Catalytically Active Au Layers Grown on Pd Nanoparticles for Direct

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Catalytically Active Au Layers Grown on Pd Nanoparticles for Direct Synthesis of HO: Lattice Strain and Charge Transfer Perspective Analyses 2

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Jin-Soo Kim, Hong-Kyu Kim, Sung-Hoon Kim, Inho Kim, Taekyung Yu, Geun-Ho Han, Kwan-Young Lee, Jae-Chul Lee, and Jae-Pyoung Ahn ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.9b01394 • Publication Date (Web): 03 Apr 2019 Downloaded from http://pubs.acs.org on April 3, 2019

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Catalytically Active Au Layers Grown on Pd Nanoparticles for Direct Synthesis of H2O2: Lattice Strain and Charge Transfer Perspective Analyses Jin-Soo Kim,†,‡ Hong-Kyu Kim,‡ Sung-Hoon Kim,†,‡ Inho Kim,§ Taekyung Yu,§ Geun-Ho Han,⊥ Kwan-Young Lee,⊥ Jae-Chul Lee,*,† and Jae-Pyoung Ahn*,‡ †Department ‡Advanced

of Materials Science and Engineering, Korea University, Seoul 02841, South Korea

Analysis Center, Korea Institute of Science and Technology, Seoul 02792, South Korea

§Department ⊥Department

of Chemical Engineering, Kyung Hee University, Yongin 17140, South Korea

of Chemical and Biological Engineering, Korea University, Seoul 02841, South Korea

*To whom correspondence should be addressed. E-mail: [email protected]; [email protected].

Abstract: Despite its effectiveness in improving the properties of materials, strain engineering has not yet been employed to endow catalytic characteristics to apparently non-active metals. This limitation can be overcome by controlling simultaneously lattice strains and charge transfer originated from the epitaxially prepared bi-metallic core-shell structure. Here, we report the experimental results of the direct H2O2 synthesis enabled by a strained Au layer grown on Pd nanoparticles. This system can benefit the individual catalytic properties of each involved material and the heterostructured catalyst displays an improved productivity for the direct synthesis of H2O2 by ~100% relative to existing Pd catalysts. This is explained here by exploring the individual effects of lattice strain and charge transfer on the alteration of the electronic structure of ultrathin Au layers grown on Pd nanoparticles. The approach used in this study can be viewed as a means of designing catalysts with multiple catalytic functions.

Keywords: catalyst, hydrogen peroxide, core-shell structure, Pd@Au, strain engineering

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Au is known to be catalytically non-active for the dissociation of molecular hydrogen (H2) at room temperature (RT). In 1995, Hammer and Nørskov1 proved this by showing that the antibonding level of Au-H bonds was lower than the Fermi energy level of bulk Au. However, it was recently observed that the minor alloying of Au with Pd improved yields in the direct synthesis of H2O2.2 Subsequent studies showed that laser-irradiated Au nanoparticles promoted the dissociation of H2 at RT under visible light,3,4 which suggests that if the electronic structure of Au is altered, it may be catalytically active for the dissociation of H2.5 The imposition of residual elastic strains on a material by adding coatings of ultrathin overlayers to a substrate material is a common method for altering a material’s electronic structure. Since the first commercial application of this method in the semiconductor industry,6 strain engineering has diverse applications in various industrial sectors as an efficient way to tune the optical, electric, magnetic, and catalytic properties of materials.7-13 Furthermore, two-element heterostructures, such as alloys, thin films, and core–shell structures, encourage charge transfer because they contain materials of differing electronegativities, which promotes additional changes in the surface electronic structure.14-17 Such heterostructures can demonstrate the individual catalytic properties of each component, thereby resulting in overall enhanced catalytic properties. Of all heterostructures, the core–shell structure is an ideal configuration that induces the largest lattice strain,7,9,16 while maximizing the use of the surface areas required for catalytic reactions. One method to obtain core–shell structures is through the electrical or chemical extraction of solute atoms from the alloy (“dealloying”). However, because of the similarity in the lattice parameters of the constituent materials, the residual strains of the shell material are usually limited to less than a few percent,9 resulting in minor changes to the 2


