© 2010 Macmillan Publishers Limited. All rights reserved.

Lattice-strain control of the activity in dealloyedcore–shell fuel cell catalystsPeter Strasser1,2*, Shirlaine Koh2, Toyli Anniyev3,4, Jeff Greeley5, Karren More6, Chengfei Yu2,

Zengcai Liu2, Sarp Kaya3,4, Dennis Nordlund4, Hirohito Ogasawara3,4, Michael F. Toney3,4

and Anders Nilsson3,4

Electrocatalysis will play a key role in future energy conversion and storage technologies, such as water electrolysers, fuelcells and metal–air batteries. Molecular interactions between chemical reactants and the catalytic surface control theactivity and efficiency, and hence need to be optimized; however, generalized experimental strategies to do so are scarce.Here we show how lattice strain can be used experimentally to tune the catalytic activity of dealloyed bimetallicnanoparticles for the oxygen-reduction reaction, a key barrier to the application of fuel cells and metal–air batteries. Wedemonstrate the core–shell structure of the catalyst and clarify the mechanistic origin of its activity. The platinum-richshell exhibits compressive strain, which results in a shift of the electronic band structure of platinum and weakeningchemisorption of oxygenated species. We combine synthesis, measurements and an understanding of strain from theory togenerate a reactivity–strain relationship that provides guidelines for tuning electrocatalytic activity.

Electrocatalytic energy-conversion processes are expected to playa major role in the development of sustainable technologies thatmitigate global warming and lower our dependence on fossil

fuels. More specifically, fuel cells that use polymer electrolyte mem-branes (electrochemical energy-conversion devices) are potentiallyuseful as power sources in the transportation sector, one of thelargest emitters of greenhouse gases and consumers of fossil fuels.Efficient and stable fuel-cell electrocatalysts, however, are largelyunavailable1–3, and so fundamental progress in the design of thesecatalysts is needed.

The ultimate goal in catalytic design is to have complete syntheticcontrol of the material properties that determine the reactivity4,5.Catalysts that consist of two metals (bimetallic) allow greater reactivityand more flexible design, and recent studies have focused on these cat-alysts6–14. There are three fundamental effects in bimetallic catalysis:ensemble, ligand and geometric. Ensemble effects arise when dissim-ilar surface atoms, individually or in small groups (ensembles), take ondistinct mechanistic functionalities, as demonstrated for palladiumatom pairs on gold for a gas-phase catalytic reaction7,15. Ligandeffects are caused by the atomic vicinity of two dissimilar surfacemetal atoms that induces electronic charge transfer between theatoms, and thus affects their electronic band structure. Finally, geo-metric effects are differences in reactivity based on the atomic arrange-ment of surface atoms and may include compressed or expandedarrangements of surface atoms (surface strain)16. Ligand and geometriceffects1,2,8,15,17–20 and, in some cases, all three effects15, are generallysimultaneously present and coimpact the observed catalytic reactivity.To date, however, no effective strategy to isolate and tune strain effectsin electrocatalytic systems has been achieved. Considering the dimen-sions across which ensemble, geometric strain and ligand effects areeffective, only geometric strain can impact surface reactivity over

more than a few atomic layers. Hence, a catalyst structure that consistsof a few atomic monometallic layers, supported on a substrate withdifferent lattice parameters (a core–shell structure), should isolate geo-metric strain effects. If the amount of strain in these structures canbe controlled, we could use the rich effects of bimetallic catalysis totune surface catalytic reactivity continuously.

The preferential dissolution (removal) of the electrochemicallymore reactive component from a bimetallic alloy (precursor) thatconsists of a less reactive (here Pt) and more reactive metal (hereCu)21–23 is commonly referred to as ‘dealloying’. We showed thatdealloyed Pt–Cu nanoparticles have uniquely high catalytic reactivityfor the oxygen-reduction reaction (ORR) in fuel cell electrodes24–28,which is of tremendous importance as it occurs at the cathode ofvirtually all fuel cells29,30, with pure Pt as the preferred catalyst.The ORR is sluggish, which is why significant amounts of Ptmetal are required in fuel cells, which makes them prohibitivelyexpensive. Dealloyed Pt catalysts, however, meet and exceed thetechnological activity targets in realistic fuel cells24, as shown inSupplementary Fig. S1. Owing to their reactivity, dealloyed Pt cata-lysts can reduce the required amount of Pt by more than 80%.Despite such importance, the mechanistic origin of their enhancedreactivity remains poorly understood. To realize similar propertyimprovements in related electrocatalytic systems, the principlesthat underlie their performance must be elucidated. Here, wedemonstrate that the concept of strain tuning the catalytic proper-ties, introduced above, provides such a unifying principle.

