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Effect of artificial aging on intergranular corrosion of extruded AlMgSi alloy with small Cu content
Corrosion Science 48 (2006) 1528–1543 www.elsevier.com/locate/corsci
Effect of artificial aging on intergranular corrosion of extruded AlMgSi alloy with small Cu content Gaute Svenningsen a,*, Magnus Hurlen Larsen a, John Charles Walmsley b,d, Jan Halvor Nordlien c, Kemal Nisancioglu a a
d
Department of Materials Technology, Norwegian University of Science and Technology, NO-7491 Trondheim, Norway b SINTEF Materials and Chemistry, NO-7465 Trondheim, Norway c Hydro Aluminium, R&D Materials Technology, NO-4265 Ha˚vik, Norway Department of Physics, Norwegian University of Science and Technology, NO-7491 Trondheim, Norway Received 25 January 2005; accepted 24 May 2005 Available online 1 September 2005
Abstract The effect of artificial aging parameters on the corrosion performance of air cooled AlMgSi(Cu) model alloy extrusions was investigated. Accelerated corrosion test revealed that the extrusions were highly susceptible to intergranular corrosion (IGC) in the naturally aged condition. However, IGC susceptibility was reduced, and finally eliminated, by artificial aging. Overaging introduced slight pitting susceptibility. EDS X-ray mapping in FE-TEM revealed Mg2Si and Q-phase (Al4Cu2Mg8Si7) grain boundary precipitates and a continuous Cuenriched grain boundary film. IGC susceptibility was related to the Cu-enriched grain boundary film. Increased IGC resistance was caused by coarsening of the grain boundary film by aging. Pitting susceptibility by over aging evolved due to coarsening of the Q-phase particles in the grain bodies. 2005 Elsevier Ltd. All rights reserved.
*
Corresponding author. Tel.: +47 73 59 40 56; fax: +47 73 59 40 83. E-mail address: [email protected] (G. Svenningsen).
0010-938X/$ - see front matter 2005 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2005.05.045
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Keywords: A. Aluminium; A. Intermetallics; B. FE-TEM; B. STEM; C. Intergranular corrosion
1. Introduction Aluminium 6000 series alloys (AlMgSi and AlMgSiCu alloys) are extensively used as extruded profiles, e.g. in architectural, automotive and marine applications [1] owing to favourable strength-to-weight ratio, compliance to surface treatment and good corrosion resistance. However, susceptibility to intergranular corrosion (IGC) may occur as a result of unfavourable heat treatment or alloying as reviewed in detail in [2,3]. Some of the important factors are the following: IGC of aluminium alloys is the result of microgalvanic cell action at the grain boundary, related to grain boundary precipitates, which are either more active or more noble than the surrounding solid solution aluminium matrix. In the case of 6000-series alloys, precipitation of phases, such as Si in alloys with excess Si relative to the stoichiometry of the Mg2Si-phase or the Q-phase (Al5Cu2Mg8Si6) in the case of Cu-containing alloys, may cause IGC. However, when in contact with water, Siprecipitates will be covered by an insulating layer of SiO2, and are therefore not effective cathodes [4]. Depletion of silicon probably renders the matrix in the particle-free zone, that is observed along both sides of the grain boundaries, active and susceptible to corrosion by microgalvanic action. In the case of Q-phase precipitation, formation of both a noble phase along the grain boundaries and depletion of the particle-free zone provides an increased driving force for the microgalvanic effect and increased danger of IGC. Copper content as low as 0.1 wt.% can cause Q-phase precipitation and IGC for certain Mg/Si ratios and thermomechanical histories, as shown in previous papers [2,3]. A model AlMgSi alloy, with Mg/Si ratio = 0.87 and containing 0.13% Cu, became susceptible to IGC as a result of slow-cooling in air after extrusion, whereas cooling at a higher speed in water produced a temper resistant to IGC. These phenomena were correlated, respectively, with the precipitation and absence of the Qphase along the grain boundaries. While artificial heat treatment to obtain the peak strength T6-condition reduced the IGC-susceptibility of the air cooled extrusion, the same treatment made the water cooled extrusion slightly susceptible to IGC. In contradiction to the foregoing results for slow-cooled extrusions, IGC is reported to increase as a result of artificial aging of certain 6000-series alloys, reaching a maximum susceptibility at peak hardness T6 condition [5–12]. Moreover, overaging reduced or eliminated IGC and introduced pitting [13–15], probably due to extensive precipitation in the matrix, thus reducing the electrochemical potential difference between the matrix and the depleted zone [16]. Additionally, coarsening of the grain boundary precipitates was possibly an important factor in reducing IGC susceptibility, especially if the presence of a continuous active or noble grain boundary film can be postulated [17]. However, overaging may not always be a desirable option because it reduces the mechanical strength. As indicated above, the water cooled AlMgSiCu extrusions are likely to be resistant to IGC since grain boundary precipitation of the Q-phase can be avoided or
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limited. Artificial aging caused precipitation of the Q-phase along the grain boundaries and thereby increased susceptibility to IGC [2]. Explanation of increased IGCresistance resulting from artificial aging of the air cooled extrusions was however less straightforward. It is, furthermore, apparent from the literature reviewed above that the role of artificial aging on the IGC of 6000-series alloys is not well understood, and a systematic investigation of the subject is not available. The objective of the present paper is therefore to investigate how artificial aging affects the IGC susceptibility of an extruded model AlMgSi alloy containing about 0.13 wt.% Cu. The study is limited to air cooled extrusions.
