Phase Behavior of Pseudobinary Precious Metal−Carbide Systems

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Phase Behavior of Pseudobinary Precious Metal-Carbide Systems John M. Gregoire,†,‡ Michele E. Tague,§ Eva H. Smith,† Darren Dale,| Francis J. DiSalvo,§ He´ctor D. Abrun˜a,§ Richard G. Hennig,† and R. Bruce van Dover*,† Departments of Materials Science and Engineering, Chemistry and Chemical Biology, and Cornell High Energy Synchrotron Source, Cornell UniVersity, Ithaca, New York 14853, United States, and School of Engineering and Applied Sciences, HarVard UniVersity, Cambridge, Massachusetts 02138, United States ReceiVed: September 20, 2010; ReVised Manuscript ReceiVed: October 27, 2010

Transition metal carbides exhibit a variety of interesting material properties, including electrochemical stability. When combined with precious metals, Ta and W carbides have shown promise as fuel cell electrode materials; yet, the phase behavior of these precious metal-carbide systems is largely unexplored. We investigated P-M-C phase behavior with P ) Pt, Pd, and Ru and M ) Ta and W using composition spread thin films. We attained limited control of the deposited carbide phase through variation of the sputter atmosphere and demonstrated decreased corrosion of W-C materials with increasing C content. A high-throughput X-ray diffraction and X-ray fluorescence experiment was employed for thin film characterization, which revealed solubility of Pt, Pd, and Ru in cubic WC. Density functional calculations of the lattice parameter dependence on carbon concentration enabled the determination of carbon concentration from the X-ray data as a function of transition metal stoichiometry. Our measurement of variations in the C stoichiometry and evolution of thin film texture with transition metal composition yielded surprising results. We detail how the combination of the composition spread technique, the high-throughput thin film characterization, and the density functional modeling of ternary carbide alloys provided a deep understanding of the chemical systems. Introduction Fuel cells represent an attractive technology to meet future energy needs due to their potential for high efficiency in converting chemical energy to electrical energy. However, to realize the energy conversion potential of fuel cells, it will be necessary to overcome current materials limitations, as exemplified by the degradation of the electrode materials. Transition metal carbides, which are known to have high mechanical and chemical stability, comprise a promising class of materials for fuel cells electrodes.2,3 In particular, carbides of W and Ta are being investigated as fuel cell electrode materials due to their high conductivity, high electrochemical stability, and catalytic properties.1,3 While carbides of Ta and W have been shown to have some catalytic activity for the reduction of oxygen4 and the oxidation of hydrogen,5,6 higher catalytic activity and electrochemical stability have been attained by using the carbide as a catalyst support and by incorporating the carbide in a ternary composite or alloy. For example, Ta and C play an important role in mitigating the electrochemical leaching of Ni in ternary Ni-Ta-C catalysts.7,8 Tungsten carbide-based materials have been more intensively explored and exhibit many desirable properties. Tungsten carbide-supported Pt has been shown to have high electrochemical stability and does not significantly oxidize in air below 450 °C.1,9 The tungsten carbide supports have also been found to mitigate Pt poisoning by CO and increase the activity for fuel oxidation due to interactions between the Pt and the carbide * To whom correspondence should be addressed. E-mail: vandover@ cornell.edu. † Department of Materials Science and Engineering, Cornell University. ‡ Harvard University. § Department of Chemistry and Chemical Biology, Cornell University. | Cornell High Energy Synchrotron Source, Cornell University.

support.1 Tungsten carbides have also been used to enhance the fuel oxidation activity of Pd catalysts, both as a support11 and as a composite material.10 Lu et al.12 studied CO tolerance in the pseudoternary system Pt-Ru-WC with thin films that were sputter-deposited from elemental and WC targets. The three materials were deposited in sequentially alternating layers and subsequently annealed at 520 °C to promote interdiffusion. The highest CO tolerance was found for an electrode with the approximate composition Pt25 Ru5 (WC) 70, but the crystallographic phase of the sample was not determined. In fact, despite early promising electrochemical results that motivate further development of carbide-based electrode materials, the phase behavior of carbide/precious metal systems has been largely unexplored. In this manuscript, we employ combinatorial materials science techniques to explore the phase behavior of the pseudobinary systems P-MC, where P is a precious metal (Pt, Ru, and Pd) and MC is a cubic transition metal carbide (WC and TaC). We additionally explore the pseudoternary system Pt-Ru-WC. While the 2:1 tungsten carbide, hexagonal W2C, has been explored as an electrode material, the higher C stoichiometry compounds demonstrate higher electrochemical stability.1 The hexagonal 1:1 phase, h-WC, is in equilibrium with W2C at low temperatures, but above 2500 °C, a C-deficient phase with the NaCl structure (c-WC) is formed.13 This phase has been observed in certain electrode preparation methods at relatively low reaction temperatures.11,14,15 Wang et al.15 recently reported preparation of nanosized c-WC materials, and the cubic phase was also attained by Jang et al.11 In the latter work, WCsupported Pd was prepared by sonochemical synthesis from inorganic salts, and the product was annealed at 450 °C. X-ray diffraction (XRD) analysis of the resulting electrodes revealed alloying of the Pd and c-WC cubic phases. This alloying during

