 4f).

Under various C‐rates during galvanostatic cycling, the SNP‐9@cMOF electrode maintained distinct reaction plateaus even at 5 C, whereas electrodes lacking crystallite alignment exhibited poorly defined plateaus because of the sluggish reaction kinetics and significantly higher polarization (Figure 4g; Figure S23, Supporting Information).^{[ 59 ]} The voltage profiles were analyzed to determine the Q₂/Q₁ ratio and polarization, revealing kinetic differences (Figure S24, Supporting Information). SNP‐9@cMOF consistently achieved the highest Q₂/Q₁ ratios at all tested C‐rates, with a 1.75 ratio at 5 C. (Figure S25a, Supporting Information). Moreover, SNP‐9@cMOF exhibits a notably smaller increase in polarization across the tested C‐rates compared to SNP and SNP@cMOF electrodes, reflecting superior electrochemical kinetics (Figure S25b, Supporting Information). Electrochemical impedance spectroscopy (EIS) further validated these findings, showing that the SNP‐9@cMOF electrode possessed the lowest charge‐transfer resistance (R_ct = 109.9 Ω), which was significantly lower than SNP@cMOF (125.4 Ω) and bare SNP (168.1 Ω). The introduction of the cMOF shell reduced R_ct by enhancing the conductivity, and further alignment of crystallites in SNP‐9@cMOF substantially boosted interfacial charge transfer (Figure S26, Supporting Information).^{[ 60} , ⁶¹ , ^{62 ]} Consequently, the rate capability tests showed that SNP‐9@cMOF exhibited superior cycling stability to bare SNP, delivering higher capacities (1263 mAh g⁻¹ at 0.2 C; 632 mAh g⁻¹ at 5 C) than SNP@cMOF and recently reported sulfur cathodes (Figure 4h; Table S2, Supporting Information). In particular, the core–shell nanoparticles preserved their original morphology after rate testing (Figure S27, Supporting Information). These performances were achieved despite the high sulfur loading (86.18 wt.%), underscoring the advantage of integrating conductive, catalytic, and selective permeability functionalities within a single architecture. In addition, the SNP‐9@cMOF electrodes showed excellent cycling stability, exhibiting high capacity retention rates of 78.2% at 3 C and 63.0% at 5 C after 500 cycles, corresponding to very low capacity fade rates of 0.044% and 0.074% per cycle, respectively (Figure 4i). At a high sulfur loading of 3.17 mg cm⁻², Li─S cells incorporating these electrodes delivered high capacities across a range of C‐rates, retaining 935 mAh g⁻¹ at 0.1 C after 100 cycles (Figure S28, Supporting Information). Translating these laboratory‐scale achievements into practical applications requires scalable and cost‐efficient synthesis methods. In this context, our synthetic approach offers significant advantages, including the use of minimal quantities of metal and ligand precursor and mild reaction conditions (65 °C), both favorable for industrial‐scale manufacturing. By optimizing reactor designs to maintain uniform temperature and reagent distribution, consistent, high‐quality coatings can be reliably produced. Thus, the proposed microstructural engineering approach bridges enhanced electrochemical performance, significantly broadening the applicability of cMOF‐based heterostructures toward the development of advanced multifunctional heterostructure platforms.

3. Conclusion

This study highlighted the pivotal role of microstructural optimization in unlocking the multifunctional potential of cMOFs as encapsulating shells in core–shell architectures. The packing density and alignment of Ni₃(HITP)₂ nanocrystallites were tuned precisely by systematically modulating the crystallization kinetics and incorporating a controlled seeding process to construct robust, multifunctional shells on spherical nanoparticle cores. The performance of the sulfur core–shell nanoparticles in Li─S batteries confirmed the efficacy of this chemical strategy because the optimized cMOF layer effectively confines polysulfides, enhances catalytic activity, and ensures exceptional cycling stability. Further exploration of the compatibility between diverse cores and cMOF shells will enable versatile microstructural designs, ultimately broadening their applicability and advancing their utility for next‐generation energy storage and environmental applications.

Conflict of Interest

The authors declare no conflict of interest.

Supporting information

Supporting Information

Park J. S., Yang S. M., Park D., et al. “Microstructural Engineering of Conductive MOF Nanocrystals for Multifunctional Heterostructures.” Small 21, no. 33 (2025): 21, 2504112. 10.1002/smll.202504112

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.