Journal of Membrane Science 375 (2011) 96–103
Contents lists available at ScienceDirect
Journal of Membrane Science
journal homepage: www.elsevier.com/locate/memsci
Electronic origins for sulfur interactions with palladium alloys for
hydrogen-selective membranes
Susanne M. Opalka a,∗ , Ole M. Løvvik b,c , Sean C. Emerson a , Ying She a , Thomas H. Vanderspurt a
a
b
c
United Technologies Research Center, East Hartford, CT 06108, USA
SINTEF Materials and Chemistry, NO-0314 Oslo, Norway
University of Oslo, Dept. of Physics, NO-0316 Oslo, Norway
a r t i c l e
i n f o
Article history:
Received 9 August 2010
Received in revised form 7 March 2011
Accepted 9 March 2011
Available online 16 March 2011
Keywords:
Pd alloys
Hydrogen membrane
DFT
Sulfur tolerance
a b s t r a c t
Atomic modeling was conducted to investigate the origin of S interactions with Pd alloy H selective
membrane candidates selected from the Pd–Cu, Pd–Ag, and Pd–Au binary systems, as well as their constitutive metals. The electronic characteristics of these alloy/metal systems played a more predominant
role in controlling S bonding behavior than surface site geometries. The electronic coupling of S p orbitals
bonding with alloy/metal d-bands in the adsorbate/slab density of states split the lower energy p bonding state and the d-band center further apart with increasing S bonding strength. A universal linear
correlation was established for increasing adsorption strength (decreasing adsorption enthalpy) of 0.25
monolayer S with the increasing density of states energy difference: [d-band center – S p bonding peak].
The S interactions predicted at higher coverage provided indications of alloy susceptibility to irreversible
S corrosion. The reversible adsorption of 1.0 monolayer S was only the most stable configuration on
the more open Pd0.5 Cu0.5 Im3̄m and P4 mmm (1 1 0) surfaces. The most competitive configuration for the
interaction of a full S monolayer with the Pd0.75 Cu0.25 Pm3̄m and Pd0.875 Au0.125 Fm3̄m surfaces was the
partial desorption and coupling of S. Partial incorporation of S to form a mixed absorbed/adsorbed S
monolayer was more favorable for the Pd Fm3̄m (1 1 1) surface, and also on the Pd0.5 Cu0.5 P4 mmm (1 0 1)
and Pd0.75 Ag0.25 Pm3̄m (1 1 1) surfaces when accompanied by Pd segregation. The combination of S incorporation and Pd segregation was interpreted to be the first step towards nucleation of irreversible Pd4 S
formation.
© 2011 Susanne M. Opalka. Published by Elsevier B.V. All rights reserved.
1. Introduction
Dense metallic Pd alloy membranes are being developed for the
separation of high purity H2 (>99.99% purity) from hydrocarbon,
biomass, or coal fuel-derived syngas. Thin Pd alloy membranes
selectively dissociate H2 in the presence of other syngas species,
and then effectively solubilize and permeate atomic H through their
lattices with a low mass transfer resistance. Fuel sulfur (S)-bearing
contaminants readily dissociate to form strongly adsorbed S species
on Pd alloy membrane surfaces, reducing H2 permeability by poisoning active surface sites for H2 dissociation and H solubilization
[1–3]. The impact of S contaminants can be minimized by designing S-tolerant Pd alloys that weakly bind S on their surfaces and
reversibly desorb S at elevated temperatures. For example, Pd–Au
and Pd–Cu alloy membranes were reported in pioneering studies
to be more S tolerant in permeating H2 than pure Pd or Pd–Ag alloy
membranes [1]. More recently, both experimentation and model-
∗ Corresponding author. Tel.: +1 8606107195.
E-mail address: OpalkaSM@utrc.utc.com (S.M. Opalka).
ing have demonstrated the S tolerance of selected Pd–Cu and Pd–Au
alloy membrane compositions [4–6]. In particular, the S tolerance
of Pd–Cu alloy H2 separation membranes is dependent on the Pd
fraction in the alloy. A 70.5 Pd: 29.5 Cu alloy (compositions given
in atomic or mole percent) membrane demonstrated strong S tolerance by exhibiting essentially no change in H2 permeance when
1000 ppm H2 S was added to the feed over the temperature range of
350–725 ◦ C [7]. On the other hand, a 47 Pd: 53 Cu alloy membrane
having a H2 flux of ∼14 cm3 cm−2 min−1 with a 90% H2 –10% He
feed at 350 ◦ C and 310 kPa stopped permeating H2 when 1000 ppm
H2 S was added to the feed [8]. Recent studies have also demonstrated S tolerance in H2 separation with Pd–Cu alloy membranes
having greater than 60 atomic % Pd. No significant changes were
observed in the H2 flux of 32.4 scfh ft−2 with a 50% H2 –50% N2 feed
at 553 ◦ C and 100 psig H2 when 487 ppm H2 S was added to the
feed [9].
Atomic modeling is being used to generate fundamental insights
into the experimentally observed S interactions with Pd alloys. It
has been shown that the dissociation of S-bearing gas contaminant species, such as hydrogen sulfide, H2 S [10–12], has relatively
low activation barriers on Pd alloy surfaces, and that the pathway
0376-7388/$ – see front matter © 2011 Susanne M. Opalka. Published by Elsevier B.V. All rights reserved.
doi:10.1016/j.memsci.2011.03.018
S.M. Opalka et al. / Journal of Membrane Science 375 (2011) 96–103
97
Table 1
Alloy and metal bulk and surface parameters.
