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Influence of constitutional liquation on corrosion behaviour of aluminium alloy 2017A
M A TE RI A L S C H A RAC TE RI ZA T ION 6 0 ( 2 00 9 ) 1 0 0 8–1 0 1 3
Influence of constitutional liquation on corrosion behaviour of aluminium alloy 2017A B. Kuźnicka⁎ Institute of Materials Science and Applied Mechanics, Wrocław University of Technology, 50-370 Wrocław, Poland
AR TIC LE D ATA
ABSTR ACT
Article history:
The purpose of this work was to investigate microstructural aspects of constitutional
Received 4 January 2008
liquation in the aluminium alloy 2017A and to determine its effect on corrosion behaviour of
Received in revised form
this alloy. Non-equilibrium melting of the alloy in the naturally aged condition was
12 February 2009
provoked by rapid heating above the eutectic temperature and immediate cooling in air.
Accepted 6 April 2009
Corrosion testing was performed by exposure to a marine onshore atmosphere. The microstructure examinations were carried out using light microscopy, scanning electron
Keywords:
microscopy, X-ray energy dispersion and X-ray diffraction analysis. It was found that, due to
Aluminium
rapid heating rate, coarse θ (Al2Cu) particles were melted by constitutional liquation and this
Intermetallic
way introduced strong susceptibility of 2017A alloy to intergranular corrosion.
SEM
© 2009 Elsevier Inc. All rights reserved.
Constitutional liquation Intergranular corrosion
1.
Introduction
Wrought aluminium alloys of 2000 series are heat-treatable, high-strength materials that are used as construction materials, primarily in aerospace applications. Copper, as the major alloying element in these alloys, increases strength but decreases their corrosion resistance in natural environments. One of corrosion problems for this group alloys is susceptibility to intergranular corrosion (IGC) [1–4]. Susceptibility to IGC is influenced by microstructure developed during hot forming, heat treatment or alloying. In particular, intermetallic particles play an important role in IGC, i.e. their chemical composition, size and distribution determined by processing route [5–7]. It is recognized [1,7] that IGC of aluminium alloys is a result of microgalvanic cells acting at grain boundaries, related to grain boundary precipitates which are either more noble or more active than the surrounding aluminium matrix being a solid solution. In the case of 2000 series alloys, depending on concentrations of Cu and Mg, precipitation of the Al2Cu phase
(cathodic relative to α-Al matrix) or the Al2CuMg and Al2 Mg3 phases (anodic relative to α-Al matrix [8]) may cause IGC. According to literature data [9–12] unfavourable heat treatment that may introduce IGC is both slow quenching after solution treatment, which increases the probability of grain boundary precipitation, and aging resulting in precipitate forming at the grain boundaries and in creating continuous Cu-enriched films at the boundaries. The first problem can be avoided by using higher quenching rates, if it is attainable in practice. The second problem can be solved by overaging that leads to coarsening of the particles at the grain boundaries and in the matrix. However, overaging may not always be a desirable option because it reduces mechanical strength. Moreover, coarse intermetallic particles reducing or removing IGC susceptibility become preferential sites for pitting [1,5,7,10,12]. As indicated above, cooling rate after solution heat treatment, aging temperature and time are important parameters determining the corrosion properties, as well as the mechanical properties of the alloy. However, the needs for proper heat
⁎ Tel.: +48 71 320 2767; fax: +48 71 321 1235. E-mail address: [email protected]. 1044-5803/$ – see front matter © 2009 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2009.04.004
M A TE RI A L S C H A RAC TE RI ZA T ION 6 0 ( 2 00 9 ) 1 0 0 8–1 0 1 3
Table 1 – Chemical composition of 2017A alloy [wt%]. Cu
Mg
Si
Mn
Fe
Cr
Zn
Al
4.5
0.80
0.39
0.79
0.28
0.01
0.02
Balance
