Although impressive accomplishments have been achieved in all aspects of proton exchange membrane fuel cells (PEMFCs), the successful commercialization of PEMFCs has been hindered by the performance, durability and cost of electrocatalysts for the oxygen reduction reaction (ORR) [1-5]. Platinum-based core-shell nanostructures are widely accepted as one of the most promising ORR electrocatalysts [6, 7]. The thickness of the Pt or Pt alloy shell coating on a non-Pt core plays an important role in the reduction of the amount of Pt used [8, 9]. Furthermore, the structure-induced lattice compressive/tensile strain and electronic effect are known to affect the ORR reactivity through the adjustment of the d-band center energy and they can be designed and tuned by manipulating the particle composition and morphology [10-12].

Core-shell nanoparticles are generally prepared through underpotential deposition (UPD), seed-mediated growth and chemical/electrochemical dealloying [10, 13, 14]. A series of non-Pt cores coated with Pt monolayer have been constructed using UPD by Adzic et al. [15]. However, the difficulty of performing a large-scale synthesis without sacrificing ORR activity has been a great challenge for establishing the widespread use of this method. Seed-mediated growth was employed by Xia et al. [16] to fabricate Pd@Pt nanocatalysts. Unfortunately, the obtained nanocatalysts are usually bigger than 20 nm, which decreases the utilization of precious metal. Furthermore, high-boiling point surfactants, such as poly-(vinylpyrrolidone) (PVP), are present during the preparation procedure, which makes it difficult to remove the surfactant completely from the surface of the catalyst. PtCu@Pt was obtained by voltammetric dealloying of Cu from Pt/Cu alloys [14]. This method complicates the preparation process and results in a waste of dealloyed metal. Guo et al. [17] obtained hierarchical platinum-cobalt nanowires with excellent ORR performance by heating oleylamine containing Pt(acac)₂, Co(acac)₃, CTAC and glucose at 160 ℃. Recently, Guo et al. [18] synthesized a biaxially strained PtPb/Pt core/shell nanoplate as an electrocatalyst by heating 1-octadecene and oleylamine containing Pt(acac)₂, Pb(acac)₂ and L-ascorbic acid at 160 ℃. The high price of the organic precursor and solvent, difficulty of removing the surfactant and high preparation temperature constrained the widespread use of this synthesis procedure for the preparation of electrocatalysts.

Herein, we propose a facile and clean one-pot method for the preparation of highly dispersed Pd@Pt nanoparticles at room temperature without the assistance of high-boiling point surfactants or solvents. The schematic diagram of the preparation of the Pd@Pt nanostructure is shown in Scheme 1. First, molecular hydrogen adsorbs on the surface of the Pd particles and then dissociates into hydrogen atoms. Next, the platinum precursor, K₂PtCl₄, is added to the solution and reduced into elemental Pt on the surface of the Pd particles. Finally, during cyclic voltage scanning, the Pd nanoparticles are completely wrapped by Pt, and the Pd@Pt structure is fabricated successfully. More importantly, molecular hydrogen can also dissociate into atoms on the surface of the platinum as a reductant, and the K₂PtCl₄ can be reduced to deposit on the surface of platinum. Hence, the shell thickness can be controlled by adjusting the amount of Pt precursor and the electrocatalytic performance can be optimized correspondingly.

Pd_core@Pt_shell supported on Vulcan XC-72 was synthesized through the reduction of K₂PtCl₄ by dissociated hydrogen atoms adsorbed on the surface of the Pd core at room temperature. To prepare Pd@Pt nanoparticles with different composition, the amount of K₂PtCl₄ and Pd/C was altered. The following procedure was used to prepare the electrocatalyst. First, 50 mg Pd/C (BASF) was dispersed in 50 mL ethanol, followed by sonication for 1 h. The resulting suspension was stirred during the synthesis procedure. Different amounts of an aqueous solution of K₂PtCl₄ (49.7 mmol/L; 645, 1290 and 2580 µL for Pd₇₁@Pt₂₉/C, Pd₅₇@Pt₄₃/C and Pd₄₀@Pt₆₀/C, respectively) were added into the reaction mixture, which was purged by H₂ at 10 mL/min for 1 h. The reaction mixture was sealed and stirred for 10 h. H₂ at 10mL/min was introduced for 1 h to ensure the complete reduction of K₂PtCl₄. The solution was centrifuged/washed repeatedly in water/ethanol and dried under vacuum. The electrochemical scanning for different samples was performed in N₂-saturated HClO₄ aqueous solution (0.1 mol/L) at 100 mV/s between 0.05 and 1.20 V (vs. RHE) for 20 cycles. The nominal Pd@Pt compositions were 71/29, 57/43 and 40/60. For comparison, Pt/C was produced by reducing K₂PtCl₄ with hydrogen in the absence of Pd_core, and PdPt/C was acquired without hydrogen.

