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Yunqing Kang, Shuangjun Li, [Ovidiu Cretu](https://orcid.org/0000-0002-1822-8172), [Koji Kimoto](https://orcid.org/0000-0002-3927-0492), Yingji Zhao, Liyang Zhu, Xiaoqian Wei, Lei Fu, Dong Jiang, Chao Wan, Bo Jiang, Toru Asahi, Dieqing Zhang, Hexing Li, Yusuke Yamauchi

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[Mesoporous amorphous non-noble metals as versatile substrates for high loading and uniform dispersion of Pt-group single atoms](https://mdr.nims.go.jp/datasets/0fc3df06-7c3d-4eff-a099-631ae53eeb99)

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Mesoporous amorphous non-noble metals as versatile substrates for high loading and uniform dispersion of Pt-group single atomsKang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e1 of 12M AT E R I A L S  S C I E N C EMesoporous amorphous non-noble metals as versatile substrates for high loading and uniform dispersion of Pt-group single atomsYunqing Kang1,2, Shuangjun Li3, Ovidiu Cretu4, Koji Kimoto4, Yingji Zhao5, Liyang Zhu5,  Xiaoqian Wei5, Lei Fu2, Dong Jiang5, Chao Wan2, Bo Jiang3, Toru Asahi5, Dieqing Zhang3,  Hexing Li3*, Yusuke Yamauchi6,7*Atomically dispersed Pt-group metals are promising as nanocatalysts because of their unique geometric struc-tures and ultrahigh atomic utilization. However, loading isolated Pt-group metals in single-atom alloys (SAAs) with distinctive bimetallic sites is challenging. In this study, we present amorphous mesoporous Ni boride (Ni-B) as an ideal substrate to uniformly disperse Pt atoms with tunable loadings (1.7 to 12.2 wt %). The effect of the morphology, composition, and crystal phase of the Ni-B host on the growth and dispersion of Pt atoms is dis-cussed. The resulting amorphous Pt-Ni-B mesoporous nanospheres exhibit superior electrocatalytic H2 evolution performance in acidic media. This strategy holds the potential to synthesize a diverse library of mesoporous amor-phous Pt-group SAAs, by leveraging functional amorphous nanostructured 3d transition-metal borides as sub-strates, thereby proposing a comprehensive strategy to control atomically dispersed Pt-group metals.INTRODUCTIONAtomically dispersed Pt-group metals have gained attention as a po-tential material for various catalytic applications because of their ultrahigh metal utilization and unique geometric structure (1–3). Although the high atomic unsaturated coordination of Pt-group metals affords them high catalytic activity when reduced to the atomic level, such as for forming metal single atoms (SAs), it also makes them prone to instability. One effective approach to address this limitation is combining SAs with specific supports that can modulate the catalytic activity by enhancing the interaction between the metal and support (4). Therefore, extensive efforts are currently underway for developing catalysts with a high loading of atomically dispersed noble metals to maximize the utilization of materials that have remarkable catalytic activity but are relatively expensive (5, 6). Thus far, various types of supports, including nitrogen-carbon com-posites (7, 8), metal oxides (9, 10), and metal sulfides (11) have been investigated extensively as ideal substrates for well-dispersed SAs. However, these supports form the oxidation states of the SAs. Another effective approach is forming single-atom alloys (SAAs), where the guest metal is isolated by another metallic substrate (12, 13), thereby enabling the SAAs to exhibit surface metallic states for both the host metal support and the atomically dispersed metal atoms and form bimetallic sites. Consequently, the guest and host metals mutually influence their catalytic performances via metal-metal interactions such as ligand, geometric, and/or lattice strain effects (14). The interaction patterns of adsorbed chemical species on individual Pt-group metal atoms within SAAs undergo a change because of the lack of robust interatomic binding locations (15), en-hancing selectivity and improving resilience against poisoning.Successful synthesis and effective catalytic performances have been demonstrated for various SAAs of M1/M2 (M1 and M2 represent an atomically dispersed noble metal and a nonprecious metal sub-strate, respectively), including Pt/Ni (16), Pt/Cu (17–19), Ir/Co (20), and Pd/Cu (21). A majority of published SAAs are confined to crys-talline NPs without specifically designed nanostructures (22). The loading amount of guest SAs in SAA catalysts is restricted, e.g., 0.25, 3.3, and 6.7 wt % for Pt/Ni (16), Pt/Co (23), and Pt/Cu (17), respec-tively, and increasing the loading content can leads agglomerate and form larger clusters or NPs with segregated phases from the host metal (17, 24). Thus, breaking the inherent crystal lattice constraints of the host metal in SAAs to achieve the uniform dispersion and high loading of guest Pt-group metals remains desired and challenging.In this study, we show that Ni boride (Ni-B) with an amorphous structure and mesoporous nanospheres (MNs) shape serves as an ideal substrate to achieve the uniform dispersion of atomically dis-persed Pt atoms produced through a galvanic replacement reaction (GRR) (table S1). This results in the formation of amorphous Pt-Ni-B (a-Pt-Ni-B) SAAs in the MNs form. The loading amount of Pt in a-Pt-Ni-B MNs can be controlled easily, reaching as high as 12.2 wt %, without changing the amorphous and mesoporous charac-teristics of the material. The effects of different morphologies, com-positions, and crystal phases (amorphous or crystalline) of Ni-B on the dispersion of Pt atoms and their catalytic properties are explored and highlighted. Further, we use the electrocatalytic hydrogen evo-lution reaction (HER) as a probe reaction to demonstrate that the prepared a-Pt-Ni-B MNs exhibit a low overpotential and high mass activity, making them highly effective catalysts. In addition, our ap-proach displays excellent versatility in synthesizing a wide range of 1Nanozyme Laboratory in Zhongyuan, Henan Academy of Innovations in Medical Science, Zhengzhou 451163, Henan, China. 2Research Center for Materials Nanoar-chitectonics, National Institute for Materials Science (NIMS), 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan. 3The Education Ministry Key Lab of Resource Chemistry, Joint International Research Laboratory of Resource Chemistry, Shanghai Frontiers Science Center of Biomimetic Catalysis, College of Chemistry and Materials Sci-ence, Shanghai Normal University, Shanghai 200234, China. 4Electron Microscopy Group, Center for Basic Research on Materials, NIMS, Namiki 1-1, Tsukuba, Ibaraki 305-0044, Japan. 5Faculty of Science and Engineering, Waseda University, 3-4-1 Okubo, Shinjuku, Tokyo 169-8555, Japan. 6Australian Institute for Bioengineering and Nanotechnology (AIBN) and School of Chemical Engineering, The University of Queensland, Brisbane, QLD 4072, Australia. 