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electronic structure. Another technique to realize the core–shell structures is epitaxial growth, which is usually applied to materials with lattice mismatches of approximately 2%.18 Despite this limitation, we showed recently that Au, having a large lattice mismatch (4.8%) with Pd, can be grown epitaxially on Pd nanoparticles (NPs).19 Core–shell structured Pd NPs with epitaxially grown Au overlayers (Pd@Au NPs) can induce large compressive residual strains on the Au shell. Furthermore, the different electronegativities of Pd and Au can promote charge transfer at the interface of this heterostructure.17,20 These features of Pd@Au NPs not only enable the systematic probing of the effects of residual strains () on their catalytic activities in the previously unexplored regime in the range of -0.05 <  < 0, but also allow the investigation of additional effects of charge transfer on catalytic activity. One application of Pd@Au NPs is to synthesize H2O2 directly from molecular H2 and O2 (Figure S1 in Supporting Information); despite the extensive studies and industrial maturity of direct electrochemical synthesis of H2O2,21-24 one problem of this process is its low H2O2 productivity. To improve the productivity of H2O2, it is necessary to promote the dissociation of H2, which has typically been achieved using alloyed Au-Pd25-28 or core-shell structured Au-Pd NPs with Pd-rich overlayers.29,30 However, these Pd-based catalytic systems, while highly active for the dissociation of H2, are also active for the dissociation of O-O bonds, including molecular oxygen (O2), which is the major reason why the productivity is less than 100% under typical reaction conditions.21,31 Unlike Pd, Au is unable to dissociate both H2 and O2 at RT.31,32 Hence, Au crystals, when combined with Pd crystals in the form of Pd@Au NPs,2,33,34 can compensate for the adverse effects of existing bare Pd and Pd-based catalysts by suppressing unintended side reactions caused by the dissociation of O-O bonds, while promoting the dissociation of H2. In this study, by structuring Pd and Au in the form of Pd@Au core-shell structures, we report the experimental results of the dissociation of H2 on 3


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strained ultrathin Au shell layers on Pd NPs for the direct synthesis of H2O2. We observed that the Pd@Au catalyst displays a 100% improvement in the productivity of H2O2 relative to existing Pd catalysts. We performed a comprehensive study on Pd@Au NPs with different shell thicknesses to investigate their high catalytic productivity. This study addresses two issues and is organized as follows: (1) We analyzed microstructural features of the Au shells of Pd@Au NPs and measured the lattice strains as a function of thickness and crystallographic orientation of the Au shell. (2) Using density functional theory (DFT), we calculated the electronic structures of the Au shells to investigate how the simultaneous action of lattice strain and charge transfer makes the Au shell catalytically active for H2 dissociation, while impeding O2 dissociation.

Results and discussion Observation of Pd@Au NPs and evaluation of their catalytic properties

Figure 1. Morphologies of Pd and Pd@Au NPs along with their catalytic properties. Brightfield TEM images of (a) Pd, (b) Pd@Au (1 nm), (c) Pd@Au (3 nm), and (d) Pd@Au (5 nm). The insets of (b-d) are high-resolution EDS images (green: Au, red: Pd) recorded from each 4


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core-shell nanoparticle. (e) Catalytic properties evaluated from various NPs in terms of H2 conversion, H2O2 selectivity, and H2O2 productivity.

Figures 1a-d show the morphologies of as-prepared pure Pd and Pd@Au NPs with various shell thicknesses (details on the crystallographic relationship between the Pd core and Au shell are explained in Figure 2). The Pd NPs are cubes of side length of ~10 nm with their facets exposed to {100} (Figure 1a). As the Au shell grows on the surfaces of Pd NPs to average thicknesses of 1-5 nm (thickness measured from the corner of the Pd cube to the Au surface), it displays intriguing changes in the growth direction and morphology. When the thickness of the Au shell grown on the Pd@Au NPs is ~1 nm, it grows uniformly on all {100} planes of Pd NPs by maintaining a cuboidal shape with truncated corners (Figure 1b). As the shell thickness increases, the morphologies of Pd@Au NPs change from a cube to faceted polyhedra with their surfaces exposed to various crystal planes. This morphological evolution is clearly seen from the mapping images of the Au shells recorded using a highresolution energy-dispersive X-ray spectroscope (EDS) equipped in a scanning transmission electron microscope (STEM), as shown in the insets of Figures 1b-d. The detailed morphologies of Pd@Au NPs with various crystal planes of the exposed facets and shell thicknesses are shown in Figures S2 and S3 (Supporting Information), respectively. The catalytic properties of Pd and Pd@Au NPs were evaluated in terms of H2 conversion, H2O2 selectivity, and H2O2 productivity through the direct synthesis of H2O2, as shown in Figure 1e (for details, see Methods). It should be noted that both Pd@Au (1 nm) and Pd@Au (3 nm) NPs display catalytic properties for H2 conversion during the direct synthesis of H2O2. This is convincing evidence that, unlike bulk Au that is catalytically nonactive for the dissociation of H2, the strained Au layer on Pd@Au NPs can dissociate H2. 5