The morphology of the nanoparticles was investigated using transmission electron microscopy (TEM; JEOL JEM-2000EX electron microscope operating at 120 kV). High resolution TEM (HRTEM) and selected area electron diffraction (SAED) images were obtained with a JEOL JEM-2100 TEM operated at 200 KV and equipped with an Oxford X-Max detector. Energy dispersive X-ray spectroscopy (EDS) was performed on a JEOL JEM-2100F TEM operated at 200 kV and equipped with an Oxford X-Max detector. The high-angle annular dark field scanning TEM (HAADF-STEM) observations were conducted on a FEI Tecnai G2 F30 microscope. X-ray powder diffraction (XRD) patterns were obtained using a PANalytical X'Pert Pro Super diffractometer with Cu K_α radiation (λ = 0.154178 nm). The bulk phase and surface compositions of Pd@Pt/C were determined by inductively coupled plasma (ICP) (Perkin-Elmer Optima 2000 DV) and X-ray photoelectron spectroscopy (XPS) (Thermo Multi-Lab 2000 spectrometer).

Rotating disk electrode (RDE) tests were performed on a CHI 730D electrochemical station (CH Instruments, Inc.) with a conventional three electrode electrochemical system in HClO₄ aqueous solution (0.1 mol/L). All electrode potentials are given with reference to the reversible hydrogen electrode (RHE). The cyclic voltammogram (CV) measurements were conducted in N₂-saturated HClO₄ aqueous solution (0.1 mol/L) at 50 mV/s. The ORR polarization curves were recorded positively at the rate of 10 mV/s in O₂-saturated HClO₄ aqueous solution (0.1 mol/L) at 1600 r/min. The specific kinetic current density was obtained according to the Koutecky-Levich equation. Accelerated degradation tests (ADT) were performed by cycling the potential between 0.60 and 1.20 V (vs. RHE) at 50 mV/s in HClO₄ aqueous solution (0.1 mol/L) under continuous N₂ bubbling for 1500 cycles. CV curves and ORR polarization curves were measured after the ADT by using the method mentioned above. The specific electrochemical active area (ECA) was calculated through the integration of the hydrogen desorption area in the CVs using 0.21 mC/cm as the conversion factor.

TEM images of Pd/C, Pd₇₁@Pt₂₉/C, Pd₅₇@Pt₄₃/C, Pd₄₀@Pt₆₀/C, Pd₃₅Pt₆₅/C and Pt/C are displayed in Fig. 1. As shown in Fig. 1(a), the Pd nanoparticles were uniformly distributed on carbon and had an approximately ellipsoidal shape. The Pd₇₁@Pt₂₉, Pd₅₇@Pt₄₃ and Pd₄₀@Pt₆₀nanoparticles were supported on carbon and had an ellipsoidal shape similar to that of the Pd nanoparticles. The favorable dispersity and ellipsoidal morphology of the Pd@Pt nanostructure was consistent with the proposed mechanism, in which the platinum was deposited on the surface of the palladium particles to form the Pd₇₁@Pt₂₉, Pd₅₇@Pt₄₃ and Pd₄₀@Pt₆₀. In comparison, the morphology of the Pd₃₅Pt₆₅in Fig. 1(e) was closer to a sphere, and the particle sizes were less uniform than those of the core-shell particles. The phenomenon above suggested that a uniform Pd@Pt nanostructure could not be obtained without the assistance of H₂. As for Pt/C in Fig. 1(f), the particles of Pt did not have a defined shape, and significant aggregation was observed. The different particle shapes of Pt and Pd@Pt demonstrated the importance of Pd in the preparation of the core-shell structure. Even though Pt can be reduced by H₂without the presence of the Pd core, the deposition of Pt on the support was random. If platinum was reduced on carbon or in solution instead of on the surface of the Pd core with the aid of dissociated H, the so-called core-shell particles scattered irregularly or even aggregated, like those of Pt. However, the consistent particle morphology and remarkable dispersibility of the core-shell structure supported the proposed mechanism in which platinum was reduced and deposited simultaneously on the surface of the core with the aid of dissociated hydrogen atoms.

