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. 2026 Jun 9;148(24):24627–24636. doi: 10.1021/jacs.5c21539

Pressurized Gas-Driven Elemental Redistribution Enables Ultrastable PtNi Catalysts for Heavy-Duty Vehicles

Xueru Zhao †, Kate Chen ‡, Zixian Jiao □, Hong Zhang □, Haoran Yu §, Xiaobo Chen ∥, Nebojsa Marinkovic ⊥, Lu Ma #, Xiaojing Wang ‡, Michael J Zachman §, David A Cullen §, Ping Liu †, Jacob S Spendelow ‡, Kotaro Sasaki †,*
PMCID: PMC13307270  PMID: 42263233

Abstract

Gas-driven element redistribution, characterized by the preferential enrichment of one element at the surface relative to the bulk, is frequently observed in multicomponent alloys. Using L10-ordered PtNi as a model system, we reveal that gas pressure plays a critical role in governing adsorption-driven surface composition during annealing in reducing gases: low pressure favors Pt surface segregation, while high pressure facilitates Ni surface enrichment. In this study, we developed a high-pressure nitriding (HPN) strategy that modulates the surface structure and composition of PtNi catalysts. The resulting HPN-PtNi exhibits enhanced performance and durability in membrane electrode assemblies for heavy-duty fuel cell applications, maintaining a high current density of 1.19 A cm–2 at 0.7 V after 90,000 voltage cycles. Through a combination of experimental and theoretical analyses, we reveal that the HPN process forms additional stabilizing Ni–N bonds and induces elemental redistribution with Ni surface enrichment and a Ni-deficient Pt subsurface. These modifications alter the atomic coordination environment of the ordered PtNi phase. This work presents a generalizable strategy to design robust and high-performing Pt-based catalysts by controlling gas-pressure-driven elemental redistribution and dopant incorporation.

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1. Introduction

The limited durability of Pt-based catalysts remains a significant obstacle to the commercialization of proton-exchange-membrane fuel cells (PEMFCs). − Alloying Pt with 3d metals can maintain a high catalytic activity while significantly reducing the noble metal content. The catalytic behavior can be tuned by alloying through various effects, making the Pt-3d metal alloys efficient oxygen reduction reaction (ORR) electrocatalysts. However, under harsh acidic heavy-duty fuel cell conditions, the excessive dissolution of transition metals (TMs) from the core can poison the membrane electrode assembly (MEA) ionomer and membrane, ultimately compromising performance. − Therefore, mitigating the leaching of TMs in Pt-based catalysts is essential to enhancing the oxygen reduction reaction (ORR) stability of Pt-based catalysts.

Surface segregation is a widely observed phenomenon in Pt-based alloy systems, where specific atoms migrate from the uniform bulk to enrich the surface. − This atomic redistribution occurs mainly due to differences in surface free energies among the alloy’s constituent elements, which are influenced by factors such as atomic radius and interactions with the environment, particularly surrounding gases. As a result, elements with energetically favorable configurations tend to accumulate on the surface. Gas-induced surface segregation has been extensively reported, particularly during thermal annealing treatments under different gaseous conditions, including oxygen, carbon monoxide, ,, and hydrogen atmospheres. , These processes typically occur under ambient or slightly reduced pressure conditions, revealing significant implications for catalytic activity, stability, and selectivity. However, a fundamental understanding of how increased gas pressure modulates elemental segregation remains elusive. This gap primarily arises from the limited availability of reactor systems capable of precisely controlled synthesis under high-pressure conditions (≥10 bar). Given the importance of gas pressure as a crucial factor influencing thermodynamic stability and kinetic pathways during surface restructuring, a detailed investigation of its role is vital, which is essential for the rational design of high-performance catalysts and nanostructured materials.

In this study, we introduce a gas-phase strategy using a high-pressure nitriding (HPN) reactor to regulate surface segregation in PtNi nanoparticles (NPs) by controlling the gas pressure. The HPN treatment induces structural reconstruction of the N-doped L10-ordered PtNi catalysts (denoted as HPN-PtNi), which were initially prepared by annealing under ambient pressure nitriding (denoted as APN-PtNi). Under HPN conditions, simultaneous Ni extraction and N incorporation collectively modulate the local bonding environment, optimize Pt–Ni interactions, and promote the formation of Ni–N bonds. As a result, the resulting HPN-PtNi catalyst achieved a high current density of 1.36 A cm–2 at 0.7 V at the beginning-of-life (BOL) in MEA testing under H2/15%-O2 conditions (selected to represent the average conditions in a HDV stack operating at high power), outperforming both the APN-PtNi and the commercial Umicore Pt/C benchmark. Furthermore, even at the end-of-life (EOL) after an accelerated durability test (ADT) consisting of 90,000 cycles from 0.6 to 0.95 V, the HPN-PtNi catalyst maintained a current density of 1.19 A cm–2 at 0.7 V, suggesting the ability to exceed the U.S. Department of Energy (DOE) target of 1.07 A cm–2 after 25,000 h equivalent durability testing. The high durability of HPN-PtNi can be attributed to its modified coordination structure and the robust “pinning effect” of Ni–N bonds, which effectively mitigate Ni dissolution and enhance the structural stability of the electrocatalyst in real HDV applications.