treatment to obtain high corrosion resistance should consider another important factor — homogenization process of the α solid solution, requiring control of age hardening heat treatment of aluminium alloys. Homogenization substantially affects microstructure of final products of age-hardened aluminium alloys, and thus their IGC immunity. The homogenization process is influenced by two parameters: heating rate above the solvus temperature and time of isothermal holding at that temperature. The 2000 series aluminium alloys rapidly heated to a temperature above the eutectic temperature and below the solidus temperature are susceptible to constitutional liquation [13]. The phenomenon of constitutional liquation exists in many commercial alloy systems. It is defined as the subsolidus, non-equilibrium eutectic-type reaction between a second-phase particle and the surrounding matrix producing a metastable solute-rich liquid film at the particle/matrix interface. Constitutional liquation occurs when, due to rapid heating, the intermetallic particles may not dissolve at the time the eutectic temperature of the matrix-intermetallic phase system is reached. Heterogeneous nucleation of metastable liquid is followed by rapid melting of the second-phase particles and a part of the surrounding matrix. During prolonged annealing, metastable liquid droplets slowly dissolve back into the surrounding matrix but if the annealing time is short, solidification microstructure depends on the two factors: cooling rate and the nucleation of the dissolved phase. Liquid formed by constitutional liquation may spread along grain boundaries through intersection of grain boundaries with liquated regions in the matrix. The solidified liquid may introduce IGC susceptibility as well as may embrittle the grain boundaries. Due to high heating rate, the phenomenon of constitutional liquation may occur during welding operations [14,15] and laser melting. In practice, very often naturally age-hardened alloys are heated up to reduce hardness before cold plastic forming. It is thus of a great industrial as well as scientific interest to understand the mechanisms of constitutional liquation and their influence on properties of aluminium alloys. The purpose of this study was to investigate microstructural aspects of constitutional liquation in the aluminium alloy 2017A and to determine its effect on corrosion behaviour of this alloy.
2.
1009
condition after natural ageing, corresponding with T4 acc. to EN 515. The microstructural aspects of constitutional liquation were studied on specimens induction-heated to 555 °C (above the eutectic temperature of 547 °C) and immediately cooled in air. On the ground of the results presented in [13], the accepted heat treatment parameters should result in melting the Al2Cu phase particles during heating up and, upon cooling, solidification of the droplets to eutectic structure. Immediate cooling of the specimens from 555 °C, without annealing, was to prevent dissolving the droplets in the matrix. Corrosion testing was performed on the specimens both in the sensitized condition and as-delivered, by exposure to service environment — a marine onshore atmosphere. The susceptibility to IGC was evaluated by comparative metallographic examination of the specimens before and after outdoor exposure for 24 months. Characterisation of microstructure was performed by an optical microscope, a Jeol-JSM 5800 LV scanning electron microscope (SEM) equipped with EDX analysis system and a Siemens D500 X-ray diffractometer (XRD). Microscopic examination was carried out on longitudinal polished sections, unetched and etched with the solution of 0.5 ml HF (40%) in 99.5 ml of distilled water.
3.
Results
The as-delivered 2017A-T4 alloy contained coarse phase particles distributed in α-Al matrix (Fig. 1). The particles were moderately aligned with the direction of hot extrusion. A part of them was arranged on grain boundaries. Results of SEM/EDX and XRD analysis showed that they were mainly round-shaped θ-Al2Cu particles of approximate size 5 to 15 µm and irregularly shaped particles (black-coloured particles in Fig. 1), mainly of SiO2, Al2 Mg3, Cu2 Mg and of the Mn-, Si- and Fe-enriched phase of approximate size 5 to 20 µm. The grains of α solid solution were elongated in both the longitudinal and
Experiments
The samples were taken from a single commercial 4 mm thick extruded tube 65 mm (O/D) × 4 mm (WT), grade EN AW2017A-T4 acc. to EN 573. Chemical analysis of chips taken from the tube was made by gravimetric method. The results are given in Table 1. Hardness measurements were taken on the surface of the tube, by Vickers method at 29.42 N, acc. to EN ISO 6507. The obtained values of 212 ± 2 HV3 indicate the
Fig. 1 – Microstructure of 2017A alloy in the naturally aged condition. The arrows indicate coarse Al2Cu particles. Longitudinal section, etched.
