Effect of high temperature heat treatment on intergranular corrosion of AlMgSi(Cu) model alloy

Corrosion Science 48 (2006) 258–272 www.elsevier.com/locate/corsci

Effect of high temperature heat treatment on intergranular corrosion of AlMgSi(Cu) model alloy Gaute Svenningsen a,*, Magnus Hurlen Larsen a, Jan Halvor Nordlien b, Kemal Nisancioglu a a

Department of Materials Technology, Norwegian University of Science and Technology, NO-7491 Trondheim, Norway b Hydro Aluminium, R&D Materials Technology, NO-4265 Ha˚vik, Norway Received 29 November 2004; accepted 15 December 2004 Available online 12 March 2005

Abstract Copper containing 6000-series aluminium alloys may become susceptible to intergranular corrosion (IGC) as a result of improper thermomechanical processing. Effect of cooling rate after solution heat treatment on the corrosion behaviour of a model AlMgSi(Cu) alloy of nominal composition (wt%) 0.6 Mg, 0.6 Si, 0.2 Fe, 0.2 Mn and 0.1 Cu was investigated. Slow cooling rates were simulated by isothermal treatment for predetermined times in lower temperature baths immediately after solution heat treatment. Treatment for 10–100 s at temperatures below 400
C introduced susceptibility to IGC. Longer heat treatment at the same temperatures introduced susceptibility to pitting. A corrosion resistant time zone was found between the zones of IGC and pitting at temperatures lower than 350
C. Quenching in water after solution heat treatment prevented IGC. IGC was related to microgalvanic coupling between the noble Q-phase (Al4Mg8Si7Cu2) grain boundary precipitates and the adjacent depleted zone. Pitting was attributed to coarse particles in the matrix. Possible mechanisms causing the corrosion resistant intermediate zone are discussed. The results indicate possible methods for obtaining increased corrosion resistance of similar alloys by proper thermal processing.  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.2004.12.003

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Keywords: A. Aluminium; A. Intermetallics; B. SEM; C. Intergranular corrosion

1. Introduction AlMgSi alloys (6000 series) are generally known to have high corrosion resistance [1–3]. However, several factors may introduce IGC susceptibility, e.g. Si in excess of the Mg2Si ratio [4], presence of Cu [5,6], and thermomechanical treatment [1,2]. It has been claimed that the susceptibility to IGC increases during aging, and it is highest at the peak hardness [6–10]. IGC susceptibility is caused by grain boundary precipitates formed during aging. If the corrosion potential of these particles are different from that of the matrix, the grain boundaries will be decorated with microgalvanic cells, and susceptibility to IGC is expected. The grain boundary precipitates can be noble and the adjacent depleted zone active, or vice versa, the grain boundary precipitates can be more active than the adjacent zone, depending on the type of phase(s) precipitated [2,3,7]. Slow quenching after solution heat treatment, which increases the probability of grain boundary precipitation [11], may also introduce IGC [2,7,8,12,13]. This problem can be avoided by using higher quenching rates. However, the quenching rates required may not always be attainable in practice, e.g. on thick sections. Additionally, high cooling rates are generally not desirable because they may introduce residual stresses in the material [12]. Overaging may be beneficial; it can reduce, or remove, IGC susceptibility at the expense of introducing pitting [7,14–19]. Reduced susceptibility appears to be the result of coarsening of the particles at the grain boundaries [20] and in the matrix [14,21]. While the former breaks the continuity of the microgalvanic cells, the latter reduces the difference in the corrosion potential between the depleted zone and the matrix because solute elements are lost to the precipitates formed. Cu addition may reduce the corrosion resistance of AlMgSi alloys. Susceptibility to IGC increases with increasing Cu content [1,5,6,22–24], apparently over a critical value of 0.1 wt% [5]. Higher critical Cu levels (0.2 wt% and 0.4 wt%) have also been reported [1,25,26]. IGC on AlMgSiCu alloys has also been related to the presence of noble SiCu and AlMgSiCu grain boundary precipitates and a microgalvanic coupling between these particles and the matrix [14–16]. In addition to the b-phase (Mg2Si) and its precursors, pure Si may precipitate if the Mg:Si ratio is low [27]. The age hardening precipitation sequence in these alloys was extensively investigated and described in detail elsewhere [2,28,29]. Grain boundary precipitates of the b-phase are not associated with IGC susceptibility, unless present as a continuous film [1]. The precipitation sequence in AlMgSiCu alloys has traditionally been assumed to be analogous to that of AlMgSi without Cu [30,31]. However, the presence of Cu may affect the precipitation sequence [32,33] and the number of possible intermetallic phases formed, such as the added precipitation of the quaternary Q-phase (Al5Mg8Si6Cu2) or the binary h-phase (Al2Cu) [1,33,34]. The presence of Q-phase in

