Volume 6 Preprint 97
The Surface Oxidation Kinetics of Cu(100) and (110) Thin Films Visualised by In situ UHV-TEM
J. C. Yang and G. W. Zhou
Keywords: Oxidation, oxygen, copper, Cu2O, in situ, TEM, UHV, kinetics, surface science
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Volume 6 Paper H054
The Surface Oxidation Kinetics of Cu(100) and
(110) Thin Films Visualised by In situ UHV-TEM
J. C. Yang and G. W. Zhou
Materials Science & Eng., University of Pittsburgh, Pittsburgh, PA,
15261, USA, firstname.lastname@example.org
In situ microscopy provides dynamic information about nucleation,
growth and coalescence of islands in real time. We utilised in situ
fundamental insights into the oxidation of Cu films. A semiquantitative model, where oxygen surface diffusion is the dominant
mechanism for transport, nucleation and growth of the copper oxide,
describes both the Cu(100) and Cu(110) initial oxidation behaviour.
Oxide island formation was observed even at atmospheric pressures,
where we speculate that nucleation and coalescence of oxide islands
explains the self-limiting behaviour of Cu passivation and apply this
concept to explain the different passivation film thickness for Cu(100)
Keywords: Oxidation, oxygen, copper, Cu2O, in situ, TEM, UHV,
kinetics, surface science.
Environmental stability is one of the most important properties for
materials exposed to air.
As dimensions of materials systems
approach nanoscale, it is critical to fundamental understanding of their
interactions with oxygen at this length scale is essential for
environmental stability [#ref1] as well as for processing thin films,
where gas-surface reactions are utilised [#ref2].
temperature corrosion, electrochemistry, catalytic reactions, sensors,
gate oxides, thin film growth and fuel reactions.
The interaction of oxygen with surfaces ranges from single oxygen
atom surface adsorption to bulk oxide growth. The improved
experimental techniques in ultra-high vacuum (UHV) now make it
feasible to investigate surfaces that are atomically clean. Many elegant
experiments have been performed using UHV scanning-tunnellingmicroscopy (STM) to watch the interaction of gases, including oxygen,
on bare metal surfaces.[#ref3], [#ref4], [#ref5], [#ref6] The STM studies
can provide many elegant insights into the atomic mechanisms of
oxygen adsorption but limited to a few monolayers. Hence, these
surface techniques do not provide insights into the nucleation and
initial growth to coalescence of the metal oxides.
On the other hand, most high temperature oxidation investigations
have been concerned with thicker, on the order of a few microns or
greater, thermodynamically stable oxide films. A classic and standard
method for measuring the rate of oxide growth is thermogravimetric
analysis (TGA), which measures the weight gain due to oxygen
consumption. The rate that the weight changes during oxidation
provides information of reaction kinetics. Thermogravimetric analysis
can provide important information about the dominant migrating
defects, such as metal cations and/or oxygen anions vacancies and/or
interstitials, the transport of the cations, anions, electrons and/or
holes for charge balance, when the measurements are performed as a
function of oxygen partial pressure. However, TGA does not provide
any structural information. Hence, nearly all classical theories assume
a uniform growing film, where structural changes are not considered
because of the lack of previous experimental methods to visualise this
non-uniform growth in conditions that allowed for highly controlled
surfaces and impurities. [#ref7], [#ref8], [#ref9]
We are presently using a novel technique, in situ ultra-high vacuum
transmission electron microscopy (UHV-TEM), that combines the
correct range in spatial resolution for bridging the gap between
nucleation and initial growth of oxide as well as provides the UHV
environment necessary for controlled surface conditions. Visualising
the oxidation process at the nanometer scale with in situ experiments
under ultra high vacuum conditions provides essential insights into the
complex kinetics and energetics of nano-oxide formation. This
experimental tool provides dynamic, unique and critical data of these
gas surface reactions in a wide pressure and temperature range
mechanisms, which enables us to understand oxidation mechanisms in
order to beneficially manipulate surface reactions.
