###### Surf. Sci. Letts., subm. 9/00

**STM-induced Void Formation at the Al~2~O~3~/Ni~3~Al(111) Interface**

# 

# N. P. Magtoto, C. Niu, M. Anzaldua and J. A. Kelber\*

## Department of Chemistry, University of North Texas, Denton, TX-76203

and

**D. R. Jennison**

### Sandia National Laboratories, Albuquerque, NM 87185-1415

#### 

#### ABSTRACT

Under UHV conditions at 300 K, the applied electric field and/or
resulting current from an STM tip creates nanoscale voids at the
interface between an epitaxial, 7.0 Å thick Al~2~O~3~ film and a
Ni~3~Al(111) substrate. This phenomenon is independent of tip polarity.
Constant current (1 nA) images obtained at +0.1 V bias and +2.0 bias
voltage (sample positive) reveal that voids are within the metal at the
interface and, when small, are capped by the oxide film. Void size
increases with time of exposure. The rate of void growth increases with
applied bias/field and tunneling current, and increases significantly
for field strengths \>5 MV/cm, well below the dielectric breakdown
threshold of 12  1 MV/cm. Slower rates of void growth are, however,
observed at lower applied field strengths. Continued growth of voids, to
\~ 30 Å deep and \~ 500 Å wide, leads to the eventual failure of the
oxide overlayer. Density Functional Theory calculations suggest a
reduction-oxidation (REDOX) mechanism: interfacial metal atoms are
oxidized via transport into the oxide, while oxide surface Al cations
are reduced to admetal species which rapidly diffuse away. This is found
to be exothermic in model calculations, regardless of the details of the
oxide film structure; thus, the barriers to void formation are kinetic
rather than thermodynamic. We discuss our results in terms of mechanisms
for the localized pitting corrosion of aluminum, as our results suggest
nanovoid formation requires just electric field and current, which are
ubiquitous in environmental conditions.

\*Corresponding author, Ph: (940)-565-3265, FAX: (940)-369-8295, email:
<kelber@bob.unt.edu>

Keywords: Scanning tunneling microscopy, density functional
calculations, oxidation, corrosion, field effect, aluminum oxide,
aluminum, nickel oxides, ceramic thin films, metal insulator interfaces.

\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\-\--

In this Letter, we report nanovoid formation at the interface between a
7.0 Å thick Al~2~O~3~ film and a Ni~3~Al(111) substrate, caused by an
STM tip under constant current feedback. First principles density
functional theory (DFT) calculations suggest a facile local
reduction-oxidation (REDOX) reaction, in which substrate metal is
incorporated into the film and surface cations are reduced and diffuse
away at room temperature.

The application of several volts across ultrathin (1-2 nm) films of
dielectric materials can lead to extremely high electric fields and
resulting tunneling currents. The behavior of ultrathin oxides under
such conditions is of significant interest in microelectronics, because
dielectric breakdown of the oxide has a direct impact on the stability
of ultra-large scale integrated circuit devices \[1\] and the
development of tunneling-based magnetoresistance devices \[2, 3\]. The
effect of high electric fields on the stability of thin oxide films is
also of critical importance in understanding possible connections
between dielectric breakdown and the onset of localized corrosion
(pitting) \[4, 5\]. The ability of STM to provide high spatial
resolution and generate very high electric fields in a controlled
ultrahigh vacuum (UHV) environment allows for the systematic assessment
of the effects of electric fields on morphology, composition, and
localized defect density (breakdown and pre-breakdown) at surfaces and
interfaces.

The morphology and location of the STM-observed voids strongly resemble
voids recently studied at the alumina/aluminum interface using positron
annihilation spectroscopy \[6\], and also by atomic force microscopy
during the dissolution of aluminum in aqueous NaOH \[7\]. In all cases,
the nanovoids are in the metal itself and, when small, are capped by the
oxide film overlayer. In the present work, however, no chemically
corrosive agent is involved in the formation of interfacial voids.

The conditions which cause rapid nanovoid growth (fields \~ 5 MV/cm) are
comparable to those generated at magnetoresistance junctions \[8, 9\]
and also made by a halide ion adsorbed on the thin passivating oxide
layer on aluminum metal in an aqueous environment \[5\]. Preliminary
results at lower field strengths, however, indicate that void initiation
and growth still occurs, but at rates difficult to observe.

