Effect of thiol self-assembled monolayers and plasma polymer films on dealloying of Cu–Au alloys

Abstract

We report on the influence of chemical surface modification on the selective dissolution and dealloying of
Cu3Au (111) in 0.1 M H2SO4. Cu–Au alloys serve as longstanding model alloys for dealloying and stress
corrosion cracking. Employing well-defined hexadecanethiol, mixed-aminobenzenethiol, and plasmapolymerized
hexamethyldisiloxane surface layers we obtain detailed atomic-scale insight in the stability of
the surfaces. The initial structural evolution of the modified surfaces is tracked by in situ X-ray diffraction
using synchrotron radiation. In comparison to the usual sequence of surface states on unmodified Cu3Au
(111) the modified surfaces develop a thicker and more stable passive-like Au-rich film below the critical
potential. In this regime we observe anodic shifts in the potentials of the surface structural transitions and
the suppression of an otherwise observed island morphology. This extreme stability of the modified
passive-like surfaces is further confirmed by ICP-MS coupled to a scanning micro-electrochemical flow cell
and ex situ SEM. Above the critical potential, the presence of the inhibiting protective layers leads to a
localized dealloying mechanism with an altered microstructure of the forming nanoporous film in the
form of micro-cracks. Understanding the stability of alloy surfaces is a prerequisite for further applications,
e.g. for lithography or sensors.

Full text

Cite this: RSC Advances, 2013, 3,
6586
Effect of thiol self-assembled monolayers and plasma
polymer films on dealloying of Cu–Au alloys
Received 20th November 2012,
Accepted 6th February 2013
DOI: 10.1039/c3ra22970j
www.rsc.org/advances
A. Pareek, G. N. Ankah, S. Cherevko, P. Ebbinghaus, K. J. J. Mayrhofer, A. Erbe and
F. U. Renner*
We report on the influence of chemical surface modification on the selective dissolution and dealloying of
Cu
3
Au (111) in 0.1 M H
2
SO
4
. Cu–Au alloys serve as longstanding model alloys for dealloying and stress
corrosion cracking. Employing well-defined hexadecanethiol, mixed-aminobenzenethiol, and plasma-
polymerized hexamethyldisiloxane surface layers we obtain detailed atomic-scale insight in the stability of
the surfaces. The initial structural evolution of the modified surfaces is tracked by in situ X-ray diffraction
using synchrotron radiation. In comparison to the usual sequence of surface states on unmodified Cu
3
Au
(111) the modified surfaces develop a thicker and more stable passive-like Au-rich film below the critical
potential. In this regime we observe anodic shifts in the potentials of the surface structural transitions and
the suppression of an otherwise observed island morphology. This extreme stability of the modified
passive-like surfaces is further confirmed by ICP-MS coupled to a scanning micro-electrochemical flow cell
and ex situ SEM. Above the critical potential, the presence of the inhibiting protective layers leads to a
localized dealloying mechanism with an altered microstructure of the forming nanoporous film in the
form of micro-cracks. Understanding the stability of alloy surfaces is a prerequisite for further applications,
e.g. for lithography or sensors.
1 Introduction
Self-assembled monolayers (SAMs) have been receiving wide
attention in the scientific community due to their variety of
technical applications.
1–4
The most common and widely used
approach to fabricate SAMs is based on the chemisorption of
alkanethiols onto Au surfaces. Atomic sulfur on gold surfaces
shows a variety of potential dependent surface reconstruc-
tions.
5
Thiol SAMs contain molecules anchored to the metal
surface through their head sulfur groups (–S) and the alkyl
chains stand upright or slightly tilted with respect to the
surface normal. By this approach well-ordered thiol SAMs are
easily fabricated through a simple chemisorption process (in
contrast to using other organic molecules
6
) and the assembled
layers show strong adhesion to metal surfaces. Furthermore,
their chemical composition and thickness can be precisely
controlled by proper selection of the adsorbates. Therefore, the
surface of metal substrates can be easily derivatized by
forming well organized adherent SAMs of thiols.
7–12
Thiol
SAMs have been widely discussed as corrosion inhibitors also
for Cu and Au systems.
12–15
In this context thiol SAMs on
copper substrates decrease the oxidation rate and it also
inhibits cathodic corrosion. Thiol SAMs on Au have been
furthermore used as model systems for corrosion protection or
for understanding the oxygen reduction reaction, thereby
addressing delamination of organic coatings on engineering
materials.
14
The defined SAM structures on both metals have
their hydrocarbon chains exhibiting a tilt towards the surface
normal (12utilt in Cu
8,16
and 30utilt in Au
1,8,16,17
). In this way
the properties of the substrate surface can be on one hand
determined on the atomic scale
18
and on the other hand
tailored and utilized for various applications including sensors
or corrosion inhibition.
Plasma deposition of polymers is another low-temperature
process widely used for producing ultra-thin organic layers
including coatings for corrosion protection.
19,20
Plasma coat-
ings have a highly cross-linked matrix and thus are largely
impermeable to humidity even in corrosive environments. In
addition to this, they also form some degree of covalent bonds
with the substrate, which facilitates good adhesion. The
plasma polymer films provide an increased degree of cross-
linking compared to the thiol films, although the bonding to
the substrate is less well-defined. Plasma polymer films have
been utilized for corrosion protection in many metals
including iron.
19
We used plasma deposition of the non-toxic
and non-explosive Si precursor hexamethyldisiloxane
(HMDSO) to form an ultra-thin protective coating on the
Cu
3
Au (111) surface and investigated its dealloying behaviour
analogous to thiol SAMs.
In regard to the substrate, the binary alloy system Cu–Au is
a classic model for investigating dealloying processes.