2. Experimental The investigated material was a flat plate (78 · 0.27 cm) extrusion of a model AlMgSiCu alloy with the chemical composition given in Table 1. The temperature at the die-exit was 575–580 C during extrusion. The extruded profile was cooled in ambient air (air cooling) after extrusion. This cooling method was quite slow, and it took about 4 min before the temperature reached a value below 200 C. The profile was stretched 0.5% in the longitudinal direction and naturally aged by storing at room temperature (T4 temper) for 3 months prior to artificial aging and corrosion testing. 2.1. Aging Samples of size 5 · 3.9 · 0.27 cm were artificially aged in two types of oil bath and a fluidised oxide bed depending on the aging temperature, as summarised in Table 2. The samples were quenched in water immediately after aging. All samples were aged directly from the naturally aged condition (T4), i.e., no solution heat treatment was applied before the artificial aging treatment.
Table 1 Chemical composition of sample material Mg
Si
Fe
Mn
Cu
Cr
Ti
Zn
Al
0.57
0.62
0.21
0.21
0.13
0.001
0.01
0.01
Balance
Table 2 Aging parameters Temperature
Aging medium
Aging time [s]
140 C 185 C 220 C
Oil bath Oil bath Fluidised oxide bed
5000–1 720 000 500–86 400 100–86 400
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Vickers hardness was measured by using 1 kg load on the as-extruded (and aged) surface, i.e. without grinding or polishing. Electronic conductivity was measured by using a Sigmatest D v 2.068 instrument operated at 60 kHz, also on the extruded surface. 2.2. Corrosion testing The samples were corrosion tested according to the British Standard BS-ISO 11846, method B [18]. This test involved degreasing in acetone, alkaline etching (5 min in 7.5 wt.% NaOH at 55–60 C) followed by 24 h immersion in an acidified salt solution (30 g NaCl and 10 ml concentrated HCl per litre). Afterwards, the samples were washed in distilled water and ethanol and then dried. The corrosion susceptibility, including maximum corrosion depth, was evaluated visually from cross sections of corrosion tested material examined in an optical microscope. This accelerated test has been widely used for qualitative ranking of Al alloys with respect to IGC since its introduction by Hug in 1941 [19]. Despite giving high level of acceleration in an aggressive environment, the test has been successfully employed for distinguishing resistant materials from those with likelihood of IGC susceptibility. In our earlier related work [2], we have compared and revalidated the accelerated method with long-term atmospheric exposure of the present alloy of interest in a marine-industrial atmosphere. 2.3. Grain boundary microstructure Samples for field emission gun transmission electron microscope (FE-TEM) investigation were prepared in the standard manner by electropolishing followed by 30–90 min argon ion-sputtering, using a Gatan PIPS model 391 ion-thinner operating at 3.5–3.0 kV at a thinning angle of 3.5–4.0. FE-TEM was performed with a JEOL 2010F instrument operating at 200 kV in scanning transmission electron microscope (STEM) mode. Energy dispersive X-ray spectroscopy (EDS) data were acquired using an Oxford Instruments INCA system. A drift compensation system allowed maps to be acquired over periods of up to several tens of minutes. Quantitative grain boundary analysis using FE-(S)TEM instruments has been reported in alloys containing several percent of copper [20–23]. Similar analytical procedures were followed here in order to obtain a qualitative relative impression in the grain boundary levels of copper in our alloy under different aging conditions.
3. Results 3.1. Corrosion Corrosion susceptibility and mode changed as a function of aging temperature and time as summarised in Table 3 and illustrated in Figs. 1–3. The as-extruded
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Table 3 Results from corrosion test as a function of artificial aging time and temperature Aging
Dominating corrosion mode (average depth)
Max corr. depth [lm]
Temp. [C]
Time [s]
140 140 140 140 140 140 140 140 140 140
5000 10 000 18 000 36 000 86 400 172 800 345 600 864 000 2 592 000 5 184 000
Uniform IGC (300 lm) Uniform IGC (250 lm) Uniform IGC (300 lm) Uniform IGC (300 lm) Uniform IGC (100–300 lm) Uniform IGC (100 lm) Uniform IGC (75 lm) Local IGC No local corrosion Pitting
550 410 480 440 420 400 150 375 75 320
185 185 185 185 185 185 185 185
500 1000 2500 5000 10 000 18 000 36 000 86 400
Uniform IGC (300 lm) Uniform IGC (250 lm) Uniform IGC (100–200 lm) Local IGC No local corrosion No local corrosion Pitting, IGC? Pitting
500 450 400 400 30 25 100 100
220 220 220 220 220 220 220 220 220 220
100 250 500 1000 2500 5000 10 000 18 000 36 000 86 400
Uniform IGC (200–300 lm) Uniform IGC (100–300 lm) Slight IGC (20–100 lm) No local corrosion No local corrosion Pitting, and possible IGC Pitting, and possible IGC Pitting and IGC Pitting Heavy pitting
430 400 250 620 610 270 300 250 275 300
Uniform IGC (300 lm)
550
No artificial aging
samples were highly susceptible to IGC (uniform IGC). The susceptibility was first reduced to a more localised IGC, confined to discrete areas on the sample (local IGC) and then finally eliminated as a result of increasing aging time (no localised corrosion). Pitting susceptibility was introduced as a result of overaging. The corrosion cross sections shown in Fig. 4 indicated that pitting resulted from grain boundary attack and crystallographic tunnelling at the pit front and subsequent corrosion of whole grains. The time required for these transitions, summarised in Fig. 5, became shorter with increasing aging temperature. The Figs. 1–3 include hardness and conductivity measurements for easy monitoring of the aging process. The highest corrosion resistance was approximately associated with the maximum hardness obtained by the particular aging procedure. Conductivity (or resistivity) measurements were used to indicate the extent of solute
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Fig. 1. Hardness, electric conductivity and corrosion susceptibility as a function of artificial aging time at 140 C.
