Catalyst design and structural characterizationPtMg/C samples were prepared through three systematic steps namely; (I) solvolysis, (II) alloying and atomic ordering, and (III) creation of Pt-rich skin. The three-step synthesis route is illustrated schematically in Fig. 1a. Briefly, platinum (IV) chloride (PtCl4), magnesium chloride (MgCl2), sodium borohydride (NaBH4), and N, N-dimethylformamide (DMF) were employed as the Pt and Mg precursors, the reducing agent, and the non-aprotic solvent, respectively. DMF is a non-aqueous aprotic solvent featuring a lower boiling point (153 °C) than most aprotic solvents available which can facilitate a faster collision between the Pt and Mg nuclei at a lower temperature. In addition, the polar nature of this solvent helps to solvate the metal chlorides. In the solvolysis step (I), all precursors aforementioned were weighed out as specified in the methods section in an argon-filled glove box, sonicated, and transferred into a Teflon-lined autoclave. Subsequently, the mixture was heated at 165 °C, which is above the boiling point of the solvent to increase the kinetic energy of the two metal nuclei for efficient collision. In the alloying and atomic ordering step (II), the obtained precipitate from the hydrothermal reaction of the initial precursors, solvent, and reducing agent at 165 °C was subjected to further reductive annealing in the presence of the fast-reducing agent; powdered potassium triethylborohydride (KEt3BH) at 760 °C in 5% H2/Ar atmosphere to obtain a structurally ordered intermetallic PtMg alloy (Fig. 1c, e). Before the Ar/H2 treatment, the furnace was purged with ammonia (NH3) until 180 °C since it provides a better protective atmosphere in the initial heating stage in comparison to the Ar/H2 atmosphere49. KEt3BH is very essential in the second reductive annealing step for the alloy formation, since the absence of this reagent results in a spectrum highly dominated by crystalline Pt and magnesium oxide (MgO) phases (Supplementary Fig. 1). Finally, in step III, the structurally ordered alloy was acid-etched in 0.5 M acetic acid to remove surface oxides and unreacted metal precursors (full details are in the methods section) and to create a Pt-rich skin (Fig. 1f). We then characterized the as-prepared samples using transmission electron microscopy (TEM), X-ray photoelectron spectroscopy (XPS), X-ray absorption spectroscopy (XAS) and powder X-Ray diffraction (PXRD). Further insights on the evolution of the alloy, composition, and temperature effects are discussed in the mechanistic study section.Fig. 1: Synthesis and structural characterization.a Schematic representation of the synthesis of PtMg/C NPs b TEM image of the as-synthesized PtMg/C-31 NPs showing uniform distribution with little agglomeration. c HAADF-STEM image and d high-resolution image of the red-marked particle showing the orderly arrangement of the Pt (bright spots) and Mg (dark spots) atoms. e PXRD patterns of PtMg/C-31 before (blue) and after (red) acid treatment. f STEM-EDS line scan of a single PtMg/C-31 NP.PtMg alloys with different precursor ratios of Pt to Mg by mass were synthesized to test the effect of composition on the structure and catalytic property. Based upon our experience with the synthesis of AEM alloys, we started with a Pt to Mg precursor mass ratio of 3:1 and 1:1, herein denoted as PtMg/C-31 and PtMg/C-11, respectively. Inductively coupled plasma-optical emission spectrometry (ICP-OES) analysis shows that the PtMg/C-11 and PtMg/C-31 have Pt to Mg atomic ratios of 65.1:34.9 and 71.0:29.0, respectively.The TEM image of PtMg/C-31 presented in Fig. 1b shows that the particles are uniformly distributed on the carbon substrate with a particle size ranging from 5 to 12 nm. The wide particle size range is attributed to the formation of two main intermetallic phases. In Fig. 1c, the high-angle annular dark-field scanning TEM (HAADF-STEM) image shows that some particles exhibit the L10 intermetallic structure (Fig. 1d) as evidenced by the alternating rows of Pt and Mg atoms, where the Pt atoms are in bright contrast to the dark Mg atoms owing to the difference in atomic numbers, while others show the L12 intermetallic structure (Supplementary Fig. 2). A critical study of the L12-ordered HAADF-STEM