Alloy/metal
Space group
Calculated
lattice
parameter (Å)
Pd0.5 Cu0.5
Pd0.5 Cu0.5
Pd0.5 Cu0.5
Pd0.75 Cu0.25
Pd0.75 Ag0.25
Pd0.875 Au0.125
Pd
Cu
Ag
Au
Pm3̄m
P4 mmm
P4 mmm
Pm3̄m
Pm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
a: 3.01
c: 3.80a
c: 3.80b
a: 3.88
a: 4.00
a: 3.98
a: 3.96
a: 3.62
a: 4.15
a: 4.17
4 layer Slab
configuration
2×2
2×2
1×1
1×1
1×1
1×1
2×2
2×2
2×2
2×2
Surface crystal
plane
Surface density
␳ (atoms/Å2 )
Surface energy
␥ (J/m2 )
Surface energy ␥seg
with 0.25 ML Pd
segregation (J/m2 )
(1 1 0)
(1 1 0)
(1 0 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
0.16
0.14
0.16
0.15
0.14
0.15
0.15
0.18
0.13
0.13
1.40
1.51
1.31
1.25
1.16
1.18
1.34
1.95
0.86
1.24
1.38
1.50
1.28
1.24
1.20
–
–
–
–
–
Corresponds to a disordered Fm3̄m structure.
Ordered P4 mmm c lattice parameter equivalent to a disordered Fm3̄m a lattice parameter. By crystallographic definition, the P4 mmm a lattice parameter = (2)−0.5 × c lattice
parameter.
a
b
to the resulting S adsorption is very favorable. Atomic modeling
has shown that S adsorbs strongly in 3-fold and 4-fold hollows
on the most stable Pd alloy surfaces [13,14]. The S adsorption
strength does not change significantly with Pd alloy composition,
and increases in adsorption strength with increasing (less negative) alloy d-band center energies [13] and increasing lattice size
[15]. The S adsorption strength decreases with increasing coverage
on Pd (1 1 1), leading to increasing S–S bonding interactions above
0.5 monolayers (ML) S coverage [14] and favorable S incorporation
(absorption) above 0.75 ML S [16].
With increasing exposure to S-bearing contaminants, atomic S
has been shown to irreversibly absorb into susceptible Pd alloys,
especially Pd and Pd–Ag alloys, to form a non-limiting corrosion
product layer [17,18]. This layer often predominantly contains
Pd4 S, which has a significantly lower H2 permeability [18]. The
absorption of the relatively large S atoms significantly distorts the
lattice to facilitate a high degree of S-metal coordination, serving as a precursor structure for formation of a Pd4 S surface layer
[16]. Here, in addition to distorting the Pd alloy lattice, the formation of such layers must be linked to ligand-induced selective
surface segregation of the Pd co-reactant. Binary transition metal
alloys are known to selectively segregate one (often lower surface
energy) alloying element to the surface, and form a sub-surface
layer enhanced with the secondary element [19,20]. In vacuum,
Pd–Ag alloys have been reported to have an Ag-rich surface layer
[21] and Pd–Cu alloys to have a Cu-rich layer [22]. This selective segregation can be enhanced or reversed by the adsorption of ligands
that bond more strongly to one alloying element than the other; for
example, the stronger adsorption interaction of H with Pd reverses
Pd0.75 Ag0.25 segregation to form a Pd surface layer [2,23,24], and
hydrocarbon [25] and S [22] interactions with Pd0.70 Cu0.30 enhance
Pd surface concentration. Since these ligand effects confound wellestablished ionic and electronic trends for alloy selective surface
segregation [22], atomic modeling plays an important role in delineating the relative tendencies for surface segregation and precursor
corrosion layer formation.
The current study will build upon the results of Alfonso [10–12]
and Hyman [15] that established fundamental S interactions with
selected Pd alloys, structural analogs, and constituent metals. Here,
atomic modeling results will be presented on an even wider range
of candidate Pd alloys for H2 selective membranes, as well as
their constituent metals, which exhibit varying S susceptibility.
The results will be interpreted to determine the influence of their
alloy/metal ionic structures and electronic characteristics on the
favorability and reversibility of S bonding interactions. Universal
trends will be established which can serve to guide the design of
S-tolerant Pd alloys.
2. Methodology
Atomic modeling calculations were made to compare S interactions with the most favorable surfaces of several Pd–Cu, Pd–Ag,
and Pd–Au alloys, as well as their constituent elements: Pd, Cu, Ag,
and Au. The alloy and element crystallographic information is given
in Table 1, with the alloy formulas specified in atomic (or mole)
percent. The Pd–Cu alloy models included the near maximum H
permeable 50 Pd: 50 Cu (Pd0.5 Cu0.5 ) Im3̄m B2 (ordered BCC) low
temperature phase, the corresponding 50 Pd: 50 Cu (Pd0.5 Cu0.5 )
P4 mmm (ordered Fm3̄m) high temperature phase, and the 75 Pd:
25 Cu (Pd0.75 Cu0.25 ) Pm3̄m phase. The analogous Pd–Ag alloy to the
latter Pd–Cu alloy, the 75 Pd: 25 Ag (Pd0.75 Ag0.25 ) Pm3̄m phase, was
selected for comparison. The high H permeability, S-tolerant Pd–Au
alloy, the 87.5 Pd: 12.5 Au (Pd0.875 Au0.125 ) Fm3̄m phase, was also
examined. These alloy phases were intentionally represented with
ordered structures, where reasonable size atomic models could
be created by imposing higher order symmetry (e.g., the ordering
of PdCu Pm3̄m disordered phase enforces a higher P4 mmm symmetry), without significantly altering the predicted atomic lattice
parameters.