7Department of Materials Process Engi-neering, Graduate School of Engineering, Nagoya University, Nagoya 464–8603, Japan.*Corresponding author. Email: y.​yamauchi@​uq.​edu.​au (Y.Y.); hexing-li@​shnu.​edu.​cn (H.L.)Copyright © 2024 The Authors, some rights reserved; exclusive licensee American Association for the Advancement of Science. No claim to original U.S. Government Works. Distributed under a Creative Commons Attribution NonCommercial License 4.0 (CC BY-NC). Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024mailto:y.​yamauchi@​uq.​edu.​aumailto:hexing-li@​shnu.​edu.​cnmailto:hexing-li@​shnu.​edu.​cnhttp://crossmark.crossref.org/dialog/?doi=10.1126%2Fsciadv.ado2442&domain=pdf&date_stamp=2024-06-21Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e2 of 12amorphous SAAs MNs with precisely controlled compositions, in-cluding a-M-Ni-B (M  =  Pt, Pd, Rh, Ru, and Ir), a-Pt-M-Ni-B (M = Rh, Ru, and Ir), and a-Pt-Co-B (tables S2 and S3). This study presents a universal approach to control the dispersion and amount of noble SAs and the catalytic properties of the SAAs catalysts by manipulating the phase and morphology of the host metal.RESULTSMorphological and structural characterizationsThe amorphous Pt-Ni-B MNs (a-Pt-Ni-B MNs) were synthesized by a wet chemical reduction with the self-assembly of di-block copoly-mer micelles, following a GRR. First, bare amorphous Ni-B MNs (a-Ni-B MNs) with a well-defined mesoporous morphology of ex-posed pores and a disordered arrangement of local metal atoms (attributed to the incorporation of B in the Ni lattice) were synthe-sized based on our previously reported strategy (fig. S1) (25). Load-ing Pt with various contents was achieved through GRR by adding the desired H2PtCl6·6H2O aqueous solution to the as-formed me-sostructured a-Ni-B/micelle suspension (see Materials and Meth-ods for details). A typical sample with a Pt mass loading of 9.1 wt % [determined by inductively coupled plasma–optical emission spec-troscopy (ICP-OES); table  S1] on a-Ni-B MNs is denoted as a-Pt9.1-Ni-B MNs.Scanning electron microscopy (SEM) images showed a well-defined mesoporous structure for a-Pt9.1-Ni-B MNs, and the ex-posed pores were observed outside the nanospheres (Fig. 1A and B). The average sizes of the nanospheres and the pores were ~118 and ~11 nm (fig. S2), respectively, as illustrated in Fig. 1C. The po-rous morphology of a-Pt9.1-Ni-B MNs was observed using transmis-sion electron microscopy (TEM) and high-angle annular dark-field scanning TEM (HAADF-STEM) images. Figure 1 (D and E) shows that abundant mesopores extend from the exterior to the interior of the spherical structure, enabling the reactants to diffuse deeply. Mesopores offer not only a substantial increase in the active surface area but also the presence of rough surfaces, creating highly favor-able conditions for efficiently distributing active sites and facilitating adsorption, diffusion, and desorption of reactants, intermediates, and products. Selected-area electron diffraction (SAED) pattern shown in the inset of Fig. 1D indicates that one a-Pt9.1-Ni-B MN is amorphous without a long-range order. The a-Pt9.1-Ni-B MNs have a core-shell structure where the Pt atoms are distributed on the out-er surface and the Ni atoms are dominant in the core (Fig. 1F and G). Electron energy-loss spectroscopy (EELS) analyses identify the existence and uniform distribution of B atoms throughout a-Pt9.1-Ni-B MNs (fig. S3).The well-defined mesoporous structure of a-Pt9.1-Ni-B MNs was confirmed using small-angle x-ray scattering (SAXS) and the N2 adsorption-desorption isotherm. As shown in fig. S4A, a single peak located at q = 0.32 nm−1 indicates the periodicity of mesopores in a-Pt9.1-Ni-B MNs. The pore-to-pore distance is calculated as ~19.6 nm. The N2 adsorption-desorption isotherm (fig. S4B) indi-cates the mesoporous structure of a-Pt9.1-Ni-B MNs, which have a Brunauer-Emmett-Teller (BET) surface area of ~32.5 m2 g−1.Powder x-ray diffraction (XRD) profiles of a-Pt9.1-Ni-B MNs show a broad peak without any diffraction peaks of face center cubic (fcc) Pt, indicating the amorphous alloying of Pt and Ni-B (Fig. 2A). As shown in Fig. 2B, no well-resolved lattice fringes are observed in the high-resolution HAADF-STEM images, thereby verifying the amorphous structure without a discernible long-range order in a-Pt9.1-Ni-B MNs, which is consistent with the XRD observations. The rich atomically dispersed Pt atoms can be distinguished from the edge region of the mesoporous walls (Fig. 2B and fig. S5), wherein some possible Pt atoms are circled in red (Fig. 2C). In contrast, the phase structure of the a-Pt9.1-Ni-B MNs was converted to the crys-talline phase in the reference crystalline Pt9.1-Ni-B (denoted as c-Pt9.1-Ni-B) MNs prepared via annealing (400°C) a-Pt9.1-Ni-B under Ar atmosphere for 2  hours (fig.  S6). The c-Pt9.1-Ni-B exhibits a single-phase alloy structure without phase segregation even after in-creasing the annealing temperature to 900°C (fig. S7), confirming Fig. 1. Structural characterization of core-shell a-Pt9.1-Ni-B MNs. (A) Low-magnification scanning electron microscopy (SEM; scale bar, 1 μm), (B) high-magnification SEM (scale bar, 100 nm), (C) illustration of a single MN, (D) transmission electron microscopy (TEM) image (scale bar, 100 nm), (E) high-angle annular dark-field scanning TEM images (HAADF-STEM; scale bar, 50 nm), (F) elemental energy-dispersive x-ray spectroscopy (EDS) maps (scale bar, 100 nm), and (G) corresponding normalized line-scan profiles of a-Pt9.1-Ni-B sample. The inset of (D) shows the SAED pattern from a single MN (scale bar, 5 nm−1).Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e3 of 12that Pt atoms are highly dispersed into the lattice of the host Ni-B. These data indicate that the bonding between heterometals, that is, Pt-Ni in our case, is stronger than that of homometallic bonding (i.e., Pt-Pt), which ensures no bonding between Pt atoms.X-ray photoelectron spectroscopy (XPS) and x-ray absorption spectra (XAS) were used to investigate the electronic state and local coordination environment of Pt and Ni atoms of a-Pt9.1-Ni-B MNs. The Pt 4f7/2 and 4f5/2 main peaks of a-Pt9.1-Ni-B are loaded at 71.3 and 74.6 eV, respectively, which confirm its dominant nature as me-tallic Pt (Fig. 2D). The peaks located at 72.7 and 76.1 eV indicate the coexistence of Pt2+. Both metallic and oxidized states of Ni are ob-served (Fig. 2D and fig. S8A). Owing to the strong attraction for O, B atoms exist in an oxidized state on the surface (i.e., B-O) (fig. S8B). Compared to the homemade Pt nanoparticles (NPs), the peaks be-longing to Pt 4f in a-Pt9.1-Ni-B MNs shift negatively (fig. S9), con-firming the effect of Ni-B on the electronic structure of Pt. XAS offers an additional understanding of the characteristics of the Pt and Ni species of a-Pt9.1-Ni-B MNs. As shown in Fig. 