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Expansion of interatomic distances in platinum catalyst nanoparticles
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Acta Materialia 58 (2010) 836–845 www.elsevier.com/locate/actamat
Expansion of interatomic distances in platinum catalyst nanoparticles K. Du b,a, F. Ernst a,*, M.C. Pelsozy c, J. Barthel d, K. Tillmann d b
a Department of Materials Science and Engineering, Case Western Reserve University, Cleveland, OH 44106, USA Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China c Department of Chemical Engineering, Case Western Reserve University, Cleveland, OH 44106, USA d Ernst Ruska-Centrum fu¨r Mikroskopie und Spektroskopie mit Elektronen (ER-C) and Institut fu¨r Festko¨rperforschung, Forschungszentrum Ju¨lich GmbH, Ju¨lich, Germany
Received 4 September 2009; received in revised form 30 September 2009; accepted 30 September 2009 Available online 31 October 2009
Abstract We study the atomistic structure of Pt catalyst nanoparticles using HRTEM (high-resolution transmission electron microscopy). The particles exhibit a faceted, cubo-octahedral shape, extended planar defects, and mono-atomic surface steps. HRTEM imaging with negative spherical aberration yielded atomic-resolution images with a minimum of artifacts. Combining digital image processing, quantitative image analysis, and HRTEM image simulations to determine local variations of the spacing between neighboring Pt atom columns, we have found an expansion of the lattice parameter in the particle core and even larger, locally varying expansion of Pt–Pt next-neighbor distances at the particle surface. The latter likely originates from an amorphous oxide on the nanoparticle surface and/or dissolution of oxygen on subsurface sites. These structural features may significantly impact the catalytic activity of Pt nanoparticles. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Nanoparticles; Catalysts; Atomistic structure; Quantitative high-resolution transmission electron microscopy; Spherical-aberration-adjusted transmission electron microscopy
1. Introduction Platinum nanoparticles play an important role in catalysis. For example, they catalyze the dissociation of hydrogen and oxygen in proton-exchange membrane fuel cells (PEMFCs). In spite of their enormous technological significance, fundamental understanding of how such particles function in catalysis is still rather incomplete. Therefore, the catalytic properties of Pt nanoparticles and potential alternatives, for example Pt-alloy nanoparticles, are an important field of ongoing research and development. One major goal is to improve the specific catalyst activity and resistance to CO poisoning [1–4]. It has been reported that local variation of the Pt–Pt next-neighbor atom distance (i.e. elastic strain) considerably influences the catalytic activity of Pt-based nanoparticles [5–7]. For example, alloying Pt with elements *
Corresponding author. E-mail address: [email protected] (F. Ernst).
of smaller atomic radius (e.g. Co, Ni, Cr) significantly improves the catalytic behavior – particularly the oxygen reduction reaction – by reducing the Pt–Pt next-neighbor distances in their vicinity [8–11]. Another approach based on the same principle is to form nanoparticles with a core–shell structure. For example, Cr (-core) nanoparticles with a Pt shell introduce compressive biaxial elastic strain in the Pt shell (dilatative volumetric strain in the Cr core). These Cr–Pt core–shell nanoparticles also demonstrated an enhanced catalytic activity on oxide reduction, as well as a significantly improved resistance to CO poisoning [12]. Similar beneficial effects of elastic strain were also reported for Pd-shell core–shell particles [11,13]. Aiming to understand effects of elastic strain and alloying on the catalytic activity of metal nanoparticles, it is hence important to determine their strain state – particularly at the particle surface – and the spatial distribution of atomic species in the particles. For both purposes, a particularly powerful experimental method is HRTEM
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.09.061
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(high-resolution transmission electron microscopy). HRTEM can resolve individual atom columns in nanoparticles [14,15]. Under ideal imaging conditions, a one-to-one correlation would exist between the position of peaks in the image intensity and the projected location of atom columns. In reality, however, effects of dynamic (as opposed to kinematic) electron diffraction and aberrations of the electron-optical system usually introduce delocalization of object information and may displace the image intensity peaks against the true positions of the atom columns they represent [16–20]. Owing to the recent development of an electron-optic device denoted as “Cs corrector”, these artifacts can now be effectively reduced by adjusting the spherical aberration coefficient Cs of the objective lens [21,22]. This technology enables quantification of the strain state of nanoparticles with unprecedented accuracy by: (i) recording HRTEM images with Cs adjusted for minimum delocalization and high contrast, (ii) determining the displacements of atom columns against their regular positions, using either real space [23–25] or reciprocal space approaches [26,27], and (iii) determining local elastic strain by differentiating the displacement field with respect to the spatial coordinates. In the present work, we have used both conventional and spherical-aberration-corrected HRTEM to obtain atomic-resolution images of Pt nanoparticles. To determine the spatial distribution of elastic strain in these particles, we have employed algorithms of digital image processing described in earlier publications [19,24]. With the aid of highly realistic HRTEM image simulations, we have carefully examined the effect of the above-mentioned imaging artifacts on the results of our strain analysis.
HRTEM images were recorded with a spherical-aberration-adjusted Titan 80-300 (FEI) and a Tecnai F30 ST (FEI). Both microscopes were operated at 300 kV accelerating voltage. With the Titan 80-300, we recorded focus series of images. By adjusting Cs to an (absolutely) small negative value, we minimized the delocalization of intensity maxima against the true atom column positions and, at the same time, obtained high contrast with the intensity maxima appearing as sharp peaks [22]. During the experiment, Zemlin tableaus using an incident beam tilt up to 24 mrad were recorded from the amorphous area adjacent to the particle in order to precisely measure residual higher-order aberrations [29]. Based on these aberration measurements, the residual aberrations up to the 3rd order were corrected to magnitudes below the respective p4-limits required for delocalization-free phase-contrast imaging with a spatial resolution close to 0.1 nm [30]. In order to measure the sampling (digitization) rate of the imaging system, a reference image of a Si single-crystal was recorded under the same experimental conditions. From the resulting reference image, the sampling rate was determined to be (0.012826 ± 0.00003) nm/pixel. We have also measured the anisotropy of the sampling rate for this imaging mode, which could originate from a geometric distortion of the image caused by aberrations of the imaging system of the electron microscope (e.g. astigmatism of post-magnifying lenses). The relative variation of the sampling rate in different directions was found to be less than 0.2%. Accordingly, the anisotropy of sampling rate is negligible for our purpose.
2. Experimental methods
In order to determine the elastic strain state of nanoparticles, we evaluated their HRTEM images to obtain quantitative and highly precise information on the lateral positions of the atom columns constituting the particle. To prepare experimental HRTEM images for this analysis, we first reduced the background contrast consisting of the speckle pattern generated by the amorphous carbon support film and noise. The filtering operation corresponds to multiplying the (complex-valued) Fourier spectrum D(g) of the image by a filter function W(g), yielding the Fourier spectrum Dw(g) of the filtered image:
2.1. Transmission electron microscopy Our study includes two modifications of Pt nanoparticles: (i) the particles of a carbon-supported catalyst and (ii) unsupported particles. The carbon-supported catalyst is commercially available (Johnson Matthey Fuel Cells, West Chester, PA, USA). It contains 38–41 wt.% Pt supported on an XC72R membrane, consisting of amorphous globular carbon. The unsupported Pt nanoparticles were produced by a synthesis method developed by Tsurumi et al. [28]. For HRTEM imaging, both materials were dispersed in ethanol with the aid of ultrasound. Subsequently, a drop of the suspension was deposited onto a holey carbon film supported by a Cu grid. For imaging the atomistic structure of these particles by HRTEM, we searched the specimens for particles that were accidentally oriented for imaging in a <1 1 0> or <1 0 0> viewing direction. While it is not possible to find particles with <1 1 0> or <1 0 0> aligned exactly parallel to the viewing direction, we examined the power spectrum of the HRTEM particle image to confirm that the deviation is tolerably small and accounted for the residual deviation in our image simulations.
