Chemical Ordering and Surface Segregation in CuâPt Nanoalloys

Article pubs.acs.org/JPCC

Chemical Ordering and Surface Segregation in Cu−Pt Nanoalloys: The Synergetic Roles in the Formation of Multishell Structures Jianfeng Tang,†,‡ Lei Deng,*,‡ Shifang Xiao,§ Huiqiu Deng,§ Xingming Zhang,‡ and Wangyu Hu*,† †

College of Materials Science and Engineering and §School of Physics and Electronics, Hunan University, Changsha 410082, China ‡ Department of Applied Physics, Hunan Agricultural University, Changsha 410128, China

Downloaded by KTH ROYAL INST OF TECHNOLOGY on September 1, 2015 | http://pubs.acs.org Publication Date (Web): September 1, 2015 | doi: 10.1021/acs.jpcc.5b06145

S Supporting Information *

ABSTRACT: We performed Monte Carlo simulations coupled with MAEAM potentials to study the surface segregation and chemical ordering patterns in Cu−Pt nanoalloy particles for a broad range of sizes, shapes, composition, and temperature. It was found that both the Cu segregation on the surface and the chemical ordering in the core are the general rules and usually compete with each other. Surface segregation of Cu is enhanced with increasing particle size or surface openness or global Cu composition. Despite their different morphologies, most of the types of ordered phases in the core region are the same as bulk alloys. Due to the modification or suppression effects of surface segregation, the degrees of chemical ordering shift to the Ptricher side and are more apparent in a large-sized particle. Particularly, at a narrow composition range, the multishell structures (onion-ring or multishell/maze-like core) form in (truncated) octahedrons, illustrating a subtle synergy between the segregated Cu {111} facets and the L11 ordering. The possibility and advantage of transformation from these multishell structures to Pt multilayer shell/single Cu-rich core ORR catalysts by selective etching of Cu were also discussed.

1. INTRODUCTION Pt-based bimetallic catalysts usually show physicochemical properties that are different from those of their parent metals, offering the opportunity to design new heterogeneous nanocatalysts with enhanced selectivity, activity, and stability.1−7 Ptbased bimetallic nanoparticles (BNPs) or nanoalloys (NAs) can be consequently tuned and engineered to meet the desired electronic (ligand) effect and geometric (ensemble or strain) effect and thus the catalytic performance by varying elemental compositions, internal structures, and particle morphology.3,5−7 Extensive experimental and theoretical works have suggested that Pt-based NA catalysts with a variety of structures can effectively improve both the catalytic performance and utilization of Pt.1−7 Even so, insufficient knowledge about the surface structure and composition of complex structured NAs hinders the development of an atomistic scenario of the structure−performance relationship. Moreover, even if NAs with complex initial structures can be fabricated by designed growth or etching strategies, they will always transform to the thermodynamically preferred configurations sooner or later. Thermodynamically, a steady structure of NAs results from the balance between bulk alloying and surface segregation. The latter is thermodynamically driven by reduction in the total free energy, which can be largely attributed to a difference in surface energy between the two components. 3,6,7 Due to the anisotropic nature of shape and large specific surface areas of NAs, surface segregation consequently becomes a critical issue © XXXX American Chemical Society

when designing NA catalysts. As a result of the obvious competition between surface segregation and bulk alloying, confined NA catalysts will exhibit complicated and novel structures, such as the core−shell, onion-ring, patchy multishell, and Janus ones, just to name a few.6,7 To engineer Pt-based nanocatalysts at the atomic level, it is of particular importance to systematically study the steady structures of NAs and understand the interplay between surface segregation and bulk alloying as well as their control factors. Among Pt-based bimetallic catalysts, Cu−Pt bulk alloys or NAs have attracted increasing attention in recent years, owing to their higher activity than monometallic catalysts in a variety of energy-related catalytic processes, such as water−gas shift,8,9 reduction of NO,10 O2,11−26 and CO2,27−30 hydrodechlorination of 1, 2-dichloroethane,31−33 oxidation of CO,34−38 methanol,39−43 and formic acid.44−48 Previous studies suggested a close relationship between the catalytic performance and the structure of Cu−Pt alloys or NAs, in terms of the electronic, geometric, or synergistic effects. For instance, Cu− Pt (111) near surface alloys (NSAs) with Pt-skin and Cu-rich subsurface or NSA-like Cu-core/Pt-shell NAs have been found to exhibit superior activities toward water−gas shift8 and oxygen reduction reaction (ORR),11−26 as well as significantly Received: June 26, 2015 Revised: August 16, 2015