ResultsWe studied six different Pt–Cu alloy nanoparticle precursors andtheir corresponding dealloyed counterparts. We used alloy precur-sors with various initial atomic Pt/Cu ratios and preparation, in

1The Electrochemical Energy, Catalysis, and Materials Science Laboratory, Department of Chemistry, Chemical Engineering Division, Technical UniversityBerlin, 10623 Berlin, Germany, 2Department of Chemical and Biomolecular Engineering, University of Houston, Houston, Texas 77204, USA, 3StanfordInstitute of Materials and Energy Sciences, SLAC National Accelerator Laboratory, Menlo Park, California 94025, USA, 4Stanford Synchrotron RadiationLightsource, SLAC National Accelerator Laboratory, Menlo Park, California 94025, USA, 5Center for Nanoscale Materials, Argonne National Laboratory,Argonne, Illinois 60439, USA, 6Materials Science & Technology Division,’ Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831-6064, USA.

*e-mail: pstrasser@tu-berlin.de

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particular the annealing temperature. We focused on two sets ofthree Pt/Cu ratios (Pt25Cu75, Pt50Cu50 and Pt75Cu25); one set wasannealed at 800 8C and the other at 950 8C. To obtain the activedealloyed catalysts, the precursors were subjected to an identicalelectrochemical dealloying protocol in which Cu was removed pre-ferentially from the precursor particles.

We first address the structure of the Pt–Cu bimetallic precur-sors and the corresponding dealloyed catalysts. These form face-centred cubic disordered alloys31. Figure 1a,b shows energy-dispersiveelemental colour-map overlays for Pt and Cu, acquired using aprobe-corrected scanning transmission electron microscope(STEM), for the Pt25Cu75 800 8C alloy precursor and the dealloyedcatalyst, respectively; the mean particle size prior to dealloying(Fig. 1a) was around 4.5 nm (see Supplementary Fig. S2e), but thedealloyed, catalytically active form (Fig. 1b) exhibited a decreasein the average particle size to �3.4 nm (Supplementary Fig. S2e).The corresponding Pt25Cu75 950 8C alloy precursor and dealloyedcatalysts had average particle sizes of 6.0 nm and 5.1 nm, respect-ively. As Fig. 1b shows, for the dealloyed catalyst, Cu is confinedto the centre of the majority of the dealloyed Pt–Cu nanoparticles.The dealloyed particles exhibit a distinct Pt-enriched layer (blue inFig. 1b) on the surfaces of the alloy Pt–Cu cores (pink in Fig. 1b).A corresponding energy-dispersive spectroscopy (EDS) line profiletaken across the diameter of typical (�4 nm) dealloyed nano-particles confirmed the presence of a �0.6 nm, Pt-enriched layeron the surface of the dealloyed particles (Fig. 1c). Hence, the Ptand Cu elemental maps and line profiles shown in Fig. 1 provideevidence for the formation of a core–shell structure in the nano-particles after dealloying, in which a Pt-enriched shell surrounds aPt–Cu core.

Furthermore, aberration-corrected, high-angle annular dark-field(HAADF) STEM images indicate a change in the Pt and Cu distri-butions within the nanoparticles, from a uniform Pt–Cu alloy(Fig. 2a) to a morphology suggestive of a core–shell structure (Fig. 2b)32.

Further evidence comes from X-ray photoelectron spectroscopy(XPS) data, which show a large enhancement in the surfaceconcentration of Pt (Table 1). Based on the STEM, EDS and XPSdata, we estimate the thickness of the Pt shell to be 0.6–1.0 nm,which corresponds to three or more Pt-rich layers, consistent withestimates from anomalous small-angle X-ray scattering of similarmaterials33. The propensity for Cu to dissolve means that dealloyed

a b

0

20

40

60

80

100

0.5 1 1.5 2 2.5 3 3.5 4 4.5

CuPt

Inte

nsity

(c.

p.s.

)

Distance (nm)

c

10 nm 10 nm

Figure 1 | Elemental maps and line profiles of Pt–Cu bimetallic nanoparticle precursors and dealloyed active catalysts. a,b, High-resolution EDS (HR-EDS)

elemental maps of a Pt25Cu75 bimetallic nanoparticle alloy precursor (a) and of the active electrocatalyst after Cu dealloying from the precursor (b). Pt is

given in blue, Cu in red and pink domains indicate well-alloyed Pt–Cu domains. The precursor alloy was prepared at 800 8C and is supported on carbon

black of high surface area. The dealloyed catalyst was obtained by voltammetric cycling of the precursor between 0.05 V and 1.2 V (with respect to a

reversible hydrogen electrode) in 0.1 M HClO4 solution. Pt–Cu alloy domains are discernible in the centres of the dealloyed particles with a Pt-enriched

surface. EDS elemental maps were acquired using a probe-corrected STEM and were extracted from spectrum images acquired from an area that measured