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commercial wrought alloys in 6000 series AlMgSi alloys with Cu additions and 2000 series AlCuMg alloys with Si additions is well known [1,31,34,35]. According to Chacrabarti and Laughlin [34] the Q-phase, metallic Si and the b-phase are the expected precipitates when the Cu content is in the range 0.2–0.5%. Grain boundary precipitates of Q and b phases were observed on AlMgSi alloys after high temperature heat treatment [36–40]. Phase diagrams calculated by thermodynamical simulation packages suggest that the Q-phase may grow at the expense of the b-phase at temperatures around 350–450
C [41,42]. This work is intended as the second part of a series of papers for reporting research results on IGC of AlMgSi(Cu) model alloys. In the first paper [43], we showed that the model AlMgSi alloy with 0.2% Cu content became susceptible to IGC if the cooling rate after extrusion was slow (air cooled). Susceptibility to IGC was associated with the precipitation of the Q-phase particles along the grain boundaries. Rapidly cooled samples (water quenched) were resistant to IGC because there was no Q-phase precipitation along the grain boundaries. However, high quenching rates may not always be attainable or desirable, respectively, due to large section thickness or the danger of introducing residual stresses. A detailed investigation of the effect of quenching rate on IGC, as is the case for, e.g., for some 2000 and 7000 series alloys [2,12], is not yet available for AlMgSi alloys containing small amounts of Cu. The objective of this work is therefore to investigate the effect of exposing such alloys to intermediate temperatures for predetermined periods after solution heat treatment on IGC susceptibility, and thereby simulate the effect of slow quenching. The data (ITT-diagram) obtained in this manner is expected to throw light on the effect of exposing extruded samples to lower temperatures for varying periods after solution heat treatment.

2. Experimental 2.1. Test material The investigated material was a flat plate profile of model alloy AlMgSi(Cu) extruded in a laboratory press as described in reference [43]. The chemical composition of the profile is given in Table 1. 2.2. Heat treatment The samples, cut into 5 · 4 · 0.27 cm plates, were solution heat treated in a salt bath for 30 min at 540
C. They were then quickly moved to another salt bath, kept Table 1 Alloy composition [wt%] Mg