A

DOI: 10.1021/acs.jpcc.5b06145 J. Phys. Chem. C XXXX, XXX, XXX−XXX

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The Journal of Physical Chemistry C greater selectivity for N2 formation during NO reduction10 than that of pure Pt catalysts. Strasser and co-workers demonstrated that the superior ORR activity of dealloyed Cu−Pt catalysts were largely attributed to the lateral strain imposed to the thick Pt shell by the Cu-rich core, which weakens the binding of the Pt surface atoms to oxygenated species.13−17 On the contrary, equivalent role of ligand effects has been demonstrated, where Cu−Pt (111) NSAs exhibited 8-fold enhancement in ORR activities over Pt (111).11,12 The Cu−Pt NAs have also been found to exhibit superior activity toward methanol and formic acid electro-oxidation, attributed to the ligand effects and a more accessible surface area.39−48 These NAs were in a rich variety of compositions and morphologies, including Cu2Pt3 nanocube or nanooctahedron, Cu3Pt nanocage, CuPt3 nanodendrite or nanoframe, and so on. The energy-dispersive X-ray elemental mapping demonstrated that both Pt and Cu were distributed evenly from surface to interior; however, the structure−activity relationship yet remains largely unexplored due to insufficient knowledge about surface and near-surface structure of catalysts with such a broad distribution of size, composition, and morphology. Experimental research on the structure and chemistry of NAs is limited by the poor availability of catalysts with a homogeneous distribution of size, morphology, and composition. Moreover, the analysis of the local structure near the surface of NAs is complicated and is limited to only a few surface sensitive characterization techniques.1,21 In contrast, force-field methods and first-principles methods based on density functional theory (DFT) can provide atomic level insights into the structure and chemistry of NAs and the controlling factors.6 To determine the structure of NA, however, a direct DFT calculation is very computationally demanding and limited to clusters with a few dozens of atoms. Recently, DFT combined with lattice models in the framework of cluster expansion method were developed to deal with the case of large-sized NAs.49−53 Alternatively, pair or many-body potentials coupled with genetic algorithms or Monte Carlo (MC) algorithms were commonly used to investigate the energetically favorable atomic arrangements in large-sized nanosystems.6,54−58 Despite recent advances in force-field studies on the Cu−Pt NAs, the surface segregation and chemical ordering patterns of Cu−Pt NAs and the effects of sizes, shapes, compositions, and temperatures have yet to be comprehensively investigated.59−62 In this perspective, we performed MC simulations with the modified analytical embedded atom method (MAEAM) potential63−69 to unravel the roles of surface segregation and chemical ordering in determining the thermodynamically stable structure of Cu−Pt NAs. Parameters such as alloy composition, particle shape and size, and system temperature were correlated with surface Cu composition and the chemical ordering. It was found that the Cu segregation on the surface and the chemical ordering in the core are the general rules and usually compete with each other. At a narrow composition range, the multishell structures form in (truncated) octahedrons, illustrating a subtle synergy between the segregated Cu {111} facets and the L11 ordering. The possibility and advantage of transformation from these multishell structures to Pt multilayer shell/single Cu-rich core ORR catalysts by selective etching of Cu were also discussed.

2. MODEL AND METHOD 2.1. NA Morphology. Metallic NPs usually display a variety of competitive morphologies, such as decahedron, icosahedron, and octahedron (OCT).6 The idealized structure motifs considered here include FCC structured OCT and its variants, cube-octahedron (CUB), and truncated octahedron (TOC). The OCT morphology is enclosed by 8 {111} facets, 12 {111}/ {111} edges, and 6 vertices. The TOC is a moderately truncated morphology of OCT and is enclosed by 6 square {100} facets, 8 hexagonal {111} facets, 12 {111}/{111} edges, 24 {111}/{100} edges, and 24 vertices. The heavily truncated CUB is enclosed by 6 square {100} facets, 8 triangular {111} facets, 24 {111}/{100} edges, and 12 vertices. We constructed a series of NAs with N atoms (4 × 102 ∼ 1 × 104) according to the magic numbers.6 Initial atomic configurations of Cu−Pt NAs were generated by randomly designating the proper mole fractions (0.125, 0.25, 0.375, 0.5, 0.625, 0.75, 0.875) as Cu atoms and the remaining as Pt atoms. The NAs’ equilibrated configurations were obtained at 10, 300, 600, and 900 K. 2.2. MAEAM Potential. The reality and accuracy of the simulation results principally depends on proper descriptions of the interatomic potentials. In the present work, MAEAM potential was adopted to describe atomic bonding in Cu−Pt NAs because of their underpinnings in quantum mechanics, mathematic simplicity, and extendibility to alloy materials. The original EAM potential,58,70 the modified one (MEAM),56,71,72 and our MAEAM one63−69 have already been successfully applied for the study of bulk, surface, and nanoparticles of metals and alloys through incorporating many-body interactions. In short, for a monometallic system, the original EAM represents the energy per atom as a sum of the pair interaction energy and the embedding energy.70 The embedding energy incorporates many-body interactions and is a nonlinear function of the host/background electron density. In the EAM, the host electron density was usually taken to be a summation of the spherical atomic electron densities. While in original MEAM, an angularly dependent term was added to include the possible influences of energy caused by the deviations of the atomic electron density from the spherical symmetry coming from each first nearest-neighbor (1NN) atom.71 The MEAM was modified once again by Lee and Baskes to partially consider second nearest-neighbor (2NN) interactions as well as the 1NN ones, overcoming some critical shortcomings in the original MEAM.72 To avoid the computational complexity of the calculation of angle factors, in our MAEAM, an extra energy-modified term was introduced to resolve the negative Cauchy pressure problem in Johnson’s model,73 and the argument of energy-modified term was presented as the sum of high orders of electron density in order to correct the discrepancy of the linear superposition of atomic electron density at the same time. For pure elements, the equilibrium physical properties fitted were the cohesive energy, unrelaxed vacancy formation energy, second-order elastic constants, and the lattice constants. The potential functions of pure elements can be found in ref 69. The final optimized parameters for Cu and Pt was given in ref 69 and ref 64, respectively. For binary alloys, the unlike-atom pair potential was determined by assuming a geometric mean between the likeatom potentials. The cross-pair potential functions of alloys also can be found in ref 69. The dilute heats of solution or formation heats of intermetallics were used to determine the B