21–28
Its
Max-Planck-Institut fu
¨r Eisenforschung, Max-Planck-Straße 1, 40237 Du
¨sseldorf,
Germany. E-mail: renner@mpie.de
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constituent elements have a wide difference in their Nernst
potentials. With anodic polarization in an acidic environment
(just above the Nernst potential of Cu) the less noble element,
here Cu, selectively dissolves, leaving the alloy surface
enriched with the noble element, i.e. here Au. The Au
enrichment on the alloy surface mimics a passive-like effect
by a purely metallic surface until a certain potential, the so-
called critical potential (E
c
), is reached (Fig. 1a). Beyond E
c
,a
breakdown of the passive-like behaviour occurs with a massive
Cu dissolution resulting in a nanoporous network of residual
Au
22,23,29
showing a number of new functionalities.
30
The
lattice parameters of pure Cu and Au differ by 13%, hence it is
easily possible to probe the surface of Cu–Au alloys during
dealloying with X-ray diffraction techniques (Fig. 1b), which
allow to follow the structural surface evolution. Thiol SAMs on
Au alloys provide thus an atomistic access to obtain a
mechanistic understanding of dealloying and alloy corrosion,
and also of their inhibition and of materials stability in
general.
We previously reported the epitaxial surface films evolving
during the passivation regime of bare Cu
3
Au (111) surfaces in
0.1 M H
2
SO
4
on an atomic scale.
23–27
With first initial
dealloying a stacking-reversed (CBA) ultra-thin passive Au-rich
layer at low overpotentials and thicker pure Au islands at
medium overpotentials is formed. Close to the critical
potential we then reported
27
a transition of these structures
to substrate oriented (ABC) Au ligaments. The distinctive
sequence of structural transformations in pure sulfuric acid
opens the way to probe the influence of changes to the
electrolyte as well as of surface modifications (Fig. 1c). The
influence of halide additives to 0.1 M H
2
SO
4
on the selective
dissolution of Cu
3
Au (111) has been addressed in earlier
studies.
25,26
Large cathodic shifts of the initial structural
transitions and of the value of E
c
were reported. The influence
of alkanethiol SAMs on Cu
3
Au (111) surfaces was first
addressed by Moffat et al.
22
where the authors utilized
voltammetry and in situ electrochemical scanning tunneling
microscopy (EC-STM). They reported a shift in the critical
potential (E
c
) by approximately 150 mV in the anodic direction,
which was explained by the effect of thiol SAMs on the surface
mobility of Au atoms. Although the study did report the
inhibition effect of thiols, no detailed information on the
structural features involved in the passivation was provided.
Our initial insight on the hexadecanethiol-covered surface has
been reported recently.
27
Here, we provide a full report on in
situ X-ray diffraction studies of the effect of several organic
model inhibition layers on initial dealloying (selective dissolu-
tion). The applied thiol SAMs suppress Au surface diffusion
and allow thus to directly visualize corrosion inhibition.
Moreover, we report new crack morphologies near E
c
which
strongly depend on the surface orientation, i.e. on differences
in the atomic arrangements. The atomistic insight is a
prerequisite for further applying structured thiol films as for
example for a possible fabrication of 3-D hierarchical
nanoporous structures.
31
The effect of the surface modifica-
tion by hexadecanethiol (HDT) and mixed aminobenzenethiol
(ABT) is discussed in detail. Also, we provide first insight on
the effect of hexamethyldisiloxane (HMDSO) plasma polymer
coated surfaces.
2 Experimental
Cu
3
Au (111) single crystals and polycrystalline Cu
3
Au foils
were obtained from MaTecK GmbH, Germany. The single
crystals were oriented to ,0.1uand mechanically polished
down to 0.03 mm roughness. The UHV surface preparation
involved several sputter-anneal cycles until the desired clean
surface was achieved. A gentle electrochemical treatment
decreased the roughness of fresh surfaces considerably as
described previously.
26
The resulting surfaces were character-
ized using Auger electron spectroscopy (AES), low energy
electron diffraction (LEED) and atomic force microscopy
(AFM). This routine characterization ensured similar and
reproducible starting surfaces to be used for electrochemical
and in situ diffraction measurements.
Chemisorption of organic self-assembled monolayers
HDT (1-hexadecanethiol, 99% purity) as well as 2-ABT (2-
aminobenzenethiol, 99% purity) and 3-ABT (3-aminobenze-
nethiol, 96% purity) from Sigma-Aldrich were used without
any further purification. The HDT-modified surface was
obtained by immersing the Cu
3
Au (111) crystal in a 1 mM
hexadecanethiol solution in ethanol for several hours (mini-
mum 24 h) before the experiment. In order to produce a mixed
Fig. 1 a) Sketch of a linear polarization curve showing the passivation region
and bulk Cu dissolution at the critical potential E
c
during anodic polarization of a
pristine Cu
3
Au (111) surface in 0.1 M H
2
SO
4
. b) A schematic view of the in situ
X-ray diffraction cell set-up. c) Sketch of a thiol-SAM on Cu
3
Au (111) after initial
selective dissolution below E
c
.
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aminobenzenethiol (m-ABT) self-assembled monolayer
32
, 0.5
mM of 2-aminobenzenethiol (o-ABT) and 3-aminobenzenethiol
(m-ABT), respectively, was added to 50 ml ethanol. Henceforth
this solution is referred to as mixed ABT or m-ABT. The Cu
3
Au
(111) crystal was immersed about 30 min. In both cases, the
samples were rinsed with ethanol and dried in a stream of
nitrogen after thiol deposition.