40 nm × 40 nm, with a beam size of �0.2 nm and a step (pixel) size of �0.2 nm. c, HR-EDS line profile across an individual �4 nm diameter dealloyed

Pt–Cu alloy active-catalyst particle. Data were acquired using a probe-corrected STEM with a probe diameter �0.2 nm and step size �0.2 nm. The thickness

of the Pt-enriched particle shell was �0.6 nm, as shown by the dashed lines. c.p.s.¼ cycles per second.

2 nm

2 nm

b

a

Figure 2 | HAADF-STEM images of Pt–Cu bimetallic nanoparticle

precursors and dealloyed active catalysts. a,b, Sub-angstrom resolution

images of individual (�4 nm) Pt–Cu nanoparticles that show a typical Pt–Cu

precursor alloy nanoparticle (a) and a typical Pt–Cu dealloyed nanoparticle

(b). The dealloyed nanoparticle exhibits an outer Pt-enriched shell (outer

shell images brightly) and a Pt–Cu alloy core (which images less brightly

than the shell). Contrast variations in the HAADF-STEM images result from

the atomic number (Z) difference between Pt and Cu (commonly referred to

as Z-contrast imaging). The thickness of the Pt-enriched particle shell in (b)

is �0.6 nm.

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Pt50Cu50 and Pt75Cu25 nanoparticles possess a similarly Pt-enrichedshell structure. From these combined, consistent results, we con-clude that the geometry of the dealloyed Pt-enriched shell on thePt–Cu core leads to the high electrocatalytic activity. Specifically,we hypothesize that lattice strain in the Pt shell controls thesurface catalytic reactivity.

We address our lattice-strain hypothesis using anomalous X-raydiffraction (AXRD), which permits independent measurements ofthe average lattice parameter and average composition of the scatter-ing nanoparticles before and after dealloying. In AXRD, diffractionpatterns are collected at a number of X-ray energies near the X-rayabsorption edge of an element of interest (here Cu). Figure 3a shows(111) diffraction profiles of a Pt25Cu75 precursor at several energies.The scattering power of Cu drops near the Cu absorption edge, andso the diffracted intensity of the alloy shows a characteristicdecrease, as indicated by the arrow in Fig. 3a. In AXRD, peakpositions yield average lattice parameters for the nanoparticles,and the relationship between scattering intensity and energy(Supplementary equation (S2)) provides the chemical compositionof the scattering alloy phase. As illustrated in Fig. 3b, this modelequation for diffraction intensity (black line) was fitted to the inte-grated (111) peak intensities (red circles), and the compositions xPtand xCu were determined. Supplementary Fig. S3a provides a directcomparison of the chemical compositions, xPt, of the alloy precur-sors derived by AXRD and by inductively coupled plasma opticalemission spectroscopy.

Given the core–shell nature of the dealloyed particles, the latticestrain in the Pt shell is most relevant for surface catalysis. To esti-mate the lattice parameter in the particle shell, we approximatedthe structure of the dealloyed particles by a simple two-phasecore–shell model, as shown schematically in Fig. 4a. We assumeda pure Pt shell with lattice parameter ashell that surrounded a

Pt–Cu alloy core with lattice parameter acore. Using the AXRD-derived nanoparticle compositions and lattice-parameter data34–36,our core–shell model allowed the determination of ashell. Thestrain, s(Pt), in the particle shell relative to bulk Pt, is given byequation (1), where aPt is the bulk Pt lattice parameter.

s(Pt) = ashell − aPt

aPt

× 100 (1)

Figure 4b shows ashell as a function of precursor composition andpreparation temperature, and s(Pt) is shown in SupplementaryFig. S4. Foremost, the data show that for all catalysts the lattice par-ameter of the Pt shell, ashell , is smaller than that of pure Pt (dottedline, Fig. 4b) and so is strained compressively. With increasing Cu inthe alloy precursor and with higher preparation temperature, ashellbecomes smaller and the magnitude of s(Pt) larger. The observedsynthesis–strain trends are understood within our core–shellmodel. The lattice mismatch between the Pt shell and the Pt–Cucore causes a reduced Pt–Pt interatomic distance in the shell. Thericher in Cu the particle core, the smaller its lattice parameter,and hence the more strain induced in the shell. A similar argumentholds for the increased strain in the high-temperature materials; forthese, the bimetallic precursor phase is alloyed more uniformly withless residual, unalloyed Cu (ref. 25), which effectively makes thealloy phase richer in Cu. Both experimental and computationalwork suggest that lattice contraction induced by particle size(surface stress) in small Pt particles only becomes significantbelow a particle diameter of 2.5 nm (refs 37,38) and can thereforebe ruled out as source of strain for the dealloyed core–shell particles.