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Furthermore, as-prepared Pd@Au (1 nm) NPs display improved H2O2 selectivity of more than 125% over Pd NPs, which indirectly indicates that the dissociation of O2 is impeded by the Au layer. The result suggests that core-shell structured catalysts consisting of Pd and Au utilized individual properties of each component and thus can resolve the low selectivity associated with O2 dissociation displayed by single-element Pd catalysts. In addition, upon observing the structures of reacted Pd@Au NPs using TEM and EDS elecmental mapping, Pd@Au NPs retained their initial structural integrity even after catalytic reaction experiments (for details, see Figure S4 in Supporting Information). Although more systematic future studies for post reaction characterization are needed, the present observation indicates that Pd@Au NPs are structurally stable and thus, can be used for repetitive operations for the synthesis of H2O2. However, as the Au layer grows thicker than 5 nm, Pd@Au NPs are no longer active for the dissociation of H2 and thus are unable to produce H2O2. Therefore, we deduce that the catalytic behaviors observed from Figure 1e are related to the structural changes of the Au layers grown on Pd NPs.

Microstructural characterization of the Au overlayers on Pd@Au NPs

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Figure 2. Orientation relationship and characterization of Pd and Pd@Au NPs. HRTEM images of (a) Pd and (b) Pd@Au (1 nm) NPs at the [001] zone axis. (c) High-resolution EDS mapping of a Pd@Au (1 nm) NP, showing the presence of the 1-nm-thick Au shell grown on Pd cube facets. (d) HRTEM images of a Pd@Au (3 nm) NP at the [011] zone axis. Note that dislocations are observed at positions indicated by the  symbol. (e) Low-angle XRD spectrum of Pd@Au (1 nm) NPs (in pink), which is superimposed on those obtained from the NPs of Pd (in red) and Au (in green). (f) EXAFS radial distance of Au-Au atomic pairs measured from the various NPs.

Figure 2a is a high-resolution TEM (HRTEM) image of the representative Pd NP recorded from the [001] zone axis. The Pd NP is the cube-shaped perfect crystal with its facets exposed to {100} planes. When an Au shell grows on Pd NPs, it initially grows uniformly by maintaining its outer surfaces parallel to {100}, as shown in Figure 2b. A close examination shows that, despite the large difference in the lattice parameters between Au (4.078 Å) and Pd (3.891 Å), the Au shell grown on the Pd {100} surfaces is observed to maintain a perfect lattice coherency (see the image in the inset of Figure 2b). The EDS line 7


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profile taken across the NP confirmed the presence of a 1-nm-thick Au layer on the surface of Pd@Au NPs, as shown in Figure 2c. As the Au shell grows thicker than ~1 nm, dislocations begin to nucleate at local regions (Figure 2d), which weakens the lattice coherency at the Pd/Au interface and thus annihilates partially the lattice strain of the Au shell (for details, see Figure S2, Supporting Information). The large difference in the lattice parameter between Pd and Au can induce residual elastic strains at the Pd/Au interfacial region of the Pd@Au NPs. Figure 2e compares the Xray diffraction (XRD) spectrum recorded from Pd@Au (1 nm) NPs, which is superimposed on those obtained from pure Au and pure Pd NPs; it is apparent that the diffraction peaks of Pd@Au (1 nm) NPs are to the left from the reference Pd peaks (denoted by the red vertical lines), suggesting that the Pd lattice in Pd@Au (1 nm) NPs is larger. A slight leftward broadening of the Pd (111) and (200) peaks, as denoted by the arrows in Figure 2e, indicates that the lattice parameter of the Au crystal is similar to that of the pure Pd crystal, suggesting that the Au lattice is in compression. The characteristic Au peaks begin to appear when the thickness of the Au shell is greater than 3 nm (Figure S5, Supporting Information), although they are shifted to the right. This suggests that the lattice parameter of the Au shell of Pd@Au (3 nm) NPs begins to retain the value characteristic of Au crystals as the shell thickness increases, albeit while under compression. For Pd@Au (5 nm) NPs, the diffraction peaks of the Au shell approach those of pure Au NPs, indicating that the Au shell of Pd@Au (5 nm) NPs is nearly free from residual strains. Dislocations observed from the 3- and 5-nm-thick Au shells (Figure S2, Supporting Information) evidence that elastic strains have been relieved via the nucleation of local deformations. To quantitatively evaluate the lattice strain of the 1-nm-thick Au shell of Pd@Au (1 nm) NPs, the interatomic spacing of Au-Au pairs was measured using extended X-ray 8


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absorption fine structure (EXAFS) spectra collected at a synchrotron (see Experimental Section), as shown in Figure 2f. The results show that the interatomic spacing of Au-Au pairs is 2.2% smaller than that of pure Au NPs, suggesting that the compressive residual strains are developed in the 1-nm-thick Au shell of Pd@Au (1 nm) NPs. However, this strain measured from EXAFS is the average value measured along all crystallographic directions. Considering that catalytic reactions occur at the surface, both the magnitude and direction of strains at the material’s surface are important parameters that affect the electronic structure and thus catalytic properties of the Au layer. Therefore, any description of strains that only consider the magnitude (measured using XRD,9,12 EXAFS,35 and simple comparison of lattice parameters9) without taking into account their direction is insufficient to establish the lattice strain versus catalytic property relationship of the core-shell structured catalysts.

