Hydrogen Adsorption at Strained Pd Nanoshells - The Journal of

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Hydrogen Adsorption at Strained Pd Nanoshells M. G. Montes de Oca,† H. Kumarakuru,‡ D. Cherns,‡ and D. J. Fermín*,† † ‡

School of Chemistry, University of Bristol, Cantocks Close, Bristol BS8 1TS, United Kingdom School of Physics, University of Bristol, HH Wills Physics Laboratory, Tyndall Avenue, Bristol BS8 1TL, United Kingdom ABSTRACT: Electrochemically induced hydrogen adsorption at Au
Pd core
shell (CS) nanostructures was investigated as a function of the Pd lattice strain. CS nanoparticles with shell thicknesses between 1 and 10 nm were prepared by selective reduction of Pd on Au nanoparticles in aqueous solution. High-resolution transmission electron microscopy images and selected-area electron diffraction patterns confirmed the formation of epitaxial Pd films on the Au cores. Examination of diffraction patterns allowed a quantitative analysis of the lattice strain as a function of the shell thickness based on Matthews theory. This analysis showed that complete strain relaxation was not achieved even for Pd layers of 10 nm. Detailed electrochemical studies of two-dimensional assemblies with controlled particle number density provided valuable information not only on the H adsorption charge in acid solution but also on the average surface roughness of individual CS nanoparticles. These results show that apparent changes in the H adsorption charge density with increasing Pd thickness above 3 nm are mainly brought about by increasing shell roughness. Comparison with literature values suggests that these roughening effects are present in extended surfaces composed of a few atomic layers. The picture emerging from these results is that the H adsorption charge (coverage) is effectively independent of the Pd lattice strain. The implications of these studies with regard to established theoretical models for predicting the reactivity of epitaxial thin films are briefly discussed.

1. INTRODUCTION Control over the interatomic distances in nanoscale materials is considered to be one of the most powerful tools for tuning their catalytic reactivity.1
3 Such control can be achieved by growing thin epitaxial films on foreign substrates. The epitaxial growth of thin metallic films on a foreign substrate can lead to substantial changes in the structure of the d band, which plays a crucial role in the catalytic activity of the material. These changes are brought about by so-called ligand and geometric (strain) effects.4
6 The former is associated with modifications in the density of states arising from electronic interactions between the overlayer and the substrate. Geometric effects are related to strain induced by the lattice mismatch of the two solids. Whereas ligand effects are particularly relevant for single epitaxial monolayers, geometric effects can still manifest themselves over several monolayers depending on the strain relaxation process. Nørskov et al. postulated that the shift of the d-band center is linearly proportional to the adsorption strength of a target molecule.3 This model, commonly referred to as the d-band model, has successfully described a number of experimental trends observed at extended metallic surfaces. One of the key challenges from the experimental point of view is to unambiguously quantify the contributions from ligand and geometric effects to the overall reactivity of the system. For instance, a very recent study by Strasser and co-workers on dealloyed Cu
Pt nanostructures demonstrated that geometric effects arising from Pt shells under compressive strain are responsible for their catalytic activity toward the oxygen reduction reaction.7 r 2011 American Chemical Society

A case study that has generated significant interest over the years involves hydrogen adsorption and evolution at epitaxial Pd layers on Au surfaces.8
11 Pd monolayers can be pseudomorphically grown on Au(111) and Au(100,) generating a lattice expanded by approximately 5%.12,13 Density functional theory (DFT) studies have concluded that ligand and geometric effects are responsible for 0.35 eV upshift of the d-band with respect to bulk Pd.5 This upshift of the d-band establishes that the H binding energy to a monolayer of Pd on Au is larger than on bulk Pd. Roudgar and Gross established from DFT analysis that the binding energy increases in the presence of a second Pd monolayer, followed by a decrease upon adding several monolayers.14 So far, a direct experimental observation of H binding energy with the thickness of Pd overlayers has not been reported. Kibler rationalized changes in the exchange current density for the hydrogen evolution reaction (HER) at epitaxial Pd films on Au(111) and Au(100) surfaces with H binding energies calculated by DFT,9 postulating that the rate of hydrogen evolution decreases as the hydrogen binding energy increases. However, the use of binding energy as the main descriptor for HER has been criticized by Trassatti and Schmickler as it disregards the reaction mechanism and the solvent effect on kinetics.15 In fact, Santos et al. have recently computed the activation energy of Received: March 5, 2011 Revised: April 14, 2011 Published: May 05, 2011 10489

