Achieving Highly Durable Random Alloy Nanocatalysts through

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Achieving Highly Durable Random Alloy Nanocatalysts through Intermetallic Cores Jocelyn T. L. Gamler,† Alberto Leonardi,‡ Hannah M. Ashberry,† Nicholas N. Daanen,† Yaroslav Losovyj,† Raymond R. Unocic,§ Michael Engel,‡ and Sara E. Skrabalak*,† †

Department of Chemistry, Indiana University, 800 East Kirkwood Avenue, Bloomington, Indiana 47405, United States Institute for Multiscale Simulation, Friedrich-Alexander Universität Erlangen-Nürnberg, Cauerstraße 3, 91058 Erlangen, Germany § Center for Nanophase Materials Sciences, Oak Ridge National Laboratory, One Bethel Valley Road, Oak Ridge, Tennessee 37831, United States ACS Nano Downloaded from pubs.acs.org by ALBRIGHT COLG on 04/08/19. For personal use only.

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ABSTRACT: Pt catalysts are widely studied for the oxygen reduction reaction, but their cost and susceptibility to poisoning limit their use. A strategy to address both problems is to incorporate a second transition metal to form a bimetallic alloy; however, the durability of such catalysts can be hampered by leaching of non-noble metal components. Here, we show that random alloyed surfaces can be stabilized to achieve high durability by depositing the alloyed phase on top of intermetallic seeds using a model system with PdCu cores and PtCu shells. Specifically, random alloyed PtCu shells were deposited on PdCu seeds that were either the atomically random face-centered cubic phase (FCC A1, Fm3m) or the atomically ordered CsCl-like phase (B2, Pm3m). Precise control over crystallite size, particle shape, and composition allowed for comparison of these two core@shell PdCu@PtCu catalysts and the effects of the core phase on electrocatalytic durability. Indeed, the nanocatalyst with the intermetallic core saw only an 18% decrease in activity after stability testing (and minimal Cu leaching), whereas the nanocatalyst with the random alloy core saw a 58% decrease (and greater Cu leaching). The origin of this enhanced durability was probed by classical molecular dynamics simulations of model catalysts, with good agreement between model and experiment. Although many random alloy and intermetallic nanocatalysts have been evaluated, this study directly compares random alloy and intermetallic cores for electrocatalysis with the enhanced durability achieved with the intermetallic cores likely general to other core@shell nanocatalysts. KEYWORDS: electrocatalysts, multimetallic, platinum alloys, nanoparticles, nanocatalysts

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specific shell thicknesses were superior catalysts for formic acid electrooxidation when compared to monometallic Pt and Pd reference catalysts of the same shape, respectively.8,9 Enhanced catalysis through core@shell architectures has been demonstrated by numerous groups.10−14 In these cases, strain introduced from the lattice mismatch shifts the d-band position of the catalytic surfaces and, in turn, strengthens or weakens surface−adsorbate interactions. We recently advanced this concept by straining random alloyed surfaces through core@shell architectures by building on top of intermetallic seeds of a different composition. This concept was demonstrated in a model PdCu B2@PtCu NP system. Significantly, the PtCu surface also provided ligand

roton-exchange membrane fuel cells (PEMFCs) are promising renewable energy devices; however, challenges remain in the development of cathode catalysts for the oxygen reduction reaction (ORR).1 Pt is traditionally used, but its high cost and susceptibility to poisoning make long-term use impractical.2−4 To overcome the high cost and unfavorable surface−adsorbate interactions of Pt, a second transition metal is often added. This modification decreases the amount of expensive metal loading and can be used to tune surface−adsorbate interactions in accordance with the Sabatier principle.2 Secondary metals can be incorporated to form dimers, random alloys, and intermetallics and even core@shell nanoparticles (NPs).5−7 In fact, the lattice mismatch between the components of a core@shell nanocatalyst can strain the surface, enhancing catalysis through tuning of surfaceadsorbate interactions. For example, our group demonstrated that core@shell Rh@Pt and Au@Pd nanocatalysts with © XXXX American Chemical Society

Received: October 19, 2018 Accepted: April 2, 2019

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Figure 1. TEM images of (a) PdCu A1 seeds and (b) PdCu A1@PtCu NPs; (c) powder XRD pattern of resulting A1@PtCu NPs showing only FCC structures, where the orange lines indicate an FCC Pt−Cu alloy and the red line the A1 phase; (d) STEM-EDX elemental mapping of A1@PtCu with (e) histogram of size distribution and (f) resulting line scan from single particle STEM-EDX with red line in d.

compared to random alloys.15 As these examples also illustrate, core@shell nanocatalysts with intermetallic cores are often compared to a Pt reference or NPs of the shell composition. However, these comparisons are insufficient to understand the true impact of the intermetallic core on catalyst performance as enhanced stability may simply arise from the core@shell architecture, for example. In addition, these comparisons often do not account for size differences between the random alloy and intermetallic NPs being compared.23 Such size differences are common due to the reliance on annealing of random alloy NPs to form intermetallic NPs, which is often accompanied by particle sintering.24 Annealing methods may also form a Pt skin17,18 that may not exist for the random alloy NPs, and thus, the catalytic comparison is not an isolation of one variable (e.g., core structure) but multiple (e.g., size and architecture). Our synthetic methods used in this manuscript avoid these limitations.