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The Journal of Physical Chemistry C Table 1. Surface Energy γ (in mJ/m2) of Low-Index Extended Surfaces for Cu and Pt

alloying parameters and thus to better reproduce the experimental ordered phases. For Cu−Pt alloys, the alloying parameters rc and μ were taken as 1.26 and 2.75. 2.3. MC Simulation. In nature, the prediction of energetically stable structures is regarded as a highly complex problem of global optimization. To obtain the equilibrium structures of the Cu−Pt NAs, off-lattice MC simulations were conducted.56,65−67 In our MC simulations, the NPT ensemble was used where the pressure, temperature, and total number of Cu or Pt atoms were kept constant.56 The random mixed initial configuration evolved via a series of configuration relaxations to achieve thermodynamic equilibrium. One of the following three types of trial was used for each MC step that corresponds to N configuration relaxations. The first one is the positional relaxation where a randomly selected atom displaces in an arbitrary direction with a fixed distance. The second is the compositional relaxation where two randomly selected atoms exchange their positions. The last is the global volume relaxation where box change along one out of the three orthogonal directions with a fixed distance, accounting for zero pressure. The attempts are accepted or rejected according to the standard Metropolis criterion.56,69 That is, a trial is accepted if the change in the configuration energy is negative; otherwise, it is accepted only if the appropriate Boltzmann factor is greater than a random number selected from a uniform distribution. During the equilibrium stage (depends on temperatures, typically after 20 000 MC steps), detailed balance is observed where there are relatively small fluctuations in total energy and the acceptance probabilities of different attempts. Then, the physical quantities of interest are obtained by averaging over 10 000 configurations.

Cu

Pt

γ(111)

γ(100)

γ(110)

1407 1952a 1170b 1409c 1266d 1730e 1190f 1180g 1387h 1560 2299a 1440b 1656c 1580d 2510e 1080f 1120g 1694h



a

DFT (ref 74). bEAM (ref 70). cMEAM (ref 71). dVC-EAM (ref 61). NRL-TB (ref 76). fTB-SMA (ref 75). gLR-EAM (ref 77). hEAM optimized by force-matching method (ref 78). e

small nanoclusters.6 Herein, we sacrificed some of the accuracy of surface energy to focus on better overall performances during fitting the current potential models. It has been confirmed that this slight deviation in surface energy values do not prevent the reasonable prediction of surface reconstruction, relaxation, and segregation of extended surfaces.70,75−78 The alloying parameters of cross-pair potential account for the physical properties of the Cu−Pt bulk alloys, particularly for the relative stability of possible ordered phases. Figure 1 shows the atomic structures and the calculated heats of formation ΔH for homogeneous solid solutions and D1/D7, L10, L11, L12, L13, and D022 long-range ordering (LRO) phases. The LRO phase structures are given in Strukturbericht designation: D7 for CuPt7, L12 or D022 or L13 for CuPt3 and Cu3Pt, L10 or L11 for CuPt, and D1 for Cu 7 Pt. For comparison, available experimental data and calculation results obtained from the generalized gradient approximation (GGA), the augmented spherical wave (ASW) method, the linearized augmented plane wave (LAPW) method, EAM, MEAM, and Bozzolo−Ferrante− Smith (BFS) potential were also shown.81−87 The phase diagram of Cu−Pt system is relatively simple and exhibits several stable intermetallics at low temperatures: L12-Cu3Pt, L11-CuPt, and L13-CuPt3 phases.88−91 Furthermore, it was theoretically suggested that D1 and D7 at 1:7 stoichiometry are the possible ground states at low temperature, which may be overlooked experimentally as the low order−disorder transition temperatures.82,90,91 The ground states L11-CuPt, L13-CuPt3, and D7-CuPt7 phases of bulk alloys are well reproduced by the present Cu−Pt potentials, whereas the stabilities of D1-Cu7Pt and L12-Cu3Pt at the Cu-rich side are underestimated. It is notable that none of the many-body potentials exactly reproduces the whole convex hull of the heats of formation provided by DFT calculation. Among these ground states, L11 and L13 LRO phases have only been observed experimentally in the Cu−Pt system. L11-CuPt consists of alternative fcc (111) layers of Pt and Cu, in contrast with the metastable L10

3. RESULTS AND DISCUSSION 3.1. Validation of the MAEAM Potential. It is understood that surface energy is an important factor determining the structural features of NAs because of their high surface-tovolume ratios. The surface energies of fully relaxed low-index surfaces were listed in Table 1. For comparison, the results obtained from DFT74 and the many-body potentials, including the tight-binding model within the second-moment approximation (TB-SMA)75 or the Naval Research Laboratory version (NRL-TB),76 as well as the original version,70,78 the modified one (MEAM),71 the Voter−Chen one (VC-EAM),61 or the long-range one (LR-EAM)77 of EAM potential were also listed. Our calculations show the correct order (γ(111) < γ(100) < γ(110)) and similar anisotropy ratios (γ(100)/γ(111) and γ(110)/γ(111)) of surface energies of the three low-index surfaces in comparison with DFT and many-body potentials. The smaller surface energy values of Cu relative to Pt have been also correctly predicted by present potential and other many-body ones, except for TB-SMA and LR-EAM ones. The average surface energy is about 1800 mJ/m2 for Cu and 2500 mJ/m2 for Pt from experiments.79,80 Although accurate experimental values of the surface energy are needed, it is found that most of the many-body potentials still more or less underestimate the surface energies, whereas the NRL-TB overestimates surface energies of Pt. One common problem with the many-body potentials is that they are not satisfactorily transferable from isolated atoms to bulk via surface, which means fitting the parameters to well reproduce the bulk properties often leads to an underestimate of the surface energy and vice versa. Although the DFT methods can reasonably reproduce both surface and bulk properties of transition metals, they are limited to very C


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The Journal of Physical Chemistry C

Figure 2. Surface configurations of (a) OCT8119, (b) TOC9201, and (c) CUB8217 NAs with different Cu:Pt ratios at 100 K. Visualization of structure was carried out using OVITO, where Cu atoms are shown in red and Pt in blue.

