Sub-4-nm Pt₄FeCoNiCu HEI nanocrystals via borophene-mediated co-anchoring for efficient and durable electrocatalysis

Abstract

Pt-based intermetallic compounds with atomically ordered arrangements are highly promising catalysts for the oxygen reduction reaction (ORR) in proton exchange membrane fuel cells. However, the mutual constraint between high-temperature ordering and small-size stability (especially at high metal loadings), as well as the corrosion of both metal particles and carbon supports under operating conditions, severely limit their mass activity and long-term durability. Here, we report a borophene-mediated multimetal site co-anchoring strategy to synthesize sub-4-nanometer high-entropy Pt₄FeCoNiCu intermetallic catalysts (MMCA-HEIMCs) supported on pristine carbon, which enables simultaneous anchoring of Pt and non-noble metals onto the carbon support. The strong metal-borophene interaction suppresses particle sintering during high temperature annealing (1000°C) and enables a high metal loading of 36 wt % at a small particle size. The resulting L1₀-ordered structure with a unique FeCoNiCu atomic stacking configuration is confirmed by electron microscopy and x-ray absorption spectroscopy. The catalyst achieves an exceptional H₂-air fuel cell peak power density of 1.055 watts per square centimeter and an ORR mass activity of 1.4 amperes per milligram of Pt at 0.9 volts in H₂-O₂ fuel cells. Owing to the high-entropy stabilization effect and the robust borophene co-anchoring effect, the catalyst retains 72% of its peak power density and 80% of its initial mass activity after 30,000 durability cycles. Moreover, the borophene interlayer mitigates Pt-catalyzed carbon corrosion, as verified by a 5000 startup/shutdown cycling test and online mass spectrometry. This work demonstrates a general strategy to overcome the activity-durability trade-off in multicomponent intermetallic catalysts through strong metal-support interactions.

INTRODUCTION

Proton exchange membrane fuel cells (PEMFCs) represent a clean and efficient energy conversion technology that holds great promise for mitigating greenhouse gas emissions and accelerating the global transition from fossil fuels to renewable energy sources (1–7). The oxygen reduction reaction (ORR) at the cathode is the kinetic bottleneck of PEMFC operation, characterized by sluggish reaction dynamics that necessitate high-loading Pt-based electrocatalysts (1–7). To address the high cost and limited durability of pure Pt catalysts, Pt-based alloys with non-noble transition metals have been extensively explored, where the ligand and strain effects induced by non-noble metals optimize the Pt electronic structure and weaken the adsorption of ORR intermediates, thus enhancing intrinsic catalytic activity (8–12). Compared with disordered Pt alloys, atomically ordered Pt-based intermetallic compounds (IMCs) with well-defined atomic arrangement and specific atomic ratios exhibit more pronounced ligand and strain effects, leading to further improved ORR activity (13–16). However, the formation of ordered IMC structures requires high-temperature annealing to overcome the energy barrier for atomic migration, which inevitably causes severe particle sintering and coarsening, which is an issue that is exacerbated at high metal loadings. This particle coarsening reduces the specific surface area and thus results in poor mass activity (MA) (17–21), creating a fundamental conflict between high-temperature ordering, high metal loading, and ultrafine particle stability in Pt-based IMC synthesis.

To alleviate particle sintering and maintain ultrafine sizes of Pt-based IMCs during high-temperature annealing, various strategies have been developed to strengthen metal-support interaction (MSI), including spatial confinement and modification of carbon supports (22–28). The spatial confinement via carbon encapsulation can effectively inhibit interparticle contact, but this approach inevitably blocks lots of active sites and impairs mass transport. Surface modification of carbon supports with heteroatoms (e.g., S, N, and P) or defects can form strong chemical bonds with Pt atoms, achieving size control of Pt-based IMCs. For example, sulfur-doped carbon supports have been used to synthesize Pt-based IMCs with average particle sizes <5 nm via robust Pt─S bonding, which resists sintering even at 1000°C (29). Nevertheless, most modified carbon supports suffer from compromised electron conductivity due to disruption of the structural integrity of the sp²-hybridized carbon skeleton, hindering electron transfer during the ORR process (30, 31). Nearly all existing MSI regulation strategies focus solely on anchoring Pt atoms while neglecting the immobilization of non-noble metal atoms in Pt-based IMCs. This oversight leads to the easy segregation and dissolution of non-noble metals, accounting for the challenges in preparing Pt-based IMCs with smaller sizes, particularly under high metal loading conditions, making it difficult to prepare ultrafine, high-loading Pt-based IMCs with uniform elemental distribution.

Durability is another critical performance metric for ORR electrocatalysts as important as activity for practical PEMFC applications. Although ordered Pt-based IMCs have a higher formation enthalpy and structural stability than disordered alloys, ultrafine Pt-based IMCs still suffer from poor operational durability due to their high surface energy and the corrosive PEMFC working environment (e.g., acidic electrolyte, high potential, and oxygen radicals) (5, 15). The construction of an ultrathin Pt shell on Pt-based IMCs is a widely adopted strategy to mitigate the dissolution of non-noble metals, yet substantial leaching of Pt and non-noble metals still occurs in such structure under long-term cycling (7, 8). Notably, the leaching of Fe²⁺ can trigger the Fenton reaction, which generates reactive oxygen species and damages the membrane electrode assembly (MEA) of PEMFCs, further limiting device lifetime (32). In addition to catalyst dissolution, carbon support corrosion is another major cause of ORR catalyst deactivation (1, 4). Modified carbon supports used for size control of Pt-based IMCs often have defective sp²-hybridized networks, which are prone to oxidative corrosion under high-potential conditions. This corrosion would cause catalyst particle detachment and agglomeration, leading to irreversible performance degradation (14, 23). Therefore, the simultaneous improvement of the intrinsic structural stability of ultrafine Pt-based IMCs, the MSI between IMCs and carbon supports, and the corrosion resistance of carbon supports has become the key to resolving the activity-durability trade-off in Pt-based ORR electrocatalysts.