Fig. 2. Hardness, electric conductivity and corrosion susceptibility as a function of artificial aging time at 185 C.
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Fig. 3. Hardness, electric conductivity and corrosion susceptibility as a function of artificial aging time at 220 C.
Fig. 4. Overaged samples after corrosion testing.
depletion as a result of precipitation of the hardening phases during heat treatment [24,25]. The conductivity measurements here were comparable to the values reported in the literature [24,25]. The conductivity in the corrosion resistant zone was in the range 28–29.5 MS/m for the three aging temperatures. This indicated possibly an
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Corrosion mode: IGC No local corrosion
Aging temperature [°C]
225
No localised corrosion
200
Pitting Pitting
175
Intergranular corrosion 150
125
100 100
1000
10000
100000
1000000
10000000
Aging time [s]
Fig. 5. Effect of artificial aging on the corrosion mode.
optimal precipitation of the hardening phases in terms of size, number density and distribution to give the observed corrosion resistance. 3.2. Microstructure The grain boundaries of the air cooled sample were decorated with large and discrete Mg2Si-phase and Q-phase particles, as characterised in detail in Refs. [2,3]. By use of FE-TEM in the present work, the presence of a thin Cu enriched grain boundary film could in addition be documented on samples both in the T4 and T6 tempers, as shown in Figs. 6 and 7, respectively. In both figures, micrographs (a) show typical Q-phase precipitates. Elemental maps of the marked grain boundaries near the particles indicated enrichment of the grain boundary by an apparently continuous Cu-rich film of nanometer scale thickness (micrographs (c) compared to (b), (d) and (e)). The presence of Cu at the grain boundaries, away from visible Q-phase precipitates was verified by EDS point analyses, as shown in Fig. 8. Copper was the only enriched element detected at the grain boundaries outside the Q-phase and Mg2Si precipitates. The Cu-enriched film was less evident on the overaged sample, as indicated by elemental mapping of the particle free regions of grain boundaries, as shown in Fig. 9. While Cu was not detectable on many grain boundaries as in the case of Fig. 9, grain boundaries still exhibiting a Cu-rich film were also present, as indicated by spot EDS analysis of such a grain boundary shown in Fig. 10. It is clear, however, that the
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Fig. 6. TEM analysis of a naturally aged sample.
amount of grain boundary Cu detectable on the overaged samples was in all cases significantly smaller than the grain boundary Cu on samples corresponding to the T4 and T6 tempers.
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Fig. 7. TEM analysis of a peak aged (18 000 s at 185 C) sample.
4. Discussion In our earlier work [2,3] grain boundary corrosion of the present material was correlated with the precipitation of the Q-phase along the grain boundaries. However, the presence of coarse Q-phase particles was not deemed adequate to explain the knife-edge type continuous corrosion attack along the grain boundaries. The present work demonstrated that this type of attack is caused by the presence of a nanoscale thick Cu-rich continuous film along the grain boundaries. Based on thermodynamic calculations presented earlier for the specific alloy composition [3], the film cannot consist of Cu alone, and the precipitation of an AlCu-phase (e.g. Al2Cu) can also be ruled out. The Cu-rich film must therefore be a precursor of the more easily detectable Q-phase. Since the analytical characterisation of the film was clearly at
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400 Cu L
350
1
Mg K
300
Grain boundary Matrix
1
Grain boundary Matrix
1
50
250 200 150
Si K
100
1
Si K
40 30 Cu K
20
1
1
10
50 0
Cu K
60
1
Counts
Counts
Mg K
0.8
1.0
1.2
1.4
1.6
1.8
2.0
0 7.6
2.2
7.8
8.0
8.2
8.4
8.6
8.8
9.0
9.2
Energy [keV]
Energy [keV]
(a) Naturally aged sample. 400 350
Cu L
1
Mg K
300
1
70
Grain boundary Matrix
1
200 Si K
100
Si K
1.0
1.2
Grain boundary Matrix
1.4 1.6 1.8 Energy [keV]
40 30 Cu K
20
1
50 0 0.8
1
50
250 150
Cu K
60
Counts
Counts
Mg K
1
1
10 2.0
2.2
0 7.6
7.8
8.0
8.2 8.4 8.6 Energy [keV]
8.8
9.0
9.2
(b) Peakaged sample (18 000 s at 185 ˚C). Fig. 8. EDS spectrum of point-analysis on grain boundaries and in the matrix.