Table 1, which provides large driving forces for segregation. Because the lattice constant of bulk Cu is smaller than Pt (3.60 Å vs 3.91 Å in MAEAM), it is reasonable and necessary to consider the strain effects. Surface sites are usually tensile and prefer the large Pt atoms. The contribution of the strain energies are less that 0.08 eV/atom (depending on alloy composition). Therefore, the decrease in surface energy by Cu segregation (0.15 to 0.28 J/m2 or 0.30 to 0.68 eV/atom depending on facets) overwhelms the increase in strain and chemical energy because of the high surface-to-volume ratio of NAs. However, it should be noted that the segregation of Cu can be suppressed/promoted at severely tensile/compressive sites. When the global Cu composition increases, Cu atoms preferentially occupy the vertex, edge, {100} facet, and {111} facet sites, opposite to Pt. This preferential segregation of Cu or Pt can be largely attributed to the different cohesive energies (in negative values) of nonequivalent surface sites, which has been found decreasing linearly with increasing of the coordination numbers (CNs).59,65−67,93 Therefore, the change in cohesive energy for exchanging a Cu atom with a Pt atom will prefer Cu atoms segregating onto the surface sites with lower CNs. The mean Cu concentrations in the topmost shell of Pt-rich NAs with different compositions and shapes are shown in Figure 3 as a function of particle sizes. We used CNs to distinguish the surface and the core of NAs. The available results of MC simulation with surface modified pair-potential were also included for comparison.59 Generally, surface Cu composition increases with global Cu composition or particle size. The key factor driving the composition-dependent pattern of surface segregation is the limited supply of Cu.65−67,93 For a fixed particle size, the surface sites available for Cu segregation are also fixed, and then an increase in the global Cu composition can promote the surface segregation. However, it is more difficult to interpret the size-dependent pattern of surface segregation. From the viewpoint of preferential segregation, Cu occupation proportion is in order of the following: vertex > edge > {100} facet > {111} facet sites.65−67 As surface sites are fully (vertices and edges) or almost ({100} facets) covered by Cu, the increase of the Cu surface concentration with the particle sizes can be mainly

Figure 1. (a) Atomic structures and (b) heats of formation ΔH for disordered solid solutions and D1/D7, L10, L11, L12, L13, and D022 ordered phases. The results of many-body potentials calculations and DFT calculations as well as experimental data can be found in refs 81− 87.

structure that has alternating (001) planes of atoms.82,91,92 L13CuPt3 has been recently confirmed by experiments90 and DFT calculation,82 which is similar to L11 structure but has alternative (111) layers of pure Pt and half Cu and Pt. It is noteworthy that the present potential correctly describes the relative stabilities of L11 to L10 for CuPt and L13 to L12 or D022 for CuPt3, alternative layers with pure or mixing Cu/Pt along directions are expected to dominate the structural features of the NAs as will be discussed later. 3.2. Surface Segregation of Cu. From the perspective of thermodynamics, the atomic arrangement of the NAs is a result of the competition and balance between surface segregation and chemical ordering, and thus, this arrangement will be highly sensitive to their compositions, particle size or shape, and operating environments.6,7,93 We first looked at the overall configurations of the lowest-energy Cu−Pt NAs. Figure 2 shows the surface configurations of OCT8119, TOC9201, and CUB8217 NAs at 100 K. Visualization of structure was carried out using OVITO,94 where red atoms are Cu and blue ones are Pt. When global Cu compositions are higher than 0.5, the surfaces of large-sized particles are fully covered by Cu and thus not presented here for clarity. Surface enrichment of Cu has been found in Cu−Pt NAs under an inert or reducing atmosphere.15,20−22 Present results show that surface (preferential) segregation of Cu is the rule more than the exception in Cu−Pt NAs. This surface segregation of Cu is not surprising as the surface energy of Cu is much less than that of Pt as listed in D