According to the TEM images of Pd/C, Pd₇₁@Pt₂₉/C, Pd₅₇@Pt₄₃/C, Pd₄₀@Pt₆₀/C, Pd₃₅Pt₆₅/C and Pt/C in Fig. 1, the sizes of the metal nanoparticles was measured and the statistical analysis results are shown in Fig. 2. Based on the statistical results, the Pd particle size varied between 2 and 5 nm. The sizes of the Pd₇₁@Pt₂₉, Pd₅₇@Pt₄₃ and Pd₄₀@Pt₆₀ particles was in the range of 3-6 nm, 3-7 nm and 4-7 nm, respectively. As expected, the particle size increased with increase of Pt/Pd ratio. Furthermore, the particle size distribution intervals of the core-shell particles were close to that of Pd, indicating that Pt deposited on the surface of Pd particles uniformly. For comparison, the particle size distribution interval of Pd₃₅Pt₆₅ (3-8 nm) was obviously larger than that of Pd₄₀@Pt₆₀ (4-7 nm), corresponding to the random formation of Pd-Pt alloy. Furthermore, the particle sizes of the Pt particles varied from 2 to 20 nm, revealing an entirely random formation of Pt. All the data above highlights the crucial role of the dissociated hydrogen atom as a reductant and Pd as a deposition core. The average particle sizes of the Pd, Pd₇₁@Pt₂₉, Pd₅₇@Pt₄₃ and Pd₄₀@Pt₆₀ particles were 3.4±0.8 nm, 4.3±0.9 nm, 4.9±1.2 nm and 5.2±0.6 nm, respectively. Therefore, the thicknesses of the Pt shells of Pd₇₁@Pt₂₉, Pd₅₇@Pt₄₃ and Pd₄₀@Pt₆₀ were estimated to be approximately 0.45, 0.75 and 0.90 nm, respectively, by excluding the size of Pd core. Considering that the interplanar spacing of the Pt(111) lattice plane was 0.23 nm, the Pd cores for Pd₇₁@Pt₂₉, Pd₅₇@Pt₄₃ and Pd₄₀@Pt₆₀ were wrapped with approximately 2, 3 and 4 layers of Pt shell, respectively. Therefore, the shell thickness could be easily manipulated by simply adjusting the amount of the K₂PtCl₄.

In comparison, the particle size of Pd₃₅Pt₆₅ prepared without the presence of H₂ was 5.5±1.4 nm, and the particle size of Pt synthesized in the absence of a Pd_core was 7.6±5.0 nm. The particle size distribution windows of Pd₃₅Pt₆₅and Pt were significantly bigger than those of the Pd_core@Pt_shell particles obtained with the aid of H₂. The remarkable monodispersity of the core-shell particles was attributed to the reducing properties of the dissociated hydrogen atom on the surface of the core.

As shown in Fig. 3, high resolution TEM (HRTEM) and selected area electron diffraction (SAED) were performed to explore the crystal nanostructure of Pd₅₇@Pt₄₃. The interplanar spacing of 0.2272 and 0.2292 nm in Fig. 3(a) corresponded to the (111) crystal plane of Pt and Pd (JCPDF #04-0802 and #46-1043), respectively. Furthermore, the crystal faces of numerous nanostructures were identified, according to the SAED image in Fig. 3(b). These results clearly indicated the remarkable crystallinity of the Pd₅₇@Pt₄₃ nanoparticles. Because of the similar crystal structures Pt and Pd, the crystal boundary between the Pd_core and Pt_shell could not be observed in the HRTEM image. The HRTEM and SAED observations suggested that Pt was reduced by the dissociated hydrogen atom on the surface of the core, and the shell was fabricated in epitaxial growth mode [13, 19].

