The Localised Corrosion Associated with Individual Second Phase Particles in AA7075-
T6: A Study by SEM, EDX, AES, SKPFM and FIB-SEM
Christopher F. Mallinson*a, Paul M. Yatesa, Mark A. Bakera, James E. Castlea, Ann Harveyb,
John F. Wattsa
aThe Surface Analysis Laboratory, Department of Mechanical Engineering Sciences,
University of Surrey, Guildford, Surrey, GU2 7XH, UK,
bLife Prediction Team, Metallurgy Group, Science Function, AWE, Aldermaston, Reading,
To investigate the role of intermetallic particles in the localised corrosion of AA7075-T6,
three particles were monitored over 16 hours immersion in 3.5 wt.% KCl solution. These
were examined using Auger electron spectroscopy, energy dispersive x-ray spectroscopy,
scanning Kelvin probe force microscopy and focused ion beam-scanning electron
microscopy. Despite similar Volta potential measurements, the corrosion microchemistry
varied significantly with composition. A Al7Cu2Fe intermetallic resulted in trenching while a
(Al,Cu)6(Fe,Cu) intermetallic showed crevice corrosion and sub-surface intergranular
corrosion and a Al12Fe3Si intermetallic appeared to be galvanically inactive but showed
crevice formation at the matrix interface and sub-surface intergranular corrosion.
Keywords: Aluminium alloy, Intermetallics, AES, SEM, Pitting corrosion
Aluminium alloys are the second most widely used engineering metallic alloys after steels
. The alloy 7075-T6, is much used within the aerospace industry because of its particularly
high strength to weight ratio . The 7xxx series (Al-Zn-Mg-Cu alloys) provide the potential
for precipitation hardening throughout a range of compositions . The alloying elements
present in aluminium are either in solid solution or segregated as secondary phase micro-
constituent particles . These particles fall into one of three categories: hardening
precipitates, dispersoids and constituent particles . Hardening precipitates range in size up
to ten nanometres and MgZn2 is a common example . Dispersoids control grain size, grow
to hundreds of nanometres, and include Al3Ti and Al3Zr . Constituent particles are the
largest growing to tens of micrometres and more than a dozen types are known to occur [1,4].
Typical compositions include Mg2Si, MgZn2 (η phase), Al2CuMg (S phase), (Al,Cu)6(Fe,Cu),
Al7Cu2Fe and Al3Fe. The most numerous of these have been identified as (Al,Cu)6(Fe,Cu)
and Al7Cu2Fe [1,4,5].
The low conductivity of the aluminium oxide film results in a susceptibility to localised
corrosion in the form of pitting when the metal is exposed to an aggressive environment .
The presence of “weak spots” which interfere with the homogeneity of the oxide film are
thought to be the cause of this susceptibility . These weak spots can consist of: grain
boundaries, dislocations, mechanical damage, boundary regions between a metal matrix and a
second phase particle or the galvanic activity of second phase particles .
The potential of these second phase particles is primarily a result of their composition. With
more noble elements such as copper and iron usually resulting in cathodic particles and more
active elements such as magnesium and zinc resulting in anodic particles with respect to the
alloy matrix. While the presence of these elements can help to predict galvanic cell formation
between particles and the matrix it is not a guaranteed method for determining the behaviour
of an individual particle.
It has been observed that the magnesium rich particles, such as Mg2Si, can be anodic or
cathodic to the matrix depending upon the type of heat treatment performed on the alloy .
This occurs as a result of the dissolution or formation of MgZn2 particles in the alloy. Their
dissolution increases the magnesium and zinc concentration within the bulk alloy making it
more anodic which can drive the Mg2Si particles to be cathodic to the matrix . The anodic
MgZn2 hardening precipitates have been observed to accumulate at grain boundaries in the
T6 temper . The dissolution of these particles can result in severe intergranular corrosion
of the T6 alloy which in turn changes to exfoliation corrosion with the non-Faradaic removal
of significant amounts of material .
The most common micro-constituent particles in A7075-T6 are (Al,Cu)6(Fe,Cu) and
Al7Cu2Fe. These are also two of the most detrimental particles to the corrosion performance
of the alloy. The presence of iron and especially copper in the second phase micro-
constituents results in these particles supporting an enhanced local oxygen reduction reaction
at their surface [11,12]. This leads to an increase in the pH of the solution around the particles
. The reaction causes a potential difference between the particles and the matrix with the
particles acting as cathodes to the aluminium alloy. The copper rich Al2Cu phase has been
found to support the reduction of water at a faster rate than pure aluminium because of the
presence of metallic copper in the oxide layer of the particle. In addition to this effect the
oxides formed on the surface contain more electrically conductive copper oxides .
Whether the various intermetallics are anodic or cathodic to the matrix the result is the same,
a pit is formed either by the dissolution of the intermetallics or by their undermining and
eventual release from the surface as the surrounding matrix is dissolved .
In order to study the pitting corrosion associated with individual particles which are usually
0.1 - 20 µm in diameter , analysis techniques with a spatial resolution similar or superior
than the particle size is required [16–19]. A combination of surface specific scanning electron
microscopy (SEM), Auger electron spectroscopy (AES) and scanning Auger microscopy
(SAM) techniques, in addition to a bulk analysis technique, energy dispersive x-ray
spectroscopy (EDX), has previously been shown to provide a great deal of understanding
about the corrosion process surrounding second phase particles in multiple metal systems
The aim of this work was to investigate the corrosion processes associated with individual
second phase particles in A7075-T6 over the course of an immersion experiment and to
monitor the initiation of corrosion accompanying particles of varying copper concentration to
study the effects of galvanic couple with the alloy matrix.
A sample 1 cm2 of AA7075-T6 was wet ground with silicon carbide papers and then polished
to a 1 μm finish using diamond paste. Afterwards it was rinsed in deionised water. The metal
coupon was stored wrapped in aluminium foil in a desiccator prior to and in between
analyses. Intermetallics were identified using reflected light microscopy and marked using an
arrangement of Vickers microhardness indentations. This enabled the intermetallics to be
relocated for further analysis. EDX was performed on a range of intermetallics to determine
their bulk composition and three intermetallics were chosen for the study of localised
corrosion phenomena. These three were selected as each had a different composition, with an
increasing copper content, representing two common particle compositions and another less
common constituent particle type.