2E, the Pt LIII-edge x-ray absorption near edge structure (XANES) data confirm that the valence state of Pt is close to the standard Pt foil with a weaker intensity of the white line, suggesting the negatively charged nature of Pt in a-Pt9.1-Ni-B MNs. The slightly electron-enriched Pt species was reported to be favorable for HER (26). The Ni K-edge, which is located between Ni foil and NiO, provides evidence for the partial oxidation of Ni in a-Pt9.1-Ni-B (fig. S10A). These observa-tions align well with the findings from the XPS results. The forma-tion of atomically dispersed Pt atoms on a-Ni–B MNs is confirmed by the absence of a Pt-Pt scattering path in the Fourier trans-forms extended x-ray absorption fine spectra (EXAFS) of Pt LIII-edge (Fig. 2F). One prominent peak (at ~2.16 Å) positioned between the Pt foil and standard PtO2 sample indicates the formation of the Pt-Ni alloy, aligning with the reported EXAFS data showing the pres-ence of isolated Pt atoms on Cu (17) or Ru (27) NPs. The presence of single Pt atom on the a-Ni-B MNs is supported by the fitting results, as shown in fig. S11 and table S4. The data reveal that a-Pt9.1-Ni-B MNs show a Pt-Ni coordination number (CN) of ~7.33 at 2.57 Å, which is different from the Pt-Pt CN of 12 at 2.77 Å in the Pt foil. In addition, the absence of detectable Pt-Pt, Pt-O, and Pt-B coordina-tion contributions in a-Pt9.1-Ni-B MNs suggests that Pt atoms are predominantly distributed as isolated entities, surrounded by Ni at-oms, as evidenced by the wavelet transform (WT) contour plot (fig. S12). The amorphous nature of Ni-B is retained after the forma-tion of the Pt-Ni SAA, as confirmed by absence of the long-range order in the EXAFS curve (fig. S10B) and the corresponding fitting data with the lower CN and shorter Ni-Ni distance compared to that of the Ni foil (table S4). The XAS results validate the existence of the Pt-Ni SAAs, characterized by the presence of atomically dispersed Pt atoms on a-Ni-B MNs.GRR process studyAchieving a uniform atomic-level dispersion of Pt atoms on Ni nanostructures is challenging because a GRR occurs rapidly during the synthesis process when metallic Ni nanostructures mix with Pt salts. The rapid GRR and large Ni-Pt lattice mismatch leads to the random atomic mixing of Ni and Pt, as well as a loss of control over the nanostructure (28–30). Our approach achieves the atomic-level dispersion of Pt effortlessly without the need to stringently control GRR conditions by adjusting the coordination environment of Ni, Fig. 2. Structural and surface characterizations of a-Pt9.1-Ni-B MNs. (A) Powder x-ray diffraction (XRD) pattern. (B) High-resolution HAADF-STEM image (scale bar, 5 nm). (C) Enlarged HAADF-STEM image (scale bar, 1 nm) of the selected area of (B). The bright spots (circled in red) represent the possible positions of atomically dis-persed Pt atoms. (D) XPS spectra of Pt 4f and Ni 3p. Pt LIII-edge (E) x-ray absorption near the edge structure and (F) Fourier transform extended x-ray absorption fine structure of a-Pt9.1-Ni-B MNs with standard references (shoulder peaks located between 1.5 and 2.0 Å of both Pt foil and a-Pt9.1-Ni-B originate from the slight surface oxida-tion of Pt during the ex situ XAS measurement process).Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e4 of 12i.e., by forming short-range ordered and long-range disordered amorphous structures. Figure 3A shows an illustration of the syn-thesis route of a-Pt-Ni-B MNs using the proposed strategy. Ini-tially, mesostructured Ni-B/micelles composites were obtained using polystyrene-poly(ethylene oxide) (PS-​b-PEO) block copolymer mi-celles as the pore-directing agent. Pt atoms were then inserted into the parent Ni-B sacrificial template by GRR, forming the highly well distribution of Pt atoms, as shown in the SEM–energy-dispersive x-ray spectroscopy (SEM-EDS) results for the a-Pt9.1-Ni-B MNs (fig. S13). Owing to the gap of standard reduction potentials (E0) between the Ni2+/Ni and [PtCl6]2−/Pt pairs (the E0 of Ni2+/Ni, [PtCl6]2−/[PtCl4]2−, and [PtCl4]2−/Pt is −0.26, 0.73, 0.76 V versus standard hydrogen electrode, SHE, respectively) (table S2) (31–33). Pt ions are rapidly reduced to metallic Pt atoms when contacted with metallic Ni atoms of a-Ni-B, thereby causing metallic Ni atoms to be oxidized to Ni2+ asPt atoms preferentially replace the outer Ni atoms because of the first contact with Ni on the surface of MNs; however, some Pt atoms probably penetrate deeper into the a-Ni-B MNs because of the presence of abundant interconnected channels within the mesopo-rous shape. The preferential replacement of the surface metallic Ni atoms of a-Ni-B MNs results in the final core-shell structure of a-Pt9.1-Ni-B with a Pt-rich shell, as identified from the above HAADF-STEM data (Figs. 1, G and H, and 2B; and fig. S5).Further, we investigated the effect of the synthesis conditions, in-cluding reaction time, temperature, and concentration/type of Pt precursor of GRR on the final morphology and disordered structure of a-Pt-Ni-B MNs. Changes in the content of Ni and Pt elements in different GRR time demonstrate that the occurrence of the GRR be-tween [PtCl6]2− ions and metallic Ni atoms (Fig. 3B). For the syn-thesis of a-Pt9.1-Ni-B MNs, the content of Pt in MNs increases gradually with an increase in GRR time, whereas the Ni content de-creases gradually, retaining the similar mesoporous morphology, as determined by the SEM and SEM-EDS results (Fig. 3B and fig. S14). The reaction temperature plays an important role in redox reactions. Predictably, the GRR rate between Ni and Pt atoms increases with an increase in the GRR temperature. Therefore, increased reaction temperatures (i.e., 40°, 60°, and 70°C) were applied to explore the effect of the reaction rate on the Pt dispersion and structure. As shown in fig. S15, the mesoporous morphology, well-distribution of Pt, and amorphous natures do not change with an increase in the GRR temperature. Unlike the sample prepared using GRR, the core-duction of the Ni2+ and [PtCl6]2− precursors with dimethylamine borane (DMAB) as a reducing agent leads to a formation of aggre-gated NPs and separated phases (fig. S16).In amorphous transition-metal borides, the wide range of adjust-able compositions between the metal and B offers promising oppor-tunities for adjusting the morphology, phase structure, electronic properties, adsorption performance, etc. of the catalysts (34–36). Therefore, in addition to various GRR preparation parameters, the effect of the content of B in the host a-Ni-B MNs on the final mor-phology of a-Pt-Ni-B was also evaluated. During the preparation process of a-Ni-B MNs, DMAB is used as both a primary reducing agent and a B source (36). The content of B in Ni-B can be adjusted [PtCl6]2−(aq)+2Ni(s)→Pt(s)+2Ni2+(aq)+6Cl−(aq) (1)Fig. 3. Synthesis of a-Pt-Ni-B MNs through GRR. (A) Illustration of the synthesis route, including the micelle-induced wet chemical reduction synthesis of mesostruc-tured a-Ni-B/micelles and the following loading of Pt through GRR. (B) Evolution content of Ni and Pt in the a-Pt-Ni-B MNs during GRR at different reaction times. (C) Re-lationship between the concentrations of Pt precursor (i.e., [PtCl6]2−) and final mass ratios of Pt/(Pt + Ni). (D) XRD patterns of a-Pt-Ni-B MNs with different loading contents of Pt: (I) a-Ni–B, (II) a-Pt1.7-Ni-B, (III) a-Pt3.9-Ni-B, (IV) a-Pt9.1-Ni-B, and (V) a-Pt12.2-Ni-B MNs. Error bars in (B) and (C) that correspond to the SD of three independent measure-ments.Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e5 of 12by changing the concentration of DMAB, which does not affect the amorphous nature of the obtained samples (fig.  S17A). Although there is no obvious aggregation of Pt NPs when the DMAB concen-tration decreases from 1.0 to 0.125 M, the surface of the MNs is wrapped by a flocculent layer at low DMAB concentrations (i.e., 0.125 and 0.25 M). At high DMAB concentrations (i.e., 0.5 and 1.0 M), the surface of the MNs is kept clean (fig. S17, B to E). These results indicate that a higher concentration of DMAB can result in more metal-metalloid (i.e., Ni-B) bonds with more metallic Ni on the surface, thereby ensuring sufficient Ni atoms to be replaced by Pt in the GRR. In this study, 0.5 M DMAB was used to prepare a-Ni-B unless stated otherwise.The loading content of Pt is controlled easily by varying the con-centrations of H2PtCl6·6H2O (1.0 ml of 2.0, 4.0, 8.0, and 16.0 mM aqueous H2PtCl6·6H2O). Four a-Pt-Ni-B samples with increased Pt mass loadings of 1.7, 3.9, 9.1 and 12.2 wt % are denoted as a-Ptx-Ni-B, where x displays the weight percent of the Pt loading (Fig. 3C), as determined by ICP-OES (table  S1). The mesoporous structure remains as the increased Pt loading ratio (fig.  S18). Further, with more Pt replaced by Ni, the broad peak of XRD gradually shifts to-ward the lower angle, demonstrating the successfully incorporated Pt with a larger lattice within the disordered matrix of Ni-B (Fig. 3D). The proportion of guest metal (M1) to host metal (M2) in SAAs is low to ensure that M1 atoms do not bind with each other (5). Although crystal M2 may contain some structural defects, such as segregation and dislocations, the strong metallic bonding (M2-M2) interactions in the host metal M2 makes it challenging to accommo-date a large loading of isolated guest M1. Thus, the loading of M1 in most SAAs is usually less than 8 wt % (22). However, in our mate-rial, the Pt/(Pt + Ni) ratio in a-Pt9.1-Ni-B MNs can easily reach 9.7 wt % (even 13.0 wt % in a-Pt12.2-Ni-B MNs) (Fig. 3C), which is considerably higher than that commonly observed in other SAAs. The aberration correction HAADF-STEM image confirms isolated Pt atoms on the Ni-B of a-Pt12.2-Ni-B MNs (fig. S19); however, fur-ther increasing the loading amount of Pt results in the appearance of some microcrystalline Pt, as confirmed by HAADF-STEM images and the XRD pattern (fig. S20). We can reasonably assume that the metallic amorphous state of M2 with local disordered atomic ar-rangement possesses the ability to break away from the lattice constraints observed in conventional crystalline metals, thereby achieving a substantial loading of atomically dispersed M1.Reference hosts were prepared using crystalline Ni-B, crystalline Ni, and nonporous amorphous Ni-B as sacrificial matrices for load-ing Pt to explore the effect of mesoporous morphology and long-range disordered arrangement of atoms of Ni-B on the GRR process. Amorphous alloys are temperature-sensitive materials susceptible to transitions to the crystalline state at high-temperature treatments. Therefore, the crystalline mesoporous Ni-B (c-Ni-B) sample was synthesized via the thermal treatment of the prepared a-Ni-B for 2 hours at 400°C under Ar atmosphere. Further, the crystalline Ni nanothorn (c-Ni) without doping B was synthesized by replacing DMAB with hydrazine hydrate. Then, the Pt atoms were deposited on the surface of the as-generated c-Ni-B and c-Ni through GRR, thereby resulting in the Pt-​c-Ni-B (fig. S21) and Pt-​c-Ni (fig. S22), respectively. Compared to a-Pt9.1-Ni-B, the distribution of Pt atoms on both c-Ni-B and c-Ni are more random, and part of the aggre-gated Pt NPs were observed (figs. S21, B to D, and S22D). The XRD patterns in figs. S21E and S22C confirm the crystalline structure of Pt-​c-Ni-B and Pt-​c-Ni with separated fcc Ni (PDF#04-0850) and fcc Pt (PDF#04-0802) phases. A nonporous Ni-B nanosphere matrix was synthesized when preparing a-Ni-B while omitting PS-​b-PEO (fig. S23, A and B). Although an amorphous structure is still formed, similar to a-Pt9.1-Ni-B MNs, after loading Pt on nonporous Ni-B nanospheres (np-Pt-Ni-B) (fig. S23C), the Pt atoms are randomly distributed around on the surface of the nonporous Ni-B, probably because of the lower surface area and considerably smoother surface of the nonporous Ni-B nanospheres compared to that of a-Ni-B MNs (fig. S23D). These results suggest that atomically distributed Pt atoms on a-Pt9.1-Ni-B MNs originated from the disordered arrange-ment of Ni atoms of a-Ni-B. Compared to the crystalline phase, the amorphous state has a large number of unsaturated coordination sites and hanging bonds, as well as a uniform distributed structure defects (37), which facilitates the dispersion of guest Pt atoms. In addition, the bumpy surface of Ni-B MNs with abundant atomic steps and kinks ensures a more homogeneous dispersion of Pt around the mesopores (Fig. 4A), thereby verifying the contribution of the porous structure to the distribution of the guest metal atoms in the GRR.Library synthesisOur strategy can be extended to synthesize various other noble–non-noble metal borides, such as a-M-Ni-B (M = Pt/Rh/Ir/Ru/Pd/Au) MNs (Fig. 4, B to E, and figs. S24 to S30), with precise composi-tional and experimental control. Despite the diverse reduction po-tentials of these noble metals, the standard reduction potentials of Mx+/M pairs of noble metals are more positive than that of the Ni2+/Ni pair (table S2), which is the driving force for the GRR. However, the potentials listed here represent the value in the standard state only. The actual value of the reduction potential may be also affected by pH, concentrations of relevant ions, coordination environment, and other nonstandard conditions (38). For example, the higher standard reduction potential (E0, 1.0 V versus SHE) for the [AuCl4]−/Au pair is more positive than that of the [PtCl6]2−/[PtCl4]2− (0.73 V versus SHE) and [PtCl4]2−/Pt (0.76 V versus SHE) pairs, resulting in crystalline-amorphous Au-Ni-B nanocomposites with crystalline Au on the shell and amorphous Ni-B core (fig. S24). This result is consistent with our previous work on the Au-CoFeB composite (39). However, although the E0 of Ir3+/Ir (1.16 V versus SHE) is higher