2.2. Determination of particle strain from HRTEM images
Dw ðgÞ ¼ W ðgÞ DðgÞ:
ð1Þ
The vector argument g denotes the location in twodimensional Fourier space. The mathematical form of the filter function is W ðgÞ ¼
P ðgÞ P bg ðgÞ ; P ðgÞ
ð2Þ
where P(g) is the (real-valued) power spectrum jD(g)j2 of the image and Pbg(g) is that of the background [31]. To obtain the denominator of (2), we proceeded as follows: the power spectrum P(g) of the image constitutes the sum of Pbg(g) and a component Pp(g) that represents the image
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of the nanoparticle to be extracted. Since the speckle pattern of the amorphous carbon and the noise are isotropic, Pbg(g) can be assumed to be rotation-symmetric, i.e. independent of the azimuth in Fourier space. Pp(g), in contrast, exclusively consists of highly-localized peaks. Assuming that their integrated intensity is small compared to the integrated intensity of Pbg(g), Pbg(g) can be obtained by rotation-averaging P(g), i.e. averaging the intensities of N power spectra Pi(g) obtained by rotating P(g) by Ni 2p about the center. The Wiener filter is superior to a Bragg filter or a “band-pass” filter, as the latter remove intensity from the Fourier spectrum of the image and, consequently, introduce artifacts, especially in images of interfaces or surfaces [32]. The effect of the Wiener filter on our analysis of particle strain is examined by a comparative analysis of simulated images of non-strained Pt nanoparticles further below. After noise reduction, the actual strain analysis was performed by the LADIA method [24]. This algorithm determines the displacement of actual column image positions versus the position of a reference lattice. From this information, we determined the local expansion/contraction of next-neighbor atom column distances in the Pt nanoparticles, i.e. the local lattice parameters in the <1 1 0> directions orthogonal to the <1 0 0> viewing direction. 2.3. Image simulations To enable detailed interpretation of experimental HRTEM images and to investigate potential artifacts in our method of strain analysis, we performed extensive HRTEM image simulations on atomistic models of cubooctahedral Pt particles with the experimentally observed ratio of {1 1 1} and {2 0 0} area. Aiming to obtain most realistic simulated images, the models included not only the atom coordinates of the Pt particle, but also a substrate of amorphous carbon with a thickness of 2.5 nm. In total, the models contained about 20,000 atoms. The atom coordinates were generated by self-written C++ code as follows. First, we set up a tetragonal simulation box centered on the origin of a Cartesian coordinate system with the z-direction corresponding to the viewing direction and the dimensions in x- and y-direction commensurate with the Pt lattice parameter, i.e. corresponding to integer multiples of the lattice parameter. To obtain the coordinates of Pt in a cubo-octahedral particle centered on the origin, we set up an fcc (face-centered cubic) lattice in the simulation box, populated it with Pt atoms, and discarded the atoms outside of a cubo-octahedron parameterized by two vectors with directions representing the normals of one of the {1 1 1} and one of the {2 0 0} facets and lengths representing the distance of the facets from the origin. Subsequently, we added an amorphous carbon film below the Pt particle. The coordinates of the carbon atoms in this film were obtained in two steps: first, we generated a 2.5 nm thick slab of carbon with the structure of diamond and a lattice parameter adjusted to yield the
experimental value for the density of amorphous carbon. Second, we randomly displaced the atom positions to obtain an amorphous structure. Image simulations were carried out using the multislice algorithm as implemented by the software package EMS [33]. The slicing was performed automatically by our C++ program, and the slice thickness was chosen small enough (0.196 nm) to ensure that in the region of the Pt particle the slices contained no more than a single layer of Pt atoms. In order to examine the effect of crystal tilt on the apparent position of atom columns in the HRTEM image, simulations were also performed with the [0 0 1] axis of the cubooctahedron inclined by 1° or 2° versus the z-axis of the simulation box. In contrast to the conventional approach, where the object tilt is introduced by inclining the propagator functions applied between the slices, we applied a more realistic approach by actually rotating the atom coordinates about the center of the simulation box prior to slicing. For the imaging part of the simulations, we used experimentally determined electron-optical parameters of the two electron microscopes employed in this study. For the Titan 80-300, the accelerating voltage was 300 kV, the spherical-aberration coefficient (Cs) was 0.013 mm, the focus spread was 3.3 nm (1/e – width of the focal distribution), and the convergence semi-angle of the primary beam was 0.2 mrad. For the Tecnai F30, the accelerating voltage was 300 kV, Cs was 1.2 mm, the focus spread was 4 nm, and the convergence semi-angle was 0.2 mrad. 3. Results 3.1. HRTEM imaging of Pt particles HRTEM images of both specimen types consistently reveal a strong tendency of the Pt nanoparticles to facet on {1 1 1} and {2 0 0} planes. Our observations of many Pt nanoparticles in different viewing directions [34] confirm that the idealized shape of the particles corresponds to a cubo-octahedron, faceted mainly on {1 1 1} planes and truncated on {2 0 0} planes. Fig. 1a was recorded in a viewing direction that corresponds to <1 1 0> in the Pt crystal. The particle exhibits well-developed, atomically flat {1 1 1} and {2 0 0} facets. The image also reveals a few mono-atomic surface steps, apparently oriented in the <1 1 0> viewing direction. The periodicity of the FCC lattice is interrupted by planar defects parallel to {1 1 1} planes. These are stacking faults and twin boundaries. The HRTEM image of another Pt nanoparticle in Fig. 1b includes further examples of internal twinning. Actually, this image reveals a complex system of three twins, which is remarkable as the long axis of the particle is no larger than 4 nm. Twin boundaries and stacking faults are common in platinum symbol nanoparticles. However, after observing several hundreds of nanoparticles, we did not find the twin configuration that is often observed in Au and Ag nanoparticles, where five twins intersect in a line [35–37].
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Fig. 1. HRTEM images of Pt particles, recorded with a Tecnai F30. The viewing direction corresponds to <1 1 0>. The particles contain planar defects parallel to {1 1 1} planes. These defects are stacking faults or twin boundaries.
Another remarkable feature of the HRTEM images in Fig. 1 is the intimate contact between the Pt particles and the globular carbon support. Instead of the very small contact area one would expect for a particle of a non-wetting metal on a carbon substrate, the image reveals an extended contact area, covering a significant fraction (10%) of the particle surface. This behavior may adversely impact catalytic performance as it reduces the active surface area of the particle and increases the length of the boundary line between the particle, the substrate, and the space around the particle. In order to determine the focus settings of individual images, we recorded focus series, i.e. series of images with a fixed focus increment between subsequent images. Fig. 2 shows a HRTEM image we have obtained with the Cs-corrected Titan 80-300 microscope at the ER-C. This image shows a Pt nanoparticle viewed in <1 0 0> projection. The particle appears to be attached to two further Pt nanoparticles, one at the bottom and one on the lower
Fig. 2. HRTEM image of a Pt nanoparticle in <1 0 0> projection, recorded with the Titan 80-300 at a focus value of 4 nm as part of a focus series. The particle appears to be attached to two other Pt nanoparticles (marked with dash lines).
right side. (In principle, the apparent proximity could also just be a projection effect, e.g. generated by one particle above the carbon support film and the other one below, but the fact that the particle images touch exactly at their periphery indicate that this is most likely not the case here.) The focus series were recorded starting at a focus of f 30 nm to 30 nm focus with a focal increment of 2 nm between subsequent images. In this publication, positive values for focus denote overfocus, i.e. the object is farther away from the lens than the conjugate plane of the image plane. First, we estimated the focus values of the individual images of the series from the focus setting yielding minimum-contrast from the carbon support film. According to a known rule of thumb, contrast minimum corresponds to about one half of the Scherzer focus, i.e. a focus of 2 nm for the Titan 80-300 microscope with the spherical aberration Cs set to 0.013 nm. In order to determine the focus values more accurately, we evaluated the contrast in the region of the Pt particle (the contrast is defined by the standard deviation of image intensity across the region [38]) for those images with estimated focus values between 18 nm and +20 nm and compared the experimental contrast values to the contrast in simulated images. Fig. 3 shows the contrast values of the experimental images as dots and the contrast values of the simulated images as a continuous line. The simulated images well reproduce the experimentally observed variation of contrast with increasing focus. Accordingly, the precise focus values of the experimental images can be determined by retrieving the focus values of the corresponding simulations in Fig. 3. Further, Fig. 3 indicates a quantitative mismatch between the absolute contrast in the simulated images and the corresponding experimental images. The fact that simulated images overestimate the contrast of experimental images is well known; the ratio of simulated to real contrast is denoted as the “Stobbs factor” [39,40]. Fig. 3 indicates a Stobbs factor 3. As far as this effect is understood to date, it can mostly be attributed to a number of effects that are
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Fig. 3. Contrast of experimental and simulated images for a series of focus values. The scale for the contrast of simulated images is shown on the left side of the figure, and that for experimental images is shown on the right side. The scale for simulated images is about 3 times of that for experimental images.
not properly taken into account in HRTEM image simulations: inelastic scattering of electrons, non-ideal modulation transfer function of the recording device, and lateral mechanical vibrations of the TEM specimen [41]. 3.2. Pt particle elastic strain distribution In order to determine the potential existence and spatial distribution of elastic strain in Pt nanoparticles, we have performed quantitative lattice distortion analysis on HRTEM images recorded in <1 0 0>. Applying the LADIA approach, we first determined the best estimate for the true positions of the intensity peaks in the HRTEM image of the particle. For this purpose, we fitted a cubic spline function [42] to the local experimental intensity distribution around each intensity peak in the image. Then, we fitted a twodimensional square reference lattice to the set of peak positions we obtained in the first step with the least squares function [43]. The fitting parameters were the reference lattice parameter, two parameters representing the (x, y) position
of the origin, and one parameter representing the orientation (rotation angle) of the reference lattice. After establishing the reference lattice, we determined the displacement vectors describing the shift of the experimental peak positions versus the corresponding lattice points of the reference lattice. The result of this procedure is a displacement field with an arbitrary constant term. By differentiating the displacement field with respect to the spatial coordinates, we eventually obtained a strain field. Provided that the intensity maxima in the experimental image truly represent the positions of corresponding atom columns in the particle (which we will confirm further below by analyzing simulated images), the strain field obtained in this way represents the <1 0 0>-projected average elastic strain in the atom column positions of the Pt nanoparticle. As the particular <1 0 0> viewing direction can be assumed to be equivalent to any other <1 0 0> direction in the particle, we may assume that the observed strain in the atom column positions represents corresponding isotropic volumetric strain. Fig. 4a shows the Pt nanoparticle of Fig. 2 in another image of the through-focus series, i.e. at a different focus setting. The conditions under which this image was obtained show less contrast in the two attached particles but better contrast in the particle viewed in <1 0 0> direction, suitable for elastic strain analysis. The result, obtained by performing the procedure described before and shown in Fig. 4b, reveals significant dilatative strain in the near-surface region of the Pt particle, i.e. an expansion of the local Pt–Pt next-neighbor distances compared to the distances observed near the center of the particle. However, expansion is only observed at the free surface and not where the particle contacts the two other particles. This observation confirms that the two nanoparticles are actually in contact, i.e. their proximity in the HRTEM image is not just a projection effect. Fig. 4b also indicates a few regions at the surface of the Pt particle–particularly at the surface steps – where the Pt–Pt next-neighbor distance actually appears to be smaller than in the center of the particle. As suggested by the
Fig. 4. (a) Experimental micrograph of the particle in Fig. 2, recorded with the Titan 80-300 at a different focus value of 10 nm (both images are part of a focus series). (b) Map of the elastic strain component measured from image intensity peaks along the [2 2 0] direction through the Pt particle.