A solution of potassium chloride, 3.5 wt.%, pH 7, was produced from high purity KCl
(purchased from Sigma Aldrich) and ultra-pure water. The sample of AA7075 was immersed,
analysis face up, in the KCl solution for time periods of 0, 0.25, 0.75, 2, 4, 8 and 16 hours,
cumulative. At each time step a fresh corrosive solution was used. The same intermetallics
were repeatedly located and analysed after each time period. Following immersion, the
surface was rinsed with deionised water before being analysed by AES and EDX. AES and
SAM were performed prior to EDX to minimise the possibility of hydrocarbon contamination
depositing onto the intermetallics before surface analysis.
The investigation was performed using a scanning Auger microscope (MICROLAB 350,
Thermo Scientific, UK) fitted with an integral Thermo Scientific EDX detector. This
configuration enables the acquisition of both SAM and EDX maps from the same region of
the surface, without the need to reposition or relocate the specimen. A primary electron beam
energy of 10 keV was used for the acquisition of point Auger spectra and a beam energy of
15 keV was used for the acquisition of EDX spectra and SAM maps. The beam current for
AES and EDX spectral acquisition was 5 nA. The analysis conditions resulted in spatial
resolutions of ~40 nm in AES and ~1 μm in EDX. The depth of analysis for AES is <10 nm
resulting in information from the surface of the sample while the analysis depth of EDX is ~1
μm resulting in more bulk information. AES survey spectra were collected over an energy
range of 20 - 1800 eV and recorded with a retard ratio of 4. SAM maps were recorded with a
retard ratio of 2.8 to increase sensitivity. 128 x 128 pixel SAM images and 256 x 256 EDX
maps were produced for each of the elements present at the surface and in the bulk alloy
respectively. For each of the intermetallics these typically included maps for: aluminium,
magnesium, zinc, silicon, iron, copper, oxygen and chromium.
To remove the hydrocarbon contamination left on the surface after polishing the sample
surface was argon ion sputtered with a beam energy of 1 keV and a sample current density of
0.05 μA cm-2. Sputtering was not repeated during the experiment in order to avoid the
removal of corrosion products or aggressive ions at the surface. Topographic effects in the
SAM maps were reduced by mapping the ratio (P-B)/(P+B) of each transition where: P is the
intensity of the Auger peak and B is the intensity of the background. Thermo Scientific’s
Avantage V4.87 data system was used for acquiring and processing the Auger data and the
Noran System Seven was used for acquiring and processing the EDX data. EDX data was
quantified using a Phi-Rho-Z correction method of the integrated peak areas.
Following 16 hours of immersion, each of the studied intermetallics was cross sectioned
using focussed ion beam (FIB) milling. Milling was performed using a dual-beam FIB
microscope (FEI nano-Nova) with a beam current of <3 nA and a beam energy of 20 keV.
The post milled intermetallics were then imaged using a scanning electron microscope (JEOL
JSM-7100F) with a beam energy of 15 keV.
Scanning Kelvin probe force microscopy (SKPFM) was performed using an AFM (Bruker
Dimension Edge) with a silicon tip coated with a platinum-iridium conductive layer. A tip-
sample distance of 100 nm and a scan rate of 0.2 Hz were employed for the acquisition of
Volta potential data. The Volta potentials were acquired from intermetallics with a highly
similar composition as the three studied in this paper. Volta potentials were extracted from
the particles using the cross section analysis tool on potential maps. Analysis of the original
three intermetallics was not possible as a consequence of instrument delivery delays.
Results – Initial intermetallic morphology and composition
Prior to analysis or immersion, SEM micrographs of each of the intermetallics of interest
were acquired. These are shown in Fig. 1 a – c. The micrographs highlight the variations in
morphology of the three particles despite their similar size. The first particle (a) is
particularly angular with sharp corners while the second (b) is rectangular in shape. Both
particles are approximately 4 x 2 μm2 in size. The third particle (c) is the largest of the three,
being approximately 6 x 5 μm2 in size and is more rounded than the other intermetallics. The
light halo around the third particle is believed to be the onset of corrosion caused by the use
of aqueous polishing media employed in the preparation of the sample .
Following light Ar+ sputtering of the surface to reduce the level of hydrocarbon
contamination, AES survey spectra were collected from each of the three intermetallics
providing information about the surface composition of the particles. The spectra are shown
in Fig. 2 a-c. Each of the spectra shows the Al KLL Auger transition with peaks for
aluminium oxide and aluminium metal at 1386.5 and 1393.0 eV respectively . The broad
Al LMM transitions are also observed at ~68 eV for all of the intermetallics. Each of the
spectra also show weak Fe LMM peaks, the most intense L3M4,5M4,5 peak was located at 702.5
eV. Intense O KLL transitions are also observed for each of the intermetallics, with the
primary KL2,3L2,3 peak located at 506.5 eV. However, a noticeable reduction in the intensity of
the O KLL for the third intermetallic is observed, indicating that the surface oxide is thinner
on this particle. The third intermetallic shows intense Cu LMM peaks and the position of the
most intense Cu L3M4,5M4,5 peak was measured as 917.5 eV, indicating the presence of CuO at
the surface [25,26]. The first intermetallic also shows weak Si KLL and the Si LMM Auger
transitions. These peaks were too weak to record the exact kinetic energy but were located at
approximately 1618 and 90 eV respectively. A weak C KLL peak is also observed at ~270 eV
on this intermetallic. EDX spectra were also collected from the centre of each of the three
intermetallics and are shown in Fig. 3 a-c.
Fig. 3a shows the EDX spectrum acquired from the centre of the first intermetallic. These
provide an indication of the bulk composition of the particles. Intense aluminium and iron
peaks are observed together with weaker silicon, chromium and copper peaks. An EDX
spectrum from Intermetallic #2 is shown in Fig. 3b. The spectrum is similar to that acquired
from Intermetallic #1. Intense aluminium and iron peaks are observed together with weaker
silicon, chromium and copper peaks. Compared to the first intermetallic there is a decrease in
the silicon and an increase in the copper peak intensity respectively. The EDX spectrum from
Intermetallic #3 is shown in Fig. 3c. The primary difference in the spectrum compared to the
previous intermetallics is the presence of intense copper peaks and the absence of silicon and
EDX maps were also acquired prior to immersion and principle component analysis
performed on the raw x-ray counts using Thermo Scientific compass software. Each of the
analyses showed that the intermetallics were homogeneous in composition and that the point
spectra collected from their centre accurately represents their entire composition.