The formation of Pd@Pt and the unique function of H₂ are discussed below. Based on the Pd@Pt structure formation mechanism, some vacancies exist on the interfaces between the Pd core and the carbon support. At high potentials, the interaction between the Pd and the support contributes to the stabilization of Pd, which restrains the dissolution of the Pd on the interfaces than on those exposed in solution. The surface free energy of the Pt shell with vacancies decreases, during which Pt atoms move on the particle surface through an adatom-substrate atom replacement mechanism [20], and the Pt spreads across the Pd core surface and repairs the vacancies. This procedure can be initiated by potential scanning in solution. In addition, Pd dissolution and deposition could be neglected because the dominating Pd atoms are either surrounded by Pt or the support. Hence, Pt rather than Pt/Pd is deposited on the Pd core.

It is well established that H₂ absorbs on the surface of Pd and Pt and dissociates into H atoms. Those dissociated H atoms cover the surface of the Pd core and show strong reduction ability. Then, PtCl₄²⁻ is reduced by the dissociated H atoms and the obtained Pt atoms are deposited on the Pd core surface, replacing the sites occupied by the H atoms. Furthermore, the similar lattice structures of Pd and Pt are reported to contribute to the layer-by-layer growth mode [13] and hence the formation of highly dispersed Pd@Pt.

The elemental line profiles of specific Pd₅₇@Pt₄₃ and Pd₃₅Pt₆₅ nanoparticle detected by HAADF-STEM-EDX analysis are illustrated in Fig. 3(c) and (d). As displayed in Fig. 3(c), the single peak of Pd and the even signal strength of Pt for Pd₅₇@Pt₄₃ indicated that Pd was enriched in the core and Pt was distributed uniformly around the particle. The element distribution features above corresponded to a core-shell nanostructure. For comparison, both the signal strength of Pd and Pt demonstrated a single peak, and the signal intensity changed in the same trend at the identical position, revealing the uniform distribution pattern of Pd and Pt in the Pd₃₅-Pt₆₅ alloy nanoparticle. The significantly different element distribution pattern above suggested the crucial role of H₂ in the formation of a core-shell nanoarchitecture. Without the presence of H₂, the galvanic replacement reaction between Pd and K₂PtCl₄ progressed, and Pd-Pt alloy with uniform distribution of both Pd and Pt in the whole particle was formed. In contrast, the Pt shell was reduced by absorbed and dissociated H₂ covering the core, and a Pd_core@Pt_shell structure developed. The high-angle annular dark field scanning TEM (HAADF-STEM) observations were conducted on a FEI Tecnai G2 F30 microscope, as displayed in Fig. 3(e). The elemental mappings of one particle showed that the distribution of Pt and Pd elements was homogeneous and overlapped, which was consistent with a highly dispersed core-shell nanostructure for the Pd@Pt with a few Pt shell layers.

X-ray powder diffraction (XRD) patterns were obtained using a PANalytical X'Pert Pro Super diffractometer with Cu K_α radiation (λ = 0.154178 nm), as shown in Fig. 4(a). According to the PDF card, it is well established that the crystal structures of Pd and Pt were similar. Hence, if separate Pd or Pt nanoparticles were present, some satellite peaks would appear in the obtained XRD patterns. However, the diffraction peaks for the different crystal planes were evident and separated completely, and no satellite peaks were observed, as shown in Fig. 4(a). The phenomenon above suggested that all the metal particles existed in the form of a highly dispersed Pd@Pt architecture.

Furthermore, the diffraction positions at approximately 2θ = 40° for different crystals are listed in Fig. 4(b). Based on the PDF card (#65-2867 and #04-0802), the diffraction position of Pd was more positive than that of Pt, indicating that Pd had a smaller lattice interplanar distance. Compared with Pd, the diffraction position of Pd₇₁@Pt₂₉ shifted negatively, which resulted from the Pt-shell-induced lattice expansion. The diffraction position of Pd₅₇@Pt₄₃ shifted to a more negative value with increasing Pt shell thickness, corresponding to an increased lattice expansion. However, the lattice strain of Pd₄₀@Pt₆₀ compared to that of Pd₅₇@Pt₄₃ remained unchanged, because the crystal structure reached a stable state and no further core-shell-induced lattice strain occurred after the thickness of the Pt was thick enough for Pd@Pt.