Experiments were carried out in a UHV surface analysis system \[10\]
(base pressure, 3 x 10^-11^ Torr) equipped with LEED, AES and STM.
Al~2~O~3~ films were grown by directly oxidizing a clean Ni~3~Al(111)
substrate at 300 K. The 7.0 ± 1 Å thick ordered Al~2~O~3~ was obtained
by annealing at 1100 K, as described previously \[10\]. Recent studies
\[11, 12\] have demonstrated that such oxide films are pure Al~2~O~3~,
and that a (111)-oriented monolayer of metallic Al atoms exists at the
Al~2~O~3~/Ni~3~Al(111) interface \[10\], just beneath a layer of
chemisorbed oxygen -- thus the actual Al~2~O~3~ material does not start
until the second O-layer. The ordered Al~2~O~3~ film was stressed either
by applying voltage pulses to the tip or holding the voltage constant at
a given location on the surface for a specified time. (Application of a
positive gap voltage to the sample denotes the tunneling of electrons
from the occupied states of the tip to the unoccupied states of the
sample.) The constant current feedback loop was active during high field
exposure in order to prevent tip-sample physical interaction.

Voids were created as follows. After imaging a large area of the sample
at +0.1 V gap voltage and 1.0 nA feedback current, the tip was
positioned at a specific location above the surface. With the feedback
loop set at 1 nA, voltage pulses were applied. In a single pulse, the
gap voltage was raised step-wise to a specified value (pulse height),
then instantaneously decreased back to the normal tunneling voltage of
0.1 V. Each incremental step required about 200 μs. Since the pulse
height was always divided into 200 steps, the whole procedure was
completed in about 40 ms. The magnitude of the electric field (F) can be
estimated by dividing the applied voltage (V) by the distance (d)
between the Al~2~O~3~/Ni~3~Al interface and the tungsten tip (F = V/d),
where d is the sum of the tip-oxide surface distance and the oxide
thickness (7 Å). The values of the tip-oxide surface distance (\~40 --
43 Å) were estimated from the plots between the gap voltage and
tip-sample separation (not shown). Electric field strengths under such
conditions are calculated to be \~ 5 -- 7 MV/cm, well below the
STM-induced dielectric breakdown threshold of 12  MV/cm recently
reported for such films under UHV conditions \[13\]. After each high
field exposure, the surface was imaged under normal tunneling conditions
(1 nA constant feedback current, +0.1 V bias voltage).

Void formation was marked by the appearance of a depression in the STM
image. Fig. 1a displays voids created by voltage pulsing to 3.5 V at
various locations at the surface. The voids at positions 1 and 2 were
produced by 2 and 8 pulses, respectively (Fig. 1a). It is evident from
the line profiles in Fig. 1b that the void in position 2 is \~2 Å deeper
and \~350 Å wider across the rim than the void in position 1; thus, the
size of the void increases with the number of voltage pulses. A similar
increase in void size was observed upon additional pulse stressing of an
existing void.

To determine whether voids were located at the Al~2~O~3~/Ni~3~Al(111)
interface or in the oxide, the surface was imaged at constant current
(1nA) at low (0.1 V) and high (2.0 V) gap voltages after stressing
\[14-18\]. Imaging at low gap voltages reveals the topography of the
oxide/metal interface. Under this tunneling condition, electrons tunnel
directly into the electronic states of the metal, with the oxide acting
as a tunneling barrier. At high gap voltages, however, electrons can
tunnel directly between the tip and electronic states in the oxide. STM
images obtained under this tunneling condition reflect accurately the
topography of the oxide overlayer. This effect has been previously
demonstrated for ultrathin films of Al~2~O~3~ on NiAl(110) \[15\] and
SiO~2~ on Si(111) \[18\]. Typical results are shown in Fig. 2. These
constant current (1 nA) images were obtained after exposing (30 pulses)
a region of the surface to a +3.5 V bias. Subsequent constant current
imaging at 0.1 V bias (Fig. 2a) clearly revealed the formation of a pit
\~ 0.8 nm deep (see line profile shown in Fig. 2b). The image of the
same region obtained at 2.0 V bias (Fig 2c), however, did not show the
presence of a pit. Since STM is more sensitive to the topographies of
the oxide/metal interface and oxide surface at low and high tunnel
voltages, respectively, these results demonstrate that the voids are
located at the Al~2~O~3~/Ni~3~Al(111) interface and within the metal
itself.