To gain insight into how the strain of the dealloyed catalysts affectsthe catalytic surface reactivity, we studied the electronic band struc-ture. The d-band model developed by Nørskov and co-workers hasbeen successful in relating the adsorption properties of rate-limitingintermediates in catalytic processes to the electronic structure of thecatalyst39–41. For simple adsorbates, such as the ORR intermediatesO and OH, this can be understood in a simple electron-interactionmodel in which the adsorbate valence p-level forms bonding and anti-bonding states with the metal d-band40,41. Population of any antibond-ing state leads to Pauli repulsion, and the bond strength is therebyweakened. A downward shift of the d-band pulls more of the anti-bonding states below the Fermi level, which results in increasing occu-pation and weaker adsorbate bonding.

2.6 2.8 3.0 3.2 3.4 3.6

(200)

8,800 eV8,890 eV8,980 eV8,990 eV (Cu edge)9,060 eV9,150 eV

Inte

nsity

(a.

u.)

Q (A–1)

(111)

8,800 8,900 9,000 9,100

800

900

1,000

|F|2

Energy (eV)

Fit

Experiments

a b

Figure 3 | AXRD-based structural and phase-composition analysis of Pt–Cu bimetallic nanoparticle precursors and dealloyed nanoparticle catalysts.

a, AXRD intensity profiles of a Pt25Cu75 alloy precursor as function of the scattering vector Q. Diffraction profiles were taken as a function of the X-ray

energy across the X-ray absorption edge of Cu at 8,990 eV. Scattering intensities of Bragg reflections (shown are the (111) and (200) reflections) decrease

(arrow) as the Cu absorption edge energy is approached. The extent of intensity decrease correlates with the Cu content of the scattering phase and can be

used to extract composition. b, Fit (black line) of the relation (Supplementary equation (S2)) between the square of the scattering amplitude |F|2 and incident

X-ray energy to experimental values (red circles) to extract the molar fractions of Pt and Cu in the scattering phase. AXRD provided the lattice parameter of

the scattering phase (structural information) as well as its actual composition (chemical information). a.u.¼ arbitrary units.

Table 1 | Composition depth profiles for Pt25Cu75.

Photoelectron energy (eV)

264 616 1,133 1,480 8,000

Pt25Cu75 precursor (Pt at %) 8 12 19 18 31Pt25Cu75 dealloyed (Pt at %) 84 68 55 56 59

Composition depth profiles for Pt25Cu75 nanoparticles annealed at 950 8C before and afterdealloying. Increasing photon energy correlates with increased probing depth. Pt at%¼ platinumatomic per cent.

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We extended and applied these ideas to single-crystal Pt surfaces bypreparing and characterizing bimetallic single-crystal model surfacesthat consist of atomic layers of Pt with various thicknesses grown ona Cu(111) substrate. This model mimics the structural and electronicenvironment of Pt layers that surround a particle core with signifi-cantly smaller lattice parameters, similar to the dealloyed Pt–Cu cata-lysts42. Figure 5a shows the photoelectron spectra of the valence bandof Pt in Pt(111) and for five monolayers (5 ML) of Pt grown onCu(111), measured at grazing electron emission to enhance thesurface sensitivity. There is no detectable Cu signal in the 5 ML Pt spec-trum, and it thereby represents pure Pt with no ligand effect from theunderlying Cu substrate. Based on low-energy electron diffraction(LEED) probing of the lattice parameter during the growth of Pt onCu(111), 5 ML Pt should give about 2.5+0.3% compressive strain,as shown in Supplementary Fig. S6. The spectrum of Pt on Cu(111)shows a much broader d-band compared to that of the Pt(111)surface, and the d-band centre is downshifted from 2.87 eV to 3.26 eVbelow the Fermi level. The broadening is related directly to the com-pressive strain, because the electronic state overlap between the metalatoms increases with shorter interatomic distances; furthermore,keeping the d-occupancy constant for a pure metallic system givesrise to a downward shift of the d-band centre43.