image in Supplementary Fig. 2b and the corresponding Fourier transformation (FFT) pattern (inset) shows that the particle is highly crystalline as revealed by the individually isolated reciprocal lattice diffraction spots. Furthermore, the presence of the (100) superlattice spot is consistent with the ordered nature of the PtMg/C-31 alloy. The d-spacings of 0.368 and 0.157 nm correspond to the (100) and (211) planes of Pt3Mg, respectively. In comparison, the literature-recorded values for the dPt-Pt in disordered (A1) Pt are 0.381 and 0.158 nm for the (100) and (211) planes, respectively. This shows that the introduction of Mg atoms introduces compressive strains along the <100> and <211> directions. Such compressive strains enhance the stability of the alloy in electrochemical conditions and reduce its adsorption strength with reaction intermediates50,51. The ordered structure is verified by the PXRD patterns presented in Fig. 1e as evidenced by the appearance of the (001) and (110) superlattice peaks. The XRD obtained conforms with the standard diffraction peak of Pt3Mg (PDF#03-065-7462). STEM-EDS line scan (Fig. 1f) of a single particle confirms that the Pt and Mg are homogeneously distributed in the particle with a ∽1 nm Pt-rich shell. All the conditions employed for the synthesis of PtMg/C-31 were maintained for PtMg/C-11 except for the varying Pt to Mg precursor mass ratio. A detailed study of the TEM, HAADF-STEM, and the corresponding EDS elemental mapping images of the PtMg/C-11 in Supplementary Figs. 3 and 4 show a similar particle size distribution to that of the PtMg/C-31. Interestingly, HAADF-STEM images and EDS elemental mapping images illustrated in Supplementary Fig. 5, show that the PtMg/C-11 also exhibits a Pt-rich shell as well as the L12 and L10 intermetallic structures. Both of the aforementioned catalysts show interesting electrochemical performances and hence are discussed in detail in the electrochemical section.The elemental distribution and oxidation states of PtMg/C-31 were studied through energy-dispersive X-ray spectroscopy (EDS) and XPS analyses. In Supplementary Figs. 6 and 7, the STEM images and the corresponding EDS elemental mapping images show that Pt and Mg are homogeneously distributed throughout the particles. This confirms the alloy formation since the Mg or any of its derivative compounds would have been washed out during the 12-hour acid-etching process. Furthermore, the STEM and the corresponding EDS elemental mapping images in Fig. 2a show that the area corresponding to the Pt signal is slightly larger than that of Mg, indicating the formation of a Pt-rich shell. The STEM image in Supplementary Fig. 2a (yellow arrows) supports the result with about 1 ∽ 3 layers of Pt observed in the outmost surface. Due to the difference in atomic radii of Pt and Mg (139 and 160 nm, respectively) and the large difference in electronegativity, the electronic structures of both Pt and Mg are expected to be modified owing to the ligand and strain effects. We used XPS analysis to explore the impact of Mg on the surface chemical state of the as-synthesized PtMg/C-31 alloy. High-resolution Pt 4 f doublets were observed for both the PtMg/C-31 and commercial Pt/C. Binding energies of 71.40 and 71.74 eV are determined for the Pt4f7/2 of PtMg/C-31 and Pt/C, respectively (Fig. 2b). The negative shift observed in the PtMg/C-31 doublet relative to commercial Pt/C is associated with the strain and the ligand effects existing between Pt and Mg4,52,53,54. The large shift may also be attributed to a significant amount of Pt oxides existing at the surface of the Pt NPs. The alloy formation between Pt and Mg prevents such a significant degree of oxidation. A similar shift is observed in the PtMg/C-11 (Supplementary Fig. 8). Furthermore, to show the existence of metallic Mg (Mg0), we deconvoluted the high-resolution Mg 1 s spectrum (Supplementary Fig. 9). Integration of the areas under the two main doublets at binding energies of 1303.7 and 1304.9 eV for Mg0 and Mg2+, respectively shows that metallic Mg dominates over the +2-oxidation state.Fig. 2: Elemental composition and oxidation state.a HAADF-STEM image and the corresponding EDS elemental mapping of a single structurally ordered PtMg/C-31 NP, showing Pt signal (red), Mg