Local strains in the Au overlayers of Pd@Au NPs

Figure 3. NBED strain maps of Pd@Au NPs measured as a function of shell thickness for various crystallographic orientations. NBED strain maps measured for a Pd@Au (1 nm) NP along a) [100] and b) [010], Pd@Au (3 nm) NP along c) [100] and d) [010], and Pd@Au (5 nm) NP along e) [100] and f) [011]. Note that the arrows indicated in the maps are the 9


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directions along which strain maps are evaluated. g) Variations in the strains measured from the Pd/Au interfacial region along the various crystallographic orientations (see the inset) as a function of shell thickness in the Pd@Au NPs. The error bars denote the maximum and minimum strain values measured during the five test runs per sample.

To investigate how the lattice residual strain influences the catalytic properties of Pd@Au NPs, it is necessary to simultaneously measure the magnitude and direction of strains induced at their surfaces. The d-spacing measurement based on HRTEM images is typically used to evaluate the local lattice strain along a given crystallographic orientation. However, due to the complexity and difficulty of aligning the zone-axis to HR images of NPs, the precise determination of strains based on HR images is very difficult. On the other hand, the Topspin® technique,36 based on nanobeam-electron diffraction (NBED) under e-beam precession, can measure local strains by comparing diffraction patterns of nanometer scale areas (for details, see Methods Section and Figure S6, Supporting Information). This allows high-precision strain mapping of local regions of a material under kinematic conditions, even without alignment to precise zone axes. In the following, residual strains were measured for the Au shell of Pd@Au NPs along the directions parallel ([100] and [010]) and perpendicular ([001]) to the facet surface, i.e., (001), of Pd cubes (a detailed orientation relationship is shown in the inset of Figure 3g). The facet planes of the Au shell in Pd@Au (1 nm) NPs and Pd NPs share the same orientation, i.e., {001}. Hence, NBED strain maps were recorded for Pd@Au (1 nm) NPs from the [001] zone axis (perpendicular to the facets of Pd NPs), as shown in Figures 3a-b. The lattice strains of the 1-nm-thick Au shell differ depending on the orientation of the facets. For example, the Au shell grown on the Pd (100) plane is nearly free from strain normal (i.e., 10


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[100]) to the Pd facet, whereas it is under compression (-4.5%) parallel (i.e., [010]) to the Pd facet (Figure 3a). In agreement with structural symmetry, strains similar to those evaluated along [010] are measured along [001] (Figure 3b). As the thickness of the Au shell increases, the compressive strains along the directions parallel to the facet planes of the Pd NP gradually decrease; for Pd@Au (3 nm) NPs, the Au shell is under compression (-2.0%) in the directions (i.e., [100], [010]) parallel to the facets of the Pd NP, as shown in Figures 3c-d. When the shell thickness becomes greater than 5 nm, the compressive strain on the Au shell reduces to the level of that of bulk Au, as shown in Figures 3e-f. It should be noted that the compressive lattice strain (-4.5%) imposed on the 1-nmthick Au shell is very close to the elastic limit (4.7%) of the perfect Au crystal.37-39 Therefore, as the Au shell grows to thicknesses greater than 1 nm, the elastic energy accumulated in the Au shell can promote plastic yielding, as demonstrated by the generation of dislocations (in Figure 2d and Figure S2, Supporting Information) that partially relieves the residual strain in the Au shell. The measured strains imposed on the Au shell are plotted for various crystallographic orientations as a function of shell thickness (Figure 3g). This unusual range of compressive lattice strains (-4.5% < ε < 0) obtained from the core-shell structure of Pd and Au makes it feasible to systematically probe into the catalytic activity for H2 dissociation on the strained Au shell in the previously unexplored regime of the residual strain.

Interfacial charge transfer in Pd@Au NPs

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Figure 4. d-orbital structures of Pd and Au in Pd@Au NPs. a) XPS spectra measured from Au, Pd, and Pd@Au NPs, showing the changes in the peak positions corresponding to the Pd 3d and Au 4d orbitals. The inset is the deconvolution result of the Pd 3d peaks of Pd@Au (1 nm) NPs, showing the three core-level peaks corresponding to Pd-O (in gr), Pd 3d (in red), and higher binding Pd 3d (in purple). b) XANES spectra measured from Au and Pd@Au NPs, showing the changes in the intensity of the Au L3 edge of Au in pure Au and Pd@Au NPs. The values measured at the vertical dotted line (Au L3 edge) in the inset denote the white line intensity, quantifying the electron density of the unoccupied outermost d-orbital of Au.