effects, which altered the electronic structure of the active sites though charge transfer from one surface atom to adjacent ones of a different composition.15 That is, this PdCu@PtCu catalyst employed strain (from core@shell) and ligand (from alloyed surface) effects together to tune the electronic structure to optimize surface−adsorbate interactions for catalysis.16 A standout feature of the PdCu B2@PtCu nanocatalyst was its enhanced catalytic activity and resistance to dissolution compared to a PtCu reference. The PdCu B2@PtCu catalyst showed a mass activity of 8.7 times that of the PtCu alloy and only saw a 12.3% loss in activity after stability testing.16 Other groups have similarly reported nanocatalysts with multimetallic or with intermetallic cores and either thin Pt or PtM alloyed shells that displayed both enhanced catalytic activity and durability when compared to Pt or PtM references.16−22 The increased durability of intermetallic nanocatalysts has been accounted for by their lower heats of formation and strong, regular heteroatomic bonding when B

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Figure 2. TEM images of (a) PdCu B2 seeds and (b) PdCu B2@PtCu NPs; (c) powder XRD pattern of resulting B2@PtCu NPs showing ordered B2 structure (blue lines) form the core and FCC (orange lines) structure from the shell. (d) STEM-EDX elemental mapping of B2@ PtCu with (e) histogram of size distribution and (f) resulting line scan from single particle STEM-EDX with red line in d.

that include core composition, shell thickness, and lattice mismatch. This work isolates the influence of the core’s crystal structure on nanocatalyst durability through precise control over composition and structure.

This work compares the electrocatalytic durability of two ORR catalysts in order to elucidate the effect of intermetallic cores. Specifically, two core@shell NP systems with identical core and shell compositions were targeted and evaluated, but the crystal structures of the cores were different. In one case, PtCu random alloyed shells were deposited on intermetallic PdCu B2 (ordered CsCl-like, Pm3m) seeds. In the second case, the same shells were deposited on PdCu A1 (random alloy FCC, Fm3m) seeds. These systems were selected because PtCu NPs have high activity for the ORR,25,26 and our prior work found that PdCu B2 seeds could effectively strain PtCu surfaces in a core@shell architecture to achieve high-performing ORR catalysts.16 These nanocatalysts were prepared by seed-mediated coreduction (SMCR), which is a versatile route to deposit alloyed surfaces on a variety of different seeds.28 With simple bimetallic NPs, catalytic durability can have many influencers including NP size, composition, and shape. When incorporating a core@shell structure, additional factors arise

RESULTS AND DISCUSSION Nanoparticle Synthesis and Structural Characterization. The two catalysts for comparison were prepared by SMCR. First, PdCu NPs, either with the random alloy (A1 FCC) or intermetallic (B2 CsCl-like) phase, were synthesized. The size and Pd-to-Cu ratio of the NPs were held constant. Then, PtCu random alloyed shells with the same Pt-to-Cu ratios and shell thicknesses were deposited on the seeds. Structural and compositional characterization of the resultant core@shell PdCu A1@PtCu and PdCu B2@PtCu nanocatalysts are shown in Figures 1 and 2, respectively. The alloy PdCu seeds (A1 FCC) have an average size of 7.3 ± 0.7 nm and were synthesized in a manner similar to the C