Next, we used X-ray spectroscopy to monitor directly the positionand atom-specific occupation of the oxygen 2p and Pt 5d antibondingstates projected onto the oxygen atom40,41. Figure 5b shows the K-edgeX-ray emission spectra (XES) and X-ray absorption spectra (XAS) ofatomic oxygen adsorbed on thin films of Pt on Cu(111) and Pt(111) toshow the occupied and the unoccupied electronic states, respectively.For oxygen on Pt(111), we observed a broad occupied bonding state inthe XES spectrum and an intense resonance related to the antibondingstate in the XAS spectrum. For the two Pt films on Cu(111), whichcorrespond to strains of around 2.8+0.3% and 3.3+0.3% (ref. 42),we observed a decrease in the intensity of the antibonding resonancewith increasing strain. Of primary interest is that, for the film withmaximum strain, the antibonding resonance vanishes in the XASspectrum and is resolved in the XES spectrum, which directly indi-cates that the antibonding state is occupied fully with a peak around1.5 eV below the Fermi level.

Pt shell

Pt–Cu core

acore

ashell

12

a b

25 50 753.70

3.75

3.80

3.85

3.90

3.95

Uni

t cel

l par

amet

er, a

shell (

Å)

Precusor Cu content (%)

Bulk Pt

Figure 4 | A simple structural two-phase core-shell model for the dealloyed nanoparticles and evaluation of their lattice parameters. a, Scheme of a

simple two-phase structural model for the dealloyed state of a bimetallic particle. Pure Pt layers surround an alloy particle core (ashell and acore represent the

mean lattice parameters in shell and core, respectively). Deviations of ashell from the bulk unit-cell parameter of pure Pt (aPt) imply strain, s(Pt). Shell and core

regions possess a face-centred cubic structure and are assumed to be structurally and compositionally uniform. Using a simple linear model (Supplementary

equation (S3)), ashell can be estimated without further assumptions about the volume of the core or shell. As Pt–Cu alloys exhibit unit-cell parameters that are

smaller than those of pure Pt, the model predicts ashell , aPt; that is, compressive strain in the shell. Unlike the model, a real core–shell particle shows a

gradient in lattice parameters, with strain in the surface layer (point 1) being less than in layers near the core (point 2). b, Determination of ashell for dealloyed

Pt–Cu bimetallic particles, plotted as a function of the alloy precursor Cu atomic composition at precursor annealing temperatures of 950 8C (blue) and

800 8C (red). Dealloyed catalysts derived from precursors richer in Cu or precursors annealed at higher temperatures display a smaller lattice parameter in

the shell. All alloys show lattice parameters below that of pure bulk Pt (dotted line). Error bars refer to resulting uncertainty involved in obtaining ashell from

experiments and relations (Supplementary equations (S2), (S3), (S17)).

10 5 0 –5

Binding energy (eV)

2.6 ML

Pt(111)

3.5 ML

Pt/Cu(111)

0%

XAS

Inte

nsity

(a.

u.)

Inte

nsity

(a.

u.)

Strain

3.3%

2.8%

XES Adsorbed oxygen

Bonding Antibonding

b

8 6 4 2 0

Binding energy (eV)

Cu(111)

5 ML Pt/Cu(111)Pt(111)

e–

15°

a

Figure 5 | Surface-science XAS and XES studies of single-crystal model

systems that mimic dealloyed bimetallic core–shell structures.

a, Photoelectron spectra of the valence-band region for Cu(111), Pt(111) and

5 ML of Pt on Cu(111) measured with a photon energy of 620 eV and with

a 158 grazing electron-emission angle. The inelastic background has been

subtracted. The spectra of the 5 ML Pt on Cu(111) is dominated completely

by the emission from Pt because no sharp Cu d-band emission is seen at

2.3 eV from the underlying substrate (similar spectra are also found for 3 ML

of Pt). b, In-plane polarized oxygen K-edge XAS and normal emission oxygen

K-edge XES of 0.2 ML of oxygen chemisorbed on 3.6 ML and 2.6 ML films

of Pt on Cu(111) and on bare Pt(111). The XAS were normalized with respect

to the high-energy region, where the cross-section is dominated by atomic

effects related to the photoionization continuum, whereas the absolute

intensity scaling between the XES and XAS spectra is arbitrary. The energy

scale is with respect to the oxygen 1s binding energy, which represents the

Fermi level38. The strain was estimated based on the Pt coverage and using

the LEED data shown in Supplementary Fig. S6.