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facet sites available for Cu segregation, instead a small amount of Pt atoms occupy the edge sites near the vertices. From the viewpoint of preferential segregation, the {111} facets have the lowest surface energy and thus are the least favorable surface sites for Cu segregation. As a result, the morphology-dependent patterns of surface segregation mainly result from the ratios of the {111} facets to the whole surface sites for considered morphologies, as shown in Figure 4. It is seen that the contribution of the {111} facets follows the order OCT > TOC > CUB for NAs with comparable sizes. Therefore, as shown in Figure 3, surface enrichment of Cu is more apparent for Cu−Pt NAs with more open surfaces. 3.3. Chemical Ordering in Core Region. As the competition or synergy between surface segregation and bulk ordering, it is reasonable to anticipate that NAs can exhibit more complicated structures than bulk systems. As an example, the midmost cross sections of energetically favorable OCT8119, TOC9201, and CUB8217 NAs at 100 K are shown in Figure 5. These cross sections were obtained by cutting the atomic configuration shown in Figure 2 by along Z directions and thus exhibit {001} facets. The atomic intermixing in the core of Cu− Pt NAs depends strongly on the alloy compositions. On the basis of the surface configurations illustrated in Figure 2 with the cross sections shown in Figure 5, it is seen that equimolar and Cu-rich NAs exhibit mixed-cores and completed Cu shells, whereas Pt-rich ones have shells where vertices, edges, and facets are preferentially occupied by Cu. At this low temperature, it is interesting to note that Cu and Pt atoms in the cores of NAs with 0.375 to 0.750 Cu compositions show obviously chemical ordering along the directions. The choice of proper parameters to yield the degree of chemical ordering will be particularly important when dealing with the NAs. The overall mixing index Ω is a useful measurement of the degree of tendency for ordering or clustering in bulk materials and NAs:95

Figure 3. Mean Cu concentrations in the topmost shell of Pt-rich NAs with different compositions and shapes as a function of particle sizes.

attributed to {111} facets. The increase of the Cu concentration in {111} facets with the particle sizes can be attributed to the competing nature between segregation and ordering. The surface-to-volume ratio decreases quickly with increasing particle sizes (from about 50% for 500 atoms to 20% for 8000 atoms), which means that the surface Cu concentration would increase greatly with particle sizes. It is not true because the strong surface segregation of Cu can be suppressed by the chemical ordering in core (see Figure 1b). Apparently, increasing the particle size also promotes the ability to accommodate the surface segregation behavior.67,93 Interestingly, the surface segregation patterns of CuPt7 NAs show anomalous dependences on the particle sizes, as shown in Figure 3. When the size increases, the surface Cu compositions increase for OCT and are almost constant for CUB and TOC. Understanding this anomaly requires a look back at the preferential Cu segregation and the morphology features in NAs. Revisiting Figure 4, when the particle sizes increase from

Ω=

CNCu − Cu + CNPt − Pt − 2CNCu − Pt CNCu − Cu + CNPt − Pt + 2CNCu − Pt

(1)

where CNCu−Cu, CNPt−Pt, and CNCu−Pt (or equivalent CNPt−Cu) are the number of Cu−Cu, Pt−Pt, and Cu−Pt bonds, which are sum over all of alike or unlike the 1NNs (or CNs) of each Cu or Pt atom. On the basis of this definition, a fully phase separated bulk system would have Ω = 1, whereas a fully mixed phase would have Ω = −1. This index is actually equivalent to the Warren−Cowley chemical short-range order (SRO) parameter αl for the first shell of neighbors.96 However, these two quantities make it somewhat difficult to identify possible ordered structures in NAs. Herein, we extended the overall mixing index Ω to the local mixing index σ for Cu and Pt atoms: cn − cn Cu − Pt σCu = Cu − Cu cn Cu − Cu + cn Cu − Pt (2) Figure 4. Ratios of different surface sites to the whole surface sites for considered morphologies.

σPt =

cnPt − Pt − cnPt − Cu cnPt − Pt + cnPt − Cu

(3)

where cnCu−Cu, cnPt−Pt, cnCu−Pt, and cnPt−Cu are the number of alike or unlike the 1NNs (or CNs) for each Cu or Pt atom. The relationship between local mixing index σi,j and overall mixing index Ω is

1 000 to 10 000 atoms, the percentages of vertices and edges decrease from 35% to 18% for CUB and TOC but 25% to 12% for OCT. Therefore, all of vertices and edges and even partial {100} facets in CUB and TOC CuPt7 NAs are occupied by segregated Cu atoms when the particle sizes increase. For OCT CuPt7 NAs, there is not enough vertex and edge and no {100}

Ω= E

1 N

∑ σi ,j × Ni ,j

(4) DOI: 10.1021/acs.jpcc.5b06145 J. Phys. Chem. C XXXX, XXX, XXX−XXX

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Figure 5. Midmost cross sections of (a) OCT8119, (b) TOC9201, and (c) CUB8217 NAs with different Cu:Pt ratios at 100 K. Cu atoms are shown in red and Pt in blue.

Table 2. Mixing Index σi,j and Ω for Typical Ordered Phases in Bulk Alloys ordered phase

σCu (XCu)

σPt (XPt)

Ω

α1

L10-CuPt L11-CuPt L12-Cu3Pt L12-CuPt3 L13-Cu3Pta L13-CuPt3 D1-Cu7Ptb D7-CuPt7

−1/3 (1/2) 0 (1/2) 1/3 (3/4) −1 (1/4) 1/3 (1/2) or 2/3 (1/4) −2/3 (1/4) 2/3 (3/4) or 1 (1/8) −1 (1/8)

−1/3 (1/2) 0 (1/2) −1 (1/4) 1/3 (3/4) −2/3 (1/4) 1/3 (1/2) or 2/3 (1/4) −1 (1/8) 2/3 (3/4) or 1 (1/8)

−1/3 0 0 0 1/6 1/6 1/2 1/2

−1/3 0 −1/3 −1/3 −1/9 −1/9 −1/7 −1/7

a For L13-Cu3Pt, σCu can be 2/3 (for Cu with two Pt and ten Cu 1NNs) or 1/3 (for Cu with four Pt and eight Cu 1NNs). bFor D1-Cu7Pt, σCu can be 2/3 (for Cu with two Pt and ten Cu 1NNs) or 1 (for Cu with 12 Cu 1NNs).