Furthermore, in the rapidly developing field of nanotechnology, the
formation of ordered structures through surface processes has been
intensively investigated because of potential applications and the
intrinsic interest in structure with reduced dimension. Controlled
nanoscale pattern formation on surfaces has become one of the most
important and challenging areas of nanoscience. Currently, various
techniques with controlled deposition, such as MBE, CVD, and PVD, are
used to grow self-assembled epitaxial nanostructures, in which
coherent island formation occurs during the growth of latticemismatched materials systems. Such self-assembled nano-structures
have been realized in semiconductor systems of Ge/Si , GeSi/Si
[#ref10], [#ref11], and InAs/GaAs.[#ref12] Nano-oxidation can be
nanostructures. The controlled formation of the oxide nanostructures
would be technologically important for their potentially novel optical,
magnetic and sensor properties. [#ref13], [#ref14]
Copper is considered to be a model system for the oxidation of metal
for the fundamental understanding of oxidation mechanisms and has
been studied extensively. [#ref15], [#ref16], [#ref17], [#ref18] Copper
forms two thermodynamically stable oxides, Cu2O and CuO. Cu2O is
simple cubic lattice (space group PN-3M) with 4Cu and 2O atoms in its
basis, and a lattice parameter of 4.22Å. The Cu atoms form a FCC
lattice and the O atoms form a BCC lattice, where each O atom is
surrounded by a tetrahedron of Cu atoms. CuO has a monoclinic
structure. In this article, we will review the results regarding the nanooxidation of Cu(100) and (110) thin films.
The microscope used in this work was a modified JEOL 200CX. [#ref19]
A leak valve attached to the column of the microscope permits the
introduction of gases directly into the microscope. In order to
minimize the contamination, a UHV chamber was attached to the
middle of the column, where the base pressure was less than 10-8 torr
without the use of a cryoshroud. Under UHV conditions, the film
surface is atomically clean at the start of the oxidation experiment,
which is extremely important for quantitative understanding of growth
kinetics. The microscope was operated at 100 KeV to minimise
irradiation effects. Single crystal 99.999% pure 700Å Cu films were
grown on single crystal (100) or (110) NaCl substrates in an UHV ebeam evaporation system, then removed from the substrate by
dissolving the NaCl in de-ionized water. The native Cu oxide was
removed inside the TEM by annealing the Cu films in methanol vapor
at a pressure of 5×10-5 torr and 350°C, which reduces the copper
oxides to copper .[#ref20]
Scientific grade oxygen gas of 99.999%
purity can be admitted into the column of the microscope through the
leak valve at a partial pressure between 5×10-5 torr and 760 torr. The
temperatures between room temperature and 1000°C. After removal
from the in situ TEM investigation, the samples can be analysed using
a Digital Instruments NanoScope IIIa scanning probe microscope.
Contacting-mode atomic force microscope (AFM) height images were
employed to assess the sample topography.
Oxygen gas molecules impinge on the metal surface and dissociate.
The chemisorbed oxygen creates a “missing-row” Cu-O surface
reconstruction on both Cu(100) and Cu(110) surfaces.[#ref21], [#ref5]
The role of surface reconstruction in the nucleation of metal oxides is
important to study since previous workers have speculated that a
surface saturated layer must form before the onset of oxidation of
Cu.[#ref22] Figure 1 is the Cu (200) dark field image after Cu film was
exposed to oxygen at 5x10-4 torr partial pressure and 4500C. It shows
pronounced strain contrast due to the anisotropic √2 x 2√2 R45
surface reconstruction. [#ref21] The oxide nucleation occurs on the
reconstructed Cu-O surface and hence the surface topology and
factors that influence the surface, should significantly alter the initial
stages of oxidation. We will speculate that the differences in surface
reconstruction explains the quantitative differences between Cu(100)
and Cu(110) oxidation behaviour.
Figure 1: Dark field image using the Cu(200) reflection, where the
vertical and horizontal strain contrast due to the anisotropic √2 x 2√2
R45 is observed.
In the following sections, the model of the oxide nucleation and
growth, based on oxygen surface diffusion, will be presented first and
then the comparison with in situ experimental data will be shown.
Next, a novel view on metal passivation, based on nucleation to
coalescence, will be described, and the implications of this model on
the passivation behaviour of Cu(100) and (110) will be discussed.