Atomic models were relaxed to their local ground state minimum with the density functional theory Vienna ab initio simulation
package (VASP) code [26,27], using hard Pd pv 4p6 4d10 , hard Cu
pv 3p6 3d10 4s1 , regular Ag 4d10 4s1 , regular Au 5d10 5s1 , and
regular S 2s2 2p4 projector augmented wave potentials [28] with
the generalized gradient PW91 exchange-correlation corrections
[29], 0.3 Å−1 or finer spacing of the k-point meshes (with only
the point in the slab z directions), and spin polarization. All
calculations were made with a planewave cut-off of 410 eV and
Methfessel–Paxton smearing using a broadening of 0.3 eV. The criterion for self-consistency was that the total energy difference
between two consecutive cycles converged to less than 0.01 meV.
Bulk and ionic minimizations were made until the atomic forces
were converged to less 0.02 eV Å−1 . The alloy and element bulk
phase lattice parameters and atomic positions were first minimized prior to slab formation. Four layer slabs were then cleaved
from the minimized bulk models to reveal the lowest energy surfaces. Each of the alloy slab configurations listed in Table 1 had
four of each type of surface site geometry, so that the adsorption of 4 atoms or molecules constituted a monolayer (ML). The
slabs were constructed with 12 Å vacuum spacing between adjacent periodic repeats. Selective segregation was modeled in the
alloys by swapping the positions of segregating element in the subsurface layer with another element in the surface layer. The slab
ionic positions were minimized with the bottom layer fixed. The
ground state slab surface energies (), were determined from the
98
S.M. Opalka et al. / Journal of Membrane Science 375 (2011) 96–103
relationship
bare surface =
Eslab − m × Ebulk
2 × slab surface area
(1)
where Eslab is the energy of the surface slab, m is the number of
atoms in the slab, and Ebulk is the energy per atom of the corresponding bulk phase. A lower value is indicative of a more stable
surface.
The S adsorption and absorption calculations profiled all possible surface and subsurface sites, respectively. The most favorable
sites were filled for the 0.25 ML S adsorption and absorption results.
Additional sites were filled to the equivalent of 1.0 ML S for the
1.0 ML S adsorption and absorption results. At both coverages, S
absorption was modeled by incorporating 0.25 ML S in the most
favorable interstitial site between the subsurface and surface layers, so the equivalent to a full 1.0 ML S coverage was a mixed
configuration with 0.25 ML S absorption and 0.75 ML S adsorption. The influence of S ligands on the selective segregation of 0.25
ML Pd was probed wherever possible. Since H2 S is the most common S-bearing gas contaminant precursor species to the adsorption
or absorption of atomic S, the average S adsorption or absorption
enthalpies, Hads/abs , in units of kJ/mol × S atom, were determined
referenced to the dissociation reaction of H2 S → H2 + *S, where
*denotes an adsorbed species. The Hads/abs values were determined with the equation
Hads/abs =
Eslab+n∗S + EnH2 − (Eslab + EnH2 S )
n
(2)
where the E values were the energies of the product phases
minus the reactant phases and n was the number of adsorbing S
atoms. The more negative Hads/abs values indicated more favorable adsorption and stronger substrate–adsorbate bond formation.
The electronic densities of states were analyzed to distinguish
the unique electronic contributions of the metal/alloy slabs to the
adsorption interactions with atomic S. The d-band breadth, d-band
center, and the number of states at the Fermi level were analyzed from the slab total density of states (DOS), including both
the occupied and unoccupied electronic levels. The d-band center
was referenced relative to the Fermi level.
3. Results and discussion
3.1. Alloy surfaces
The H selective membrane Pd alloys and their constitutive metals provided an opportunity to probe the influence of their varying
lattice, electronic, and surface properties on the binding of atomic
S. The minimized metal/alloy structural and energetic properties
are given in Table 1. The atomic surface density (atoms/Å2 ), determined from the lattice parameter a, has a direct manifestation on
the surface energy (). Here, the most closed-packed Pd0.75 Cu0.25
and Pd0.75 Ag0.25 Pm3̄m (1 1 1); Pd0.875 Ag0.125 Fm3̄m (1 1 1); Pd, Cu,
Ag, and Au Fm3̄m (1 1 1); and Pd0.5 Cu0.5 P4 mmm (1 0 1) surfaces
showed a clear inverse correlation of the lattice parameter and surface density with the surface energy, as depicted in Fig. 1. The four
metal surface energies: Cu, Pd, Ag, and Au, are comparable to the
predicted and measured results reported in an earlier study [30],
and show decreasing surface energy with increasing lattice parameter. Typically for transition metals, a periodic trend of increasing
atomic size and decreasing bond strength occurs going down a
group within the periodic table (for example here, the group IB
coinage alloying elements in the current study), which is manifested as increasing lattice size and decreasing lattice energy. So,
since the surface energy is determined approximately (see Eq. (1))
as the ratio of (energy to break bonds)/(surface area), this ratio
decreases as one moves down within a group.
