Hydrogen Solubility of Magnetron Co-Sputtered ... - ACS Publications

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Hydrogen Solubility of Magnetron Co-Sputtered FCC and BCC PdCuAu Thin Films Bruno Honrado Guerreiro, Manuel. H. Martin, Lionel Roué, and Daniel Guay* INRS Énergie, Matériaux et Télécommunications, 1650 Lionel Boulet Boulevard, Varennes (QC) Canada, J3X 1S2 S Supporting Information *

ABSTRACT: Palladium−copper−gold (PdCuAu) ternary alloy membranes are promising alternatives to pure palladium for the purpose of hydrogen separation processes. In the current work, PdCuAu ternary alloys were prepared by magnetron cosputtering deposition. All as-deposited samples showed face-centered cubic (fcc) structure in the composition range of study: 30 ≤ [Pd] ≤ 100, 0 ≤ [Cu] ≤ 66, and 0 ≤ [Au] ≤ 6 (in at%). The body-centered cubic (bcc) phase was prepared from the fcc samples by annealing for 4 h at 400 °C under Ar and was identified in alloys with Pd content between 40 and 46 at%. The hydrogen solubility was evaluated for both fcc and bcc phases using an electrochemical method and an electrochemical in situ X-ray diffraction method, respectively. The alloys’ hydrogen solubility increases with the lattice parameter and palladium content, and decreases with crystallite size and phase transition to bcc. In the same range of composition, the hydrogen solubility of the fcc and bcc phases vary in the same way with palladium content, suggesting that the latter factor is an important determinant of the hydrogen solubility in each phase. higher hydrogen flux when compared to PdCu or pure palladium.8 Both PdCu and PdAu alloys also show higher resistance to hydrogen embrittlement than pure Pd. The combination of PdCu and PdAu in a single ternary alloy PdCuAu may be beneficial, in the sense that the ternary alloy would gather the high permeability of PdCu alloys with the high H2S poisoning resistance of the PdAu alloys. To date, most studies on the application of PdCuAu ternary alloys for hydrogen purification have concentrated on alloys with face-centered cubic (fcc) structures. Experimental measurements of both solubility12,13 and permeability14,15 have been presented and compared to theoretical calculations made with, for example, density functional theory.16,17 For the fcc phase, it has been demonstrated that the hydrogen solubility of the ternary alloys is always lower than that of pure palladium, and palladium content is the decisive factor governing hydrogen solubility in this alloy type.12,13 Furthermore, in our previous work,12 we demonstrated that replacing copper with gold at constant palladium content results in increased solubility of the fcc phase. Recently, Tarditi et al. 15 demonstrated that the use of gold also promotes an increase in permeability at constant palladium contents. Indeed, of the two alloys, Pd70Cu25Au5 and Pd69Cu14Au17, the latter presents the higher permeability (1.9 vs 8.7 × 10−9 mol·m−1·s−1 Pa−0.5 at 400 °C). The permeability values are nonetheless lower than those of pure Pd (1.2 × 10−8 mol·m−1·s−1 Pa−0.5). The ternary alloys were also tested at 673 K in the presence of 100 ppm of

1. INTRODUCTION Hydrogen gas is an important commodity worldwide, with applications in the chemical, oil, and food industries. 1 Moreover, hydrogen gas is a promising energy carrier for the delivery of clean energy,2 as the combustion of hydrogen yields only water. Hydrogen production is primarily reliant on steam reforming of natural gas,3 a method that requires a final, costly purification step. Purification can be achieved by cryogenic distillation or pressure swing adsorption.4 Alternately, a membrane module located downstream of the reactor where hydrogen is produced could simplify and reduce the cost of the hydrogen purification process.5 Dense palladium-based membranes are particularly attractive for hydrogen purification. However, their widespread industrial application is impaired mostly by the high cost of palladium, the risk of hydrogen sulfide poisoning, and hydrogen embrittlement.3,6,7 The use of palladium alloys may be the key to overcoming the problems of pure palladium. Palladium− copper or palladium−gold alloys are promising candidates in this regard. In fact, a Pd47Cu53 (in at%) membrane shows higher hydrogen flux than a pure palladium membrane (0.208 vs 0.142 standard cubic feet per hour (scfh), respectively) of the same thickness and at the same pressure and temperature conditions.8 This represents a substantial cost decrease, as palladium content is reduced by 60% (mass-based) and copper is much cheaper than palladium. The high hydrogen flux through a membrane of PdCu alloys with low Pd content is attributed to the existence of a body-centered cubic (bcc) phase, which is observed at Pd concentrations between 34 and 47 at%, for which the diffusion of hydrogen is facilitated.9−11 Conversely, in the presence of H2S, PdAu alloys maintain © 2016 American Chemical Society

Received: November 2, 2015 Revised: February 25, 2016 Published: February 25, 2016 5297

DOI: 10.1021/acs.jpcc.5b10711 J. Phys. Chem. C 2016, 120, 5297−5307

Article

The Journal of Physical Chemistry C

to a deposition rate of 0.9 Å·s−1. Deposition rates were estimated with a microbalance using the equipment’s tooling factors. Films were 300 nm thick on average. 2.3. Characterization. The composition of the films was obtained by energy dispersive spectroscopy (EDX, from Oxford Link ISIS on JEOL JSM-6300 microscope). EDX measurements were performed at three different locations on the film at a 2500× magnification. Film composition was then calculated as the average of these measurements. In all cases, the standard deviation of the measurements was below 2%. The surface composition was determined by X-ray photoelectron spectroscopy (XPS) using VG Escalab 220i-XL instrument equipped with a Al Kα (1486.6 eV) monochromatic source. A survey from 0 to 1300 eV (100 eV pass energy) was first acquired, and high-resolution core level spectra (20 eV pass energy) were then recorded (Pd 3d 330−360 eV; Cu 2p 925−970 eV; Au 4f 75−95 eV; and C 1s 275−295 eV). The energy scale was calibrated for adventitious carbon by adjusting the C 1s core level peak at 284.5 eV.18 Element quantification was achieved by fitting the core level spectra with mixed Gaussian− Lorentzian functions considering a Shirley-type baseline, assisted by CasaXPS software. In some cases, Ar etching was performed. An estimation of the etch rate was made using the following equation:19