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marked areas in Fig. 5, this could be associated with a reconstruction of the atomistic structure at the surface, leading to the formation of a 5-member polyhedron of Pt atom columns and – consequently – shortening of the Pt–Pt next-neighbor distances. We performed a detailed analysis of local elastic strain in the surface regions marked by A, B, and C in Fig. 4a. The result, displayed in Fig. 6, shows local Pt–Pt (column) next-neighbor distance expansions of 2% in the first two layers below the particle surface. 3.3. Lattice parameter expansion of Pt particles To investigate the lattice parameter of the Pt nanoparticles, i.e. the average value of next-neighbor distance of atom columns compared to the corresponding value in bulk Pt, we evaluated the Titan 80-300 HRTEM image
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of five Pt particles in Fig. 7. For each particle, positions of Fourier components (corresponding to Bragg peaks) were determined with a sub-pixel precision [44] from the power spectrum (an example is shown as an inset in Fig. 7). Table 1 presents the spacings dhkl of lattice planes {h k l} observed in the particles. To facilitate comparison of measured spacings, the table also the corresponding pincludes ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi lattice parameter d 100 ¼ d hkl h2 þ k 2 þ l2 that the measured spacing implies. The data reveal that the lattice of the Pt particles is significantly distorted. In particle A, for example, the lattice spacing d200 is 3.3% larger than d020. Similar findings were reported in Pd nanoparticles by Sun et al. [45]. In addition, the lattice spacing d ð1;1;1Þ is about 1.6% larger than d ð1;1;1Þ in particle B. This suggests that the distortion of the crystal structure is not limited to a tetragonal distortion.
Fig. 5. Regions of the HRTEM image of Fig. 4, zooming in on steps at the regions marked “1” and “2”. Apparently, the surface reconstructs to form polyhedral configurations of atoms.
Fig. 6. Elastic strain profiles from surface regions A, B and C in Fig. 4a along a <2 2 0> direction.
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Particle
(h k l)
dhkl (nm)
d 0100 (nm)
d 0100 a0 a0
A
(020) (2 0 0)
0.199 0.205
0.398 0.410
1.3 4.6
B
ð1 1 1Þ ð0 0 2Þ ð1 1 1Þ
0.235 0.200 0.231
0.407 0.401 0.401
3.8 2.1 2.1
C
ð2 2 0Þ ð0 2 2Þ ð2 0 2Þ
0.144 0.141 0.142
0.408 0.400 0.403
3.9 2.0 2.8
Fig. 1a may have profound implications for the catalytic activity of such particles. Pronounced faceting has been observed for nanoparticles of Pt and other metals by many researchers, and it is known that metal surfaces of different crystallographic orientation possess different catalytic activity. What has been discussed much less in the literature, however, is how the catalytic activity is influenced by extended structural defects in the particle interior. The particle image in Fig. 1 suggests that stacking faults and twin boundaries, where they intersect with the particle surface, cause notable surface corrugation. It seems likely that the line of structural disturbance around the particle introduced by such extended planar defects is a site of enhanced catalytic activity. The same goes for the mono-atomic steps that Fig. 1a exhibits in the {1 1 1}-oriented facets. If these conclusions are correct, the catalytic activity of corresponding nanoparticles might be enhanced by processing them to increase the formation probability of stacking faults, twin boundaries, and steps on the surface. The existence of such defects also has implications for the interpretation of HRTEM images. In HRTEM images, which project the object structure in a particular crystallographic viewing direction, stacking faults and twin boundaries inclined versus the viewing direction may give rise to imaging artifacts that manifest themselves as a broadening of the contrast features representing the atom columns – corresponding to the projected displacement vector of the respective defect. In order to eliminate artifacts from inclined extended structural defects, we have performed HRTEM image simulations of nanoparticles with inclined stacking faults and studied their effect on the resulting HRTEM image. Based on the results of these studies, we can rule out that our interpretation of the HRTEM image in Figs. 2 and 4a suffers from artifacts related to inclined planar defects.
D
ð3 1 1Þ ð3 1 1Þ
0.124 0.122
0.411 0.404
4.8 2.9
4.2. Reliability of atomistic structure determination
E
ð0 2 2Þ ð111Þ
0.143 0.230
0.406 0.398
3.4 1.5
Fig. 7. HRTEM image showing five Pt particles used for the measurement of their lattice parameters. The inset displays the power spectrum of particle B.
Table 1 Distance between crystal planes (h k l) measured from power spectra of five particles shown in Fig. 7. The different hypothetical (1 0 0) spacings d 0100 were calculated for the individual (h k l) spacings. (%)
Based on this analysis, we observe a significant expansion of the lattice parameter of the Pt nanoparticles compared to the lattice parameter of bulk Pt (a0 = 0.3923 nm [46]). The measured value is 1.3–4.8%. Although particles A and B are larger than particles C, D and E, there is no significant difference between their lattice parameters. Therefore, the lattice parameter expansion appears to be insensitive to the particle size in the observed regime of particle diameters less than 5 nm. 4. Discussion 4.1. Significance of nanoparticle defect structure The typical nanoparticle morphology and microstructure (structural defects) that become apparent from
In order to examine the influence of experimental imaging parameters and chosen image processing methods on the accuracy of the strain measurement, we applied the same procedure we had applied to our experimental images to corresponding simulated images. In particular, we applied the Wiener filter and LADIA strain analysis in exactly the same way as for the experimental images. The structure model we used for the simulation, however, did not include any strain. Fig. 8 shows an example of a corresponding simulated image. The focus was set to 10 nm, the same value as in the experimental image of Fig. 4. The analysis of the simulated image (Figs. 8b and 9) shows no detectable strain near the surface of the Pt particle–in striking contrast to the expanded Pt–Pt next-neighbor distance observed in the experimental image. Under the conditions we employed, accordingly, the HRTEM imaging process does not introduce artifacts that could be misinterpreted as significant local expansion or contraction of nextneighbor atom column distances. That the local variation
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Fig. 8. (a) Simulated HRTEM image of a Pt particle supported by a 2.5 nm-thick carbon film. The imaging conditions correspond to those of Fig. 4. (b) Elastic strain component measured from image intensity peaks along the [2 2 0] direction through the Pt particle.
Fig. 9. Elastic strain profiles from surface regions of the simulated particle in Fig. 8a along a <2 2 0> direction.
of strain we observe in the experimental images is real is further confirmed by the absence of strain in the region where the particle of Fig. 2 is attached to neighboring particles. In conclusion, the strain distributions shown in Figs. 4b and 6 are trustworthy and accurate. 4.3. Physical origin of strain One remarkable result of the present work is the expansion of the lattice parameter of 1.3–4.8% we observe comparing the nanoparticle core to bulk Pt. From some of the existing literature on metal nanoparticles [47–53], one might actually expect no expansion, but a contraction of the lattice parameter. In a continuum model, such contraction may be attributed to the increasing influence of surface tension with increasing surface-area-to-volume ratio, i.e. decreasing particle size [50]. Of course, as particle sizes approach atomistic dimensions, continuum models become less useful. Moreover, crystallographic effects on the particle shape (facetting) need to be taken into account [50]. Finally, there are also some known cases in which the lattice parameter of metal nanoparticles was observed to be expanded with regard to bulk material. In FePt nanoparticles, for example, the expansion can be as much as 9% [54].
One potential mechanism by which the lattice parameter in the core of a nanoparticle might expand is the presence of solute atoms on interstitial sites. For example, the expansion of the lattice parameters that Jacobs et al. observed in Pd nanoparticles was attributed to inward diffusion of oxygen atoms when the particles were exposed to air [55]. In principle, the particles we have studied may contain e.g. interstitially dissolved hydrogen or oxygen (expected to dissolve interstitially from atomic-size consideration [56]). Owing to small atomic size, dissolved hydrogen should not cause significant lattice parameter expansion. Dissolved oxygen, in contrast, could in principle lead to a measurable effect. However, different from the case of Pd, which dissolves relevant concentrations of oxygen (and hydrogen), the solubility of oxygen in (bulk) Pt is so small that dissolved oxygen can hardly explain the observed lattice parameter expansion. Another explanation for the lattice parameter expansion in the nanoparticle interior arises from the even larger expansion of Pt–Pt next-neighbor distances observed at the particle surface. The Pt particles are so small that introduction of misfit dislocations cannot be an energetically favorable mechanism for relaxing misfit stress (this would be a case of strong bonding where the shear strength of bonds across the lattice-mismatched interface is comparable to the shear modulus of the material on either side [57]). Further, since the volume is small compared to the surface area, supporting the misfit stress introduced by a layer with expanded next-neighbor distances covering the entire surface will likely lead to considerable strain in the particle interior. The remaining question is what actually causes the substantial expansion of Pt–Pt next-neighbor spacings at the particle surface. Close inspection of the HRTEM image in Fig. 4a reveals a distinct amorphous layer on the crystalline core of the particle, discernible from the speckle pattern of the ultra-thin carbon support film by somewhat higher contrast and somewhat lower average intensity. Fig. 10 shows the HRTEM image of the nanoparticle in
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Fig. 10. The HRTEM image of Fig. 4 after adjusting the brightness and contrast to increase the visibility of the amorphous layer on the particle surface (indicated by the markers at the top left).