The EDX quantification of the spectra acquired at the centre of the intermetallics is shown in
Table 1, in both weight and atomic percent. The quantification was normalised to 100% and
oxygen and carbon were omitted from the quantification. The table also shows the
composition of the alloy as measured by EDX. The ratios of the higher mass transition metal
elements have been used to estimate the composition of the intermetallics types because of
the contribution of the surrounding aluminium matrix to the aluminium signal intensity by x-
ray fluorescence .
Corrosion of Intermetallic #1
The AES survey spectrum acquired from the surface of the intermetallic and the
corresponding EDX spectrum are shown in Fig. 2a and Fig. 3a, providing surface and bulk
information respectively. Based upon the EDX quantification the intermetallic is believed to
be Al12Fe3Si. These particles are known to be a minor constituent in the alloy . Electron
backscattered diffraction was attempted on each of the intermetallics to confirm their crystal
structure but Kikuchi lines of sufficient intensity could not be obtained for positive
SEM micrographs and SAM maps of the Auger transitions from elements that were present in
the survey spectra were acquired after each immersion time from the first intermetallic. These
are shown in Fig. 4. Arranging the images in a row allows the morphological and
compositional changes brought about by corrosion processes to be observed at a specific
immersion time. The SEM micrographs are shown in the first row. They reveal that there is
little significant change to the surface of the intermetallic over the immersion time. After 15
minutes the transition metal peaks become less intense and are no longer observed in the
maps after 45 minutes. However, a noticeable change occurs after 8 hours of immersion. The
aluminium and oxygen maps show a reduction in the signal from the matrix surrounding the
particle after 8 hours. This is even more apparent after 16 hours. The reduction in intensity
shows that the surface becomes rougher as surface material is removed by general corrosion
over the course of the investigation. However, the deposition of corrosion products is not
observed on the surface of the particle. The roughening of the surface causes changes in the
electron and x-ray counts because of the changing topography. A flat surface will generate
electrons and x-rays in all directions with equal intensity but a rough surface changes the
distribution of emitted radiation. A surface facing away from the x-ray/electron detector will
result in a lower count rate than a surface facing the detector. In addition to the general
roughening of the surface a dark halo is observed at the interface between the particle and the
matrix in the oxygen and aluminium maps. This lack of signal intensity suggests the
formation of a crevice where Auger electrons are unable to escape from and reach the
AES survey spectra acquired from the centre of the intermetallic at three different immersion
times are shown in Fig. 5. They illustrate how the surface of the intermetallic changes as the
immersion time progresses. These results provide clearer data than the SAM maps but only
for the point they are collected from. After 45 minutes the surface of the intermetallic is still
similar to the spectrum acquired prior to immersion. The Fe LMM peaks are clearly observed
as are the Si KLL, Al KLL and O KLL peaks. After 4 and 16 hours of immersion the intensity
of the Fe LMM peaks drops significantly compared to 45 minutes of immersion. It is possible
that the thickness of the aluminium oxide at the surface has increased slightly or there is a
slight dissolution of iron from the surface of the intermetallic as the ratio of the Fe LMM to
the Al KLL and O KLL peaks has decreased. There is however, no evidence for the deposition
of significant amounts of corrosion products on the surface as the Fe LMM peaks are still
visible and would be quickly attenuated by a few nanometres of material on the surface. This
fits well with the pattern observed in the iron SAM map, where the intensity on the
intermetallic appears to decrease significantly after 2 hours of immersion. AES spectra
acquired from the matrix close to the particle do not show the presence of iron and so if iron
is dissolved from the surface of the particle it is not redeposited onto the local matrix.
Therefore, it is believed that the reduction in the intensity of the iron Auger transition is
caused by an increase in the thickness of the aluminium oxide/hydroxide layer at the surface
of the particle.
EDX maps acquired from the particle and these are shown in Fig. 6. They reveal that the
bulk, within the top ~ 1 µm, of the intermetallic is unchanged throughout the immersion time,
while the surrounding alloy is being corroded. The shape of the intermetallic in the transition
metal maps and the magnesium and aluminium maps does not change. After 8 hours of
immersion the intensity in the oxygen and aluminium maps become more diffuse, showing a
drop in intensity from the surface surrounding the intermetallic as the alloy begins to undergo
corrosion at the periphery of the particle. This is most clearly observed in the oxygen maps as
the O Kα x-ray is of significantly lower kinetic energy than the Al Kα x-ray and is more
easily attenuated. The surface becomes rougher with the development of small pits observed
in the SEM micrographs of the intermetallic after 8 and 16 hours of immersion. These small
pits do not appear to be associated with any heterogeneity in the surface that is detectable by
EDX or AES. The dark halo that was shown in the oxygen SAM map is considerably more
pronounced in the 16 hour oxygen x-ray map, completely encircling the particle, indicating
significant crevice formation at the interface.
Each of the three intermetallics was cross sectioned, following 16 hours of immersion, using
focussed ion beam milling. This enabled information to be gathered from the sub-surface
matrix, a region not usually available to the analyst using more traditional analysis
techniques. The SEM micrographs collected from the intermetallics after ion beam milling
are shown in Fig. 7. The first intermetallic to be examined was intermetallic #2 and the
deposition of a protective platinum strap was performed as shown in Fig.7b. However,
deposition of the strap appeared to slightly erode the surface of the intermetallic and matrix
and so deposition of protective straps on intermetallics #1 and #3 was not performed.