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The Journal of Physical Chemistry C HER on pseudomorphic Pd monolayers concluding that Pd on Au exhibits faster kinetics than on bulk Pd.16 This theoretical analysis appears consistent with experimental results reported by other authors.17
19 In the present report, we describe a detailed study of hydrogen adsorption at thin epitaxial Pd layers on Au nanostructures as a function of the lattice strain. The use of core
shell nanostructures generated by colloidal chemical synthesis is not only highly relevant for the design of fuel cell electrocatalysts, but also allows detailed structural examination by electron diffraction employing high-resolution transmission electron microscopy. Electron diffraction patterns obtained from single nanoparticles revealed a decrease in the lattice strain with increasing Pd thickness following well-established macroscopic models (Matthews theory). Two-dimensional arrays with controlled particle number density were electrostatically assembled at poly(L-lysine)- (PLL) modified indium-doped tin oxide (ITO) electrodes. Detailed electrochemical analysis allowed deconvolution of the potential range of hydrogen adsorption and absorption, as well as changes in the average surface roughness of the nanostructures. Clear evidence of surface roughening was obtained as the Pd shell thickness increased above 3 nm, which could be linked to strain relaxation by lateral surface diffusion. After normalization by the average particle roughness factor, our results show that a decrease of the lattice strain from ca. 3.5% to 1% has no effect on the average H adsorption charge (coverage) at the Pd nanoshells. The charge obtained for all of the Au
Pd core
shell nanostructures is very close to those reported for pseudomorphic Pd monolayers on Au(111) surfaces. The implications of these observations with regard to established theoretical models for predicting reactivity of epitaxial thin films are briefly discussed.

2. EXPERIMENTAL SECTION 2.1. Synthesis of Nanostructures. The synthesis of various nanostructures was performed using the following high-purity reagents: gold(III) chloride trihydrate (HAuCl4 3 3H2O, 99.9%), sodium hexapalladate(IV) tetrahydrate (Na2PdCl4 3 4H2O, 98%), palladium chloride (PdCl2, 99.9999%), sodium citrate dihydrate (C2H5Na3O7 3 2H2O, 99.5%), and L-ascorbic acid (C6H8O6, 99.7%). All solutions were prepared with ultrapure Milli-Q water (18.2 MΩ cm). Gold nanoparticles were prepared following the well-established colloidal method.20,21 The one-step method consisted of adding 10 cm3 of 1% trisodium citrate to 190 cm3 of 2.5  10
4 mol dm
3 HAuCl4 3 3H2O under reflux and vigorous agitation. The reaction mixture was allowed to reflux for 1 h. As discussed further below and in a previous work,22 the average particle diameter obtained by this method is 19.3 ( 1.2 nm. Au
Pd core
shell (CS) nanoparticles were synthesized by selective reduction of H2PdCl4 on Au seeds in the presence of ascorbic acid.23
25 Aqueous solution of H2PdCl4 was prepared by diluting Na2PdCl4 in HCl solution. Different shell thicknesses were obtained by varying the amount of 0.1 mol dm
3 H2PdCl4 added to 50 mL of solutions of the as-grown Au nanoparticles placed in an ice bath and under vigorous stirring. This step was followed by adding an excess amount of ascorbic acid (0.1 mol dm
3, 6 mL) dropwise during 1 h to avoid the formation of isolated Pd clusters. The reaction was continued for another 30 min. The relationship between the total particle diameter (DCS) and the added volume of H2PdCl4 (VPdCl42-) is approximately