2. Results and Discussion

2.1. Structure and Morphology

N-doped L10-ordered PtNi NPs (APN-PtNi) were initially synthesized by impregnating Pt and Ni acetylacetonate precursors (Pt­(acac)2 and Ni­(acac)2) in a Ketjenblack (KB) support and reducing them through annealing in a flowing ammonia (NH3) atmosphere at ambient pressure. Following this, HPN-PtNi NPs were produced by thermally annealing APN-PtNi in NH3/N2 (10 vol % NH3/90 vol % N2) gas under high-pressure conditions of 800 psi (∼55 bar). Elemental maps extracted from scanning transmission electron microscopy–energy-dispersive X-ray spectrum images (STEM-EDS) show that the HPN-PtNi NPs have a Ni-rich surface (Figures a and S2), while the APN-PtNi NP exhibits a Pt-rich surface (Figure b). STEM–EDS line–scan profiles acquired across individual nanoparticles further corroborate the proposed surface compositional gradients (Figure S3). The representative high-resolution, high-angle annular dark-field (HAADF)-STEM images reveal distinct surface compositions for two catalysts: HPN-PtNi displayed a Ni-rich surface, including an outer NiO layer (Figure c), which is primarily formed via airborne oxidation following the HPN treatment. In contrast, APN-PtNi exhibited a Pt-enriched surface (Figure d), confirming that the HPN process extracts Ni atoms from the core of the APN-PtNi NPs (Figure e). The measured lattice spacings of 0.37 and 0.27 nm correspond to the (001) and (110) planes of the L10-type PtNi structure (Figures c–d and S4), consistent with the X-ray diffraction (XRD) results. As shown in Figure S1b, the XRD patterns of both APN-PtNi and HPN-PtNi can be indexed to the L10-ordered structure of PtNi (refer to PDF#65-2797). The presence of additional superlattice reflections at 33.5°, corresponding to the (110) plane, further confirms the formation of the ordered intermetallic phase in both catalysts. The ordering degrees for HPN-PtNi and APN-PtNi are 36.7% and 38.1%, respectively, which are calculated by the ratio of the (110) peak area to the combined area of the (111) and (200) peaks. The HAADF-STEM and TEM images (Figure S1c–e) show that both HPN-PtNi and APN-PtNi NPs are well dispersed on the KB carbon support. STEM-EDS analysis further reveals a higher nitrogen content in HPN-PtNi compared with APN-PtNi (Table S1), indicating that the HPN treatment incorporates additional N atoms into the PtNi NPs.

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Morphology and structure properties of the HPN-PtNi and APN-PtNi catalysts. (a) Representative STEM-EDS elemental maps of HPN-PtNi NPs. (b) Representative STEM-EDS elemental maps of APN-PtNi catalysts. (c) Representative atomic-resolution HAADF-STEM image of the HPN-PtNi particle. (d) Representative atomic-resolution HAADF-STEM image of the APN-PtNi particle. (e) Schematic illustration of the structure of HPN-PtNi and APN-PtNi. (f) XANES spectra at the Pt L3-edge of HPN-PtNi, APN-PtNi, and Pt foil. (g) XANES spectra at the Ni K-edge of HPN-PtNi, APN-PtNi, and Ni foil. (h) Pt L3-edge FT-EXAFS spectrum of HPN-PtNi and APN-PtNi with a first-shell fit together with Pt–Pt and Pt–Ni contributions. (i) Ni K-edge FT-EXAFS spectrum of HPN-PtNi and APN-PtNi with a first-shell fit together with Ni–Ni, Ni–Pt, and Ni–N contributions. (j) The coordination number and bond length of Pt–Pt, Pt–Ni, Ni–Ni, Ni–Pt, and Ni–N for HPN-PtNi and APN-PtNi.