Survey and high-resolution Pt 4f XPS profiles of Pd₇₁@Pt₂₉/C, Pd₅₇@Pt₄₃/C, Pd₄₀@Pt₆₀/C and Pd₃₅Pt₆₅/C are shown in Fig. 5. The existence of Pt and Pd was identified in the survey spectra, as in Fig. 5(a). According to Kitchin et al. [21], the electronic and chemical properties of Pt surfaces can be modified by subsurface 3d transition metals. Interactions between the Pt surface and the subsurface 3d metals broaden the Pt surface d-band and lower the Pt d-band energy, leading to weaker dissociative adsorption energy of oxygen on the surface. The surface d-band center of Pd@Pt can be determined with the aid of XPS [22]. According to Fig. 5(b), the Pt 4f binding energies of Pd₇₁@Pt₂₉/C, Pd₅₇@Pt₄₃/C, Pd₄₀@Pt₆₀/C and Pd₃₅Pt₆₅/C were 71.5, 71.7, 71.7 and 71.4 eV, respectively. As for core-shell structures, the Pt 4f binding energy increased by 0.2 eV when the Pt/Pd ratio increased from 29:71 to 43:57. The increase of the binding energy of Pt 4f from Pd@Pt resulted from the broadening of the Pt valence band owing to the incorporation of the Pd valence states to the Pt valence band structure [23]. The phenomenon above leads to a shift of the d-band center of Pt away from the Fermi level, which has a positive impact on ORR activity. Furthermore, when the Pt/Pd ratio increased from 43:57 to 60:40, the Pt 4f binding energy remained unchanged. This was because the electronic interaction effect between Pt and Pd weakened with the increase of the shell thickness. In comparison, the Pt 4f binding energy of the Pd₃₅-Pt₆₅ alloy was 0.1 eV smaller than that of Pd₇₁@Pt₂₉. This phenomenon demonstrated that the electronic effect of the Pd_core@Pt_shell structure was apparently more significant than that of the alloy structure.

To investigate the oxidation state of Pt in different samples, the Pt 4f XPS spectra were deconvoluted into two pairs of peaks, which were assigned to Pt⁰ and Pt²⁺, respectively [24]. The deconvolution results are shown in Fig. 5(c-f) and Table 1. The binding energies of the Pd@Pt electrocatalysts increased in the order Pd₇₁@Pt₂₉(71.6 eV) < Pd₅₇@Pt₄₃ (71.7 eV) < Pd₄₀@Pt₆₀ (71.8 eV). This trend was consistent with the increase in the shell thickness. As Pt was more electronegative than Pd, the binding energies of Pt⁰ 4f_7/2 for different Pd@Pt structures increased with the increase of the Pt shell thickness. The binding energy of Pt⁰ 4f_7/2 for Pd₃₅Pt₆₅ (71.5 eV) was even smaller than that for Pd₇₁@Pt₂₉ (71.6 eV), even though the Pt/Pd ratio of Pd₃₅Pt₆₅ was significantly bigger than that of Pd₇₁@Pt₂₉. The phenomenon above demonstrated that the role of the core-shell architecture was more prominent than that of the electronegativity of the as-synthesized Pd@Pt and PdPt alloy. With the increase of shell thickness, the metal-Pt ratio increased in the same order: Pd₇₁@Pt₂₉(80.6%) < Pd₅₇@Pt₄₃ (82.8%) < Pd₄₀@Pt₆₀ (83.6%). The oxidation rate of the electrocatalysts above was obviously lower than Pt/Pd₃Co (35%) [25], Pt (28%) [26], PtPd (65%) and PtFe (40%) [27] and PtNi (26%) and PtCoCr (24%) [28]. Based on the reported result by Arico et al. [29], the lower oxidation-state ratio of Pt in the samples contributed to a higher electrocatalytic performance towards oxygen reduction reaction.