signal (green), and the overlapped signal. b High-resolution Pt 4 f XPS profiles of PtMg/C-31 and commercial Pt/C. c Pt L3 XANES and d Pt L3 K2 – weighted EXAFS spectra of PtMg/C-31, Pt foil, and PtO2.X-ray absorption spectroscopy (XAS) was conducted to understand the coordination environment and the valence state in Pt foil, PtO2, and PtMg/C-31. X-ray absorption near edge structure (XANES) spectra and the insert (Fig. 2c) show that the absorption edge of PtMg/C-31 is located between those of the Pt-foil and PtO2 but is much closer to the Pt foil. This suggests that the oxidation state of Pt in the PtMg/C-31 is close to the metallic state. The white line peak intensity, which is associated with the transition of an electron from the core 2p 3/2 orbital to the unoccupied 5d orbitals of the Pt L3 edge is large for PtO2 but shows almost similar low intensity for Pt foil and PtMg/C-31, suggesting a low degree of catalyst oxidation. Fourier transform of the higher energy regions gives rise to extended X-ray absorption fine structure (EXAFS) spectra which provide information about the coordination environment of the Pt atoms. The fitting results provided in Supplementary Fig. 10 and Supplementary Table 1 show that the Pt-Pt bond distance in PtMg/C-31 (2.75 Å) is shorter than in the Pt foil (2.76 Å). This indicates that in the first coordination shell, the Pt-Pt bonds are compressively strained as a result of alloying with Mg. The EXAFS results shown in Fig. 2d also show the presence of a Pt-Mg bond with a bond length of (2.763 Å). It must be emphasized that the bond lengths of Pt-Pt and Pt-Mg in the first coordination shell are similar.Electrochemical half-cell evaluationThe electrochemical sensitivity and reactivity of our as-synthesized catalysts towards oxygen reduction reaction (ORR) were determined by cycling the PtMg/C electrocatalysts deposited on glassy carbon with an active area of 0.071 cm2 in an N2 or O2-saturated 0.1 M HClO4 electrolyte. The rotating disk electrode (RDE) cell configuration consisted of glassy carbon coated with either the as-synthesized catalyst or commercial Pt/C as the working electrode, with a Pt wire as the counter electrode and Ag/AgCl as the reference electrode (details are provided in the supplementary information)55. The metal weight percent as determined on the carbon support was 16.0 and 19.4 wt%, respectively for PtMg/C-31 and Pt/C (Supplementary Fig. 11) according to thermogravimetric analysis (TGA) data.First, we tested the electrocatalytic performance of PtMg/C-31, and the performance showed a half-wave potential (E1/2) of 0.87 V vs RHE, which is superior to commercial Pt/C with an E1/2 of 0.84 V vs RHE as evidenced by the positively shifted onset potential and LSV curve (Fig. 3a). Furthermore, the CVs presented in Fig. 3b show a positive shift to higher potential of the Pt-oxide formation peak (red arrow) for the PtMg/C-31 as compared to Pt/C. This signifies that the PtMg/C-31 is less likely to suffer from catalyst oxidation, which would retard its dissolution and subsequently decrease Ostwald ripening. In addition, the oxygen reduction peaks in the potential region of 0.7 to 0.9 V vs RHE (black arrow) show a similar positive shift, suggesting that the interaction between Pt and the oxygen intermediates in the PtMg/C-31 is weaker than that of the Pt/C. Per Sabatier’s principle, a catalyst with too strong binding may lead to difficulty in desorption, hence, resulting in the lower electrocatalytic performance observed for the commercial Pt/C. Further information obtained from the CV curves shows that the hydrogen adsorption and desorption regions that characterize the electrochemically active surface areas for Pt-based catalysts were smaller for PtMg/C-31 than Pt/C with values of 33 and 56 m2g−1, respectively (Fig. 3c). Despite the lower ECSA of PtMg/C-31, the enhanced activity is thought to originate from the strain and ligands effects of the Mg on the Pt rather than the particle size effect. Similar trends have also been observed in other alloys51.Fig. 3: Liquid half-cell tests.a Linear sweep voltammetry (LSV) curves, b cyclic voltammograms (CVs) for PtMg/C-31 and commercial Pt/C recorded in O2-saturated solution. c Electrochemically active surface areas (ECSAs) and d mass activities