1010
M A TE RI A L S C H A RAC TE RI ZA T ION 6 0 ( 2 00 9 ) 1 0 0 8–1 0 1 3
Fig. 2 – X-ray diffraction pattern of 2017A alloy in the naturally aged condition.
circumferential directions and flattened in the radial direction. After rapid heating to 555 °C followed by immediate cooling, the quantitative phase composition of the alloy, determined by XRD method, was not changed in comparison to the initial condition (Fig. 2). However, as can be seen in Fig. 3, significant changes occurred in microstructure with respect to shape, size and arrangement of the phases. Coarse particles of θ phase were transformed to a eutectic-like structure. The eutectic particles arranged inside the grains assumed spherical shape (Fig. 4a), but those located on grain boundaries were elongated, in places turning into narrow bands of divorced eutectic (Fig. 4b). Along the grain boundaries, the eutectic mixture was coarser than that distributed within the matrix. The only molten phase was the θ phase. In addition, as can be seen in Figs. 3 and 4, during heating up the dispersive hardening phases precipitated inside the solid solution grains. As a result of overaging, hardness dropped to 125 ± 3 HV3.
Fig. 4 – Eutectic particles formed during rapid heating to 555 °C followed by cooling in air: (a) in matrix; (b) along the grain boundary; (c) on partially melted θ particles (SEM).
Fig. 3 – Microstructure of 2017A alloy after rapid heating to 555 °C followed by cooling in air (without annealing): 1 — eutectic particles, 2 — divorced eutectic particles. Longitudinal section, etched.
The SEM micrographs (Fig. 4a–c) demonstrate that eutectic particles on the boundaries and inside the grains are asymmetrically surrounded by α solid solution. It is visible that the surrounding α phase is slightly porous and free of dispersive hardening particles. Moreover, some particles are either of featureless morphology (Fig. 4a), or are partially eutectic and partially without any internal microstructure (Fig. 4c). Qualitative analysis by EDX proved that the eutectic particles contain Al and Cu and some of them additionally Si, while the featureless particles contain Al, Cu, Mn, Si and Fe.
M A TE RI A L S C H A RAC TE RI ZA T ION 6 0 ( 2 00 9 ) 1 0 0 8–1 0 1 3
1011
Corrosion test revealed that the 2017A alloy in naturally aged condition (as delivered) was resistant to IGC. After 2-year exposure in marine atmosphere, no signs of corrosion were found on both the surfaces and cross-sections of the specimens. However, rapid heating and cooling of the alloy introduced a strong IGC susceptibility. As soon as after 12 months of exposure, the specimens were covered with white corrosion products, and intensive corrosion attack of grain boundaries was observed on longitudinal sections (Fig. 5). As elongated grains are beneficial to the propagation of corrosion channels, laminar exfoliation corrosion was produced. After 2-year exposure, corrosion depth from outside of the tubular specimen reached approximately 0.5 mm. The presence of eutectic particles Al–Cu and Al–Cu–Si was found on the widely attacked (dissolved) grain boundaries (Fig. 6).
4.
Discussion
The obtained results demonstrated that rapid heating of the artificially aged 2017A alloy above the eutectic temperature resulted in constitutional liquation of large θ-Al2Cu particles. Size of the θ particles ranged from 5 to15 µm. As results from computational investigation of constitutional liquation in Al–Cu alloys [16], this size precipitates can react with the matrix by eutectic-type reaction during heating at higher rate than 2 × 10− 2 and 2 × 10− 3 °C/s, respectively. The applied heating rate was
Fig. 5 – Intergranular corrosion of the heat treated 2017A alloy after marine atmosphere exposure for 2 years: (a) general view, and (b) magnified detail (SEM). Longitudinal section.
Fig. 6 – Al–Cu- and Al–Cu–Si-containing eutectic particles left in dissolved grain boundaries of corroded 2017A alloy. Longitudinal section (SEM).