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given by, DCS ¼ Dcore

V Pd ½PdCl4 2
V PdCl 4 2
1þ V Au ½AuCl4
V AuCl 4


!1=3 ð1Þ

where Dcore and V correspond to the diameter of the Au seed and the molar volume of the corresponding metal, respectively. The concentration terms relate to stock solutions of the corresponding precursors. The H2PdCl4 solution volume was adjusted to tune the thickness of the Pd layer between 1 and 10 nm according to the different ratios of Pd
Au nanoparticles. 2.2. Electrostatic Assembly of Nanoparticles. The electrostatic assembly of nanostructures was performed following previously established methods.22,26,27 CS nanoparticles were assembled on indium-doped tin oxide (ITO) electrodes modified with poly(L-lysine hydrobromide) (PLL, Mw 30000
70000). ITO electrodes were cleaned by sequential sonication in acetone, ethanol, and ultrapure Milli-Q-water for 15 min in each solvent and dried in a stream of pure argon. The clean surfaces were modified by being dipped in a solution of PLL (1 mg cm
3) for 10 min, rinsed copiously with Milli-Q-water and dried under a high-purity argon flow. Electrostatic adsorption of nanoparticles was achieved by dipping the freshly prepared PLL
ITO material into colloidal solutions for 1 h. After the ITO had been removed from the nanoparticle solution, the surface was again rinsed with Milli-Q water and dried in a stream of pure argon. The electrochemical studies reported below provide evidence of a high electrical connectivity in the nanoparticle assembly. 2.3. Electrochemistry, Microscopic, and Diffraction Pattern Characterization. A two-compartment electrochemical cell was used incorporating a Pt wire and a KCl-saturated silver/silver chloride electrode (KCl-saturated Ag/AgCl) as the counter and reference electrodes, respectively. The reference and working electrodes were placed in different compartments connected by a Luggin capillary. All potentials in this work were measured and are quoted with respect to the Ag/AgCl reference electrode. The electrochemical cell was placed in a Faraday cage to isolate it from environmental electronic noise. Measurements were carried out in a solution containing high-purity 0.1 mol dm
3 perchloric acid (HClO4, 99.999%). Solutions were purged with highly purified argon for at least 20 min prior to a series of experiments. Cyclic voltammograms were recorded with an Autolab PGSTAT30 instrument. The CS assemblies were subjected to consecutive cycling in 0.1 mol dm
3 HClO4 between
0.35 and 1.0 V at 100 mV s
1, starting at the open-circuit potential and scanning toward the negative limit. This potential range induces desorption of citrate and ascorbate molecules physically bound to the nanostructures while monitoring the integrity of the assembly as probed by the palladium oxide signal. After two cycles, the voltammetric signal became unaffected by the number of cycles. It should be noted that the structure of the Pd layer can be severely compromised upon potential cycling above 1.2 V. Nanoparticle size and size distribution were estimated by transmission electron microscopy (TEM) using a JEOL JEM 1200 EX MKI instrument and the image analysis software analySIS 3.0 from Soft Imaging Systems GmbH. Samples for TEM images were produced by placing 10 μL drops on a carboncoated copper grid of 3-mm diameter. Excess solution was absorbed by filter paper, and the sample was dried in air at room temperature. The composition of bimetallic nanostructures was determined by energy-dispersive X-ray (EDX) spectroscopy. 10490

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Figure 1. TEM images and selected-area electron diffraction patterns (SADPs) of the core
shell nanoparticles featuring a 19.3 ( 1.2 nm Au core coated with Pd shells with thicknesses of (A) 1.3 ( 0.9, (B) 2.7 ( 1.0, (C) 5.1 ( 0.9, and (D) 9.9 ( 1.0 nm. Scale bars in the TEM images correspond to 5 nm. Electron diffraction patterns were dominated by the lattice spacing at the projected center of the nanoparticles. The yellow and green dotted lines indicate the positions of the diffraction spots of Au and Pd, respectively.

Selected-area electron diffraction patterns (SADPs) were recorded using a Philips EM430 TEM instrument operating at 200 kV. The SADPs were taken from single isolated particles, using a parallel incident electron beam, with the area defined by a 10-μm selected-area aperture. The particles were oriented to show strong two-beam diffracting conditions from the particle centers, verified by observing the image contrast. The sample holder allows tilting by (60° in one direction and at least (30° in the perpendicular direction. Acoustic images of nanoparticle arrays were obtained with ex situ atomic force microscopy (AFM) measurements in acoustic mode (Veeco Multimode instrument with Nanoscope V controller with Picoforce Extender). AFM tips used in these experiments typically exhibited a radius of curvature below 20 nm, a 40 N/m force constant, and a 300 kHz resonance frequency. The AFM images were processed with Gwyddion program software.