determined at 0.9 V for PtMg/C-31 and commercial Pt/C before and after 30,000 accelerated durability tests (ADT). LSVs recorded before and after the ADT e for PtMg/C-31 and f for Pt/C. Note: All electrochemical measurements were conducted in a 0.1 M HClO4 solution at room temperature. The durability tests were performed according to the DOE protocol described in the Supplementary Information.The accelerated durability test (ADT) was evaluated in a N2-saturated 0.1 M HClO4 within the potential range from 0.60 and 0.95 V vs RHE at a scan rate of 100 mV s−1 for 30,000 potential cycles according to the DOE protocol56. In Fig. 3d, the BOL mass activity (MA) initially recorded for PtMg/C-31 was 0.62 A mgPt−1, with EOL mass activity of 0.56 A mgPt−1 after 30,000 cycles (only 10% degradation). In contrast, commercial Pt/C shows BOL and EOL mass activities of 0.31 and 0.15 A mgPt−1, respectively. In full agreement with the above results, the improved stability of PtMg/C-31 is visualized in Fig. 3e, where the LSV curves before and after 30,000 ADT cycles almost overlap while that of the commercial Pt/C decreases significantly with a half-wave potential loss of 0.35 mV (Fig. 3f). The specific activity (SA) was evaluated by normalizing the mass activities with their corresponding ECSAs. The specific activities for PtMg/C-31 and Pt/C at the BOL were 1.8 and 0.55 mA cm−2 respectively. This indicates efficient utilization of the PtMg/C-31 ECSA (Supplementary Fig. 12a). In addition, the CVs recorded before and after the 30,000 ADT overlapped with only a small change in the hydrogen desorption area (Supplementary Fig. 12b). The resulting ECSA presented in Fig. 3c shows a 19% drop for PtMg/C-31 while a more significant loss of 42% was recorded for Pt/C. These half-cell data demonstrate the high durability of PtMg/C-31 for ORR.We also tested the electrocatalytic performance of the PtMg/C-11, and the performance was good as both the onset and half-wave potentials are shifted to more positive potentials relative to the reference Pt/C (Supplementary Fig. 13a). The ADT results conducted for the PtMg/C-11 show that there is only a small change in the onset potentials of the LSVs but the half-wave potential before and after the 30k ADT cycles is maintained (Supplementary Fig. 13c). A loss in ECSA similar to that recorded for the PtMg/C-31 can be observed in Supplementary Fig. 13d.Mechanistic studyTo explain the origin of these results, we took a step-by-step approach to explore the effect of composition and heating temperature on the formation of the PtMg alloys. First, we studied the effect of the composition on the degree of ordering. When all the synthesis conditions reported in the methods section were maintained except for varying the precursor composition of Pt to Mg, it can be realized from Supplementary Fig. 14 that the degree of ordering increases with increasing amounts of Mg. This is evidenced by the appearance and intensification of the (110) and (001) peaks as we move from PtMg/C-31, PtMg/C-11 to PtMg/C-13 (Pt: Mg precursor mass ratio of 1:3). The corresponding TEM images presented in Supplementary Fig. 15 show that all the PtMg alloy NPs are evenly distributed on the carbon substrate. The average particle size distribution is 5.81 and 5.76 nm for PtMg/C-11 and PtMg/C-31, respectively. In contrast, the TEM image of PtMg/C-13 (Supplementary Fig. 15c) reveals the presence of highly aggregated particles, as corroborated by sharp peak intensities observed in the PXRD patterns (Supplementary Fig. 14). ICP-OES analysis indicates that PtMg/C-13 exhibits a Pt:Mg atomic ratio of 63.4:36.6. These findings suggest that, at the optimal temperature of 760 °C, varying the initial Mg content does not significantly influence the composition of the nanoparticles; rather, it primarily affects the degree of ordering as seen in Supplementary Fig. 14. This phenomenon explains the marginal differences observed in the atomic ratios.Similarly, the effect of temperature on the PtMg/C-31 was also studied by maintaining all other conditions specified in the methods section (Supplementary Fig. 16). XRD patterns obtained after the first reduction process at 165 °C conform with that of crystalline Pt. This suggests that the Mg reacts with the Pt in some way, and