higher (order of 10− 1 °C/s), so all the Al2Cu particles visible at the light microscope magnifications were subject to non-equilibrium eutectic melting. The equilibrium particles of the hardening phase became apparent in the form of rod-shaped dispersoids (see Fig. 4), characteristic for overaged condition. This is comprehensible, considering the fact that in the naturally aged condition the hardening phase particles are in the form of coherent G–P zones, which are subject to dissolving before nucleation of the equilibrium θ phase. So, it can be assumed that, during heating before reaching the complete solid-state dissolution temperature, precipitation and coagulation of θ phase take place. Under rapid heating condition, dissolution temperature increases with increasing heating rate and with increasing initial particle size [17]. This means that, in the case when in the matrix co-exist phases with particles differing by at least one size order, at a temperature above the eutectic point, incipient melting occurs at the interface of large particles, while the dispersed hardening particles are subject to dissolution in the α-matrix. The time required for the incipient melting is decidedly shorter than that required for dissolving the particles in the solid solution. It is interesting that although, as can be seen in Fig. 1, the grain boundaries were not especially privileged with regard to θ phase arrangement, after the applied heating and cooling cycle the fraction of the eutectic and divorced eutectic particles solidified along grain boundaries (Fig. 3) is clearly larger than it could be expected. This phenomenon well explains the grain boundary penetration mechanism wherein liquid film from the liquating particles infiltrates and migrates along the grain boundary regions. As the metastable intergranular liquid produced by constitutional liquation of θ phase reacts with the adjacent solid grains through back-diffusion of solute atoms across the solid–liquid interface, the solid–liquid interfacial energy is considerably low. For this reason, the liquid can effectively wet the grain boundaries and spread along intergranular regions. Since the liquid phase was formed by melting the Al2Cu phase particles with the surrounding α matrix, its chemical composition was close to the eutectic composition. Therefore, during cooling, the liquid phase solidified by epitaxial growing
1012
M A TE RI A L S C H A RAC TE RI ZA T ION 6 0 ( 2 00 9 ) 1 0 0 8–1 0 1 3
from the pre-existing α particles and, finally, solidified as lamellar eutectic or divorced eutectic (Figs. 4 and 5). A similar solidifying course of droplets in contact with the α matrix was observed in [13–15,18]. According to the authors of [13], the critical concentration depends on efficiency of the nucleation site: the α/liquid interface or the partially melted θ-Al2Cu particles. In their opinion, featureless θ particles in the matrix can be only parts of the original Al2Cu left in the liquid phase. But, as can be seen in Fig. 4a, the presence of α-phase regions with no dispersive precipitates solidified around the featureless particles and the particles with the dendritic-like internal structure, as well as around the particles with mixed structure (Fig. 4c), can speak in favour of massive solidification of Al2Cu [18]. According to this mechanism, undercooled liquid in contact with α phase should solidify as supersaturated solid solution with a gradient concentration of Cu, and in sufficient undercooling situation should yield a massive θ phase or θ phase with interdendritic segregation. In the light of the above results, grain boundary corrosion of the present material is evidently due to the presence of noble Al2Cu phase along the grain boundaries, either in normal form or as elongated divorced eutectic. Electrochemically based mechanisms for intergranular corrosion of Al-Cu and Al–Cu–Mg alloys are based on the existence of a preferential anodic path along the grain boundaries. This active region along the grain boundaries is thought to be caused by either solute depleted zones or anodic precipitates [1,3]. The IGC susceptibility of these alloys is a function of the aging time, where the resistance to IGC increases with overaging [3,9–12]. The authors of [3] proved that intergranular corrosion of Al–Cu alloys is not due to a difference in potentials between grain boundaries and grain bodies but to a difference in their breakdown potentials. As this mechanism conditions corrosion occurrence to the presence of Cu-depleted zones along grain boundaries with the breakdown potential lower than that of the grain bodies (in the underaged condition with high concentration of Cu in the matrix), it does not explain the investigated case of grain boundaries corrosion in the 2017A alloy. In the marine atmosphere (containing Cl− anions capable to breakdown passivity of the Al), the 2017A alloy in naturally aged condition (underaged condition), proved resistance to both intergranular and pitting corrosion. This means that, after hyperquenching in water and natural ageing, the alloy behaves as a monophase alloy and is cathodically protected. This confirms the conclusions resulting from the research presented in [9–12] that, as opposed to air cooling, hyperquenching in water ensures IGC resistance to Cu-rich alloys. Rapid heating of the alloy to a temperature above the eutectic point with immediate air cooling, resulting in constitutional liquation, generated the Cu concentration gradient in the grain boundary regions different than in the case of a proper heat treatment. Considering size of precipitates inside the grains and low hardness of the alloy, Cu content in the matrix was reduced to 0.1–0.2% [3,11]. The nonequilibrium heterogenic solidification of α phase ended by eutectic or massive solidification, following non-equilibrium partial melting of grain boundaries, had to result in increasing Cu concentration from 0.1–0.2% to the eutectic concentration
of 33.2% or Al2Cu content of 54% [15,17]. Qualitative EDX maps of Cu distribution along a grain boundary (Fig. 7a, b) in some fragments proved the concentration of Cu on grain boundaries and no depleted zone along the grain boundaries. This does not exclude the occurrence of a nanoscale thick Cu-rich film, whose detection requires applying FE-(S) TEM instruments [10]. This would mean low Cu content in the matrix between dispersive precipitates, increasing at the boundaries, and high concentration in the eutectic and θ particles arranged along
Fig. 7 – EDX elemental maps of grain boundary region: (a) SEM image, (b) Al distribution, and (c) Cu distribution.