3. RESULTS AND DISCUSSION 3.1. Pd Lattice Strain As a Function of the Shell Thickness. Representative TEM images of Au
Pd CS nanoparticles are presented in Figure 1, showing an increase of the shell thickness from 1.3 ( 0.9 nm (Figure 1A) to 9.9 ( 1.1 nm (Figure 1D). As the core size was kept constant in the synthesis, the increase of the shell thickness manifested itself by an increase of the overall particle size (the scale bar corresponds to 5 nm in all of the images). Examination of a large quantity of particles and different synthetic batches confirmed the continuous growth of Pd shells over Au cores that were single crystals or multiply twinned. The average diameters of the core
shell structures obtained from at least 200 particles per sample and their elementary compositions estimated from EDX measurements are summarized in Table 1. The mass ratio from the EDX data is highly consistent with the composition of the synthesis bath, demonstrating that Pd

Table 1. Nanoparticle Diameter (DCS), Pd Thickness (δ), and Atomic Composition as Estimated from TEM
EDX Analysis particle

DCS (nm)

δ (nm)

mass ratio (EDX)

Au

19.3 ( 1.2







CS1

21.8 ( 1.1

1.3 ( 0.9

18.7 ( 2.4: 81.3 ( 2.4

CS3

24.7 ( 1.3

2.7 ( 1.0

38.1 ( 1.1: 61.9 ( 1.1

CS5

29.5 ( 1.2

5.1 ( 0.9

57.8 ( 0.9: 42.2 ( 0.9

CS10

38.9 ( 1.5

9.9 ( 1.1

82.1 ( 0.7: 17.9 ( 0.7

nucleation exclusively occurs at Au core surfaces. The core
shell particles are electrostatically stabilized in aqueous solution due to weakly adsorbed ionized ascorbic acid molecules. Selected-area electron diffraction patterns (SADPs) from isolated nanoparticles are displayed in Figure 1. Because diffraction angles in TEM are relatively small (1
2°), SADPs depend on diffracting planes that are nearly parallel to the incident electron beam (i.e., perpendicular to the nanoparticle surface at its projected center) and parallel to the surface at the nanoparticle edge. In the limit in which the Pd shell thickness (δ) is much smaller than the particle radius, a tensile mismatch strain in Pd in the interfacial plane leads to a Poisson contraction in the perpendicular direction. This phenomenon means that SADPs should see an increased spacing between diffraction planes in the Pd layer toward the projected center (region A) and reduced spacing toward the edges (region B).28 To separate these contributions, diffraction patterns were recorded with the nanoparticle oriented such that region A was at the exact Bragg condition, with regions closer to the nanoparticle edges being off the exact Bragg condition owing to plane rotations. Consequently, the diffraction patterns in Figure 1 are dominated by the lattice spacing at the projected center of the particles. This approach generates remarkably clear diffraction patterns of CS 10491

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Figure 2. Surface strain (ε) as a function of the average Pd thickness (δ) as calculated from SADPs. The dotted line describes the strain relaxation for epitaxial layers of Pd on extended Au(111) surfaces as predicted from Matthews theory.

nanostructures, demonstrating the epitaxial Pd growth and no evident surface alloying. It is important to highlight that conventional electron diffraction patterns from ensembles of particles, which show diffuse rings arising from the polycrystalline nature of these systems, do not distinguish diffraction from the nanoparticle center and edge regions, thus introducing significant uncertainty in lattice spacing estimations. The diffraction spacing obtained from the SADPs allows the Pd lattice strain (ε) to be determined as a function of thickness, as shown in Figure 2. The Pd diffraction spot (green dotted line) is observed at larger diffraction spacing than for Au (yellow dotted line), indicating that the Pd lattice spacing in the interfacial plane was reduced compared to Au. This observation suggests that some lattice relaxation occurred in all of the CS samples, although Pd remained under considerable tensile strain. A progressive increase in the intensity of the Pd diffraction spot was also observed as the Pd/Au atomic ratio increased in the CS nanostructures. Tails on the Pd spots in directions obliquely inclined to the g vector indicate strain relaxation on differently inclined facets. To a first approximation, the misfit-thickness relationship can be rationalized within the framework of the socalled Matthews theory.29 For the simplest case of thin Pd films on thick flat Au(111) substrates, where ε is relieved by the introduction of pure-edge dislocations with Burgers vectors b of 1 /2 < 110> lying in the (111) plane, this gives     δ ln þ1 b ð2Þ ε¼b 8πð1 þ νÞδ where ν is Poisson’s ratio. Equation 2 (continuous line in Figure 2) predicts the introduction of misfit dislocations at a critical thickness of 0.6 nm, in agreement with experimental results for Pd/Au(111).30 It can be concluded from the trends in Figure 2 that a measurable lattice strain on the Pd shell occurs even at thicknesses close to 10 nm. 3.2. Two-Dimensional Assemblies of the CS Nanostructures through Electrostatic Adsorption. Illustrative AFM images of two-dimensional assemblies of CS1 and CS10 obtained by electrostatic adsorption at PLL-modified Si(111) surfaces are