H2S for 24 h. Under these conditions, both ternary alloys showed less decrease in permeability than pure Pd (50−57 vs 85%, respectively). PdCuAu alloys exhibit a bcc phase (CsCl-type lattice) for palladium content between 29.5 and 45.8 at%.12 The permeability of bcc-PdCuAu alloys is expected to be as high as that of bcc-PdCu - if not higher−due to the aforementioned beneficial effect of gold. Using the same argument, the solubility of bcc-PdCuAu should also be higher than that of bcc-PdCu. However, contrary to bcc-PdCu alloys, the potential of the bccPdCuAu alloys remains unexplored, and comprehensive studies on solubility or permeability are lacking. In the present work, we address this issue by studying the solubility of bcc-PdCuAu alloys in greater detail. Ternary alloys were prepared by magnetron sputtering and by simultaneously operating the three metal targets (palladium, copper, and gold), a versatile one-step method for alloy production. The advantage of this method is that it allows deposits to be grown simultaneously on several substrates, thus, guaranteeing a sample set with precisely identical compositions. All as-deposited samples have fcc phases, but a heat treatment allows for the preparation of samples with bcc phase. Following characterization, the hydrogen solubility of the fcc and bcc samples was measured by electrochemical and electrochemical in situ XRD (E in situ XRD) methods, respectively. The latter method enables access to low hydrogen solubility values (lower than 1%), which is essential to access the hydrogen solubility of PdCuAu bcc phase. Finally, hydrogen solubility was shown to increase with palladium content, which also leads to an increase in the lattice parameter. The observed hydrogen solubility increase with the Pd content is similar for both fcc and bcc phases, despite the fact that the variation of the size of the interstitial sites with the palladium content is larger for the fcc than for the bcc phase. Consequently, the palladium content is the main factor governing solubility in each phase for the set of alloys that were studied here.

etch rate(nm·s−1) =

IYw ρAeN

(1)

where I is the etching beam current, Y is the sputter yield, w is the atomic weight of the ablated species and ρ is its density, A is the area of the sputtered surface, e is the electron charge, and N is Avogadro’s constant. X-ray diffraction (XRD) was used to elucidate the crystallographic structure of the films and was performed on a Bruker AXS D8 instrument with Cu Kα radiation (weighted average wavelength of 1.54184 Å) operating at 40 kV and 40 mA. The diffractograms were obtained with a grazing incidence angle of 5°. The angular step size was 0.05°, and the data acquisition time was 5 s/step. As-prepared samples were analyzed from 2θ = 30−100°. In some cases, films were heat-treated at 400 °C for varying periods of time under argon atmosphere. Afterward, the samples were slowly cooled to room temperature (typically overnight). The structure of the annealed samples was also assessed by XRD from 2θ = 20−100°. The position of the peaks and their full width at half-maximum (fwhm) were calculated by fitting the peaks with a Voigt function with OriginLab software. 2.4. Hydrogen Solubility Measurements. 2.4.1. Electrochemical Measurements. A custom cell was used for the electrochemical measurements, essentially consisting of a glass tube 7 cm in height, with an outer diameter of 1.2 cm and inner diameter of 0.9 cm, placed vertically on top of the working electrode (the PdCuAu film). Full details, including a schematic, can be found in the literature.12,20 Potential step chronoamperometry was used to measure hydrogen solubility via a Biologic SP300 potentiostat and EC-lab V10.02 software. All measurements were performed in 0.1 M NaOH (Fisher Scientific). A mercury/mercuric oxide (Hg/HgO) electrode with a Luggin capillary was used as the reference electrode and a platinum wire was used as the counter-electrode. To minimize the amount of dissolved hydrogen molecules, the NaOH solution was continuously replenished with a solution extensively purged with Ar through two peristaltic pumps (variable flow mini-pump, Control Company, Fisher Scientific);

2. EXPERIMENTAL SECTION 2.1. Substrate Preparation. Sputter deposition was performed on titanium 1.5 × 1.5 cm2 substrates cut from commercially available titanium foils 0.25 mm thick (Alfa Aesar). The substrates’ surface was first roughened with SiC sandpaper (BuehlerMetII 320/P400 grit from Buehler) to improve the adhesion of the sputtered films and prevent delamination upon hydrogen absorption. A simple cleaning of the substrates followed, which involved an initial sonication in acetone for 15 min and then a second 15 min sonication in isopropanol. Finally, the substrates were rinsed with isopropanol and quickly dried with a nitrogen stream. 2.2. Sputter Deposition. Thin films of PdCuAu ternary alloys on the Ti substrates were prepared by DC/RF magnetron Ar cosputtering with a background pressure of 7 × 10−7 Pa and 0.5 Pa deposition Ar pressure. The substrates were rotated at 10 rpm throughout the entire deposition. All depositions were performed for 3 min by sputtering all three metal targets simultaneously. Different film compositions were obtained by varying the power applied to each metallic target. For Pd (DC power source), the power varied from 120 to 340 W, which is equivalent to increasing the deposition rate from 5 to 16 Å·s−1. In the case of Cu (RF power source), the power varied between 100 and 500 W, which corresponds to a deposition rate of 4.1 and 7.3 Å·s−1, while the power was kept constant at 80 W for Au (RF power source), which corresponds 5298