10.1021/jp1092465  2010 American Chemical Society Published on Web 11/15/2010

Phase Behavior of Precious Metal-Carbide Systems electrode preparation indicates that further systematic development of ternary (or higher order) carbide catalysts may be enhanced by understanding the phase behavior of P-(c-WC) systems (P ) Pt, Pd, and Ru). The P-MC thin films presented in this work also exhibit alloying of c-WC with the precious metal(s) in each chemical system. Alloying of Pt or Ru with c-WC has not been reported in fuel cell electrode synthesis, but a few bulk cubic alloys with unknown (or poorly determined) carbon contents have been prepared, including Pt0.4W0.6Cz,16 Pt0.5W0.5C0.1, and Ru0.5W0.5C0.33.17 These single-composition alloys and the composition spread studies in this work demonstrate that P-W-C alloys can be formed in the fcc phase, but we note that explorations of the phase equilibria of h-WC and Pt-W alloys at 2000 °C revealed low solubility of these species.18 Further investigations of the ternary and quaternary systems are necessary for a more complete understanding of the equilibrium phase diagrams. The phase behavior of precious metal/TaC systems has not been discussed in the fuel cell electrode literature, but Holleck’s investigation of bulk samples indicated little solubility of TaC with Pt at 1500 °C, little solubility of TaC with Pd at 1300 °C, and a low solubility of C in Ru3Ta.18,19 These results are in reasonable agreement with our thin film studies, as we find very limited solubility of Pt, Pd, and Ru with TaC. Experimental Section Composition Spread Deposition. Films were prepared in a custom-built combinatorial sputter deposition system, described previously.20 Each P1-xMxCz library was generated by first depositing a 12 nm Ti adhesion layer (underlayer) onto a 76.2 mm diameter Si substrate. During and after this deposition, the substrate was radiatively heated and maintained at the library deposition temperature of 400 °C. For a given library, each transition element was sputtered from separate magnetron sputter sources (Angstrom Sciences) in a 0.66 Pa Ar/CH4 atmosphere. With the aid of a cryoshroud,20 the background pressure during deposition remained in the 10-5 Pa range. The geometrical relation of the deposition sources with respect to the Si substrate provided a deposition gradient from each source that upon codeposition resulted in a continuous variation in transition metal composition across the substrate. The range of x represented in a given library is dictated by the deposition profiles of the individual sources and the relative power delivered to these sources. The deposition rate from each source at substrate center was 1 × 10-9 mol/s/cm2, measured with a quartz crystal monitor just prior to deposition of the library. Carbon was incorporated into the growing film through reaction of the transition metals with the CH4 and its plasmainduced derivatives, both at the surface of the deposition source and at that of the growing film’s surface.21 The C concentration in the thin film monotonically increases with the CH4 concentration in the sputter atmosphere, and at a given substrate position (a given transition metal composition), the value of z is dictated by the reactivity of the transition metals with the carbon species. While CH4 is potentially a source for hydrogen incorporation into the thin film, at the growth temperature of 400 °C, the metal-hydrogen phase diagrams indicate that the hydrogen concentrations will likely be well below 1 at. % for all transition metal compositions.13 Thin Film Characterization. Diffraction patterns of the WCz thin films were acquired using a Bruker GADDS diffractometer. The crystallographic structure and composition of the composition spread thin films were characterized by a high-throughput XRD and X-ray fluorescence (XRF) technique. These experi-