The local structure and coordination environments of the two PtNi catalysts were investigated by using ex situ X-ray absorption spectroscopy (XAS). Compared to a reference Pt foil and Ni foil, both the X-ray absorption near-edge structure (XANES) spectra of the Pt L3-edge and Ni K-edge for HPN-PtNi and APN-PtNi shift to higher energies (Figure f,g), indicating the formation of a PtNi alloy. Additionally, the increased intensity of the white line (WL) for the Pt L3 (Figure f) and Ni K (Figure g) edges in HPN-PtNi, relative to APN-PtNi, suggests that the HPN process has tuned the electronic and structural properties of Pt and Ni, as discussed below. Fourier-transform extended X-ray absorption fine structure (FT-EXAFS) analysis at the Pt L3-edge (Figure h) and Ni K-edge (Figure i) was carried out to probe the local atomic structure of HPN-PtNi and APN-PtNi, with fitting results detailed in Table S2. The Pt L3-edge fitting included contributions from Pt–Pt and Pt–Ni interactions, while the Ni K-edge fitting accounts for Ni–Ni, Ni–Pt, and Ni–N contributions, all showing good agreement with the experimental spectra. Figure j clearly illustrates the differences in the coordination number (CN) and bond length (d) between the HPN-PtNi and APN-PtNi catalysts. The CN for Pt–Pt pairs in the APN-PtNi fitting is 5.4 ± 1.0, which is comparable to our previous findings for Int-PtNiN/KB (5.9 ± 0.3), supporting the formation of thin Pt shells. Following HPN treatment, the HPN-PtNi sample shows a reduced CN value for Pt–Pt pairs (3.6 ± 1.1) coupled with enhanced Pt–Ni and Ni–Ni coordination. These changes highlight the redistribution of Pt and Ni and the modified local coordination environment induced by the HPN process. We hypothesize that Ni atoms migrate from the core into the Pt-rich surface shell of APN-PtNi during the HPN process, partially replacing Pt–Pt bonds with Pt–Ni–Pt linkages, thereby decreasing CN­(Pt–Pt) and contracting the residual Pt–Pt distance. The formation of Ni-rich surface layers accounts for the increased CN­(Ni–Ni). This effect also leads to the smaller bond lengths for Pt–Pt and Pt–Ni interactions in HPN-PtNi compared to those of APN-PtNi (Figure j), which are known to induce lattice strain and establish a critical foundation for the enhancement of ORR performance. −

Moreover, the CN of Ni–N for HPN-PtNi, derived from the EXAFS fitting, is 4.1 ± 1.5 (Table S2), which is higher than that of APN-PtNi (3.0 ± 1.8), confirming the enhancement of Ni–N interactions during the HPN process. The increased N dopants are also evidenced by soft XAS at the Ni L3-edge. As shown in Figure S5, the Ni L3-edge spectra of HPN-PtNi show higher WL peaks and a slight positive energy shift in both total electron yield (TEY) and total fluorescence yield (TFY) modes compared to those of APN-PtNi, consistent with a higher Ni-oxidation state. These features suggest partial electron transfer from Ni to N atoms and concomitant enhancement of the d-band vacancy of Ni, , confirming the presence of more Ni–N bonds. We consider that this electronic redistribution further promotes electron transfer from Pt to Ni, reflected by the enhanced WL intensity of the Pt L3-edge spectrum in HPN-PtNi (Figure f). , X-ray photoelectron spectroscopy (XPS) spectra confirm the coexistence of Pt, Ni, and N in both catalysts, as shown in Figure S6. Compared with APN-PtNi, HPN-PtNi exhibits a positive shift in the Ni 2p binding energy and a negative shift in the N 1s peak, indicating enhanced electron transfer from Ni to N due to stronger Ni–N interaction. The increased N 1s peak intensity further confirms the higher N content in HPN-PtNi (Figure S6c). Meanwhile, the Pt 4f peak of HPN-PtNi shifts to a higher binding energy, suggesting a downshift of the Pt d-band center. This electronic modulation is expected to optimize oxygenated intermediate adsorption and contribute to enhanced ORR activity. These findings demonstrate that the HPN process not only introduces additional N atoms but also modifies the coordination environments of both Pt and Ni. These changes are expected to play a critical role in the ORR performance, as discussed below.

DFT calculations were performed to elucidate the effect of controlling the gas pressures on tuning the atomic structure of PtNi observed experimentally. Based on the XRD and EXAFS observations (Figures and S1), L10-Pt2Ni2N and L10-Pt2Ni2N2 were constructed to describe the core of synthesized APN-PtNi and HPN-PtNi particles, respectively (Figure S7c,d). The shell structure with three monolayers (MLs) of Pt (111)-like shell supported on the L10-Pt2Ni2N­(111) surface was used to represent the APN-PtNi configuration featuring a Pt-rich surface (Pt2Ni2N@Pt, Figures a and S7a,b). To capture the observed Ni-rich surface of HPN-PtNi (Figure a), an antisegregation of Ni atoms from the core–shell interface was also considered. Thus, the HPN-PtNi system was represented by an L10-Pt2Ni2N2(111) core and 1 ML of the Ni (111)-like shell, which was interconnected by 4 MLs of Pt interlayers (Pt2Ni2N2@Pt@Ni, Figures b and S7c,d).