Fig. 7a shows the micrograph from Intermetallic #1, which reveals the presence of a
significant crevice at the matrix/intermetallic interface. Additionally, in the bottom left corner
of the micrograph the onset of intergranular corrosion is observed. The matrix surrounding
the intermetallic does not appear to have corroded at an accelerated rate, while the matrix at
the interface has been severely attacked. The surface of the intermetallic remained smooth
and in the as-polished state with no deposition of corrosion products. It is possible that any
loose or poorly adhered corrosion products became dislodged during rinsing of the sample
after each corrosion step. Fig. 7b shows the micrograph from Intermetallic #2 the particle is
outlined by the dashed line in the micrograph. It shows severe crevicing around and
underneath the intermetallic. This also appears to have progressed to intergranular corrosion
of the matrix beneath the surface. The matrix shows significantly more subsurface corrosion
than that observed around Intermetallic #1. The lumps of material at the bottom of the
micrograph are matrix material that was redeposited during ion milling. The apparent internal
contrast of the particle in the image is caused by slight charging during imaging and the effect
of sharpening the image to improve the clarity of features in the micrograph. The SEM
micrograph from Intermetallic #3 after FIB milling is shown in Fig. 7c. Peripheral pitting or
'trenching' is observed on the matrix surrounding the intermetallic . The trench appears to
extend up to 2.5 μm into the alloy and almost 2 μm down along the matrix/intermetallic
interface, with the intermetallic standing proud of the surface. Examination of the corrosion
products shows the open mesh like structure they form upon deposition. The surrounding
matrix in the trench has a similar morphology.
Measurement of the Volta potential and diagnosis of galvanic activity by cation
The Volta potentials of three different intermetallics with a highly similar composition as
those studied in this paper were recorded after repolishing the specimen. It has previously
been shown that the Volta potential of intermetallics in aluminium, measured in air by
SKPFM, has a linear relationship to their open circuit potential measured in aqueous solution
. The SKPFM maps, together with their cross section analysis, are shown in Fig. 8 a - f.
Despite the significant variations in composition, the Volta potentials were found to be
almost equal for the three intermetallics. These had highly similar compositions as
Intermetallics #1, #2 and #3 and were measured to have Volta potentials differences of 310
mV, 328 mV and 309 mV compared to the matrix. These values are in line with those
previously recorded from iron and copper rich intermetallics in AA7075-T6 [9,10,29,30]. The
small difference in the values indicates that any difference in the microchemistry associated
with these particles is dominated by the kinetics of the corrosion process. The Volta potential
values have been inverted in line with other authors [10,28,30,31].
The use of Volta potential values as a measure for galvanic activity has been cautioned and so
to further explore the corrosion kinetics, the three particles were tested for their cathodic
current density [32,33]. To estimate the cathodic current density the sample was sequentially
exposed to solutions of MgCl2, pH 7, with concentrations of 0.001 M, 0.01 M and 0.1 M for
15 minutes . Mg(OH)2 is formed and deposited onto the cathodic surface when [Mg+]
[OH-]2 exceeds the solubility product. The concentration of [OH-] depends on the balance
between cathodic current density and the rate at which the OH- ions diffuse away into the
bulk solution. The latter was the same for all inclusions in this test and in the actual corrosion
exposure. The observation of Mg(OH)2 by AES enables bracketing of the cathodic current
density of the particles . To calculate the cathodic current density range that is bracketed
by magnesium deposition Equation 1 is used.
=Fδ (DOH (( K sp
[ Mg ] )0.5
[ H ] )+DH ([ H ]−K w( [ Mg ]K sp )
Where: F is the Faraday constant, DOH and DH are the diffusion coefficients for hydroxyl ions
and hydrogen ions respectively, Ksp is the solubility product for magnesium hydroxide, Kw is
the ionic product for water and δ is the diffusion distance, taken as twice the radius of
intermetallic. It has been shown that for near neutral solutions only the first term of Equation
1 is important and so the equation can be rearranged to Equation 2 . This reveals the
concentration at which magnesium is deposited for a given current density.
[ Mg ]=K sp
( iδAF DOH
[ H ] )2 (2)
AES spectra acquired from the surface of the intermetallics exposed to MgCl2 showed the
presence of magnesium after exposure to the different solutions. AES was used as it provided
information regarding the initial deposition of magnesium hydroxide onto the particle surface
while a magnesium signal in EDX would not be observed until after substantial deposition.
The first intermetallic showed the Mg KLL peak after exposure to 0.01 M, the second also
after 0.01 M although with greater peak intensity and the third after exposure to 0.001 M
MgCl2. The kinetic energy of the Mg KL2,3L2,3 peak in each case was ~1180 eV, which is
consistent with magnesium hydroxide or oxide and not the metallic magnesium present in the
alloy . The Mg KLL Auger transition was not detected in high resolution spectra acquired
from the intermetallics prior to immersion in the magnesium chloride solutions. The presence
of magnesium hydroxide on each of the intermetallic particles shows that they are able to
support a cathodic reaction capable of generating hydroxyl ions. In each case for the three
intermetallics radius of the particles was taken from the average of the x and y axis lengths of
the intermetallics. Additionally, the diffusion coefficient for hydroxyl ions was taken as 0.52
x 10-8 m2 s-1 and a magnesium hydroxide solubility product of 5 x 10-11 was used .
Discussion Intermetallic #1
Based upon the EDX quantification Intermetallic #1 is believed to be Al12Fe3Si, which has
been identified as a minor particle type in AA7075 [4,9]. Examination of the SEM
micrographs shows that no material was deposited onto the surface of the intermetallic during
the investigation. The AES point spectra acquired from the intermetallic after 45 min, 4 hours
and 16 hours of immersion show that the surface of the intermetallic does not change
significantly. However, there is a drop in the intensity of the Fe LMM peaks indicating that
the oxide/hydroxide layer thickness has increased slightly or a thin layer of material is
coating the intermetallic.
Examining the AES point spectra, SAM maps and EDX maps shows there is no evidence of
dissolution of the intermetallic or deposition of corrosion products onto the surface of the
intermetallic. This evidence supports the belief that this intermetallic was not acting as a
significant cathode that was coupled to the matrix in a manner resulting in the accelerated
anodic dissolution of the matrix, despite the significant cathodic Volta potential measured
from this particle composition. Corrosion of the surrounding alloy surface is observed but
does not appear to be enhanced by the presence of the intermetallic as no trench with sloping
sides from the matrix to the intermetallic interface was observed. However, the FIB-SEM
micrograph and the oxygen/aluminium EDX maps clearly show that a crevice formed at the
A weak Mg KLL peak was observed in the AES point spectrum acquired from an
intermetallic of the same composition and similar size after exposure to a 0.01 M MgCl2
solution. This brackets the cathodic current density of this particle type to 0.0061 - 0.0019 A
cm-2 . The behaviour of this intermetallic can be explained by considering it to have a low
electrochemical activity. Despite being apparently nobler than the matrix, with a corrosion
current greater than that of the alloy, the particle cannot sustain a sufficient cathodic current
to severely impact the corrosion kinetics. Therefore, as a result of the data gathered from the
surface, a particle of this composition might not be expected to have a significant impact on
the pitting corrosion performance of the alloy. However, the presence of the crevice and the
apparent onset of sub-surface intergranular corrosion suggest that even particles which do not
form active galvanic couples with the matrix can act as sites for localised corrosion
processes. It is possible that the presence of the crevice resulted in detachment of the particle
from the matrix resulting in poor electrical contact decoupling it from the matrix. Therefore,
the behaviour of the particle studied in this work may not be representative of this
composition as a whole in the alloy.