than that of the [AuCl4]−/Au pair, amorphous Ir-Ni-B (i.e., a-Ir6.0-Ni-B) MNs are produced without observing the crystalline state of Ir (Fig. 4C). Furthermore, when an aqueous Na2PdCl4 pre-cursor was used for GRR, the crystalline Pd presented on the surface of the Ni-B MNs despite a lower reduction potential (0.59 V versus SHE for [PdCl4]2−/Pd) than that of the Pt-based pair (fig. S25). Fur-ther, amorphous Pd-Ni-B (i.e., a-Pd7.8-Ni-B) MNs were synthesized when Pd(acac)2 dissolved in acetone was used instead of that dis-solved in aqueous Na2PdCl4 (Fig.  4E). These results demonstrate that the driving force of GRR does not simply depend on the E0 gap between guest metal and parent substrate; the crystal phase and morphology of the host, precursor of the guest metal, and reaction conditions also need to be considered. Thus, besides a-Pt-Ni-B MNs, a library synthesis of four other core-shell amorphous meso-porous Mx-Ni-B (i.e., a-Mx-Ni-B, M = Rh, Ir, Ru, Pd; x represents the corresponding mass loadings; table S3) alloys were performed using a similar strategy. Figure 4 (B to E) shows the SEM, HAADF-STEM, and corresponding EDS elemental mappings, as well as the high-resolution TEM (HRTEM) images of a-Rh2.0-Ni-B, a-Ir6.0-Ni-B, a-Ru3.1-Ni-B, and a-Pd7.8-Ni-B MNs, respectively. The Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e6 of 12well-defined mesoporous shapes of a-Rh2.0-Ni-B, a-Ir6.0-Ni-B, a-Ru3.1-Ni-B, and a-Pd7.8-Ni-B MNs are confirmed by the SEM and HAADF-STEM results. The corresponding EDS maps demonstrate the core-shell structure with a Ni-B–dominant core and the noble metal–rich shell in all a-Mx-Ni-B MNs. The HRTEM images and the XRD patterns suggest the amorphous natures of all a-Mx-Ni-B MNs (Fig. 4, B to E, and fig. S26).The composition of noble metals on a-Ni-B MNs is not limited to a single noble metal element but can include two noble metals elements, as exemplified by a-Pt-Ru-Ni-B, a-Pt-Ir-Ni-B, and a-Pt-Rh-Ni-B (figs.  S27 to S29). Further, the mesoporous host can be substituted with other amorphous 3d transition-metal borides, such as a-Co-B MNs, thereby forming a-Pt-Co-B (fig. S30). With appro-priate control over the preparation conditions, the composition of both the host mesoporous amorphous non-noble metal borides and guest noble metals can be considerably adjusted. Therefore, this ap-proach offers notable command over both the composition and mesoscopic structure/morphology of disordered noble–non-noble metal borides, making it possible to discover interesting new disor-dered alloys and tailor the performance in their midst.Electrocatalytic HER performanceElectrocatalytic water splitting, driven by renewable electricity, of-fers a green path for producing highly pure hydrogen through the cathode HER, thereby providing a promising approach to realize a carbon-neutral energy society (40, 41). Mainstream electrolysis sys-tems for water splitting include alkaline water electrolysis and acidic water electrolysis (42). Compared to the alkaline environment, acid-ic water electrolysis has advantages of high partial load range and high current density; however, the corrosive acidic environments require the utilization of costly Pt-based noble catalysts, which in-creases the overall stack cost (43). The preparation of atomic-dispersed Pt catalysts not only improves the atomic utilization of Pt to reduce cost but also exhibits excellent HER reactivity (44, 45).In this study, we evaluated the HER performance of the prepared a-Pt-Ni-B catalysts in an Ar-saturated 0.5 M H2SO4 solution using a three-electrode system. The commercial carbon-supported Pt (Pt/C) containing 20 wt % Pt was used as the reference catalysts for com-parison. The linear sweep voltammetry (LSV) measurements un-veils that Pt atoms are the active center for the acidic HER, and the optimal catalyst is a-Pt9.1-Ni-B MNs (Fig. 5A and fig. S31). When normalized on a geometrical surface area, a-Pt9.1-Ni-B requires an overpotential of 11 mV to reach a current density of 10 mA cm−2, which is greater than the 18 mV of the Pt/C catalyst (Fig. 5B). The mass activity, normalized by the loading amount of Pt, of a-Pt9.1-Ni-B at an overpotential of 50 mV is 30.0 A mgPt−1 (20.7 times higher than that of the commercial Pt/C), which indicates the high intrinsic catalytic of a-Pt9.1-Ni-B. Such a low overpotential and high mass ac-tivity over a-Pt9.1-Ni-B still have high levels relative to other report-ed atomically dispersed Pt-based electrocatalysts in acidic media (Fig. 5C and table S5) (44, 46–61). Cyclic voltammetry (CV) studies reveal that our a-Pt9.1-Ni-B catalyst exhibits a different hydrogen adsorption/desorption curve in comparison to Pt/C (fig.  S32) because of the low loading amount of Pt on the electrode. The Fig. 4. Library synthesis of a-M-Ni-B MNs. (A) Schematic of the effect of host structure on the dispersion of guest metals. SEM (scale bars, 100 nm), HAADF-STEM and EDS maps (scale bars, 40 nm), merge of Ni (red) and M (blue) (scale bar, 20 nm), HRTEM images (scale bars, 1 nm) of a-M-Ni-B MNs: (B) a-Rh2.0-Ni-B, (C) a-Ir6.0-Ni-B, (D) a-Ru3.1-Ni-B, and (E) a-Pd7.8-Ni-B.Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e7 of 12enhanced HER kinetics of a-Pt9.1-Ni-B are verified through Nyquist plots and Tafel analyses. The Nyquist plots show that a-Pt9.1-Ni-B had a smaller semicircular diameter in the electrochemical impedance spectra compared to that of Pt/C, suggesting the lower charge trans-fer resistance and faster HER kinetics of a-Pt9.1-Ni-B (fig. S33A) (62). In addition, a-Pt9.1-Ni-B has a lower Tafel slope of 27 mV dec−1 compared to that of the Pt/C catalyst (29 mV dec−1) (fig. S33B). The values of the Tafel slopes suggest that the well-accepted Volmer-Tafel mechanism with the Tafel-step (i.e., hydrogen desorption) lim-ited pathway in the acidic environment occurred on both a-Pt9.1-Ni-B and Pt/C (63). The experimental measurements of H2 production closely align with the theoretical values, indicating a high Faradaic efficiency (FE) of ~100% (fig.  S34) for a-Pt9.1-Ni-B. Further, the higher HER activity of a-Pt9.1-Ni-B than that of np-Pt-Ni-B, c-Pt9.1-Ni-B, and Pt-​c-Ni catalysts verifies its structural and compositional advantages (fig. S35).The comparatively high catalytic stability of a-Pt9.1-Ni-B MNs for the acidic HER was confirmed by chronopotentiometry (CP) and CV cycles tests. Our CP studies reveal that the overpotential increased by only 9 mV over the a-Pt9.1-Ni-B catalyst in a 20-hour continuous test. For comparison, Pt/C catalyst exhibits a 78-mV overpotential increase under the same conditions (Fig. 5D). In addi-tion, after 200 and 2000 CV cycling tests, the a-Pt9.1-Ni-B catalyst exhibits superior preservation of its original catalytic activity com-pared to that with the Pt/C catalyst (fig. S36). We characterized the a-Pt9.1-Ni-B catalyst using TEM, SAED, HAADF-STEM, and XPS measurements to investigate the structural transformation of the catalyst after CV stability testing (figs. S37 to S39). The results show that, after the 200 CV cycling test, the a-Pt9.1-Ni-B catalyst largely maintained its initial mesoporous morphology and amorphous structure (fig.  