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(a) Optimized structure, Pt2Ni2N@Pt, to describe the APN-PtNi model. (b) Optimized structure, Pt2Ni2N2@Pt@Ni, to describe the HPN-PtNi model. (c) Surface phase diagram at 773 K as a function of partial pressures of NH3 and H2 under a gas mixture of 10% NH3 (99.99%) and 90% N2. (d) Optimized structure of Pt–1H with 1 ML of hydrogen adsorbed on Pt2Ni2N@Pt. (e) Optimized structure of Ni-0.75N with 1 ML of nitrogen adsorbed on Pt2Ni2N2@Pt@Ni. (f) Surface phase diagram at 773 K as a function of NH3 and N2 partial pressures under 100% NH3 (99.99%). Note: The red and blue dashed lines denote the partial pressures of NH3 and N2, respectively; the black intersection in (c) is located in the thermodynamically preferred phase under experimental conditions (55 bar, 773 K, 10% NH3 + 90% N2) according to DFT calculation results; the red star is the lowest pressure that can reach the Ni-rich shell phase.

To evaluate the effect of pressures of NH3 and N2 applied experimentally on the thermodynamic stability of Pt-rich and Ni-rich PtNi surfaces, surface diagrams were mapped based on the DFT-calculated Gibbs free energy (G ph) over diverse Pt2Ni2N@Pt and Pt2Ni2N2@Pt@Ni core–shell model surfaces (Figure a,b and Table S3). Herein, the coverage of surface adsorbates resulting from NH3 and N2 dissociations, *H and *N at the Pt or Ni hollow sites (Figure S8), was found to vary depending on the applied environment. In a vacuum, Pt2Ni2N@Pt is more stable than Pt2Ni2N2@Pt@Ni due to the lower surface energy of Pt than Ni. At 10% NH3 (bal. N2) gas and 773 K (Figure c), increasing the partial pressure of NH3 promotes a shift in preference from Pt2Ni2N@Pt to *H-covered Pt2Ni2N@Pt (Pt-0.25H: 0.25ML *H, Pt-1H: 1ML *H, Figure d) and *N-covered Pt2Ni2N2@Pt@Ni (Ni-0.75H: 0.75ML *N, Pt-1N: 1ML *N, Figure e). Specifically, the most stable configuration is Pt–1H under ambient pressure with a formation energy of −1.48 eV (Figure c and Table S3), where the Pt-rich shell of Pt2Ni2N@Pt remains intact.

Under experimental HPN conditions (55 bar, 773 K; indicated by the black spot in Figure c), Ni-0.75N is preferred, corresponding to a formation energy of −2.25 eV (Figure c and Table S3). In this case, the *N coverage is sufficiently high to thermodynamically enable the anti-segregation of Ni from the core–shell interface to the topmost shell. This process is driven by the formation of stronger Ni–N interactions than that of Pt–N, which releases an energy of 0.79 eV (shown in Table S3) on going from Pt-0.75N and Ni-0.75N under HPN conditions (55 bar, 773 K, 10% NH3 + 90% N2). These results describe well the experimental observations by HAADF-STEM/EDS and EXAFS, showing the transition from the Pt-rich to Ni-rich shell on going from APN to HPN conditions (Figure ). By contrast, the lower *N coverage under the APN conditions hinders the displacement of Ni and the formation of Ni–N bonds; instead, the *H species dominate the Pt surface. Wherein, the difference in Ni–H and Pt–H interaction is not significant enough to drive the displacement of Ni from the inner core to the topmost shell, and Pt-1H is 0.84 eV more preferred than Ni-1H according to the formation energy (shown in Table S3). Notably, the minimum NH3/N2 pressure required to drive the transition from a Pt-rich to a Ni-rich shell configuration is approximately 6.0 bar (red star, Figure c).

We further evaluated the influence of the NH3/N2 gas ratio by considering 100% NH3. Herein, the Pt2Ni2N@Pt → Pt2Ni2N2@Pt@Ni transition can also occur. However, it requires an extremely high pressure of at least 2074 bar (the red star in Figure f). This elevated pressure threshold arises from the increased chemical potential of N2 in the absence of dilution by N2, which shifts the equilibrium boundaries toward Pt2Ni2N@Pt surfaces covered by *H. The DFT results collectively indicate that tuning of pressure and ratio of gas environments can effectively drive elemental redistribution in the PtNi system, thereby potentially modifying ORR performances under fuel cell operating conditions, as demonstrated below.