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to the end of the experiment. The cell’s position was then carefully adjusted and the XRD measurements begun anew. Subsequently, the cell was pulled back down to position 1 and the potential switched to +300 mV to promote oxidation of the absorbed H. This potential was maintained for 15 min, after which the cell was moved upward into position 2 and a new XRD experiment was run. In summary, three diffractograms (XRD1, pristine sample; XRD2, charged sample; and XRD3, discharged sample) were taken per sample. This procedure was repeated for all samples. In one case, a diffractogram from 20 to 100° while maintaining −600 mV was recorded to verify for other possible changes in crystallographic structure.

one was used to remove the hydrogen-containing NaOH solution from the electrochemical cell, and the other to replenish the cell with an Ar-purged NaOH solution. In each case, the samples were first activated by applying Enegative = −1400 mV for tnegative = 3 min, followed by Epositive = −400 mV for tpositive = 3 min. This procedure was repeated five times. In order to measure the electrochemical hydrogen isotherms, Enegative was varied from −650 to −1400 mV in 50 mV steps (with tnegative = 3 min) while maintaining Epositive = −400 mV (with tpositive = 3 min). The oxidation current measured in the above conditions originates from oxidation of the hydrogen atoms previously absorbed in the metal alloy.20 Therefore, from the charge associated with the oxidation current (Qox), the number of moles of hydrogen (nH) may be determined by dividing Qox by Faraday’s constant. The detection limit of this technique is 1.5 mC,20 which H corresponds to M ≅ 1%. The determination of the mass of the deposit was performed by neutron activation analysis (NAA) at the SLOWPOKE NAA laboratory (École Polytechnique, Montreal, Canada). The average mass of the deposits in contact with the electrolyte was 0.151 mg, roughly corresponding to 1.6 × 10−6 mol of alloy. 2.4.2. Electrochemical In Situ XRD. This technique was used for alloys with an oxidation current lower than the detection limit of the electrochemical method. A custom cell that works both as an electrochemical cell and as a reflection unit in X-ray diffraction was used. A full description of its construction and operation is available in the literature.21−23 In this instance, the measurements were performed in 0.1 M H2SO4, using a Pt wire as the counter-electrode and a hydrogen-saturated palladium wire (Pd−H) as the pseudoreference electrode. The cell was then filled with the electrolyte and a 7.5 μm Kapton film (SPEXSamplePrep) was used to seal the top. The greater stability of the Kapton film in acidic versus alkaline media justifies the use of H2SO4 electrolyte instead of NaOH, as it was used in the electrochemical measurements of hydrogen solubility. There are two possible configurations for the operation of the E in situ XRD cell: position one, when the working electrode is fully immersed in the electrolyte; and position two, when the working electrode is above the solution level and in contact with the Kapton window. Hydrogen absorption and desorption occur in position 1 and XRD measurements occur in position 2. When in position 2, there is a thin film of electrolyte between the working electrode and the Kapton window that guarantees a complete electrical circuit. Before each XRD measurement, the cell’s position was carefully adjusted to the point where half of the X-ray beam’s intensity is detected. At this point, the working electrode was at the reference plane of the diffractometer. Prior to any electrochemical experiments, the cell was set to position 2 and a first diffractogram at open circuit potential (OCP) was recorded. The diffractograms were obtained with a grazing incidence angle of 5°, from 2θ = 75 to 81°. The angular step size was 0.04° and the data acquisition time was 6 s/step. Each diffractogram was repeated four times. After the first XRD measurement, the cell was pulled down to position 1, where the PdCuAu alloy films are saturated in hydrogen, by applying −600 mV versus a palladium hydride reference electrode. After 15 min, the cell was pushed upward to position 2 and was placed in contact with the Kapton window. The reduction current consequently decreased and remained virtually constant

3. RESULTS AND DISCUSSION 3.1. Characterization of as-Deposited Samples. The diffractograms of the as-deposited samples show two sets of peaks (Figure 1): one from the titanium substrate and the other

Figure 1. Diffractograms of representative samples indicating that an alloy was formed and peak shift occurs with changes in composition. The palladium content decreases from top to bottom and the composition of the samples is as follows: (A) Pd87.2Cu8.2Au4.6; (B) Pd76.2Cu19.7Au4.1; (C) Pd67.3Cu29.1Au3.6; (D) Pd59.1Cu37.6Au3.3; (E) Pd43.9Cu52.4Au3.7.

corresponding to a single-phase face-centered cubic (fcc) alloy. For example, in Curve A, the peaks of the fcc phase appear at 2θ = 40.2 (111), 46.8 (200), 68.4 (220), 82.4 (311), and 86.9° (222). Furthermore, the peaks shift with composition. For example, peak (200) is at 2θ = 48.2° at the lowest palladium content (Pd 43.9 at%, Figure 1, curve E) and shifts by 1.4° when changing palladium concentration to 87.2 at% (Figure 1, curve A). From the peaks’ position, the lattice parameter of the alloy may be determined by using the equation: a=

λ h2 + k 2 + l 2 2 sin θ

(2)

where a is the lattice parameter (in Å), λ is the wavelength of the Cu Kα radiation (1.5418 Å weighted value for copper Kα1 and Kα2), h, k, and l are the Miller indexes, and θ is the diffraction angle.24 Using the equation above, it is possible to 5299