S37). The EDS maps and corresponding line-scan profiles confirm the out-shell predominantly distributed Pt, while some microcrystallines are observed on the edge site, probably be-cause of the aggregation of Pt atoms after leaching metallic Ni under acidic condition measurements (fig. S38). However, the surface of the catalyst remains composed of metallic Pt, as confirmed by the XPS results (fig. S39). In contrast, the decline in the HER activity of Pt/C catalyst can be attributed to the evident particle agglomeration and growth after the CV stability test, possibly resulting from ag-gregation and/or Ostwald ripening (fig. S40).We conducted density functional theory (DFT) calculations to analyze the d-band center shift and hydrogen adsorption free energy (∆GH*) on the Pt and Pt-Ni-B model (fig.  S41, A and B) to gain deeper insights into the origin of the high HER catalytic perfor-mance of a-Pt9.1-Ni-B MNs. The density of states of the topmost Pt atomic layer of the Pt and Pt-Ni-B surfaces are shown in fig. S41C. The Fig. 5. Electrochemical HER performance in acidic media. (A) Polarization curves after the manual iR correction of a-Pt-Ni-B MNs with various Pt contents and com-mercial Pt/C in the 0.5 M H2SO4 electrolyte at a scan rate of 2 mV s−1. (B) Comparison of overpotentials (@10 mA cm−2) and mass activities normalized by Pt loading on the electrode at −50 mV for a-Pt9.1-Ni-B MNs and Pt/C. Error bars correspond to the SD based on three independent measurements. (C) Comparison of overpotentials at 10 mA cm−2 and mass activities at overpotential of 50 mV with a-Pt9.1-Ni-B that reported Pt-based HER catalysts in acidic electrolytes. (D) CP recorded on a-Pt9.1-Ni-B MNs and Pt/C at a constant current density of 10 mA cm−2.Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e8 of 12d-band center of Pt-Ni-B is −2.23 eV, which was downshifted in comparison to that of bare Pt (−2.16 eV). According to the d-band center theory (64), the lower d-band center indicates a weaker inter-action binding between the active sites and surface hydrogen adsorp-tion. Therefore, the modulating ∆GH* is determined to be −0.2 eV on Pt-Ni-B, which is closer to zero compared to the value of −0.26 eV observed on Pt (fig. S41D). In HER, a lower value of |∆GH*| (i.e., closer to zero) indicates enhanced adsorption and desorption ca-pabilities, leading to excellent HER performance (65). Moreover, the Bader charge analysis confirms the electron-enriched Pt sites on Pt-Ni-B, indicating that the modulated electronic structure of Pt (fig. S42) is consistent with the XPS and XANES results. Therefore, the above theoretical results can potentially clarify why our synthesized material, a-Pt9.1-Ni-B, demonstrates su-perior electrochemical HER activity compared to that of pure Pt in the experiments.DISCUSSIONWe showed that fabricating amorphous and mesoporous structures of Ni-B can help achieve atomically dispersed Pt atoms, thereby forming SAAs. Both local long-range disordered atomic arrange-ment Ni atoms and rich structural steps/kinks of Ni-B ensure the uniform dispersion and easily controllable loading content of Pt (ranging from 1.7 to 12.2 wt %). The prepared a-Pt-Ni-B samples have isolated Pt atoms, amorphous phases, bimetallic surface states, and well-defined mesoporous shape. Owing to the high density of Pt active sites and metal-metal interactions in SAAs, typical a-Pt9.1-Ni-B MNs exhibit a modified electronic structure and reduced hydrogen binding energy, as confirmed from experimental and DFT calculation results. In acidic HER, the a-Pt9.1-Ni-B catalyst demonstrates a low overpotential (11 mV@10 mA cm−2) and high mass activity (30.0 A mgPt−1@−50 mV), thereby surpassing com-mercial Pt/C and most single Pt-based catalysts. This synthetic method can be extended to other mesoporous amorphous noble–non-noble metal borides with atomically dispersed noble metals, such as a-M-Ni-B (M  =  Pt, Pd, Rh, Ru, and Ir), a-Pt-M-Ni-B (M  =  Rh, Ru, and Ir), and a-Pt-Co-B. Our results indicate that when preparing SAAs using GRR, besides considering the redox potential gap between the host and guest metals, the loading amount and dispersion of the guest metal depend greatly on the phase and morphology of the host.MATERIALS AND METHODSChemicalsAll reagents were used as purchased without further purification. Nickel(II) acetate tetrahydrate [Ni(OAc)2·4H2O], cobalt(II) acetate tetrahydrate [Co(OAc)2·4H2O], iron (III) chloride hexahydrate (FeCl3·6H2O), chloroplatinic acid hexahydrate (H2PtCl6·6H2O), palladium(II) acetylacetonate [Pd(acac)2], sodium hexachlororhodate(III) (Na3RhCl6), ruthenium(III) chloride (RuCl3), iridium(III) chloride hydrate (IrCl3·xH2O), tetrabutylphosphonium bromide (Bu4PBr), sodium borohydride (SBH), DMAB, Pluronic F-127 (F127), hydrazine hydrate (N2H4, 50 to 60%), and Nafion perfluori-nated resin solution (5 wt % in a mixture of lower aliphatic alcohols and water; containing 45% water) were purchased from Sigma-Aldrich. The block copolymer poly(styrene)5000-​b-poly(ethylene ox-ide)2000 (PS-​b-PEO) (average molecular weight of the corresponding blocks is shown in subscript numbers) was obtained from Polymer Source. N,N-dimethylformamide (DMF) and acetone were pur-chased from Nacalai Tesque Inc.Synthesis of a-Ni-B MNsThe a-Ni-B MNs were synthesized by a wet chemical reduction with an assembly of the diblock copolymer micelles method (25). In brief, 10-mg PS-​b-PEO was completely dissolved in 3.0 ml of DMF in a double-neck flask to form a uniform solution under sonication. Then, 3.0 ml of H2O, 4.0 ml of aqueous 60 mM Ni(OAc)2·4H2O (containing 0.3-g Bu4PBr), and 6.0 ml of aqueous 0.5 M DMAB solution were added in the above solution. The mixed solution in the flask was purged with Ar and kept in an oil bath at 40°C under gentle stirring for 5.0 min. Next, 30 μl of freshly prepared SBH aqueous solution (1.0 mg/ml) was injected into the flask to initi-ate the reduction process. The solution gradually developed bub-bles and transitioned from light green to black. Subsequently, it reacted for 1.0 hour, resulting in the a-Ni-B/micelle suspension. Nonporous amorphous Ni-B nanospheres were synthesized simi-lar to a-Ni-B MNs by omitting the block copolymer.Synthesis of a-Pt-Ni-B MNsThe a-Pt-Ni-B MNs were synthesized by a GRR. The prepared a-Ni-B/micelle suspension was cooled to room temperature (~22°C), and then 1.0  ml of the desired concentration of H2PtCl6·6H2O (i.e., 2.0, 4.0, 8.0, and 16 mM) aqueous solution was added drop-wise with stirring and kept for 1.0 hour to prepare a-Ptx-Ni-B (x displayed the weight percent of Pt loading) MNs. Last, the prod-uct was collected