In our previous work, we reported a hydrogenated borophene (HB)–assisted synthesis of Pt/B/C catalysts, where borophene and Pt form Pt-B coordination to anchor Pt particles and suppress sintering, which were further used as precursor to prepare high-entropy intermetallic catalysts (SPA-HEIMCs) through an impregnation and high-temperature sintering approach (33). Although these catalysts show promise, this preconstructed single–Pt site anchoring strategy suffers from a clear limitation—the metal-support interaction lacks anchoring for non-noble metal atoms, leading to their easy segregation and dissolution during cycling. To address this limitation, we herein develop an upgraded borophene-mediated multimetal site co-anchoring (MMCA) strategy that enables the simultaneous and dense anchoring of both Pt and non-noble metal (Fe, Co, Ni, and Cu) onto pristine carbon supports via borophene-bridged M-B coordination. This strategy allows the synthesis of sub-4-nm L1₀-ordered Pt₄FeCoNiCu high-entropy IMCs (MMCA-HEIMCs) with high metal loading of 36 wt % even after annealing at 1000°C. The borophene interlayer not only establishes robust MSI to suppress particle sintering and elemental dissolution but also physically isolates Pt from the carbon support to inhibit Pt-catalyzed carbon corrosion. The high-entropy effect of the FeCoNiCu atomic stacking further enhances the intrinsic structural stability of IMCs by increasing the vacancy formation energy of metal atoms. Benefiting from the synergistic effects of regulated Pt electronic structure, high-entropy stabilization, robust borophene-mediated co-anchoring, and enhanced carbon corrosion resistance, the MMCA-HEIMCs exhibit exceptional ORR activity and long-term durability in both rotating disk electrode (RDE) tests and practical PEMFC measurements. This work not only provides a class of high-performance Pt-based ORR electrocatalysts but also establishes a general strategy for the rational design of ultrafine, high-loading, and stable multicomponent intermetallic catalysts for energy conversion technologies.

RESULTS

Synthesis and structural characterization

MMCA-HEIMCs involved two synthetic steps: codeposition of metal salts with HB and a high-temperature triggered ordering process (Fig. 1A). The HB methanol suspension was prepared with the ion exchange method in previous literature (33, 34). First, the HB methanol suspension and carbon (Ketjen 600) were mixed evenly. H₂PtCl₆ and other precursors were mixed with a Pt/Fe/Co/Ni/Cu molar ratio of 4:1.55:1.05:1.05:1.05 in a methanol solution and further added the above suspension, stirring for 3 days at room temperature to form precursor/B/C compound. The additional Fe feedstock was used to facilitate the formation of a high-quality L1₀ PtM phase during the annealing stage (fig. S1). Subsequently, the compound was annealed at 1000°C in an H₂/Ar atmosphere, and the product was further subjected to acid treatment and annealing. The transmission electron microscopy (TEM) image shows that the obtained MMCA-HEIMCs represent an average diameter of 3.4 nm with a narrow size distribution (Fig. 1C). X-ray diffraction (XRD) pattern of MMCA-HEIMCs shows a highly ordered structure and high crystallinity similar to that of L1₀-PtFe IMCs, which adopted a cubic structure in space group Pm-3m (Fig. 1B). The total metal loading in the MMCA-HEIMCs was determined to be as high as 36 wt % based on inductively coupled plasma mass spectrometry (ICP-MS) analysis, achieved by using an optimized amount of HB (table S1 and fig. S2). The elemental mapping via energy dispersive x-ray (EDX) spectroscopy shows that Pt and non-Pt elements were distributed uniformly in individual MMCA-HEIMCs, revealing the homogeneously distributed elements at the nanoscale without phase segregation (fig. S3). Aberration-correction high-angle annular dark-field scanning transmission electron microscopy (AC-HAADF-STEM) was conducted to reveal the atomic structure of MMCA-HEIMCs. The AC-HAADF-STEM image displays a characteristic atomic resolution image of an individual nanoparticle, in which atoms with high and low contrasts exhibit close-packed staggered arrangement along the [001] direction (Fig. 1D). Because the atomic contrast is related to the atomic number, the bright spots represent the heavy Pt atoms, while lower contrast spots represent the randomly distributed non-noble metal atoms in MMCA-HEIMCs. Concurrently, the core-shell structure of MMCA-HEIMCs can be further found, which has a Pt₄FeCoNiCu intermetallic core and an ultrathin Pt-rich shell of about two atomic layers. The formation of the core-shell structure in MMCA-HEIMCs should be attributed to the surface segregation of Pt atoms with lower surface free energy and the subsequent acid treatment. The lattice fringe spacings are determined to be 0.27 and 0.37 nm, which correspond to the (110) and (001) planes of L1₀-Pt₄FeCoNiCu IMCs, respectively. Besides, atoms with high and low contrast in another representative particle further demonstrate the unique structure of the L1₀-PtM IMC (Fig. 1E). Combined with the AC-HAADF-STEM image and corresponding atomically resolved elemental mapping, the distribution of Pt atoms at bright spots and the random distribution of other metals at dark spots can be clearly observed (Fig. 1, F and G, and fig. S4). The above results effectively confirm the successful preparation of ultrafine MMCA-HEIMCs at both the nanoscale and atomic scales.