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from zero-point shift or sample displacement by computing the position of the Ti(101) diffraction peak used as an internal reference. Indeed, over all the samples investigated in this study, the Ti(101) peak occurs at 2θ = (40.14 ± 0.01) Å. In comparison, the maximum shift of the (111) peak compared to that expected from Vegard’s law for a compound with the same composition is 0.26 Å, clearly indicating that the error that can be introduce due to zero-point shift or sample displacement is negligible. This deviation may be related to isotropic deformation energy, which in turn is related to the relative size dispersion of the atoms in a substitutionally disordered crystal. Previously, it was demonstrated for PdCuAu alloys prepared by electrodeposition that a maximum difference is observed at (3.82 ± 0.04) Å.12 In the current work, a maximum deviation is observed at approximately 3.80 Å, which means that the deviation observed from the Vegard’s law in the current case is also a consequence of the different sizes of the atoms that constitute the alloy. Ar inclusion during film preparation, a well-known phenomenon in sputter deposition,27 could also result in added uniform strain to the lattice and contribute to the larger than expected lattice parameters. This issue will be further discussed when presenting the results of the annealed samples. Elemental profile distribution was performed on a crosssection of the films perpendicular to the substrate, using Electron Energy Loss Spectroscopy (EELS). Examples are shown in Figure S1 in the Supporting Information. The elemental maps of Pd and Cu show both a layered structure (∼5 nm thick layers), while Au has a homogeneous distribution throughout the film. The oscillation in the composition of palladium and copper is a consequence of substrate rotation.28−30 3.2. Characterization of the Annealed Samples. The PdCuAu ternary alloys were annealed for varying time periods (4, 8, or 48 h) under Ar atmosphere at 400 °C. Samples were subsequently analyzed by XRD, cyclic voltammetry, XPS, and EDX. Depending on alloy composition, a phase transition may occur with annealing. For alloys with palladium content ranging between 30.1 and 43.9 at% (49.9 ≤ Cu ≤ 65.7 at% and 3.5 ≤ Au ≤ 6.1 at%), a 4 h heat treatment at 400 °C gives rise to a phase transition from fcc to bcc (CsCl-type lattice; an example is given in Figure S2 in the Supporting Information). Outside this composition range, the studied samples remained in fcc phase. The figures provided here are EDX compositions of the as-deposited samples. Later in this article, we will comment that the as-deposited sample compositions may not be representative of the bcc phase. Accordingly, a discussion of the stability domain of the bcc phase will await these measurements. Changes are also observed in the diffractograms of samples that remained fcc, namely, in the position and the full width at half-maximum (fwhm) of the peaks. As illustrated in Figure 3 for alloy with as-deposited composition Pd67.3Cu29.1Au3.7, the position of peak (220) changes from at 2θ = 69.45 (0 h), 69.52 (4 h), and 69.49° (8 h) to 68.56 (48 h), and the fwhm ranges from 1.05 (0 h), 0.44 (4 h), 0.42° (8 h) to 1.02 (48 h). For comparison, the titanium (103) peak is located at (70.58 ± 0.02)° in all four diffractograms, indicating that the smallest shift observed above (0.03°) is significant. The much larger shift observed with the 48 h heat-treated sample suggests that a new alloy is formed. Later in this article, we will present justification for such a shift; for now, we will only deal with samples heat-treated for 0, 4, and 8 h. The peaks’ position and the fwhm were used to calculate the lattice strain and crystallite

determine a lattice parameter for each [hkl] direction. In Figure 2, the average experimental lattice parameters calculated over

Figure 2. Comparison between the experimental lattice parameter and the calculated lattice parameter using Vegard’s law. The average experimental lattice parameters calculated over the different [hkl] directions are shown.

the different [hkl] directions are compared to expected values calculated using the alloy composition determined by EDX measurements and the following relation (Vegard’s law): aexpected =

y 100 − (x + y) x aPd + aCu + aAu 100 100 100

(3)

where aPd, aCu, and aAu are the lattice parameters of pure palladium (3.89 Å), copper (3.62 Å), and gold (4.07 Å), respectively, and x and y are the atomic percent compositions in Pd and Cu. Table S1 in the Supporting Information section summarizes the composition and the experimentally measured a lattice parameter of all samples investigated in the present study. Two comments can be made regarding Figure 2. First, for every sample investigated, the standard deviation of the lattice parameters determined from each specific [hkl] direction is lower than 0.006 Å. This is contrary to what was observed with PdCu alloys prepared by pulsed laser deposition, where the lattice parameters determined from the [200] direction are higher by about 0.03 Å in comparison with the other directions, indicating anisotropic lattice deformation.22 Similar results were reported for PdCuAu alloys also prepared with pulsed laser deposition.23 Galipaud and co-workers attributed this deformation to strain in the lattice arising from the deposition technique, and to the different Zener anisotropy ratios of the depositing metals. Pulsed laser deposition is a much more energetic process (up to 100 eV)25 than magnetron sputtering (5−10 eV),25,26 which means the extra energy in the former process is contributing to an additional deformation of the lattice following direction [200]. This is not observed in the present study and there is no evidence of anisotropic lattice deformation. The second comment regarding Figure 2 is that there is a positive deviation from Vegard’s law and the experimental lattice parameter is higher than that calculated from Vegard’s law. We made sure this deviation is real and is not originating 5300


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Figure 3. XRD patterns of fcc Pd67.3Cu29.1Au3.7 annealed at 400 °C for varying time periods.

size. Using a uniform stress deformation model31,32 and considering that both broadening effects (strain and crystallite size) are better deconvoluted by Gauss and Cauchy curves,33 respectively, the following equation can be written: 2 βhkl cos2 θ =

2σβhkl sin(2θhkl) λ + 2 Ehkl τ

Figure 4. Variations in lattice parameter (A), crystallite size (B), and uniform stress (C) as a function of the Pd content. The samples were heat-treated at 400 °C for various time periods. Crystallite sizes and uniform stress were calculated using eq 4.