Figure 3. AFM images (1 μm  1 μm) of (A) CS1 and (B) CS10 assemblies on PLL-modified Si wafers. (C) Average particle number density (ΓCS) and geometrical surface area of the CS array (Aarray) as functions of Pd thickness.

presented in parts A and B, respectively, of Figure 3. The initial PLL layer adsorbed on the Si substrate has an extremely low corrugation with a thickness below 1 nm.31 For all of the CS nanostructures, electrostatic adsorption on PLL-modified surfaces results in two-dimensional assemblies of nanostructures randomly distributed across the surface with a low density of aggregates. As exemplified in Figure 3, Si(111) surfaces featuring a native oxide layer are ideal substrates for quantifying the particle number density of electrostatic assemblies at ultrathin PLL films. By keeping both the particle concentration in the colloidal solution and the adsorption time constant, we accurately established the number of particles in the assembly (NCS) for all samples. Such accurate determination is somewhat more complex on ITO surfaces, which are characterized by flakes and globular features with dimensions similar to those of the CS nanostructures. As demonstrated in our previous studies, particle number density and spatial distribution in electrostatic assemblies are not strongly affected by the underlying substrate topography.22,26,31,32 The average particle number density, ΓCS, estimated from a large number of samples as a function of the Pd shell thickness is shown in Figure 3C. It can be clearly seen that ΓCS decreases with increasing particle size. This trend can be rationalized in terms of two aspects: (i) a decrease in the transport properties of the particles with increasing particle size and (ii) a decrease in the concentration of the as-prepared colloidal particles with increasing particle size. To generate thicker Pd layers, larger volumes of the Pd precursor solution should be added, diluting the synthesis 10492

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Figure 5. (A) Cathodic and (B) anodic charges associated with H-loading normalized by the corresponding palladium oxide reduction charge (qcO) for CS1 (solid green circles), CS3 (hollow blue squares), CS5 (solid red triangles), and CS10 (hollow black circles) as a function of the potential sweep limit. The normalized charges converge to the same value at potentials more positive than
0.20 V, indicating the transition from H absorption to adsorption.

Figure 4. (A) Cyclic voltammograms of two-dimensional electrostatic assemblies of CS1, CS3, CS5, and CS10 at 100 mV s
1 in 0.1 mol dm
3 HClO4. (B,C) Cyclic voltammograms recorded at 5 mV s
1 of (B) CS1 and (C) CS10 assemblies with negative potential limits set at (black)
0.15, (red)
0.20, (green)
0.25, and (blue)
0.30 V vs Ag/AgCl.

bath. The geometrical surface area of the CS array (Aarray = NCSaCS) is also included in Figure 3C. This parameter was calculated assuming that the average area of individual particles (aCS) could be approximated using a sphere of diameter D (see Table 1). These relationships are crucial in our analysis, allowing the normalization of the electrochemical responses with the geometric area of the Pd assembly. 3.3. Hydrogen Adsorption and Absorption in Two-Dimensional Assemblies of the Au
Pd CS Nanostructures. Figure 4A presents cyclic voltammograms of the various core
shell particles electrostatically assembled at PLL-modified ITO electrodes in 0.1 mol dm
3 HClO4 solution at 100 mV s
1. The current is normalized with respect to the geometric surface area of Pd shown in Figure 3C. The current responses at negative potentials are associated with H adsorption, absorption, and evolution, which are partially convoluted at this scan rate. As discussed further below, voltammograms recorded at lower scan rates allow for a more effective deconvolution of these responses in the potential scale. The broad increase in the current between 0.8 and 1.0 V in the positive scan corresponds to the formation of palladium oxide, whereas the reduction manifests itself by a peak at a potential close to 0.75 V in the reverse scan. The oxide reduction peak exhibits a small shift toward more positive