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ACS Nano report by Yu and co-workers (Figure 1a).24 Powder X-ray diffraction (XRD) is consistent with the A1 FCC phase (ICDD 01-071-7854) (Figure S1). SMCR was then used to deposit PtCu shells on these seeds. Specifically, PtBr2 and Cu(acac)2 (1:1 molar ratio) were co-reduced in the presence of 1,2dodecandiol (DDD), 1-octadecene (ODE), and oleylamine (OLA). The resulting core@shell NPs have an average size of 8.5 ± 0.7 nm (Figure 1b). The bulk composition was determined by energy-dispersive X-ray spectroscopy (EDX) interfaced with a scanning electron microscope (SEM-EDX) and found the NPs to contain 19 atomic % Pt; ICP-MS indicated 21.1 atomic % Pt (Table S1). XRD shows peaks consistent with the alloy PdCu seeds as well as shoulders to lower 2θ that are consistent with a Pt−Cu alloy (Figure 1c). The Pt−Cu alloy reference was calculated on the basis of the Pt/Cu ratios determined by SEM-EDX and Vegard’s law and is in good agreement. The spatial distribution of elements was evaluated by scanning transmission electron microscopy-EDX (STEM-EDX) and reveals that Pd is localized in the interior, Pt on the exterior, and Cu is distributed throughout (Figure 1d). This distribution is consistent with the desired core@shell architecture, which is further supported by the increase in average particle size after SMCR (Figure 1e) and line scan analysis where the highest intensity for Pt is at the edges of the particles and Pd at the center and the Cu signal spans the whole particle (Figure 1f). The PdCu A1@PtCu NPs were compared to PdCu B2@ PtCu NPs that were synthesized similarly to what was previously reported.16 Characterization of the PdCu B2@ PtCu particles is shown in Figure 2. The B2 PdCu seeds (Figure 2a) were 7.6 ± 0.3 nm in size, and the resulting core@ shell NPs (Figure 2b) had an average size of 8.8 ± 0.4 nm, giving a similar size and shell thickness (0.6 nm, approximately 3−5 monolayers depending on the exposed facet) to the PdCu A1@PtCu nanocatalysts. The bulk composition of PdCu B2@ PtCu evaluated by SEM-EDX indicated the NPs were 18 atomic % Pt and 18.7% by ICP-MS (Table S1), similar to the PdCu A1@PtCu catalyst. As anticipated, the XRD pattern for the PdCu B2@PtCu (Figure 2c) can be indexed with the B2 (ICDD 01-078-4406) CsCl-like intermetallic phase associated with the PdCu core (Figure S2) and the PtCu shell as an FCC structure. The shoulders at lower 2θ support the deposition of FCC shells on the intermetallic core. The elemental distribution shown by STEM-EDX (Figure 2d) shows the same elemental distribution as the PdCu A1@PtCu NPs. Once again, an increase in average particle size (Figure 2e) and the resulting line scan (Figure 2f) confirm the core@shell architecture. Atomic resolution images were also obtained for both samples by high angle annular dark field-STEM (HAADFSTEM) in an effort to provide greater understanding of the core@shell interface of these two core@shell nanocatalysts (Figure S3). However, precise information about faceting and interfacial mixing was not possible due to the small NP sizes and lack of Z-contrast that arises from their complex compositions. Our prior work, however, with slightly larger PdCu B2@PtCu NPs revealed epitaxial deposition in a manner that minimizes strain, and similar deposition likely occurs in these systems as well.16 Additionally, classical molecular dynamics simulations discussed later in this manuscript provide insight into the impact any Pd migration from the core to the shell may have on nanocatalyst durability as such

metal migration has been reported in Pd@Pt nanocatalysts.11,22 As the compositional and architectural features of these core@shell nanostructures are similar, they are an ideal platform to evaluate the impact atomic ordering of the core has on catalytic performance. Compressively strained Pt can destabilize strongly bound oxygenated intermediates, increasing the catalytic activity for the ORR compared to unstrained Pt surfaces; however, too weak of a surface−adsorbate interaction is also unfavorable. Surface−adsorbate interactions can be optimized by tuning the lattice mismatch in core@shell systems or through alloying surfaces.2,11,27 The lattice mismatch between the cores and shells in these systems was calculated on the basis of nearest neighbor distances (Figure S4). The lattice mismatch between the PdCu B2@PtCu (−4.6%) is higher than that of the PdCu A1@PtCu (−2.1%). Such differences in lattice mismatch are inherent to a study of this nature because the crystal structures of the cores are different by design. Such differences can give rise to different strain profiles and faceting and potentially differences in catalytic activity. However, our electrochemical characterization discussed shortly reveals similar surface properties and activity despite differences in lattice mismatch, allowing the impact of core ordering (i.e., random alloy or intermetallic) on nanocatalyst durability to be elucidated. Electrochemical Evaluation. PtCu surfaces have been shown to be superior catalysts for the ORR compared to Pt references due to ligand effects,30 and when strained appropriately through a core@shell architecture, even higher performing catalysts have been reported.16 Here, the influence of core ordering on durability is evaluated. Both core@shell nanocatalysts were compared to each other and a Pt reference (20% Pt on carbon, Fuel Cell Store). The core@shell NPs were loaded onto a carbon support (EC-600JD, Ketjen; labeled B2@PtCu/C and A1@PtCu/C) and cleaned to remove capping ligands through a washing series. No structural changes were evident, and the particles were well dispersed (Figure S5). In an Ar-purged 0.1 M HClO4 solution, cyclic voltammograms were collected and used to determine the electrochemically active surface areas (ECSAs) based on hydrogen adsorption experiments (Hads, Figure S6). The A1@PtCu/C displayed an ECSA that was 1.4 times larger than the B2@PtCu/C. However, Hads should not be solely used to estimate ECSAs for Pt−M surfaces, so the surfaces were also evaluated by comparing the charge densities from CO stripping experiments and those obtained from Hads (QCO/ QH).31The values obtained for the Pt/C reference are close to 1 as expected from a purely Pt surface. Similar, and expectedly higher, ratios were obtained from the A1@PtCu/C and B2@ PtCu/C samples, showing that they have similar surface adsorption properties consistent with PtM compositions (Figure S7). Ultimately, these results suggest there are similar surface compositions and active surface areas for the two samples, allowing for the core’s influence on durability to be isolated experimentally. The nanocatalysts were then tested for ORR activity in an O2-saturated 0.1 M HClO4 solution. The polarization curves (Figure 3a) show that the half-wave potential (E1/2) for the A1@PtCu/C is shifted 10 mV versus the reversible hydrogen electrode (RHE) more positive than the B2@PtCu/C. Both core@shell nanocatalysts had a more positive E1/2 than the Pt reference by 25 mV and 15 mV for A1@PtCu/C and B2@ PtCu/C, respectively. The catalytic activity was evaluated at D