M A TE RI A L S C H A RAC TE RI ZA T ION 6 0 ( 2 00 9 ) 1 0 0 8–1 0 1 3
the grain boundaries. So, non-equilibrium melting of grain boundaries of the 2017A alloy results in the conditions, in which the grain boundary corrosion must be due to differences in potentials between Cu-rich grain boundaries and Cudepleted grain bodies. This is evidenced by the wide corroded zone along grain boundaries (Fig. 6) and the progression with time dissolving of the grain matrix (Fig. 5b).
5.
Conclusions
In the microstructure of the extruded commercial alloy 2017A in naturally aged condition, a significant portion of the hardening Al2Cu phase exists in the form of coherent precipitates and the remaining part in the form of large particles arranged on grain boundaries and inside the grains. In the circumstances of rapid heating to a temperature above the eutectic point, large particles are subject to constitutional liquation. Non-equilibrium melting of particles on grain boundaries results in spreading the Cu-rich liquid along grain boundaries. At the same time overaging of the alloy occurs by precipitating of the hardening phase inside the grains and, as a consequence, depletion of the matrix in Cu. During cooling, the liquid phase solidifies like hypoeutectoid alloy by forming α phase on the interface α-matrix/liquid and ended by forming Al2Cu phase by eutectic or massive transformation. As a result, high Cu concentration in the areas adjacent to the grain boundary, compared to low concentration of the depleted grain interior, creates the condition under which a high-cathodic area is in contact with a high-anodic area. Thus, the main factor producing severe intergranular corrosion of the examined alloy in marine atmosphere is the difference in electrode potentials, rather than the difference in pitting potentials. Susceptibility of Cu-containing aluminium alloys to constitutional liquation and, as a consequence, severe reduction of their resistance to intergranular corrosion should be taken into account in the cases of applying the processes requiring heating to high temperatures, like homogenizing treatment (when heat-treating equipment, which provides high heating rates, is employed), welding or laser treatment.
REFERENCES [1] Guillaumin V, Mankowski G. Localized corrosion of 2024 T351 aluminium alloy in chloride media. Corros Sci 1999;41:421–38. [2] Liu TY, Robinson JS, McCarthy MA. The influence of hot deformation on the exfoliation corrosion behaviour of aluminium alloy 2025. J Mater Proc Technol 2004;153–154:185–92.
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[3] Galvele JR, de De Micheli SM. Mechanism of intergranular corrosion of Al–Cu alloys. Corros Sci 1970;10:795–807. [4] Robinson MJ, Jackson NC. Exfoliation corrosion of high strength Al–Cu–Mg alloys: effect of grain structure. Br Corros J 1999;34:45–9. [5] Campestrini P, van Westing EPM, van Rooijen HW, de Wit JHW. Relation between microstructural aspects of AA2024 and its corrosion behaviour investigated using AFM scanning potential technique. Corros Sci 2000;42:1853–61. [6] Blanc C, Lavelle B, Mankowski G. The role of precipitates enriched in copper in susceptibility to pitting corrosion of the 2024 aluminium alloy. Corros Sci 1997;39:495–510. [7] Blanc