otherwise, the Mg would have been lost during washing. The absence of additional peaks corresponding to MgCl2 or other Mg moieties such as MgO or Mg(OH)2 further supports this postulation. When the temperature was raised in the annealing process, the results presented in the same figure show that the crystalline Pt phase remains dominant at temperatures as high as 600 °C. However, when the temperature was increased to 650 °C, the structurally ordered PtMg alloy was observed as evidenced by the appearance of the (110) superlattice peak. Increasing the temperature from 650 to 800 °C resulted in the appearance of the (001) diffraction peaks and an increase in the intensity of the (110) superlattice peaks. Among all the samples prepared at the different annealing temperatures, the sample obtained at 760 °C showed the best catalytic performance. Hence, all control samples were also synthesized at the same temperature. The resulting TEM images and particle size distributions of the temperature-variant samples are presented in Supplementary Fig. 17. In short, we found the formation of the PtMg alloy to be influenced by temperature and composition. Increasing the Pt to Mg ratio beyond a certain threshold value (for example, a sample with Pt to Mg precursor mass ratio of 3:0.5) could limit the alloy formation as indicated in Supplementary Fig. 18, resulting in poor catalytic performance. The optimum Pt to Mg mass ratio was determined to be 3:1.Membrane electrode assembly (MEA) full cell evaluationThe performance obtained from half-cell RDE test with PtMg/C samples prompted us to evaluate the performance in a practical fuel cell by employing each of the two electrocatalysts (PtMg/C-31 and PtMg/C-11) as the cathode of the MEA under H2-O2/air conditions. Under the same test conditions, the commercial Pt/C (19.4 wt% Tanaka, Japan) was used as a reference catalyst for efficient comparison. All polarization curves presented in this section were determined under iR-free conditions. The polarization curves obtained at 80 °C under the H2-O2 gas flow in Supplementary Fig. 19a show that the PtMg/C-31 exhibits higher operating voltages for the same current densities. The initial power density documented for the PtMg/C-31 (1575 mW cm−2) is higher than that of the commercial Pt/C (1465 mW cm−2). Similarly, the performance of PtMg/C-31 and Pt/C was evaluated under H2-air conditions in Supplementary Fig. 20. The initial peak power density recorded for both samples followed a comparable trend as those measured under H2-O2 conditions with values of 850 and 660 mW cm−2 for PtMg/C-31, and Pt/C respectively. Detailed values of the measurements are reported in Supplementary Table 2.The stability of the electrocatalysts was determined by the potential-dependent cycling on the MEA between 0.60 and 0.95 V for 30,000 cycles according to the DOE accelerated durability test (ADT) protocol. Typical in the H2-O2 condition, the current density @ 0.8 V only dropped by 31 mA cm−2 (from 476 to 445 mA cm−2) with a 6.5% loss for the PtMg/C-31 cathode. In contrast, at the same voltage, the Pt/C shows a ∽ 49% loss in current density from 287 to 147 mA cm−2 (Supplementary Fig. 19b). The polarization curves recorded in Fig. 4a, b and the histogram of the peak power densities in Fig. 4h show that PtMg/C-31 is stable with only a 6% drop (1575 to 1480 mW cm−2) after 30,000 ADT cycles, while the Pt/C exhibits a 30% loss (1465 to 1030 mW cm−2) for the same number of cycles. Interestingly, the peak power density for the PtMg/C-31 at its EOL is higher than that of Pt/C at its BOL (Fig. 4h). In terms of overall Pt usage, the PtMg/C-31 reaches a BOL mass activity of 0.50 A mgPt−1 and EOL mass activity of 0.48 A mgPt−1 at 0.9 V, corresponding to only a 3% decrease after the 30,000 ADT cycles, surpassing the DOE 2020-2025 targets in terms of BOL and EOL mass activities (0.44 and 0.26 A mgPt−1 respectively). This performance overwhelms that of the Pt/C catalyst which shows a 49.6% loss in initial mass activity (from 0.22 A mgPt−1 (BOL) to 0.11 A mgPt−1 (EOL)), a performance well below the DOE targets (Fig. 4g).Fig. 4: Fuel cell performance and durability tests.Polarization and power density curves before and after 30,000 cycles for a PtMg/C-31, b Pt/C, and c PtMg/C-11 in H2-O2 condition. Polarization