Fig. 1. Structure characterization of MMCA-HEIMCs.

(A) Schematic illustration of the preparation of MMCA-HEIMCs. (B) Powder XRD patterns of MMCA-HEIMCs. (C) TEM image and size distribution histogram of MMCA-HEIMCs. (D and E) AC-HAADF-STEM images of MMCA-HEIMCs. (F) Enlarge image of Fig. 1E and (G) corresponding atomically resolved elemental mapping of MMCA-HEIMCs. a.u., arbitrary units.

Electronic structure and coordination environment analysis

The x-ray photoelectron spectroscopy (XPS) technique was performed to evaluate the chemical valence state of the element in MMCA-HEIMCs. The Pt 4f spectra are composed of two main peaks 4f_5/2 and 4f_7/2 with Pt⁰ as the main state (Fig. 2A). Compared to Pt reference, the 4f peak position of Pt⁰ in MMCA-HEIMCs shifts to a lower binding energy (Δ = 0.15 eV), implying the electron transfer from non-noble metals to Pt. In contrast, the XPS signals of transition metals exhibit mixed valence states, such as Fe²⁺ and Fe⁰ (Fig. 2B and figs. S5 to S7). Moreover, the XPS spectra exhibit a weak B 1s signal in the energy range of 190 to 194 eV, which can be attributed to the combined contributions of B─C and B─O bonds (fig. S8). The formation of B─C bonds can be attributed to high-temperature annealing, confirming robust borophene bonding to the carbon substrate. To further get the fine electronic structure and coordination environment of each element in MMCA-HEIMCs, x-ray absorption fine structure (XAFS) measurements were collected to investigate. The normalized x-ray absorption near edge structure (XANES) spectra of five elements in MMCA-HEIMCs and reference samples were analyzed. The “white line” intensity at the edge of Pt L₃-edge is slightly lower than that of Pt foil, indicating an electron-rich state of Pt, which is consistent with the XPS analysis results (Fig. 2C). The EXAFS results show a shorter Pt-M distance in MMCA-HEIMCs than Pt-Pt in Pt foil, with a lattice compression of ∼5%, according to fitting result (Fig. 2E, fig. S9, and table S2). The compressive strain is proven to tailor the electronic state of the Pt surface through downshifting the d-band center, which is beneficial to ORR activity by weaker binding energy of intermediates (i.e., OH*) compared to pure Pt. Moreover, the K-edge XANES spectra of Fe and other transition metals in MMCA-HEIMCs are slightly higher than those of the metal foils, which is ascribed to their lower electronegativity with respect to Pt, causing the donation of electrons to neighboring Pt atoms (Fig. 2D and figs. S10 to S12). EXAFS analysis suggests that a longer Fe-M coordination in MMCA-HEIMCs than the Fe-Fe coordination in Fe foil (Fig. 2F) and similar elongation are observed for the Co-M, Ni-M, and Cu-M coordinations, which agrees well with the contraction of the Pt-M coordination (figs. S13 to S15). The wavelet transform (WT) analysis of Fourier-transformed EXAFS in the R space further identifies the bonding information of Pt/Fe/Co/Ni/Cu in MMCA-HEIMCs and the standard sample. Two wavelets maximum of Pt in MMCA-HEIMCs can be observed at y axis 2.5 and 2.0, attributing to the Pt-M and Pt-B coordinations (Fig. 2G and fig. S16). In addition, Fe-B, Co-B, Ni-B, and Cu-B coordinations for Fe, Co, Ni, and Cu elements of MMCA-HEIMCs are also observed, indicating that each of these metals participates in M-B coordination with the borophene interlayer (Fig. 2G and figs. S17 to S24). Notably, the comparative samples of SPA-HEIMCs and none-metal-sites anchoring HEIMCs (NMA-HEIMCs) synthesized via an HB-free method exhibit notably larger particle size (figs. S25 to S28). WT analysis of EXAFS confirms that the Pt-B coordination in the obtained SPA-HEIMCs is weak, and there is no detectable M-B (M = Fe, Co, Ni, and Cu) coordination (fig. S29). This highlights the critical role of the borophene-mediated MMCA strategy in forming M-B coordination and thereby suppressing particle coarsening.

Fig. 2. Characterization of electronic structure and coordination structure.

(A) Pt 4f and (B) Fe 2p XPS spectra of MMCA-HEIMCs. (C) L₃-edge XANES spectra of Pt and (D) K-edge XANES spectra of Fe in MMCA-HEIMCs. EXAFS spectra of (E) Pt and (F) Fe in MMCA-HEIMCs. (G) WT-EXAFS plots of Pt, Fe, Co, Ni, and Cu in MMCA-HEIMCs.