potentials as δ increases. More importantly, the current density associated with oxide formation and reduction also increases with increasing particle size. The dependence of the oxide current density on δ is a manifestation of changes in the CS surface roughness. Representative cyclic voltammograms of CS1 and CS10 assemblies recorded at 5 mV s
1 are contrasted in Figure 4B, C. In the range between 0.0 and
0.25 V, CS1 is characterized by a weak dependence of the current on the applied potential (Figure 4B). This response involves contributions from H adsorption as well as absorption into the ultrathin Pd shell. On the other hand, CS10 exhibits significant changes in the current density, particularly at potentials more negative than
0.2 V (Figure 4C). The higher current densities observed for CS10 assemblies at the negative end of the potential scale are consistent with an enhancement of H absorption into the thicker Pd shell. The cathodic (qcH) and anodic (qaH) charges associated with H loading and stripping as a function of the applied potential are plotted in Figure 5. These charges were calculated by the integration of cyclic voltammograms at 5 mV s
1 for at least five different samples and normalized with respect to the charge associated with the palladium oxide reduction (qcO). This normalization allows for changes in the average surface roughness of the nanoparticles to be taken into account.33,34 Two distinctive regions can be observed in this potential range: (i) 0.0 > E >
0.2 V, where the charges are little affected by the Pd shell thickness, and (ii)
0.20 > E >
0.3 V, where the charges are strongly dependent on the Pd thickness. These trends strongly suggest that contributions from H absorption to the measured charges are significant at potentials more negative than
0.15 V. Furthermore, the cathodic and anodic charges exhibit similar values, suggesting that contributions from hydrogen evolution can be neglected to a first approximation in this potential range. This onset potential for H absorption is highly consistent with data reported for Pd overlayers at extended surfaces.33
36 10493

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Figure 6. Average H absorption charge density (Qabs) and H/Pd ratios at
0.3 V as functions of Pd thickness in core
shell nanostructures.

Figure 8. H adsorption charge density (Qad) as a function of Pd thickness (δ) in core
shell nanostructures (red symbols). The adsorption charge was normalized by the geometrical (CSgeo, solid red circles) and real (CSreal, hollow red circles) surface areas of the CS nanoarrays. Our data are compared to values reported for Pd films electrodeposited on Au nanoparticles40 [Au(nano), hollow black circles], as well as on extended Au(111) surfaces (gray,13 blue35, and green triangles35). The increase of Qad with increasing δ is linked to roughening of the Pd film.

Figure 7. Oxide reduction charge (qcONCS
1) as a function of the average geometrical surface area per CS particle (aCS). The solid line corresponds to the geometrically predicted charge based on the oxide reduction charge density reported for polycrystalline Pd films (424 μC cm
2)37.

The average hydrogen absorption charge density (Qabs) and H/Pd atomic ratio as functions of the Pd shell thickness are contrasted in Figure 6. Qabs values correspond to the integration of the voltammograms in the range between
0.15 and
0.30 V, followed by normalization using the geometric area of the twodimensional assembly (Ageo). Consequently, Qabs and H/Pd correspond to average parameters defined for single CS nanostructures. As expected, Qabs monotonically increases with increasing Pd thickness, whereas the H/Pd ratios are within the range expected for β-phase characteristic of macroscopic palladium hydride.12 The apparent increase of the H/Pd ratio has also been observed in thin Pd films grown on extended Au surfaces and was rationalized in terms of high H loading in subsurface traps and grain boundaries.34 The key conclusion from this analysis is that normalization of the electrochemical responses by Ageo is self-consistent as far as H absorption is concerned. To quantify changes in the average surface roughness of the CS nanostructures, the oxide reduction charge as a function of