Plasma polymer deposition
The deposition of hexamethyldisiloxane (HMDSO) plasma
polymer film was performed in a chamber with a linear
microwave source (Roth und Rau, Germany) at a power of 300
W and frequency of 2.46 GHz. The width of the plasma zone
was 150 mm and little plasma could be observed outside of
this zone. For the plasma deposition, hexamethyldisiloxane
(HMDSO, (CH
3
)
6
Si
2
O, purity .98.0%) procured from Fluka,
Germany was used as a liquid monomer without further
purification. For the deposition, we used ‘‘remote plasma’’
conditions i.e., a fine metal grid was used to separate the
samples from the direct plasma glow. Before the actual plasma
polymer deposition, the plasma chamber was evacuated down
to 10
22
mbar and refilled with argon several times to avoid
contamination. An Argon flow rate of 80 sccm was used during
operation. The samples were then plasma cleaned by moving
them above argon plasma at the speed of 1 mms
21
. Usually,
this procedure is performed in oxygen plasma,
19,20
however,
here we used Argon in order to avoid further oxidation on the
Cu
3
Au surface. Immediately after cleaning and activation, a
mixture of Argon and HMDSO in the ratio of 20 : 5 was
produced at a chamber pressure of 0.04 mbar. The HMDSO
precursor was heated at 40 uC and then injected into the
chamber. During plasma deposition the substrate was then
moved at a constant speed of 70 mm s
21
. This high speed was
chosen in order to restrict the film thickness to a few nm (y2–
3 nm). The details of the plasma chamber are described
elsewhere.
20
In situ surface X-ray diffraction
The samples were transferred in air to a thin-film in situ X-ray
diffraction cell (Fig. 1b). The details of the cell are described
elsewhere.
33,34
A thin Mylar foil (6 mm, SPI supplies) was used
to seal the cell leaving a thin layer of electrolyte above the
crystal surface during measurements. At regular intervals and
while changing potentials the electrolyte volume above the
sample was increased to allow better electrochemical condi-
tions. The potential of the sample was varied by a computer-
controlled potentiostat (PAR273). Commercial Ag/AgCl micro-
electrodes and Pt wire were used as reference and counter
electrodes, respectively. All potentials indicated in the X-ray
data are referred with respect to this Ag/AgCl reference
electrode at 298 K. After crystallographic alignment of the
sample with respect to the incident X-ray beam (Fig. 1b), the in
situ thin film cell was filled with electrolyte at an initial
potential of 2100 mV (starting potential well below the Cu
Nernst potential of +100 mV). Cu
3
Au crystallizes in the cubic,
fcc-like (a
0
= 375.3 pm), L1
2
structure at temperatures below an
order-disorder transition temperature T
c
= 390 uC. The cubic
crystal structure is obtained by the stacking of close-packed
atomic planes (ABCABC…stacking). X-Ray diffraction measure-
ments were performed using the 6-circle diffractometer of the
beamline ID32 at the European Synchrotron Radiation Facility
(ESRF) and at ANKA, Karlsruhe, Germany. At the ESRF, an
X-ray energy of 23 keV, which corresponds to wavelength of l=
53.9 pm was used. The beam was focused to a vertical spot size
at the sample of 30 mm by a set of compound refractive lenses
(CRL) at a focal distance of 15 m from the sample. A grazing
incidence angle of 0.2utowards the surface plane was chosen.
At ANKA, the energy was set to 10 keV (l= 124 pm). The
scattering vector qis defined as q=(k
f
2k
i
), where (k
i
) and (k
f
)
are the incoming and diffracted wave vectors, respectively. To
describe qand the reciprocal space of the L1
2
–Cu
3
Au crystal,
and the scattering vector q, we chose the commonly used
surface unit cell with two unit vectors a
1
and a
2
lying in the
surface plane and the third unit vector a
3
pointing along the
surface normal (a
1
=a
2
=!2a
0
= 531 pm, a
3
=!3a
0
= 650 pm, a=
b=90uand c= 120u). This leads to a hexagonal reciprocal
lattice unit cell of size a
1
*=a
2
*= 13.670 and a
3
*= 9.666 nm
21
,
and a*=b*=90uand c*=60u.
Scanning flow cell ICP-MS
The copper dissolution was traced by a scanning flow cell
(SFC) coupled with inductively coupled plasma mass spectro-
meter (ICP-MS, NexION 300X, Perkin Elmer), described in
more details elsewhere.
35
The electrolyte flow rate was 139 mL
min
21
and the geometrical contact area of the working
electrode was 0.01 cm
2
. The transient signals of the dissolved
63
Cu are monitored versus an internal standard of 7.5 ppb
89
Y
at 50 ms dwell time and 5 sweeps per reading.
3 Results and discussion
3.1 Initial electrochemical dealloying behaviour
The schematic anodic polarization curve (Fig. 1a) obtained
from Cu
3
Au (111) in 0.1 M H
2
SO
4
aqueous solution includes
also the nomenclature used in this paper. When the pristine
Cu
3
Au (111) alloy is contacted to the electrolyte, selective
dissolution of Cu starts above the Cu Nernst potential (+100
mV). As was shown in earlier works, the initial selective
dissolution of Cu from clean alloy surfaces in 0.1 M H
2
SO
4
leads to the formation of an ultra-thin (3 ML) Au-rich film,
which protects (passivates) the surface from fast dissolution. A
further increase in potential to y300–400 mV immediately
increases the thickness of this passive-like Au-rich layer and
leads to Au-islands (10–15 ML). The formed Au layer protects
the surface until a breakdown at the critical potential (E
c
)
occurs. At E
c
, massive (bulk) Cu dissolution sets in. The initial
sweep of a cyclic voltammogram of a HDT-modified surface
showed an anodic shift of y100 mV in its critical potential
and surface cracks.
27
Fig. 2a shows a cyclic voltammogram
obtained on an m-ABT-modified Cu
3
Au (111) surface. In the
first cycle of the voltammogram the ‘‘critical’’ potential with a
value of about 950 mV is shifted anodic by about 110 mV
compared to the behavior in pure sulfuric acid
23,24
where the
critical potential was measured to be about 840 mV. The steep
rise in current indicates a massive Cu dissolution from the
alloy surface at E
c
.
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At this potential, dissolution sites (cracks) are created on the
alloy surface, exposing a fresh non-modified, thiol-free alloy
below the surface. Therefore, in the second cycle, a new value
of E
c
appears back at lower potential and the current rises
steeply already after reaching 890 mV instead of 950 mV.