J. Phys. Chem. C, Vol. 114, No. 49, 2010 21665 ments were conducted at the Cornell High Energy Synchrotron Source (CHESS) and were described in detail in a separate publication.22 Briefly, monochromatic 60 keV X-rays impinged the thin film with ∼1 mm2 spot size, and the diffraction image was attained in transmission geometry by a 345 mm diameter image plate (Mar Research). Integration of the images provided diffraction patterns that are compared to patterns of known phases in the powder diffraction file (PDF).23 Sputter deposition commonly yielded textured thin films in which the constituent crystallites were randomly oriented with respect to substrate azimuth but aligned with respect to substrate normal.24 The average crystallite orientation in the film was determined by analyzing the diffraction images.22 XRF spectra were obtained with a Rontec X-Flash energy-dispersive silicon-drift detector oriented orthogonal to the direct beam and analyzed with a custom algorithm.25 These measurements were not sensitive to elements as light as C, and thus, the XRF measurements provided the stoichiometry of the transition metals in the thin films. XRD Data Processing. The algorithms for data processing and peak detection of the XRD patterns are described in detail elsewhere.22,26 Lattice constants of identified phases were calculated from the fitted positions of Bragg reflections. To attain minimum uncertainty, only high intensity, nonoverlapped peaks were considered. For Ru-containing systems, comparisons between hcp and cubic phases were made by calculating Vatom1/3, the cube root of the average atomic volume of the transition metal. For a cubic lattice, Vatom1/3 differs from the lattice constant by a proportionality constant, and this quantity provides meaningful comparison between cubic and noncubic lattices, such as hcp. Electrochemical Stability. A high-throughput screening of the stability of the WCz films was performed by using the entire film as the working electrode in a specially designed electrochemical cell.27 The potential applied to the film was swept positive at 5 mV/s, and a large area Au coil served as the counterelectrode. The testing solution was near pH 7 with 0.1 M potassium triflate as the supporting electrolyte. Oxidation of the film results in a local decrease in pH, and in the pH and potential range of interest, W forms a water-soluble oxide.28 The working electrode current, which is attributed to thin film oxidation and corrosion, was recorded, and because all of the films were mirror-flat in appearance, the oxidation current density was calculated using the geometric area of the substrate. Density Functional Theory (DFT) Calculations. While Pt and c-WC are both fcc structures, c-WC has the NaCl structure type with interpenetrating fcc lattices of W and C. We treat the Pt1-xWxCz alloy observed in the thin films as a single-phase fcc alloy (NaCl structure) where the transition elements fully occupy the Na sites and the C atoms partially occupy the Cl sites. The site occupancy of the Cl sites by C is denoted by z. DFT modeling of these NaCl structures was performed with VASP,29,30 a density functional code using a plane-wave basis and the projector-augmented wave method.31,32 The generalized gradient approximation (GGA) of Perdew, Burke, and Ernzerhof33 was used with a Xe core for Pt and W and a He core for C. A planewave kinetic energy cutoff of 500 eV ensured energy convergence to 2 meV/atom, and the k-point meshes for the different structures were chosen to guarantee an energy accuracy of 0.5 meV/atom. Cell volume, cell shape, and ionic positions were relaxed until the total energy changed by less than 0.1 meV between relaxation steps. To determine the effect of carbon concentration, z, on the energy and structure of Pt1-xWxCz, supercells consisting of eight cubic NaCl unit cells were generated with random occupancy

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Figure 1. Phase identification and electrochemical oxidation of WCz films deposited at different concentrations of CH4 in Ar. The diffraction patterns (top) are overlaid with the peak positions for the corresponding phases, as reported in the respective PDF entries (bcc-W, 04-0806; W2C, 35-0776; and c-WC, 20-1316). The thin film oxidation current is shown for each WCz film and a pure bcc-W film. Each potential sweep proceeded at 5 mV/s, and shift in corrosion onset potential demonstrates increased stability with increasing z.

of Pt and W on the 32 Na sites and random population of C on 32z of the Cl sites. All supercells maintained a roughly cubic shape after relaxations, and the lattice constant of each relaxed supercell was estimated as the cube root of the unit cell volume. The standard elemental structures used in the heat of formation calculations were fcc-Pt, bcc-W, and graphite. In the graphite simulation, the in-plane lattice parameter, a, was allowed to relax, but the c/a ratio was fixed at the experimental value34 due to the well-known shortcomings of GGA functionals for modeling dispersion interactions. With these standard atomic energies, EX, for each element X, the heat of formation of the supercells containing 32(1 + z) atoms was calculated as

Hf ) EPt1-xWxCz - 32[(1 - x)EPt + xEW + zEC]

(1)

where EPt1-xWxCz is the total energy of the supercell. Results and Discussion Tungsten Carbide Phases and Stability. Four WCz thin films were deposited from a single deposition source in atmospheres of 4.8, 7.7, 12, and 20% CH4 in Ar. Diffraction patterns of the four films, shown in Figure 1, indicate that the crystallographic phase of the thin film changes with increasing CH4 concentration. At the lowest concentration, the bcc-W phase is obtained with a slightly increased lattice constant, presumably due to incorporation of C into bcc interstitials. The hexagonal W2C phase is obtained with 7.7% CH4, but hexagonal h-WC is not observed at higher CH4 concentrations. Instead, cubic c-WC is obtained, which in bulk WCz forms above 2500 °C with z j 0.82.13 As seen in Figure 1, c-WC is obtained at both 12 and 20% CH4, and increasing the CH4 concentration in this range results in a reduced grain size but not necessarily an increased

Figure 2. XRD intensity map for Pt-Ta composition spread thin films deposited in pure Ar (top) and 12% CH4 in Ar (bottom). The compositions corresponding to the diffraction measurements are indicated with horizontal red arrows on the left ordinate axis, and the diffraction intensity of the diffraction patterns is plotted in a logarithmic color scale. The peak positions for fcc-Pt (PDF entry 04-0802) and NaCl-TaC (PDF entry 4-007-4123) are also shown.