Fig. 4 after adjusting brightness and contrast to increase the visibility of the amorphous layer. As indicated by the markers on the top left, the thickness of this layer roughly corresponds to one times the lattice parameter of Pt, i.e. 0.5 nm. This layer could be an early stage of a platinum oxide scale. Indeed, HRTEM image simulations we obtained from models including a corresponding coverage of the particle surface with amorphous oxide qualitatively reproduce the observed contrast. The increase of Pt–Pt next-neighbor spacings at the surface of the crystalline core of the particle as well as the observed surface reconstruction yielding 5-member polyhedra of Pt atom columns would then find their explanation in the formation of Pt–O–Pt bonds expanding the Pt–Pt distance. (The presence of oxygen cannot be directly observed in HRTEM images. Owing to the relatively small electron scattering potential of oxygen compared to that of Pt, the HRTEM image mainly shows Pt atoms, i.e. no considerable contrast is expected from oxygen atoms.) The 5-member polyhedra of Pt atom columns in Fig. 5 seem to indicate a soft structural transition between the ordered crystal structure of Pt and the amorphous oxide, similar to what has been observed at the interface between crystalline grains and glassy grain boundary films in Si3N4 ceramics [58,59]. The hypothesis of expanded Pt–Pt next-neighbor distances by formation of Pt–O bonds agrees well with recent findings of Imai et al., who have observed the formation of platinum oxides at the surface of platinum nanoparticles in situ using time-resolved hard X-ray XRD (X-ray diffractometry) and XAS (energy-dispersive X-ray absorption spectroscopy) [60]. Similarly, Su et al. [61] have observed an 10% lattice parameter expansion near the surface of Ag nanoparticles and related this expansion to the formation of silver oxide at the surface of the particles. However, in spite of the expansion of Pt–Pt distances at the particle surface, and in apparent contradiction to our observations, Imai et al. find the core of the particles to be totally relaxed, i.e. unstrained and exhibiting the lattice parameter of bulk
Pt. However, the oxide layers in that case were thicker, and their detailed morphology is not known (e.g. small islands versus a continuous scale). Unfortunately, XRD and XAS cannot image the atomistic structure of the oxide–metal interface to reveal how misfit stresses are accommodated. In addition to the formation of an amorphous oxide scale, oxygen may expand the Pt–Pt next-neighbor spacings at the nanoparticle surface by dissolving on subsurface sites. Recent observations of Sachs et al. [62] of hydrogen in Pd suggest that the solubility of interstitial solutes at the surface of nanoparticles may be considerably higher than it is in corresponding bulk material. In conclusion, the most realistic explanation for the observed increase of Pt–Pt next-neighbor spacings at the Pt nanoparticle surface and the increase of the lattice parameter in the nanoparticle core is that Pt atoms at the crystallite surface bond to oxygen and/or oxygen is dissolved on subsurface sites of the nanoparticles. The resulting increase of the Pt–Pt next-neighbor spacing introduces coherency stresses that cannot relax by the introduction of misfit dislocations and therefore have to be supported by the relatively small particle core, causing its lattice parameter to expand as well. According to the results of other studies, an expansion of the Pt–Pt next-neighbor atom distance (i.e. dilatative elastic strain) can have considerable influence on the catalytic activity of Pt-based nanoparticles [5–7]. The expansion of 2% we observe near the nanoparticle surface is expected to reduce the catalytic activity compared to that of corresponding surfaces of bulk Pt [8–11]. On the other hand, the regions near steps, where the Pt–Pt distance is smaller than in the particle center, may be highly catalytically active. It needs to be stressed, of course, that the details of the particle structure can (and will) change during service, as observed in the beautiful experiment by Imaiet al. [60]. 5. Conclusion Our study shows that HRTEM, and in particular quantitative evaluation of HRTEM images with adjusted spherical aberration can make a major contribution to understanding performance-critical structural aspects of catalyst nanoparticles. Compared to other techniques that have been applied for studying catalyst nanoparticles, such as XRD (X-ray diffractometry) or XAS (X-ray absorption spectroscopy), HRTEM has the great advantage of providing highly resolved, local information on individual nanoparticles. The observation of expanded inter-atomic spacings at the particle surface and surface corrugation caused by stacking faults, twin boundaries, and surface steps suggests that the impact of strain and extended structural defects on catalytic activity warrants further investigation. The complex correlation we observed between the spatial variation of inter-atomic spacings, surface scales, particle shape, and contact to other particles demonstrates that techniques that average over a large ensemble of particles may lead to non-realistic conclusions. A clear dis-
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advantage of HRTEM is its limitation to ex situ studies, i.e. studies performed not in the same environment and at the same temperature at which respective catalysts operate. A possible route to a deeper understanding could be to combine TEM studies with in situ techniques, such as XRD and XAS, in order to enable more detailed interpretation of XRD and XAS in situ data in terms of structural details of individual particles. Acknowledgments We thank Stephen Campbell for stimulating discussions. We acknowledge financial support from the Department of Energy (contract No. DE-FC36-03-GO13107) and from the Army Research Office (MURI grant DAAD19-03-10169). Kui Du also thanks the Natural Sciences Foundation of China for support under Grant No. 50871116. Further, we thank Maya Bar-Sadan, Lothar Houben, Markus Lentzen, Martina Luysberg, and Andreas Thust at the Ernst Ruska Centre (ER-C) in Ju¨lich, Germany, for outstanding support and collegiality. References [1] Antolini E. J Appl Electrochem 2004;34:563. [2] Mukerjee S, Urian RC, Lee SJ, Ticianelli EA, McBreen J. J Electrochem Soc 2004;151:A1094. [3] Gasteiger HA, Kocha SS, Sompalli B, Wagner FT. Appl Catal B 2005;56:9. [4] Raimondi F, Scherer GG, Kotz R, Wokaun A. Angew Chem Int Ed 2005;44:2190. [5] Mukerjee S, Srinivasan S, Soriaga MP, McBreen J. J Electrochem Soc 1995;142:1409. [6] Mavrikakis M, Hammer B, Norskov JK. Phys Rev Lett 1998;81:2819. [7] Schlapka A, Lischka M, Gro A, Kasberger U, Jakob P. Phys Rev Lett 2003;91:016101. [8] Uribe FA, Zawodzinski TA. Electrochim Acta 2002;47:3799. [9] Salgado JRC, Antolini E, Gonzalez ER. J Phys Chem B 2004;108:17767. [10] Antolini E, Salgado JRC, Giz MJ, Gonzalez ER. Int J Hydrogen Energy 2005;30:1213. [11] Bezerra CWB, Zhang L, Liu HS, Lee KC, Marques ALB, Marques EP, et al. J Pow Sour 2007;173:891. [12] Alayoglu S, Nilekar AU, Mavrikakis M, Eichhorn B. Nat Mater 2008;7:333. [13] Suo Y, Zhuang L, Lu J. Angew Chem Int Ed 2007;46:2862. [14] Wallenberg L, Bovin JO, Petford-Long A, Smith D. Ultramicroscopy 1986;20:71. [15] Hofmeister H, Tan GL, Dubiel M. J Mater Res 2005;20:1551. [16] Marks LD. Ultramicroscopy 1985;18:33. [17] Coene W, Jansen AJEM. Scan Microsc Suppl 1992;6:379. [18] Zandbergen HW, Tang D, Van Dyck D. Ultramicroscopy 1996;64:185. [19] Du K, Phillipp F. J Microsc 2006;221:63. [20] Crozier PA, Tsen SCY, Liu J, Cartes CL, Perez-Omi JA. J Electron Microsc 1999;48:1015. [21] Kabius B, Haider M, Uhlemann S, Schwan E, Urban K, Rose H. J Electron Microsc 2002;51:S51.
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[22] Jia CL, Lentzen M, Urban K. Microsc Microanal 2004;10:174. [23] Bierwolf R, Hohenstein M, Phillipp F, Brandt O, Crook GE, Ploog K. Ultramicroscopy 1993;49:273. [24] Du K, Rau Y, Jin-Phillipp NY, Phillipp F. J Mater Sci Technol 2002;18:135. [25] Houben L, Thust A, Urban K. Ultramicroscopy 2006;106:200. [26] Hytch MJ, Snoeck E, Kilaas R. Ultramicroscopy 1998;74:131. [27] Hytch MJ, Putaux JL, Penisson JM. Nature 2003;423:270. [28] Tsurumi K, Sugimoto H, Yamamoto N, Nakamura T, Stonehart P. US Patent 5,275,999; 1994. [29] Zemlin F, Weiss K, Schiske P, Kunath W, Herrmann KH. Ultramicroscopy 1978;3:49. [30] Uhlemann S, Haider M. Ultramicroscopy 1998;72:109. [31] Castleman KR. Digital image processing. Upper Saddle River: Prentice Hall; 1996. [32] Marks LD. Ultramicroscopy 1996;62:43. [33] Stadelmann PA. Ultramicroscopy 1987;21:131. [34] Pelsozy M. Investigation of hydrogen peroxide production and transport in a proton exchange membrane fuel cell and the atom resolved micro-characterization of its catalyst. Ph.d. thesis, Case Western Reserve University; 2008. [35] Marks LD. Philos Mag A 1984;49:81. [36] Hall BD, Flueli M, Monot R, Borel JP. Phys Rev B 1991;43: 3906. [37] Buffat PA, Flueli M, Spycher R, Stadelmann P, Borel JP. Faraday Discuss 1991;92:173. [38] Smith DJ. Rep Prog Phys 1997;60:1513. [39] Hytch MJ, Stobbs WM. Ultramicroscopy 1994;53:191. [40] Thust A. Phys Rev Lett 2009;102:220801. [41] Du K, von Hochmeister K, Phillipp F. Ultramicroscopy 2007;107:281. [42] PV-WAVE 7.00. Visual Numerics Inc., Houston (TX); 2007. [43] Press WH, Teukolsky SA, Vetterling WT, Flannery BP. Numerical recipes in FORTRAN: the art of scientific computing. 2nd ed. Cambridge: Cambridge University Press; 1992. [44] de Ruijter WJ, Sharma R, McCartney MR, Smith DJ. Ultramicroscopy 1995;57:409. [45] Sun Y, Frenkel AI, Isseroff R, Shonbrun C, Forman M, Shin KW, et al. Langmuir 2006;22:807. [46] Swanson H, Tatge E. Natl Bur Stand (US) Circ 1953;1:1. 539. [47] Apai G, Hamilton JF, Stohr J, Thompson A. Phys Rev Lett 1979;43:165. [48] Zubov VI. Nanostruct Mater 1995;5:571. [49] Lamber R, Wetjen S, Jaeger NI. Phys Rev B 1995;51:10968. [50] Qi WH, Wang MP. J Nanoparticle Res 2005;7:51. [51] Qi WH, Wang MP, Su YC. J Mater Sci Lett 2002;21:877. [52] Huang Z, Thomson P, Di S. J Phys Chem Solids 2007;68:530. [53] Nanda KK, Maisels A, Kruis FE. J Phys Chem C 2008;112: 13488. [54] Wang RM, Dmitrieva O, Farle M, Dumpich G, Ye HQ, Poppa H, et al. Phys Rev Lett 2008;100:017205. [55] Jacobs JWM, Schryvers D. J Catal 1987;103:436. [56] Velho L, Bartlett R. Metall Trans 1972;3:65. [57] Ernst F, Raj R, Ruhle M. Z Metall 1999;90:961. [58] Ernst F, Kienzle O, Ruhle M. J Eur Ceram Soc 1999;19:665. [59] Kleebe HJ, Cinibulk MK, Cannon RM, Ruhle M. J Am Ceram Soc 1993;76:1969. [60] Imai H, Izumi K, Matsumoto M, Kubo Y, Kato K, Imai Y. J Am Chem Soc 2009;131:6293. [61] Su DS, Jacob T, Hansen TW, Wang D, Schlogl R, Freitag B, et al. Angew Chem Int Ed 2008;47:5005. [62] Sachs C, Pundt A, Kirchheim R, Winter M, Reetz MT, Fritsch D. Phys Rev B 2001;64:075408.