Ch, Freulon A, Marie-Christine Lafont MCh, Kihn Y, Georges Mankowski G. Modelling the corrosion behaviour of Al2CuMg coarse particles in copper-rich aluminium alloys. Corros Sci 2006;48:3838–51. [8] Buchheit RG. A compilation of corrosion potentials reported for intermetallic phases in aluminium alloys. J Electrochem Soc 1995;142:3994–6. [9] Svenningsen G, Larsen MH, Nordlien JH, Nisancioglu K. Effect of high temperature heat treatment on intergranular corrosion of AlMgSi(Cu) model alloy. Corros Sci 2006;48:58–272. [10] Svenningsen G, Larsen MH, Walmsley JCh, Nordlien JH, Nisancioglu K. Effect of artificial aging on intergranular corrosion of extruded AlMgSi alloy with small Cu content. Corros Sci 2006;48:1528–43. [11] Little DA, Connolly BJ, Scully JR. An electrochemical framework to explain the intergranular stress corrosion behaviour in two Al–Cu–Mg–Ag alloys as a function of aging. Corros Sci 2007;49:347–72. [12] Svenningsen G, Lein JE, Bjørgum A, Nordlien JH, Yu Y, Nisancioglu K. Effect of low copper content on intergranular corrosion of model AlMgSi alloys. Corros Sci 2006;48:226–42. [13] Reiso O, Øverlie HG, Ryum N. Dissolution and melting of secondary Al2Cu phase particles in an AlCu alloys. Metall Trans A 1990;21A:1689–95. [14] Sato YS, Park SHC, Michiuchi M, Kokawa H. Constitutional liquation during dissimilar friction stir welding of Al and Mg alloys. Scri Mater 2004:501233–6. [15] Huang C, Kou S. Partially melted zone in aluminium welds — liquation mechanism and directional solidification. Supl to Weld J May 2000:113–20. [16] Jiang C, Liu ZK. Computational investigation of constitutional liquation in Al–Cu alloys. Acta Mater 2003;51:4447–59. [17] Bjorneklett BJ, Grong O, Myhr OR, Kluken AO. Additivity and isokinetic behaviour in relation to particles dissolution. Scripta Mater 1998;46:6257–66. [18] Chattopadhyay K, Swamy VT, Agarwala SL. The solidification behaviour of undercooled Al–Cu alloys in contact with the primary phases. Acta Metall Mater 1990;38:521–31.
Influence of constitutional liquation on corrosion behaviour of aluminium alloy 2017A B. Kuźnicka⁎ Institute of Materials Science and Applied Mechanics, Wrocław University of Technology, 50-370 Wrocław, Poland
AR TIC LE D ATA
ABSTR ACT
Article history:
The purpose of this work was to investigate microstructural aspects of constitutional
Received 4 January 2008
liquation in the aluminium alloy 2017A and to determine its effect on corrosion behaviour of
Received in revised form
this alloy. Non-equilibrium melting of the alloy in the naturally aged condition was
12 February 2009
provoked by rapid heating above the eutectic temperature and immediate cooling in air.
Accepted 6 April 2009
Corrosion testing was performed by exposure to a marine onshore atmosphere. The microstructure examinations were carried out using light microscopy, scanning electron
Keywords:
microscopy, X-ray energy dispersion and X-ray diffraction analysis. It was found that, due to
Aluminium
rapid heating rate, coarse θ (Al2Cu) particles were melted by constitutional liquation and this
Intermetallic
way introduced strong susceptibility of 2017A alloy to intergranular corrosion.
SEM
© 2009 Elsevier Inc. All rights reserved.
Constitutional liquation Intergranular corrosion
1.