For the Pd-Au system, charge transfer is expected to occur from Pd to Au, because the electronegativity is smaller for Pd (= 2.2) than Au (= 2.5). Previous studies on Pd-Au binary alloys show that Pd loses s and p electrons and gains d electrons, when charges are 12


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transferred from Pd to Au; whereas, Au gains s and p electrons and loses d electrons.40,41 However, the charge transfer behavior observed in Pd-Au alloys may not be the case for Pd@Au core-shell heterostructures, because of differing structural configurations.17 The alteration of the d-orbital structures of Pd and Au by charge transfer was confirmed by X-ray photoelectron spectroscopy (XPS). Figure 4a shows the XPS spectra of Pd@Au NPs with differing shell thicknesses, showing the positions of core-level peaks corresponding to the Au 4d and Pd 3d orbitals. For comparison, the core-level peaks were also obtained from pure Au and pure Pd NPs. The Au 4d core-level peaks for Pd@Au (5 nm) NPs and those of Au NPs coincide with each other, indicating that charge transfer from the Pd core to the outer Au layer is negligible. For Pd@Au (3 nm) NPs, the Au 4d core-level peaks are shifted slightly to the lower energy side, indicating that Au is in the reduction state by charge transfer from Pd to Au.40 For Pd@Au (1 nm) NPs, the electronic structures of the Au layer undergo a marked change such that the core-level peaks corresponding to the Au 4d orbital are no longer observed from Pd@Au (1 nm) NPs. Rather, when comparing the XPS spectrum of Pd@Au (1 nm) NPs with that of Pd NPs, the major peaks of Pd@Au (1 nm) NPs correspond to the Pd 3d peaks. The result indicates that the electronic structure of Au in the shell is transformed to Pd-like structures. Although the mechanism of this peculiar transformation is largely unknown, one possible explanation for this behavior is that the lattice parameter of the Au crystal of Pd@Au (1 nm) NPs becomes similar to that of the Pd crystal, as shown by the XRD spectrum (Figure 2e) and strain maps (Figures 3a and b). Systematic future studies are necessary for a comprehensive investigation of this issue. A close examination of the Pd 3d peaks of Pd@Au (1 nm) NPs shows that they are characterized by a tail on the higher energy side. Upon deconvolution (see the inset graph in Figure 4a), the peaks are separated into three characteristic peaks. The 13


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lowest intensity peaks (in gray) correspond to the Pd-O core-level peaks, indicating the presence of minor quantities of oxide layers formed probably during hydropyrolysis of Pd@Au NPs on the surface of Pd NPs. The appearance of the Pd 3d peaks of higher binding energy (in purple) is a typical result of the strong interaction between Pd and other transition metals including Au,42,43 which signifies the electron loss from Pd in the Pd@Au core-shell structure due to the difference in electronegativities. These electrons are considered to move to the unoccupied Au 5d orbitals as explained below. The Pd 3d core level peak is indicated in red. X-ray absorption near-edge structure (XANES) is an absorption spectroscopy based on synchrotron X-ray sources, which permits the determination of the electron density at the core-level orbital of an element.44 Here, the state of the Au 5d outermost orbital was quantified by measuring the XANES of the Au L3 edge (i.e., the energy required to excite electrons from the 2p to 5d orbitals, 11.9237 keV) of Au and Pd@Au NPs to determine whether charges are transferred from Pd to Au. Figure 4b shows the XANES spectra recorded from Au and Pd@Au NPs, showing that the white line intensity corresponding to the Au L3 edge becomes lower as the shell thickness of the Pd@Au NPs thins. This indicates that electrons of the Pd outermost orbital move to Au to fill the initially unoccupied Au 5d orbital of the Au shell constructing Pd@Au NPs. Synthesizing the results shown in Figures 4a and 4b, we conclude that the differing electronegativities of Pd and Au comprising the core-shell structure induces electrons in the Pd core to move toward the Au shell to fill the unoccupied Au 5d orbital.

Density functional theory calculations

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Figure 5. Effects of strain and charge transfer on the DOS of Au-H and Au-O bonds. Density of states (DOS) of (a-b) Au-H and (c-d) Au-O bonds calculated for pure Au and Pd/Au heterostructure models as a function of strain.