the detection limit of the FE-TEM instrument, possible enrichment of the light elements Mg and Si could probably not be detected. The detection sensitivity of these elements is lower than for Cu, and their EDS peaks overlap with the adjacent intense Al peak (Fig. 8). It can thus be postulated that IGC propagates along the narrow anodic solute depleted zone with the Cu-rich film and the corroded grain walls acting as the cathodes. The exposed grain walls are expected to act as cathodes because they are still solute rich, and they are therefore nobler than the depleted zone. Moreover, Cu from the Cu-rich film and the Q-phase particles can dissolve and redeposit on the exposed grain walls in the presence of an acidic chloride environment [26–28], as illustrated in Fig. 11. Although direct evidence could not be demonstrated, we attribute the increased resistance to IGC by artificial aging to a coarsening process, by which the Cu-rich film retracts, at least locally, to form patches of film and discrete platelets along the grain boundaries as sketched in Fig. 12. In particular, analysis of a three dimensional volume of a TEM foil sample at a grain boundary by EDS could fail to detect the gaps between the patches of film as shown in Fig. 13. Existence of grain boundaries along which the continuity of the film is disrupted, as envisaged in Figs. 12 and 13 would be sufficient to arrest IGC before the corrosion front reaches the next patch
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Fig. 9. TEM analysis of a overaged (86 400 s at 185 C) sample.
70
35
60
Cu K
Grain boundary Matrix
1
25
Counts
Counts
50 40 Cu K
30
1
10
5 8.0
8.2
8.4
8.6
Energy [keV]
8.8
9.0
9.2
1
15 10
7.8
Cu K
20
20
0 7.6
Grain boundary Matrix
30
0 7.6
Cu K
7.8
8.0
8.2
8.4
8.6
8.8
1
9.0
9.2
Energy [keV]
Fig. 10. EDS spectrum of point-analysis on grain boundaries and in the matrix. Overaged sample (86 400 s at 185 C).
of Cu-rich film. It probably also impedes initiation of IGC as indicated by the transition from uniform to more localised type of IGC, followed by a more passive condition of the entire surface with increasing time of aging.
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Re-deposited Cu (cathodic) Remnant of Q-phase (detached)
Aluminium matrix
Solute depleted zone (active) Cu-containing film (cathodic)
Q-phase grain boundary precipitate (cathodic)
Fig. 11. Conceptual sketch of the IGC mechanism.
Fig. 12. Conceptual sketch of the effect of aging on the microstructure and the corrosion mode.
The present results are in conflict with the general contention [5–12] that IGC susceptibility of Cu-containing 6000-series alloys increases as a result of artificial aging. Our data showed that this conclusion applied to rapidly cooled (water quenched) extrusions [2], but the opposite was true for the slowly cooled (air cooled) extrusions.
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Plan view of grain boundary
Cross sectional view of grain boundary
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Grain boundary
Continuous film
Coarsened film (discontinuous) Aging time
Fig. 13. Conceptual sketch showing coarsening of the Cu enriched grain boundary film.
Thus, the importance of the cooling rate as an important parameter in the IGC of such materials was emphasised. Continuing aging and overaging leads to coarsening of precipitates in the bulk of grains along with increased solute depletion of the surrounding aluminium matrix. The resulting increased electrochemical driving force between the intermetallic particles favours pitting corrosion.
5. Conclusions • The model AlMgSi alloy tested in this work (Mg/Si ratio = 0.87, Cu content = 0.13%) became susceptible to IGC as a result of slow cooling in air after extrusion. • Artificial aging reduced the susceptibility to IGC, and close to peak hardness, the material became resistant to localised corrosion. Overaging introduced pitting susceptibility. Thus, the need for proper heat treatment to obtain high corrosion resistance was demonstrated. • The effect of aging temperature and time, which are important parameters in determining the corrosion properties, as well as the mechanical properties of the metal, was systematically documented. • IGC susceptibility was related to the presence of a nanoscale Cu-rich film along the grain boundaries along with the presence of coarse Q-phase particles and the microgalvanic coupling between these cathodic sites and the adjacent solute depleted zone. • Increased resistance against IGC obtained by artificial aging was attributed to the coarsening of the Cu-rich film along with increased solute depletion of the surrounding matrix by precipitation of the hardening phases.
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Acknowledgements This work was part of a Norwegian national research program entitled ‘‘Light Metal Surface Science’’, supported by The Norwegian Research Council, Hydro Aluminium, Profillakkering AS, Norsk Industrilakkering AS, NORAL AS, Jotun Powder Coatings AS, Electro Vacuum AS and DuPont Powder Coatings.