In addition, the surface Pt/Pd atomic ratio obtained from XPS measurements was 0.41, 1.06 and 2.28 for the Pd₇₁@Pt₂₉, Pd₅₇@Pt₄₃ and Pd₄₀@Pt₆₀ nanoparticles, respectively. The Pt/Pd atomic ratio obtained by ICP detection was 0.40, 0.74 and 1.52 for the Pd₇₁@Pt₂₉, Pd₅₇@Pt₄₃ and Pd₄₀@Pt₆₀ nanostructures, respectively, which was consistent with the increase in Pt shell thickness. The results above confirmed that Pt was enriched on the surface of the Pd@Pt nanostructures. More importantly, the Pt surface enrichment became more severe with the increase of the Pt shell thickness. In contrast, the Pt/Pd atomic ratio of Pd₃₅Pt₆₅ in bulk (0.53) was greater than that on the surface (0.34), indicating a Pd-enriched surface.

Fig. 6(a) displays the CV curves of Pt/C(JM), Pd₇₁@Pt₂₉/C, Pd₅₇@Pt₄₃/C, Pd₄₀@Pt₆₀/C, Pt/C and Pd₃₅Pt₆₅/C. Pd/C showed distinct hydrogen desorption peaks between 0.05 and 0.11 V [13]. As for Pt/C(JM), two significant hydrogen desorption peaks were observed between 0.05 and 0.4 V [30]. For Pd₇₁@Pt₂₉/C, which had two layers of Pt, the characteristic hydrogen desorption peak of Pd was much weaker than the Pd/C, and the characteristic Pt hydrogen desorption peak appeared, indicating the successful coating of Pt on the surface of Pd. The Pd hydrogen desorption peak was not observed for Pd₅₇@Pt₄₃/C, which had three layers of Pt; only one Pt hydrogen desorption peak was observed. The hydrogen behavior of Pd₄₀@Pt₆₀/C, between 0.05 and 0.4 V was identical to that of Pt/C, demonstrating the complete coverage of Pd by Pt. Furthermore, the peak potential of the surface metal oxides reduction peak (Por) in the oxygen desorption area on the CV profiles revealed the adsorption strength of the OH_ads species on the surface of the electrotatalysts [19, 30]. The order of the Por was Pd₄₀@Pt₆₀/C (784 mV) > Pd₅₇@Pt₄₃/C (776 mV) > Pd₇₁@Pt₂₉/C (772 mV) > Pt/C (746 mV). The OH radical is one of the most crucial intermediates during the ORR, and its adsorption on Pt is widely believed to block oxygen adsorption sites and hinder the ORR [19, 31]. The positive shift of the peak potential of the surface oxides reduction peak corresponded to the decreasing desorption strength of the oxide adsorption species, contributing to enhancement of the ORR performance [30]. The polarization curves of different samples are shown in Fig. 6(b). The half-wave potential of different samples decreased in the order Pd₄₀@Pt₆₀/C ≈ Pd₅₇@Pt₄₃/C > Pd₇₁@Pt₂₉/C ≈ Pt/C(JM) > Pd₃₅Pt₆₅/C > Pt/C. This order indicated superior ORR activity of the Pd_core/Pt_shell nanostructure, consistent with the correlation between electrocatalytic activity and the adsorption strength of OH_ads species on the surface of electrocatalysts discussed above.

As shown in Fig. 6(c), the hydrogen desorption area of Pd₅₇@Pt₄₃ was significantly bigger than that of Pt/C(JM), demonstrating a larger electrochemical surface area (ECSA) of Pd₅₇@Pt₄₃. The increased ECSA of Pt for Pd₅₇@Pt₄₃ benefited from the high dispersity of Pt in the shell of the core-shell structure. Furthermore, the higher Pt utilization for Pd₅₇@Pt₄₃ after degradation tests was still prominent. Based on Fig. 6(d), the half-wave potential of both Pd₅₇@Pt₄₃/C and Pt/C(JM) after the degradation test shifted negatively. The half-wave potential of Pd₅₇@Pt₄₃/C after the stability test still exceeded that of Pt/C(JM) before the stability test, indicating the remarkable electrocatalytic activity and stability of Pd₅₇@Pt₄₃/C compared with that of Pt/C(JM).