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is consistent with Pt metal rather than PtOx.35 The increased intensity of the Pt 4f5/2 peak can be attributed to the presence of Cu, as the Cu 3p peaks overlaps with the Pt 4f5/2 peak. There is about a 0.45 eV shift for the Pt 4f for both A1@PtCu/ C and B2@PtCu/C to higher binding energy compared to pure Pt (Table S3 and Figures S8 and S9). We also recorded about a 0.6−0.7 eV downshift of the valence band edge for both core@shell samples (Figure S8c), which can be attributed to both Pt being compressively strained by the core and ligand effects from the bimetallic surface. These results suggest that there is a weaker surface−adsorbate interaction with the core@ shell catalysts compared to bulk Pt, which is advantageous for the ORR.24 Photoelectrons contributing to the Cu 2p and Cu 3p originate from substantially different depths and thus provide different atomic % ratios, with the set of transitions containing the Cu 3p level being more bulklike (within the XPS probe depth limit) and the set of transitions containing the Cu 2p level being more surface sensitive. This analysis reveals that more Pt resides on the surface and more Cu is located in the core, consistent with the core@shell architecture (Table S4). The measured binding energies for Cu 3p also suggests a higher oxidation state for the A1@PtCu sample compared to B2@PtCu sample (Table S3), which may either be an indication of its lower durability (as samples were analyzed at different lengths of time after synthesis) or a contribution to its lower durability. In addition to achieving increased activities with core@shell catalysts, a major challenge with electrocatalysts is metal leaching during extended testing in corrosive environments. There are several examples where intermetallic catalysts outperform and are shown to be more durable than random alloyed nanocatalysts; however, the impact of core crystallinity has not been directly evaluated.16−18,36 Accelerated durability testing (ADT) was performed by cycling the catalyst 5000 times between 60 and 1100 mV (vs RHE) in an O2-purged 0.1 M HClO4 solution. The dashed plots in Figure 3 show the resulting polarization curves (Figure 3d) and the CO-stripping results (Figure 3c). There is clearly a larger E1/2 shift for the A1@PtCu/C (−25 mV) than for the B2@PtCu/C catalysts (−5 mV). Both core@shell catalysts are more durable than the Pt/C reference, with a post ADT activity of the A1@PtCu/C (0.096 mAcm−2 at 0.9 V) dropping off by 58%, whereas there is only an 18% decrease in the B2@PtCu/C specific activity to 0.18 mAcm−2 at 0.9 V (vs RHE) (Figure 3e and f). Evident in the CO-stripping experiments post ADT, there is a negative shift (−15 mV) in the CO oxidation peak for the B2@PtCu/C, indicating weaker surface binding after 5000 cycles. The A1@ PtCu/C catalyst showed a large shift (+20 mV) in the CO oxidation peak, indicating degradation of the catalytically active surface. The negative shift in the CO-stripping experiments after stability testing for the B2@PtCu/C could be due to the formation of a Pt skin on the surface, which would be even more strained than the PtCu alloyed surface that existed initially due to the increased lattice mismatch between the Pt skin and the Cu-rich subsurface.37 Finally, PtCu polyhedra were used as a comparison (Figure S10) and their activity decreased by 92% after 5000 cycles (Figure S11). Taken together, the NPs with core@shell architectures have the highest durability, with the nanocatalyst with the intermetallic core being most durable. The carbon-supported catalysts were characterized by atomic resolution HAADF-STEM as well as STEM-EDX

Figure 3. (a) Polarization curves collected in O2-saturated 0.1 M HClO4 and (b) initial specific activities. (c) CO-stripping curves before and after ADT for core@shell catalysts and (d) polarization curves after ADT compared to the initial curves. (e) Specific and (f) mass activities at 0.85 and 0.9 V before and after ADT.