Electrocatalytic ORR performance

The electrocatalytic performance of the as-synthesized MMCA-HEIMCs for ORR was evaluated and benchmarked with commercial 20 wt % Pt/C using the RDE tests with a three-electrode system in 0.1 M HClO₄ electrolyte at room temperature. Cyclic voltammetry (CV) curves of MMCA-HEIMCs and commercial Pt/C were recorded in the N₂-saturated 0.1 M HClO₄ electrolyte (fig. S30). The electrochemical active surface area (ECSA) of MMCA-HEIMCs derived from hydrogen underpotential deposition (H_UPD) is 56.6 m² g_Pt⁻¹, higher than some previously reported Pt-based IMCs with larger particle sizes (table S3) (35–42). Linear sweep voltammetry curves were recorded in the O₂-saturated 0.1 M HClO₄ electrolyte. MMCA-HEIMCs shows a high half-wave potential (E_1/2) of 0.926 V versus reversible hydrogen electrode (RHE) while that of commercial Pt/C is 0.868 V versus RHE (Fig. 3A). Meanwhile, the MMCA-HEIMCs present improved MA and specific activity (SA) of 1.68 A mg_Pt⁻¹ and 2.97 mA cm⁻² at 0.9 V, which is 7.3-fold and 8.7-fold higher than those of Pt/C. The ORR performance of MMCA-HEIMCs is also better than those of NMA-HEIMCs and SPA-HEIMCs, highlighting the advantage of ultrafine IMC structures in MMCA-HEIMCs (fig. S31).

Fig. 3. Electrocatalytic performance.

(A) ORR polarization curves of MMCA-HEIMCs and Pt/C recorded in O₂-purged 0.1 M HClO₄ electrolyte (rotating speed: 1600 rpm) before and after 30,000 ADT. (B) Quantitative comparisons of the MA and SA of MMCA-HEIMCs and Pt/C before and after 30,000 ADT. (C) The leached metal concentration of MMCA-HEIMCs, MMCA-PtFeIMCs, and Pt/C in the electrolyte after 30,000 ADT. (D) Cathodic MA Tafel plots and corresponding MA at 0.9 V_iR-free for MMCA-HEIMCs and Pt/C in a H₂-O₂ fuel cell. (E) Comparison of MA and MA retention of H₂/O₂ MEA performance metrics for MMCA-HEIMCs versus the state of the art in the literatures. (F) H₂-air fuel cell polarization and power density plots for MMCA-HEIMCs and Pt/C. (G) Enlarged H₂-air fuel cell polarization for MMCA-HEIMCs and Pt/C. (H) The ESCA tendency during startup/shutdown cycling test of MMCA-HEIMCs and Pt/C. (I) The integrated area of DEMS signal of CO₂ monitored in an O₂-saturated 0.1 M HClO₄ electrolyte of MMCA-HEIMCs and Pt/C. BOL, beginning of life; EOL, end of life. NWs, nanowires.

Accelerated durability tests (ADTs) were also performed to assess the ORR durability of the as-prepared MMCA-HEIMCs from 0.6 to 1.0 V (versus RHE) in an O₂-saturated 0.1 M HClO₄ electrolyte. The MMCA-HEIMCs exhibited negligible degradation of the ECSA and a minor E_1/2 attenuation of only 4 mV after 30,000 cycles of ADTs. In contrast, the commercial Pt/C catalyst underwent a notable negative shift of 28 mV in E_1/2 (Fig. 3A and fig. S30). Furthermore, the MMCA-HEIMCs maintains 88.7 and 93.6% of the initial MA and SA, respectively, which is substantially higher than those of the commercial Pt/C (52.2/ 79.4%) (Fig. 3B). TEM was further used to characterize MMCA-HEIMCs after ADTs. TEM images show that the structure of MMCA-HEIMCs is well preserved, with no obvious particle detachment or agglomeration observed, while commercial Pt/C experienced severe particle agglomeration (figs. S32 and S33). This is primarily attributed to the intrinsic stability of Pt-based IMCs, the anchoring effect of the borophene-modified carbon on MMCA-HEIMCs, and the narrow particle size distribution that suppresses Ostwald ripening. ICP-MS was used to quantify the metal dissolution amount in the electrolyte after ADTs. Compared with commercial Pt/C, both MMCA-HEIMCs and carbon supported multimetal site co-anchored PtFe IMCs (MMCA-PtFeIMCs) exhibit markedly reduced Pt dissolution, which is attributed to the enhanced stability of ordered Pt-based IMCs. The dissolution levels of both non-noble metals and Pt in MMCA-HEIMCs is much lower than those in MMCA-PtFeIMCs, indicating that the high-entropy FeCoNiCu atomic layer may further improve the structural and compositional stability of MMCA-HEIMCs (Fig. 3C). Furthermore, the in situ ICP-MS analysis reveals that both the absolute amount and the increase in dissolution of non-noble metals and Pt in SPA-HEIMCs are much higher than that in MMCA-HEIMCs (fig. S34). This result robustly confirms that, compared to the single-Pt site anchoring strategy, the borophene-mediated MMCA strategy can firmly anchor both Pt and non-noble metals onto the pristine carbon support simultaneously via borophene-bridged M-B coordination, thereby further enhancing corrosion resistance and achieving higher ORR durability.