2.2. Electrochemical Performance

The electrochemical properties of HPN-PtNi were evaluated and compared with commercial Pt/C (46.4%) on both rotating disk electrode (RDE) and MEA measurements. Figure a presents the iR-corrected linear sweep voltammetry (LSV) curves in O2-saturated 0.1 M HClO4. The HPN-PtNi catalyst delivers a more positive half-wave potential (E 1/2) of 931 mV (Table S5) in comparison to APN-PtNi (927 mV) and commercial Pt/C (899 mV). As illustrated in Figure b, the ORR activity of HPN-PtNi, including the mass activity (MA) and specific activity (SA), surpasses both APN-PtNi and commercial Pt/C electrocatalysts. The corresponding electrochemically active surface areas (ECSA) for HPN-PtNi and APN-PtNi are 61.4 and 65.9 m2 g–1, respectively (Figure S9 and Table S5), which are slightly lower than that of Pt/C (83.3 m2 g–1). Notably, the slightly lower ECSA of HPN-PtNi may arise from minor surface reconstruction or residual NiOx species that were not completely dissolved in the electrochemical pretreatment and thus could partially decrease the electrochemically accessible Pt sites. Furthermore, analysis of the Koutecky–Levich (K–L) plots for HPN-PtNi at various potentials yielded an electron transfer number (n) of approximately 3.99 (Figure S9d). This indicates a four-electron pathway for the ORR, suggesting a high catalytic efficiency. These findings together corroborate that the HPN-PtNi catalyst possesses a strong ORR performance.

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Electrocatalytic and MEA performance of the HPN-PtNi catalysts. (a) ORR polarization curves of HPN-PtNi, APN-PtNi, and commercial Pt/C in O2-saturated 0.1 M HClO4. (b) Mass activity and specific activity of HPN-PtNi, APN-PtNi, and commercial Pt/C. (c) H2/air fuel cell performance of different catalysts at BOL and EOL. (d) Current density of different catalysts at 0.7 V under BOL and EOL. Test conditions: cathode/anode loading of 0.25/0.05 mgPt/cm2, 80 °C, 100% RH, 150 kPaabs, NC700 membrane, D2020 ionomer. (e) H2/15%-O2 fuel cell performance of different catalysts at BOL and EOL. (f) Current density of different catalysts at 0.7 V under BOL and EOL. Test conditions: cathode/anode loading of 0.25/0.05 mgPt/cm2, 90 °C, 85% RH, H2/15%-O2, 250 kPaabs.

For comparison, solid-solution HPN-PtNi (ss-HPN-PtNi) and solid-solution APN-PtNi (ss-APN-PtNi) were synthesized (Figure S10) and evaluated by RDE measurements. As presented in Figure S11, ss-HPN-PtNi exhibits higher ORR activity and durability in comparison to ss-APN-PtNi, indicating that the HPN treatment could enhance catalytic performance. Importantly, L10-ordered HPN-PtNi further outperforms its solid-solution counterpart, demonstrating that atomic ordering provides an additional performance benefit, consistent with our previous report.

To assess the practical applications of the HPN-PtNi catalyst in fuel cells, we fabricated MEAs designed for HDV conditions with a commercial Umicore Pt/C (Elyst Pt50 0550) and TEC10E40E Pt/C as the benchmark. The polarization curves of various catalysts at BOL and EOL (based on 90,000 voltage cycles) were recorded under different environments. In both the H2/air and H2/15%-O2 testing, the HPN-PtNi exhibits a higher MEA performance than the APN-PtNi and Umicore Pt/C (Figure c–f), which is in accordance with the results obtained under the RDE measurements. Specifically, the HPN-PtNi catalyst achieves a current density of 1.36 A cm–2 at 0.7 V under the H2/15%-O2 condition (90 °C, 85% RH, 250 kPaabs). In contrast, both APN-PtNi and Umicore Pt/C exhibit a current density of 1.29 A cm–2 at 0.7 V (Table S6). Even under H2-air conditions (80 °C, 100% RH, 150 kPaabs), the HPN-PtNi catalyst maintained a robust current density of 1.29 A cm–2 (Table S6), suggesting that the HPN process endowed improved MEA activity. After 90,000 ADT cycles, the HPN-PtNi catalyst demonstrates a current density of 1.10 A cm–2 at 0.7 V and 0.33 A cm–2 at 0.8 V under H2/air conditions (Figure d). Similarly, under H2/15%-O2 conditions, the HPN-PtNi achieves a current density of 1.19 A cm–2 at 0.7 V and 0.41 A cm–2 at 0.8 V (Figure f), further outperforming the DOE’s durability goal. When compared to the APN-PtNi catalyst, the HPN-PtNi exhibits enhanced stability across both H2/air and H2/15%-O2 environments. This improvement may be attributed to the optimized coordination environments and the increased formation of Ni–N bonds via the HPN treatment, which inhibits the dissolution of Ni atoms and enhances durability. Notably, both the MA loss at 0.9 V (H2–O2) and the hydrogen underpotential deposition (HUPD)-ECSA loss for HPN-PtNi are limited to 40% after 90k ADT cycles (Figure S12), which exceed the Umicore Pt/C (MA loss of 47% and ECSA loss of 50%) and TEC10E40E Pt/C (MA and ECSA loss of 69%).