Voids were also created by holding the voltage constant for a specified
time at a given location on the oxide surface. The dependence of void
growth rate on the electric field was characterized by measuring the
cross sectional area of the void after exposing the surface to a
specified bias voltage at 300 seconds. Results are shown in Fig. 3a for
void cross sectional area vs. electric field. Under these experimental
conditions, the electric field varies approximately linearly with the
applied gap voltage. The data in Fig. 3a indicate that the rate of void
growth increases rapidly above \~5 MV/cm. The threshold for void growth
suggested by the data in Fig. 3a may be more apparent than real, since
voids can be created even at field strengths below 4 MV/cm provided the
surface is exposed to a field for a very long time (≥ 2700 sec).
Experiments at constant voltage and varying feedback limit indicate that
void growth rate also increases with tunneling current at a given
applied voltage (Fig 3b), although the growth rate apparently approaches
an asymptotic value at higher currents.

Upon reaching a certain size, the void may induce a failure of the oxide
overlayer. This is shown in Fig. 4. The constant current image of the
metal substrate (Fig. 4a) shows the presence of a void \~ 30 Å in depth
and 500 Å wide. Subsequent imaging at 1.0 V gap voltage (Fig. 4b) also
shows a gap in the oxide overlayer. This result indicates that continued
exposure to high bias voltage causes the void to grow wider and deeper
into the metal, eventually causing the (assumed) collapse of the oxide
overlayer. This behavior was also observed in an aggressive aqueous
environment where the oxide film suffered a local collapse once the void
grew to a critical size \[19\], larger than here because of a thicker
oxide layer.

The apparent lack of a discernable voltage threshold suggests the
process is not stimulated by a specific electronic excitation. In
addition, bias polarity-independence and current-dependence argue
against field-assisted diffusion across the interface. Field induced
vacancy diffusion from within the metal can be ruled out because the
applied field does not extend far into the conducting substrate. The
migration of vacancies from within the oxide is similarly improbable
because the oxide film is of high quality and the data clearly
demonstrate void growth into the metal.

A possible explanation is localized heating due to inelastic
electron-phonon scattering, enabling the system to overcome a kinetic
barrier. This mechanism would show a current and field dependence of
void growth rate, consistent with the data in Fig. 3. A similar
mechanism was proposed for Si-H bond-breaking at Si surfaces \[20\],
while electronically stimulated processes within SiO~2~ films have a
discernable threshold \[21\].

To help define a mechanism, we turn to *ab initio* theory. This work
used DFT \[22\] and slab calculations. Because energetics are compared,
we used the generalized gradient approximation (GGA) known as "PW91"
\[23\], as implemented in the Vienna *Ab-Initio* Simulations Package
(VASP) \[24\]. Ultrasoft Vanderbilt pseudopotentials \[25\] accurately
removed the core electrons with a plane wave cutoff of only 270 eV.
Geometric relaxation used a damped molecular dynamics algorithm until
all forces were less than 0.05 eV/Å. The slab had five layers of
aluminum metal, with the bottom two frozen at the bulk GGA lattice
constant of 4.035 Å. Because of long-range electrostatic interactions,
the periodic vacuum gap due to the plane-wave basis always exceeded 18
Å.

As a model for the real film on Ni~3~Al(111) (which cannot be directly
studied because of lattice-mismatch relief by domain rotation \[11\]),
computations were performed with three and four O-layer commensurate
alumina systems on Al(111), all having chemisorbed oxygen at the
interface. This interface, first proposed for ultrathin alumina films on
metals made by high temperature annealing in oxygen rich conditions
\[12\], receives further support from our angle resolved XPS results
together with those of others \[26\]: We find two types of oxygen
(chemisorbed and oxidic) to be present.

Above the layer of chemisorbed oxygen, two different phases of alumina
were explored. The first is motivated by the recent observation of
θ-alumina on NiAl(100) by A. Stierle, et al. \[27\]. This structure,
found using X-ray scattering, supported a computational result \[12\]
predicting that the normal preference for octahedral site Al-ions is
reversed at the interface; thus in this extreme model, all Al ions
occupy tetrahedral sites. In the second model, we used the recently
determined structure for the κ-phase \[28\], which has ¼ tetrahedral and
¾ octahedral site Al-ions. This structure is similar to a recent DFT
structure for the second O-layer in \~ 5 Å films on close-packed
surfaces \[29\], and is a more realistic model for the so-called \[11\]
γ'-films. We find that the qualitative results are *independent* of the
details of oxide film structure.