Given the excellent ORR performance of MMCA-HEIMCs in RDE tests, the catalyst was further integrated into a MEA to evaluate its practical performance in PEMFC. The MEA was prepared by directly coating the catalyst ink onto the proton exchange membrane via ultrasonic spraying. MMCA-HEIMCs or commercial Pt/C was used as the cathode with Pt loadings of 0.05 or 0.1 mg_Pt cm⁻², respectively, while commercial Pt/C served as the anode with a fixed Pt loading of 0.05 mg_Pt cm⁻². In H₂-O₂ text, the MMCA-HEIMC–based MEA achieves a maximum power density of 1.6 W cm⁻², higher than that of the commercial Pt/C (1.3 W cm⁻²; fig. S35). At 0.9 V, the MA of MMCA-HEIMCs reaches 1.4 A mg_Pt⁻¹, outperforming the 2025 US Department of Energy (DOE) target (MA >0.44 A mg_Pt⁻¹ at 0.9 V), commercial Pt/C (0.2 A mg_Pt⁻¹), and most reported Pt-based catalysts (Fig. 3, D and E, and tables S4 and S5) (18, 19, 23, 42–59). In H₂-air test, the maximum power densities of the MMCA-HEIMC–based MEA reaches 1.055 W cm⁻², whereas that of the commercial Pt/C based MEA is only 0.838 W cm⁻² (Fig. 3F). Furthermore, the cell voltage at a current density of 0.8 A cm⁻² is enhanced to 749 mV for the MMCA-HEIMC–based MEA, much higher than that of the Pt/C (666 mV) (Fig. 3G). Regarding the durability test conducted following DOE protocols, the MA of MMCA-HEIMCs decreases from 1.4 A mg_Pt⁻¹ at the beginning of life to 1.12 A mg_Pt⁻¹ at the end of life, corresponding to a mere 20% decay after 30,000 cycles, which meets the DOE 2025 target (MA <40% attenuation after 30,000 cycles). In comparison, the commercial Pt/C catalyst undergo obvious degradation with an MA attenuation of 55% after 30,000 cycles, showing poor stability (insert in Fig. 3D). Moreover, the MMCA-HEIMC–based MEA exhibits only a 28% decline in power density (to 0.758 W cm⁻²) and a minimal voltage drop of 39 mV at a current density of 0.8 A cm⁻² after ADTs. These results indicate that the ORR durability of MMCA-HEIMCs is superior to that of commercial Pt/C, which suffered a 50% reduction in power density (to 0.412 W cm⁻²) and a substantial voltage drop of 159 mV at 0.8 A cm⁻² (Fig. 3, F and G).

As is well known, carbon support corrosion can trigger particle detachment or aggregation. To investigate carbon corrosion in the catalysts, two complementary approaches were used. For the startup/shutdown cycling test, the electrode potential was modulated using a square wave, stepping between 1.0 and 1.5 V versus RHE with a 3-s hold time at each potential. A total of 5000 potential cycles were performed under an O₂ atmosphere, and intermediate CV measurements were performed after 100, 250, 500, 1000, 2500, and 5000 cycles to record the ECSA losses. It can be observed that the MMCA-HEIMCs retain most of its initial ECSA, whereas commercial Pt/C undergoes severe ECSA degradation (Fig. 3H and figs. S36 and S37). To gain fundamental insights into the structural degradation of carbon supports during cycling, Raman spectra were used. The intensity ratio of the D and G bands (I_D/I_G) was calculated as a structural descriptor for carbon disorder. Compared with MMCA-HEIMCs, a marked decrease in I_D/I_G can be observed in commercial Pt/C during cycling, which is attributed to the oxidative corrosion of amorphous carbon under high potential conditions (fig. S38). TEM analysis further confirms severe particle agglomeration in commercial Pt/C resulting from carbon corrosion, whereas MMCA-HEIMCs still maintain the initial morphology (figs. S39 and S40). To elucidate the origin of the enhanced carbon corrosion resistance of MMCA-HEIMCs, commercial Pt/C was also annealed at 1000°C and subsequently subjected to the same startup/shutdown cycling test. The annealed Pt/C exhibits a higher degree of graphitization and lower ECSA loss compared to the pristine Pt/C yet remains inferior to MMCA-HEIMCs (fig. S41). Previous studies have demonstrated that direct contact between Pt and the carbon support catalyzes carbon corrosion (60, 61). Building on the preceding analysis, the enhanced carbon corrosion resistance of MMCA-HEIMCs can be primarily attributed to the physical separation between MMCA-HEIMCs and the carbon support by the borophene layer. For the differential electrochemical mass spectroscopy (DEMS) test, the electrochemical operations were performed using a three-electrode system in 0.1 M HClO₄, using four CV protocols with different upper potential limits (UPLs) ranging from 1.0 to 1.6 V versus RHE. During potential cycling, the evolution of CO₂ progressively increases with increasing UPLs for both MMCA-HEIMCs and commercial Pt/C, indicating accelerated carbon corrosion at elevated potentials (Fig. 3I). No CO signal is detected, suggesting that the main product of carbon carrier corrosion is CO₂ (fig. S42). Notably, commercial Pt/C shows a much stronger CO₂ signal intensity, approximately twice that of MMCA-HEIMCs (fig. S43). This difference further widens with increasing UPLs. The DEMS results are consistent with those from the startup/shutdown cycling tests, further confirming the superior carbon corrosion resistance of MMCA-HEIMCs.