and power density curves before and after 30,000 cycles for d PtMg/C-31, e Pt/C, and f PtMg/C-11 in H2-air condition. Histograms of g mass activities, h peak power densities of PtMg/C-31, Pt/C, and PtMg/C-11 before and after 30,000 ADT cycles in H2-O2 gas flow, and i cell voltages recorded at 0.8 A cm−2 in H2-air for the three catalysts before and after the 30,000 ADT (DOE Target <30 mV drop after ADT). Note: Detailed measurement conditions are indicated in the methods section. Error bars indicate the standard deviation for three independent measurements.The stability test performed under the H2-Air conditions (Fig. 4d, e) exhibited similar degradation patterns to those measured under H2-O2 as evidenced by the 6% drop in peak power density for PtMg/C-31 (850 to 800 mW cm−2) and 29% for commercial Pt/C (660 to 470 mW cm−2) (Supplementary Fig. 20b). The voltage drop at 0.8 A cm−2 was 9 and 75 mV for PtMg/C-31 and Pt/C, respectively in the H2-air condition. (Fig. 4i). This voltage drop for PtMg/C-31 satisfies the DOE target of <30 mV drop after 30,000 ADT cycles. Interestingly, in contrast to numerous articles that report that the RDE tests are significantly different from the PEMFC tests, here, we show that similar performances reported in the half-cell tests are successfully repeated in the full-cell MEAs. In particular, the CVs recorded before and after the 30,000 stability test cycles in both the half-cell (Supplementary Fig. 12b) and full-cell (Supplementary Fig. 21) show similar degradation patterns despite the higher loss in ESCA observed in the hydrogen desorption region of the full cell MEA measurement.Due to the similar activities of PtMg/C-31 and PtMg/C-11 recorded in liquid half-cell tests, we evaluated the PEMFC performance of the PtMg/C-11 under the same conditions. The polarization curves recorded in the H2-O2 condition in Fig. 4c show that the PtMg/C-11 exhibits similar performance to the commercial Pt/C sample in the high current density region, with a peak power density of 1430 mW cm−2. However, interestingly, the kinetics of this sample in the low current density region was much improved as shown by the BOL mass activity of 0.40 A mgPt−1 (Fig. 4g). This sample also displayed stability similar to that observed for the PtMg/C-31 sample. The power density dropped by only 5% (from 1430 to 1355 mW cm−2), while almost the same percentage drop was also recorded from the BOL to the EOL mass activity (0.38 A mgPt−1). The corresponding measurement under H2-air conditions is presented in Fig. 4f. This result is in contrast to that recorded in H2-O2 conditions as evidenced by the higher power density of 700 mW cm−2 compared to that of Pt/C with 660 mW cm−2. These results emphasize the stability derived by alloying Pt with Mg. Supplementary Fig. 22 illustrates the H2-O2 fuel cell polarization and power density curves of the two PtMg/C samples and Pt/C recorded before the durability test. We postulate that the introduction of Mg atoms stabilizes the Pt atoms since both the stability and activity in the PtMg/C samples were higher than commercial Pt/C.It is also worth noting that despite the similar electrocatalyst preparation method, the performance of PtMg/C-31 was higher than PtMg/C-11. The difference in performance between the two electrocatalysts may be visualized in the N2-derived CVs shown in Supplementary Fig. 23. In particular, the nature of the humps observed in the hydrogen adsorption and desorption regions (green-shaded region) under both half-cell (Supplementary Fig. 23a) and full-cell (Supplementary Fig. 23b) clearly shows that the adsorption on the PtMg/C-31 is significantly different from that of PtMg/C-11. This indicates that the exposed sites on the two electrocatalysts may be considerably different. In addition, the electrochemical surface area of PtMg/C-31 is substantially higher than that of PtMg/C-11. This is probably due to the higher exposed sites of Pt in the PtMg/C-31 in comparison to that of PtMg/C-11. Hence the difference in performance between the two electrocatalysts may be attributed to the relative composition of the catalyst, the exposed Pt sites, and the difference in ECSA. We also evaluated the performance of additional control samples, specifically, a PtMg alloy with a Pt to Mg