Fig. 2b shows an SEM image (inset is a higher magnification
image) of the surface of a polycrystalline Cu
3
Au alloy after
cyclic voltammograms to potentials above E
c
. The surface is
finally covered with a high density of cracks of about 1–2 mmin
diameter. In contrast to the homogeneous nanoporosity
usually obtained on clean surfaces
28,36
a pitting-like behaviour
thus occurs. Since there is a large volume change during
dealloying and formation of nanoporous Au,
37
the microcracks
also reflect the formation of large strain created close to the
localized nanoporous ‘‘pit’’ or defect sites. Clearly, the SAMs
greatly stabilize the surface at large as no porosity is observed
on the remaining thiol-covered crack-free surface areas even
after applying such high anodic potentials. The behaviour of
the m-aminobenzenthiol-modified surfaces is thus nearly
identical to the hexadecane-modified surfaces where initial
results have been described earlier.
27
The passivation break-
down which is observed from the steep current rise at E
c
is
visibly caused by dissolution from defect or crack sites and the
true E
c
maybe eventually even higher. At that point the exact
mechanism of crack initiation and nature of the location
where the cracks initiate is not entirely clear yet. Illustrative is
the striking dependence of the crack density on the surface
orientation which is evident on respective grains of different
orientation on a polycrystalline sample as shown in Fig. 2b.
Apparently it is the detailed structure of the specific thiol SAM
on different crystalline surface orientations of the grains that
largely influences the actual crack density. It may be therefore
concluded that neither an intrinsic defect type of the bulk
alloy, nor an impurity to the thiol molecule supply is the
determining factor in the crack nucleation. Patterns of thiol-
film areas can be also applied by m-contact imprinting (m-cp).
Using m-cp we showed in a first quick test with a new in situ
AFM cell
31
that dealloying could be localized in a controlled
way and we could avoid crack formation (Fig. 2c).
Detailed substrate-thiol-SAM interface structures are so far
mainly addressed for Au (111) surfaces and a few studies have
been done on Au (100). The structures are mostly still critically
debated.
38–41
The stability of the thiol film itself depends, due
to steric effects during its formation, critically on the atomic
next-neighbour distances along the surface or the step density,
i.e. on the surface orientation.
The inhibition effect of the thiol layers and the stability of
the forming Au-rich surface films have been in addition
addressed by ICP-MS coupled to a micro-electrochemical SFC.
The design of this setup, which is described in detail
elsewhere
35
is following earlier approaches
42,43
but is opti-
mized for direct online monitoring of dissolution with the
prerequisite that the actual elemental dissolution rate not
exceed limits set by the sensitive detection of the mass
spectrometer. For this reason, the anodic polarizations were
limited to values slightly below the critical potential. Fig. 3
shows the measured elemental Cu ions dissolved from Cu
3
Au
(111) surfaces into the electrolyte during the initial potential
sweep for a bare surface as well as surfaces modified with
hexadecane- and mixed aminobenzenethiol. All measured
dissolution rates are relatively small and correspond at initial
contact with electrolyte for the bare surfaces in total to about
2–3 monolayers (ML) of Cu
3
Au (i.e. about half a monolayer of
residual pure Au might form here on the surface) while the
hexadecane-thiol and the mixed aminobenzenethiol-modified
Fig. 2 a) CV obtained on a m-ABT modified (stabilized) Cu
3
Au (111) surface in
0.1 M H
2
SO
4
at a scan rate of 10 mV s
21
. Cracks formed at the end of the 1
st
cycle expose fresh, thiol-free alloy surfaces leading to a lower E
c
in the 2
nd
cycle.
b) SEM image of a polycrystalline Cu
3
Au surface after similar cyclic voltammo-
grams. The inset shows a typical microcrack. c) First results of controlling
localized dealloying by application of thiol m-cp micropatterns. Arrows mark the
dealloyed sites.
Fig. 3 The copper dissolution profile as measured by the ICP-MS on a bare and
modified-Cu
3
Au (111) surfaces in deaerated 0.1 M H
2
SO
4
.
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surfaces show no measurable dissolution at the lower over-
potentials, respectively. At low and intermediate overpotentials
(above the Cu equilibrium potential Cu
eq
) the dissolution from
the bare surfaces corresponds to an additional 5–7 ML of alloy
material. Compared to the thiol-modified surfaces, the bare
Cu
3
Au (111) shows a relatively high initial Cu dissolution peak.
While for the thiol-modified surfaces, an initial Au-rich layer
had already formed (due to the selective adsorption of the thiol
to Au) and the bare surface forms the entire ultra-thin Au-rich
film in the first potential sweep. The lowest rate is seen for the
case of hexadecane-modified surfaces with no observed
dissolution even at intermediate potentials. Here, apparently
the equilibrium is reached in the slowest way which matches
to later shown (slower) in situ X-ray diffraction data, which in
turn show proceeding structural changes for HDT-modified
surfaces compared to the more stable case of m-ABT. The scan
rate for the voltammograms and the SFC-ICP-MS is much
larger than for the below in situ X-ray experiments.
3.2 Structural evolution of hexadecanethiol-modified (HDT)
Cu
3
Au surfaces
Fig. 4a depicts the reciprocal space coordinate system (HKL)
with respect to the Cu
3
Au (111) surface unit cell, with the H
and Kcoordinates in the surface plane and the L-coordinate
oriented along the surface normal. To follow the development
of the different structural features, radial scans and L-scans
were performed at several potentials along the in-plane (220)
and out-of-plane (201) directions. Fig. 4b shows a set of in-
plane HK scans (H=K= 1.7 to 2.2) close to the substrate Bragg
peak (2, 2, 0), recorded for the hexadecanethiol-modified
(HDT) surface. The potentials were gradually increased in
steps of 50 mV and several in-plane scans were recorded. For
clarity, only a few scans with evident shifts are shown in Fig. 4
and Fig. 5. At 100 mV, i.e. at the Cu Nernst potential, a broad
peak at (1.9, 1.9, 0.05) is observed as a result of Au segregation
due to thiol adsorption and initial Cu dissolution, which is the
signature of the stacking inverted ultra-thin Au-rich layer.