Although EDS data (Figure 2c) showed the presence of ultrathin Au layers on the surface of as-prepared Pd@Au NPs, these layers can restructure even at RT.45 In addition, reactive environments such as the reaction temperature, pressure, and atmosphere (e.g., O2, H2, O and H atoms) can cause Pd to preferentially segregate to the surface of Au shells (socalled fluxionality).46-49 Pd atoms, even with very small quantities, can introduce significant catalytic activity by hydrogen spillover mechanisms.50-52 Therefore, to elucidate the existence of Pd atoms in the Au shell, additional analyses were performed on Pd@Au NPs using electron energy loss spectroscopy (EELS) and three-dimensional atom probe (3DAP) tomography. Experimental results obtained from EELS (Figure S7 in Supporting Information) and 3DAP (data not shown) did not show any clear evidence for the existence of Pd atoms on the outermost Au shell. However, this result alone is insufficient to ensure that the outermost surface of Au shells of Pd@Au NPs is free of Pd atoms. This is because the 15


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analytical technique employed in this study has a limitation in detecting a very small amount (1-5%) of Pd atoms that might exist on the surface of Pd@Au NPs. Moreover, the vacuum environment used for STEM-EELS can induce enrichment of Au at the surface of Pd@Au NPs due to its low surface energy, rendering this problem more complex. Therefore, the issue on whether Pd atoms reside on the Au surfaces remains unsettled and more advanced analyses, such as in-situ infrared spectroscopy experiments in the reactive gas environment, are needed.20,45,53 Despite the uncertainty about the presence of Pd atoms in the Au layer of Pd@Au NPs, the catalytic activity of the Au layer was evaluated using density functional theory calculations by considering the compressive strain and charge transfer effect, while ruling out the potential effect of Pd atoms. Both the compressive strains imposed on the Au shell and charge transfer occurring at the Pd/Au interface can alter the electronic structures of the Au shell. This in turn leads to changes in the molecular orbitals of Au-H and Au-O bonds, affecting the ability of Au to dissociate H2 and O2 and thus catalytic activities for the direct synthesis of H2O2. This was clarified here by evaluating the individual effects of the in-plane lattice strain and interfacial charge transfer on the electronic structures of pure Au and Pd/Au heterostructure models replicating Pd@Au NPs (for the model construction see Methods). First, in order to evaluate the effect of compressive strains alone on H2 dissociation, the density of state (DOS) is calculated for the Au-H molecular orbital as a function of strain imposed on the Au lattice, as shown in Figure 5a. The results show that as the value of the compressive strain increases, the antibonding energy level (ABEL) of Au-H bonds also increases. When a compressive strain greater than -4% is imposed on the Au crystal, the ABEL of Au-H bonds increases by 0.2 eV and exceeds the Fermi energy level (see the orange line in the inset of Figure 5a). This suggests that the imposition of compressive strains on pure Au would force 16


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the energy barrier for H2 dissociation to decrease. When evaluating the changes in the energy associated with the dissociative adsorption process of H2 using nudged elastic band (NEB) calculations, the energy barrier for the dissociation of H2 decreases as the value of the compressive strain imposed on pure Au increases; it reduces from 1.20 eV for unstrained pure Au to 1.02 eV for strained pure Au ( = -5%) (see Figure S8, Supporting Information). This explains why the dissociative adsorption of H2 can be activated by the strained Au layer on Pd@Au NPs. Having identified the effect of strain on the DOS structures of Au-H bonds, we next evaluated the combined effects of lattice strain and charge transfer on H2 dissociation by calculating the DOS structures as a function of strain imposed on the Au slab in the Pd/Au heterostructured model using procedures similar to those discussed previously. According to calculations performed on the Pd/Au heterostructured model (Figure 5b), the ABEL of Au-H bonds is 0.15 eV higher in the Pd/Au heterostructure than in pure Au, even in the absence of strain (see the span enclosed by the two vertical green lines in Figures 5a-b). When lattice strain is imposed on this heterostructure, the ABEL of Au-H bonds begins to increase and exceeds the Fermi energy level of the Pd/Au composite when  = -2%, indicating that the Pd/Au heterostructure becomes catalytically active for H2 dissociation. This analysis is consistent with the experimental observation that Pd@Au (3 nm) NPs with  = -2% dissociate H2. In summary, lattice strains imposed on Au on the Pd/Au heterostructure increases the ABEL of Au-H bonds, whereas interfacial change transfer causes an additional increase in the ABEL. This interpretation satisfies the theoretical condition for H2 dissociation proposed by Hammer and Nørskov.1 Despite the suitability of existing Pd catalysts for the direct synthesis of H2O2, they easily dissociate O2 during catalytic reactions, causing them to display low H2O2 selectivity 17