Acta Materialia 58 (2010) 836–845 www.elsevier.com/locate/actamat
Expansion of interatomic distances in platinum catalyst nanoparticles K. Du b,a, F. Ernst a,*, M.C. Pelsozy c, J. Barthel d, K. Tillmann d b
a Department of Materials Science and Engineering, Case Western Reserve University, Cleveland, OH 44106, USA Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China c Department of Chemical Engineering, Case Western Reserve University, Cleveland, OH 44106, USA d Ernst Ruska-Centrum fu¨r Mikroskopie und Spektroskopie mit Elektronen (ER-C) and Institut fu¨r Festko¨rperforschung, Forschungszentrum Ju¨lich GmbH, Ju¨lich, Germany
Received 4 September 2009; received in revised form 30 September 2009; accepted 30 September 2009 Available online 31 October 2009
Abstract We study the atomistic structure of Pt catalyst nanoparticles using HRTEM (high-resolution transmission electron microscopy). The particles exhibit a faceted, cubo-octahedral shape, extended planar defects, and mono-atomic surface steps. HRTEM imaging with negative spherical aberration yielded atomic-resolution images with a minimum of artifacts. Combining digital image processing, quantitative image analysis, and HRTEM image simulations to determine local variations of the spacing between neighboring Pt atom columns, we have found an expansion of the lattice parameter in the particle core and even larger, locally varying expansion of Pt–Pt next-neighbor distances at the particle surface. The latter likely originates from an amorphous oxide on the nanoparticle surface and/or dissolution of oxygen on subsurface sites. These structural features may significantly impact the catalytic activity of Pt nanoparticles. Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Nanoparticles; Catalysts; Atomistic structure; Quantitative high-resolution transmission electron microscopy; Spherical-aberration-adjusted transmission electron microscopy
1. Introduction Platinum nanoparticles play an important role in catalysis. For example, they catalyze the dissociation of hydrogen and oxygen in proton-exchange membrane fuel cells (PEMFCs). In spite of their enormous technological significance, fundamental understanding of how such particles function in catalysis is still rather incomplete. Therefore, the catalytic properties of Pt nanoparticles and potential alternatives, for example Pt-alloy nanoparticles, are an important field of ongoing research and development. One major goal is to improve the specific catalyst activity and resistance to CO poisoning [1–4]. It has been reported that local variation of the Pt–Pt next-neighbor atom distance (i.e. elastic strain) considerably influences the catalytic activity of Pt-based nanoparticles [5–7]. For example, alloying Pt with elements *
Corresponding author. E-mail address: [email protected] (F. Ernst).
of smaller atomic radius (e.g. Co, Ni, Cr) significantly improves the catalytic behavior – particularly the oxygen reduction reaction – by reducing the Pt–Pt next-neighbor distances in their vicinity [8–11]. Another approach based on the same principle is to form nanoparticles with a core–shell structure. For example, Cr (-core) nanoparticles with a Pt shell introduce compressive biaxial elastic strain in the Pt shell (dilatative volumetric strain in the Cr core). These Cr–Pt core–shell nanoparticles also demonstrated an enhanced catalytic activity on oxide reduction, as well as a significantly improved resistance to CO poisoning [12]. Similar beneficial effects of elastic strain were also reported for Pd-shell core–shell particles [11,13]. Aiming to understand effects of elastic strain and alloying on the catalytic activity of metal nanoparticles, it is hence important to determine their strain state – particularly at the particle surface – and the spatial distribution of atomic species in the particles. For both purposes, a particularly powerful experimental method is HRTEM
1359-6454/$36.00 Ó 2009 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2009.09.061
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(high-resolution transmission electron microscopy). HRTEM can resolve individual atom columns in nanoparticles [14,15]. Under ideal imaging conditions, a one-to-one correlation would exist between the position of peaks in the image intensity and the projected location of atom columns. In reality, however, effects of dynamic (as opposed to kinematic) electron diffraction and aberrations of the electron-optical system usually introduce delocalization of object information and may displace the image intensity peaks against the true positions of the atom columns they represent [16–20]. Owing to the recent development of an electron-optic device denoted as “Cs corrector”, these artifacts can now be effectively reduced by adjusting the spherical aberration coefficient Cs of the objective lens [21,22]. This technology enables quantification of the strain state of nanoparticles with unprecedented accuracy by: (i) recording HRTEM images with Cs adjusted for minimum delocalization and high contrast, (ii) determining the displacements of atom columns against their regular positions, using either real space [23–25] or reciprocal space approaches [26,27], and (iii) determining local elastic strain by differentiating the displacement field with respect to the spatial coordinates. In the present work, we have used both conventional and spherical-aberration-corrected HRTEM to obtain atomic-resolution images of Pt nanoparticles. To determine the spatial distribution of elastic strain in these particles, we have employed algorithms of digital image processing described in earlier publications [19,24]. With the aid of highly realistic HRTEM image simulations, we have carefully examined the effect of the above-mentioned imaging artifacts on the results of our strain analysis.
HRTEM images were recorded with a spherical-aberration-adjusted Titan 80-300 (FEI) and a Tecnai F30 ST (FEI). Both microscopes were operated at 300 kV accelerating voltage. With the Titan 80-300, we recorded focus series of images. By adjusting Cs to an (absolutely) small negative value, we minimized the delocalization of intensity maxima against the true atom column positions and, at the same time, obtained high contrast with the intensity maxima appearing as sharp peaks [22]. During the experiment, Zemlin tableaus using an incident beam tilt up to 24 mrad were recorded from the amorphous area adjacent to the particle in order to precisely measure residual higher-order aberrations [29]. Based on these aberration measurements, the residual aberrations up to the 3rd order were corrected to magnitudes below the respective p4-limits required for delocalization-free phase-contrast imaging with a spatial resolution close to 0.1 nm [30]. In order to measure the sampling (digitization) rate of the imaging system, a reference image of a Si single-crystal was recorded under the same experimental conditions. From the resulting reference image, the sampling rate was determined to be (0.012826 ± 0.00003) nm/pixel. We have also measured the anisotropy of the sampling rate for this imaging mode, which could originate from a geometric distortion of the image caused by aberrations of the imaging system of the electron microscope (e.g. astigmatism of post-magnifying lenses). The relative variation of the sampling rate in different directions was found to be less than 0.2%. Accordingly, the anisotropy of sampling rate is negligible for our purpose.
2. Experimental methods
In order to determine the elastic strain state of nanoparticles, we evaluated their HRTEM images to obtain quantitative and highly precise information on the lateral positions of the atom columns constituting the particle. To prepare experimental HRTEM images for this analysis, we first reduced the background contrast consisting of the speckle pattern generated by the amorphous carbon support film and noise. The filtering operation corresponds to multiplying the (complex-valued) Fourier spectrum D(g) of the image by a filter function W(g), yielding the Fourier spectrum Dw(g) of the filtered image:
2.1. Transmission electron microscopy Our study includes two modifications of Pt nanoparticles: (i) the particles of a carbon-supported catalyst and (ii) unsupported particles. The carbon-supported catalyst is commercially available (Johnson Matthey Fuel Cells, West Chester, PA, USA). It contains 38–41 wt.% Pt supported on an XC72R membrane, consisting of amorphous globular carbon. The unsupported Pt nanoparticles were produced by a synthesis method developed by Tsurumi et al. [28]. For HRTEM imaging, both materials were dispersed in ethanol with the aid of ultrasound. Subsequently, a drop of the suspension was deposited onto a holey carbon film supported by a Cu grid. For imaging the atomistic structure of these particles by HRTEM, we searched the specimens for particles that were accidentally oriented for imaging in a <1 1 0> or <1 0 0> viewing direction. While it is not possible to find particles with <1 1 0> or <1 0 0> aligned exactly parallel to the viewing direction, we examined the power spectrum of the HRTEM particle image to confirm that the deviation is tolerably small and accounted for the residual deviation in our image simulations.