DFT calculation of the performance enhancement mechanism

Density functional theory (DFT) calculations were conducted to theoretically elucidate the mechanism underlying the enhanced ORR performance of MMCA-HEIMCs. Following model screening and structural optimization, the Pt (111) and HEIMCs (111) surfaces, which have the lowest energy, were ultimately selected as the final computational models (fig. S44). The density of states (DOS) of the d-band center for Pt atoms on the HEIMCs (111) surface downshifts by 0.7 eV relative to that on the Pt (111) surface, demonstrating a regulated electronic structure of Pt in the HEIMCs (111) surface (Fig. 4A). In addition, the work function of the HEIMCs (111) surface was calculated to be lower than that of the Pt (111) surface, indicating the strong electron supply capability of the former (Fig. 4B). To investigate the ORR mechanism in detail, the Gibbs free energies of the intermediates involved in the ORR process were calculated (figs. S45 to S47). The free energy diagrams at U = 1.23 V reveal that the rate-determining step (RDS) of ORR on the HEIMCs (111) surface shifted from *OH + H⁺ + e⁻ → H₂O [on the Pt (111) surface] to *O + H⁺ + e⁻ → *OH, with a lower free energy change (ΔG = 0.19 eV versus 0.57 eV). This indicates that ORR is more thermodynamically favorable on the HEIMCs (111) surface (Fig. 4C).

Fig. 4. Mechanistic insights by DFT.

(A) DOS of the d-band center for Pt atoms on the Pt (111) and HEIMCs (111) surfaces. (B) The work function of the Pt (111) and HEIMCs (111) surfaces. (C) Calculated free energy diagrams for ORR on the Pt (111) and HEIMCs (111) surface. (D) Charge density difference of *OH adsorption on the Pt (111) and HEIMCs (111) surfaces. (E) The vacancy formation energy of Pt or Fe atoms on the Pt (111), intermetallic PtFe (111), and HEIMCs (111) surfaces. (F) The binding energy of HEIMCs/borophene, HEIMCs/graphene, and Pt/borophene models.

To elucidate the mechanism of the reduced *OH desorption energy barrier on the HEIMCs (111) surface, we further calculated the charge density difference of *OH adsorption on the Pt (111) and HEIMCs (111) surfaces, as well as the projected crystal orbital Hamilton population (COHP) for the Pt─OH bond. The results show that the HEIMCs (111) surface transferred more electrons (0.17 e) to *OH with respect to the Pt (111) surface (0.06 e), illustrating the changed adsorption behaviors of *OH (Fig. 4D). Moreover, the integrated COHP value of the Pt─OH bond on the HEIMCs (111) surface was calculated to be 0.97 eV, which is lower than the 1.14 eV of the Pt (111) surface, indicating relatively weaker Pt─OH bonding (fig. S48). On the basis of the above analysis, the modified Pt electronic structure of the HEIMCs (111) surface can weaken the adsorption of strongly adsorbed *OH on Pt atoms, thereby optimizing the ORR pathway and accelerating ORR dynamics. This provides a theoretical explanation for the higher SA of MMCA-HEIMCs compared to Pt/C.

The stability of HEIMCs was also investigated via DFT calculations. The vacancy formation energy of Pt or Fe atoms on the HEIMCs (111), intermetallic PtFe (111), and pure Pt (111) surfaces were calculated (Fig. 4E). Compared to the Pt (111) surface, removing a Pt atom from the PtFe (111) and HEIMCs (111) surfaces is more energetically unfavorable. In addition, removing either a Pt atom or a Fe atom from the HEIMCs (111) surface is slightly more difficult than from the PtFe (111) surface. These thermodynamic analyses indicate that the ordered structure of IMCs and high-entropy atomic stacking can effectively stabilize Pt and non-noble metal atoms (represented by Fe) against dissolution. To further explore the interaction between borophene and HEIMCs, HEIMCs/borophene and HEIMCs/graphene models were constructed. The optimized structures reveal that in the HEIMCs/borophene model, the borophene layer undergoes obvious deformation due to strong interactions with the HEIMCs, whereas the graphene in the HEIMCs/graphene model retains its original planar two-dimensional structure (Fig. 4F and fig. S49). Binding energy (E_bind) calculations indicate that the E_bind of HEIMCs/borophene was 65.22 eV, which is higher than the 7.02 eV of HEIMCs/graphene. Additional computational results display that the E_bind of HEIMCs/borophene closely matched to that of Pt/borophene (65.20 eV), indicating that borophene can form strong M─B bonding with both Pt and non-noble metals. These results provide valid support for the enhanced interaction between borophene and HEIMCs, which enables more robust immobilization of HEIMCs and effectively suppressing their migration, detachment, and aggregation under both high-temperature annealing and electrochemical ORR conditions.

DISCUSSION

In summary, we successfully synthesized MMCA-HEIMCs with well-distributed, sub-4-nm small particles using a borophene-mediated MMCA strategy. The unique high-entropy L1₀ structure of MMCA-HEIMCs and the borophene-induced strong anchoring structure of MMCA-HEIMCs were confirmed by atomically resolved elemental mapping and EXAFS analysis. Toward acidic ORR, the MA and SA of MMCA-HEIMCs reached 1.68 A mg_Pt⁻¹ and 2.97 mA cm⁻² at 0.90 V versus RHE, respectively, which are 7.3-fold and 8.7-fold higher than those of commercial Pt/C. In addition, the MA of MMCA-HEIMCs retains 88.7% after 30,000 cycles, with a negligible E_1/2 loss. In H₂-O₂ fuel cells, the catalyst yields superior ORR MA (1.4 A mg_Pt⁻¹ at 0.9 V), and in H₂-air fuel cell, it delivered exceptional performance with a maximum power density of 1.055 W cm⁻². Owing to the high-entropy stabilization and robust borophene anchoring, 80/72% of the initial MA/power density was retained after 30,000 cycles. Startup/shutdown cycling and DEMS tests further demonstrate the outstanding carbon corrosion resistance of MMCA-HEIMCs, which is primarily attributed to the suppression of Pt-catalyzed carbon corrosion by the borophene interlayer. Theoretical studies reveal that the modified Pt electronic structure, the high-entropy stabilization of metal atoms, and the enhanced interaction between borophene and HEIMCs collectively contribute to superior ORR activity and durability of MMCA-HEIMCs. Briefly, this work adopts a strong MSI induced by borophene-bridged M-B coordination, which inhibits catalyst coarsening during high-temperature synthesis and stabilizes the robust anchoring of metal particles on the corrosion-resistant carbon support, thereby achieving the simultaneous improvement of MA and stability.