23,24
Upon further increase in potential, the intensity of the peak
increases (with a narrowing of the peak width) and a shift
towards lower H,Kvalues, reaching at 750 mV at the (1.84,
1.84, 0.05) HKL position corresponding to pure Au. This
potential is about 350 mV more anodic than the corresponding
potential (400 mV) obtained on a thiol-free, bare Cu
3
Au (111)
surface.
23,24
The lateral size of these structural features can be
estimated by considering the peak width of the radial scans.
The respective size of the domain deduced from the first curve
at 100 mV is y6 nm which increases to y10 nm at 750 mV
and to y12–13 nm at still higher potentials.
The position and width of the respective peak in the radial
scans provides information about the lateral structure and size
of the Au-rich passive layer. At specific positions the stacking
sequence can be distinguished along the out-of-plane L
direction (0KL or H0L), i.e. in directions parallel to the (111)
surface normal. Fig. 4c shows reciprocal space (KL) maps
around the (0, 2, 1) and (0, 2, 2) substrate Bragg peak positions
at three different conditions (dry sample, at 350 mV and at 750
mV polarizations).
The fine streak of intensity at K= 1.9 r.l.u. in the first map
(dry sample) indicates an already existing initial ultra-thin Au-
rich film (thickness) (CBA). The existence of this initial film at
the beginning can be explained by the segregation of Au
during the SAM chemisorption process, as the Au–S bond is
preferred over a Cu–S bond. At 350 mV, the peak intensity
corresponding to the Au-rich film further grows (CBA intensity,
peak here at Ly1), accounting for the increase in film
thickness. However, no further new peak is observed at this
potential. At 750 mV a new, sharp intensity peak appears at
Ly2, a position which corresponds here to pure epitaxial,
ABC-stacked (substrate-oriented) Au. The respective ABC
(Ly2) Bragg peaks are sharper in L-direction than the CBA
(Ly1) peaks and are almost isotropic in shape. Therefore, the
newly emerged ABC crystallites (ligaments of a nanoporous
structure) are much thicker than the initial, inverted Au CBA
ultrathin layer and their isotropic peak width points to the
porous Au structures that form above the critical potential.
28,44
The structural evolution was also followed by measuring a
set of out-of-plane L-scans with gradual increase in potential.
27
Fig. 4 In situ X-ray diffraction results of a hexadecanethiol modified Cu
3
Au (111)
surface. a) Sketch of reciprocal space of Cu
3
Au (111). H,Kand Ldirections are
shown along with the positions of ultra-thin Au rich layer (CBA) and substrate
oriented (ABC) stacking peaks in HKL space. b) In-plane scans recorded at
different potentials showing the growth of a thin Au-rich (passive) layer. With
an increase in potential, the position of the layer peak shifts towards lower H,K
values (towards pure Au). c) Out-of-plane reciprocal space maps showing the
epitaxial CBA-Au and ABC-Au positions along with the Cu
3
Au (0, 2, 2) peak. An
ultra-thin Au-rich layer with inverted stacking (CBA-Au) is present at 0 mV,
grows at 350 mV, and substrate oriented Au-ligaments (ABC- peak) appear at
750 mV. Even at these high potentials, the initial CBA-Au layer is stable and co-
exists with ABC-Au ligaments.
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Fig. 5a shows here the out-of-plane scans along the (H,0,L)
intensity rod. Corresponding H-values in (H,0,L) were
determined from the peak positions of radial scans at the
respective potentials. The first scan at low potential (200 mV)
shows the intensity at Ly2 corresponding to the first CBA-
stacked ultra-thin Au-rich film. The peak width estimates to
about DL= 1, which corresponds to approximately 3 mono-
layers (ML) of Cu
3
Au (1 ML = 0.2167 nm). The thickness of the
initial passive Au-rich film on the thiol-modified surface is
similar to that of the bare Cu
3
Au (111) surface, which indicates
that initial Cu dissolution can proceed homogeneously also
with the presence of thiol SAMs. With successive increase in
potential, we observed a small change in thickness of the CBA
stacked film, from 3 ML at 200 mV to approximately 4 ML at
700 mV. This observation is different from the bare Cu
3
Au
(111) surface, where at medium overpotentials, the CBA layer
grows to form (CBA-stacked) Au-islands with a larger thickness
of approximately 13–15 ML.
24
However, a new peak at Ly1 emerged at 750 mV, which
corresponds to the substrate-oriented ABC Au ligaments. As
was evident in the shown KL maps, the new ABC peak at Ly1
is intense and sharper than the peak corresponding to the
ultrathin CBA film. In the presence of HDT SAMs, the
substrate-oriented crystallites (ABC) are finally emerging 300
mV higher than in the case of a bare Cu
3
Au (111) surface.
24
The ABC stacked Au ligaments emerge as a final growth stage,
very close to the critical potential. The cyclic voltammogram
on a HDT-modified surface showed an anodic shift of y100
mV in the critical potential.
27
Even at still higher potentials i.e.
850 mV, the thickness of inversely stacked (CBA) Au-rich
passive layer remains nearly constant (y6–7 ML) and it co-
exists with substrate-oriented (ABC) Au ligaments. However, in
the case of thiol-free, bare Cu
3
Au (111) surfaces, the ABC
crystallites continue growing while CBA stacked Au-islands
(y15 ML) are then vanishing, at higher overpotentials very
close to the critical potential.