(4)

where βhkl is the fwhm for the peak with hkl Miller indexes, τ is the crystallite size, σ is the uniform stress (in TPa), and Ehkl is Young’s modulus in the [hkl] direction. Details on how to calculate Young’s modulus may be found elsewhere.32 This model considers that stress is uniform in all directions and is related to the microstrain through Hooke’s law: σ ϵhkl = Ehkl (5)

from the cyclic voltammograms (CVs) recorded before and after hydrogen solubility measurements was deemed crucial to understand how annealing affects PdCuAu alloy structures. In Figure 5, the CVs of fcc phase Pd67.3Cu29.1Au3.6 alloy in NaOH 0.1 M are shown for different annealing times (0, 4, 8, and 48 h) in panels A, B, C, and D, respectively. The heat treatment was performed at 400 °C. The CV of the as-deposited sample (panel A, solid line curve) exhibits the same features as the CV taken after the hydrogen solubility measurement (dashed line curve). The major difference between the two curves is the position and intensity of the hydrogen oxidation peak, which shifts from −0.588 V to −0.651 V versus Hg/HgO after the hydrogen solubility measurement. In addition, the intensity of the peak increases from 1.616 to 3.649 mA. These changes in the hydrogen oxidation peak justify the need for an activation procedure prior to measuring hydrogen solubility. Similar results were obtained with the sample annealed for 4 h (panel B). However, after 8 and 48 h of annealing (panels C and D, respectively), the initial CV is remarkably different from the final one and, most notably, the hydrogen oxidation peak is absent. This suggests that the surface composition changes with the annealing, and the hydrogen solubility measurements in NaOH restore the initial surface composition, in the sense that the hydrogen desorption peak reappears. XPS and EDX measurements were performed to better understand the observed changes. The results for two different alloys, alloy 1 (not heat-treated) and alloy 2 (48 h annealed at 400 °C), are shown in Figure 6. Figure 6 shows the variation of the Au, Cu and Pd concentrations obtained from EDX and XPS measurements taken before (columns in black) and after (columns in orange) the hydrogen solubility measurements for Alloy 1 (as-deposited Pd50.8Cu46.0Au3.2, panels A, C, and E) and Alloy 2 (heat-treated Pd67.3Cu29.1Au3.6, panels B, D, and F). Overall, there is less accuracy in the quantification of gold due to its lower content

where ϵhkl is the microstrain in the [hkl] direction. eq 4 predicts that β2hkl cos2 θ varies linearly with 2βhkl sin(2θhkl)/Ehkl. Indeed, application of eq 4 to our data provides linear relationships with R2 higher than 0.9. Peak (222), with an intensity less than 4% that of peak (111), was not considered in these calculations due to the higher error associated with its fwhm measurement. From this analysis, it is possible to calculate the crystallite size and uniform stress on each sample, measured from the intercept and slope values of the linear regression, respectively. These results along with the experimental lattice parameters are shown in Figure 4 for three compositions of the films and for the 0, 4, and 8 h heat treatments. As observed in Figure 4, the lattice parameter contracts slightly, by 0.007 Å on average. Chalamala and Reuss27 demonstrated in their work that Ar gas trapped in sputter-deposited films is released upon annealing at 400 °C. Removal of Ar from the film would result in the release of strain and, consequently, in a reduction of lattice parameter. This is exactly what is observed in the present case, suggesting indeed that the lattice parameter contraction upon annealing is due to the release of Ar from the film. The crystallite size of the three samples increases with annealing, while their stress level decreases, as displayed in Figure 4. The increase of the crystallite size with annealing was already described in the literature for palladium thin films34 and palladium copper alloys.35 The effect of crystallite size on hydrogen solubility will be discussed later. In section 3.3, we will analyze how these ternary alloys respond to hydrogen absorption. However, information derived 5301


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Figure 5. Cyclic voltammograms (50 mV·s−1) taken before (black line) and after (red dashed line) hydrogen solubility measurements in NaOH 0.1 M. All samples shown were prepared simultaneously under the same conditions except for annealing time, which varied from 0 (A), 4 (B), 8 (C) or 48 h (D). All samples have fcc crystallographic structure and the as-deposited composition is Pd67.3Cu29.1Au3.6.

Figure 6. Variation of the Au (A, B), Cu (C, D), and Pd (E, F) content as a function of the different treatments. The content was determined by EDX and XPS. Two different ternary PdCuAu alloys were investigated: Alloy 1 at left (A, C, and E; Pd50.8Cu46.0Au3.2) with no heat treatment; and Alloy 2 at right (B, D, and F; Pd67.3Cu29.1Au3.6), annealed at 400 °C for 48 h.

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The Journal of Physical Chemistry C in the alloy in comparison with the other elements. Hence, the behavior of palladium and copper will be analyzed instead. First, let us examine the results for as-deposited Alloy 1. Prior to the hydrogen experiment), the bulk Pd (EDX measurements) and the surface Pd (XPS measurements) compositions are the same and equal to 50.8 at%. The same is true for the Cu content (46 at%). However, differences arise when the sample is immersed in the electrolyte and undergoes hydrogen absorption. Indeed, from XPS measurements, we note the surface becomes enriched with palladium (77.3 at%), while the bulk Pd content remains the same (50.4 at%) as that observed in the as-deposited state. Concurrently, the copper content at the surface of the sample is reduced (18.5 at%) when compared to the Cu bulk content (46.4 at%). Etching the surface for 24 nm decreases (increases) the surface Pd (Cu) content, with both values drawing closer to their respective bulk values, suggesting the change in composition occurs only on the film’s outermost layers. Looking at the results for the heat-treated Alloy 2 in Figure 6 (the same alloy whose CV is shown in Figure 5, panel D), we note a drastic difference between the bulk and the surface composition after 48 h of annealing and before the H solubility measurements. Indeed, the bulk palladium composition is 69.9 at% while the surface content is only 1.5 at%. The reverse is true for copper; there is a surface enrichment in copper when compared to the bulk (98.5 and 26.5 at%, respectively). In other words, the surface is becoming poorer (richer) in Pd (Cu) with the annealing treatment, suggesting that copper is diffusing to the surface (or Pd to the bulk). After the hydrogen solubility measurements, both bulk and surface become richer (poorer) in Pd (Cu). However, the final bulk copper content (21.8 at%) is lower than the initial bulk content of the asdeposited sample (26.5 at%). This strongly suggests that copper was dissolved by the electrolyte during the measurement. The CV results in Figure 5 now become easier to interpret. In fact, the CVs for the 8 and 48 h annealed alloy do not show the characteristic hydrogen oxidation peak because the surface was enriched in copper as a result of the annealing process. As copper does not absorb hydrogen to any great extent, the peak associated with the oxidation of absorbed hydrogen is not observed on the positive scan. Following, copper dissolution in the electrolyte during the measurement process, a new surface richer in Pd is exposed and the final CVs exhibit the features typical for this kind of alloy. Overall, the data supports two main thoughts: first, copper is migrating to the surface upon annealing, and second, copper is being dissolved by the electrolyte. At the pH and potential conditions used for this study, some copper is expected to form Cu(OH)2− and, consequently, dissolve into solution;36 however, this effect is more significant when there is already an enriched layer of copper at the electrode’s surface. This occurs with the alloy heat-treated for 48 h, and is the reason peak (220) shown in Figure 3 exhibits the greater shift. In fact, the segregation of copper at the surface led to the formation of a new alloy with increased palladium content (67.3 vs 74.5 at%) and consequently a larger lattice parameter (3.81 vs 3.87 Å). Copper segregation occurs not only with samples that remain fcc after annealing, as demonstrated above, but also with samples that undergo a phase transition to bcc. This is illustrated in Figure 7 with the CVs of an alloy with asdeposited composition Pd43.9Cu52.4Au3.7. In panel A, the CVs of the alloy before and after hydrogen solubility are shown for the as-deposited sample, while in panel B, those of the annealed