2.3. Electrocatalytic Mechanism

To evaluate the performance enhancement of HPN-PtNi in MEA testing, we examined the morphology and structure of both HPN-PtNi and APN-PtNi catalysts at BOL (fully broken-in) and EOL (after 90k ADT cycles). As shown in Figure a, the average particle size of HPN-PtNi increases modestly from 3.4 ± 2.1 nm at the BOL to 4.5 ± 2.7 nm at the EOL based on TEM analysis. The representative images for each sample are shown in Figure S13. In contrast, APN-PtNi particles grow significantly, from 3.5 ± 2.7 nm to 5.8 ± 3.3 nm, suggesting that the optimized PtNi structure and additional Ni–N bonds in HPN-PtNi help prevent nanoparticle coarsening, contributing to enhanced durability. The HADDF-STEM images show that the HPN-PtNi NPs possess a thicker Pt shell at BOL, although the coating is not perfectly symmetric (Figures b and S14), compared to that of APN-PtNi (Figure c). This thicker Pt shell may provide enhanced protection for the interior structure. After 90,000 ADT cycles, the outer regions of HPN-PtNi remain Pt-rich, while the core maintains a half portion of the original Ni (12 at. % Ni at EOL), as shown in Figure b and Table S7. By contrast, APN-PtNi shows more Ni loss (5 at. % Ni at EOL) and less uniform distribution after cycling (Figure c). STEM-EDS analysis at BOL shows average compositions of Pt3.2Ni for HPN-PtNi and Pt1.5Ni for APN-PtNi (Table S7). This indicates that the Ni-rich surface layers on the HPN-PtNi surfaces were dissolved during the break-in process, resulting in a higher Pt/Ni molar ratio for HPN-PtNi compared to APN-PtNi. At EOL, the compositions change to Pt7.3Ni for HPN-PtNi and Pt19Ni for APN-PtNi (Table S7). The higher retained Ni content in HPN-PtNi at the EOL must be a consequence of its enhanced durability performance. Although the N contents in both catalysts at EOL were below the detection limit of STEM-EDS analysis, we consider that an increased number of N-dopants in the core of HPN-PtNi, coupled to the thicker Pt shell, suppressed the dissolution of Ni during MEA testing compared to APN-PtNi.

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(a) Particle size distribution of HPN-PtNi and APN-PtNi at BOL and EOL. The numbers of analyzed particles are 1273, 1184, 2711, and 851 for HPN-PtNi-BOL, HPN-PtNi-EOL, APN-PtNi-BOL, and APN-PtNi-EOL, respectively. (b) STEM-EDS elemental maps of HPN-PtNi at BOL and EOL. (c) STEM-EDS elemental maps of APN-PtNi at BOL and EOL. The scale bar in (b–c) is 2 nm. (d) Schematic of the structural evolution pathway from APN-PtNi to HPN-PtNi (step 1) and HPN-PtNi-BOL after the break-in process (step 2). Red arrows indicate the migration direction of Ni atoms, purple arrows indicate the N incorporation, and orange arrows indicate the dissolution of surface Ni atoms. (e) In situ Pt L 3-edge XANES spectra of the HPN-PtNi catalyst. The inset is the change of the Pt absorption edge peaks (Δμ) of the XANES spectra (relative to Δμ at 0.42 V) for HPN-PtNi and APN-PtNi as a function of potential obtained in 0.1 M HClO4. (f) In situ Ni K-edge XANES spectra of the HPN-PtNi catalyst in 0.1 M HClO4.

XANES analysis was performed to explore the structural evolution of the HPN-PtNi catalyst during MEA testing. After 90,000 cycles, the WL intensity of the Pt L3-edge for HPN-PtNi increased while that of the Ni K-edge decreased (Figure S15a,d), compared to those at BOL. This change reflects the formation of Pt oxide and the Ni dissolution, respectively. Partial Ni dissolution during MEA operation would reduce Ni–N moieties, implying a corresponding decrease in associated N content, consistent with the in situ XAS data shown below. FT-EXAFS results revealed an increase in the CN­(Pt–Pt) for both HPN-PtNi and APN-PtNi after 90k cycling (Figures S15 and S16), while the CN­(Pt–Ni) (or CN­(Ni–Pt)) decreased. This is likely to result from Ni dissolution and a relative increase in the Pt content during the ADT, causing Pt–Pt lattice expansion (Tables S8 and S9). It should be noted that although HPN-PtNi-EOL exhibits a slightly higher CN­(Pt–Pt) than APN-PtNi-EOL (9.4 ± 0.4 vs 8.8 ± 0.7) despite its lower overall Pt/Ni ratio, this difference is insufficient to support a clear structural conclusion, given the limited sensitivity of EXAFS-derived CN to subtle variations in particle size and local atomic configuration, particularly above ∼4 nm. An intriguing feature is the increase in Ni–N coordination for HPN-PtNi, which rises from 1.2 ± 0.7 at the BOL to 2.5 ± 1.4 at the EOL (Table S8). This enhancement in Ni–N coordination likely reflects an increased contribution from Ni–O bonding associated with Ni oxide formation; however, distinguishing between Ni–N and Ni–O contributions in EXAFS fitting remains challenging. Although the CN­(Ni–N) values may include contributions from Ni–O, we believe that HPN-PtNi possesses more Ni–N bonds than APN-PtNi (0.5 ± 0.3 at BOL and 0.6 ± 1.2 at EOL in Table S9), and it is therefore inferred that the stronger “pinning effect” by Ni–N bonding in HPN-PtNi effectively suppresses Ni dissolution, contributing to its improved durability during MEA testing.