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and thus low H2O2 productivity. As such, to develop excellent catalysts for the direct synthesis of H2O2, it is also necessary to impede the breaking of O-O bonds, while promoting H2 dissociation during catalytic reactions. Au, in addition to its ability to impede the dissociation of O2 under most circumstances, can dissociate H2 under strain and thus can be active in the direct synthesis of H2O2. Hence, we evaluated the ability of Au for impeding O2 dissociation as a function of strain. Figure 5c shows the DOS structures of Au-O bonds calculated for pure Au. The ABELs of Au-O bonds always lie below the Fermi energy level regardless of the values of in-plane strain, indicating that pure Au is unable to dissociate O2. Furthermore, unlike the case of Au-H molecular orbitals, the ABELs of Au-O bonds are shifted lower as compressive strain increases, making Au more non-active for O2 dissociation. To evaluate the combined effects of lattice strain and charge transfer on the DOS structures of Au-H bonds, we repeated calculations for the Pd/Au heterostructure model (Figure 5d), using procedures similar to those applied to pure Au. As seen previously, compressive strains tend to shift the ABELs of Au-O bonds below the Fermi energy level. Other consequences of the compressive strain imposed on the Au shell are the widening of the d-band and the reduction of the d-band center energy (see Figure S9, Supporting Information), which promotes wider separation of ligand splitting upon the reaction between Au and H and make the dissociation of O2 even more difficult.54

Conclusion In summary, Pd@Au core-shell NPs induce the development of in-plane compressive strains on the Au shell, while promoting interfacial charge transfer from Pd to Au. The simultaneous action of compressive strains imposed on the ultrathin Au shell 18


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and interfacial charge transfer from Pd to Au promote H2 dissociation, while inhibiting O2 dissociation. This induces the ultrathin Au overlayer (~1 nm) on Pd realized by epitaxial growth to exhibit the catalytic properties of both Pd and Au in the heterostructured catalyst, resulting in ~100% improvement in productivity for the direct synthesis of H2O2 relative to existing Pd catalysts. The present work can be viewed as an alternative route to transform catalytically non-active materials to active ones and provide them with multiple catalytic functions by engineering their electronic structures via the simultaneous control of lattice strain and charge transfer.

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Methods Synthesis of cube-shaped Pd NPs and Pd@Au NPs Pd nanoparticles (NPs) were synthesized by the hydropyrolysis method. Polyvinyl pyrrolidone (PVP, MW = 55,000, 105 mg, Aldrich), L-ascorbic acid (60 mg, Aldrich), and potassium bromide (KBr, 300 mg, Fisher) were dissolved in 8 mL of deionized (DI) water at 80 °C while stirring with a magnetic stirrer. The solution was then mixed with the Pd precursor solution (i.e., 3 mL of an aqueous solution containing sodium tetrachloropalladate (Na2PdCl4, 57 mg, Aldrich)) for 3 h at 80 °C. Synthesized Pd NPs were collected by centrifugation followed by washing several times in acetone, ethanol, and water and were redispersed in DI water (8 mL). To synthesize Pd@Au NPs, 1 mL of redispersed Pd NPs suspension, PVP (5 mg), and L-ascorbic acid (6 mg) were dissolved in DI water (8 mL) and heated to 95 °C while stirring in air. The thickness of the Au shell was controlled by dissolving differing amounts of HAuCl4. After all synthesis, the aqueous solution was cooled to room temperature (RT). Other details for synthesizing Pd@Au NPs are reported elsewhere.19

Measurement of catalytic properties Synthesized NPs were dispersed in a DI water solution containing an appropriate amount of silica gel (Sigma-Aldrich) followed by mixing at RT and in air for 24 h. After impregnation, the product was collected by centrifugation and redispersed in acetic acid solution (30 vol.% in DI water) while stirring at RT for 8 h to remove PVP from the surface of NPs. The recovered catalysts were centrifuged and dried at 60 °C for 10 h. Before activity tests, the catalysts were reduced at 40 °C for 3 h in a flowing H2 gas (10 vol.% in N2 gas) at a rate of 50 mL/min. 20


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The catalytic properties of synthesized Pd@Au NPs were measured for the direct synthesis of H2O2 using a double-walled glass-jacket reactor containing 150 mL of aqueous solution mixed with ethyl alcohol (20 vol.%, Aldrich) and phosphoric acid (H3PO4, 0.03 M, Aldrich). The reactant gas was supplied at a rate of 22 mL/min by maintaining 2 mL/min of H2 and 20 mL/min of O2 flow. Each catalyst (0.2 g) was then added into the reactor to promote the reaction at 20 °C and 1 atm for 3 h. During the reaction, the H2 concentration of the outflow gas was measured by gas chromatography (Younglin, ACME 6000), whereas the H2O2 concentration was measured after the reaction by the iodometric titration method. The catalytic properties such as H2 conversion, H2O2 selectivity, and H2O2 production rate were calculated according to the standard equations provided elsewhere.55