both 0.9 and 0.85 V (vs RHE) by Koutchey−Levich analysis.13 The initial specific activities (Figure 3b) were very similar for both A1@PtCu/C and B2@PtCu/C catalysts, where at 0.9 V the activities were 0.23 and 0.22 mAcm−2, respectively (Table S2). CO-stripping experiments were conducted to determine the relative d-band positions for the various metal surfaces.32,33 In such experiments, the more negatively shifted the peak potential, the more weakly bound CO is to the surface. The weakening of CO binding indicates a lowering of the d-band and is expected for Pt surfaces under compressive strain. A1@ PtCu/C displayed a more negative peak potential (−10 mV) compared to B2@PtCu/C, with both being shifted more negative than the Pt/C reference (Figure 3c and Figure S7). This finding is consistent with the anticipated compressive strain in these systems. One might also anticipate the A1@ PtCu/C sample to display greater activity than B2@PtCu/C on account of this result; however, we also note asymmetry to the peaks, which may arise from nonuniform strain across the surfaces of the core@shell catalysts.16 This complexity makes a full analysis of activity as a function of strain challenging. Such compressive strain is supported by X-ray photoelectron spectroscopy (XPS), with differences in the electronic structures revealed. Specifically, the Pt 4f peaks were analyzed as corresponding electrons and demonstrated high sensitivity to the surface Pt coverage in the submonolayer regime.8,34 Moreover, the valence band originating from d-electrons is sensitive to surface strain and shifts away from the Fermi level if the lattice strain increases.4 Thus, the Pt 4f5/2 and 4f7/2 peaks were fit with a doublet, and the asymmetric shape of the peaks E

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Figure 4. Comparison of the stability of PdCu A1@PtCu and PdCu B2@PtCu nanocatalysts using simulation. (a) Powder X-ray diffraction profiles obtained by solving the Debye scattering equation. (b) Atom potential energies as a function of the shell number counted from the outside. Potential energies are shown separately for each atom species. (c−h) Distribution of atom potential energies in the models for Cu (c and d), Pd (e and f), and Pt (g and h). We analyze potential energy distributions (c, e, g) and central (111) cut-sections colored according to the atom potential energy (d, f, h). Values for CuPd A1@CuPt and CuPd B2@CuPt are shown side by side. All data are time averaged to reduce noise.

(Figure 4a) and in situ (Figures 1c and 2c) powder XRD profiles. In our comparison, we considered that simulated profiles do not account for peak broadening contributions such as thermal-dynamic disorder, sample size distribution, and inhomogeneity. The characteristic sequence of peaks, their relative integral intensity area, and 2θ peak positions are in reasonable agreement (Figure 4a). In particular, unit cell parameters estimated from the peaks in the 40−50° 2θ range are in agreement with literature data (Table S5), and the interplanar distances of CuPd B2 {110} (0.207 nm) and CuPt A1 {111} (0.218 nm) for the PdCu B2@PtCu model agree with the high-magnification STEM observations (0.212 and 0.218 nm, respectively) from our prior report of the intermetallic system.16 The core@shell nanostructure was also assessed by evaluating the local crystal structure across the NP models (Table S6). Whereas a uniform fcc crystal symmetry spans over the entire CuPd A1@CuPt NP model, a continuous bcc−fcc transition from the core through the surface-shell is unveiled for the CuPd B2@CuPt model despite the uniform bcc structure of the nanostructure as-built before MD equilibration (Figure S15). Elemental composition and structural order of both the nanostructures are not altered by elevated simulation temperature, up to the melting condition (Figure S16). Potential energy was used as a measure of stability. We observed that the average potential energy for the model with an ordered intermetallic CuPd B2 core region is lower by about 0.3% than with a random alloy CuPd A1 core region, irrespective of the equilibration (environmental) temperature (Figure S16a). Interestingly, numerical models that included Pd in the shell showed the same behavior of lower potential

after stability testing in an effort to elucidate any changes in composition and architecture to the nanocatalysts (Figures S12−14 and Table S1). Unfortunately, the small NP sizes and lack of Z-contrast between the core and shell compositions limits the analysis, although a Pt-containing shell is evident by STEM-EDX after ADT for both samples. The STEM-EDX results obtained from NP ensembles before and after ADT for each sample also reveal changes in composition (Table S1). Notably, there is a 50% decrease in the Cu/Pt ratio after ADT for the A1@PtCu/C sample, while only a 10% decrease was measured for the B2@PtCu/C sample. The large decrease in Cu percentage for the A1@PtCu sample suggests significant leaching of Cu from the catalyst’s surface. We note that STEMEDX is the most suitable means of analysis as the low mass loading on the electrodes coupled with the inability to recover most of the used catalyst precludes ICP-MS, SEM-EDX, and related techniques. However, we believe that STEM-EDX is a reasonable means to evaluate changes in elemental composition after stability testing as comparison of the initial STEMEDX values are reasonably close to those obtained by ICP-MS. Investigation of Nanoparticle Stability by Simulation. Although the durability provided by the intermetallic core is evident from the ADT, the physical origin of this durability is not immediately evident. Computational methods are often used to understand the nature of the surface electronics when experimental methods are inadequate to probe these types of phenomena.27 Thus, we used classical molecular dynamics simulations to probe the durability enhancement of PdCu B2@ PtCu compared to PdCu A1@PtCu NPs. Thermodynamic equilibrium models of the two nanostructures were validated against experimental data based on the match of simulated F