25–27
3.3 Structural evolution of mixed-aminobenzenethiol-modified
(m-ABT) surfaces
The dealloying behaviour of mixed-aminobenzenethiol-cov-
ered (m-ABT) Cu
3
Au (111) surfaces was investigated similarly
to that of the HDT-covered Cu
3
Au (111) surface. Fig. 5b shows
the corresponding set of L-scans measured for m-ABT covered
surfaces. The first curve at 250 mV also shows the intensity at
Ly2 corresponding to CBA stacked Au-rich film. The thickness
of this initial CBA Au-rich passive layer, estimated at this
potential, is around y4 ML, which is comparatively larger
than at the same stage on a HDT-covered Cu
3
Au (111) surface.
At successive potentials, we first observe only a slight increase
in the intensity of the CBA-stacking peak (at Ly2). It was only
at very high overpotentials of 900 mV that the peak at Ly1
(ABC stacking) emerged. This new ‘‘ABC’’ Bragg peak is
sharper than the respective peak for the inversely stacked
‘‘CBA’’ layer, however a comparison with Fig. 5a reveals that
this peak is less intense than the corresponding ABC Bragg
peak obtained from the HDT-covered Cu
3
Au (111) surface.
Furthermore, it emerges at a potential even 150 mV higher
than the corresponding potential on the HDT-covered Cu
3
Au
(111) surface and is even above the critical potential of the
thiol-free Cu
3
Au (111) surface. This result indicates a stronger
inhibition effect in the presence of m-ABT self-assembled
layers. Similar to the HDT-covered surface, the initial passive
Au-rich film (CBA stacked) co-exists with the newly formed
substrate-oriented (ABC) Au crystallites and does not vanish
even at very high potentials of 1 V (Fig. 5b). Interestingly, the
width of the respective peak is larger and indicates a smaller
corresponding ligament size. This may be related to higher
dissolution kinetics at the sudden breakthrough at higher
potential.
Fig. 6a shows respective reciprocal space (HL) maps around
the (2, 0, 1) and (2, 0, 2) substrate positions at 0, 900 and 1000
mV. The HL-plane is rotated by 60ucompared to the KL-plane
presented above. The first map measured at 0 mV from a
m-ABT modified surface already shows a small intensity at
Ly2 corresponding to CBA stacked Au-rich passive layer,
which is similar to the HDT-covered Cu
3
Au (111) surface. This
initial layer as well points to Au segregation during the thiol
chemisorption process. The second map was measured at 900
mV, where we observed intensity at Ly1 corresponding to ABC
stacked Au-crystallites. The epitaxial peak at Ly1 shows a
pronounced mosaic spread indicating that not all domains of
the newly formed Au crystallites are of the same orientation.
This observation might be related to the above mentioned
smaller ligament size. The difference compared to the HDT-
covered surface film is clearly visible in the maps of Fig. 4 and
6. A rocking scan is a scan in the transverse direction (see
arrow in second map) and represents the so-called mosaic
spread of the respective peak rather than the corresponding
crystal size (Fig. 6b). The peak width of the rocking scan is thus
a measure of an orientation distribution of the new crystallites
which contribute to the scattering intensity. Fig. 6b shows the
respective rocking scans on the CBA (Ly2) and the ABC (Ly1)
peaks. The rocking scan measured across the CBA peak has a
width of approximately 1uwhile that across the ABC peak has a
much larger width of approximately 4u.
Fig. 5 Out-of-plane Lscans recorded as a function of applied potentials on the
a) hexadecanethiol modified and b) mixed-aminobenzenethiol modified Cu
3
Au
(111) surface in 0.1 M H
2
SO
4
. The broad peak corresponds to inversely stacked
(CBA) ultra-thin Au rich layer. The peak corresponding to substrate-oriented Au
crystallites (ABC Au) appears narrow at 750 mV in hexadecanethiol and at 900
mV (broad) in mixed-aminobenzenethiol.
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3.4 Structural evolution of plasma-polymer-modified surfaces
Plasma-polymerized HMDSO films provide covalent binding to
the substrate (good adhesion) but the binding is less
pronounced than for thiol SAMs. Plasma polymer films
provide another interesting corrosion-inhibiting system and
we present here initial results on their influence on dealloying
of Cu
3
Au (111) in comparison to the above applied thiol
systems. We deposited an ultra-thin HMDSO coating (,3 nm)
and verified the deposition by optical spectroscopy. Fig. 7a
shows a series of out-of-plane (0, K,L) scans for the HMDSO
plasma-polymerised Cu
3
Au (111) surface, where (0, K, 1) and
(0, K, 2) are Bragg peak positions for inverse stacking (CBA)
and substrate stacking (ABC), respectively. The KL-plane is
rotated by 60ucompared to the HL-plane. Similar to thiol-
modified Cu
3
Au (111) surfaces, at 200 mV, the curve shows
small intensity at the inverted CBA (here Ly1) peak position.
In Fig. 7a, the intensity at this position (CBA) for the first 2–3
curves is still small and a streak of background intensity from
the close Cu
3
Au Bragg peak is easily visible. However, from the
peak-width, the thickness of this initial CBA Au-rich passive
layer was roughly estimated to be y3 ML at 200 mV, which is
in accordance with the previously discussed results on thiol-
modified surfaces. At further, higher potentials, the intensity
of the CBA stacking peak continues to grow, with a
corresponding thickness increasing from y3toy5MLat
550 mV. The first scan measured at 550 mV does not show any
additional intensity. However, after a duration of y2h,we
observe a small intensity appearing at Ly2 (see second scan at
550 mV). At successive higher potentials, the peak correspond-
ing to the substrate orientation (ABC, Ly2) continues to grow
in intensity. A closer look at scans measured from 550–650 mV
is shown in Fig. 7b, where the emergence of a new ABC peak at
550 mV is clearly evident at higher L-values (Fig. 7b, Scan 2 at
550 mV). At 550 mV, also a sudden reduction in width of CBA
stacking peak was also observed, which corresponds to a
change in thickness from y5toy9 ML between the two
scans. With a further increase in potential beyond 600 mV, the
CBA stacking peak does not show any significant change in
intensity or width. Similar to the thiol-modified surfaces, the
CBA layer is, in this case, as stable at higher overpotentials and
co-exists with the additional substrate-oriented (ABC) crystal-
lites.