Fig. 1. Predicted surface energy and surface density versus lattice parameters for
the selected Pd alloys listed in Table 1.
The compilation of the alloy and metal electronic characteristics
in Table 2 only shows a linear trend of the (1 1 1) and (1 0 1) surface
plane d-band breadth, not of the d-band center or the density of
the states at the Fermi level, with lattice parameter (Fig. 2). Thus,
it can be seen by comparing Figs. 1 and 2 that there is a strong
correlation of surface energy with d-band breadth. The alloying
of Pd with 3rd row transition metal Cu with a lower d-band center results in a smaller lattice and a proportionately higher surface
energy. The addition of the even more electronically stable 4th row
transition metal Ag or the 5th row transition metal Au with even
lower d-band centers increases the lattice size and decreases the
surface energy. The slightly higher index Pd0.5 Cu0.5 Im3̄m (1 1 0)
and Pd0.5 Cu0.5 P4 mmm (1 1 0) surfaces do not follow this trend
and have intermediate surface energies. The correlation of these
characteristics with the lattice parameter could be attributed to
geometric effects. Atomic modeling of geometric effects, such as
isotropic lattice strain, typically shows that metal d-band characteristics: center, breadth, and number of electrons at the Fermi level
track one another with varying lattice parameters. Such modeling
showed that the d-band broadened, lowering the d-band center as
the metal lattice was compressed [15]. The absence of complete
d-band correlation with lattice size alludes to the predominating
influence of electronic characteristics for controlling the surface
reactivity of these Pd alloys that will be revealed in the following
discussion.
Fig. 2. Correlation with the Pd alloy electronic density of states d-band centers and
d-band breadths with their lattice parameters.
S.M. Opalka et al. / Journal of Membrane Science 375 (2011) 96–103
99
Table 2
Slab Electronic characteristics determined from density of states analyses.
Alloy/metal
Model space group
Calculated lattice
parameter (Å)
Surface crystal
plane
Total d-band center
(eV)
Total d-band
breadth ( eV)
Density of occupied
states at Fermi
level (number of
electrons)
Pd0.5 Cu0.5
Pd0.5 Cu0.5
Pd0.5 Cu0.5
Pd0.75 Cu0.25
Pd0.75 Ag0.25
Pd0.875 Au0.125
Pd
Cu
Ag
Au
Im3̄m
P4 mmm
P4 mmm
Pm3̄m
Pm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
a: 3.01
c: 3.80
c: 3.80
a: 3.88
a: 4.00
a: 3.98
a: 3.96
a: 3.62
a: 4.15
a: 4.17
(1 1 0)
(1 1 0)
(1 0 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
−1.39
−1.86
−1.84
−1.59
−1.92
−1.63
−1.59
−2.44
−3.97
−3.21
16.41
15.39
16.02
15.42
14.16
15.14
12.16
18.03
13.27
14.98
10.26
9.96
4.12
13.21
11.17
10.08
20.56
0.00
0.36
2.50
The stability of the alloy surfaces with respect to preferential
segregation of one alloying element during annealing at elevated
temperatures can have an impact on the composition of local surface sites available for ligand adsorption. All of the bare alloy
surfaces were predicted to be relatively stable with respect to
preferential elemental segregation at the ground state conditions,
which approximate a vacuum atmosphere. The segregation of
0.25 ML Pd was predicted to be only slightly more stable for the
Pd0.5 Cu0.5 Im3̄m (1 1 0), Pd0.5 Cu0.5 P4 mmm (1 0 1), and Pd0.75 Cu0.25
Pm3̄m (1 1 1) surfaces, as shown by the lower surface energies in
Table 1. However, these differences may not be significantly distinguishable within the error of the calculation methodology, which
is typically in the range of 5–15 kJ/mol [31]. The segregation of the
secondary alloy element, either Cu or Ag, was not predicted to be
favorable for any of the Pd alloy surfaces. Complete Pd termination
of the Pd0.875 Au0.125 Fm3̄m (1 1 1) was slightly more favorable than
for partial Au termination, so that further Pd segregation could not
be evaluated. The predicted stability of the ordered alloy surfaces
with respect to segregation, contrast with experimental reports of
selective surface segregation after exposure of the related Pd alloys
to different gas atmospheres [4,25,32]. Here, segregation in the
latter may be attributed to local concentration gradients or heterogeneities in surfaces of the disordered alloys. All of the alloys, with
the exception of the Pd0.5 Cu0.5 ordered Im3̄m (B2) phase, exist as
disordered phases with both alloying elements occupying the same
lattice sites.
3.2. Atomic S adsorption
The modeled trends in atomic S adsorption given in Table 3,
for the most part, showed a strong linear dependence on lattice
parameter for the different types of surfaces. A linear trend for H2 S
adsorption enthalpies with lattice parameter, and a direct correlation between H2 S and S adsorption enthalpies was reported for
selected Pd alloys and chemical analogs [15]. The trend of adsorption enthalpy becoming more negative (more favorable) linearly
with increasing lattice parameter was especially apparent for the
adsorption of 0.25 ML S at the most favorable metal/alloys sites
on the (1 1 0) and on the (1 0 1)/(1 1 1) families of surfaces (with
the exception of the Ag and Au (1 1 1) surfaces). If geometric effects
alone were operative, then the stronger S binding would be a direct
manifestation of increased reactivity due to the narrowing of the
d-band, the shift of electronic states closer to the Fermi level, and
increases in the number of unoccupied states, with increasing lattice parameter [15]. However, such trends are superseded by the
electronic effects introduced by the wide range of alloy chemistries.