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(Cu(ac)2), palladium(II) bromide (PdBr2), and copper acetylacetonate (Cu(acac)2) were purchased from Aldrich and unaltered. Perchloric acid solution (HClO4, 1.0 M) was used as received from Sigma-Aldrich. Oleic acid (90%, OA) and perchloric acid were purchased from Ion Power. Nanopure water (18.2 MΩ·cm) was used for all electrochemical testing. All chemicals were used without further purification. Preparation of PdCu B2 Seeds. The PdCu B2 seeds were synthesized as previously described from the Skrabalak laboratory.16 To a 100 mL round-bottom flask (RBF), 0.100 mmol of Cu(ac)2, 0.100 mmol of PdBr2, and 9.0 mL of OLA and 10 μL OA was added, and this was heated under vacuum to 110 °C and held for 30 min with stirring. Then 50 μL of TOP was added, and the solution was heated to 235 °C under Ar and held at this temperature for 20 min. The solution was allowed to cool to room temperature naturally. A size selection was necessary to obtain size monodispersity. The solution was transferred to a centrifuge tube along with 20 mL of hexane. The solution was spun at 3.9 rpm for 10 min to remove the larger particles. Then the rest of the particles were precipitated out by acetone and collected by centrifugation. The particles were washed three times with a mixture of ethanol and hexane (5:1 vol). The particles were suspending in hexane for further use. Preparation of PdCu A1 Seeds. PdCu alloyed spheres were prepared similarly to what was reported by Gao and co-workers.29 PdBr2 (18.5 mg) and Cu(acac)2 (27.4 mg) were added to a 50 mL RBF with 5 mL of OLA. This was heated to 100 °C under Ar and held for 30 min. The temperature was then increased to 200 °C, and this was held for 30 min. The solution was allowed to cool naturally and collected the same as previously described. The seeds were redispersed in hexane for further use. Preparation of PdCu (A1/B2)@PtCu. The PdCu seeds were added to a 100 mL RBF containing 0.010 mmol of PtBr2, 0.010 mmol of Cu(acac)2, DDD in a 20:1 DDD/Pt ratio, 7.0 mL of ODE, and 2.0 mL of OLA. This was heated to 110 °C under vacuum and held at that temperature for 30 min. The temperature was then increased to 235 °C under Ar and was allowed to incubate for 25 min. The particles were cooled and collected as previously described. Preparation of PtCu/C Reference. PtCu polyhedra were synthesized similarly to what was reported by Xu and co-workers.38 Briefly, 19.8 mg of Pt(acac)2, 13.2 mg of Cu(acac)2, and 101.2 mg of DDD were dissolved in 7.0 mL of ODE and heated to 110 °C under an Ar stream for 15 min. Then a mixture of ODE/OLA (0.308 mL of OLA and 1.29 mL of ODE) and a mixture of DDT/ODE (0.0082 mL of DDT and 0.242 mL of ODE) were injected into the reaction solution. The solution was heated to 230 °C and incubated for 20 min. The solution was then cooled to room temperature and washed as previously described. Characterization. All NPs were characterized on JEOL JEM 1010 TEM, which operated at 80 keV. TEM samples were drop-cast onto the carbon-coated Cu grids after the entire sample was washed two to three times as described below. Additional characterization with STEM-EDX was completed with JEOL JEM 3200FS operating at 300 keV and a JEOL 2200FS operating at 200 keV. The TEM samples were prepared as described above, but carbon-coated Au grids were used. Atomic resolution HAADF STEM imaging was conducted using the aberration corrected Nion UltraSTEM 100 (operating at 100 keV). SEM-EDX was performed with a FEI Qunata 600F Environmental SEM operating at 30 kV with a spot size of 3. SEM-EDX was also used to determine the Pt loading for the working electrode. The ICP-MS analysis was carried out using PerkinElmer SCIEX ELAN DRC-e ICP-MS at the University of Illinois Microanalysis Laboratory. The power diffraction (XRD) was collected on a PANalytical Empyrean instrument with Cu Kα radiation and an X’Celerator linear strip detector. X-ray photoelectron spectroscopy (XPS) measurements were collected using a PHI 5000 Versa Probe II scanning X-ray microprobe under ultrahigh vacuum conditions with a monochromatic Al Kα X-ray source. Electrochemical Evaluation. A Pine rotating disk electrode was used to collect ORR measurements with AfterMath 1.3.7060 software. A typical three-electrode cell was used where the working electrode