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To quantify the relationship between surface strain, O and OHbinding energies and the catalytic ORR reactivity, we carried outdensity functional theory (DFT) calculations to predict the changes inthe Pt–O surface bond energy for a strained Pt(111) model surface.Our computations showed a linear relationship between lattice strainand the adsorbate bond energy, consistent with the experimental X-ray spectroscopy data and previous computational analyses16,44. Bycombining the strain-bond energy relationships with a microkineticmodel, originally developed by Hammer and Nørskov39, for the electro-reduction of oxygen18,45, we derived a ‘volcano’ relation between the pre-dicted ORR rate and the strain (the dashed line in Fig. 6). The volcanoshape implies that compressive strain first enhances the overall ORRactivity by reducing the binding energy of intermediate oxygenatedadsorbates and, thereby, lowers the activation barriers for proton- andelectron-transfer processes. Beyond a critical strain, however, thebinding becomes too weak and the catalytic activity is predicted todecrease because of an increased activation barrier for either oxygen dis-sociation or the formation of a peroxyl (OOH) intermediate39.

We also plotted (Fig. 6) the experimentally measured ORR electro-catalytic activities for our Pt–Cu catalysts as a function of s(Pt) (elec-trochemical currents are reported in Supplementary Table S1). Wedid not, however, observe a decrease in the experimental activityvalues on the left side of the volcano curve, as predicted by theory;this is probably related to compressive-strain relaxation in the Ptshells. Pt atoms adjacent to the Pt–Cu cores (Fig. 4a, point 2) adopta lattice parameter closer to that of the cores32, but outer Pt shellatoms (Fig. 4a, point 1) relax towards the lattice constant of bulk Pt(ref. 42). Hence, the surface strain is less than that represented byashell, which is an average strain in the Pt shell; if the surface strainwas plotted in Fig. 6, we would expect a shift of all experimentaldata points to the right. Furthermore, it is difficult to prepare dealloyed

nanoparticles of sufficiently high surface strain to access the truemaximum of the volcano. When the strain passes a critical point,surface relaxation probably relieves further strain and thereby limitsthe accessibility of the high-strain side of the volcano. Figure 6 alsoreveals that the set of dealloyed nanoparticle catalysts prepared atthe higher annealing temperature exhibits reduced activity at com-parable lattice strain in the particle shells, presumably because ofdifferences in the mean particle size. Hence, our strain-related con-clusions generally refer to particles of comparable size.

DiscussionDealloyed fuel-cell catalysts show unprecedented electrocatalyticactivity for the electroreduction of oxygen, yet we did not have afundamental understanding of the mechanistic origin of the cataly-tic enhancement. Here, we clarify the origin, on an atomic scale, ofthe exceptional electrocatalytic activity of dealloyed Pt–Cu nanopar-ticles. We present microscopic and spectroscopic evidence for theformation of a Pt–Cu alloy core–shell nanoparticle structure usingSTEM elemental maps (Fig. 1), XPS depth profiling (Table 1) andanomalous small-angle X-ray scattering33. Given the thickness ofthe pure Pt shell and considering the limited range of ligandeffects46, we conclude that compressive-strain effects rather thanligand effects are responsible for the exceptional reactivity of theparticle surface. This is in contrast to other ORR electrocatalystsystems, such as Pt monolayer20 or Pt skin17 catalysts, in whichstrain and ligand effects are always convoluted. Using X-ray diffrac-tion, we measured and quantified the presence of compressive latticestrain in the Pt shells of the dealloyed particles.

To further corroborate the lattice-strain hypothesis in core–shellstructures, we studied a surface-science core–shell model system.Our goal was to verify experimentally the predicted effects on theband structure for compressively strained Pt layers. In contrast toprevious reports of correlations in reactivity-band structure19, inwhich band-broadening and band-centre shifts were adopted fromcomputational predictions, we demonstrated experimentally a con-tinuous change of the oxygen 2p and Pt 5d antibonding state fromabove to below the Fermi level as additional compressive strain wasapplied, which resulted in a weakening of the adsorbate bond. Thisrepresents the first direct experimental confirmation of the compu-tational prediction of band shifts of adsorbate-projected band struc-ture. Finally, we correlated experimental synthesis–strain–activitydata of dealloyed core–shell particles (Fig. 6). The resultingactivity–strain relationships provide experimental evidence thatthe deviation of the Pt-shell lattice parameter from that of bulkPt, that is the lattice strain in the shell, is the controlling factor inthe catalytic enhancement of dealloyed Pt nanoparticles; in particu-lar, these relationships are consistent with computational predic-tions that compressive strain enhances ORR activity.