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Effect of artificial aging on intergranular corrosion of extruded AlMgSi alloy with small Cu content Gaute Svenningsen a,*, Magnus Hurlen Larsen a, John Charles Walmsley b,d, Jan Halvor Nordlien c, Kemal Nisancioglu a a
d
Department of Materials Technology, Norwegian University of Science and Technology, NO-7491 Trondheim, Norway b SINTEF Materials and Chemistry, NO-7465 Trondheim, Norway c Hydro Aluminium, R&D Materials Technology, NO-4265 Ha˚vik, Norway Department of Physics, Norwegian University of Science and Technology, NO-7491 Trondheim, Norway Received 25 January 2005; accepted 24 May 2005 Available online 1 September 2005
Abstract The effect of artificial aging parameters on the corrosion performance of air cooled AlMgSi(Cu) model alloy extrusions was investigated. Accelerated corrosion test revealed that the extrusions were highly susceptible to intergranular corrosion (IGC) in the naturally aged condition. However, IGC susceptibility was reduced, and finally eliminated, by artificial aging. Overaging introduced slight pitting susceptibility. EDS X-ray mapping in FE-TEM revealed Mg2Si and Q-phase (Al4Cu2Mg8Si7) grain boundary precipitates and a continuous Cuenriched grain boundary film. IGC susceptibility was related to the Cu-enriched grain boundary film. Increased IGC resistance was caused by coarsening of the grain boundary film by aging. Pitting susceptibility by over aging evolved due to coarsening of the Q-phase particles in the grain bodies. 2005 Elsevier Ltd. All rights reserved.
*
Corresponding author. Tel.: +47 73 59 40 56; fax: +47 73 59 40 83. E-mail address: [email protected] (G. Svenningsen).
0010-938X/$ - see front matter 2005 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2005.05.045
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Keywords: A. Aluminium; A. Intermetallics; B. FE-TEM; B. STEM; C. Intergranular corrosion
1. Introduction Aluminium 6000 series alloys (AlMgSi and AlMgSiCu alloys) are extensively used as extruded profiles, e.g. in architectural, automotive and marine applications [1] owing to favourable strength-to-weight ratio, compliance to surface treatment and good corrosion resistance. However, susceptibility to intergranular corrosion (IGC) may occur as a result of unfavourable heat treatment or alloying as reviewed in detail in [2,3]. Some of the important factors are the following: IGC of aluminium alloys is the result of microgalvanic cell action at the grain boundary, related to grain boundary precipitates, which are either more active or more noble than the surrounding solid solution aluminium matrix. In the case of 6000-series alloys, precipitation of phases, such as Si in alloys with excess Si relative to the stoichiometry of the Mg2Si-phase or the Q-phase (Al5Cu2Mg8Si6) in the case of Cu-containing alloys, may cause IGC. However, when in contact with water, Siprecipitates will be covered by an insulating layer of SiO2, and are therefore not effective cathodes [4]. Depletion of silicon probably renders the matrix in the particle-free zone, that is observed along both sides of the grain boundaries, active and susceptible to corrosion by microgalvanic action. In the case of Q-phase precipitation, formation of both a noble phase along the grain boundaries and depletion of the particle-free zone provides an increased driving force for the microgalvanic effect and increased danger of IGC. Copper content as low as 0.1 wt.% can cause Q-phase precipitation and IGC for certain Mg/Si ratios and thermomechanical histories, as shown in previous papers [2,3]. A model AlMgSi alloy, with Mg/Si ratio = 0.87 and containing 0.13% Cu, became susceptible to IGC as a result of slow-cooling in air after extrusion, whereas cooling at a higher speed in water produced a temper resistant to IGC. These phenomena were correlated, respectively, with the precipitation and absence of the Qphase along the grain boundaries. While artificial heat treatment to obtain the peak strength T6-condition reduced the IGC-susceptibility of the air cooled extrusion, the same treatment made the water cooled extrusion slightly susceptible to IGC. In contradiction to the foregoing results for slow-cooled extrusions, IGC is reported to increase as a result of artificial aging of certain 6000-series alloys, reaching a maximum susceptibility at peak hardness T6 condition [5–12]. Moreover, overaging reduced or eliminated IGC and introduced pitting [13–15], probably due to extensive precipitation in the matrix, thus reducing the electrochemical potential difference between the matrix and the depleted zone [16]. Additionally, coarsening of the grain boundary precipitates was possibly an important factor in reducing IGC susceptibility, especially if the presence of a continuous active or noble grain boundary film can be postulated [17]. However, overaging may not always be a desirable option because it reduces the mechanical strength. As indicated above, the water cooled AlMgSiCu extrusions are likely to be resistant to IGC since grain boundary precipitation of the Q-phase can be avoided or
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limited. Artificial aging caused precipitation of the Q-phase along the grain boundaries and thereby increased susceptibility to IGC [2]. Explanation of increased IGCresistance resulting from artificial aging of the air cooled extrusions was however less straightforward. It is, furthermore, apparent from the literature reviewed above that the role of artificial aging on the IGC of 6000-series alloys is not well understood, and a systematic investigation of the subject is not available. The objective of the present paper is therefore to investigate how artificial aging affects the IGC susceptibility of an extruded model AlMgSi alloy containing about 0.13 wt.% Cu. The study is limited to air cooled extrusions.