Introduction
Wrought aluminium alloys of 2000 series are heat-treatable, high-strength materials that are used as construction materials, primarily in aerospace applications. Copper, as the major alloying element in these alloys, increases strength but decreases their corrosion resistance in natural environments. One of corrosion problems for this group alloys is susceptibility to intergranular corrosion (IGC) [1–4]. Susceptibility to IGC is influenced by microstructure developed during hot forming, heat treatment or alloying. In particular, intermetallic particles play an important role in IGC, i.e. their chemical composition, size and distribution determined by processing route [5–7]. It is recognized [1,7] that IGC of aluminium alloys is a result of microgalvanic cells acting at grain boundaries, related to grain boundary precipitates which are either more noble or more active than the surrounding aluminium matrix being a solid solution. In the case of 2000 series alloys, depending on concentrations of Cu and Mg, precipitation of the Al2Cu phase
(cathodic relative to α-Al matrix) or the Al2CuMg and Al2 Mg3 phases (anodic relative to α-Al matrix [8]) may cause IGC. According to literature data [9–12] unfavourable heat treatment that may introduce IGC is both slow quenching after solution treatment, which increases the probability of grain boundary precipitation, and aging resulting in precipitate forming at the grain boundaries and in creating continuous Cu-enriched films at the boundaries. The first problem can be avoided by using higher quenching rates, if it is attainable in practice. The second problem can be solved by overaging that leads to coarsening of the particles at the grain boundaries and in the matrix. However, overaging may not always be a desirable option because it reduces mechanical strength. Moreover, coarse intermetallic particles reducing or removing IGC susceptibility become preferential sites for pitting [1,5,7,10,12]. As indicated above, cooling rate after solution heat treatment, aging temperature and time are important parameters determining the corrosion properties, as well as the mechanical properties of the alloy. However, the needs for proper heat
⁎ Tel.: +48 71 320 2767; fax: +48 71 321 1235. E-mail address: [email protected]. 1044-5803/$ – see front matter © 2009 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2009.04.004
M A TE RI A L S C H A RAC TE RI ZA T ION 6 0 ( 2 00 9 ) 1 0 0 8–1 0 1 3
Table 1 – Chemical composition of 2017A alloy [wt%]. Cu
Mg
Si
Mn
Fe
Cr
Zn
Al
4.5
0.80
0.39
0.79
0.28
0.01
0.02
Balance
treatment to obtain high corrosion resistance should consider another important factor — homogenization process of the α solid solution, requiring control of age hardening heat treatment of aluminium alloys. Homogenization substantially affects microstructure of final products of age-hardened aluminium alloys, and thus their IGC immunity. The homogenization process is influenced by two parameters: heating rate above the solvus temperature and time of isothermal holding at that temperature. The 2000 series aluminium alloys rapidly heated to a temperature above the eutectic temperature and below the solidus temperature are susceptible to constitutional liquation [13]. The phenomenon of constitutional liquation exists in many commercial alloy systems. It is defined as the subsolidus, non-equilibrium eutectic-type reaction between a second-phase particle and the surrounding matrix producing a metastable solute-rich liquid film at the particle/matrix interface. Constitutional liquation occurs when, due to rapid heating, the intermetallic particles may not dissolve at the time the eutectic temperature of the matrix-intermetallic phase system is reached. Heterogeneous nucleation of metastable liquid is followed by rapid melting of the second-phase particles and a part of the surrounding matrix. During prolonged annealing, metastable liquid droplets slowly dissolve back into the surrounding matrix but if the annealing time is short, solidification microstructure depends on the two factors: cooling rate and the nucleation of the dissolved phase. Liquid formed by constitutional liquation may spread along grain boundaries through intersection of grain boundaries with liquated regions in the matrix. The solidified liquid may introduce IGC susceptibility as well as may embrittle the grain boundaries. Due to high heating rate, the phenomenon of constitutional liquation may occur during welding operations [14,15] and laser melting. In practice, very often naturally age-hardened alloys are heated up to reduce hardness before cold plastic forming. It is thus of a great industrial as well as scientific interest to understand the mechanisms of constitutional liquation and their influence on properties of aluminium alloys. The purpose of this study was to investigate microstructural aspects of constitutional liquation in the aluminium alloy 2017A and to determine its effect on corrosion behaviour of this alloy.
2.
1009
condition after natural ageing, corresponding with T4 acc. to EN 515. The microstructural aspects of constitutional liquation were studied on specimens induction-heated to 555 °C (above the eutectic temperature of 547 °C) and immediately cooled in air. On the ground of the results presented in [13], the accepted heat treatment parameters should result in melting the Al2Cu phase particles during heating up and, upon cooling, solidification of the droplets to eutectic structure. Immediate cooling of the specimens from 555 °C, without annealing, was to prevent dissolving the droplets in the matrix. Corrosion testing was performed on the specimens both in the sensitized condition and as-delivered, by exposure to service environment — a marine onshore atmosphere. The susceptibility to IGC was evaluated by comparative metallographic examination of the specimens before and after outdoor exposure for 24 months. Characterisation of microstructure was performed by an optical microscope, a Jeol-JSM 5800 LV scanning electron microscope (SEM) equipped with EDX analysis system and a Siemens D500 X-ray diffractometer (XRD). Microscopic examination was carried out on longitudinal polished sections, unetched and etched with the solution of 0.5 ml HF (40%) in 99.5 ml of distilled water.