Figure 6. Midmost cross sections of (a) OCT8119, (b) TOC9201, and (c) CUB8217 NAs with different Cu:Pt ratios at 100 K. The atoms are indexed to the local mixing index σ as shown in the gradient color scheme.

where Ni,j is the numbers of Cu or Pt atoms at the equivalent symmetry site i with the CNs j. For ordered alloys, such as L13 and D1, Cu or Pt atoms can occupy different symmetry sites. For the NAs, Cu or Pt atoms on the surface may have various CNs. The mixing index σi,j and Ω for typical ordered phases in bulk alloys were summarized in Table 2. The Warren−Cowley SRO parameter of the first shell α1 was also listed for comparison. It is seen that αl are 1 for even shells of neighbors and −1/3 for odd shells for both perfect L10 and L12 LRO phases, zero for each shell for both perfect L11 and disordered solid solution; however, the overall mixing index Ω is zero for perfect L11, L12, and disordered solid solution. Only the local mixing index σ can well distinguish these ordered phases from each other. It should be noted that the present local mixing indexes are unavailable to distinguish L12 phase from

corresponding long period structures (such as D022 and D023) and then may overestimate somewhat the amount of L12 phase. An accurate characterization of the long period structures can be obtained by considering at least the 2NNs. Furthermore, the low-index surface atoms of ordered phases can also be distinguished by these indexes if the feature values are available. As shown in Figure 2, however, the surface configurations of Cu−Pt NAs exhibit little ordered features because surface compositions are far away from stoichiometry due to Cu preferential segregation. In the following analyses, we focused primarily on quantifying the chemical ordering in the core region of Cu−Pt NAs by the local mixing index σ. The OCT8119, TOC9201, and CUB8217 NAs were further analyzed for understanding the composition and morphology effects on the chemical orderings. The midmost cross sections F


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Figure 7. Statistical results of σCu and σPt for OCT8119 NAs with different global Cu compositions.

and CUB NAs, surface enrichment of Cu is more apparent for Cu−Pt NAs with more open surfaces, but only CuPt and Cu5Pt3 NAs exhibit apparently morphology-dependent L11 ordering, in the following order OCT > TOC > CUB. Nevertheless, the modification or suppression effects of surface segregation on bulk ordering in NAs further highlight the limited validity of Cu−Pt bulk phase diagrams for NAs. 3.4. Multishell Structures. As mentioned above, both the surface segregation of Cu and the chemical ordering are the general rules in Cu−Pt NAs and usually compete with each other. Even so, counterintuitive synergy effects between surface segregation and bulk ordering are also possible under proper conditions. Figure 8 shows the midmost cross sections of OCT NAs with different sizes and selected compositions, which were colored by atomic type in Figure 8a or the local mixing index σ in Figure 8b. Furthermore, the shell-by-shell Cu distributions in these OCT NAs were shown in Figure 8c. The n-th shell was used to distinguish the surface (the first shell) and the internal core of NAs. In Figures 8a−c, a striking feature is found that a multishell onion-ring structure (MSORS) was formed in the range of sizes considered, which consists of a pure Cu surface shell, pure Pt or Pt-rich subshell, Cu-rich third shell, and so on (see Movie S1 and S2 for details). The selected alloy compositions correspond to those of the perfect MSORS, socalled the “magic composition”. The formation mechanism of the MSORS in OCT NAs with the “magic composition” was proposed to be a result of the synergy between surface segregation and chemical ordering under proper conditions. For these NAs, strong surface enrichment of Cu helps the formation of pure Cu surface shell, whereas near equimolar compositions in the core regions prefer the L11 bulk ordering characterized by alternative stacking (111) layers of Pt and Cu. Most importantly, Cu {111} facets on the surface can largely induce the alternative stacking of Cu and Pt {111} shells instead of (111) layers and hence the transition from the L11 bulk ordering to the MSORS. This strong surface-induced quasi-L11 ordering can induce the formation of MSORS in large enough OCT NAs. We take the OCT8119, TOC9201, and CUB8217 NAs as examples to illustrate the effects of morphology and

of these NAs at 100 K were shown again in Figure 6, but the atoms were indexed to the local mixing index σ, as shown in the gradient color scheme. The statistical results of σCu and σPt for OCT8119 were shown in Figure 7. Despite the morphology differences among these NAs, one can see that several ordered phases form in the core region, depending strongly on the alloy composition. Specifically, D7 (only for CuPt7), L12 (for CuPt3 and Cu3Pt5), L13 (for CuPt3, Cu3Pt5, and Cu3Pt), and L11 (for CuPt, Cu5Pt3, and Cu3Pt) ordered phases can be found in large-sized NAs. It is noteworthy that there are some geometrical frustrations because of a competition among different {111} or {001} facets. The presences of ordered phases at specific compositions can be attributed to the heats of formation of the bulk phases given in Figure 1. For example, the L13 and L12 structures are the stable and metastable LRO phases in CuPt3 bulk alloys, respectively. The difference in heats of formation between the L13 and L12 structures is only several meV/atom. As a result, L13 and L12 structures coexist in CuPt3 NAs with comparable probability. In fact, all large NPs with thousands of atoms characterize the chemical ordering of bulk alloys, in fair agreement with experiment.15,22−24 In a confined system such as shaped NAs, the chemical ordering in core region would be modified or suppressed by surface segregation.7,67,93 The modification or suppression effect is general more than especial, as surface segregation of Cu is the rule more than the exception in Cu−Pt NAs. Due to surface enrichment of Cu, the core regions of NAs have slightly lower Cu compositions than global ones, and thus, the corresponding ordering tendency shifts to the Pt-richer side. For instance, the core regions of Cu5Pt3 and Cu7Pt NAs have nearly 1:1 and 3:1 stoichiometric composition, respectively. Thus, as shown in Figure 7, L11 ordering can be frequently found in Cu5Pt3 NAs but rarely in CuPt ones. Similarly, L13 ordering is predominant in Cu7Pt NAs but not found in Cu3Pt ones. Furthermore, the D1 phase has been suppressed for Cu7Pt NAs and should be found in more Pt-deprived NAs. As mentioned previously, the smaller the NAs are, the more obvious the surface Cu enrichments are. Hence, the modification or suppression effects are more apparent in smaller NAs (not shown here for clarity). Among OCT, TOC, G