2. Experimental The investigated material was a flat plate (78 · 0.27 cm) extrusion of a model AlMgSiCu alloy with the chemical composition given in Table 1. The temperature at the die-exit was 575–580 C during extrusion. The extruded profile was cooled in ambient air (air cooling) after extrusion. This cooling method was quite slow, and it took about 4 min before the temperature reached a value below 200 C. The profile was stretched 0.5% in the longitudinal direction and naturally aged by storing at room temperature (T4 temper) for 3 months prior to artificial aging and corrosion testing. 2.1. Aging Samples of size 5 · 3.9 · 0.27 cm were artificially aged in two types of oil bath and a fluidised oxide bed depending on the aging temperature, as summarised in Table 2. The samples were quenched in water immediately after aging. All samples were aged directly from the naturally aged condition (T4), i.e., no solution heat treatment was applied before the artificial aging treatment.
Table 1 Chemical composition of sample material Mg
Si
Fe
Mn
Cu
Cr
Ti
Zn
Al
0.57
0.62
0.21
0.21
0.13
0.001
0.01
0.01
Balance
Table 2 Aging parameters Temperature
Aging medium
Aging time [s]
140 C 185 C 220 C
Oil bath Oil bath Fluidised oxide bed
5000–1 720 000 500–86 400 100–86 400
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Vickers hardness was measured by using 1 kg load on the as-extruded (and aged) surface, i.e. without grinding or polishing. Electronic conductivity was measured by using a Sigmatest D v 2.068 instrument operated at 60 kHz, also on the extruded surface. 2.2. Corrosion testing The samples were corrosion tested according to the British Standard BS-ISO 11846, method B [18]. This test involved degreasing in acetone, alkaline etching (5 min in 7.5 wt.% NaOH at 55–60 C) followed by 24 h immersion in an acidified salt solution (30 g NaCl and 10 ml concentrated HCl per litre). Afterwards, the samples were washed in distilled water and ethanol and then dried. The corrosion susceptibility, including maximum corrosion depth, was evaluated visually from cross sections of corrosion tested material examined in an optical microscope. This accelerated test has been widely used for qualitative ranking of Al alloys with respect to IGC since its introduction by Hug in 1941 [19]. Despite giving high level of acceleration in an aggressive environment, the test has been successfully employed for distinguishing resistant materials from those with likelihood of IGC susceptibility. In our earlier related work [2], we have compared and revalidated the accelerated method with long-term atmospheric exposure of the present alloy of interest in a marine-industrial atmosphere. 2.3. Grain boundary microstructure Samples for field emission gun transmission electron microscope (FE-TEM) investigation were prepared in the standard manner by electropolishing followed by 30–90 min argon ion-sputtering, using a Gatan PIPS model 391 ion-thinner operating at 3.5–3.0 kV at a thinning angle of 3.5–4.0. FE-TEM was performed with a JEOL 2010F instrument operating at 200 kV in scanning transmission electron microscope (STEM) mode. Energy dispersive X-ray spectroscopy (EDS) data were acquired using an Oxford Instruments INCA system. A drift compensation system allowed maps to be acquired over periods of up to several tens of minutes. Quantitative grain boundary analysis using FE-(S)TEM instruments has been reported in alloys containing several percent of copper [20–23]. Similar analytical procedures were followed here in order to obtain a qualitative relative impression in the grain boundary levels of copper in our alloy under different aging conditions.
3. Results 3.1. Corrosion Corrosion susceptibility and mode changed as a function of aging temperature and time as summarised in Table 3 and illustrated in Figs. 1–3. The as-extruded
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Table 3 Results from corrosion test as a function of artificial aging time and temperature Aging
Dominating corrosion mode (average depth)
Max corr. depth [lm]
Temp. [C]
Time [s]
140 140 140 140 140 140 140 140 140 140
5000 10 000 18 000 36 000 86 400 172 800 345 600 864 000 2 592 000 5 184 000
Uniform IGC (300 lm) Uniform IGC (250 lm) Uniform IGC (300 lm) Uniform IGC (300 lm) Uniform IGC (100–300 lm) Uniform IGC (100 lm) Uniform IGC (75 lm) Local IGC No local corrosion Pitting
550 410 480 440 420 400 150 375 75 320
185 185 185 185 185 185 185 185
500 1000 2500 5000 10 000 18 000 36 000 86 400
Uniform IGC (300 lm) Uniform IGC (250 lm) Uniform IGC (100–200 lm) Local IGC No local corrosion No local corrosion Pitting, IGC? Pitting
500 450 400 400 30 25 100 100
220 220 220 220 220 220 220 220 220 220
100 250 500 1000 2500 5000 10 000 18 000 36 000 86 400
Uniform IGC (200–300 lm) Uniform IGC (100–300 lm) Slight IGC (20–100 lm) No local corrosion No local corrosion Pitting, and possible IGC Pitting, and possible IGC Pitting and IGC Pitting Heavy pitting
430 400 250 620 610 270 300 250 275 300
Uniform IGC (300 lm)
550
No artificial aging
samples were highly susceptible to IGC (uniform IGC). The susceptibility was first reduced to a more localised IGC, confined to discrete areas on the sample (local IGC) and then finally eliminated as a result of increasing aging time (no localised corrosion). Pitting susceptibility was introduced as a result of overaging. The corrosion cross sections shown in Fig. 4 indicated that pitting resulted from grain boundary attack and crystallographic tunnelling at the pit front and subsequent corrosion of whole grains. The time required for these transitions, summarised in Fig. 5, became shorter with increasing aging temperature. The Figs. 1–3 include hardness and conductivity measurements for easy monitoring of the aging process. The highest corrosion resistance was approximately associated with the maximum hardness obtained by the particular aging procedure. Conductivity (or resistivity) measurements were used to indicate the extent of solute
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Fig. 1. Hardness, electric conductivity and corrosion susceptibility as a function of artificial aging time at 140 C.