Figure 7. Cyclic voltammograms (50 mV·s−1) taken before (black line) and after (red dashed line) hydrogen solubility measurements in NaOH 0.1 M for a sample with composition Pd43.9Cu52.4Au3.7. Results for both as-deposited (A) and annealed for 4 h at 400 °C (B) samples are shown. Alloys have fcc and bcc structure before and after annealing, respectively.

sample are displayed. Before annealing (panel A), the alloy presents a hydrogen oxidation peak at −0.662 and −0.797 V versus Hg/HgO before and after the hydrogen solubility measurements, respectively. After annealing (panel B), no H oxidation peak is observed before hydrogen solubility, while it is present at −0.774 V versus Hg/HgO after the hydrogen solubility measurements. These results suggest again that annealing promotes copper segregation, and surface copper is dissolved when performing the hydrogen solubility measurements. However, industrial hydrogen separation is performed in the gas phase, so dissolution of copper would not be an issue in that situation. Tarditi and co-workers37 studied the surface composition of fcc and bcc PdCuAu ternary alloys. They observed that the near-surface region characterized by XPS is enriched in Pd upon heat treatment under hydrogen at 500 °C. This is contrary to what we observed (copper-enriched surface as measured by XPS) and the difference is most probably related to the atmosphere conditions used during annealing (Ar vs H2). This is confirmed by heat-treating the Pd39Cu54Au7 (EDX composition) sample at 400 °C under Ar with 5% H2, which results in surface Pd enrichment (XPS composition is Pd60Cu27Au13). The atmosphere under which annealing is performed clearly has an influence on metal segregation. As noted above, the dissolution of copper can occur to such an extent that even the bulk composition of the alloy will change. This change in composition is more relevant for thin films in the nm range. Figure 8 further explores the changes in copper composition occurring in the samples after hydrogen solubility. This figure outlines the bulk copper content before and after hydrogen solubility measurements for samples which underwent different heat treatments, with and without phase 5303


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Figure 8. Comparison between the copper content measured by EDX before and after hydrogen solubility measurements in NaOH 0.1 M (for the fcc phase) or 0.1 M H2SO4 (bcc phase). The dashed line shows the case where the final and the initial compositions are the same.

Figure 9. Influence of the palladium content on lattice parameter of the bcc phase. A comparison is drawn between the experimental values for PdCuAu alloys and literature values for PdCu.9 Dashed and dotted lines represent the best linear regression to each data set (PdCuAu3.5 and PdCu alloys, respectively).

higher than 1%) when compared to the others (standard deviation is lower than 0.6%). For alloys with constant Au content, the lattice parameter increases with Pd, in the same way that it increases for PdCu alloys. This is illustrated in Figure 9 for alloys with 3.5 at% Au. Both dotted and dashed lines represent the best linear fit to the PdCu and PdCuAu3.5 data sets, respectively. The most remarkable aspect is that both lines are parallel to each other. In the construction of Figure 9, the Pd content of the PdCuAu alloys is measured after hydrogen absorption occurred. If in Figure 9 the lattice parameters of the as-deposited samples had been used instead of those obtained after the hydrogen solubility measurements, no linear relation would have been found. The congruence of the data with published results on lattice parameter reinforces the hypothesis that Cu is dissolved in electrolyte, and the true composition of the bcc alloy is the one obtained after hydrogen solubility has taken place. 3.3. Hydrogen Solubility Measurements. 3.3.1. Electrochemical Measurements. The hydrogen solubility of the PdCuAu ternary alloys was assessed by an electrochemical method in aqueous alkaline media at room temperature. As discussed above, the surface copper dissolves in the electrolyte. For fcc samples that did not undergo annealing, the change in bulk composition is not significant, which may be the reason why this feature went unnoticed in previous publications.12,20,22 The method used is fully described in the literature.12,20 In short, H-absorption is promoted at a significantly negative potential, that is, from −650 to −1400 mV (vs Hg/HgO). After applying the reduction potential for 3 min, the potential is increased to −400 mV (vs Hg/HgO) for a further 3 min to ensure all of the H absorbed in the Pd alloy film is oxidized. The oxidation current is then integrated, and this charge is converted into a certain number of moles of absorbed hydrogen. Dividing this value by the number of moles of metal alloy results in the H solubility in that alloy, expressed as a percentage. The hydrogen solubility of PdCuAu fcc alloys (asdeposited and annealed) is presented in Figure 10 as a function of the lattice parameter (panel A) and palladium content (panel B). All alloys presented in Figure 10 have fcc crystallographic structure. In general, and similarly to what was previously