Oxygen Surface Diffusion Model
After surface reconstruction, further impinging oxygen molecules
most likely dissociate into oxygen atoms, and then diffuse across the
O-chemisorbed surface, where they may be lost to re-evaporation,
form new oxide nuclei, or be captured by an existing nuclei.
Assuming oxygen surface diffusion is the dominant transport
mechanism for the nucleation of copper oxides, then the probability of
an oxide nucleation event is proportional to the fraction of the
available surface area outside these “zones of oxygen capture” and the
oxide nucleus density can be determined to be
(1 − e )
− kL2d t
where L2d is the area of the zone of oxygen capture, 1 / L2d is the
saturation island density, Ld is much larger than the diameter of the
k is the initial nucleation rate, which depends on the
probability for Cu and O to form Cu2O, and t is the oxidation time.
Because of higher mobility of oxygen at higher temperatures, the
attachment to existing island is more probable than the nucleation of
new nuclei. Hence, it is reasonable to expect many small islands form
at low temperatures, whereas for high temperatures less island density
but larger average island size is observed. Therefore, the saturation
island density dependence on temperature could follow an Arrhenius
N s ~ e− Ea / kT
where k is Boltzmann constant, T is the oxidation temperature. This
activation energy, Ea, of the nucleation depends on the energies of
nucleation, absorption and/or desorption,[#ref23], [#ref24], [#ref25]
and not necessarily on the oxygen surface diffusion energy only. By
measuring the island density at different temperatures, then the
activation energy, Ea, for this surface-limited nucleation process can be
developed an oxidation model based on the assumption that oxygen
surface diffusion should play a major role in the initial growth of the
metal oxide. They assumed that the oxide islands grew on the metal
surface, i.e. 2-dimensional (2D) and obtained a parabolic growth rate
law if oxygen surface diffusion and impingement on the island’s
perimeter is the dominant transport mechanism.
= c s f s πr 2
where N(t) is the number of oxygen atoms in Cu2O island at time t, cS is
the sticking coefficient, fS is the diffusive flux of oxygen and r is the
radius of the circular profile of an island.
The formation of oxide is accompanied the conversion of copper
atoms from the substrate to Cu2O islands, therefore, the oxide islands
should grow 3-dimensionally (3-D) into the substrate. We have
extended Orr’s model to incorporate the 3-D growth. For 2-D lateral
growth of a disk-shaped island, with thickness a, then by solving the
differential equation, eqn. (3), the cross-sectional area increases
parabolically with respect to time.
Following a similar analysis for 3-
D growth of a spherical island, then the cross-sectional area, A, of the
oxide islands, is
A( t ) = πΩc s f s (t − t0 )
where Ω is the volume occupied by one O atom in Cu2O. The power
law dependence, t2 for 2-D and t for 3-D, is independent of the shape
of the island. For growth rates faster than linear for 3-D or t2 for 2-D
growth, then other mechanisms, such as direct impingement, could
also contribute to the growth of the oxide island.
In situ UHV-TEM of Cu Oxidation
Figure 2(a) is a dark field image taken from the Cu2O(110) reflection
after the Cu film has been cleaned. No oxide islands are visible in this
region. Figure 2(b-c) shows subsequent dark field images, taken from
the Cu2O(110) diffraction spot of the same area as shown in the
Figure2(a) of the Cu film, at successive 10 minute time increment after
O2 was leaked into the column at a partial pressure of 5x10-4 torr and
sample temperature at 3500C. After the exposure to O2, no oxide
reconstruction. The oxide islands were then seen to nucleate rapidly
followed by growth. These TEM images show the Cu2O islands that
formed on both surfaces of Cu film. The SAD (selected area diffraction)
pattern of the Cu film after oxidation can be indexed as (001) Cu2O,
where the relative orientation between the Cu2O and Cu film is (001)
Cu//(001) Cu2O and Cu//Cu2O.
Figure 2: Dark field images of Cu(100) oxidation at P(O2) ~ 5x10-4 torr
and T= 350°C, for (a) 0 min, (b) 10 min and (c) 20 min.