MATERIALS AND METHODS

Materials

MgB₂ (99.9%, 200 mesh) and ion-exchange resin (Amberlite IR120, H⁺ form) were obtained from Macklin Reagent. Methanol (analytical reagent, 99.9%) was purchased from Sinopharm Group Chemical Reagent. Graphitized carbon black (Ketjen black, ECP-600JD) was supplied by Cabot Carbon Ltd. Nafion solution (5 wt %) and HClO₄ were purchased from Sigma-Aldrich. H₂PtCl₆·6H₂O were purchased from Kunming Institute of noble Metals. Fe(NO₃)₃·9H₂O, Co(NO₃)₂·6H₂O, Ni(NO₃)₂·3H₂O, and Cu(NO₃)₂·6H₂O were obtained from Macklin Reagent. Commercial Pt/C (20 wt %) was purchased from Alfa Aesar Chemical Co Ltd. All the chemicals were analytical grade and used without further purification. Deionized water (18.25 MΩ/cm) was used throughout the experiments.

Synthesis of HB methanol dispersion

HB methanol dispersion was prepared by the ion exchange method (34). Typically, 60 mg of MgB₂ powder was added to 200 ml of methanol, followed by ultrasonication for 30 min in an ice bath. The prepared suspension was then added to a methanol suspension (100 ml) containing the ion-exchange resin (30 ml) under nitrogen at room temperature. After stirring at 250 rpm for 3 days under nitrogen atmosphere, the supernatant was evaporated on an oil bath at 343 K under nitrogen. The HB methanol dispersion was obtained until the volume of the supernatant evaporated to about 10 ml.

Synthesis of MMCA-HEIMCs

Twenty milliliters of HB solution and 0.4 g of carbon black (Ketjen black, ECP-600JD) were dispersed in 200 ml of methanol and stirred for 5 hours, and then H₂PtCl₆·6H₂O and all kinds of transition metal precursor were added with a Pt/Fe/Co/Ni/Cu molar ratio of 4:1.55:1.05:1.05:1.05, stirring 3 days at room temperature. After the reaction, the excess methanol was evaporated. Then, the black power was grounded, annealed in a 5% H₂/Ar atmosphere with a heating rate of 5°C min⁻¹, held at 1000°C for 1 hour, and then cooled to room temperature naturally. Last, the as-prepared alloys were treated in 0.1 M HClO₄ at room temperature for 2 hours and annealed at 400°C. For comparison, SPA-HEIMCs were synthesized by using presynthesized Pt/B/C as precursor, which is further added with transition metal precursors and annealed after grinding; NMA-HEIMCs were synthesized by using NaBH₄ to deposit all metal elements onto the carbon support, followed by high-temperature annealing.

Characterizations

TEM was taken on a JEM-2100 microscope. EDX spectroscopy was acquired on a Talos F200X microscope. Atomic-resolution STEM images were acquired on a Hitachi HF5000 microscope. Atomic-resolution STEM-EDX mapping was performed on a Spectral 300 (Thermo Fisher Scientific) equipped with double spherical aberration correctors and a Super-X G2 detector. XRD patterns were acquired on a LabX XRD-6100 x-ray diffractometer by using Cu Kα radiation source (λ = 1.5406 A), operating at 40 kV and 30 mA. The mass loadings of Pt in the as-prepared catalysts were determined by an inductively coupled plasma mass spectrometer with NexION 350D. XPS was carried out on a Thermo Fisher Scientific ESCALAB Xi⁺ spectrometer with an Al Kα radiator. Pt L-edge analysis was performed with Si (111) crystal monochromators at the BL14W1 beamlines at the Shanghai Synchrotron Radiation Facility (Shanghai, China). Before the analysis at the beamline, samples were pressed into thin sheets with 1 cm in diameter and sealed using Kapton tape film. The XAFS spectra were recorded at room temperature using a four-channel Silicon Drift Detector Bruker 5040, operated at 3.5 GeV with injection currents of 140 to 210 mA. We used IFEFFIT software (Athena) to correct the background signal, calibrate the energy scale, and normalize the intensity of shock wave. Reliable parameters for the high Z (Pt, Fe, Co, Ni, and Cu) and low Z (B and O) contributions were determined by multiple-shell fitting in R space with application of k³ and k¹ weightings in the Fourier transformations by software (Artemis) (62).

Atomic-resolution EDX mapping

Plasma cleaning of the sample and vacuum stabilization within the electron microscope are required before shooting to ensure optimal sample acquisition. The energy spectrum acquisition process requires the degree of sample drift to be controlled within a single atomic layer, as well as a clear zone axis and atomic structure. At a magnification of >5.1 Mx, a screen current of 100 pA, a dwell time of 5 μs, and a shaping time of 3 μs, the atomic-level energy spectra were acquired for roughly 10 to 15 min.