value of z. We expect that we have reached a saturation in the reactivity of the CH4, much like that observed in the deposition of nitride thin films in Ar-N2 atmospheres.35 In both examples of reactive sputtering, the grain size in the thin film decreases with increasing reactive gas concentration above the saturation level. Figure 1 shows that the stability of the films toward electrochemical oxidation increases with C stoichiometry, and the films that contain a binary carbide phase are significantly more stable than the pure W film. In particular, the onset potential of oxidation of the c-WC thin films is at least 250 mV higher than that of W. For TaCz films, similar trends with respect to CH4 concentration were observed both in the XRD data and in the electrochemical stability testing. Because of the relatively small XRD line width and desirable stability properties obtained with 12% CH4, this concentration was used for the deposition of the P1-xMxCz libraries discussed in the following sections. Pseudobinary TaC-Precious Metal Systems. Figure 2 shows the XRD characterization of Pt-Ta and Pt-Ta-C libraries. The Pt1-xTax diffraction map indicates that in addition to a fcc-Pt alloy, the binary thin film contains three binary intermetallic phases, as discussed in our recent publication.36 The diffraction map of the Pt1-xTaxCz film is remarkably different as only the two end-member phases are identified. While cosputtered Pt and Ta react in the absence of carbon, the Pt-TaC library contains segregated phases with the absence of chemical alloying indicated by the invariance of the Bragg peak positions with respect to x. This latter behavior of the chemical system is quantified in Figure 3, which shows the lattice constants of the identified phases for the Pt-TaC, Pd-TaC, and Ru-TaC systems with respect to the end-member lattice constants,

Phase Behavior of Precious Metal-Carbide Systems

Figure 3. Lattice parameter (or Vatom1/3) for P1-xTaxCz libraries with P ) Pt, Pd, and Ru. These values are calculated from individual Bragg peaks from identified phases, as noted in the respective legends. The carbon stoichiometries z are unknown, but the value and invariance of the lattice constants with x are indicative of very limited solubility of each precious metal and TaC.

which were obtained from the respective PDF cards. We note that aTaC ) 0.4465 nm was obtained from PDF card 4-004-5303, which reports the nearly stoichiometric compound TaC0.997.23 For all three pseudobinary systems, the solubility of the end-member phases appears to be less than 5 at. %, and no other phases are identified. The insolubility may be due to the large difference in the lattice constants (or Vatom1/3), which is J10% for each system. Another possible explanation for the insolubility is that since C is known to be insoluble in Pt, Pd, and Ru, the effective repulsion of these elements with C corresponds to a positive free energy for the substitution of precious metals on the M sites of high C stoichiometry MCz phases. While this simple explanation of the P-Ta-C phase behavior seems plausible, the following section demonstrates that very different behavior is observed in the P-W-C systems. Pseudobinary Precious Metal-WC Systems. DFT EWaluation of Pt-W-C NaCl Supercells. Pt1-xWxCz NaCl supercells were processed as described above for many values of x and z. For nearly every supercell, the heat of formation was found to be positive and commonly on the order of 0.1 eV/atom. Such high values of Hf indicate the presence of large entropic contributions to the free energy of this phase, which are not accounted for in our zero-temperature DFT calculations. We note that appreciable configurational entropy may exist due to the relatively low C occupancy z on the Cl sites but resolve that without consideration of entropy, the computational model does not reflect thermodynamic equilibrium and the constitution of the NaCl phase with the known h-WC and W2C phases cannot be determined. As such, we use the DFT modeling of NaCl supercells to establish compositional trends in the heat of formation and lattice constant. For each value of x used in DFT modeling, we found that Hf varies linearly with z in the range

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Figure 4. (a) From Pt1-xWxCz supercells at five values of x; the slopes of Hf vs z (0.3 < z < 0.7) are shown, indicating the energetic preference for C incorporation with increasing x. (b) For each of three Pt1-xWxCz compositions, Hf values of 12 randomly generated supercells are shown along with the strong trend of increased energetic stability with average population of W on C nearest neighbor sites. (c) A C atom (blue) in an octahedral interstitial site of an fcc Pt/W (yellow) parent lattice is shown with its six nearest-neighbor bonds. Pt and W randomly occupy the yellow sites of this NaCl structure, and z is the C occupancy of the blue sites.

0.3 < z < 0.7. The corresponding slope provides a zero temperature chemical potential of C, which is shown in Figure 4a. This energy for addition of C on a Cl site varies linearly with the composition on the Na sites and increases with increasing Pt concentration, as expected from the known insolubility of C in fcc-Pt. The linear variation of Hf with z and strong dependence of the slope on x indicate that C-(Pt, W) interactions dominate the enthalpy in this system. Further insight is gained by considering the variation in Hf in the randomly generated supercells at fixed composition. Figure 4c shows a NaCl cubic unit cell in which the six neighbors of the central C atom are indicated. At fixed, intermediate values of x and z, the average W population of these nearest neighbor sites varies appreciably among the randomly generated supercells due to their finite size. For three such x, z choices, the variation in Hf with this average W-C coordination is shown in Figure 4b, demonstrating that the average occupation of W on the C neighbor sites is an important proxy for the total supercell energy. One may expect that this energetic importance of the W-C bond would drive ordering into a lower symmetry phase with increased W-C coordination. While we find no evidence for such a phase in our diffraction data, further computational and experimental investigations are required to evaluate the existence of such a ternary intermetallic phase. Model for the NaCl Lattice Constant. Unlike the nearly stoichiometric c-TaC phase, the highest reported C stoichiometry in the c-WC phase is z ) 0.82 (PDF entry 4-007-4123).23 The associated lattice constant aWC0.82 ) 0.4215 nm is taken as the end member of the P1-xWxCz NaCl systems. While fcc-W has not been synthesized, the lattice constant aW ) 0.3988 nm was calculated by the DFT method. Because this lattice constant is usedalongsideexperimentalvalues,thePerdew-Burke-Ernzerhof exchange-correlation functional revised for solids37 was employed, and relative agreement to the other end member phases was confirmed by reproducing the known lattice constants with analogous DFT calculations.