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oscillations more or less. It is not surprising that both the amplitude and perturbation deepness in OCT are slightly larger than TOC but much larger than CUB because the ratios of {111} facets (Figure 4) follow a similar tendency. Moreover, the MSORS only forms at a narrow composition range (from CuPt to Cu5Pt3 as shown in Figures 8 and 9) where quasi-L11 ordering can be largely induced by {111} facets. As a consequence, the MSORSs are almost prefect in OCT, whereas they are slightly less perfect in small TOC (see Movie S3 and S4 for example) due to the negative contribution of {100} facets. For large TOC NAs, the multishell feature is retained in near surface shells but the core region is degenerated to a mazelike structures that exhibits more antisite defects or antiphase boundaries than MSORS. The formation of multishell and maze-like core structure (MSMCS) can be attributed to the strong alloying effects and weak surface effects in large TOC NAs. Notably, regardless of these differences in atomic stacking, both the MSMCS and MSORS are characterized by the quasiL11 bulk ordering. Furthermore, there are little such multishell features in CUB due to the largely negative contribution of {100} facets. We further take the OCT NAs with different sizes as examples to illustrate the thermodynamic stability of MS structures. The temperature and size dependence of the midmost cross sections, shell-by-shell Cu distributions and the L11 LRO of OCT489, OCT891, and OCT8119 NAs with the “magic composition” (nearly 1:1) were shown in Figure 10. First, it is seen that both the LRO and the damped oscillation of Cu concentration in different shells are gradually diminished with increasing temperature, despite the different sizes. At elevated temperatures, the MSORSs are degraded into random intermixing of Cu and Pt atoms because of the entropy effects. Meanwhile, as shown in Figure 10a−c, different surface sites (facet, edge, and vertex) are less distinguishable because of larger inward relaxations at the vertices and edges compared to the rest of the facets. Especially, some degree of surface premelting can be found on the surface of OCT489 and OCT891 NAs. Second, the particle size also affects the order−disorder phase-transition of Cu−Pt NAs. Compared with the bulk alloys, the order−disorder transition in NAs is seen to be more smooth for the range of sizes considered. That is to say, there is no clear sign of the abrupt first-order phase-transition behavior in OCT NAs. Considering the competing nature of the chemical ordering and surface segregation, it is not surprising that possible LRO in NAs will be largely suppressed by surface segregation because of the high surface-to-volume ratio of NAs. Although the order−disorder transition is continuous in nature, the transition is more abrupt in larger particles. At low temperatures (below 250 K), the multishell structure is perfect and free from defect, independent of the particle size. However, the larger particles have lower proportion of defects (revisiting Figure 8b) and thus exhibit higher degrees of LRO. At elevated temperatures, the boot is on the other foot. This reversal mainly origins from the higher core compositions of Pt in the smaller particles having higher surface Cu concentration. This crossover of LRO again confirms that the transition is more abrupt in larger particles. Similarly, continuous transition and more abrupt transition in larger particles had been observed in other NAs, such as the Fe−Pt and Au−Pd ones.57,58 Theoretical and experimental works demonstrated that similar MSORS had been obtained in other NA systems. For example, force-field and DFT calculations suggested that ORS can be found in Pt0.75Ni0.25 NAs with OCT shape,56 Au−Cu,

Figure 8. Midmost cross sections of differently sized OCT NAs with atoms indexed by (a) the atomic type and (b) the local mixing index σ as well as (c) the shell-by-shell Cu distribution in these OCT NAs. The first shell corresponds to the surface of NAs.

composition on the formation of MSORS. The shell-by-shell Cu distributions of these NAs were shown in Figure 9. Among these shaped NAs, all of the NAs exhibit composition

Figure 9. Shell-by-shell Cu distributions in OCT8119, TOC9201, and CUB8217 NAs with different global Cu compositions. H


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Figure 10. Temperature and size dependence of (a−c) the midmost cross sections, (d−f) shell-by-shell Cu distributions, and (g) the L11 LRO of OCT489, OCT891, and OCT8119 NAs with the “magic composition”.