Overall, the structural characterization at both the BOL and EOL reveals that the HPN-PtNi catalyst exhibits greater stability than APN-PtNi in the acidic MEA environment. As illustrated in Figure d, we propose a structural evolution pathway from APN-PtNi to HPN-PtNi and the subsequent changes after the break-in conditioning step. During HPN treatment (step 1), the Ni atoms diffuse toward the surface mainly via the vacancy mechanism, where Ni atoms exchange positions with neighboring vacant lattice sites, leading to the formation of a Ni-rich shell and an intermediate Pt-rich sublayer. In the Pt-rich sublayer, Pt–Pt linkages could be altered to Pt–Ni–Pt due to the ingress of outward Ni atoms, resulting in the decrease in CN­(Pt–Pt) and the contraction of Pt–Pt and Pt–Ni bond lengths (Figure and Table S2). Meanwhile, additional N dopants are incorporated into Ni in both the surface layer and core by occupying interstitial sites between metal atoms, leading to the formation of more Ni–N bonds, as confirmed by the increased CN­(Ni–N) (Figure and Table S2). Electron transfer from Ni to N, driven by their electronegativity difference, downshifts the Ni d-band center, which is demonstrated by the higher WL intensity of Ni K and L3 edges in XAS spectra of HPN-PtNi compared to APN-PtNi (Figures g and S5, respectively), since HPN-PtNi contains more N-dopants. In turn, this tunes Pt–Ni coupling and consequently lowers the Pt d-band center to optimize the adsorption strength of ORR intermediates, particularly for HPN-PtNi, in agreement with our previous finding. During step 2 (in the break-in process), partial leaching of surface Ni produces a Pt-rich shell and increased CN­(Pt–Pt), as evidenced by EXAFS fitting (Figure S15 and Table S8). Importantly, while APN-PtNi undergoes continuous Ni redistribution, HPN-PtNi retains stable Ni–Ni coordination, highlighting a “pinning effect” that suppresses Ni dissolution from the core. The optimized Pt–Ni coordination environment, together with the strengthened Ni–N interactions introduced by the HPN treatment, plays a pivotal role in suppressing Ni leaching and thereby ensures long-term structural stability in acidic electrolytes.

To further support the substantial stability of the HPN-PtNi catalyst as a fuel cell cathode, we conducted in situ XAS in a 0.1 M HClO4 solution. As displayed in Figure e, the XANES spectrum of the Pt L3-edge for HPN-PtNi gradually shifts toward higher energy as the potential increases from 0.42 to 1.12 V, accompanied by a slight rise in the WL intensity, implying Pt oxide formation during the ORR process. In contrast, the APN-PtNi NPs exhibit a much larger increase in the WL intensity under the same operational condition (Figure S18a). The difference becomes clearer when comparing the relative change in WL intensity [(ΔμE – Δμ0.42V)/Δμ0.42V] of the Pt L3-edge spectra for both catalysts across different potentials. The WL intensity for the HPN-PtNi catalyst rises at a higher potential than for the APN-PtNi catalyst (inset in Figure e), confirming the higher oxidation resistance of Pt in the HPN-PtNi catalyst. This evidence directly supports the enhanced stability of Pt within the HPN-PtNi system. Furthermore, the XANES spectrum of the Ni K-edge for the HPN-PtNi catalyst shows a negligible shift in edge features (Figure f), and the WL intensity only slightly decreases at a potential exceeding 1.02 V. In comparison, the APN-PtNi catalyst experiences a decrease in WL intensity at lower potentials (0.72 V), as shown in Figure S18b. The decrease could be attributed to decreasing N atoms accompanied by Ni dissolution, as shown in the decrease in Ni–N coordination numbers for APN-PtNi in Figure S17c. This suggests that the HPN-PtNi catalyst has improved resistance to acid corrosion and oxidation during the ORR.