Fig. 7 a) Out-of-plane Lscans close to Au Bragg positions on the HMDSO
plasma-polymerised Cu
3
Au (111) surface in 0.1 M H
2
SO
4
. The broad peak
corresponds to inversely stacked (CBA) ultra-thin Au rich layers. The peak
corresponding to substrate-oriented Au crystallites (ABC Au) appears at 550 mV
and grows. b) Extended L-scans recorded at 550–650 mV. c) Out-of-plane
reciprocal space maps showing the epitaxial CBA-Au and ABC-Au positions
along with the Cu
3
Au (0, 2, 2) peak of a HMDSO plasma-polymerised Cu
3
Au
(111) surface. At 2100 mV, mainly the (0, 2, 2) peak is visible. The CBA-Au peak
is visible in the map at 550 mV, and both inversely stacked (CBA-Au) and
substrate oriented Au-ligaments (ABC- peak) are co-existing at 800 mV. d) SEM
image of a rough area of initial localized dealloying. The arrow points at the
unaltered area surrounding the corroded spot.
Fig. 6 a) Out-of-plane reciprocal space maps showing the epitaxial CBA-Au and
ABC-Au positions along with the Cu
3
Au (0, 2, 2) peak in a mixed-
aminobenzenethiol modified Cu
3
Au (111) surface. An ultra-thin Au-rich layer
with inverted stacking (CBA-Au) is present at 0 mV. The CBA-Au peak continues
to grow, while substrate oriented Au-ligaments (ABC- peak) appear at 900 mV.
Both peaks still co-exist at 1000 mV. b) Rocking scans recorded at CBA-Au
(narrow, left) and ABC-Au (broad, right) at the positions obtained from the HL
map at 900 mV.
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In Fig. 7c we show corresponding reciprocal space maps
(here: KL) for plasma-polymerised Cu
3
Au (111) surfaces at
three different potentials. Immediately after filling up the cell
with electrolyte, at 2100 mV, clear intensity was observed only
at Cu
3
Au Bragg peak positions. However, some diffuse low-
intensity streaks are here initially present and point to some
limited structural rearrangement of the surfaces after applica-
tion of the ultra-thin HMDSO coating. At 550 mV, the intensity
at the CBA Au Bragg peak position is clearly visible (Ly1),
however, in this map we do not observe any intensity with
respect to the substrate-oriented ABC Au position. The map
was recorded immediately after increasing the potential to 550
mV. The first scan at 550 mV, shown in a set of out-of-plane
scans (Fig. 7b, Scan 1 at 550 mV) is along the 0KL direction,
crossing the peak positions of the corresponding map. Our
experiments are potential-dependent as well as time-depen-
dent.
24,25
The third map was recorded at a higher overpotential
of 800 mV, where the peak corresponding to ABC substrate-
oriented crystallites is present at Ly2. The peak at Ly2is
again isotropic in shape. The corresponding ligaments are
thus again thicker than the CBA-stacked layer and can be
again associated to forming nanoporous Au. Rough and later
porous areas develop emerging from localized spots of
dealloying. The respective areas show first small nano-sized
cracks when they grow larger as exemplified in Fig. 7d.
4 Discussion
The stability of the self-assembled monolayers (SAMs) on the
Cu
3
Au (111) surface plays a crucial role for their functionality,
i.e. here the inhibition of the dealloying process. At the same
time, SAMs on single-crystal surfaces are well-defined interface
systems which provide key test cases for mechanistic under-
standing and atomic scale insights. The strong variation of the
crack density on different grain surfaces of a polycrystalline
substrate shown above (Fig. 2b) directly illustrated that the
surface orientation and therefore the local atomic arrange-
ments of substrate and SAM are important parameters.
Impurities or structural crystal defects are thus not the
determining factors in the initial crack formation. Structural
models for thiol-based SAMs are available mainly only for Au
(111) and Au (100) surfaces
38
and are still often debated. The
formation and reactivity is also potential-dependent.
14,45
Although addressing the electrochemical stability and reactiv-
ity of thiol-covered surfaces is thus challenging, modern
atomic scale characterization as well as theoretical simulation
techniques will be able to provide increasing insight and
understanding.
Surface-sensitive X-ray diffraction employing synchrotron
radiation is, in addition to real space (localized) scanning
probe microscopy, an important in situ characterisation
technique. Electrochemical interface structures depend criti-
cally on potential or chemical environment and thus require
often in situ studies.
Surface reconstructions,
46–48
adlayer structures
49,50
as well
as initial stages of electrodeposition
51,52
or selective dissolu-
tion
53,54
can be naturally addressed by diffraction techniques.
Due to the strong sulphur–Au bond and pronounced inter-
molecular interactions within the SAM, thiol films can
substantially decrease the surface diffusion rate which is
probably the most important mechanism of dealloying
reactions.
21–24,27,44,55
Moffat
22
first studied thiol adlayers on
Cu–Au surfaces and their effect on dealloying, but did not
provide detailed structural analysis at the time as is possible
now with modern characterization techniques. During deal-
loying of Cu
3
Au (111), the initial Au-rich surface layer forms
with an inversed stacking sequence or rotated film orientation
which requires substantial rearrangement of the surface and
near-surface atomic structure. In contact with pure sulphuric
acid or with halide additives, the initial ultra-thin Au film
transforms at intermediate potentials (still far below E
c
)to
thicker Au islands covering the surface. The island structure is
following the inversed orientation of the initial surface film.