Fig. 2. Hardness, electric conductivity and corrosion susceptibility as a function of artificial aging time at 185 C.
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Fig. 3. Hardness, electric conductivity and corrosion susceptibility as a function of artificial aging time at 220 C.
Fig. 4. Overaged samples after corrosion testing.
depletion as a result of precipitation of the hardening phases during heat treatment [24,25]. The conductivity measurements here were comparable to the values reported in the literature [24,25]. The conductivity in the corrosion resistant zone was in the range 28–29.5 MS/m for the three aging temperatures. This indicated possibly an
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Corrosion mode: IGC No local corrosion
Aging temperature [°C]
225
No localised corrosion
200
Pitting Pitting
175
Intergranular corrosion 150
125
100 100
1000
10000
100000
1000000
10000000
Aging time [s]
Fig. 5. Effect of artificial aging on the corrosion mode.
optimal precipitation of the hardening phases in terms of size, number density and distribution to give the observed corrosion resistance. 3.2. Microstructure The grain boundaries of the air cooled sample were decorated with large and discrete Mg2Si-phase and Q-phase particles, as characterised in detail in Refs. [2,3]. By use of FE-TEM in the present work, the presence of a thin Cu enriched grain boundary film could in addition be documented on samples both in the T4 and T6 tempers, as shown in Figs. 6 and 7, respectively. In both figures, micrographs (a) show typical Q-phase precipitates. Elemental maps of the marked grain boundaries near the particles indicated enrichment of the grain boundary by an apparently continuous Cu-rich film of nanometer scale thickness (micrographs (c) compared to (b), (d) and (e)). The presence of Cu at the grain boundaries, away from visible Q-phase precipitates was verified by EDS point analyses, as shown in Fig. 8. Copper was the only enriched element detected at the grain boundaries outside the Q-phase and Mg2Si precipitates. The Cu-enriched film was less evident on the overaged sample, as indicated by elemental mapping of the particle free regions of grain boundaries, as shown in Fig. 9. While Cu was not detectable on many grain boundaries as in the case of Fig. 9, grain boundaries still exhibiting a Cu-rich film were also present, as indicated by spot EDS analysis of such a grain boundary shown in Fig. 10. It is clear, however, that the
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Fig. 6. TEM analysis of a naturally aged sample.
amount of grain boundary Cu detectable on the overaged samples was in all cases significantly smaller than the grain boundary Cu on samples corresponding to the T4 and T6 tempers.
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Fig. 7. TEM analysis of a peak aged (18 000 s at 185 C) sample.
4. Discussion In our earlier work [2,3] grain boundary corrosion of the present material was correlated with the precipitation of the Q-phase along the grain boundaries. However, the presence of coarse Q-phase particles was not deemed adequate to explain the knife-edge type continuous corrosion attack along the grain boundaries. The present work demonstrated that this type of attack is caused by the presence of a nanoscale thick Cu-rich continuous film along the grain boundaries. Based on thermodynamic calculations presented earlier for the specific alloy composition [3], the film cannot consist of Cu alone, and the precipitation of an AlCu-phase (e.g. Al2Cu) can also be ruled out. The Cu-rich film must therefore be a precursor of the more easily detectable Q-phase. Since the analytical characterisation of the film was clearly at
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400 Cu L
350
1
Mg K
300
Grain boundary Matrix
1
Grain boundary Matrix
1
50
250 200 150
Si K
100
1
Si K
40 30 Cu K
20
1
1
10
50 0
Cu K
60
1
Counts
Counts
Mg K
0.8
1.0
1.2
1.4
1.6
1.8
2.0
0 7.6
2.2
7.8
8.0
8.2
8.4
8.6
8.8
9.0
9.2
Energy [keV]
Energy [keV]
(a) Naturally aged sample. 400 350
Cu L
1
Mg K
300
1
70
Grain boundary Matrix
1
200 Si K
100
Si K
1.0
1.2
Grain boundary Matrix
1.4 1.6 1.8 Energy [keV]
40 30 Cu K
20
1
50 0 0.8
1
50
250 150
Cu K
60
Counts
Counts
Mg K
1
1
10 2.0
2.2
0 7.6
7.8
8.0
8.2 8.4 8.6 Energy [keV]
8.8
9.0
9.2
(b) Peakaged sample (18 000 s at 185 ˚C). Fig. 8. EDS spectrum of point-analysis on grain boundaries and in the matrix.