energy for the intermetallic system (Figures S17 and S18); this finding is significant because it suggests the enhanced durability is not simply from Pd migration from the core to the shell and may be a more general property of intermetallic core systems. Notably, potential energy of atoms belonging to the NP core region are strongly affected by the order/disorder phase switch. Given that the energies of Cu atoms in the core can be identified based on the lower stability (i.e., higher energy) of those interacting with Pd compared to Pt (Figure 4d), we observe a balancing effect. In the B2 configuration, the potential energy of Cu atoms (Figure 4c,d) is 0.20 eV higher than in the A1, whereas the potential energy is 0.25 eV lower for Pd atoms in the B2 compared to the A1 (Figure 4e, f). Potential energy fluctuations narrow significantly because of local compositional order. Interestingly, the potential energy of Cu atoms in the surface shell is insensitive to the core phase structure, while the potential energy of Pt atoms slightly decreases with the A1-to-B2 phase switch (Figure 4g,h). Pt atoms at the core−shell interface have lower potential energy as they interact with the ordered B2 phase than with the disordered A1 phase. Furthermore, the increase of the potential energy of Pt atoms observed with Pd impurities in the surface-shell (Figure S18f, l) suggests that ordering of Cu atoms in the interface improves the binding between the core and the surface shell. The potential energy of Cu atoms in the CuPd B2@CuPt and CuPd A1@CuPt NP models become indistinguishable in the interface region where the intermetallic order is progressively lost (Figure 4b). Pd and Pt atoms have otherwise a lower potential energy for B2 than A1 core phase structure over some of the atom layers closest to the core region. These results demonstrate that the core directly affects the stability of the shell. Overall, our molecular dynamics simulation supports that the CuPd B2@CuPt configuration has to be expected to be more durable than the CuPd A1@CuPt.

CONCLUSIONS In summary, we directly compared core@shell catalysts for the ORR where the shape, size, and composition were constant and only the crystal structure of the core was altered. Both core@shell catalysts showed enhanced catalysis when compared to a Pt/C reference. Initially, the two core@shell catalysts showed similar activity, which can be attributed to lattice mismatch between the core and shell as well as ligand effects that can arise from alloyed surfaces. After ADT, the B2@PtCu/C catalyst maintained its activity to a much greater extent than the A1@PtCu/C catalyst. These results show that the durability of a core@shell catalyst is not only dependent on the surface structure but on the properties of the core. Atomistic simulations were used to help understand the durability enhancements as experimental tools do not exist to probe this phenomena. The simulation results support the idea that the alloyed surface is more stable on the intermetallic core than on the random alloyed core. This work systematically examined the influence of core structure on catalytic durability. We envision that this idea can be applied to other core@shell systems to develop more durable catalysts for the ORR and other electrocatalytic processes. METHODS Materials. Oleylamine (70%, OLA), trioctylphosphine (97%, TOP), 1,2-dodecanediol (90%, DDD), 1-octadecene (90%,ODE), dodecanthiol (DDT), platinum bromide (PdBr2), copper(II) acetate G

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ACS Nano was glassy carbon (diameter 5.5 mm, area 0.238 cm2), the reference electrode was Ag/AgCl (4 M KCl), and a Pt coil was used as the counter electrode. All NPs were even dispersed on a carbon support (Ketjen EC-600JD). The NPs were massed and then resuspended in 1 mL of hexane. The NP were then added dropwise to a slurry of the carbon support in 5 mL of acetone and 2.5 mL of hexane. The mass ratio between the NPs and carbon support was 1:1. Once the NPs were added, the solution was sonicated for 1 h and then allowed to stir overnight. The solvents were then evaporated, and the resulting catalyst was dried in vacuo overnight. The working electrode was prepared by creating a 2 mg mL−1 ink solution of the catalysts, where the solution consisted of 1:3.98:0.020 volumetric ratio of 2-propanol/ H2O/Nafion. The catalyst ink was sonicated for 30 min and 10 μL was dropcast onto the glassy carbon electrode and dried. The Pt loading was determined from the metal ratios from SEM-EDX in combination with the catalyst preparation method in which a 1:1 mass ratio of carbon support-to-nanocatalyst was used. Normalization of the electrode by sweeping 200 times between 60 mV and 1200 mV vs RHE was performed first to remove any excess surfactant or surface contamination. Hydrogen desorption experiment were used to determine the ESCA and were performed by cycling between 60 and 1000 mV vs RHE in an Ar purged 0.1 M HClO4 solution. The area under the hydrogen desorption peaks were measured and using a conversion factor of 210 μC cm−2 was converted to the active surface area. The value obtained was then normalized by the Pt loading on the electrode. Linear sweep voltammetry (LSV) was performed in an O2saturated solution of 0.1 M HClO4 between 60 to 1200 mV (vs RHE) at a scan rate of 10 mV/s at 1600 rpm. The kinetic current was obtained by collecting LSV curves at several different rotation speeds (400, 900, 1600, 2025, 2500 rpm). The specific (normalized by surface area) and mass (normalized by Pt loading) were determined at 850 and 900 mV (vs RHE). CO-stripping experiments were conducted by purging a 0.1 M HClO4 with CO gas while holding the open circuit potential for 30 min. After the potential hold, the solution was purged with Ar for 30 min to removed excess CO gas. The CO-stripping curves were obtained by cycling between 60 and 1200 mV (vs RHE) at a scan rate of 50 mV/s. Atomistic Simulation. Numerical models of nanocatalysts were built by carving cuboctahedral nanocrystals out of either an ideal fcc lattice (lattice constant a0 = 0.377 nm) or an ideal bcc lattice (a0 = 0.300 nm), according to the crystal structure in the core region. A family of {h00} planes was inserted after every six planes in the surface shell of the CuPd B2@CuPt model to mimic the lattice mismatch described by Wang et al.16 Some atoms in the surface shell ended up less than 0.20 nm apart from core atoms. To improve the model, we randomly selected about one-third of these atoms and removed them. After the geometric construction, atomic types were assigned according to the experimental chemical composition of the NPs. We also tested the effect of Pd impurities in the surface shell by comparing CuPd@CuPt and CuPd@CuPdPt2 element compositions. Such impurities might appear due to segregation from the core to the shell or via Ostwald ripening phenomena. As-built configurations were equilibrated at room temperature (300 K) with classical molecular dynamics simulations using the LAMMPS software package.39 Atom-pair interactions were computed with longrange Finnis−Sinclair potentials40,41 based on the quantum Sutton− Chen many-body force field.42 No constraints to the mobility of atoms were applied to impose the expected elemental composition across the nanocrystals. After energy minimization, the system was equilibrated for 1 ns in subsequent NVE and NVT ensemble runs using a Nosè−Hoover thermostat with a 1 fs time step. The Langevin dynamic was employed in the initial stage to smoothen the evolution of the system. After equilibration, NVE ensemble simulations recorded trajectories of 500 independent configurations sampled at 2 ps interval time. Time-average (TA) microstructures were obtained by averaging trajectory configurations to cancel the thermal atom vibrations out of the crystalline lattice.43