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Figure 5. Shifting of the fcc {220} peak with transition metal stoichiometry. The compositions corresponding to the diffraction measurements are indicated by horizontal red arrows on the left ordinate axis, and the diffraction intensity of the diffraction patterns is plotted in a logarithmic color scale. The fitted peak positions used for lattice parameter determination are indicated with white stars. The variation in the intensity of the peak is due both to the variable film thickness and to the evolution of texture in the thin film. While this small window of scattering vector magnitude shows a single fcc Bragg peak, the corresponding shift in peak position was observed for >10 fcc peaks.

Given these end-member lattice constants, the Vegard model for this single phase system presumes that the lattice constant of the ternary alloy will vary linearly with x and z. The corresponding Vegard model for the C stoichiometry is

zV )

a - (1 - x)aPt - xaW (aWC0.82 - aW)/0.82

(2)

As discussed in the Supporting Information, the DFT-calculated lattice constants of the Pt1-xWxCz supercells deviate from this linear model. The deviation is consistent with a decrease in the W-C bond length with increasing x and z, which is due to the attraction described in Figure 4. While eq 2 is by construction correct at the end-member compositions, the variation in the effective size of W and C results in a correction factor to the Vegard model:

z ) zV /(1.521 - 0.521x)

(3)

In the following data interpretation, we use this modified Vegard model to calculate the C stoichiometry from the XRD measurements of a. We note that the magnitude of the correction factor in eq 3 is comparable to our experimental uncertainty, and thus, the qualitative results presented below can be obtained when the linear model of eq 2 is employed. In this manuscript, only eq 3 is used, and because the W-C interactions give rise to this correction factor, the same model is applied to the Pd-W-C and Ru-W-C systems. The Pt-W-C System. As in the P-Ta-C systems, the diffraction maps of the P-W-C libraries indicate the presence of only the end-member transition metal and c-WC phases. In contrast to the P-Ta-C systems, chemical alloying is apparent in each of the P-W-C diffraction maps. For example, Figure 5 shows the diffraction map of the Pt1-xWxCz library for a small range of scattering vector magnitude and indicates that the fcc {220} peak position shifts continually, but not linearly, across the entire composition range. In fact, all of the fcc peaks shift in a corresponding manner, indicating that the film crystallized into this phase with a unique lattice constant at each value of x.

Figure 6. fcc lattice parameter (top) and inferred C stoichiometry (bottom, eq 3) for Pt1-xWxCz libraries deposited at 12 and 20% CH4 in Ar. For each XRD measurement, a single fcc phase was identified in the diffraction pattern, and the lattice parameter was calculated from the noted Bragg reflection. The site occupancy model of eqs 4-6 is shown for two values of the critical number of W neighbors.

The variation in lattice constant with transition metal composition x is shown in Figure 6 and exhibits a marked departure from the straight line connecting Pt and c-WC, the end members obtained in thin films deposited from the respective single deposition source. The XRD data indicate that this Pt1-xWxCz library is single phase, and we conjecture that the nonlinearity in the lattice constant with respect to x is due to a variation in the z:x ratio, as dictated by eqs 2 and 3. The calculated z profile of Figure 6 indicates that the C incorporation in the thin film increases rapidly with increasing W concentration in the 0.4 < x < 0.8 range. As noted above, we predominantly used the 12% CH4 atmosphere for P-M-C phase exploration, but Figure 6 includes data from an additional film deposited in 20% CH4. While the Bragg line widths were larger for this film, the calculated z profile is essentially identical to that of the 12% CH4 film, and we thus conclude that in this range of CH4 concentration, the carbon incorporation into the fcc lattice is saturated. Simple Model for z. Simple alloying of the Pt and WC0.82 end members would correspond to z ) 0.82x for the entire range of x. To explore the chemical behavior responsible for the highly nonlinear trend in Figure 6, we consider the importance of W-C nearest neighbor bonds demonstrated in Figure 4b. A simple model for z is obtained by asserting that at a given x, the probability Pocc(n) that C occupies a Cl site is a function of only the population n of W on the nearest neighbor Na sites. The ensemble average of Cl site occupancy is then

z)

∑ Pocc(n)Pneigh(n, x)

(4)

n

where Pneigh(n, x) is the probability that a Cl site has n W neighbors at composition x. There are six such nearest neighbor sites (see Figure 4c), and under the assumption that Pt and W randomly populate the Na sites, combinatorial probability dictates that

Phase Behavior of Precious Metal-Carbide Systems

Pneigh(n, x) ) xn(1 - x)6-nC(6, n)