Similar to the surface energy trends shown in Fig. 1, the trend seen
for the most close-packed (1 1 1) and (1 0 1)-type surfaces was separate from that for the more open (1 1 0) surfaces. The only excursion
from these trends was for the very weak S adsorption on the Ag
and the Au (1 1 1) surfaces with the completely filled d-bands. This
trend of decreasing S adsorption enthalpy also corresponds to the
trend of decreasing surface energy with increasing lattice parameter within a transition metal group of the periodic table. Decreasing
Table 3
Structural characteristics and adsorption enthalpies, Hads , for the most favorable 0.25 monolayer S adsorbed on the selected Pd alloys and constituent metals.
Alloy/metal composition
Space group
Surface crystal
plane
Most favorable site
for 0.25 ML S
adsorption
0.25 ML S Hads
(kJ/(mol × atom))
Most favorable site for 0.25
ML S adsorption after 0.25
ML Pd segregation
0.25 ML S with Pd
segregation Hads
(kJ/(mol × atom))
Pd0.5 Cu0.5
Im3̄m
(1 1 0)
−148
P4 mmm
(1 1 0)
Pd0.5 Cu0.5
P4 mmm
(1 0 1)
−109
4-fold, 3Pd and 1 Cu
surface, 1Pd below
4-fold, 3Pd and 1 Cu
surface, 1Cu below
3-fold fcc Pd3 surface
−122
Pd0.5 Cu0.5
Pd0.75 Cu0.25
Pm3̄m
(1 1 1)
−120
3-fold fcc Pd3 surface
−101
Pd0.75 Ag0.25
Fm3̄m
(1 1 1)
−142
3-fold fcc Pd3 surface
−119
Pd0.875 Au0.125
Fm3̄m
(1 1 1)
−142
N.A.
N.A.
Pd
Cu
Ag
Au
Fm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
4-fold, center of
hemi-( octahedraa
4-fold, center of
hemi-( octahedrab
3-fold fcc Pd2Cu
surface
3-fold fcc Pd3
surface
3-fold hcp Pd3
surface
3-fold fcc Pd3
Surface
3-fold fcc
3-fold fcc
3-fold fcc
3-fold fcc
−139
−101
−42
−37
N.A.
N.A.
N.A.
N.A.
N.A.
N.A.
N.A.
N.A.
a
b
Hemi-␣ octahedra-2 Pd and 2 Cu surface, 1 Pd below.
Hemi-␤ octahedra-2 Pd and 2 Cu surface, 1 Cu below.
−175
−179
−112
100
S.M. Opalka et al. / Journal of Membrane Science 375 (2011) 96–103
reactivity or adsorbate binding strength is typically observed going
down a transition metal group, especially within the group IB metals [33]. This may seem to be counter-intuitive, since studies most
often examine catalytic trends going across the transition metal
rows in the periodic table, where typically adsorption enthalpies
become less negative with decreasing surface energy (increasing
surface stability) going from left to right.
It is well known that stronger ligand interactions with one
element in an alloy can induce selective segregation of that alloying element. However, the adsorption of S was not predicted to
destabilize the surface alloy compositions with respect to selective
segregation. Partial segregation of an additional 0.25 ML Pd to the
surface in the presence of 0.25 ML adsorbed S was only slightly
more favorable for both the Pd0.5 Cu0.5 P4 mmm (1 1 0) and (1 0 1)
surfaces, and was distinguishably unfavorable for the other most
stable alloy surfaces. Note that experimentally the high temperature Pd0.5 Cu0.5 phase is a disordered face-centered cubic phase,
where Pd and Cu do not occupy separate sublattices. This may contribute to the difference in the predicted and the experimentally
observed segregation behaviors.
The electronic densities of states for 0.25 ML S adsorbed on the
alloy/metal slabs were examined in order to discern whether the S
binding was controlled by changes in the slab electronic structure.
A linear correlation of S adsorption enthalpy was found with the
d-band centers of all slab surfaces, as shown in Fig. 3, akin to the
relationships reported in a previous S adsorption study [13]. This
follows the typical trend for transition metal series, of the inverse
correlation of atomic adsorbate binding with the d-band center,
where the adsorption weakens with the lowering of the d-band
center [34]. A very weak correlation, if any, was observed with the
d-band breadth or the density of states at the Fermi level.