If the immersion time were to be extended it is likely that the intermetallic would be
undermined and removed from the matrix creating a pit . The corrosion behaviour of this
particle is similar to the behaviour observed for Al3Fe type intermetallics particles in Al-Fe
binary alloys and AA6061-T6 [13,36].
Corrosion of Intermetallic #2
Based upon the EDX quantification of Intermetallic #2, this particle may be Al7Fe2Cu, which
has been observed in AA2024 but is not widely reported as a common micro-constituent
particle type in AA7075 [4,5]. This particle type is known to occur in cast aluminium alloys
containing iron and copper . However, the EDX quantification shows the Cu/Fe ratio to
be 0.4 which is consistent with the range of 0.2 to 0.7 expected for the well-known and
numerous (Al,Cu)6(Fe,Cu) type intermetallics . Hence the particle is considered most
likely to be (Al,Cu)6(Fe,Cu).
As performed for the Intermetallic #1 SAM maps and SEM micrographs were acquired from
the intermetallic as a function of immersion time and these are shown in Fig. 9. The SEM
micrographs show that there is no apparent change to the bulk shape of the intermetallic over
16 hours. However, after 16 hours some corrosion product deposition is observed adjacent to
the particle around another feature and small clusters of material are observed on the surface
of the particle. The aluminium and oxygen maps are fairly consistent with each other and
only noticeably change after 8 hours of immersion. After which their signals become more
diffuse, with the region showing the depleted signal intensity region expanding out from the
intermetallic onto the matrix. The formation of a halo around the intermetallic is also
observed after 8 hours of immersion and is even more noticeable after 16 hours. The
matrix/intermetallic interface is the region where corrosive attack is most aggressive. The
reduced intensity in the oxygen and aluminium maps in this halo indicates the formation of a
crevice at the interface. A feature is observed in the initial chromium SAM map which
appears as a ring around three sides of the intermetallic and this is reflected in the oxygen
SAM map as a region of low signal intensity. This is not seen in the corresponding EDX
maps indicating it is a surface specific feature which does not extend down into the surface.
After 45 minutes of immersion the transition metal Auger peaks are no longer sufficiently
intense to be detected in SAM analysis, as was also observed for Intermetallic #1. This
indicates that the aluminium oxide/hydroxide layer on the intermetallic has grown thicker or
that there is slight deposition of general corrosion products over the sample surface leading to
attenuation of the signal.
Significant sub-surface volume expansion appears to have led to the formation of a blister to
the right of the intermetallic, as observed in the 8 and 16 hour SEM micrographs in Fig. 9 and
more closely in Fig. 10. The outline of the blister feature is highlighted by the dashed line in
the micrograph. This blister and the associated material distort the SAM mapping signal
because of the resulting complex topography at the sample surface. However, magnesium and
silicon are observed to be present as part of the surrounding material. The blister itself
appears to be a thin film of aluminium oxide. AES point spectra that were acquired from
Points 1, 2 and 3 in Fig. 10, are shown in Fig. 11. They show that the surface of the blister
appears to be aluminium oxide and the bright material in the bottom right of the SEM
micrograph consists of magnesium and silicon. The relative intensity of the Al KLL peak is
reduced, whereas the O KLL peak is much increased. Thus it appears that the Si and Mg are
present as oxides or hydroxides.
AES survey spectra acquired from Intermetallic #2 after 45 minutes, 4 hours and 16 hours are
shown in Fig. 12. After 45 minutes the surface of the intermetallic is still similar to the
spectrum acquired prior to immersion. The Fe LMM peaks are clearly present as are the Si
KLL, Al KLL and O KLL peaks. Between 45 minutes and 4 hours the intensity of the peaks
appears to remain fairly constant although a slight decrease in the signal to noise of the Fe
LMM peaks is observed. The Fe LMM peaks are no longer observed after 16 hours, while the
Al and O KLL peaks increased significantly in intensity indicating that the surface of the
particle was covered in initial corrosion products.
The SEM micrographs together with the EDX maps for the second intermetallic are shown in
Fig. 13. The bulk shape of the intermetallic, as detected and outlined by EDX, is unchanged
throughout the immersion time while the surrounding alloy surface is corroded, as shown by
the SEM micrographs and the oxygen SAM map.
After 8 hours of immersion a similar change occurs for the oxygen EDX map as was
observed for the oxygen SAM map, with the signal becoming more diffuse on the
intermetallic and increasing in intensity on the alloy surface. The surrounding alloy is
undergoing corrosion with the surface becoming significantly rougher and the development
of small pits which are observed in the SEM micrograph and by the increase in oxygen
intensity at localised points in the oxygen EDX maps and a decrease in intensity in the SAM
oxygen maps. Examination of the EDX maps in Fig. 13, after 16 hours of immersion, shows
four discrete regions. The metal matrix (Al and Zn maps), the intermetallic (Fe, Cu and Cr
maps), a magnesium and silicon containing particulate (Mg and Si maps) and a region
depleted in matrix x-rays (O and Cl maps). Based on the SEM micrograph and the
SAM/EDX elemental maps, this latter region is believed to be a blister. The loose particulate
type material observed in the bottom right of the SEM micrograph after 16 hours of
immersion, shown larger in Fig. 10, appears to consist of MgSi2 and Al2O3 and is believed to
be loose material not to be associated with the blister. EDX quantification from this region is
consistent with these materials.