the geometric surface area of individual particles (aCS) is plotted in Figure 7. The charge per particle corresponds to the ratio between qcO and the total number of particles in the assembly (NCS). The experimental trend is compared to the charge corresponding to a polycrystalline Pd film, which is characterized by a charge density of 424 μC cm
2.37 Similar charge densities have also been reported for Pd overlayers on Au(111) surfaces.12 It can be seen that CS1 exhibits comparable charges to a macroscopic flat surface, indicating that the geometrical area is rather close to the real surface area of nanostructures. However, the average charge per particle significantly deviates from the geometrically predicted values as δ increases. The nonlinear increase of QcO with increasing geometrical particle surface provides clear evidence of significant surface roughening as the Pd thickness increases. Accurate estimation of roughness factors by cyclic voltammetry in the metal oxide potential range is a well-documented method.34,37,38 The approach illustrated in Figure 7 is rather unique in the sense that the geometrical area of the nanostructured assemblies is independently determined from extensive AFM analysis. Furthermore, dimensional factors such as Γ and aCS are self-consistent with estimations of the average H absorption in the Pd nanoshells (see Figure 6). Other approaches that can be utilized in the determination of roughness factors at the nanoscale involve surface-sensitive probes as reviewed by Feliu and co-workers.39 Our previous studies on Te underpotential deposition at electrostatic assemblies of Au nanoparticles, identical to those used as cores in this work, concluded that the average Te coverage per particle is similar to those reported on extended Au surfaces.22 Based on these previous observations, it can be concluded that the roughness factor of the Au cores is very small. This is consistent with the small roughness factor estimated for CS1 from the analysis in Figure 7. Figure 8 contrasts the charge density associated with H adsorption (Qad) as a function of δ in the various CS nanostructures 10494

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The Journal of Physical Chemistry C with literature values reported for Au
Pd nanostructures and extended monocrystalline surfaces. These results correspond to the average charge density for individual core
shell nanoparticles obtained from integration down to
0.15 V, taking into account the geometric surface area (CSgeo) and the real surface area (CSreal). The results show a monotonic increase with increasing Pd thickness of Qad normalized by CSgeo from 110 to 260 μC cm
2. Such an increase in Qad is comparable to the behavior reported for epitaxial Pd layers at Au(111) surfaces in HClO4 and H2SO4.13,35 A notable issue is that the observed changes in Qad in CS nanostructures occur at significantly higher thicknesses than at extended surfaces. The charge density previously reported for “relaxed” Pd films on Au(111) is 240 μC cm
2,12 a value close to that obtained for CS10. Indeed, our SADP analysis (Figure 1) confirms that the strain in the Pd nanoshells is rather small for CS10 particles. If we consider that strain relaxation in core
shell nanostructures obeys macroscopic models such as Matthews theory, it can be concluded that virtually all previous electrochemical studies carried out at extended surfaces involved a significant degree of strain in the Pd layer. Figure 8 also shows that Qad appears to be only slightly dependent on the Pd thickness after normalization of the charge density taking into account the roughness factor (CSreal). Furthermore, the average charge density obtained for all CS nanostructures corresponds to (1.0 ( 0.5)  10
4 C cm
2. This value is very close to the charge densities reported for a pseudomorphic monolayer of Pd on Au(111) surface.12 Key additional experimental evidence can be extracted from recent data published by Savinova and coworkers on thin Pd films electrodeposited on Au nanoparticles.40 They reported a charge density of approximately 1  10
4 C cm
2 that was independent of the Pd thickness for values larger than 1 monolayer (see Figure 8). The overall trend clearly shows that no significant changes in H adsorption charge density (i.e., coverage) take place as the Pd thickness is increased from a single monolayer to values of 10 nm. Consequently, it can be concluded that the H coverage is effectively independent of the Pd lattice strain. The contrasting behavior between extended single-crystal electrodes and nanoscale core
shell systems can be rationalized in terms of roughening of the epitaxial Pd layer induced by lateral surface diffusion. Roughening of strained epitaxial layers is a rather well-understood process, particularly in the context of semiconductor surfaces.41,42 As the thickness of the epitaxial layer increases above the Matthews critical thickness, strain relaxation can lead to surface undulations with a characteristic wavelength of λc. Although estimations of λc are nontrivial, the key parameter to be considered is the lattice mismatch strain (ε0). In the case of heteroepitaxial layers of GeSi alloys on Si(100) exhibiting a ε0 value close to that of Pd and Au, values of λc could be on the order of 100 nm.41 This wavelength is higher than the dimensions of the core
shell nanostructures, implying that the surface roughening is reduced compared to that of extended surfaces. Consequently, the characteristic value of Qad ≈ 100 μC cm
2 observed for nanoscopic shells and extended pseudomorphic Pd films on Au(111) can be defined as the geometric H adsorption charge density limit for atomically smooth epitaxial layers. The increase of the Pd thickness will promote surface roughening, which occurs at different length scales in extended films and nanostructured domains.