Examination of the density of state plots reveals the changes in
electronic configuration upon S adsorption on the metal and alloy
surfaces. The p orbitals of adsorbed S atoms hybridize with the transition metal substrate d-band, forming lower energy bonding states
and higher energy anti-bonding states [15]. These bonding and anti-
Fig. 3. Adsorption enthalpies, Hads , for adsorption of 0.25 monolayers S on the
most favorable surfaces of the selected Pd alloys and their constituent metals versus
d-band center for the electronic density of states.
bonding states are repelled by the d-band, shifting to increasingly
lower and higher energies, respectively, with increasing adsorbate
binding [34]. An example of this trend can be seen by comparing
the densities of states for 0.25 S adsorbed on the Pd0.75 Cu0.25 and
Pd0.75 Ag0.25 Pm3̄m (1 1 1) surfaces to that for the bare surfaces in
Fig. 4. Sulfur is adsorbed more weakly on the Pd0.75 Cu0.25 Pm3̄m
(1 1 1) surface (Fig. 4a and c) resulting in a smaller splitting of the
DOS anti-bonding and bonding S p peaks, compared to S adsorption on the Pd0.75 Ag0.25 Pm3̄m (1 1 1) surface (Fig. 4b and d). This
repulsion, or shifting, is a function of the position of the d-band
center. It can be quantified as the difference between the d-band
center (DBC) and the hybridized position of the p bonding peak of
the S adsorbate partial density of states (Spb ), [DBC-Spb ]. An even
stronger correlation was found between the enthalpy of adsorption,
Fig. 4. Comparison of the densities of states for 0.25 S adsorbed on the Pd0.75 Cu0.25 Pm3̄m (1 1 1) [left panels (a) and (c)] and the Pd0.75 Ag0.25 Pm3̄m (1 1 1) [right panels (b)
and (d)] surfaces. The models with adsorbed S (top panels) are compared to the pure models without S (lower panels). The overlap of the S p partial density of states with
anti-bonding and bonding peaks are clearly visible. The Fermi level, d-band center (DBC) and S p bonding peak (Spb ) are marked with solid, dashed, and dotted vertical lines,
respectively.
S.M. Opalka et al. / Journal of Membrane Science 375 (2011) 96–103
Fig. 5. Adsorption enthalpies, Hads , for adsorption of 0.25 monolayers S on the
most favorable surfaces of the selected Pd alloys and their constituent metals versus
the difference [d-band center − S p bonding peak], [DBC−Spb ], determined from
the electronic density of states.
Hads 0.25ML S and the difference, [DBC-Spb ], as shown in Fig. 5.
This provided strong evidence that the S binding was strongly controlled by an electronic chemistry effect. As exemplified in the work
of Hyman et al. [15], the geometric influences of lattice strain were
manifested clearly in the d-band breadth, while the electronic (ligand) effects of varying alloy compositions were mainly reflected in
trends of the d-band center.
The adsorption enthalpy, Hads , trend for a full 1.0 ML S in
Table 4 is parallel to that for 0.25 ML S, but is shifted to significantly
higher (less favorable) Hads values. The ability for the closed
packed alloy/metal surfaces to geometrically and electronically
accommodate increasing adsorbed atomic S coverage increases
with lattice parameter. The adsorption of a full S monolayer was
favorable on the more open four-fold sites on the Pd0.5 Cu0.5 Im3̄m
and P4 mmm (1 1 0) surfaces. The former surface was significantly
corrugated after relaxation. On the closed packed (1 0 1) and (1 1 1)
surfaces, the formation of a stable monolayer was only slightly
favorable on the largest Pd0.75 Ag0.25 alloy lattice, and not on any
of the respective metals. On the smaller lattice close-packed alloy
surfaces, the S interatomic interactions became stronger than the
S metal interactions, so that S coupling became more favorable. A
full S monolayer was not stable on the Pd0.5 Cu0.5 P4 mmm (1 0 1)
and Pd0.875 Au0.125 Pm3̄m (1 1 1) surfaces, and S coupling occurred
spontaneously. Sulfur coupling also dramatically increased the
favorability of adsorption of the equivalent to a full S monolayer
on the Pd0.75 Cu0.25 and Pd Fm3̄m (1 1 1) surfaces. Sulfur coupling
was previously identified for atomic modeling examining various
S overlayers on Pd (1 1 1) [14].
101
favorable subsurface sites for the absorption of 0.25 ML S within Pd
and its alloys, without and with concomitant Pd segregation. The
calculated S absorption enthalpies, Habs , are given in Table 5 and
can be compared with the adsorption enthalpies, Hads , given in
Table 3. The absorption of a low S coverage (equivalent to 0.25 ML
S) was not competitive to surface adsorption and was not significantly enhanced by Pd segregation. The Habs values for 0.25 ML
S absorption increased in favorability nearly linearly with increasing lattice parameter, becoming exothermic for the (1 1 1) surfaces
formed from the larger Pd0.875 Au0.125 , Pd0.75 Ag0.25 , and Pd lattices.
The most favorable subsurface site in each alloy was that which
maximized the absorbed S coordination with Pd. Typically, the S
was coordinated with 3 atoms in the subsurface layer and 3 or 4
atoms in the surface layer. In most alloys, S was positioned over
a fcc-type threefold hollow site in the subsurface layer [16]. The
exception was S in the Pd0.5 Cu0.5 P4 mmm (1 1 0) surface, where
the S was positioned over a fourfold hollow site in the subsurface
layer, coordinated to four atoms in the subsurface layer and five
atoms in the surface layer. Since the interstitial sites were not large
enough, S incorporation increased the spacing between the surface and subsurface layers, and dramatically corrugated the surface.
The Pd–S interatomic distances (often approaching 2.3 Å) of S penetrated into the subsurface layers closely mimicked the shortest
2.34 Å Pd–S bonds in the Pd4 S phase shown in Fig. 6a, where each
Pd is coordinated to two S and each S is coordinated to eight Pd
atoms [35].
Partial S absorption only becomes favorable in some cases at
higher coverages, where increasing S adsorbate repulsion induces
surface restructuring and aids S penetration into the subsurface.
