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,

RG7 4PR

*c.mallinson@surrey.ac.uk

Abstract

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

Introduction

Aluminium alloys are the second most widely used engineering metallic alloys after steels

[1]. The alloy 7075-T6, is much used within the aerospace industry because of its particularly

high strength to weight ratio [2]. The 7xxx series (Al-Zn-Mg-Cu alloys) provide the potential

for precipitation hardening throughout a range of compositions [3]. The alloying elements

present in aluminium are either in solid solution or segregated as secondary phase micro-

constituent particles [2]. These particles fall into one of three categories: hardening

precipitates, dispersoids and constituent particles [3]. Hardening precipitates range in size up

to ten nanometres and MgZn2 is a common example [3]. Dispersoids control grain size, grow

to hundreds of nanometres, and include Al3Ti and Al3Zr [1]. 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 [6].

The presence of “weak spots” which interfere with the homogeneity of the oxide film are

thought to be the cause of this susceptibility [7]. 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 [8].

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 [9].

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 [10]. The anodic

MgZn2 hardening precipitates have been observed to accumulate at grain boundaries in the

T6 temper [10]. 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 [10].

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

[13]. 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 [14].

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 [15].

In order to study the pitting corrosion associated with individual particles which are usually

0.1 - 20 µm in diameter [10], 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

[16,20–23].

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.

Experimental

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 [3].

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 [24]. 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

chromium peaks.

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 [19].

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 [9]. 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

identification.

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

detector.

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 [27]. 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

precipitation

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

[28]. 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 [23]. 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 [23]. To calculate the cathodic current density range that is bracketed

by magnesium deposition Equation 1 is used.

iA

=Fδ (DOH (( K sp

[ Mg ] )0.5

−K w

[ H ] )+DH ([ H ]−K w( [ Mg ]K sp )

0.5

)) (1)

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 [34]. This reveals the

concentration at which magnesium is deposited for a given current density.

[ Mg ]=K sp

( iδAF DOH

+Kw

[ 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 [35]. 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 [23].

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

matrix/intermetallic interface.

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 [23]. 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 [13]. 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 [37]. 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 [9]. 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 [38]. 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 [41].

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 [9].

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

particle.

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 [42].

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 [21]. 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 [10]. 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 [9].

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.

Conclusions

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

types.

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

trenching.

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.

Acknowledgments

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

respectively.

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

Graphical abstract

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.