4. CONCLUSIONS The present study provides clear evidence that H adsorption in epitaxial Pd layesr on Au particles is effectively independent of

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the lattice strain. Electron diffraction studies on Au
Pd core
shell nanostructures obtained by controlled colloidal synthesis allowed the strain in the Pd film to be estimated as a function of the shell thickness. An innovative aspect of our approach is that its rotation of the particles with respect to the incident electron beam allows diffraction from single planes in the nanoparticles to be achieved. This produces remarkably clear diffraction patterns showing a decrease in the lattice strain with increasing Pd thickness. This trend closely follows a well-established model for strain relaxation at single-crystal substrates. Detailed electrochemical studies performed at two-dimensional assemblies of the core
shell nanostructures provided quantitative information on the average H adsorption charge density per individual particle. Estimation of the adsorption charge in core
shell nanoparticles featuring a Pd thickness normalized by its geometrical surface area of ca. 1 nm provided values similar to those obtained for a pseudomorphic Pd monolayer on an extended Au(111) surface. Increasing δ in the core
shell nanostructures generates an increase in the H adsorption charge, which originates from surface roughening of the particles. Surface roughening was independently quantified by examining the charge density of oxide reduction under identical conditions. Previous studies performed at Au(111) surfaces have shown a similar increase in the charge density but for considerably thinner films. The difference in the thickness dependence indicates that surface undulations characteristic of macroscopic surfaces are reduced in nanoscale domains. Taking into account the average roughness factor per particle, our results show that H adsorption charge density is independent of the film thickness. An outstanding issue is whether calculations of H adsorption binding energies based on the d-band model are compatible with the experimental observations in this work. Intuitively, it would be expected that increasing the H binding energy with increasing Pd lattice strain would lead to higher H coverages for the thinnest Pd films.10 However, we should emphasize that our experimental findings specifically concerns H coverage rather than binding energy. Based on these observations, we propose that H adsorption itself induces substantial changes in the surface structure of Pd layers that could exert a stronger influence than the Augenerated lattice expansion. Recent in situ nanogravimetric studies by Stafford and Bertocci showed that H adsorption on Pd overlayers generates a sizable compressive surface stress.43 Consequently, theoretical predictions of the catalytic activity of epitaxial layers should consider geometric and electronic effects induced not only by the substrate but also by the target molecule under operating conditions. These observations can have important implications in the design of Pd-based catalysts for fuel cell applications.

’ AUTHOR INFORMATION Corresponding Author

*Tel.: þ44 117 9288981. Fax: þ44 117 9250612. E-mail: David. [email protected]. URL: http://www.chm.bris.ac.uk/pt/ electrochemistry/.

’ ACKNOWLEDGMENT M.G.M.d.O. thanks the Mexican National Council for Science and Technology (CONACyT) for a studentship grant. H.K. acknowledges a University of Bristol ORS grant award. D.J.F. acknowledges financial support from the Swiss National Science 10495

dx.doi.org/10.1021/jp2021209 |J. Phys. Chem. C 2011, 115, 10489–10496

The Journal of Physical Chemistry C Foundation (Project PP002-68708) and the University of Bristol. The authors acknowledge the valuable contributions by Dr. Daniela Plana and Mr. J. A. Jones (University of Bristol), Dr. Paola Quaino (University of Ulm, Ulm, Germany), and Ms. Veronica Celorrio (CSIC-Zaragoza, Zaragoza, Spain). This project is part of the COST Action D36/005/06 program “Molecular structure
performance relationship at the surface of functional materials”.

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