Si

Fe

Mn

Cu

Cr

Ti

Zn

Al

0.57

0.62

0.21

0.21

0.13

0.001

0.01

0.007

Balance

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there for a certain time, and finally quenched in water. The temperature and hold time in the second salt bath were fixed in the range 250–450
C and 10–104 s, respectively. For comparison, some samples were water quenched and some cooled in air directly after the treatment in the first salt bath. It should be emphasised that the isothermal heat treatment procedure applied is a research test method, and it is not commonly used in practice. 2.3. Corrosion testing All sample variants were corrosion tested according to BS 11846, method B [44]. This test involved degreasing in acetone and ethanol, 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). After immersion the samples were washed in water and ethanol, and then dried. The susceptibility to IGC, including maximum corrosion depth, was evaluated by examination of transverse cross sections of corroded samples under optical microscope. Weight loss caused by corrosion was measured. 2.4. Microstructure Grain structure was investigated by optical microscope examination of cross sections prepared in the standard manner. Field emission gun scanning electron microscope (FE-SEM), operated at 5 kV, was used to investigate the grain boundary microstructure. At this accelerating voltage, the theoretical diameter of the emission volume was approximately 0.3 lm. The examined samples were prepared by grinding, metallographic polishing through 1 lm diamond paste and finally electropolishing. The electropolishing solution contained two parts by volume methanol and 1 part concentrated HNO3. The temperature of the solution was maintained between
32 and
37
C by cooling with liquid nitrogen. Electropolishing was performed at an applied potential of 12 V for 2 min. The samples were washed with ethanol directly after electropolishing and then dried. Electropolishing rendered the grain boundary phases easily visible in FE-SEM, and it was used for preparing surfaces for morphological examination of the grain boundary microstructure. The advantage gained in observing morphology on a large area was however at the expense of analytical characterisation as discussed in the earlier paper [43]. 2.5. Phase diagram An equilibrium phase diagram, showing the relative fraction of the equilibrium phases, was constructed using the FactSageTM software together with a thermodynamical database for aluminium alloys (Al-TT) from ThermoTech. Basis for the calculation was cooling of an aluminium melt of composition 0.57 wt% Mg, 0.62 wt% Si, 0.13 wt% Cu, 0.21 wt% Fe and 0.21 wt% Mn from 700 to 25
C. The melt composition used was thus nearly identical to that of the model alloy.

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3. Results 3.1. Corrosion testing The corrosion behaviour, according to the BS 11846 method B immersion test, was strongly dependent on the heat treatment route, as summarised in Table 2, and shown in Figs. 1 and 2. Samples heat treated above 400
C exhibited uniform etching (Fig. 1), characterised by shallow, dense pits tightly spread over the entire sample surface. IGC, characterised by sharp and narrow attacks along the grain boundaries, was the dominating corrosion mode for the samples treated for 10–100 s below 400
C. Longer heat treatment changed the corrosion mode to pitting. At 350 and 400
C, pitting became predominant at hold times equal to or larger than 1000 s. Longer hold times were needed for the transition at the lower temperatures. In relation to uniform etching, pitting was characterised by larger, deeper and more localised Table 2 Results from corrosion test Hold temp. [
C]

Hold time [s]

Corrosion mode

Max corr. depth [lm]

Weight loss [%]

450 450 450 450

10 100 1000 10 000

Etching Etching Etching Etching

260 280 300 235

3.7 3.7 3.2 4.5

400 400 400 400

10 100 1000 10 000

Etching Uniform IGC (200–300 lm) IGC and pitting Etching/pitting

300 570 190 160

3.5 4.7 3.6 5.1

350 350 350 350 350 350

10 100 300 600 1000 10 000

IGC (175 lm) IGC (225 lm) IGC and pitting Pitting (150 lm) Pitting (100 lm) (IGC?) Pitting (100 lm)

300 450 350 200 150 150

3.4 3.6 NA NA 4.8 5.7

300 300 300 300

10 100 1000 10 000

Uniform IGC (150 lm) A few local IGC attacks No local corrosion Pitting (100 lm)

350 50 <15 250

4.9 1.2 0.9 3.8

250 250 250 250 250 250 250

10 100 300 600 1000 10 000 100 000

IGC (200 lm) IGC (75 lm) Slight IGC IGC. Pitting? No localised corrosion No localised corrosion Pitting

380 250 70 350 50 60 150

5.1 3.2 NA NA 1.0 0.6 NA

Uniform IGC (200 lm) Etching

425 335

NA NA

Air cooled Water quenched

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Fig. 1. Cross-sectional examination of corroded specimens which were subjected to two step heat treatment.

corrosion. IGC may have also occurred, as indicated by IGC attack propagating ahead of the pit front. However, full grains corroded at a sufficiently large rate, such that pitting was the dominating corrosion mode. At temperatures below 350
C between the time zones for IGC and pitting (100– 10 000 s hold time), there was a region characterised by virtually no localised corrosion with the possible exception of edge corrosion. In comparison to the corrosion zones, in which the samples underwent 3–5% weight loss (Table 2), the samples in the corrosion resistant zone exhibited less than 1% weight loss, attributed to edge corrosion. Samples, which were water quenched directly after treatment at 540
C exhibited uniform etching (Fig. 2a), while the air cooled samples were susceptible to IGC

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Fig. 2. Cross-sectional optical micrographs after corrosion test (a,b) together with the corresponding grain structure (c,d) and grain boundary microstructure (e,f).