The EXAFS fitting results (Figures S17–S18 and Tables S10–S12) reveal distinct differences in the coordination environment and bond lengths between HPN-PtNi and APN-PtNi. For instance, the CN of Pt–Pt, Ni–Ni, and Ni–N at 0.42 V for APN-PtNi is 5.6 ± 1.0, 5.3 ± 0.8, and 0.7 ± 0.3, respectively, while the HPN-PtNi shows a similar Pt–Pt coordination (5.2 ± 1.0), a reduced Ni–Ni coordination (4.0 ± 0.8), and an increased Ni–N coordination (2.1 ± 0.5). Taking into account that (i) HPN-PtNi NPs initially have a Ni-rich surface and (ii) these Ni atoms on the surface can easily be leached into an electrolyte during conditioning CV before the in situ XAS measurements, Pt-enriched surfaces on the HPN-PtNi NPs are subject to dissolution. Therefore, the coordination value of Pt–Pt obtained by the HPN-PtNi fitting at 0.42 V is consistent with the formation of thin Pt shells, which is similar to that of APN-PtNi. The lower CN­(Ni–Ni) in HPN-PtNi compared to that in the APN-PtNi catalyst confirms the fact that the composition of Ni in the core is modified through Ni extraction following the HPN process. Moreover, the increased Ni–N coordination in HPN-PtNi compared to APN-PtNi further verifies the incorporation of extra N atoms. The enhanced “pinning effect” from the increased number of Ni–N bonds strengthens the durability of the HPN-PtNi catalyst. Additionally, as the potential rises from 0.42 to 1.12 V, the CN and d for both HPN-PtNi and APN-PtNi remain largely unchanged except for the Ni–N for APN-PtNi (Figure S17c). This stability under varying potential highlights the robust structural integrity of the N-doped L10-ordered PtNi catalysts during the ORR.

3. Conclusion

This study illustrates the strategic application of gas-pressure engineering in the rational design of advanced Pt-based catalysts tailored to high-performance fuel cell systems. By utilizing a HPN strategy, we not only achieved precise control over the surface composition and electronic structure of L10-PtNi catalysts but also uncovered a robust pathway for enhancing both activity and durability through enhanced Ni–N interactions. Beyond achieving benchmark-level performance under HDV conditions, the HPN-PtNi catalyst highlights the importance of interfacial chemistry and atomic-scale restructuring in dictating long-term stability. More broadly, our findings underscore the potential of pressure-dependent gas-driven treatments as powerful and scalable tools for materials design. This concept can be extended to a wide range of multimetallic systems, offering new opportunities for fine-tuning surface energetics and unlocking performance in electrochemical energy conversion technologies.

Supplementary Material

ja5c21539_si_001.pdf (3.1MB, pdf)

Acknowledgments

This manuscript has been authored by employees/guests of Brookhaven Science Associates, LLC, under contract no. DESC0012704 with the U.S. DOE. The publisher, by accepting this manuscript for publication, acknowledges that the United States government retains a nonexclusive, paid-up, irrevocable, worldwide license to publish or reproduce the published form of this manuscript or allow others to do so for United States government purposes. Additional work was performed by the Million Mile Fuel Cell Truck (M2FCT) Consortium, with technology managers Greg Kleen and Dimitrios Papageorgopoulos, which is supported by the U.S. DOE, Office of Energy Efficiency and Renewable Energy, Hydrogen and Fuel Cell Technologies Office, under contract no. DE-SC0012704 (BNL). The DFT calculations were supported by the U.S. DOE, Division of Chemical Sciences, Geosciences and Biosciences, Office of Basic Energy Sciences, and carried out at Brookhaven National Laboratory, operated under contract DE-SC0012704 (FWP-CO40). The calculations used the high-performance SeaWulf computing system of Stony Brook Research Computing and Cyberinfrastructure and the Institute for Advanced Computational Science at Stony Brook University, which was made possible by a National Science Foundation grant (no. 1531492) and the computing resources at the Center for Functional Nanomaterials, a U.S. DOE Office of Science Facility at Brookhaven National Laboratory under contract no. DE-SC0012704. The research used the beamlines 7-BM (QAS) and 23-ID-2 (IOS) at NSLS II, U.S. DOE Office of Science user facilities. Work at the QAS beamline was supported in part by the Synchrotron Catalysis Consortium (SCC), U.S. Department of Energy Grant No. DE-SC0012335. STEM-EDS conducted as part of a user project at the Center for Nanophase Materials Sciences (CNMS), which is a US Department of Energy, Office of Science User Facility at Oak Ridge National Laboratory. The authors would like to thank Dr. Arephin Islam from Brookhaven National Laboratory for valuable discussions on XPS results.

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/jacs.5c21539.

  • Experiment methods; additional HAADF-STEM images; additional data for electrochemical and ex/in situ spectroscopic characterization; additional MEA results; and additional DFT calculation results (PDF)

The authors declare no competing financial interest.

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