After 16 hours of immersion, the EDX chlorine map shows a strong enrichment in the region
of the blister. An EDX point spectrum acquired from the centre of this region revealed an
approximate 1:1 ratio of aluminium and oxygen in addition to ~5 at.% chlorine. The
increased depth of analysis of EDX enables it to probe further into the further below the
blister while AES provides the information about the capping layer. The lack of alloy x-ray
counts in the elemental EDX maps from this region, in particular aluminium, suggests that
the blister is covering a void from which few x-rays are generated. These are consistent with
the main blister material consisting of a thin layer of Al2O3, underneath which, acidic
aluminium products and aluminium oxychlorides are present . Although chlorine is
detected in the EDX map, and EDX point spectra from this region it is not observed in the
AES point spectra or the SAM maps, this is likely because the AlCl3 is concentrated beneath
the Al2O3 blister “skin” and any AlCl3 present at the surface would be removed when the
sample is rinsed to remove excess KCl solution from the surface.
In the manner previously described, an intermetallic of the same composition and similar size
to Intermetallic #2 was exposed to a 0.01 M MgCl2 solution. A weak Mg KLL peak was
observed in the Auger spectra acquired from this intermetallic which brackets the cathodic
current density of this intermetallic type to 0.0043 - 0.0135 A cm-2. From the FIB-SEM
micrograph in Fig. 7b, it can be seen that a significant crevice was formed at the
matrix/intermetallic interface and there is significant sub-surface removal of material leading
to the onset of integranular corrosion. The particle itself is highlighted by a dashed white line
in this image. A local cathode would accelerate the “unzipping” of grains through the
dissolution of the anodic MgZn2 precipitates dispersed along the grain boundaries [39,40].
Discussion Intermetallic #2
The second intermetallic behaves differently to the first intermetallic. Although the AES
point spectra acquired from the particle after 45 min and 4 hours exposure showed little
change, the spectrum acquired after 16 hours exposure revealed the complete attenuation of
the iron Auger transitions. Examination of the larger SEM micrograph in Fig. 10 reveals
small clusters of material deposited or formed on the surface of the intermetallic. These likely
explain the attenuation of the iron peaks. The micrograph and oxygen EDX maps also
highlight the more localised dissolution of the matrix immediately adjacent to the particle.
After 16 hours of exposure, a blister was observed which initiated in the matrix adjacent to
the intermetallic. The Auger/EDX data and SEM micrographs indicate that the blister “skin”
is an aluminium oxide layer, under which a substantial AlCl3 corrosion product has formed.
The SEM/AES/EDX results are similar to those previously observed for metastable pitting
attack at MnS inclusions in stainless steel caused by undercutting of the metal oxide .
The results from exposure of this intermetallic type to MgCl2 reveals that it has a cathodic
current density higher than that observed for Intermetallic #1. It is sufficiently cathodically
active compared to the matrix such that it can support accelerated removal of the local matrix.
However, as will be shown in the next section, the cathodic current density is significantly
less than that of the copper rich Intermetallic #3 and so the local matrix is not as severely
corroded. Despite the reduced removal rate of the adjacent matrix it is believed that the
severity of the intergranular sub-surface corrosion around the particle is a result of galvanic
coupling with the intermetallic particle.
Corrosion of Intermetallic #3
The AES survey spectrum initially acquired from the surface of the intermetallic is shown in
Fig. 2c. The spectrum shows considerably more intense iron and copper Auger transitions
than the previous two intermetallics. It also shows a significantly weaker O KLL peak,
indicating that the surface oxide is thinner on this particle. The EDX spectrum from the
intermetallic is shown in Fig. 3c. Based upon the EDX quantification, the Cu/Fe ratio is 2.3
which is consistent with the range 1.5 - 2.5 expected for Al7Cu2Fe type particles. As such this
intermetallic is believed to be Al7Cu2Fe .
The SEM micrographs and the SAM maps of the elements present in the third intermetallic
are shown in Fig. 14. Unlike the previous two intermetallics, the micrographs show a
considerable change in appearance over the immersion time. Up until 45 minutes the surface
of the particle appears smooth. However, after 2 hours of immersion the deposition of
corrosion products on the intermetallic is evident. As the exposure continues, the amount of
material gradually increases until, after 16 hours, the intermetallic is completely covered.
Corrosive attack and removal of the alloy material is most aggressive at the
matrix/intermetallic interface, with a pronounced increase in the observed roughness of the
surrounding surface. The corrosion products appear to consist of aluminium, oxygen, silicon
and, after longer immersion times, zinc. As the immersion time continues the area covered by
these products grows and they appear to be deposited primarily in clusters on the edge of the
The copper SAM maps show that the surface of the intermetallic remains rich in copper and
is not coated by corrosion products until after 2 hours of immersion. The aluminium maps
initially show low intensity from the region of the intermetallic. After 45 minutes, the
decrease in intensity in aluminium associated with the edge of the intermetallic becomes less
defined and after 2 hours the region of low intensity grows as corrosion products deposit. At
this point it is also observed that the aluminium intensity in the centre of the intermetallic
increases and remains up to 16 hours of immersion. The oxygen maps show a decrease in
intensity at the matrix/intermetallic interface surrounding the intermetallic as the surface is
attacked. After 2 hours of immersion the aluminium and oxygen SAM maps, Fig. 14, and
Auger point spectra, Fig. 15, show the corrosion products deposited on and around the
intermetallic to be aluminium hydroxide based, as would be expected.
The zinc SAM maps do not appear to show any intensity changes until after 45 minutes of
immersion. At this point an increase in the intensity all over the whole surface of the
intermetallic surface is observed. After 2 and 4 hours of immersion, no zinc is observed in the
maps. It reappears after 8 hours at the centre of the intermetallic and is also apparent in the
map acquired after 16 hours of immersion.
AES survey spectra were acquired from the centre of the intermetallic after 45 minutes, 4
hours and 16 hours and these spectra are shown in Fig. 15. After 45 minutes the surface of the
intermetallic is still similar to the spectrum acquired prior to immersion. However, small
amounts of zinc and silicon are now apparent which were not observed in the initial
spectrum. The spectra after 4 hours show that the intensity of the Si, Al and O KLL peaks
have significantly increased, while the Cu and Fe LMM peaks exhibit particularly low
intensities. This fits well with the pattern observed in the copper and iron SAM maps, where
the intensity of these elements on the intermetallic appears to significantly decrease after 2
hours of immersion. The intensity of the Si and Al KLL peaks after 4 hours suggests that the
first corrosion products deposited on the surface of the intermetallic are SiO2 and Al(OH)3.