Measurements of the thickness and local strains at Au shells High-resolution microstructures of synthesized samples were observed using transmission electron microscopy (TEM, FEI Titan TEM) performed on Pd@Au NPs by placing them on ultrathin carbon-mesh copper grids at 300 kV. The average thickness of the Au shell was measured using high-angle annular dark field scanning TEM (HADDF-STEM, FEI Talos) equipped with an energy-dispersive spectroscope (EDS, Super-X EDS) at 200 kV. Average strains of Au shells were measured using extended X-ray absorption fine structure (EXAFS) spectra corresponding to the Au L3 edge (the energy required to excite electrons from 2p3/2 to 5d5/2, E0 = 11923.7 eV) collected from a synchrotron at the 1D beamline of Pohang Accelerator Laboratory (PAL, South Korea). The incident beam was monochromatized by a Si (311) double-crystal monochromator and detuned by 30% to minimize higher harmonics arising from the contamination of Si. The EXAFS data were 21


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collected in fluorescence modes and analyzed with ATHENA-ARITEMIS software, which enabled the Fourier transform of EXAFS data to r-space.56 Nanobeam electron diffraction (NBED) strain maps were measured using a FEI TECNAI G2 transmission electron microscope equipped with Nanomegas® Digistar hardware at 200 kV. The e-beam precession and scanning for NBED measurements were controlled by the Topspin® software platform, which enabled the acquisition of NBEDs for every scanning point.57 The strain states of the NP samples were then evaluated by comparing the NBED patterns recorded from the NPs with those corresponding to the unstrained/reference sample. For this purpose, the NBEDs were recorded by scanning NP samples with a scanning step size of 1 nm and a precession angle of 0.13°.

Charge transfer analyses Charge transfer from Pd to Au was confirmed by assessing the oxidation state of Pd using X-ray photoelectron spectrometry (XPS, PHI 5000 VersaProbe Ulvac). XPS was performed and exposed a monochromated Al Kα (1486.6 eV) with a spot size of 100 × 100 µm2 to surfaces prepared by placing the aqueous suspension of NPs on the Si wafer (0.5 × 0.5 cm2). The binding energy of the XPS spectra was calibrated to that of C 1s (284.6 eV). The reduction state of Au was confirmed by comparing the intensity change of X-ray absorption near edge structure (XANES) spectra corresponding to the Au L3 edge (E0 = 11923.7 eV) collected for Pd@Au NPs at the 1D beamline of PAL.

Atomic simulations To study the effect of strains and charge transport on the H2 and O2 dissociation, the 22


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electronic structure of the Au shell was evaluated using ab initio density-functional theory calculations performed on the Pd/Au bilayer composite models replicating Pd@Au NPs; three layers of Pd with a (100) surface were fixed as a substrate, on which four layers of the Au slab were placed such that Au(100) planes were in contact with Pd(100) by maintaining the lattice coherency. Lattice strains were then applied along the in-plane direction within the range of -6 <  < 2. To calculate the density of states (DOS) projected over the molecular states of H and O in H- and O-Au shell bonds, 0.5 monolayer of H and O atoms were added to a clean Au surface. A 20-Å-thick vacuum layer was added above the free Au surfaces used for all calculations. All DOS calculations were performed using density functional theory as implanted in Quantum-Espresso.58,59 The Kohn–Sham equations were solved with the Perdew–Burke–Ernzerhof (PBE) exchange-correlation functional approach.60 The interactions between electrons and ions were described by the projector augmented wave (PAW) method.61 The electron wave functions were expanded using a plane-wave basis set up to an energy cutoff of 625 eV. A Monkhorst–Pack grid of (9 × 9 × 1) k-points was used for Brillouin zone integration. Furthermore, to investigate the energy changes required for the adsorption and dissociation of H2 molecule according to the compressive strain, the nudged elastic band (NEB) method62 was used with an additional 13 images to interpolate between the first and last states. All conditions in the NEB calculations were equal to the other DFT calculations, except using larger (3 × 3) supercell for hindering the interaction between H2 molecules in periodic cell.

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Associated Content Supporting Information The details of the morphologies and crystallographic orientations (HRTEM images), the shell thickness and compositional analysis (HAADF-STEM and EDS mapping images), lattice constant analysis (XRD spectra), strain analysis (NBED strain map images), and calculated dband structures are provided.

Author Information Corresponding Authors * E-mail: [email protected]. * E-mail: [email protected]. Author contributions J.-S.K. and H.-K.K. contributed equally to this work. T.Y., K.-Y.L., J.-C.L., and J.-P.A. designed this study. I.K. carried out the synthesis of the core-shell nanoparticles. J.-S.K., S.H.K., and G.-H.H. carried out the experiments. H.-K.K. carried out the DFT simulations. J.S.K. and H.-K.K. wrote this manuscript with contribution from all the authors. Conflict of interest The authors declare no conflict of interest.

Acknowledgements We acknowledge financial support from the National Research Foundation of Korea (NRF) grant funded by the Korean government (MSIP) (NRF-2016M3D1A1021140).

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