Figure 3 is the experimental data and theoretical fit to equation (1). A
good match is noted where the fit parameters, k = 0.17 µm-2 min-1
and Ld = 1.09 µm. Hence, the initial nucleation rate is 0.17µm-2 min-1
and the saturation island density, 1/Ld2 is 0.83µm-2 . [#ref28]
Indentations in the Cu film were noted from the thickness fringes
taken from the Cu (200) dark field images, which is sensitive to the
morphology of the Cu film. This clearly demonstrated the 3
dimensional growth, where the Cu2O islands grow into the Cu film.
The cross-sectional areas of several individual oxide islands were
measured as a function of oxidation time. Figure 4 shows the cross-
sectional area of the oxide island versus time for oxidation at 350°C at
5x10-4 torr O2. The best power law fit to our data was 1.30 ± 0.04.
This power law is slightly higher than t —the predicted power law
dependence for 3-D growth by oxygen surface diffusion. To account
for the slight deviation from a linear power law, other mechanisms
besides oxygen surface diffusion were considered.
possibility is that the oxygen that directly impinges onto the oxide
islands is incorporated into the Cu2O.
By an extension of our previous derivation, the combination of oxygen
surface diffusion and direct impingement gives the following time
− 2S ln 1 + B
πk B k B
c S fS Ω
A = t − to
c Bf BΩ
and: cS is the sticking coefficient
for oxygen surface diffusion
mechanism, fS is the surface diffusive flux of oxygen, cB is the sticking
coefficient for direct impingement, fB is the flux of oxygen for direct
impingement mechanism and Ω is the atomic volume of O in Cu2O.
We obtain an excellent fit to the combined surface diffusion and direct
impingement model, as shown in figure 3; kS is 173±12 and kB is
3.2±0.2. Hence, the growth of the Cu2O islands is initially dominated
by the surface diffusion of oxygen.[#ref29]
Figure 3: Cu2O island density as a function of oxidation time at
constant oxygen partial pressure of 5×10-4 and temperature of 350°C.
The effect of oxidation pressure and temperature were also
investigated. Figure 5 shows the Arrhenius dependency of the
saturation island density to temperature, where an activation energy of
1.1eV is noted.
Altering oxidation pressure only changed the rate of the reaction, but
dramatically different morphologies of oxide nanostructures on
Cu(100) can be achieved by modifying the oxidation temperature.
Figure 6 is a sequence of bright field images of Cu(100) oxidized at
various temperatures at the same oxidation pressure, P(O2) ~ 5x10-5
torr.[#ref30] The oxidation temperature will affect diffusion,
interfacial strain, surface and interfacial energies, and elastic
properties, which all play a significant role in the development of the
oxide morphology to pyramids, domes or terrace-layered structure.
Of particular interest was the formation of elongated islands at 600°C ,
which bear a striking resemblance to the nanorod formations of Ge on
Si. The in situ observation data on the elongation of Cu2O islands
agree with the theoretical model proposed by Tersoff and Tromp,
initially developed to explain nanorod formation of Ge on Si, thereby
demonstrating the increased universality of the elastic strain relief
Figure 4: Cu2O island cross-sectional area as a function of oxidation
time at constant oxygen partial pressure of 5×10-4 and temperature of
Figure 5: The saturation island density as function of temperature,
where an activation energy of 1.4 eV is obtained.
Figure 6: The morphology of Cu2O islands formed during in situ
oxidation of Cu(001) at an oxygen partial pressure of 5x10-4 torr and
oxidation temperatures of (a) 350°C, (b) 500°C, (c) 600°C, (d) 750°C
and (e) 1000°C.
Figure 7 (a) is a dark field TEM image after the copper film has been
cleaned with methanol. No oxide islands are visible in this region.