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(5)

and these distributions are shown in Figure 7. If Figure 4b contained free energies rather than Hf, Pocc(n) could be modeled by the corresponding Boltzman probabilities. Instead, we consider a simple step probability function at a critical value n ) n*:

Pocc(n) )

{

1 if n g n* 0 if n < n*

(6)

Equations 4-6 provide a model value for z, which is plotted for two values of n* in Figure 6. Comparison with the XRDdetermined values suggests that on average, C will occupy an fcc interstitial site if at least five of its nearest neighbors are W. We note that this modelsand in fact the interpretation of Figure 4bscan be equivalently explained in terms of Pt-C coordination instead of W-C coordination. We have chosen the W-C terminology as we generalize the discussion to the P-W-C systems in the following section. Pd-W-C and Ru-W-C Systems. The above interpretation of the XRD data was also applied to the Pd1-xWxCz and Ru1-xWxCz libraries, as shown in Figures 8 and 9, respectively. The Pd-W-C system is very similar to the Pt-W-C system with the single NaCl phase observed at every library composition. The assessed variation of z with x correlates well with the probabilistic C site occupancy model with a critical W coordination of 4. The higher values of z at lower x likely correspond to a weaker “repulsion” of Pd and C as compared to Pt and C. The Ru-W-C system is fundamentally different from the previous two examples as the end members have different structure types. The two end-member phases are observed in the XRD patterns, and the chemical alloying in each phase is analyzed with the model of eqs 2 and 3 using Vatom1/3 instead of the cubic lattice parameter. The variation in z with x indicates that the C stoichiometry is fairly constant over the 0.5 < x < 0.9 interval. While there is no measurable incorporation of C in hcp-Ru interstitials, these results indicate that in the NaCl structure, Ru is more carbon-phyllic than Pt or Pd. Pseudoternary Pt-Ru-W-C System. While the previous section presented Pt-W-C and Ru-W-C ternary alloys, the prevalence of Pt-Ru alloy catalysts in fuel cells prompted our investigation of the four-component, pseudoternary system Pt-Ru-W-C. The composition spread thin film was generated by deposition from the three metal sources in 12% CH4, and XRD/XRF analysis was performed at 62 film locations. Only the fcc and hcp phases were observed, and the calculated Vatom1/3 for each phase is shown in Figure 10. By extending the model

Figure 9. Vatom1/3 for the fcc and hcp phases (top) and inferred C stoichiometry (bottom, eq 3) for the Ru1-xWxCz library. Data are provided for x values at which the respective phase is robustly identified in the diffraction pattern. The site occupancy model of eqs 4-6 is shown for two values of the critical number of W neighbors, and the x ) z line is shown in black, demonstrating that the carbon stoichiometry surpasses the tungsten stoichiometry in this pseudobinary system.

Figure 7. In the NaCl structure with random Na site occupancy by Pt1-xWx, the plots show the probability distribution for n of the six neighbors of a given Cl site being W. As noted, the probability distributions are plotted for all seven values of n.

of eq 2 to a single phase, four-component system, z was determined for each phase. As in the Ru-W-C system (Figure 9), z was below 0.03 and within the uncertainty of zero for the hcp phase. In the fcc phase, the composition dependence of z is in strong agreement with Figures 6 and 9, with the carbon stoichiometry decaying more rapidly with increasing Pt than with increasing Ru. Deviations from Thermodynamic Equilibrium and the Role of Surfaces. The phase behavior of our thin films, which were deposited at 400 °C, is not necessarily representative of thermodynamic equilibrium at 400 °C or at room temperature. The materials reported in this manuscript were synthesized in a highly athermal environment where energetic atoms condense on a growing film’s surface, which interacts with various high-

Figure 8. fcc lattice parameter (top) and inferred C stoichiometry (bottom, eq 3) for the Pd1-xWxCz library. The site occupancy model of eqs 4-6 is shown for two values of the critical number of W neighbors.

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Figure 11. Diffraction intensity of the fcc {111} reflection is shown for each sample from the Pt1-xWxCz library (12% CH4) as a function of the angular displacement Φ from substrate normal. While an equiaxed film would exhibit a flat Φ distribution, the observed peaks indicate that the film contains (111) texture for small x and (100) for large x. The Φ values labeled with two fcc crystalline directions indicate the expected Φ position of the second direction’s Bragg peak provided that the first direction is at substrate normal (texture direction). Small deviations between the data peaks and these noted values are expected in our deposition geometry, as discussed in ref 22.