Powder XRD profiles were simulated for the TA microstructures by solving the Debye scattering equation.44 The profiles were plotted in 2θ space assuming Cu Kα radiation. The local crystal structure was evaluated using the polyhedral template matching method.45 In agreement with characterization of powder XRD data, only bcc and fcc lattice symmetries were considered.

ASSOCIATED CONTENT S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.8b08007. Additional characterization, electrochemical evaluation, and simulation information (PDF)

AUTHOR INFORMATION Corresponding Author

*E-mail: [email protected]. ORCID

Raymond R. Unocic: 0000-0002-1777-8228 Michael Engel: 0000-0002-7031-3825 Sara E. Skrabalak: 0000-0002-1873-100X Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS S.E.S, J.T.L.G., N.N.D., Y.B.L, and H.M.A. acknowledge financial support from Indiana University and U.S. DOE BES Award DE-SC0018961. Access to the powder diffractometer was provided by NSF CHE CRIF 1048613 and to the XPS by NSF DMR MRI 1126394. We also thank the IU Electron Microscopy Center and Nanoscale Characterization Facility for access to the necessary instrumentation. A.L. and M.E. acknowledge funding from Deutscht Foschungsgemeinschaft through the Cluster of Excellence Engineering of Advanced Materials (EXC 315/2). Supercomputer time for this research was supported in part by the Lilly Endowment, Inc., through its support for the Indiana University Pervasive Technology Institute, and in part by the Indiana META Cyt Initiative. The Indiana META Cyt Initiative at IU was also supported in part by the Lilly Endowment, Inc. A portion of the electron microscopy characterization was conducted as part of a user proposal at Oak Ridge National Laboratory’s Center for Nanophase Materials Sciences, a U.S. Department of Energy Office of Science User Facility. REFERENCES (1) Sui, S.; Wang, X.; Zhou, X.; Su, Y.; Riffat, S.; Liu, C. A Comprehensive Review of Pt Electrocatalysts for the Oxygen Reduction Reaction: Nanostructure, Activity, Mechanism and Carbon Support in PEM Fuel Cells. J. Mater. Chem. A 2017, 5, 1808−1825. (2) Stamenkovic, V. R.; Mun, B. S.; Arenz, M.; Mayrhofer, K. J. J.; Lucas, C. A.; Wang, G.; Ross, P. N.; Markovic, N. M. Trends in Electrocatalysis on Extended and Nanoscale Pt-Bimetallic Alloy Surfaces. Nat. Mater. 2007, 6, 241−247. (3) Greeley, J.; Stephens, I. E. L.; Bondarenko, A. S.; Johansson, T. P.; Hansen, H. A.; Jaramillo, T. F.; Rossmeisl, J.; Chorkendorff, I.; Nørskov, J. K. Alloys of Platinum and Early Transition Metals as Oxygen Reduction Electrocatalysts. Nat. Chem. 2009, 1, 552−556. (4) Stephens, L. I. E.; Bondarenko, S. A.; Grønbjerg, U.; Rossmeisl, J.; Chorkendorff, I. Understanding the Electrocatalysis of Oxygen Reduction on Platinum and Its Alloys. Energy Environ. Sci. 2012, 5, 6744−6762. H

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