2.2. Determination of particle strain from HRTEM images
Dw ðgÞ ¼ W ðgÞ DðgÞ:
ð1Þ
The vector argument g denotes the location in twodimensional Fourier space. The mathematical form of the filter function is W ðgÞ ¼
P ðgÞ P bg ðgÞ ; P ðgÞ
ð2Þ
where P(g) is the (real-valued) power spectrum jD(g)j2 of the image and Pbg(g) is that of the background [31]. To obtain the denominator of (2), we proceeded as follows: the power spectrum P(g) of the image constitutes the sum of Pbg(g) and a component Pp(g) that represents the image
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of the nanoparticle to be extracted. Since the speckle pattern of the amorphous carbon and the noise are isotropic, Pbg(g) can be assumed to be rotation-symmetric, i.e. independent of the azimuth in Fourier space. Pp(g), in contrast, exclusively consists of highly-localized peaks. Assuming that their integrated intensity is small compared to the integrated intensity of Pbg(g), Pbg(g) can be obtained by rotation-averaging P(g), i.e. averaging the intensities of N power spectra Pi(g) obtained by rotating P(g) by Ni 2p about the center. The Wiener filter is superior to a Bragg filter or a “band-pass” filter, as the latter remove intensity from the Fourier spectrum of the image and, consequently, introduce artifacts, especially in images of interfaces or surfaces [32]. The effect of the Wiener filter on our analysis of particle strain is examined by a comparative analysis of simulated images of non-strained Pt nanoparticles further below. After noise reduction, the actual strain analysis was performed by the LADIA method [24]. This algorithm determines the displacement of actual column image positions versus the position of a reference lattice. From this information, we determined the local expansion/contraction of next-neighbor atom column distances in the Pt nanoparticles, i.e. the local lattice parameters in the <1 1 0> directions orthogonal to the <1 0 0> viewing direction. 2.3. Image simulations To enable detailed interpretation of experimental HRTEM images and to investigate potential artifacts in our method of strain analysis, we performed extensive HRTEM image simulations on atomistic models of cubooctahedral Pt particles with the experimentally observed ratio of {1 1 1} and {2 0 0} area. Aiming to obtain most realistic simulated images, the models included not only the atom coordinates of the Pt particle, but also a substrate of amorphous carbon with a thickness of 2.5 nm. In total, the models contained about 20,000 atoms. The atom coordinates were generated by self-written C++ code as follows. First, we set up a tetragonal simulation box centered on the origin of a Cartesian coordinate system with the z-direction corresponding to the viewing direction and the dimensions in x- and y-direction commensurate with the Pt lattice parameter, i.e. corresponding to integer multiples of the lattice parameter. To obtain the coordinates of Pt in a cubo-octahedral particle centered on the origin, we set up an fcc (face-centered cubic) lattice in the simulation box, populated it with Pt atoms, and discarded the atoms outside of a cubo-octahedron parameterized by two vectors with directions representing the normals of one of the {1 1 1} and one of the {2 0 0} facets and lengths representing the distance of the facets from the origin. Subsequently, we added an amorphous carbon film below the Pt particle. The coordinates of the carbon atoms in this film were obtained in two steps: first, we generated a 2.5 nm thick slab of carbon with the structure of diamond and a lattice parameter adjusted to yield the
experimental value for the density of amorphous carbon. Second, we randomly displaced the atom positions to obtain an amorphous structure. Image simulations were carried out using the multislice algorithm as implemented by the software package EMS [33]. The slicing was performed automatically by our C++ program, and the slice thickness was chosen small enough (0.196 nm) to ensure that in the region of the Pt particle the slices contained no more than a single layer of Pt atoms. In order to examine the effect of crystal tilt on the apparent position of atom columns in the HRTEM image, simulations were also performed with the [0 0 1] axis of the cubooctahedron inclined by 1° or 2° versus the z-axis of the simulation box. In contrast to the conventional approach, where the object tilt is introduced by inclining the propagator functions applied between the slices, we applied a more realistic approach by actually rotating the atom coordinates about the center of the simulation box prior to slicing. For the imaging part of the simulations, we used experimentally determined electron-optical parameters of the two electron microscopes employed in this study. For the Titan 80-300, the accelerating voltage was 300 kV, the spherical-aberration coefficient (Cs) was 0.013 mm, the focus spread was 3.3 nm (1/e – width of the focal distribution), and the convergence semi-angle of the primary beam was 0.2 mrad. For the Tecnai F30, the accelerating voltage was 300 kV, Cs was 1.2 mm, the focus spread was 4 nm, and the convergence semi-angle was 0.2 mrad. 3. Results 3.1. HRTEM imaging of Pt particles HRTEM images of both specimen types consistently reveal a strong tendency of the Pt nanoparticles to facet on {1 1 1} and {2 0 0} planes. Our observations of many Pt nanoparticles in different viewing directions [34] confirm that the idealized shape of the particles corresponds to a cubo-octahedron, faceted mainly on {1 1 1} planes and truncated on {2 0 0} planes. Fig. 1a was recorded in a viewing direction that corresponds to <1 1 0> in the Pt crystal. The particle exhibits well-developed, atomically flat {1 1 1} and {2 0 0} facets. The image also reveals a few mono-atomic surface steps, apparently oriented in the <1 1 0> viewing direction. The periodicity of the FCC lattice is interrupted by planar defects parallel to {1 1 1} planes. These are stacking faults and twin boundaries. The HRTEM image of another Pt nanoparticle in Fig. 1b includes further examples of internal twinning. Actually, this image reveals a complex system of three twins, which is remarkable as the long axis of the particle is no larger than 4 nm. Twin boundaries and stacking faults are common in platinum symbol nanoparticles. However, after observing several hundreds of nanoparticles, we did not find the twin configuration that is often observed in Au and Ag nanoparticles, where five twins intersect in a line [35–37].
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Fig. 1. HRTEM images of Pt particles, recorded with a Tecnai F30. The viewing direction corresponds to <1 1 0>. The particles contain planar defects parallel to {1 1 1} planes. These defects are stacking faults or twin boundaries.
Another remarkable feature of the HRTEM images in Fig. 1 is the intimate contact between the Pt particles and the globular carbon support. Instead of the very small contact area one would expect for a particle of a non-wetting metal on a carbon substrate, the image reveals an extended contact area, covering a significant fraction (10%) of the particle surface. This behavior may adversely impact catalytic performance as it reduces the active surface area of the particle and increases the length of the boundary line between the particle, the substrate, and the space around the particle. In order to determine the focus settings of individual images, we recorded focus series, i.e. series of images with a fixed focus increment between subsequent images. Fig. 2 shows a HRTEM image we have obtained with the Cs-corrected Titan 80-300 microscope at the ER-C. This image shows a Pt nanoparticle viewed in <1 0 0> projection. The particle appears to be attached to two further Pt nanoparticles, one at the bottom and one on the lower
Fig. 2. HRTEM image of a Pt nanoparticle in <1 0 0> projection, recorded with the Titan 80-300 at a focus value of 4 nm as part of a focus series. The particle appears to be attached to two other Pt nanoparticles (marked with dash lines).
right side. (In principle, the apparent proximity could also just be a projection effect, e.g. generated by one particle above the carbon support film and the other one below, but the fact that the particle images touch exactly at their periphery indicate that this is most likely not the case here.) The focus series were recorded starting at a focus of f 30 nm to 30 nm focus with a focal increment of 2 nm between subsequent images. In this publication, positive values for focus denote overfocus, i.e. the object is farther away from the lens than the conjugate plane of the image plane. First, we estimated the focus values of the individual images of the series from the focus setting yielding minimum-contrast from the carbon support film. According to a known rule of thumb, contrast minimum corresponds to about one half of the Scherzer focus, i.e. a focus of 2 nm for the Titan 80-300 microscope with the spherical aberration Cs set to 0.013 nm. In order to determine the focus values more accurately, we evaluated the contrast in the region of the Pt particle (the contrast is defined by the standard deviation of image intensity across the region [38]) for those images with estimated focus values between 18 nm and +20 nm and compared the experimental contrast values to the contrast in simulated images. Fig. 3 shows the contrast values of the experimental images as dots and the contrast values of the simulated images as a continuous line. The simulated images well reproduce the experimentally observed variation of contrast with increasing focus. Accordingly, the precise focus values of the experimental images can be determined by retrieving the focus values of the corresponding simulations in Fig. 3. Further, Fig. 3 indicates a quantitative mismatch between the absolute contrast in the simulated images and the corresponding experimental images. The fact that simulated images overestimate the contrast of experimental images is well known; the ratio of simulated to real contrast is denoted as the “Stobbs factor” [39,40]. Fig. 3 indicates a Stobbs factor 3. As far as this effect is understood to date, it can mostly be attributed to a number of effects that are
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Fig. 3. Contrast of experimental and simulated images for a series of focus values. The scale for the contrast of simulated images is shown on the left side of the figure, and that for experimental images is shown on the right side. The scale for simulated images is about 3 times of that for experimental images.