To investigate the tendency for S incorporation at high coverage,
mixed mode S absorption and adsorption were modeled to originate from the equivalent of 1 ML S adsorption, also without and
with Pd segregation. Here, 0.25 ML S (1 atom) was absorbed in the
most favorable subsurface site and 0.75 ML S remained adsorbed
on the surface, previously shown to be favorable in the Pd (1 1 1)
surface [16]. Comparison of Tables 4 and 5 shows that this mixed
mode S absorption into the surface was exothermic and competitive to the formation of 1.0 ML S adsorbed layer for all of the Pd
and Pd alloy (1 1 1) or (1 0 1) closed packed surfaces, but not on the
more open Pd0.5 Cu0.5 (1 1 0) surfaces. However, S coupling at full
S coverage competes with S incorporation, and was found to be
the most favorable scenario for the Pd0.75 Cu0.25 and Pd0.875 Au0.125
(1 1 1) surfaces. The accompaniment of S incorporation with partial
Pd segregation was even more favorable, and alluded to the favorability of Pd4 S nucleation. The Pd segregation enabled the adsorbed
S to assume a similar coordinative environment to that in Pd4 S,
as shown for the 0.25 ML S absorbed/0.75 ML S adsorbed Pd segregated layer formed on the Pd0.75 Ag 0.25 (1 1 1) surface in Fig. 6b.
The calculations predict that partial S penetration will predominate
under conditions leading to high S coverage on the Pd (1 1 1) surface, and also on the Pd0.50 Cu0.50 P4 mmm (1 0 1) and Pd0.75 Ag0.25
Pm3̄m (1 1 1) surfaces when accompanied by Pd segregation.
3.3. Atomic S absorption
3.4. Overview
The formation of a S-bearing corrosion product surface layer
is contingent upon the initial penetration of S into the subsurface layer. Since the S atom is much larger than the Pd metal and
the Pd alloy interstitial sites, the lattices must be distorted by an
activated chemisorption process to accommodate S atoms below
the surface. In addition, the formation of a specific S scale phase
must be accompanied by selective segregation of one or more elements to the surface, such as Pd segregation for the formation of
a Pd4 S surface layer. This chemical and structural modification of
the surface is irreversible and can ultimately significantly block
H solubilization and diffusion processes [18]. To investigate this
phenomenon, atomic models were first made to identify the most
It has been demonstrated by the atomic modeling in this study
that most of the interactions of S with metal surfaces can be understood from the electronic structure of the host lattice, which is
derived from both the surface alloy composition as well as the
overall lattice structure. The mixed absorption and adsorption of
the equivalent of a S monolayer can, when combined with Pd segregation, result in the formation of a Pd–S overlayer, resembling
Pd4 S in structure. Pd–Ag alloys and pure Pd can favorably, irreversibly form the Pd4 S corrosion product. However, the mixed
adsorption/absorption configuration is not favorable for the equivalent of a S monolayer on most Pd–Cu and Pd–Au alloys. In most
102
S.M. Opalka et al. / Journal of Membrane Science 375 (2011) 96–103
Table 4
Structural characteristics and adsorption enthalpies, Hads , in kJ/(mol × atom) for 1.0 monolayer S adsorbed on the selected Pd alloys and constituent metals.
Alloy/metal composition
Space group
Surface crystal
plane
1.00 ML S Hads
(kJ/(mol × atom))
1.0 ML S with Pd
segregation Hads
(kJ/(mol × atom))
1.0 ML S equivalent
with S2 coupling Hads
(kJ/(mol × atom))
Pd0.5 Cu0.5
Pd0.5 Cu0.5
Pd0.5 Cu0.5
Pd0.75 Cu0.25
Pd0.75 Ag0.25
Pd0.875 Au0.125
Pd
Cu
Ag
Au
Im3̄m
P4 mmm
P4 mmm
Pm3̄m
Pm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
Fm3̄m
(1 1 0)
(1 1 0)
(1 0 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
−65
−104
Not stable
38
−3
14
18
73
76
Not stable
−60
−105
Not stable
41
24
N.A.
N.A.
N.A.
N.A.
N.A.
Not stable
−84
7
−28
−3
−19
−14
N.D.
N.D.
23
Table 5
Structural characteristics and adsorption enthalpies, Habs , for the absorption of 0.25 ML monolayer S on the selected Pd alloys and constituent metals alone and in a mixed
S monolayer configuration.
Alloy/metal composition
Space group
Surface crystal
plane
0.25 ML S Habs
(kJ/(mol × atom))
0.25 ML S with Pd
segregation Habs
(kJ/(mol × atom))
0.25 ML S absorb
and 0.75 ML S
adsorb Habs/ads
(kJ/(mol × atom))
0.25 ML S absorb and
0.75 ML S adsorb with
Pd segregation
Habs/ads
(kJ/(mol × atom))
Pd0.5 Cu0.5
Pd0.5 Cu0.5
Pd0.5 Cu0.5
Pd0.75 Cu0.25
Pd0.75 Ag0.25
Pd0.875 Au0.125
Pd
Im3̄m
P4 mmm
P4 mmm
Pm3̄m
Pm3̄m
Fm3̄m
Fm3̄m
(1 1 0)
(1 1 0)
(1 0 1)
(1 1 1)
(1 1 1)
(1 1 1)
(1 1 1)
36
62
40
40
−43
−8
−26
46
Not stable
43
−8
−12
N.A.