(Fig. 2b). The corrosion results are summarised in an isothermal time transformation diagram (ITT-diagram) shown in Fig. 3.

G. Svenningsen et al. / Corrosion Science 48 (2006) 258–272 600

Temperature [°C]

Corrosion mode No local corrosion IGC Pitting Uniform etching

ITT diagram

500

265

Uniform etching

400

Pitting

IGC

300

No localised corrosion

200

Uniform etching

100

0 1

10

100

1000 Hold time [s]

10000

100000

1000000

Fig. 3. ITT diagram for the dominant corrosion modes.

3.2. Phase diagram According to the FactSage computer simulations (Fig. 4) a-phase (AlMnFeSi), Si, b-phase and Q-phase could be expected as equilibrium phases in the model alloy. The a-phase was predicted as stable below 600
C. The b-phase, however, started to form at 500
C, and reached a maximum concentration around 330
C. The Qphase formed at about 330
C, and increased in concentration at the expense of the b-phase below 330
C. This latter phenomenon has also been calculated by other authors [41,42]. 3.3. Microstructure Short time heat treatment (10 and 100 s) below 400
C gave the same grain boundary microstructure (Figs. 5 and 6) as slow cooling in air after heat treatment in the 540
C-bath (Fig. 2f), characterised by large, discrete grain boundary precipitates, which appeared after only 10 s hold time. The size and density of the particles increased slightly during the first 100 s and remained essentially unchanged with increasing annealing time. EDS analysis of the grain boundary precipitates revealed the presence of two different particle types, AlMgSiCu and SiO. As shown in the earlier paper [43] and the calculated phase diagram (Fig. 4), these corresponded,

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500

Temperature [°C]

400

β

α

300

Q 200

Si

100

0 0

0.2

0.4

0.6

0.8

1

1.2

1.4

1.6

Phase fraction [weight %]

Fig. 4. Calculated phase diagram for the composition of the test material.

respectively, to the Q and b-phases, whose compositions were altered by selective dissolution of the more active components in the electropolishing bath. The apparent composition of the Q-phase particles measured on the electropolished samples was not affected by the length of heat treatment period at temperatures investigated below 400
C. The precipitates in the bulk of the grains (matrix) became visible after 100 s of hold time below 400
C, and their size increased during prolonged heat treatment. On samples treated for 10 000 s at 350
C rod-shaped particles of length approximately 1 lm could be visualised. EDS analysis of these precipitates under FE-SEM revealed that they were of type AlMgSiCu, i.e., the Q-phase. These precipitates were further oriented along the crystallographic planes in the matrix (Figs. 5d and 6d), which is a known characteristic of the precipitating Q-phase [35,45]. The b-phase or its precursors, if present in the bulk grains, should also be oriented along the crystallographic planes [2], and the possible presence of the b-phase was predicted by the phase diagram. However, qualitative compositional analysis of the matrix particles by EDS could not verify the presence of the b-phase. All analysed matrix precipitates contained Si and Cu, and many also Mg, indicating these to be the Q-phase. The use of transmission electron microscopy (TEM) for verifying the exact nature of the matrix particles was outside the present scope.

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Fig. 5. FE-SEM images of intermetallic phase morphology and distribution on samples heat treated at 350
C.