the detection limit of the FE-TEM instrument, possible enrichment of the light elements Mg and Si could probably not be detected. The detection sensitivity of these elements is lower than for Cu, and their EDS peaks overlap with the adjacent intense Al peak (Fig. 8). It can thus be postulated that IGC propagates along the narrow anodic solute depleted zone with the Cu-rich film and the corroded grain walls acting as the cathodes. The exposed grain walls are expected to act as cathodes because they are still solute rich, and they are therefore nobler than the depleted zone. Moreover, Cu from the Cu-rich film and the Q-phase particles can dissolve and redeposit on the exposed grain walls in the presence of an acidic chloride environment [26–28], as illustrated in Fig. 11. Although direct evidence could not be demonstrated, we attribute the increased resistance to IGC by artificial aging to a coarsening process, by which the Cu-rich film retracts, at least locally, to form patches of film and discrete platelets along the grain boundaries as sketched in Fig. 12. In particular, analysis of a three dimensional volume of a TEM foil sample at a grain boundary by EDS could fail to detect the gaps between the patches of film as shown in Fig. 13. Existence of grain boundaries along which the continuity of the film is disrupted, as envisaged in Figs. 12 and 13 would be sufficient to arrest IGC before the corrosion front reaches the next patch
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Fig. 9. TEM analysis of a overaged (86 400 s at 185 C) sample.
70
35
60
Cu K
Grain boundary Matrix
1
25
Counts
Counts
50 40 Cu K
30
1
10
5 8.0
8.2
8.4
8.6
Energy [keV]
8.8
9.0
9.2
1
15 10
7.8
Cu K
20
20
0 7.6
Grain boundary Matrix
30
0 7.6
Cu K
7.8
8.0
8.2
8.4
8.6
8.8
1
9.0
9.2
Energy [keV]
Fig. 10. EDS spectrum of point-analysis on grain boundaries and in the matrix. Overaged sample (86 400 s at 185 C).
of Cu-rich film. It probably also impedes initiation of IGC as indicated by the transition from uniform to more localised type of IGC, followed by a more passive condition of the entire surface with increasing time of aging.
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Re-deposited Cu (cathodic) Remnant of Q-phase (detached)
Aluminium matrix
Solute depleted zone (active) Cu-containing film (cathodic)
Q-phase grain boundary precipitate (cathodic)
Fig. 11. Conceptual sketch of the IGC mechanism.
Fig. 12. Conceptual sketch of the effect of aging on the microstructure and the corrosion mode.
The present results are in conflict with the general contention [5–12] that IGC susceptibility of Cu-containing 6000-series alloys increases as a result of artificial aging. Our data showed that this conclusion applied to rapidly cooled (water quenched) extrusions [2], but the opposite was true for the slowly cooled (air cooled) extrusions.
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Plan view of grain boundary
Cross sectional view of grain boundary
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Grain boundary
Continuous film
Coarsened film (discontinuous) Aging time
Fig. 13. Conceptual sketch showing coarsening of the Cu enriched grain boundary film.
Thus, the importance of the cooling rate as an important parameter in the IGC of such materials was emphasised. Continuing aging and overaging leads to coarsening of precipitates in the bulk of grains along with increased solute depletion of the surrounding aluminium matrix. The resulting increased electrochemical driving force between the intermetallic particles favours pitting corrosion.
5. Conclusions • The model AlMgSi alloy tested in this work (Mg/Si ratio = 0.87, Cu content = 0.13%) became susceptible to IGC as a result of slow cooling in air after extrusion. • Artificial aging reduced the susceptibility to IGC, and close to peak hardness, the material became resistant to localised corrosion. Overaging introduced pitting susceptibility. Thus, the need for proper heat treatment to obtain high corrosion resistance was demonstrated. • The effect of aging temperature and time, which are important parameters in determining the corrosion properties, as well as the mechanical properties of the metal, was systematically documented. • IGC susceptibility was related to the presence of a nanoscale Cu-rich film along the grain boundaries along with the presence of coarse Q-phase particles and the microgalvanic coupling between these cathodic sites and the adjacent solute depleted zone. • Increased resistance against IGC obtained by artificial aging was attributed to the coarsening of the Cu-rich film along with increased solute depletion of the surrounding matrix by precipitation of the hardening phases.
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Acknowledgements This work was part of a Norwegian national research program entitled ‘‘Light Metal Surface Science’’, supported by The Norwegian Research Council, Hydro Aluminium, Profillakkering AS, Norsk Industrilakkering AS, NORAL AS, Jotun Powder Coatings AS, Electro Vacuum AS and DuPont Powder Coatings.
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