transition. For samples that remained fcc after heat treatment, hydrogen solubility was measured in NaOH 0.1 M, while the solubility of samples that underwent a phase transition to bcc were measured in H2SO4 0.1 M, as explained in the Experimental Section. If no changes in composition were observed, the data points fall on the dashed line in Figure 8. The most significant changes in bulk copper content are in samples with fcc structure annealed for lengthy periods (48h) and for bcc samples with very high initial copper contents. The fact that annealing promotes copper segregation at the surface gives significant consequences for characterizing the bcc phase. Indeed, the true composition of the bcc phase should be revealed through the bulk composition measured after the hydrogen solubility experiments, and not the composition of the corresponding as-deposited fcc phase. Consequently, the bcc phase was identified in the current work using the final palladium content, which was between 40.3 and 45.6 at%. No further attempts were made to prepare bcc alloys with lower palladium contents because the alloys with 40 at% Pd already exhibited near-zero hydrogen solubility (section 3.3.2). In Figure 9, the experimental lattice parameter a[211] for the bcc phase is plotted as a function of the palladium composition obtained after the hydrogen solubility measurements. In the same figure, a comparison is made with lattice parameter values for PdCu9 alloys in the literature. For the PdCu alloys, the dashed line represents the general trend, that is, an increase in palladium content results in an increase in the lattice parameter. In the case of the PdCuAu alloys, the analysis is more complex, as it must also account for Au content. In general, an increase in the lattice parameter is observed when copper is replaced with gold, while the palladium content is kept constant. Moreover, the higher the gold content, the larger the lattice parameter. For example, the Pd40.3Cu53.7Au6.0 alloy has a larger lattice parameter (2.976 ± 0.001 Å) than alloy Pd40.4Cu56.1Au3.5 (poorer in gold) whose lattice parameter is (2.966 ± 0.002) Å. Only one exception was observed, alloy Pd41.2Cu54.4Au4.4, which exhibited the lowest lattice parameter of all the PdCuAu bcc alloys in Figure 9 even though it does not have the lowest gold content. This may be due to greater uncertainty in measuring the composition of this alloy (standard deviation is 5304


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transition was already demonstrated in the literature for PdCu alloys.20 However, the decrease was more significant in the PdCu alloys (factor of 5; fcc-5.2%; bcc-1.0% for Pd41Cu59), whereas in the present case, the decrease in solubility is only affected by a factor 1.6. This may be related to the presence of gold which is known to increase hydrogen solubility.12 Accordingly, a thorough investigation of the hydrogen solubility of bcc PdCuAu alloys was performed. 3.3.2. Electrochemical In Situ XRD Measurements. The introduction of hydrogen in the lattice of the alloy will cause its volume to expand.42 This creates the possibility of calculating the amount of absorbed hydrogen simply by measuring the lattice parameter before hydrogen absorption takes place and with the alloy fully charged. As shown by J. Galipaud et al.,22 lattice volume expansion and hydrogen solubility are related through eq 6: H ΔV (%) = × 100 M VH × NM

(6)

where ΔV is the volume expansion, VH is the hydrogen molar volume (2.9 Å3 per atom), and NM is the number of metal atoms per cell in the lattice (four in the case of a fcc lattice, and two in the case of a bcc lattice). This method was used for samples exhibiting an oxidation charge below the electrochemical method’s detection limit, that is, 1.5 mC. X-ray diffractograms of the samples were recorded for 2θ = 75−81° (an example is given in Figure S3 in the Supporting Information), before and after hydrogen absorption is promoted. In this region, peaks of the titanium substrate (112) and (201), and peak (211) of the alloy are visible. In the present work, the titanium peak (112) was used as a reference and all diffractograms were calibrated for this peak. Peak fitting was performed using 2 G functions on each peak that reflect the contribution of Cu Kα1 and Cu Kα2 to the diffraction peak. By measuring the alloy peak position before and after hydrogen absorption, ΔV in eq 6 can be calculated and, consequently the hydrogen solubility can be known. The solubility values obtained using this method are outlined in Figure 11.

Figure 10. Hydrogen solubility of the fcc phase as a function of the lattice parameter (panel A) and palladium content (panel B). The results are for both as-deposited samples and samples annealed at 400 °C for varying time periods.

described in the literature for PdCuAu alloys,12 the increased palladium content (or lattice parameter) for nonannealed samples results in an increase in hydrogen solubility. Figure 10 also shows that annealing influences hydrogen solubility. For example, the Pd67.3Cu29.1Au3.7 alloy composition has a hydrogen solubility of 26.7, 25.5, 24.4, and 23.5% when annealed for 0, 4, 8, and 48 h, respectively. The annealing process promotes changes in bulk composition and crystallite size. As previously noted, changes in bulk composition are more significant for samples annealed for long periods of time. In this case, final palladium contents are 68.4, 68.1, 69.5, and 74.5 at% for the 0, 4, 8, and 48 h annealed samples, respectively. Crystallite size was shown in section 3.2 to increase with annealing. The calculated values of the crystallite size are 219, 325, 331, and 1225 Å for 0, 4, 8, and 48 h annealed samples, respectively. In the case of the 48 h annealed sample, the crystallite was calculated considering the Williamson-Hall equation.38 It appears that larger crystallite sizes are associated with lower hydrogen solubility, possibly because there are fewer grain boundaries where H could potentially be dissolved as well. These results agree with those of Adams et al.39 who studied hydrogen sorption in PdCd nanostructures. In the latter case, when the crystallite size increased from 9.16 to 25.81 nm, the charge for hydrogen desorption decreased from 54.38 (H/ M = 4.7%) to 3.20 mC·cm−2 (H/M = 0.3%). Nonetheless, in pure Pd, an increase of the grain size may lead to both an increase of the maximum α phase solubility and a decrease of the minimum β-phase solubility.40,41 The 48 h annealed sample actually suggests that crystallite size is important in determining hydrogen solubility. However, no other samples with comparable crystallite size and different palladium contents were studied, so no trend may be found. This issue will need further investigation. In the following, we will make sure to compare samples with approximate similar crystallite sizes. An alloy’s crystallographic structure also influences hydrogen solubility. For example, the hydrogen solubility of alloy Pd43.9Cu52.4Au3.7 fcc is 3.2%, whereas the bcc phase with the same composition shows a hydrogen solubility of 2.0%. This decrease in hydrogen solubility caused by the fcc-to-bcc phase