by centrifugation and washed several times with acetone/ethanol. For comparison, np-Pt-Ni-B nanospheres were synthesized similar to a-Pt-Ni-B MNs without using the block copolymer.Synthesis of c-Pt-Ni-B MNsThe reference c-Pt9.1-Ni-B MNs sample was prepared by thermally treating the pristine typical amorphous a-Pt9.1-Ni-B MNs at 400°C in Ar atmosphere for 2.0 hours.Synthesis of Pt-​c-Ni-B compositesThe c-Ni-B sample was prepared by thermally treating the dried a-Ni-B MNs powder at 400°C in Ar atmosphere for 2.0  hours. The obtained black powder was redispersed on 10  ml of acetone, fol-lowed by adding 1.0 ml of 8.0 mM H2PtCl6·6H2O aqueous solution (Ar prepurged) dropwise. After 1.0 hour of reaction, the final Pt-​c-Ni-B composites were collected by centrifugation and washed sev-eral times with acetone.Synthesis of Pt-​c-Ni compositesThe Pt-​c-Ni composites were prepared as another comparison sam-ple. To this end, 100 mg  of F127 was dissolved in 3.0  ml of H2O under sonication, followed by adding 3.0 ml of H2O, 4.0 ml of 60 mM Ni(OAc)2·4H2O, and 6.0  ml of N2H4, in sequence. After purging with Ar for 5.0 min, the reaction solution was heated to 80°C in an oil bath and vigorous stirred for 4.0 hours. Then, the produced suspension was cooled down to 22°C and 1.0  ml of 8.0 mM H2PtCl6·6H2O aqueous solution (Ar prepurged) was added in dropwise manner. After 1.0 hour of reaction, the final product was collected by centrifugation and washed several times with acetone/ethanol.Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e9 of 12Synthesis of a-M-Ni-B (M = Rh, Ir, Ru, and Pd) MNsThe a-M-Ni-B (M = Rh, Ir, Ru, and Pd) MNs were prepared similar to a-Pt-Ni-B MNs; however, the noble metal precursors (tempera-ture for GRR) were changed to 1.0 ml of aqueous 8.0 mM Na3RhCl6 (22°C), aqueous 8.0 mM IrCl3·xH2O (22°C), aqueous 8.0 mM RuCl3 (22°C), and 8.0 mM Pd(acac)2 (0°C) (dissolving in acetone) to syn-thesize a-Rh2.0-Ni-B, a-Ir6.0-Ni-B, a-Ru3.1-Ni-B, and a-Pd7.8-Ni-B, respectively.Synthesis of a-Pt-M-Ni-B (M = Ru, Rh, and Ir) MNsThe a-Pt-M-Ni-B MNs were prepared similar to a-Pt-Ni-B MNs; however, the metal precursor of H2PtCl6·6H2O (8.0 mM, 1.0  ml) was changed to H2PtCl6·6H2O (8.0 mM, 0.8 ml), followed by adding RuCl3 (8.0 mM, 0.2 ml), Na3RhCl6 (8.0 mM, 0.2 ml), and IrCl3·xH2O (8.0 mM, 0.2 ml) for a-Pt-Ru-Ni-B, a-Pt-Rh-Ni-B, and a-Pt-Ir-Ni-B MNs, respectively.Synthesis of a-Pt-Co-B MNsThe a-Pt-Co-B MNs were prepared in the same manner as a-Pt-Ni-B; however, the a-Co-B/micelle suspension was used instead of the a-Ni-B/micelle suspension. The a-Co-B/micelle suspension was prepared similar to the a-Ni-B/micelle suspension, wherein Co(OAc)2·4H2O was used instead of Ni(OAc)2·4H2O according to the published method (25).CharacterizationsAn ICP-OES was performed on Agilent 5800 to determine the com-positions of M-Ni-B (M = Pt, Rh, Ir, Ru, and Pd) MNs. The mor-phologies of the samples were characterized using a field-emission SEM (Hitachi SU-8000, 10 kV) and a TEM (JEOL JEM-2100F; ac-celerating voltage, 200 kV). SEM-EDS was observed using a flat quad EDS (5060F, Bruker). The beam spot size for TEM-EDX map-ping was set to 0.5 nm. HAADF-STEM images were obtained using a Thermo Fisher Scientific Titan microscope with an acceleration voltage of 300 kV and a probe current of ~30 pA. The EELS spectra and spectrum images were collected using a Gatan Continuum spectrometer. The XPS spectra were obtained on a PHI Quantera SXM (ULVAC-PHI) under an excitation source of focused mono-chromatic Al Kα x-ray. The energies for all high-resolution spectra were calibrated using the C 1s main peak as 284.8 eV. The samples for XPS were fabricated in a glove box after drying in vacuum to avoid surface oxidation. Powder XRD measurements were collected with a Smart lab x-ray diffractometer (RIGAKU) with a step size of 2° min−1 using a Cu Kα radiation (40 kV, 30 mA) source. The dif-fraction peaks corresponding to the pore-to-pore distance were ob-tained by the XRD pattern in a small angle region using SAXS measurements (Rigaku NANO-Viewer). N2 adsorption-desorption isotherms were acquired from a BELSORP-mini (BEL, Japan) at 77 K and the pore size distributions were calculated based on the Barrett-Joyner-Halenda model.XAS measurements and data processingEx situ XAS measurements used to characterize the a-Pt9.1-Ni-B MNs at the Ni K-edge and Pt LIII-edge were performed at the BL14W1 station in the Shanghai Synchrotron Radiation Facility (SSRF) (Shanghai, China). The electron storage ring of SSRF was operated at 3.5 GeV with a maximum current of 200 mA. The XAS data were collected using a fixed-exit Si (111) double-crystal mono-chromator. A Lytle detector was used to collect the fluorescence signal, and the energy was calibrated using a metal foil. The obtained XAS data were processed in the ATHENA module of the IFEFFIT software package. The EXAFS contributions were separated from dif-ferent coordination shells using a Hanning windows (dk = 1.0 Å−1). Subsequently, quantitative curve fittings were performed in the R-space (1.0 to 3.0 Å) with a Fourier transform k-space range of (3 to 12.5 Å−1) using the module ARTEMIS of IFEFFIT. During the curve fitting, the overall amplitude reduction factor S02 was fixed to the best-fit values of 0.83 and 0.80 determined from fitting the data of the metal Pt foil and Ni foil, respectively. Structural parameters such as the CN, interatomic distance (R), Debye-Waller factor (σ2), and edge-energy shift (ΔE0) were allowed to vary during the fitting pro-cess. For the WT analysis, the χ(k) exported from ATHENA was imported into the Hama Fortran code. The parameters were R range = 0 to 6 Å, k range = 0 to 14 Å−1, and k weight = 2. The Mor-let function with κ = 10 and σ = 1 was used as the mother wavelet to provide the overall distribution.Electrochemical measurementsElectrochemical measurements (using CHI 660EZ) were conducted using a three-electrode system with graphite rod and Ag/AgCl (sat-urated KCl) as the counter and reference electrodes, respectively. The reference electrode was calibrated with respect to the reversible hydrogen electrode (RHE) using a high-purity hydrogen saturated three-electrode system (i.e., ERHE = EAg/AgCl + 0.23). The working electrode was prepared by dropping the ink solution onto a rotating disk electrode (RDE; 0.0706 cm2). The ink solution was prepared by dispersing the as-prepared Pt-Ni-B–based samples (1.0 mg) and Vulcan XC-72 carbon (4.0 mg) in a solution of 980 μl of ethanol and 20 μl of Nafion under sonication, and 3.2 μl of the ink was dropped onto the RDE. The catalyst was first activated by CV scanning for 200 cycles, and then LSV was conducted at a scan rate of 2 mV s−1 in an Ar-saturated electrolyte (0.5 M H2SO4, pH =  0.2) stirred at 1600 rpm without ohmic drop correction. Tafel plots were derived from the overpotential at different range versus the log (current