precursor mass ratio of 2 to 1 (PtMg-21). Our findings reveal a consistent trend in stability akin to that observed for PtMg/C-31 and PtMg/C-11. Notably the 4.5% drop in mass activity after 30k cycles underscores its stability under H2-O2 conditions (Supplementary Fig. 24). However, despite its stability, the mass activity at 0.9 V stands at 0.44 A mgPt−1 compared to the 0.50 A mgPt−1 recorded for PtMg/C-31. Furthermore, to emphasize the stability of the PtMg alloys is not solely attributable to the particle size effects, we synthesized Pt/C under identical synthesis conditions to PtMg/C except for the omission of Mg (Supplementary Fig. 25). The resulting polarization curves under H2-O2 and H2-air conditions show values of 1340 and 630 mW cm−2 respectively, with a 15% degradation in peak power density after 30,000 cycles which are much lower than those found for PtMg/C-31. The TEM analyses (Supplementary Fig. 25c, d) show a transformation in the particle size of Pt/C-760 from 4.0 to 7.0 nm before and after ADT. These findings underscore the enhancement in both activity and stability conferred by Mg on Pt in PtMg alloys.In addition to the electrochemical performance reported above, structural characterization conducted after the 30,000 ADT shows that the Pt/C particles increased from their initial average size of 1.8 nm to 4.0 nm, while that of the PtMg/C-31 increased from 5.8 to 6.0 nm (Supplementary Figs. 26 and 27). Significantly, the STEM and corresponding EDS elemental mapping images (Supplementary Figs. 28–30) of PtMg/C-31 NPs obtained before and after 30,000 ADT cycles show that the Pt and Mg are evenly distributed throughout the particles with an average Pt to Mg atomic ratio of 84: 16 after ADT, which represents a small decrease in Mg compared to the initial Pt to Mg average atomic ratio of 78 to 22. Furthermore, post-ADT XPS analysis in Supplementary Fig. 31 shows that the Pt4f doublets obtained after 30,000 cycles are shifted by 0.32 eV to higher binding energies relative to that initially recorded in Fig. 2b. We hypothesize that this shift in binding energy can be attributed to two factors; changes in the electronic structure of Pt atoms on the surface of the catalyst or the surface oxidation of Pt atoms. The change in the electronic structure of Pt could lead to a further downward alteration in the d-band center which would move the antibonding orbitals closer to the Fermi level. The binding energies are related to the energy required to remove an electron from a specific energy level in a material. Hence, as the Fermi-level moves closer to the antibonding states (energetically unfavorable state), the density of states (DOS) around the Fermi-level decreases, resulting in higher binding energy for the Pt 4 f electrons. As the anti-bonding orbitals move closer to the Fermi level, it is expected that there would be a weak interaction between the Pt atoms and the various reaction intermediates (O*, OH*, OOH*). This probably accounts for the further positive shift in the CVs recorded in Supplementary Fig. 21. On the other hand, during the ADT cycles, some Mg leaches out of the nanoparticles (average concentration changes from ∽22 to ∽16%, Supplementary Figs. 28 and 29). Such a loss in Mg concentration could lead to a decreased interaction between Pt and Mg atoms, which would make the surface Pt atoms more susceptible to oxidation. It is not clear in our present study which of the two factors mentioned earlier contributes the most to the observed shift. However, we are certain the shift is a result of a combination of the aforementioned factors. A similar shift in binding energy was also observed for PtMg/C-11 after ADT (Supplementary Fig. 32). The activity and stability exhibited by the PtMg/C-31 cathode place it amongst the most competitive platinum-alkaline earth metal, platinum-rare earth metal, and platinum-early transition metal ORR catalysts reported to date (Supplementary Table 2). The activity and stability of this catalyst are also comparable to some platinum-transition metal catalysts measured under H2-air conditions (Supplementary Table 3).Computational resultsDFT calculations were performed to elucidate the origin of the higher ORR activity and durability of PtMg NPs compared to Pt NPs. From the XRD and TEM data (Fig.1, Supplementary Fig. 2), the particles are mainly composed of L12 Pt3Mg structures with a Pt-rich shell exposing (111) facets. Therefore, we use this structural model and refer to it as Pt3Mg/Pt(111) in this section. A Bader charge analysis reveals charge transfer from Mg to Pt ( ≈ 0.54 e- per Pt atom) leading to charge separation and stronger ionic bonding, resulting in compressive strain in the Pt3Mg lattice as compared to Pt (0.37%, as shown in Fig. 5d which agrees with experiment; Supplementary Fig. 2). The LSV experiments in Fig. 3a show that Pt3Mg has a higher on-set potential (lower over-potential) as compared to commercial Pt/C. To gain insight into the origin of the low overpotential, we considered the 4-step pathway mechanism for ORR considering the intermediates; OOH*, O*, OH* as shown in Fig. 5a. We constructed free energy diagrams at three different potentials U = 0 V, 1.23 V and at the onset potential. The onset potential is the lowest electrode potential at which all the reaction steps are downhill in the potential energy diagram (each step is exothermic). Our computational results from Fig. 5b, c show that the on-set potentials for Pt(111) and Pt3Mg/Pt(111) are 0.66 and 0.73 V, respectively, indicating higher ORR activity for Pt3Mg/Pt(111) which is in agreement with our experimental results. The higher on-set potential (lower overpotential) of Pt3Mg is due to a combination of compressive strain and ligand effects resulting in weaker binding energies of OOH*, O* & OH* adsorbates compared to Pt. To further investigate the contributions of strain and ligand effects, we compared O* binding energy (where the optimal \({E}_{O}\) is set to zero) on Pt3Mg/Pt1L(111) vs strained Pt(111) as shown in Fig. 5d. In Fig. 5d, strained Pt monolayer (0.37% compressive strain) of Pt3Mg/Pt1L(111) has an O binding energy of \(\Delta {E}_{O}=- \! 0.015 \, {\text{eV}}\) which is higher than the \(\Delta {E}_{O}=- \! 0.17 \, {\text{eV}}\) on Pt(111) slab with the same compressive strain. We can then infer that strain effects play a minor role and ligand effects are more important for the weak O binding on Pt3Mg, which in turn is responsible for lower overpotential for ORR on Pt3Mg.Fig. 5: DFT Calculations.a ORR 4-step pathway for the different intermediates on Pt3Mg/Pt1L(111). ORR free energy diagrams for b Pt(111), c Pt3Mg/Pt1L(111), d \(\Delta {E}_{O}\) vs compressive strain of (111) surface, where \(\Delta {E}_{O}={E}_{O}-{E}_{{O}_{{optimal}}}\), e Mg vacancy formation energies for Mg atoms from the surface (1L) and sub-surface (2 L and 3 L) layers of Pt3Mg(111), f Correlation between vacancy formation energies and formation energies for Pt3Mg vs previously reported L10-PtM systems4.Another key feature of the Pt3Mg NPs is their high durability and their activity retention compared to commercial Pt/C. To understand the high durability, we have calculated Mg vacancy formation energies for Pt3Mg/Pt(111) where a higher vacancy formation energy indicates it is thermodynamically less favorable to form the vacancy (more durable). In Fig. 5e, the Mg vacancy formation energies are shown for the different layers of Pt3Mg which increases as Mg vacancies are formed from the inner layers (Mgvac−1L = 2.72 eV, Mgvac−2L = 3.04 eV, Mgvac−3L = 3.17 eV). This indicates that it will be more difficult to remove Mg atoms from the Pt3Mg once we have 1-2 monolayers of Pt covering Pt3Mg NPs which agrees with our elemental mapping results from Fig. 2a. We subsequently compared the vacancy formation energy of Pt3Mg with the other most stable forms of PtM alloys (M=Zn, Fe, Co, Ni) [values reported in J. Liang et al4.] and plotted them against their corresponding formation energies in Fig. 5f. Since Pt3Mg has higher vacancy formation energy compared to previously synthesized L10-PtMs, we can conclude that Pt3Mg is more durable than other PtM alloy NPs. More DFT data are presented in Supplementary Figs. 33 and 34. In particular, the Pt vacancy formation energies presented in Supplementary Fig. 34 show that it is more difficult to create a vacancy on the Pt(111) surface of Pt3Mg as compared to that of pure Pt. This can account for the lower surface oxidation and more durable ORR performance of PtMg/C-31 in comparison to that of commercial Pt/C (Fig. 2b).