Pd−Pt, and Fe−Pt NAs with ICO or DEC shapes,55,97−99 as well as Cu0.5Pt0.5, Pd0.56Pt0.44, and Pt0.75Ni0.25 NAs with TOC shape.56,61,100 The common features of these NAs exhibiting MSORS include coexistence of surface segregation and chemical ordering, small size that usually contain less than 2 000 atoms, proper shapes that are almost entirely (for OCT, ICO, and DEC) or mainly (for TOC) enclosed by {111} facets. It is interesting to note that the corresponding bulks of these NA systems (except for the Cu−Pt system) prefer the L10 chemical ordering. Although a complete analysis of the formation mechanism of MSORS in different NA systems is beyond the scope of this work, it is apparent that the formation

of MSORS in Cu−Pt NAs under proper conditions can be largely attributed to the synergy effects between surface segregation and chemical ordering in core. 3.5. Implications for Dealloyed ORR Catalysts. The ability to control the structure and near-surface composition of NAs opens up new possibilities for optimizing their catalytic performance. As mentioned previously, the most prevalent structure in dealloyed Cu−Pt NAs gave Cu-rich core/Pt shell (0.5−1 nm or 2−5 monolayer) structure and exhibited superior ORR activity due to the stain and ligand effects.13−24 Meanwhile, experimental works demonstrated that the existence of three distinctly different size-dependent morpholI


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by selective etching out of near-surface Cu, which are expected to exhibit high activity and electrochemical stability for ORR.

ogy regimes in dealloyed Cu−Pt NAs: single-core/shell structures for particles 30 nm.17,20,21 Considering the novel structure features of multishell Cu−Pt NAs, we speculate that they can be conveniently transformed toward single Cu-rich core/Pt shell structures by selective (electro)chemical etching combined with proper heat treatment.7,13−24 First, the multishell NAs can be preannealed at room temperature in a reducing or inert atmosphere. Second, in the size range considered in present works, the single Cu-rich core/Pt shell structure can be conveniently obtained by selective dissolving the Cu species of as-prepared multishell NAs. Above room temperature, the as-prepared CuxPt1−x (0.5 ≤ x ≤ 0.64) NAs will exhibit multishell structures that are not as perfect as at low temperature (see Figure 10). After dissolution of the Cu skin, Cu atoms in the Pt-rich second and Cu-rich third shell can be further dissolved, providing a possible way for the formation of the multilayer Pt shell structures. Chemical stability of dealloyed Cu−Pt ORR catalysts is also of critical importance. As suggested in recent reports, to introduce a partially ordered core was a way toward producing robust ORR catalysts.22−24 The Cu50Pt50 ordered core of present dealloyed NAs are most thermodynamically stable, as indicated in Figure 1. Furthermore, the close shell or network of Pt formed in MSORS or MSMCS (Figures 5 and 8) will effectively protect the Cu species from further dissolution and thus promote the electrochemical stability of dealloyed ORR catalysts. Further in-depth research work along this direction is currently under way.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.5b06145. Series of images for cross sections of the onion-ring structure formed in OCT891 NAs; Cu atoms are shown in red and Pt in blue (AVI) Series of images for cross sections of the onion-ring structure formed in OCT891 NAs; the atoms are indexed to the local mixing index σ with the gradient color scheme the same as Figure 6 (AVI) Series of images for cross sections of the onion-ring structure formed in TOC586 NAs; Cu atoms are shown in red and Pt in blue (AVI) Series of images for cross sections of the onion-ring structure formed in TOC586 NAs; the atoms are indexed to the local mixing index σ with the gradient color scheme the same as Figure 6 (AVI)

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AUTHOR INFORMATION

Corresponding Authors

* (W.H.) E-mail: [email protected]. Tel/Fax: +86-73184618071. * (L.D.) E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was financially supported by the National Natural Science Foundation of China (Nos. 51301066, 51201063, and 51371080), the Natural Science Foundation of Hunan Province (Nos. 13JJ4071 and 14JJ2080), the Talents Foundation of Hunan Agricultural University (No. 12YJ04), and the Chinese National Fusion Project for ITER (No. 2013GB114001). We would also like to acknowledge the support of the Atomic Simulation Lab of Hunan University and the computation platform of the National Super-Computer Center in Changsha (NSCC).

4. CONCLUSIONS This work presented a detailed study of the surface segregation and chemical ordering patterns of Cu−Pt NAs by utilizing MC simulations coupled with MAEAM potentials that well reproduce the surface energy, heats of formation, and ground states of bulk alloys. Parameters such as alloy composition, particle shape and size, and system temperature were correlated with surface Cu composition and the chemical ordering. It was found that both the surface segregation of Cu and the chemical ordering in core are the general rules and usually compete with each other. On one hand, Cu atoms were found to significantly segregate onto the surface and preferentially occupy the lowercoordinated sites. This surface segregation is enhanced with increasing particle size or surface openness or global Cu composition. On the other hand, the types and degree of chemical ordering in the NAs’ core were quantitatively characterized by the local mixing index. The degrees of chemical ordering in the core region are similar among different morphologies, and most of the types of ordered phases are the same as bulk alloys. However, the ordering patterns are modified or suppressed by surface segregation of Cu, so that the corresponding ordering tendency shifts to the Pt-richer side or large-sized particle. Specially, multishell structures with alternative Cu or Pt shells form in a narrow composition range (from CuPt to Cu5Pt3), suggesting a subtle synergy between the completed surface segregation of Cu and the L11 bulk ordering. Due to this Cu {111} facet-induced quasi-L11 ordering, the multishell structures are preferred for OCT and TOC NAs which are mainly bound by {111} facets. Furthermore, the multishell structured Cu−Pt NAs can be conveniently transformed to single Cu-rich core/Pt shell ones

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REFERENCES

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