4. Discussion The heat treatment procedure used here (interrupted isothermal heat treatment) was a convenient method for simulating slow cooling after solution heat treatment. Although not realistic in practice, we believe that the heat treatment procedure used in this work gave relevant data in understanding the effect of low cooling rates on corrosion properties. This work provided new information about the nature of grain boundary phases causing IGC susceptibility. All sample variants, which were found to be susceptible to IGC, exhibited large and discrete Q-phase grain boundary precipitates. The susceptibility must result from microgalvanic coupling between the Cu-containing Qphase grain boundary precipitates (noble) and the adjacent depleted zone (active). It is possible that the AlMgSiCu and SiCu grain boundary phases, previously claimed to cause IGC [14–16], were actually the Q-phase possibly modified during sample preparation. The corrosion performance of the model alloy investigated was highly dependent on the cooling rate. As little as 10 s hold time below 450
C caused Q-phase grain boundary precipitation, which gave IGC susceptibility. However, hold times in

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Fig. 6. FE-SEM images of intermetallic phase morphology and distribution on samples heat treated at 300
C.

excess of 100 s resulted in pitting. Between the regions of IGC and pitting below 350
C, a region of high resistance to localised corrosion was observed, indicating that localised corrosion of this type of materials can be prevented altogether by judicious thermal treatment during fabrication. Pitting susceptibility was associated with the coarsening of the Q-phase precipitates in the grain bodies as well as the grain boundaries, giving a uniform distribution of noble cathodes on the surface in contrast to the prevalence of the grain-boundary precipitation observed on sample variants susceptible to IGC. Although not yet fully understood, the possible mechanisms which provide nearly full resistance against localised corrosion deserves further discussion. Coarsening of grain boundary precipitates during heat treatment has been reported to reduce IGC on certain alloys [20]. Because the precipitates observed were already quite coarse and discrete in the susceptible variants, coarsening is not an adequate explanation in the light of the present data. Furthermore, the presence of such coarse Q-phase particles is not a fully adequate explanation of the sharp, narrow attacks along the grain boundaries of the IGC-susceptible variants. The existence of a more narrow and continuous zone (film) of noble precipitates along the grain boundary

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would be a more likely explanation of the IGC morphology observed. Coarsening of such precipitates as a result of heat treatment would further explain resistance developed against IGC [20]. The resolution of the FE-SEM may have been limited in detecting such a grain boundary film, or such a film may have dissolved as a result of sample preparation. A more detailed investigation of the grain boundaries by field emission gun transmission electron microscope (FE-TEM) is in progress [46]. Another possible hypothesis for the corrosion resistance developed is the passivation of the Q-phase by encapsulation by the b-phase as a result of heat treatment. Passivation would result by the selective oxidation of the active Mg component in contact with the aqueous solution and formation of an insulating SiO2 coating [47,48]. However, such a film, if present, could not be detected by the present FESEM analysis. Coarsening and enhanced formation of the precipitates in the matrix by overaging has also been claimed to reduce IGC due to the consumption of the solute elements in the matrix and ensuing reduction of the compositional differences between the matrix and the solute depleted grain boundary zone [14,15,17,19,21]. Coarsening by overaging would particularly apply to the Q-phase rather than the MgSi-phase(s), as indicated by the present data. Since this would lead to depletion of the solid solution Cu in the bulk of the grains as well as along the grain boundaries, levelling of the corrosion potential gradients, which cause IGC, would be expected [14,15,17,19,21]. 5. Conclusions 1. Cooling rate after solution heat treatment at 540
C had a strong influence on the corrosion performance. Slow cooling in air rendered the model AlMgSi(Cu) alloy investigated susceptible to IGC, whereas water quenching gave resistance against IGC. 2. Hold times as short as 10 s below 400
C following the solution heat treatment caused IGC susceptibility, while pitting was introduced by longer than 100 s heat treatment. 3. It was possible to avoid localised corrosion by heat treatment at temperatures below 350
C following the solution heat treatment. Hold time at these temperatures had to be long enough to prevent IGC and short enough to avoid pitting. An ITT diagram was constructed to give a more detailed picture of the temperature time dependence of obtainable corrosion resistance. 4. IGC of such materials can be avoided by devising the correct thermal route during fabrication, while incorrect thermal route may cause unexpected IGC.

Acknowledgement Arve Johansen and Lothar Lo¨chte of Hydro Aluminium are acknowledged for their contribution in the calculation of the phase diagram (Fig. 4).

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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, DuPont Powder Coatings and GSB.

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