This second step of island formation below E
c
is completely
missing for the thiol-modified, as well as the plasma polymer-
coated surfaces. Instead, the initial ultra-thin films are
stabilized and protect the surface up to much higher
potentials. We associate the substrate-oriented Au structures
with initial ligaments of a nanoporous morphology forming
close to the critical potential. Surface diffusion is the most
important mechanism in dealloying. The formation of
substrate-oriented Au is firstly observed as marked by arrows
in Fig. 8a at 900 mV (for m-ABT), 750 mV (for HDT), and at 550
mV (for HMDSO), respectively, compared to 400 mV in pure
0.1 M H
2
SO
4
and even lower with halide additives.
25,26
We thus
observe that the cross-linked mixed-ABT film allows even less
lateral surface diffusion than a HDT film. With a similar
breakdown potential E
c
, the nanoporous structures are then
formed more abruptly with an applied m-ABT film, a fact that
might be then also reflected in the clearly visible larger mosaic
spread. Although the breakdown potential is similar for both
employed thiol films and only about 100 mV higher than for a
clean surface, the X-ray data and the ICP-MS experiment
reveals, especially for m-ABT, a larger stability below the
critical potential. Below the critical potential, surface diffusion
plays the sole important role while at the critical potential also
interlayer exchange processes become important for the Cu–
Au system.
27
Interesting to note is also the significant growth
of the initial layer thickness at the point where the substrate-
oriented ligaments start to form. Fig. 8a shows the peak width
analysis of L-scans through CBA Bragg peaks in the case of
modified surfaces and the deduced number of Au monolayers.
The maximum thickness of the initial structurally-inversed
layer before is limited to 3–5 ML but with the occurrence of the
substrate-oriented signal this layer grows up to 6 ML for HDT
and 9 ML for HMDSO. For m-ABT the increase in thickness is
hardly visible. We propose here that with the first porosity
formed by dissolution of Cu from deeper layers, additional Au
atoms are reaching the layer surface and are incorporated in
the film in addition to passivating the dissolution site. For this
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process, a diffusion mechanism of the additional Au atoms
along the metal–organic interface may be crucial. The
presently discussed surface modification might thus contri-
bute to understand the stability of nanoparticles, e.g. in
catalysis
56
. Here, the largest effect is seen for the Au-film/
HMDSO interface followed by Au-film/HDT. There is no visible
increase for the Au-film/m-ABT interface (Fig. 8b).
In the case of mixed-ABT, the two used isomers o-ABT and
m-ABT have similar molecular size and the Cu
3
Au sample was
immersed in the solution for a very short duration of 30 min.
Therefore, both types of molecules will adsorb on the Cu
3
Au
(111) surface in equal probability. Hence, on average it is
reasonable to assume that the o-ABT and m-ABT molecules
adsorb alternatively on the surface. Fig. 8b shows the
schematic of such an adsorption process on the Cu
3
Au (111)
surface. At higher overpotentials (700–800 mV), the alternate
sequence of m-ABT and o-ABT molecules will polymerise by
electrochemical oxidation as was reported by Kuwabata et al.
32
This electrochemical polymerisation is regarded as synon-
ymous with that of the aniline (polyaniline formation) i.e. head
and tail coupling at the - para positions. Fig. 8c shows the
schematic of this polymerisation through their aniline units.
Such a polymerisation at the higher overpotentials could
explain the observed anodic shift and higher stability as well
as the lower permeability of additional Au atoms along the Au-
layer/m-ABT interface. Using such layers further in litho-
graphic approaches is a well-established scenario.
3
5 Conclusions
We reported on the dealloying behaviour of chemically
modified Cu
3
Au (111) surfaces in 0.1 M H
2
SO
4
solution and
provided details on an atomic-scale interplay of a functiona-
lized surface (here: an applied inhibition layer) with a well-
defined single-crystal (grain) surface. The development of
surface crystallographic features, namely the Au-rich passive-
like reversed (CBA) film and the substrate-oriented Au-
ligaments (ABC) with respect to potential, was followed by
using in situ X-ray diffraction on Cu
3
Au (111) surfaces
modified with hexadecanethiol (HDT), mixed aminobenze-
nethiol (m-ABT) and a HMDSO plasma polymer. We observed
an immense stabilisation of the passive-like region. The shift
in potential below E
c
of the first emergence of porosity was 350
mV for hexadecanethiol, 500 mV for m-ABT, and 150 mV for
plasma polymer layers with respect to the behavior of a bare
Cu
3
Au (111) surface. For all chemically modified Cu
3
Au (111)
surfaces, the initial passive Au-rich layer does not vanish and
remains stable even at high overpotentials. Surface or interface
diffusion is suppressed most efficiently by the m-ABT layer. In
addition, the usual thicker Au-islands at medium overpoten-
tials were not observed. The results corroborate the suppres-
sion of surface diffusion by thiol self-assembled layers and
highlight at the same time its role and importance in free
dealloying. The application of thiol and plasma polymer films
provides a route to obtain localized dealloying.
Acknowledgements
G.N.A. acknowledges a scholarship by IMPRS-SURMAT.
Synchrotron radiation beamtime was provided by ESRF,
Grenoble,andANKA,Karlsruhe.Weacknowledgethetechnical
supportofL.Andre,H.Isern,J.Roy,andJ.Zegenhagenat
beamline ID32. We also acknowledge the assistance provided by A.
Stierle in using the MPI-MF surface diffraction beamline at ANKA.
Fig. 8 a) Peak width analysis for the initial CBA film (from left to right) in hexadecanethiol covered, mixed-aminobenzenethiol covered and plasma covered Cu
3
Au
(111) surface as a function of potential. The FWHM and the estimated number of Cu
3
Au (111) monolayers are plotted versus potential. b) Schematic representation of
the inhibition effect of a mixed aminobenzene thiol self-assembled layer. c) Schematic of the polymerisation reaction of a mixed aminobenzene thiol self-assembled
monolayer adapted from Kuwabata et al.
32
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