To compare the electrocatalytic activity per active site and mass for different samples quantitatively, the area-specific activity (SA) and mass-based activity between 0.85 and 0.90 V were calculated and are displayed in Fig. 7(a) and (b). The SA of Pt/C(JM) was obviously smaller than those of the Pd_core@Pt_shell nanoparticles, which was attributed to the electron interaction impact between the Pd_core and Pt_shell, as discussed in the XPS results, and consistent with the results reported by other researchers [19, 30]. Specifically, the SA of the core-shell particles increased in the sequence Pd₇₁@Pt₂₉ < Pd₅₇@Pt₄₃ < Pd₄₀@Pt₆₀. This result agreed well with the increase of the Pt shell thickness, indicating that a thicker Pt-shell contributed to improved SA. For Pt nanoparticles bigger than 2 nm, the concentration of surface terraces increased with the growth of the particle size [32]. Compared with the surface step platinum atoms, terrace platinum atoms possessed bigger coordination numbers, a bigger d-band width, a lower energy of the d-band center, weaker adsorption strength to ORR intermediates, and therefore more desirable ORR performance [33-36]. Based on the discussion above, the impressive SA of Pt/C resulted from the particle size effect. Finally, the SA of Pd₃₅Pt₆₅ was evidently inferior to that of Pt/C(JM). This indicated that the specific activity per active site of Pd-Pt alloy prepared without the presence of H₂ was lower than that of the Pt surface.

The specific mass activity (MA) is the combined result of both platinum dispersity and activity per active site. Based on Fig. 7(b), the MA of the Pd@Pt structure was apparently greater than that of Pt/C(JM). This was attributed to two facts. (1) The core-shell architecture contributed to the high dispersity of platinum, and (2) the electronic effect between the Pd_core and Pt_shell improved the area specific activity of platinum. As expected, the MA of the core-shell structure increased in the order: Pd₄₀@Pt₆₀ < Pd₅₇@Pt₄₃ < Pd₇₁@Pt₂₉, which was consistent with the decrease of the shell thickness. Therefore, it clear that the platinum dispersity increased and the interaction between the Pd_core and Pt_shell strengthened when the shell thickness decreased. The smaller MAs of Pd₃₅Pt₆₅ and Pt were because of the adverse SA and inferior usage of Pt, respectively.

The SA and MA at 0.9 V of the electrocatalysts before and after the accelerated degradation test (ADT) are listed in Fig. 7(c) and (d). As expected, the SA of Pd@Pt improved after ADT, which resulted from structure transformation during the periodic voltage scan. In detail, the dissolution and redeposition of Pd exposed to the solution led to a shell surface completely wrapped by a Pt-enriched surface with a subsurface Pd-Pt alloy [19, 37]. The Pt skin with a subsurface alloy architecture had the benefit of improving its electrocatalytic activity and stability [30, 38-40]. For comparison, the Pt/C(JM) after ADT showed reduced electrocatalytic performance owing to (1) particle migration and/or coalescence and (2) electrochemical Ostwald ripening of the Pt particles [41]. The SA and MA of Pd₇₁@Pt₂₉ after ADT were 3.2 and 4.0 times higher, respectively, than that of Pt/C(JM) after ADT.

A highly dispersed Pd_core@Pt_shell nanostructure was constructed through the reduction of K₂PtCl₄ by dissociated hydrogen atoms at room temperature without the protection of surfactants or a high-boiling-point solvent. The shell thickness of these architectures, determined by TEM, varied from 2 to 4 layers of Pt. HRTEM and SAED demonstrated the remarkable crystallinity and the epitaxial growth mode of the Pd_core@Pt_shell nanostructure. The surface segregation of Pt was proven using ICP and XPS. The enhanced ORR electrocatalytic performance was assigned to the weak adsorption strength of OH_ads, resulting from the electron impact between the Pd_core and Pt_shell. For comparison, Pt/C synthesized with the absence of a Pd_core and PdPt/C was prepared without using hydrogen and showed unfavorable nanoparticle sizes and adverse electrocatalytic activity. After degradation tests, the Pd@Pt nanostructure experienced surface reconstruction, resulting in a stable Pt-rich surface with a subsurface Pd-Pt alloy and improved ORR performance. This facile and clean preparation method can also be used to prepare other core-shell architectures under mild preparation conditions, such as Ni_wire@Pt_skin, whose structure and electrochemical performance are currently under investigation in our group.