Figure 7 (b-c) show the corresponding dark field images at the same
area as shown in Figure 7(a) of the copper film at successive 10 minute
time increments after oxygen was leaked into the column of the
microscope. The partial pressure of oxygen was 5×10-4 torr and the
temperature of the copper film was held at 350°C. After the
introduction of oxygen gas, the nuclei appear after an incubation
period of several minutes. After the oxidation of about 22 minute, no
new islands formed. The selected area electron diffraction pattern of
the Cu2O island and underlying Cu(110) substrate revealed that the
oxide island is
110 Cu// 110 Cu2O
epitaxial with the underlying Cu film, i.e.
and (001)Cu//(001)Cu2O. A similar epitaxial
relationship was noted for Cu(100), where (001)Cu//(001)Cu2O and
Figure 8 shows the experimental data and theoretical fit to equation
(1), where the ranges of the error bars are based on the measured
oxide island density obtained from several experimental runs and
different regions on the Cu surface. A good match is noted where the
fit parameters are: the initial nucleation rate, k, is 1.7432 µm-2 min-1,
and the saturation island density, 1 / L2d , is 9.01159 µm-2.[#ref32]
Figure 7: In situ dark field TEM images taken as a function of oxidation
time, (a) 0 min, (b) 10 min, (c) 20 min, (d) 30 min at constant oxygen
partial pressure of 5×10-4 and temperature of 350°C.
We measured the saturation density of the nuclei as a function of
oxidation temperature, from 300°C to 450°C, at constant oxygen
pressure of 5×10-4 torr. Figure 9 shows the saturation density of nuclei
versus inverse oxidation temperature, where the activation energy, Ea,
which is equal to the slope, was determined to be 1.1±0.2eV.
comparison, the activation energy, Ea, for Cu(100) was measured to be
Figure 8: Cu2O island density as a function of oxidation time at
constant oxygen partial pressure of 5×10-4 and temperature of 350°C.
Figure 9: Cu2O saturation island density versus inverse temperature.
The absolute value of the slope is the Ea for the surface-limited
The evolution of cross section area of the islands is recorded in situ,
where a sequence of images focusing on the growth of individual
islands was obtained, where the Cu(110) film was oxidised at 5×10-4
torr and 450°C. About a couple of minutes after the introduction of
oxygen gas, Cu2O islands were observed to nucleate rapidly followed
by growth of these islands. After the initial nucleation of the oxide
islands, ~5 minutes, the saturation density of the island nuclei was
reached and no new nucleation event was observed, which is much
faster than the oxidation at 350°C, where the saturation is reached
after 22 min oxidation.
Figure 10 is the comparison of the experimental data of the crosssectional area of the oxide islands to this surface diffusion model by
using 3-D growth of oxide island, Eqn. (4). The kinetic data on the
evolution of cross section area of the islands agree well with the model
of surface diffusion of oxygen. Therefore, the excellent agreement of
evolution of cross section area of the islands with the kinetic model
validates that the growth of the 3-dimensional Cu2O islands occurs by
oxygen surface diffusion and impingement on the oxide perimeter.
Since the initial oxidation stages are surface processes, it is reasonable
to expect that the crystallographic orientation of the underlying metal
will have a major effect on the nucleation behaviour, growth rate and
the orientation of the oxide film, and this should significantly impact
bulk oxidation behavior, such as passivation. Table 1 summarises the
values of the fit parameters for Cu(100) and Cu(110) oxidation.
Differences in the rate, oxide island shapes, and fit parameters to the
surface models were noted between Cu(100) and Cu(110). Specifically,
The initial nucleation rate on Cu(110) is much faster than Cu(100). The
saturation density on Cu(110) surface is 11 times larger than that on
Cu(100), although the oxide islands reached their saturation density
after similar oxidation time for the two orientations. Hence, the active
zone of oxygen capture around each island on Cu(110) is much
smaller than that on Cu(100). The activation energy, Ea, for the
nucleation process on Cu(110), was measured to be 1.1±0.2eV, which
is also smaller than that on Cu(100) surface, 1.4±0.2eV.
Figure 10: Comparison of the experimental data of the time evolution
of the cross-sectional areas of six individual oxide islands and the
theoretical function for the surface diffusion for the 3-D growth of
Table 1: Comparison of fit parameters of Cu (110) and (100) oxidation
The observed different rates of oxidation between Cu(100) and
should effect the passivation behaviour of these two
orientations. Previous research by K.R. Lawless and A.T. Gwathmey
 in the oxidation of a spherical single crystal of Cu with its
multitude of surface orientation demonstrated the anisotropy of the
oxidation rate on different faces. Figure 11is a reproduction of their
experimental data on Cu(100) and Cu(110) in which Cu(100) forms a
much thicker passive oxide than Cu(110).[#ref18]
Figure 11: Oxidation kinetics of copper single crystals at 523K and 1
atm of oxygen from Gwathmey et al. [#ref18]
We first describe the new model of passivation based on nucleation to
coalescence and then discuss the implications of this model to the
effect of crystallographic orientation effect on the passivation
behaviour of Cu.