In all cases it is found that the REDOX reaction (Fig. 5) is preferred
energetically. In fact, for Al(111) with this interface, *the entire
first layer of Al* prefers to be taken up into the oxide, even at the
expense of becoming non-stoichiometric, with extra Al atoms reduced to
adsorbed Al (Fig. 5). This movement is preferred by 0.15 (0.21) eV per
Al atom in the tetrahedral (κ-phase) film. While these results agree
with independent experimental observations of the alumina/aluminum
system \[26\], that show a preference for the incorporation of
chemisorbed oxygen into aluminum oxide islands, the results cannot be
directly applied to the present case: starting with a perfect interface,
the first REDOX reaction breaks 6 Al-Al and 3 Al-Ni bonds, vs. 9 Al-Al
in the model systems. Given the melting temperatures of Al metal (660 C)
and Ni~3~Al (1390 C) vs. pure Ni (1455 C), Ni-Al bonds are much stronger
than Al-Al, thus reducing the exothermicity. However, these results are
consistent with a kinetically limited REDOX mechanism, and one reason
for the above energetics is the strong binding of Al adatoms to alumina
\[30\].

Overall, the above results suggest nanovoid formation in the presence of
electric field and induced current is a likely critical step in the
corrosive pitting of aluminum (and possibly of other metals) and might
significantly affect the durability of alumina-based tunneling
junctions. Remaining to be addressed are the fundamental mechanisms of
total mass transport (both interfacially and *within* the void) and
whether only Al or both Al and Ni move across the interface. In
addition, it is not known how formation and growth are related to the
cohesive energy of the substrate material. However, with respect to the
existing experimental and theoretical results, we make the following
conjectures:

\(1\) *The presence of nanovoids at the alumina/aluminum (or aluminum
alloy) interface is ubiquitous*. Voids are produced by non-uniform
electric fields and currents in the passivating oxide layer; such occur
in an electrochemical environment and are supported by the resulting
oxide point defects \[4,5\]. The induced fields are similar in magnitude
to those in the present study \[5\], while the metal cohesive energy is
lower. This conjecture is also consistent with recent positron studies
of alumina/aluminum interfaces \[6\] and with new experimental results
indicating that Cl^-^ anions do not penetrate existing oxide films under
open circuit conditions, even though pitting can occur under such
conditions \[31-33\].

\(2\) *The transition from nanovoids to microscopic corrosion pits is
induced by the collapse of, and/or the presence of microcracks in, the
oxide,* when void growth causes local mechanical failure from factors
such as strain. Cracks allow the transport of anionic species into the
void, that grow into pits because metal is etched into soluble
compounds. Local fluid flow conditions, vs. the formation of insoluble
scales and/or new oxide films at the pit surface, determine the balance
between growth and passivation.

**Acknowledgements**

The authors gratefully acknowledge informative discussions with J.
Sullivan, N. Missert, J. C. Barbour and K. Hebert. Work at UNT was
supported in part by the U.S. DOE, Office of Basic Energy Sciences,
under Grant No. DE-FG03-93ER45497, in part by the Robert Welch
Foundation, under Grant No. B-1356, and in part by NSF, under Grant No.
CHE-9714580, which are gratefully acknowledged. Sandia is a
multi-program laboratory operated by Sandia Corporation, a
Lockheed-Martin Company, for the U. S. DOE under Contract
DE-AC04-94AL85000. This work was partially supported by a Laboratory
Directed Research and Development Project and by the U. S. DOE Office of
Basic Energy Sciences, Metals and Ceramics Division.

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Figure Captions:

Fig. 1. (a) STM constant current (0.1 nA, 0.1 V bias) images of pits
formed into a "U" with varied numbers of pulses to (see text) +3.5 V
(sample positive) bias. The substrate is a 7.0 Å thick Al~2~O~3~ film on
a Ni~3~Al(111) substrate. (b) Cross sectional line scans of different
regions of the "U" after application of 2 and 8 pulses to 3.5 V,
respectively.

Fig. 2. STM constant current (1 nA) images of a region of the
Al~2~O~3~/Ni~3~Al(111) film after 30 pulses to 3.5 volts. (a) Image
acquired at 0.1 V bias, with tunneling between tip and metal states; (b)
Image of the same region acquired at 2.0 V bias, with tunneling between
tip and oxide states.

Fig. 3. (Top) Void cross sectional area after 300 sec exposure vs. the
electric field strength . (Bottom)Void cross sectional area , after 300
sec exposure, vs tunneling current.

Fig. 4. STM constant current images showing a large void and collapse of
the oxide overlayer. (a) Constant current image (1nA, 0.1V bias) showing
the void present at the oxide/metal interface; (b) Constant current
image (1nA, 1.0 V bias) of the same region showing a gap (presumed
collapse) in the oxide overlayer.

Fig. 5. Schematic diagram indicating the proposed REDOX mechanism. Atoms
are oxygen in white, Al metal in gray, Al ions in black. After the first
atom goes, it is easier for the next because of reduced coordination.
The reduced Al adatom height is shown.

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