Figure 10. Vatom1/3 (top, middle) and inferred z (bottom) for the Pt-Ru-W-C library are shown in the Pt-Ru-W ternary composition space. The Vatom1/3 plots use the same color scale, on which the values for the end member phases are noted. The approximate positions of the hcp and fcc phase boundaries are shown, while Vatom1/3 values are noted only where robust calculations were possible. The interpolated map of inferred z includes data from only the fcc phase, as z remains negligibly small in the hcp phase.

energy species in the sputter ambient.21,24 After deposition, the low bulk diffusion rates of the transition metals inhibit crystalline grain growth and may create a barrier for phase transformations. As indicated in Figure 1, our reactive sputtering of W in CH4 results in the formation of the c-WC phase, which exists in the bulk W-C phase diagram in the limited 2530-2750 °C range.13 This nonequilibrium phase has been synthesized by a number of chemical methods relevant to fuel cell electrode preparation.11,14,15 While further investigation is required to understand the preferred growth of the fcc phase, an important insight may be gained from our XRD experiments. In addition to crystalline phase information, the diffraction images contain information related to the average crystal orientation in the film.22 For example, Figure 11 shows the spatial distribution of the fcc {111} reciprocal lattice vectors with respect to substrate normal for the Pt1-xWxCz library. Careful analysis of the spatial distribution for this and other families of reciprocal lattice vectors reveals that the thin film has (111) texture for Pt-rich compositions and (001) texture for W-rich compositions. The development of texture in a growing film arises from crystalline anisotropy, and the texture direction corresponds to the fastest-growing or lowest-energy surface of the crystallographic phase.24 The composition dependence of

texture in the Pt1-xWxCz library suggests that while the elements alloy into a single fcc phase across the composition range, the anisotropic fcc surface energy changes considerably. The low surface energy of the {100} facets of c-WC could play a role in the nucleation and preferential growth of this athermal phase in some chemical synthesis methods. Further studies are needed to confirm this assertion and to better understand the thermodynamics and kinetics of the growth of carbide and carbide alloy materials. We additionally consider the applicability of our thin filmbased results to the phase behavior of bulk P-M-C materials by comparing our results with ternary carbide bulk samples. The reported bulk samples discussed in the Introduction include an arc melted Pt-W-C sample with x ) 0.6, a ) 0.4039 nm; a Pt-W-C sample annealed at 2270 °C with x ) 0.5, z ≈ 0.1, and a ) 0.399 nm; and a Ru-W-C sample annealed at 2000 °C with x ) 0.5, z ≈ 0.33, and a ) 0.402 nm.16,17 The data from the Pt-W-C samples are in strong agreement with our thin film data of Figure 6, and the bulk Ru-W-C sample has a 2% lower fcc lattice constant and correspondingly lower C concentration as compared to our thin film sample of Figure 9. Because our pseudobinary carbide libraries reflect a kinetically stable (if not thermodynamically stable) equilibrium state of the respective materials system, it is useful to consider how these systems relate to the respective ternary phase diagrams. The invariance of z with increased CH4 concentration (see Figure 6) indicates that our P1-xWxCz libraries represent a binary cut through the ternary P-W-C composition space that follows a contour of maximal carbon stoichiometry. That is, we propose that in each P-M-C ternary phase diagram, a given equi- x line contains a phase boundary near our reported z value, and at higher z values, the P-M-C fcc alloy and carbon coexist. At lower z values, the phase equilibria of the fcc ternary alloy with W2C and the pure-metal phases should be explored, and at all compositions, the equilibria with known (e.g., h-WC) and other potential ternary phases should be considered. Electrochemical Stability of Ternary Carbide Alloys. The electrochemical stability of the (Pt, Pd, and Ru)-W-C libraries was assessed in neutral media up to 200 mV vs Ag/AgCl, as described above for the WCz films (see Figure 1). In the W-rich regions of these libraries (x < 0.2), oxidation similar to that of the WCz films is observed (see the Supporting Information).

Phase Behavior of Precious Metal-Carbide Systems The remainder of the composition range appears to be stable under these operating conditions, and further testing will be required to determine the electrochemical properties of these ternary carbide alloys. Conclusions The phase behaviors of precious metal/transitional metalcarbide systems are explored with Pt, Pd, and Ru and cubic TaC and c-WC end members. While there is little solubility between the precious metals and the TaC, complete solid solubility is observed with the fcc precious metals and c-WC. Similarly, high solubility of Ru in c-WC is demonstrated. Carbon stoichiometries are inferred from measured lattice constants using a model, which is established through DFT calculations. In the ternary carbide alloys, the carbon stoichiometry varies rapidly at a critical W stoichiometry, and we demonstrate that this behavior is well-modeled by a statistical, bond-counting analysis of the NaCl ternary alloy. The solubility of precious metals in c-WC has important implications for fuel cell electrode fabrication and suggests that investigations of new ternary carbide alloys as fuel cell catalysts are a fruitful direction for future research. These results demonstrate the effective exploration of inorganic phase diagrams using composition spread thin films and high-throughput crystallography. Combined with the DFT calculations, this array of techniques provides a platform for understanding complex inorganic systems. Acknowledgment. We thank Alexander Kazimirov and Anna Legard for assistance with the XRD/XRF experiments and Sophie Cahen for assistance with the electrochemical experiments. This material is based upon work supported as part of the Energy Materials Center at Cornell (EMC2), an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences under Award Number DE-SC0001086. This work is based upon research conducted at the Cornell High Energy Synchrotron Source (CHESS), which is supported by the National Science Foundation and the National Institutes of Health/National Institute of General Medical Sciences under NSF award DMR0936384. E.S. acknowledges support from a National Science Foundation Graduate Research Fellowship. Supporting Information Available: Derivation of the factor in eq 3 that gives rise to the modified Vegard model used in the interpretation of XRD data and electrochemical stability of the Pd-W-C library. This material is available free of charge via the Internet at http://pubs.acs.org. References and Notes (1) Ham, D. J.; Lee, J. S. Energies 2009, 2, 873. (2) Antolini, E.; Gonzalez, E. R. Appl. Catal., B 2010, 96, 245.