Electrochemical measurement

All ORR electrochemical tests were conducted on the CHI 760e equipped with a three-electrode system which used a glassy carbon electrode (area, 0.196 cm²) as the working electrode and saturated Ag/AgCl and a platinum foil as a reference and counter electrodes, respectively. The reference electrode potential was calibrated to the RHE potentials under hydrogen-saturated 0.1 M HClO₄ solution before the test. Catalyst ink was prepared by mixing 2 mg of catalyst with 400 μl of ultrapure water, 1580 μl of isopropanol, and 20 μl of Nafion solution (5 wt %) and ultrasonicating for 30 min. Then, the catalyst ink was dropped onto the glassy carbon RDE with a loading amount of 10.2 μg cm⁻² for electrochemical measurements. All the data have been corrected with ohmic-drop (iR) correction. In addition, we replaced the electrolyte immediately after testing each electrode to avoid being contaminated.

Fuel cell test

The catalyst inks were prepared by mixing catalysts with 5% Nafion dispersion solution, isopropanol, and deionized water (with an ionomer-to-carbon ratio of 0.6). The ink (0.1 or 0.05 mg_Pt cm⁻² for 40% Pt/C or MMCA-HEIMCs as the cathode catalysts) was ultrasonic sprayed onto 6.25-cm² proton exchange membrane (a GORE membrane, 8 μm), and then Pt/C catalyst ink was deposited (0.05 mg_Pt cm⁻², 40% Pt/C) on the anode side. Last, the MEA assembled with gas diffusion layer (SGL 22BC) for hydrogen fuel cell tests. The electrochemical performances were measured with a fuel cell tester (850e with 880 Frequency Response Analyzer, Scribner Associates, USA) with H₂ (0.5 liters min⁻¹), O₂ (0.5 liters min⁻¹), and air (2.0 liters min⁻¹) at 80°C, with 100% relative humidity. Following the guidance of DOE, the ADT was carried out at 80°C, 100 kPa_abs, and 100% relative humidity with H₂/N₂ for the anode/cathode for 30,000 cycles of square wave with each cycle holding the MEA at a voltage of 0.6 V for 3 s and then 0.95 V for 3 s.

DEMS test

An electrochemical workstation (760e) connected to a mass spectrometer (PrismaPro) was used for the DEMS analysis. The DEMS study was recorded during the CV protocol which comprised four slow cycles with different UPLs of 1.0 to 1.6 V versus RHE in oxygen-saturated 0.1 M HClO₄. A gold foil with a radius of 0.5 cm was used as the working electrode, on which ink (contain 30 μg_Pt) was dropped.

DFT calculations

All the DFT computations were performed using the Vienna Ab initio Simulation Package. Perdew-Burke-Ernzerhof functional within the generalized gradient approximation was used to model the exchange-correlation energy (63, 64). The ionic cores were described by projector-augmented wave pseudopotentials with an energy cutoff of 450 eV. Grimme’s DFT-D3 method was incorporated to implement the van der Waals correction (65). All calculations are spin polarized. The convergence tolerance of force and energy for each atom were 0.03 eV/Å and 10⁻⁴ eV, respectively. The k-points used the Monkhorst-Pack grid with the separation set to 0.05 Å. On the basis of the elemental ratios of Pt, Co, Ni, Fe, and Cu determined by ICP, 150 structures were generated by randomly substituting Fe atoms in a 3 × 3 × 3 PtFe supercell. After structural optimization, the model with the lowest energy was selected as the high-entropy alloy (HEA) structure (with 23 Fe atoms, 54 Pt atoms, 15 Co atoms, 8 Ni atoms, and 8 Cu atoms). On the basis of the results from HRTEM and XRD, we constructed the HEA (111) surface model, in which the top two atomic layers are allowed to relax. To avoid interlayer interactions, a vacuum layer of 15 Å is used in the Z direction. Furthermore, we replaced all atoms in the topmost atomic layer of the HEA (111) surface with Pt atoms to simulate a Pt skin. We also constructed Pt (111) and PtFe(111) surfaces as control groups.

The vacancy formation energy (E_v) is calculated using the following equation

$$E_v=E_{\text{defect}}-E_{\text{perfect}}+{\mathrm{\mu }}_{\text{atom}}$$ (1)

where E_defect is the total energy of the system containing the vacancy; E_perfect is the total energy of the perfect crystal (without defects); μ_atom is the chemical potential of the removed atom, taken as the average energy per atom in the perfect crystal.

The free energy (ΔG) of each ORR step is calculated according to the following equation

$$\mathrm{\Delta }G={\mathrm{\Delta }E}_{\text{ZPE}}+\mathrm{\Delta }E-T\times \mathrm{\Delta }S$$ (2)

where ΔE_ZPE is the zero-point energy at 298.15 K; ΔE is the binding energy of the intermediates; T is the experimental temperature (298.15 K); and ΔS is the entropy change (66).

The binding energy (E_bind) is calculated according to the following equation

$$E_{\text{bind}}=E\left(\mathit{AB}\right)-E\left(A\right)-E\left(B\right)$$ (3)

E(AB) is the total energy of the system after binding, and E(A) and E(B) correspond to the respective energy before binding.

Supplementary Materials

This PDF file includes:

Figs. S1 to S49

Tables S1 to S5