After 16 hours of immersion a significant change in the spectrum is observed. The intensity
of the Si KLL has significantly reduced and intense Zn LMM, Al KLL and O KLL peaks are
observed. This indicates that corrosion products rich in aluminium and zinc are deposited
onto the surface after the initial silicon and aluminium containing deposits. It is possible that
these products are zinc hydrotalcite .
The EDX maps from the third intermetallic are shown in Fig. 16 revealing any changes in the
bulk of the particle and surrounding matrix. The aluminium and magnesium maps correlate
well up to 4 hours of immersion. The depleted intensity region in the aluminium and
magnesium EDX maps is clearly associated with the position of the intermetallic. After 4
hours, the depleted signal region increases in area. This is the result of matrix material being
corroded and removed from the surface.
The oxygen EDX map follows a similar pattern as that observed for the previous two
intermetallics. After 45 minutes of exposure, the high intensity oxygen signal is associated
with the intermetallic site and a small ring around the intermetallic. After 2 hours, the oxygen
intensity on the intermetallic increases significantly in small clusters (corrosion deposits) and
the presence of these corrosion deposits in the oxygen map remains essentially the same
throughout the rest of the exposure. In addition to the clusters of corrosion deposits, after 2
hours there is a slight increase in oxygen intensity on the matrix surrounding the
intermetallic. This becomes more clearly observable after 8 hours of immersion. The oxygen
and silicon EDX maps show a matching relationship. It is observed that after 45 minutes the
silicon maps show the same small clusters of intensity as the oxygen maps during the 16
hours of exposure. Although the maps are highly similar the difference in the maps is the
bright halo observed around the particle in the oxygen map that gradual increases in size
throughout the investigation as the matrix is corroded. This feature is not observed in the
silicon EDX maps. Silicon was not observed in the EDX spectrum from the intermetallic
prior to exposure, Fig. 3b.
The intensity from the intermetallic in the zinc EDX map is seen to increase after 8 and 16
hours exposure. This is in accordance with the increased zinc intensity observed in the AES
point spectrum after 16 hours exposure, Fig. 15, indicating that the zinc is present on the
particle surface to a greater thickness than initially indicated by AES. The copper, iron,
chromium and nickel EDX maps all remain unchanged during the exposure period. A slight
drop in the signal intensity for each element is observed at the position of the corrosion
product clusters, as a consequence of these products attenuating the x-ray signal.
The results from exposure of this intermetallic type to MgCl2 reveals that it has a cathodic
current density higher than that observed for Intermetallics #1 and #2. The Mg KLL peak was
observed in AES point spectra acquired from an intermetallic of the same composition after
exposure to a 0.001 M MgCl2, pH 7, solution, bracketing the critical concentration of
magnesium hydroxide precipitation to <0.001 M. Therefore, the cathodic current density can
be estimated to be >6.7 x 10-3 A cm-2. This value is in general agreement with previous work
using a microcapillery cell [27,29,33].
Discussion Intermetallic #3
The third intermetallic was shown to be galvanically active to the surrounding matrix.
Corrosion products containing oxygen and silicon were observed to accumulate in clusters on
the surface after 2 hours of immersion. It is interesting to note that silicon rich oxide deposits
have previously been observed in SEM/AES/EDX studies of corrosion initiation at mixed
oxide inclusion sites on stainless steels . Also, in that case, silicon was not observed in
the inclusions, yet a significant concentration of oxidised silicon was observed close to the
pitting site. The origin of the silicon is unknown, but could originate from MgSi2
intermetallics which are known to be anodic in the T6 temper . Their dissolution would
result in silicon in solution.
Zinc was found within the intermetallic at a concentration of around 2 at.% as shown in Table
1. Zinc was observed to deposit onto the intermetallic after longer immersion times (8 and 16
hours). It seems probable that the zinc is being selectively dissolved from the intermetallic, as
a consequence of its higher electrochemical activity than the other transition metal elements
present and then deposited as a corrosion product, likely zinc hydroxide. It is believed that
corrosion products containing silicon and aluminium are initially deposited onto the surface
followed by products containing zinc and aluminium. The zinc may also originate from the
gradual dissolution of anodic MgZn2 hardening precipitates located at grain boundaries
within the alloy as the alloy begins to corrode.
The oxygen EDX maps and micrographs indicate that corrosive attack is concentrated around
the matrix surrounding the intermetallic, resulting in dissolution of the aluminium alloy. The
formation of a trench surrounding the intermetallic is clearly revealed by the FIB-SEM
images shown in Fig.7c and the depletion in the intensity of the oxygen and aluminium
AES/EDX maps surrounding the intermetallic. The clusters of deposition products are
concentrated on the perimeter of the particle. This is a result of the increased pH in this
region caused by the generation of hydroxyl ions from the cathodic reaction on the surface of
the intermetallic. Aluminium, magnesium and zinc cations diffusing from the dissolving
matrix will reach a fixed distance across the intermetallic surface before they encounter a
critical pH at which the metal hydroxides will be deposited. The deposition of products at this
critical value results from the corrosion products being initially deposited at the outer edges
of the intermetallic and then later covering more of the particle surface as hydroxyl ion
generation is limited by the deposited material.
The difference in the observed corrosion behaviour and the enhanced corrosion of the matrix
adjacent to the particle compared to Intermetallics #1 and #2 is a result of the greater galvanic
coupling between the particle and the matrix. The Al7Cu2Fe type intermetallics are generally
found to be more noble by ~100 mV than the (Al,Cu)6(Fe,Cu) intermetallic particles .
Consequently, the results confirm the cathodic nature of the Al7Cu2Fe type intermetallic
particle with respect to the surrounding alloy matrix [4,27,30].
Despite the increased nobility and severity of corrosion surrounding the third intermetallic
compared to the second intermetallic. No evidence for intergranular corrosion was observed
in the sub-surface matrix surrounding intermetallic following cross sectioning. It is believed
that the severity of the galvanic couple with the matrix results in sufficiently faster corrosion
such that a crevice cannot form at the particle interface. As such, the reduced current density
of the first two intermetallics may promote crevice formation instead of severe trenching.
Based upon the observations of the sub-surface matrix around the second intermetallic it is
possible that despite their reduced galvanic activity these intermetallics can have a significant
influence on the corrosion performance of the alloy as they promote intergranular corrosion.