The Self-Limiting Oxidation Behaviour of Metal Passivation
Nearly all metals form a passivation film due to oxidation in air at
ambient temperature. This film acts as a diffusion barrier to protect
the materials from further corrosion. The classical theory of CabreraMott [#ref7] describes the passivation film formation on metals, where
they predict that this film grows as a uniform layer due to a field -
enhanced ionic transport mechanism. Cabrera-Mott model predicted
an inverse logarithmic growth rate for metal oxides, which formed due
to outward cation diffusion (eqn 6).
= A B ln t
where, x is the thickness of the oxide film, t is time and A and B are
the fit parameters. Figure 11 is a reproduction of Young et al.'s
experimental data for Cu combined with a fit to the above equation.
[#ref18] We have shown that nucleation and growth of oxide islands
occurs even at atmospheric pressure .[#ref33] Thus here the Cabrera-
Mott model, which assumes a uniformly growing film, can not be valid.
We have proposed that the self-limiting oxidation of Cu is due to the
coalescence of islands, which 'switches-off' the surface diffusion route
and needs much slower bulk diffusion for further oxidation. Figure 12
(a-c) is a sequence of dark field images for which the Cu(100) film was
oxidized at 0.1 torr and 3500C. Here the oxide island formation
followed by coalescence is clearly evident at atmospheric pressure.
One standard theory to describe coalescence of thin films is the
Johnson-Mehl-Avrami-Kolmogorav (JMAK) theory, which assumes that
the nucleation and coalescence of thin films is due to surface
processes.[#ref34] It presumes that the coverage will follow an
exponential dependency on time (eq.7):
X(t) = 1-exp(-ktn)
where X is the coverage, t is time and k and n are fit parameters that
depend on the surface mechanisms of transport, nucleation and
growth. The parameter, n, is usually an integer or half-integer. For
three-dimensional growth, the fit parameters are k= π/3 and n=4 ,
assuming constant nucleation and uniform radial growth rate. This
describes the coverage with respect to time. We have observed that the
coverage does follow this exponential law where the fit parameters
are: k =1.9x10-4 and n = 2 ± 0.005, and the goodness of fit is R =
0.99.[#ref35] Experimentally, we found the copper system to be more
complex than the simplistic view proposed by JMAK yielding the above
values for n and k. Cu(100) has been shown to have a radial growth
rate of r ∝t
that is slower than the linear growth rate. It was
observed that nucleation of copper oxide islands occurs almost
instantly and not linearly with time. Both of these factors will lead to a
slower nucleation to coalescence as compared to the JMAK calculation
for 3-D growth. Another possible reason for such a small k is that the
growth rate is very slow for copper oxidation. It was noted that the
sticking coefficient is extremely small indicating that only a few
oxygen adatoms are contributing to the growth of the copper oxides.
Figure 12 : Dark field images from the Cu2O reflection showing Cu2O
island nucleation (a), growth (b) and then coalescence (c), when
Cu(001) was oxidized at 0.1 torr at 350°C.
Figure 13 is the plot of the fractional coverage of the copper oxide
with respect of time. The circles are the experimental data. The solid
line is the best fit to the JMAK formula (see eqn.7 above).
The comparison of the nucleation and growth behaviour of Cu(110)
with that of Cu(100) at the same oxidation temperature (350°C) and
oxygen pressure (5×10-4 torr) revealed that the same surface diffusion
models, originally developed to describe Cu(100) oxidation, explains
the experimental data on Cu(110) quite well and demonstrates a
greater generality of this oxygen surface diffusion model.