N.A.
−54
−71
−11
−5
−6
−6
−27
−60
−98
−21
−18
−25
N.A.
N.A.
cases, the coupling and partial desorption of S is more favorable on
these alloys. Many of these alloys exhibit a much larger S tolerance
than Pd–Ag alloys and pure Pd.
In addition to the electronic contributions to S interactions,
there are a number of other possible factors that may come into play
for Pd alloy membrane performance. Additional atomic modeling
is underway to evaluate finite temperature and pressure contributions on the relative configurations and coverage of S-containing
species. Another effect which can potentially influence the poisoning effect of S, is the occurrence of S in mixed gas compositions,
especially where H2 is the predominant gas species. A future publication will address S interactions with Pd alloys in mixed gas
compositions under finite temperature and pressure conditions
relevant to the selective separation and purification of H2 .
Additional factors that are more challenging to investigate by
atomic modeling may also play an important role. Kinetic restrictions could also play a role by preventing S forming Pd4 S layers at
some surfaces. This may render certain geometric structures inaccessible due to high diffusion barriers. For example, it is well-known
that defects like steps and kinks are important for the catalytic
activity of surfaces. However, such defects can easily correspond
to very large periodic unit cells, making accurate studies from first
principles prohibitively computationally expensive. This is also to
a certain extent the case for reconstructions, which can give similar effects. The only reconstruction we studied in this paper was
single atom Pd segregation, but reconstructions containing more
atoms and layers could easily contribute to the formation of more
complete Pd4 S layers.
Fig. 6. Comparisons of (a) Pd4 S bulk structure with (b) Pd–S overlayer structure formed by the absorption of 0.25 ML S and adsorption of 0.75 ML S, and Pd segregation on the
Pd0.75 Ag0.25 Pm3̄m (1 1 1). Only S-metal bonds are drawn, with a cut-off of 3 Å. In (b) bonds between 3 Å and 3.4 Å are shown as dashed lines. S, Pd and Ag atoms are shown as
light yellow, medium grey and dark blue balls, respectively. Only a part of the slab unit cell is shown in (b), to emphasize the formation of Pd4 S structures in the subsurface
layer. Thin solid lines designate the unit cells. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)
S.M. Opalka et al. / Journal of Membrane Science 375 (2011) 96–103
4. Conclusions
Atomic modeling was conducted to investigate the nature and
origin of S interactions with representative Pd alloy H selective
membrane candidates selected from the Pd–Cu, Pd–Ag, and Pd–Au
binary systems and their constitutive elements. In most cases, S
binding increased in strength with increasing alloy/metal lattice
parameter for a given crystallographic surface termination. This
trend corresponded to a decrease in surface energy. However, the
S interactions were not predicated on the alloy/metal surface site
geometries or lattice strain effects. The wide range of alloy/metal
chemistries confounded the complete correlation of their electronic properties, embodied in the density of states valence d-band,
with lattice size. Only the d-band breadth was found to track with
lattice size, and not the other electronic characteristics, like the
d-band center and number of electrons at the Fermi level.
Further analyses revealed S interaction with these alloys/metals
was predominantly controlled by electronic factors. A strong linear
correlation was found for increasing adsorption strength (decreasing adsorption enthalpy) of 0.25 ML S with increasing d-band center
energy towards the Fermi level. The electronic coupling of S p
orbitals bonding with alloy/metal d-bands in the adsorbate/slab
density of states split the lower energy p bonding state and the dband center further apart with increasing S bonding strength. Thus,
an even stronger correlation was found for the increasing adsorption strength with the increasing density of states energy difference
[d-band center – S p peak], [DBC-Spb ]. This universal relationship
provides an important capability to predict the S interactions of
closely related alloy/metal systems.
The modeled S interactions at higher coverages were indicative
of the susceptibility of the alloys to S corrosion. The adsorption of
1.0 ML S followed parallel trends to the adsorption of 0.25 ML S,
but was shifted to higher energies. A full ML of adsorbed S was
only stable on the open Pd0.5 Cu0.5 Im3̄m and P4 mmm (1 1 0) and
the Pd0.75 Ag0.25 Pm3̄m surfaces, and could not be accommodated
on the other surfaces. Partial S desorption and S coupling became
competitive at full coverage, especially for the Pd0.75 Cu0.25 Pm3̄m,
Pd0.875 Au0.125 Fm3̄m, and Pd Fm3̄m (1 1 1) surfaces. Partial absorption and incorporation of S also became competitive at full S ML
coverage for some alloy surfaces, especially when accompanied by
Pd segregation. This first step towards the irreversible nucleation of
Pd4 S formation was the most favorable for the Pd0.5 Cu0.5 P4 mmm
(1 0 1), Pd0.75 Ag0.25 Pm3̄m (1 1 1), and Pd Fm3̄m (1 1 1) surfaces.
The insights gained in this study are an important step towards
understanding the factors governing S tolerance in Pd based membrane materials. It will be complemented by a forthcoming study
in which temperature and pressure effects are taken explicitly into
account.
Acknowledgements
This publication was based in part upon work conducted by
the UTRC Advanced Palladium Membrane Team under the support
of the United States Department of Energy award number DEFC26-07NT43055. A grant of computational time from the NOTUR
consortium is acknowledged by OML. The UTRC authors would like
to acknowledge useful discussions with other team members, especially Zissis Dardas.
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