The corrosion microchemistry associated with three different intermetallic particles with
varying composition, (Intermetallic #1 Al12Fe3Si, Intermetallic #2 (Al,Cu)6(Fe,Cu) and
Intermetallic #3 - Al7Cu2Fe) were studied by means of SEM, EDX, AES and SKPFM.
1. Volta potential measurements from particles with highly similar compositions to the
three studied here revealed only slight differences in the apparent nobility of the
particles despite significant variations in their cathodic current density. This
highlights the importance of including additional analytical techniques to extrapolate
the corrosion behaviour and corrosion kinetics of individual second phase particle
2. The cathodic current density varied in the order: particle #1 < #2 < #3, in line with the
intermetallic particles copper content. The severity of the micro corrosion process
associated with each of the particle types showed the same relationship.
3. The high iron content, Al12Fe3Si intermetallic studied in this work had the lowest
current density of the particles at 0.0061 - 0.0019 A cm-2. No evidence for the
traditional trenching associated with cathodic particles was found. However, crevice
corrosion was observed at the particle interface with an indication of the sub-surface
initiation of intergranular corrosion.
4. The (Al,Cu)6(Fe,Cu) intermetallic was more cathodic than the iron containing particle
with a current density of 0.0043 - 0.0135 A cm-2 and resulted in significant sub-
surface intergranular corrosion at the particle interface and signs of the onset of
5. While the copper rich, Al7Cu2Fe, intermetallic was found to be significantly more
cathodic than the other particles, with a current density of >6.7 x 10-3 A cm-2, no
evidence for the sub-surface initiation of intergranular corrosion was observed around
the particle studied in this work.
The authors wish to thank the Atomic Weapons Establishment and the University of Surrey
for funding this research.
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Fig. 1 SEM micrographs of Intermetallics #1 - #3 (a - c), prior to immersion in the
corrosion solution. The numbered points show the positions where AES and EDX point
analyses were performed.
Fig. 2 AES survey spectra collected from each of the three intermetallics prior to
immersion in the corrosion solution. All of the major Auger transitions present in the
spectra have been labelled. The spectra a – c refer to Intermetallics #1 - #3 respectively.
Fig. 3 EDX spectra acquired from the centre of the three intermetallics studied in this
investigation prior to exposure to the corrosion solution. Their compositions are shown
in Table 1. The major peaks have been identified. The spectra a – c refer to
intermetallics #1 - #3 respectively.
Fig. 4 SEM micrographs (first row) and the SAM maps of Al, O, Fe, Cr and Si at
different immersion times for the first intermetallic. The different immersion times and
AES element maps are highlighted with the titles in the columns and rows respectively.
Fig. 5 AES survey spectra from the centre of the first intermetallic after 45 min, 4 h and
16 h immersion in the corrosion solution.
Fig. 6 SEM micrographs (first row) and the EDX maps of Al, O, Fe, Cu, Cr, Zn, Si and
Mg at different immersion times for the first intermetallic. The different immersion
times and x-ray element maps are highlighted with the titles in the columns and rows
Fig. 7 a - c SEM micrographs of the three intermetallics following cross sectioning by
FIB milling. The micrographs a – c refer to intermetallics #1 - #3 respectively. The
dashed line in b highlights the particle. The micrographs were acquired with different
incident angles as a consequence of the trench orientations and instrument geometry
constraints. The micrographs have been sharpened to highlight features.
Fig. 8 a - c Volta potential maps of three second phase particles with highly similar
compositions as the three studied in this paper, z range 750 mV and 20 µm scan size.
Fig. 8 d - f cross section analysis through each of the Volta potential maps with the
measurement positions marked. Volta potential values are inverted.
Fig. 9 SEM micrographs (first row) and the SAM maps of Al, O, Fe, Cu, Cr, Mg and Si
at different immersion times for Intermetallic #2.
Fig. 10 SEM micrograph of Intermetallic #2 after 16 hours of immersion. The positions
from which AES point spectra were acquired from a feature of interest, highlighted by
the dashed line, are labelled Points 1, 2 and 3.
Fig. 11 AES point spectra from Points 1, 2 and 3, on the region surrounding
Intermetallic #2 after 16 hours of immersion in the corrosion solution.
Fig. 12 AES point spectra from the centre of Intermetallic #2 after 45 min, 4 hours and
16 hours immersion in the corrosion solution.
Fig. 13 SEM micrographs (first row) and the EDX maps of Al, O, Fe, Cu, Cr, Zn, Si, Mg
and Cl at different immersion times for Intermetallic #2.
Fig. 14 SEM micrographs (first row) and the SAM maps of Al, O, Cu, Fe and Zn at
different immersion times for Intermetallic #3.
Fig. 15 AES survey spectra from the centre of Intermetallic #3 after: 45 min, 4 hours
and 16 hours immersion in the corrosion solution.
Fig. 16 SEM micrographs (first row) and the x-ray maps of Al, O, Fe, Cu, Cr, Zn, Si
and Mg at different immersion times for Intermetallic #3.
Table 1 EDX quantification from each of the three intermetallics investigated in this
work and the aluminium alloy. Values shown in weight and atomic percent
(normalised), ND means not detected
Region/Composition Al Fe Cu Si Cr Ni Mg Zn
Alloy (wt.%) Bal 0.3 1.5 0.1 0.2 ND 2.4 5.4
Intermetallic 1 (wt.%) 58.3 27.9 3.1 3.3 5.9 ND ND 1.5
Intermetallic 1 (at.%) 72.5 16.7 1.6 5.1 3.7 ND ND 0.4
Intermetallic 2 (wt.%) 61.0 18.7 9.2 0.3 5.2 0.6 0.5 4.6
Intermetallic 2 (at.%) 76.7 11.3 4.9 0.3 3.4 0.3 0.7 2.4
Intermetallic 3 (wt.%) 50.6 12.033.
9ND ND 1.5 0.2 3.0
Intermetallic 3 (at.%) 71.1 8.018.
3ND ND 0.9 0.3 1.7
This paper presents the results of a multi technique analysis of the localised corrosion associated with three individual intermetallic particles in AA7075-T6. The three intermetallics: Al12Fe3Si, (Al,Cu)6(Fe,Cu) and Al7Cu2Fe showed a significant increase in the severity of the corrosion process in line with the cathodic current density measured from each individual intermetallics.