In conclusion, a coherent picture of the origin of the exceptionalelectrocatalytic reactivity for the ORR of dealloyed Pt–Cu particles isnow established. Strain forms in Pt-enriched surface layers (shells)that are supported on an alloy particle core with a smaller latticeparameter. The compression in the shell modifies the d-band struc-ture of the Pt atoms, and thereby weakens the adsorption energy ofreactive intermediates compared to unstrained Pt and results in anincrease in the catalytic reactivity, consistent with DFT-based pre-dictions. A unique feature of the class of dealloyed catalysts is theexperimental control over the extent of dealloying (shell thickness)and the alloy core composition (the upper limit for strain in theshell). The noble and the non-noble constituents can be adjustedin the alloy precursor, such that both expansive and compressivestrain can be achieved to control the strengthening or weakeningof surface bonds. This enables continuous tuning of catalytic reac-tivity. We demonstrate explicitly such strain-related tuning for theORR, and this phenomenon will probably offer control over theactivity of other important electrocatalytic reactions that require

–4 –3 –2 –1 0

–0.02

0.00

0.02

0.04

0.06

0.08

0.10

0.12

Pt25Cu75

Pt25Cu75

Pt50Cu50

Pt50Cu50

Pt75Cu25

Pt75Cu25

Strain (ashell–aPt)/aPt (%)

Pt

OR

R a

ctiv

ity (

eV)

Figure 6 | Experimental and predicted relationships between

electrocatalytic ORR activity and lattice strain. The experimental ORR

activity (in units of kT ln( js,alloy/js,Pt), T¼ 298 K), based on relative surface

area, of two families of dealloyed Pt–Cu bimetallic core–shell nanoparticles

plotted as a function of strain, s(Pt), in the particle shell (red and blue

triangles denote dealloyed Pt–Cu precursors prepared at annealing

temperatures of 800 8C and 950 8C, respectively). The ORR activity is

proportional to the logarithm of the ratio of the ORR current density of the

dealloyed particles to that of pure Pt nanoparticles; this quantity is the

effective difference in ORR activation energies (‘reaction driving force’). The

experimental curves are upper bounds of the strain in the surface layer. The

dashed line is the DFT-predicted, volcano-shaped trend of the ORR activity

for a Pt(111) single-crystal slab under isotropic strain. Moderate compressive

lattice strain is predicted to enhance the rate of ORR catalysis. Absolute and

relative values are provided in Supplementary Table S1. Vertical error bars

refer to experimental uncertainty involved in measuring catalytic activities

from electrochemical systems; horizontal errors bars refer to errors involved

in obtaining ashell from experiments and modelling.

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modification of the adsorption energy of reactive intermediates,such as the electrooxidation of small organic molecules, includingethanol, methanol and related species.

MethodsSynthesis. Pt–Cu binary electrocatalyst precursors were synthesized using a liquidmetal–salt precursor impregnation method followed by freeze-drying and thermalannealing. This method was applied previously to the synthesis of binary25,47 andternary24 Pt alloys. The nanoparticle-catalyst precursors were prepared by addingappropriate stoichiometric amounts of a solid Cu precursor (Cu(NO3)2

.2.5H2O,Sigma-Aldrich) to weighted amounts of commercial Pt electrocatalyst powder, of30 weight per cent platinum nanoparticles supported on carbon with a high surfacearea (TEC10E30E from Tanaka Kikinzoku).

Electrochemical measurements. The voltammetric response of the electrocatalystswas measured first during the initial three cyclic voltammetry (CV) scans at100 mV s21 to obtain the initial rapid Cu-dissolution profiles. The catalysts werefurther pretreated using 200 CV scans between 0.05 V and 1.0 V at a scan rate of500 mV s21, during which a large amount of Cu was lost from the alloynanoparticles. Thereafter, the Pt electrochemical surface area (Pt-ECSA) wasdetermined by cycling the treated catalysts at 100 mV s21 between 0.05 V to 1.2 Vand integrating the Faradaic charge associated with stripping of underpotentiallydeposited hydrogen (Supplementary Fig. S1a). Pt-ECSA measurements using COstripping resulted in comparable values of the surface area.

Linear-sweep voltammetry (LSV) measurements were conducted under an oxygenatmosphere by sweeping the potential from 0.06 V anodically to the open-circuitpotential (�1.0 V) at a scan rate of 5 mV s21 (Supplementary Fig. S1b). The ORRactivities of the dealloyed, activated catalysts were corrected for mass-transportlimitation using equation (5) in Gasteiger et al.3. Specific activities for mass and surfacearea were established at 900 mV at room temperature (Supplementary Fig. S1c,d).

Electron microscopy (STEM and EDS). EDS imaging was carried out in STEMmode at 200 kV with a Philips CM200FEG equipped with an EDAX detector/pulseprocessor and an Emispec Vision system (Supplementary Fig. S2). Pt-L and Cu-Kelemental maps were extracted from the spectral data and Pt–Cu (blue and red,respectively) colour map overlays were produced (Fig. 1a,b). High-angle annulardark-field (HAADF) STEM images of individual Pt–Cu nanoparticles were recordedat sub-angstrom resolution using a JEOL 2200FS Cs-corrected STEM (CEOShexapole aberration-corrector) operated at 200 kV.