not properly taken into account in HRTEM image simulations: inelastic scattering of electrons, non-ideal modulation transfer function of the recording device, and lateral mechanical vibrations of the TEM specimen [41]. 3.2. Pt particle elastic strain distribution In order to determine the potential existence and spatial distribution of elastic strain in Pt nanoparticles, we have performed quantitative lattice distortion analysis on HRTEM images recorded in <1 0 0>. Applying the LADIA approach, we first determined the best estimate for the true positions of the intensity peaks in the HRTEM image of the particle. For this purpose, we fitted a cubic spline function [42] to the local experimental intensity distribution around each intensity peak in the image. Then, we fitted a twodimensional square reference lattice to the set of peak positions we obtained in the first step with the least squares function [43]. The fitting parameters were the reference lattice parameter, two parameters representing the (x, y) position
of the origin, and one parameter representing the orientation (rotation angle) of the reference lattice. After establishing the reference lattice, we determined the displacement vectors describing the shift of the experimental peak positions versus the corresponding lattice points of the reference lattice. The result of this procedure is a displacement field with an arbitrary constant term. By differentiating the displacement field with respect to the spatial coordinates, we eventually obtained a strain field. Provided that the intensity maxima in the experimental image truly represent the positions of corresponding atom columns in the particle (which we will confirm further below by analyzing simulated images), the strain field obtained in this way represents the <1 0 0>-projected average elastic strain in the atom column positions of the Pt nanoparticle. As the particular <1 0 0> viewing direction can be assumed to be equivalent to any other <1 0 0> direction in the particle, we may assume that the observed strain in the atom column positions represents corresponding isotropic volumetric strain. Fig. 4a shows the Pt nanoparticle of Fig. 2 in another image of the through-focus series, i.e. at a different focus setting. The conditions under which this image was obtained show less contrast in the two attached particles but better contrast in the particle viewed in <1 0 0> direction, suitable for elastic strain analysis. The result, obtained by performing the procedure described before and shown in Fig. 4b, reveals significant dilatative strain in the near-surface region of the Pt particle, i.e. an expansion of the local Pt–Pt next-neighbor distances compared to the distances observed near the center of the particle. However, expansion is only observed at the free surface and not where the particle contacts the two other particles. This observation confirms that the two nanoparticles are actually in contact, i.e. their proximity in the HRTEM image is not just a projection effect. Fig. 4b also indicates a few regions at the surface of the Pt particle–particularly at the surface steps – where the Pt–Pt next-neighbor distance actually appears to be smaller than in the center of the particle. As suggested by the
Fig. 4. (a) Experimental micrograph of the particle in Fig. 2, recorded with the Titan 80-300 at a different focus value of 10 nm (both images are part of a focus series). (b) Map of the elastic strain component measured from image intensity peaks along the [2 2 0] direction through the Pt particle.
K. Du et al. / Acta Materialia 58 (2010) 836–845
marked areas in Fig. 5, this could be associated with a reconstruction of the atomistic structure at the surface, leading to the formation of a 5-member polyhedron of Pt atom columns and – consequently – shortening of the Pt–Pt next-neighbor distances. We performed a detailed analysis of local elastic strain in the surface regions marked by A, B, and C in Fig. 4a. The result, displayed in Fig. 6, shows local Pt–Pt (column) next-neighbor distance expansions of 2% in the first two layers below the particle surface. 3.3. Lattice parameter expansion of Pt particles To investigate the lattice parameter of the Pt nanoparticles, i.e. the average value of next-neighbor distance of atom columns compared to the corresponding value in bulk Pt, we evaluated the Titan 80-300 HRTEM image
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of five Pt particles in Fig. 7. For each particle, positions of Fourier components (corresponding to Bragg peaks) were determined with a sub-pixel precision [44] from the power spectrum (an example is shown as an inset in Fig. 7). Table 1 presents the spacings dhkl of lattice planes {h k l} observed in the particles. To facilitate comparison of measured spacings, the table also the corresponding pincludes ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi lattice parameter d 100 ¼ d hkl h2 þ k 2 þ l2 that the measured spacing implies. The data reveal that the lattice of the Pt particles is significantly distorted. In particle A, for example, the lattice spacing d200 is 3.3% larger than d020. Similar findings were reported in Pd nanoparticles by Sun et al. [45]. In addition, the lattice spacing d ð1;1;1Þ is about 1.6% larger than d ð1;1;1Þ in particle B. This suggests that the distortion of the crystal structure is not limited to a tetragonal distortion.
Fig. 5. Regions of the HRTEM image of Fig. 4, zooming in on steps at the regions marked “1” and “2”. Apparently, the surface reconstructs to form polyhedral configurations of atoms.
Fig. 6. Elastic strain profiles from surface regions A, B and C in Fig. 4a along a <2 2 0> direction.
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Particle
(h k l)
dhkl (nm)
d 0100 (nm)
d 0100 a0 a0
A
(020) (2 0 0)
0.199 0.205
0.398 0.410
1.3 4.6
B
ð1 1 1Þ ð0 0 2Þ ð1 1 1Þ
0.235 0.200 0.231
0.407 0.401 0.401
3.8 2.1 2.1
C
ð2 2 0Þ ð0 2 2Þ ð2 0 2Þ
0.144 0.141 0.142
0.408 0.400 0.403
3.9 2.0 2.8
Fig. 1a may have profound implications for the catalytic activity of such particles. Pronounced faceting has been observed for nanoparticles of Pt and other metals by many researchers, and it is known that metal surfaces of different crystallographic orientation possess different catalytic activity. What has been discussed much less in the literature, however, is how the catalytic activity is influenced by extended structural defects in the particle interior. The particle image in Fig. 1 suggests that stacking faults and twin boundaries, where they intersect with the particle surface, cause notable surface corrugation. It seems likely that the line of structural disturbance around the particle introduced by such extended planar defects is a site of enhanced catalytic activity. The same goes for the mono-atomic steps that Fig. 1a exhibits in the {1 1 1}-oriented facets. If these conclusions are correct, the catalytic activity of corresponding nanoparticles might be enhanced by processing them to increase the formation probability of stacking faults, twin boundaries, and steps on the surface. The existence of such defects also has implications for the interpretation of HRTEM images. In HRTEM images, which project the object structure in a particular crystallographic viewing direction, stacking faults and twin boundaries inclined versus the viewing direction may give rise to imaging artifacts that manifest themselves as a broadening of the contrast features representing the atom columns – corresponding to the projected displacement vector of the respective defect. In order to eliminate artifacts from inclined extended structural defects, we have performed HRTEM image simulations of nanoparticles with inclined stacking faults and studied their effect on the resulting HRTEM image. Based on the results of these studies, we can rule out that our interpretation of the HRTEM image in Figs. 2 and 4a suffers from artifacts related to inclined planar defects.
D
ð3 1 1Þ ð3 1 1Þ
0.124 0.122
0.411 0.404
4.8 2.9
4.2. Reliability of atomistic structure determination
E
ð0 2 2Þ ð111Þ
0.143 0.230
0.406 0.398
3.4 1.5
Fig. 7. HRTEM image showing five Pt particles used for the measurement of their lattice parameters. The inset displays the power spectrum of particle B.
Table 1 Distance between crystal planes (h k l) measured from power spectra of five particles shown in Fig. 7. The different hypothetical (1 0 0) spacings d 0100 were calculated for the individual (h k l) spacings. (%)
Based on this analysis, we observe a significant expansion of the lattice parameter of the Pt nanoparticles compared to the lattice parameter of bulk Pt (a0 = 0.3923 nm [46]). The measured value is 1.3–4.8%. Although particles A and B are larger than particles C, D and E, there is no significant difference between their lattice parameters. Therefore, the lattice parameter expansion appears to be insensitive to the particle size in the observed regime of particle diameters less than 5 nm. 4. Discussion 4.1. Significance of nanoparticle defect structure The typical nanoparticle morphology and microstructure (structural defects) that become apparent from
In order to examine the influence of experimental imaging parameters and chosen image processing methods on the accuracy of the strain measurement, we applied the same procedure we had applied to our experimental images to corresponding simulated images. In particular, we applied the Wiener filter and LADIA strain analysis in exactly the same way as for the experimental images. The structure model we used for the simulation, however, did not include any strain. Fig. 8 shows an example of a corresponding simulated image. The focus was set to 10 nm, the same value as in the experimental image of Fig. 4. The analysis of the simulated image (Figs. 8b and 9) shows no detectable strain near the surface of the Pt particle–in striking contrast to the expanded Pt–Pt next-neighbor distance observed in the experimental image. Under the conditions we employed, accordingly, the HRTEM imaging process does not introduce artifacts that could be misinterpreted as significant local expansion or contraction of nextneighbor atom column distances. That the local variation
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Fig. 8. (a) Simulated HRTEM image of a Pt particle supported by a 2.5 nm-thick carbon film. The imaging conditions correspond to those of Fig. 4. (b) Elastic strain component measured from image intensity peaks along the [2 2 0] direction through the Pt particle.
Fig. 9. Elastic strain profiles from surface regions of the simulated particle in Fig. 8a along a <2 2 0> direction.
of strain we observe in the experimental images is real is further confirmed by the absence of strain in the region where the particle of Fig. 2 is attached to neighboring particles. In conclusion, the strain distributions shown in Figs. 4b and 6 are trustworthy and accurate. 4.3. Physical origin of strain One remarkable result of the present work is the expansion of the lattice parameter of 1.3–4.8% we observe comparing the nanoparticle core to bulk Pt. From some of the existing literature on metal nanoparticles [47–53], one might actually expect no expansion, but a contraction of the lattice parameter. In a continuum model, such contraction may be attributed to the increasing influence of surface tension with increasing surface-area-to-volume ratio, i.e. decreasing particle size [50]. Of course, as particle sizes approach atomistic dimensions, continuum models become less useful. Moreover, crystallographic effects on the particle shape (facetting) need to be taken into account [50]. Finally, there are also some known cases in which the lattice parameter of metal nanoparticles was observed to be expanded with regard to bulk material. In FePt nanoparticles, for example, the expansion can be as much as 9% [54].