The differences in the fit parameters (see table1) obtained from the
oxygen surface diffusion models should be due to the crystallography
effects. To explain these differences, we consider the effects of
kinetics and energetics on both surfaces. The nucleation and growth of
oxide islands are nonequilibrium processes depending on both the
energetics, such as surface energies of the system, and the kinetics, in
particular, on the diffusion process.
Figure 13: Experimental data of the Cu2O fractional area coverage, X(t),
versus time, when Cu(001) was oxidized at 0.1 torr at 350°C, and the
comparison to the JMAK formula, X(t) = 1-exp(-ktn) (dashed line),
where k =1.98x10-4 sec-2 , n = 2.
It should be noted that the diffusion of oxygen is on the
reconstructured Cu surface.
Previous investigators have elegantly
demonstrated that Cu(100) and (110) surfaces are unreconstructed,
and then transform into “missing-row” or “adding-row” reconstruction
when exposed to oxygen.[#ref36], [#ref37], [#ref38], [#ref39] Oxygen
chemisorption on Cu surface is the first step for the oxidation. We
speculate that the impinging oxygen molecules dissociate into oxygen
atoms, and the dissociated oxygen is adsorbed at the Cu surface to
reconstruction, the oxygen atoms diffuse on this reconstructed
surface, and nucleation occurs on the reconstructed Cu-O surface.
Arriving oxygen adatoms will diffuse on the surface and can either
nucleate new oxide islands by reacting with copper atoms or attach to
an existing island, causing growth. Qualitatively a larger diffusion
coefficient for oxygen should yield a lower number density of stable
islands. Since the path length of oxygen surface diffusion depends on
the atomic structure of the substrate plane, different nucleation
behaviour of Cu2O islands is therefore expected for different
orientations of the Cu. The reconstructed
2x2 2 R45º O-Cu(100)
surface has a more compact surface structure than (2×1)O-Cu(110)
surface which has a corrugated structure (Figure 14). Therefore, it is
reasonable to expect the activation barrier of surface diffusion of the
dissociated oxygen will be higher on the Cu(110) surface, and thus
have a shorter path length. The shorter diffusion path length will give
rise to a smaller capture zone of oxygen and create a higher number
density of oxide nuclei. This is confirmed by our results where the
active zone of oxygen capture around each island on Cu(110) is
0.3331µm, which is significantly smaller than that on Cu(100),
Figure 14: Schematic diagram of the reconstructed
( 2x2 2 )R45º O-
Cu(100) surface (a), and (2×1)O-Cu(110) surface (b) due to oxygen
chemisorption. Filled circles: O atoms; open circles: top layer Cu
atoms; shaded circles: second layer Cu atoms.
The surface energy of Cu(100) is 1280mJ/m2, which is lower than the
If all other interface and surface
energies were equal, then, it would be expected that Cu(110) surface
will be less stable under the oxidizing atmosphere and the nucleation
of oxide islands will be facilitated, leading a higher nuclei density of
oxide and smaller overall activation energy for the nucleation of the
oxide islands as determined by our measurements. As yet, the surface
energies of Cu2O and the interfacial energies between Cu2O/Cu(100)
and Cu2O/Cu(110) are unreported. The interfacial strain energy could
depend on orientation since it is a function of the Possion ratio (ν) and
shear modulus (µ) of the substrate, as well as oxide island bulk stress
(σb), and is also not known in Cu2O/Cu system. Hence, determination
of the surface energies and metal-oxide interfacial energies, such as
by theoretical modelling, are critical to fundamentally quantitative
understanding of oxidation.
The oxygen surface model assumes homogeneous nucleation, not
heterogeneous nucleation. One particularly interesting question is the
role of defects, such as dislocation and steps in the initial oxide
nucleation. In the oxidation of Cu(100), no preferential nucleation sites
at dislocations or surface steps were observed. [#ref28], [#ref41]
Similar as in the oxidation of Cu(110), repeated oxidation, reduction,
followed by oxidation experiments were performed, but no nuclei
appeared at the same positions. Furthermore, the oxide island density
was observed to decrease with increasing temperature, following as
Arrhenius dependence, indicates that nucleation was homogeneous. If
the nucleation mechanism was heterogeneous, then it would be
reasonable to expect similar island density at different temperatures,
which was not observed.