density) according to the corresponding LSV curves (η = a + b log | j |, where η, a, b, and j represent the overpotential, exchange current density, Tafel slope, and current density, respectively). Electrochem-ical impedance spectroscopy was measured at the different poten-tials with a frequency range from 105 to 0.01 Hz. The resistance of the solution (Rs) resulting from the Nyquist plot was used to correct the ohmic drop (iR-correction) for the HER measurement. The cor-rected potential could be obtained by Ecorrected = ERHE  − iRs. Electro-catalytic durability tests were conducted by subjecting the ink catalyst loaded on carbon fiber paper (with a loading of 1.0 mg cm−2) to CP measurements. Before CP measurements, the catalyst underwent 200 CV cycles at a scan rate of 50 mV s−1 to establish stable catalytic performance. Gas products generated at the cathode side were collected in a gas-tight H-type cell to determine the FE of the catalyst. The gas product (i.e., H2) was confirmed using gas chro-matography (Shimadzu Tracera, BID-2010), and the experimental amount of H2 was measured using a drainage method. Then, the FE was calculated by comparing the experimental H2 production with the theoretically calculated value.Computational detailsThe calculations were performed in the framework of the DFT with the projector augmented plane-wave method, as implemented in the Vienna ab  initio simulation package (66). The generalized Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e10 of 12gradient approximation proposed by Perdew-Burke-Ernzerhof was selected for the exchange correlation potential (67). The long-range van der Waals interaction was described by the DFT-D3 approach (68). The Pt-Ni-B system was constructed based on a 2 × 2 × 2 of the primitive cell of the fcc phase where five Pt atoms were substituted by three Ni atoms, and two additional B atoms were added to the octahedral sites. The number of atoms of the model was set based on the compositions of stabilized Pt-Ni-B catalyst after HER, deter-mined by ICP-OES. The Pt-Ni-B structure was fully relaxed until the residual forces on the atoms declined to less than 0.02 eV/Å. The cut-off energy for the plane wave was set to 450 eV. The energy cri-terion was set to 10−5 eV in the iterative solution of the Kohn-Sham equation. Brillouin zone integration was performed using a 10 × 10 × 10 k-mesh. The Pt-Ni-B (111) surface was constructed based on 2 × 2 of its primitive cells, containing eight atomic layers (10 Pt, six Ni, and four B atoms). The Pt (111) surface was construct-ed based on 2 × 2 of its primitive cells, containing four atomic layers (16 Pt atoms). A vacuum layer of 15 Å was added perpendicular to the slabs to avoid artificial interaction between periodic images. For the Pt-Ni-B surface, the upper three layers were allowed to relax, whereas the bottom five layers were fixed. For Pt (111) surface, the upper two layers were allowed to relax, whereas the bottom two lay-ers were fixed. For calculating the surfaces, the symmetry was re-moved, and Brillouin zone integration was performed using a 6 × 6 × 1 k-mesh, thereby leading to 36 k-points for k-points sam-pling. Further, the dipole correlation was considered for the surface calculations.The ∆GH* (change in the Gibbs free energy before and after hy-drogen adsorption) was used as an indicator for evaluating the per-formance of HER. ∆GH* = Gsurface + H − Gsurface + GH2/2, Gsurface + H was the Gibbs free energy of the surface with hydrogen adsorption, Gsurface was the Gibbs free energy of the surface, and GH2 was the Gibbs free energy of the hydrogen gas. The value of |∆GH*| closer to zero suggested better HER performance. The hydrogen adsorption Gibbs free energy under standard atmospheric pressure was defined as Gsurface + H = Esurface + H + ∆G(T), where Esurface + H was the total energy of the surface with hydrogen adsorption calculated by the DFT; ∆G(T) = ∆EZPE + ∆H(T) − ∆S(T), where ∆EZPE, ∆H(T), and ∆S(T) denote the contributions of the zero-point energy, enthalpy, and entropy to the Gibbs free energy attributed to hydrogen adsorp-tion, respectively. Gsurface was approximated as Esurface because the correction of the ∆G(T) to the Gibbs free energy of the surface could be cancelled out before and after hydrogen adsorption. The calculation of ∆G(T) was performed by the VASPKIT Package (69), and T was used as 298.2 K.Supplementary MaterialsThis PDF file includes:Figs. S1 to S42Tables S1 to S5REFERENCES AND NOTES  1.  Y. Chen, J. Lin, B. Jia, X. Wang, S. Jiang, T. Ma, Isolating single and few atoms for enhanced catalysis. Adv. Mater. 34, e2201796 (2022).  2.  Y. Yao, S. Hu, W. Chen, Z.-Q. Huang, W. 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This work was also supported by the National Key Research and Development Program of China (SQ2020YFA020032), the National Natural Science Foundation of China (22236005, 22108238, and 22376142), and Shanghai Government Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024Kang et al., Sci. Adv. 10, eado2442 (2024)     21 June 2024S c i e n c e  A d v a n c e s  |  R e s e ar  c h  A r t i c l e12 of 12(22dz1205400, 23520711100, and 22010503400). Author contributions: Y.Y. and H.L. conceived and supervised the research. Y.K. and Y.Y. designed the experiments and wrote the paper. Y.K. performed most of the experiments and data analysis. O.C. and K.K. performed the TEM and EELS characterization and data analysis. S.L., X.W., and D.Z. processed and analyzed the XAF data. B.J. supported the theoretical calculations. Y.Z., L.Z., L.F., D.J., C.W., and T.A. participated in the experiments and discussions. All authors discussed the results and commented on the manuscript. Competing interests: The authors declare that they have no competing interests. Data and materials availability: All data needed to evaluate the conclusions in the paper are present in the paper and/or the Supplementary Materials.Submitted 24 January 2024 Accepted 14 May 2024 Published 21 June 2024 10.1126/sciadv.ado2442Downloaded from https://www.science.org at National Institute for Materials Science on August 29, 2024 Mesoporous amorphous non-noble metals as versatile substrates for high loading and uniform dispersion of Pt-group single atoms INTRODUCTION RESULTS Morphological and structural characterizations GRR process study Library synthesis Electrocatalytic HER performance DISCUSSION MATERIALS AND METHODS Chemicals Synthesis of a-Ni-B MNs Synthesis of a-Pt-Ni-B MNs Synthesis of c-Pt-Ni-B MNs Synthesis of Pt-​c-Ni-B composites Synthesis of Pt-​c-Ni composites Synthesis of a-M-Ni-B (M = Rh, Ir, Ru, and Pd) MNs Synthesis of a-Pt-M-Ni-B (M = Ru, Rh, and Ir) MNs Synthesis of a-Pt-Co-B MNs Characterizations XAS measurements and data processing Electrochemical measurements Computational details Supplementary Materials This PDF file includes: REFERENCES AND NOTES Acknowledgment