One potential mechanism by which the lattice parameter in the core of a nanoparticle might expand is the presence of solute atoms on interstitial sites. For example, the expansion of the lattice parameters that Jacobs et al. observed in Pd nanoparticles was attributed to inward diffusion of oxygen atoms when the particles were exposed to air [55]. In principle, the particles we have studied may contain e.g. interstitially dissolved hydrogen or oxygen (expected to dissolve interstitially from atomic-size consideration [56]). Owing to small atomic size, dissolved hydrogen should not cause significant lattice parameter expansion. Dissolved oxygen, in contrast, could in principle lead to a measurable effect. However, different from the case of Pd, which dissolves relevant concentrations of oxygen (and hydrogen), the solubility of oxygen in (bulk) Pt is so small that dissolved oxygen can hardly explain the observed lattice parameter expansion. Another explanation for the lattice parameter expansion in the nanoparticle interior arises from the even larger expansion of Pt–Pt next-neighbor distances observed at the particle surface. The Pt particles are so small that introduction of misfit dislocations cannot be an energetically favorable mechanism for relaxing misfit stress (this would be a case of strong bonding where the shear strength of bonds across the lattice-mismatched interface is comparable to the shear modulus of the material on either side [57]). Further, since the volume is small compared to the surface area, supporting the misfit stress introduced by a layer with expanded next-neighbor distances covering the entire surface will likely lead to considerable strain in the particle interior. The remaining question is what actually causes the substantial expansion of Pt–Pt next-neighbor spacings at the particle surface. Close inspection of the HRTEM image in Fig. 4a reveals a distinct amorphous layer on the crystalline core of the particle, discernible from the speckle pattern of the ultra-thin carbon support film by somewhat higher contrast and somewhat lower average intensity. Fig. 10 shows the HRTEM image of the nanoparticle in
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Fig. 10. The HRTEM image of Fig. 4 after adjusting the brightness and contrast to increase the visibility of the amorphous layer on the particle surface (indicated by the markers at the top left).
Fig. 4 after adjusting brightness and contrast to increase the visibility of the amorphous layer. As indicated by the markers on the top left, the thickness of this layer roughly corresponds to one times the lattice parameter of Pt, i.e. 0.5 nm. This layer could be an early stage of a platinum oxide scale. Indeed, HRTEM image simulations we obtained from models including a corresponding coverage of the particle surface with amorphous oxide qualitatively reproduce the observed contrast. The increase of Pt–Pt next-neighbor spacings at the surface of the crystalline core of the particle as well as the observed surface reconstruction yielding 5-member polyhedra of Pt atom columns would then find their explanation in the formation of Pt–O–Pt bonds expanding the Pt–Pt distance. (The presence of oxygen cannot be directly observed in HRTEM images. Owing to the relatively small electron scattering potential of oxygen compared to that of Pt, the HRTEM image mainly shows Pt atoms, i.e. no considerable contrast is expected from oxygen atoms.) The 5-member polyhedra of Pt atom columns in Fig. 5 seem to indicate a soft structural transition between the ordered crystal structure of Pt and the amorphous oxide, similar to what has been observed at the interface between crystalline grains and glassy grain boundary films in Si3N4 ceramics [58,59]. The hypothesis of expanded Pt–Pt next-neighbor distances by formation of Pt–O bonds agrees well with recent findings of Imai et al., who have observed the formation of platinum oxides at the surface of platinum nanoparticles in situ using time-resolved hard X-ray XRD (X-ray diffractometry) and XAS (energy-dispersive X-ray absorption spectroscopy) [60]. Similarly, Su et al. [61] have observed an 10% lattice parameter expansion near the surface of Ag nanoparticles and related this expansion to the formation of silver oxide at the surface of the particles. However, in spite of the expansion of Pt–Pt distances at the particle surface, and in apparent contradiction to our observations, Imai et al. find the core of the particles to be totally relaxed, i.e. unstrained and exhibiting the lattice parameter of bulk
Pt. However, the oxide layers in that case were thicker, and their detailed morphology is not known (e.g. small islands versus a continuous scale). Unfortunately, XRD and XAS cannot image the atomistic structure of the oxide–metal interface to reveal how misfit stresses are accommodated. In addition to the formation of an amorphous oxide scale, oxygen may expand the Pt–Pt next-neighbor spacings at the nanoparticle surface by dissolving on subsurface sites. Recent observations of Sachs et al. [62] of hydrogen in Pd suggest that the solubility of interstitial solutes at the surface of nanoparticles may be considerably higher than it is in corresponding bulk material. In conclusion, the most realistic explanation for the observed increase of Pt–Pt next-neighbor spacings at the Pt nanoparticle surface and the increase of the lattice parameter in the nanoparticle core is that Pt atoms at the crystallite surface bond to oxygen and/or oxygen is dissolved on subsurface sites of the nanoparticles. The resulting increase of the Pt–Pt next-neighbor spacing introduces coherency stresses that cannot relax by the introduction of misfit dislocations and therefore have to be supported by the relatively small particle core, causing its lattice parameter to expand as well. According to the results of other studies, an expansion of the Pt–Pt next-neighbor atom distance (i.e. dilatative elastic strain) can have considerable influence on the catalytic activity of Pt-based nanoparticles [5–7]. The expansion of 2% we observe near the nanoparticle surface is expected to reduce the catalytic activity compared to that of corresponding surfaces of bulk Pt [8–11]. On the other hand, the regions near steps, where the Pt–Pt distance is smaller than in the particle center, may be highly catalytically active. It needs to be stressed, of course, that the details of the particle structure can (and will) change during service, as observed in the beautiful experiment by Imaiet al. [60]. 5. Conclusion Our study shows that HRTEM, and in particular quantitative evaluation of HRTEM images with adjusted spherical aberration can make a major contribution to understanding performance-critical structural aspects of catalyst nanoparticles. Compared to other techniques that have been applied for studying catalyst nanoparticles, such as XRD (X-ray diffractometry) or XAS (X-ray absorption spectroscopy), HRTEM has the great advantage of providing highly resolved, local information on individual nanoparticles. The observation of expanded inter-atomic spacings at the particle surface and surface corrugation caused by stacking faults, twin boundaries, and surface steps suggests that the impact of strain and extended structural defects on catalytic activity warrants further investigation. The complex correlation we observed between the spatial variation of inter-atomic spacings, surface scales, particle shape, and contact to other particles demonstrates that techniques that average over a large ensemble of particles may lead to non-realistic conclusions. A clear dis-
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advantage of HRTEM is its limitation to ex situ studies, i.e. studies performed not in the same environment and at the same temperature at which respective catalysts operate. A possible route to a deeper understanding could be to combine TEM studies with in situ techniques, such as XRD and XAS, in order to enable more detailed interpretation of XRD and XAS in situ data in terms of structural details of individual particles. Acknowledgments We thank Stephen Campbell for stimulating discussions. We acknowledge financial support from the Department of Energy (contract No. DE-FC36-03-GO13107) and from the Army Research Office (MURI grant DAAD19-03-10169). Kui Du also thanks the Natural Sciences Foundation of China for support under Grant No. 50871116. Further, we thank Maya Bar-Sadan, Lothar Houben, Markus Lentzen, Martina Luysberg, and Andreas Thust at the Ernst Ruska Centre (ER-C) in Ju¨lich, Germany, for outstanding support and collegiality. References [1] Antolini E. J Appl Electrochem 2004;34:563. [2] Mukerjee S, Urian RC, Lee SJ, Ticianelli EA, McBreen J. J Electrochem Soc 2004;151:A1094. [3] Gasteiger HA, Kocha SS, Sompalli B, Wagner FT. Appl Catal B 2005;56:9. [4] Raimondi F, Scherer GG, Kotz R, Wokaun A. Angew Chem Int Ed 2005;44:2190. [5] Mukerjee S, Srinivasan S, Soriaga MP, McBreen J. J Electrochem Soc 1995;142:1409. [6] Mavrikakis M, Hammer B, Norskov JK. Phys Rev Lett 1998;81:2819. [7] Schlapka A, Lischka M, Gro A, Kasberger U, Jakob P. Phys Rev Lett 2003;91:016101. [8] Uribe FA, Zawodzinski TA. Electrochim Acta 2002;47:3799. [9] Salgado JRC, Antolini E, Gonzalez ER. J Phys Chem B 2004;108:17767. [10] Antolini E, Salgado JRC, Giz MJ, Gonzalez ER. Int J Hydrogen Energy 2005;30:1213. [11] Bezerra CWB, Zhang L, Liu HS, Lee KC, Marques ALB, Marques EP, et al. J Pow Sour 2007;173:891. [12] Alayoglu S, Nilekar AU, Mavrikakis M, Eichhorn B. Nat Mater 2008;7:333. [13] Suo Y, Zhuang L, Lu J. Angew Chem Int Ed 2007;46:2862. [14] Wallenberg L, Bovin JO, Petford-Long A, Smith D. Ultramicroscopy 1986;20:71. [15] Hofmeister H, Tan GL, Dubiel M. J Mater Res 2005;20:1551. [16] Marks LD. Ultramicroscopy 1985;18:33. [17] Coene W, Jansen AJEM. Scan Microsc Suppl 1992;6:379. [18] Zandbergen HW, Tang D, Van Dyck D. Ultramicroscopy 1996;64:185. [19] Du K, Phillipp F. J Microsc 2006;221:63. [20] Crozier PA, Tsen SCY, Liu J, Cartes CL, Perez-Omi JA. J Electron Microsc 1999;48:1015. [21] Kabius B, Haider M, Uhlemann S, Schwan E, Urban K, Rose H. J Electron Microsc 2002;51:S51.
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