Since we are interested in the initial oxidation mechanisms, then it is
essential to isolate the effect of the electron beam on the oxidation
kinetics. In order to minimise the possible radiation damage, the TEM
was operated only at 100KeV. Careful oxidation experiments with and
without the electron beam irradiation were conducted. Qualitatively,
the structural changes appeared similar as with the electron beam on.
However, the effect of the electron beam was to reduce the reaction
rate slightly, but the saturation density of the oxide islands was
observed in both cases, indicative of a surface diffusion limited
nucleation mechanism. This is again similar to the Cu(100) oxidation.
As discussed previously, Cu(110) shows a faster nucleation rate of
oxide than Cu(100) in the initial stage of oxidation, and the oxidation
on Cu(110) surface results in a higher number density of oxide nuclei
under the same oxidation conditions. Therefore, there is a faster
coalescence of oxide islands on Cu(110) than that on Cu(100). In the
oxidation of Cu(100), the oxidation time is longer for the coalescence
of oxide islands due to smaller number density of oxide nuclei. As a
result, Cu(100) develops a thicker passive oxide than Cu(110) at the
later stage of oxidation as confirmed by Gwathmey’s work.[#ref18]
Surprisingly, orientations that form a higher density oxide nuclei and
have a faster initial oxidation rate, may have a slower long-term
growth rate due to the rapid coalescence of the oxide which switches
the oxide growth mechanism from surface diffusion to the slower
diffusion through an oxide scale. It is interesting to speculate on the
poor oxidation characteristics of Cu and other metals which form
crystalline epitaxial oxide passivation films.
These films, when
coalesced, may have a tendency to contain microcracks and pinholes,
for reasons of mismatch strain relief, which provides improper
passivation. In contrast, good passivation layers, such as those which
form on Al and Si, are amorphous.
Surface diffusion of oxygen is the dominant transport mechanism for
the oxide nucleation and growth on Cu(100) and Cu(110). Cu(110)
surface shows a highly enhanced oxidation rate than Cu(100) in the
initial stage of oxidation. We speculate that the faster oxidation rate is
explained by the O-chemisorbed corrugated roughness of Cu(110) and
larger surface energy of Cu(110). Detailed experimental data along
with a sophisticated theoretical treatment of the initial stages of a
chemical treatment on a metal surface is essential to a fundamental
understanding of oxidation.
The dynamic observation obtained from these in situ experiments
enable us to gain basic insight into the kinetics/energetics in the initial
stages of oxidation.
These observations provide fundamental
understanding of the formation, distribution, and morphological
evolution of oxide islands and demonstrate the possibility of
controlled oxidation as a processing tool for creating specific
Furthermore, nucleation and growth of islands occur even at
atmospheric pressure conditions.
Thus models of the self-limiting
(“passivation”) behaviour of oxidation for Cu, such as that of Cabrera
and Mott, which assume a uniform growing film, cannot be valid. We
suggest that a simple phenomenological explanation for the selflimiting oxidation is due to the coalescence of islands, which
“switches-off” the surface diffusion route and requires much slower
bulk diffusion for further oxidation. The microstructure of the scale at
the onset of oxidation plays an influential role in the oxidation rate at
later stages, including passivation. In light of this phenomenological
model, the thinner passive oxide of Cu(110) is due to the increased
nucleation rate and density of oxide islands on Cu(110) as compared
to Cu(100). A complete understanding of oxidation properties must
include the investigation of the initial stages of oxidation, i.e., the
nucleation, growth to coalescence.
Since oxidation is a surface process, then understanding and control
of surface structure, conditions and reactions are essential in tailoring
oxidation behaviour. Our results are based on in situ transmission
electron microscopy of the oxidation of Cu thin films, but we suspect
that the results are more general.
This research project is funded by the National Science Foundation (#
9902863), Department of Energy−Basic Energy Science and a National
experiments were performed at the Materials Research Laboratory,
University of Illinois at Urbana-Champaign, which is supported by the
U.S. Department of Energy (#DEFG02-96-ER45439).
kindly thank I. Petrov, R. Twesten, M. Marshall, K. Colravy, and N.
Finnegan for their help.
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