Anomalous X-ray diffraction. Synchrotron-based XRD was used to characterizePt-alloy electrocatalyst precursor powders as well as electrochemically treatedactivated catalyst films using X-ray energies from 8,900 eV, through the CuK-adsorption edge (8,979 eV) to 9,150 eV. Diffraction measurements wereconducted at the Stanford Synchrotron Radiation Lightsource (SSRL) beamline 2-1.A detailed description of the analysis of the AXRD results is provided in theSupplementary Information.

Computational methods. Computational analysis was carried out using DACAPO(ref. 48), a total-energy calculation code. All calculations were performed on afour-layer slab with a 2 × 2 unit cell. Full relaxation of the oxygen adsorbate andof the first two metal layers was allowed. For strained Pt slabs, uniform expansion(or contraction) of the Pt(111) lattice was allowed in all three Cartesian directions,and no corrections to the interlayer Pt distance were included. This model providesa reasonable representation of the compression found in the Pt-base metal alloysthat form the substrate of the Pt samples.

X-ray photoelectron, X-ray emission and X-ray absorption spectroscopy. XPS,XES and XAS measurements were performed in an ultrahigh vacuum end-station with abase pressure better than 10–10 torr at beamline 13-2 at SSRL, which contains anelliptically polarized undulator that allows control of the direction of the photon Evector about the propagation direction. An electron energy analyser (VG-ScientaSES-100 or R3000), mounted perpendicular to the incoming light, was used for theXPS measurements. This was also equipped with a partial-yield detector for X-rayabsorption measurements. Samples were mounted on a rotatable sample rod at grazingangle (�58 for the photon incidence angle) with respect to the incoming light. Theindependent rotation of the sample and the photon polarization (E vector) allowed forselection of arbitrary angles between the E vector and the sample surface and any choiceof detection angle with respect to the sample surface. Composition of the catalysts wasdetermined by measuring the ratio of Pt 4f to Cu 3p XPS intensities normalized to theirrespective subshell photoionization cross-sections49 for different photon energies. Thekinetic energy of the photoelectron defines the inelastic mean free path and the probingdepth of the analysis. We varied the photoelectron kinetic energy by changing theincident photon energy to obtain the composition at different probing depths (seeTable 1). The estimated probing depths are 0.6 nm, 1 nm, 1.5 nm, 1.8 nm and 7 nm atphoton energies of 250 eV, 620 eV, 1,130 eV, 1,480 eV and 8,000 eV, respectively.Further experimental details are provided in the Supplementary Information.

Received 27 October 2009; accepted 11 March 2010;published online 25 April 2010

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AcknowledgementsThis project was supported by the Department of Energy, Office of Basic Energy Sciences,under the auspices of the President’s Hydrogen Fuel Initiative. Acknowledgment is alsomade to the National Science Foundation (grant #729722) for partial support of thisresearch. P.S. acknowledges support from the Cluster of Excellence in Catalysis (UNICAT)funded by the German National Science Foundation (Deutsche Forschungsgemeinschaft)and managed by the Technical University Berlin, Germany. Portions of this research werecarried out at the Stanford Synchrotron Radiation Lightsource, a national user facilityoperated by Stanford University on behalf of the US Department of Energy, Office of BasicEnergy Sciences. Use of the Center for Nanoscale Materials was supported by the USDepartment of Energy, Office of Science, Office of Basic Energy Sciences, under contractNo. DE-AC02-06CH11357. We acknowledge computer time at the Laboratory ComputingResource Center (LCRC) at Argonne National Laboratory, the National Energy ResearchScientific Computing Center (NERSC) and the EMSL, a national scientific user facilitysponsored by the Department of Energy’s Office of Biological and Environmental Researchand located at Pacific Northwest National Laboratory. Microscopy research supported byORNL’s SHaRE User Program, which is sponsored by the Scientific User Facilities Division,Office of Basic Energy Sciences, US Department of Energy. The authors thank L. Petterssonfor reading the manuscript.

Author contributionsP.S., M.F.T., J.G. and A.N. designed the research and co-wrote the paper, S.K., T.A., K.M.,C.Y., Z.L., S.K., D.N. and H.O. performed the experiments and analysed the data, and J.G.performed the theoretical calculations.

Additional informationThe authors declare no competing financial interests. Supplementary informationaccompanies this paper at www.nature.com/naturechemistry. Reprints and permissioninformation is available online at http://npg.nature.com/reprintsandpermissions/.Correspondence and requests for materials should be addressed to P.S.

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