J. Phys. Chem. C, Vol. 114, No. 49, 2010 21671 (3) Zoltowski, P. Electrochim. Acta 1986, 31, 103. (4) Ishihara, A.; Doi, S.; Mitsushima, S.; Ota, K. Electrochim. Acta 2008, 53, 5442. (5) Yang, X. G.; Wang, C. Y. Appl. Phys. Lett. 2005, 86, 224104. (6) Nagai, M.; Yoshida, M.; Tominaga, H. Electrochim. Acta 2007, 52, 5430. (7) McIntyre, D. J. Power Sources 2002, 108, 1–7. (8) Yang, R.; Bonakdarpour, A.; Dhan, J. R. J. Electrochem. Soc. 2007, 154, B1. (9) Chhina, H.; Campbell, S.; Kesler, O. J. Power Sources 2007, 164, 431. (10) Hu, F.; Cui, G.; Wei, Z.; Shen, P. Electrochem. Commun. 2008, 10, 1303. (11) Jang, J. Electrochim. Acta 2009, 55, 485. (12) Lu, G.; Cooper, J. S.; Mcginn, P. J. J. Power Sources 2006, 161, 106. (13) Villars, P., Ed. ASM Alloy Phase Diagrams; ASM International: Materials Park, OH, 2006. (14) Hara, Y.; Minami, N.; Matsumoto, H.; Itagaki, H. Appl. Catal., A 2007, 332, 289. (15) Wang, Y.; Song, S.; Shen, P. K.; Guo, C.; Li, C. M. J. Mater. Chem. 2009, 19, 6149. (16) Laswons, A. C. J. Less-Common Met. 1971, 23, 103. (17) Holleck, H. Z. Naturforsch. 1979, 34b, 647. (18) Holleck, H. Metall 1983, 37, 703. (19) Holleck, H. Metall 1983, 37, 475. (20) Gregoire, J. M.; van Dover, R. B.; Jin, J.; DiSalvo, F. J.; Abrun˜a, H. D. ReV. Sci. Instrum. 2007, 78, 072212. (21) Thornton, J. A.; Penfold, A. S. Cylindrical Magnetron Sputtering. In Thin Film Processes; Vossen, J. L., Kern, W., Eds.; Academic Press: New York, 1978. (22) Gregoire, J. M.; Dale, D.; Kazimirov, A.; DiSalvo, F. J.; van Dover, R. B. ReV. Sci. Instrum. 2009, 80, 123905. (23) Powder Diffract. File, JCPDS International Centre Diffract. Data, PA, 2004. (24) Ohring, M. Materials Science of Thin Films, 2nd ed.; Academic Press: Boston, 2002. (25) Gregoire, J. M.; Dale, D.; Kazimirov, A.; DiSalvo, F. J.; van Dover, R. B. J. Vac. Sci. Technol., A 2010, 28, 1. (26) Gregoire, J. M.; Dale, D.; van Dover, R. B. ReV. Sci. Instrum. 2010, 81. (27) Gregoire, J. M.; Kostylev, M.; Tague, M. E.; Mutolo, P. F.; van Dover, R. B.; DiSalvo, F. J.; Abruna, H. D. J. Electrochem. Soc. 2009, 156, B160. (28) Pourbaix, M. Atlas of Electrochemical Equilibria in Aqueous Solutions; Cebelcor: Brussels, 1974. (29) Kresse, G.; Furthmuller, J. Phys. ReV. B 1996, 54, 11169. (30) Kresse, G.; Hafner, J. Phys. ReV. B 1993, 47, 558. (31) Kresse, G.; Hafner, J. Phys. ReV. B 1994, 50, 17953. (32) Kresse, G.; Joubert, D. Phys. ReV. B 1999, 59, 1758. (33) Perdew, J. P.; Burke, K.; Ernzerhof, M. Phys. ReV. Lett. 1996, 77, 3865. (34) Baskin, Y.; Meyer, L. Phys. ReV. 1955, 100, 544. (35) Bacon, D. D.; English, A. T.; Nakahara, S.; Peters, F. G.; Schreiber, H.; Sinclair, W. R.; van Dover, R. B. J. Appl. Phys. 1983, 54, 6509. (36) Gregoire, J. M.; Tague, M. E.; Cahen, S.; Khan, S.; Abruna, H. D.; DiSalvo, F. J.; van Dover, R. B. Chem. Mater. 2010, 22, 1080. (37) Perdew, J. P.; Ruzsinszky, A.; Csonka, G. I.; Vydrov, O. A.; Scuseria, G. E.; Constantin, L. A.; Zhou, X.; Burke, K. Phys. ReV. Lett. 2008, 100, 136406.

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