Figure 11. Hydrogen solubility of the bcc phase as a function of lattice parameter (panel A) and palladium content (panel B). In panel A, the dashed lines represent the best linear fit for each data set. In panel B, the values beside the data points represent the alloy’s gold content. 5305


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tetrahedral sites is 2.18 × 10−4 Å/at% . This represents a 34% difference of the size variation of the interstices between the fcc and bcc phases. As mentioned earlier, the variation of the hydrogen solubility with the palladium content does not differ by more than 3% between the fcc and the bcc phases. As a consequence, the size of the interstices is not as important as the palladium content in defining the hydrogen solubility of the alloy. In other words, the Pd content seems to be the primary factor determining hydrogen solubility in PdCuAu ternary alloys with constant Au content (approximately 3.5 at%) and in the range of study of palladium content. In this sense, it would be beneficial to extent the limits of the bcc phase toward higher palladium content in order to maximize the solubility and hence the permeability of the membrane.

Along with the results of E in situ XRD, Figure 11 also shows the results for the bcc sample obtained with the electrochemical measurements. In panel A, hydrogen solubility is plotted as a function of the lattice parameter (measured before and after the hydrogen solubility measurements), while in panel B, hydrogen solubility is plotted as a function of palladium content. Similarly to results from the fcc phase, hydrogen solubility increases with palladium content, as shown in panel B. The effect of gold is unclear, as only one sample with differing gold content was studied. However, from the results presented, the effect of gold is secondary to that of palladium. The samples with identical gold content (3.5 at%) were plotted in panel A as a function of the lattice parameter. Not only is the variation in hydrogen solubility based on lattice parameter linear, but the final lattice parameter is also larger than the initial one, and the observed increase is similar for all three samples (on average 0.06%). It would appear that hydrogen absorption is causing an irreversible plastic distortion of the alloy films. However, this may also be an indirect consequence of hydrogen absorption itself. Indeed, a titanium hydride phase was observed in the full XRD scan (not shown). The titanium hydride has a lattice volume larger that of titanium. This may justify why the alloy films do not return to their initial lattice parameter prior to hydrogen absorption, as they accompanied the expansion of the substrate lattice. In Figure 12, a comparison is made between the hydrogen solubility of the bcc phase and that of the fcc phase as a

4. CONCLUSION PdCuAu ternary alloys were successfully prepared by magnetron cosputtering. All as-deposited samples have fcc crystallographic structures, even for compositions where the thermodynamically stable phase is bcc. Phase transition to bcc is promoted by annealing, provided alloy composition falls within the range where the bcc phase is stable. Other changes occur upon annealing, namely the crystallite sizes increase and the surface of the alloys becomes enriched with copper in comparison to the bulk. The hydrogen solubility of the fcc and bcc alloys was evaluated by electrochemical and electrochemical in situ X-ray diffraction methods, respectively. After measuring hydrogen solubility, the alloy surface is enriched with Pd, most likely due to copper leaching to solution. Copper dissolution is more significant in fcc samples that were previously annealed for long periods as well as for bcc samples with high initial copper content. Three factors were identified as determinants for the hydrogen solubility of the alloys, namely, palladium content (both fcc and bcc phases), crystallographic structure (fcc vs bcc), and crystallite size (fcc phase). Palladium content affects the solubility of the fcc and bcc phases to equal extents.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.5b10711. Summary table of samples with composition and lattice parameters, elemental maps of Pd, Cu, and Au obtained with electron energy loss spectroscopy at a vertical crosssection of a Pd50.8Cu46.0Au3.2 thin film, XRD diffractograms of alloy with composition Pd32.2Cu63.3Au4.5 taken before and after annealing at 400 °C for 4 h under Ar, and electrochemical in situ XRD of bcc alloy with composition Pd40.4Cu56.1Au3.5 (PDF).

Figure 12. Comparison between the hydrogen solubility of the bcc phase and fcc phase as a function of palladium content. Dashed lines represent the best linear fit for each data set. Values beside the data points represent the alloy’s gold content.

function of palladium content at approximately the same gold concentration. Both data sets indicate that hydrogen solubility varies linearly with palladium content. The dashed lines in Figure 12 represent linear regressions of the data. The slopes calculated are 0.295 and 0.286 for the bcc and the fcc phases, respectively, which represents a 3% difference between the two phases. H in the fcc phase occupies the octahedral interstices, while in the bcc phase the tetrahedral sites are preferred.43 An estimation of the size of the lattice interstices, roctahedral or rtetrahedral can be made from the lattice parameter a, using the geometrical

relations

roctahedral =

0.414 × 2 a 4

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors would like to thank A. Korinek (McMaster University, Hamilton, Canada) for the work with the TEM/ EELS images. This work was supported by the Natural Sciences and Engineering Research Council of Canada (NSERC)

and

0.291 × 3 a

rtetrahedral = . From these relations, the size variation 4 of the fcc-octahedral sites with the palladium content is 3.35 × 10−4 Å/at%, while the corresponding increase of the bcc5306


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through the Strategic Grant Program, the NSERC Hydrogen Canada (H2CAN) Strategic Research Network, and Air Liquide Canada.

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