3.
Results
The as-delivered 2017A-T4 alloy contained coarse phase particles distributed in α-Al matrix (Fig. 1). The particles were moderately aligned with the direction of hot extrusion. A part of them was arranged on grain boundaries. Results of SEM/EDX and XRD analysis showed that they were mainly round-shaped θ-Al2Cu particles of approximate size 5 to 15 µm and irregularly shaped particles (black-coloured particles in Fig. 1), mainly of SiO2, Al2 Mg3, Cu2 Mg and of the Mn-, Si- and Fe-enriched phase of approximate size 5 to 20 µm. The grains of α solid solution were elongated in both the longitudinal and
Experiments
The samples were taken from a single commercial 4 mm thick extruded tube 65 mm (O/D) × 4 mm (WT), grade EN AW2017A-T4 acc. to EN 573. Chemical analysis of chips taken from the tube was made by gravimetric method. The results are given in Table 1. Hardness measurements were taken on the surface of the tube, by Vickers method at 29.42 N, acc. to EN ISO 6507. The obtained values of 212 ± 2 HV3 indicate the
Fig. 1 – Microstructure of 2017A alloy in the naturally aged condition. The arrows indicate coarse Al2Cu particles. Longitudinal section, etched.
1010
M A TE RI A L S C H A RAC TE RI ZA T ION 6 0 ( 2 00 9 ) 1 0 0 8–1 0 1 3
Fig. 2 – X-ray diffraction pattern of 2017A alloy in the naturally aged condition.
circumferential directions and flattened in the radial direction. After rapid heating to 555 °C followed by immediate cooling, the quantitative phase composition of the alloy, determined by XRD method, was not changed in comparison to the initial condition (Fig. 2). However, as can be seen in Fig. 3, significant changes occurred in microstructure with respect to shape, size and arrangement of the phases. Coarse particles of θ phase were transformed to a eutectic-like structure. The eutectic particles arranged inside the grains assumed spherical shape (Fig. 4a), but those located on grain boundaries were elongated, in places turning into narrow bands of divorced eutectic (Fig. 4b). Along the grain boundaries, the eutectic mixture was coarser than that distributed within the matrix. The only molten phase was the θ phase. In addition, as can be seen in Figs. 3 and 4, during heating up the dispersive hardening phases precipitated inside the solid solution grains. As a result of overaging, hardness dropped to 125 ± 3 HV3.
Fig. 4 – Eutectic particles formed during rapid heating to 555 °C followed by cooling in air: (a) in matrix; (b) along the grain boundary; (c) on partially melted θ particles (SEM).
Fig. 3 – Microstructure of 2017A alloy after rapid heating to 555 °C followed by cooling in air (without annealing): 1 — eutectic particles, 2 — divorced eutectic particles. Longitudinal section, etched.
The SEM micrographs (Fig. 4a–c) demonstrate that eutectic particles on the boundaries and inside the grains are asymmetrically surrounded by α solid solution. It is visible that the surrounding α phase is slightly porous and free of dispersive hardening particles. Moreover, some particles are either of featureless morphology (Fig. 4a), or are partially eutectic and partially without any internal microstructure (Fig. 4c). Qualitative analysis by EDX proved that the eutectic particles contain Al and Cu and some of them additionally Si, while the featureless particles contain Al, Cu, Mn, Si and Fe.
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Corrosion test revealed that the 2017A alloy in naturally aged condition (as delivered) was resistant to IGC. After 2-year exposure in marine atmosphere, no signs of corrosion were found on both the surfaces and cross-sections of the specimens. However, rapid heating and cooling of the alloy introduced a strong IGC susceptibility. As soon as after 12 months of exposure, the specimens were covered with white corrosion products, and intensive corrosion attack of grain boundaries was observed on longitudinal sections (Fig. 5). As elongated grains are beneficial to the propagation of corrosion channels, laminar exfoliation corrosion was produced. After 2-year exposure, corrosion depth from outside of the tubular specimen reached approximately 0.5 mm. The presence of eutectic particles Al–Cu and Al–Cu–Si was found on the widely attacked (dissolved) grain boundaries (Fig. 6).
4.
Discussion
The obtained results demonstrated that rapid heating of the artificially aged 2017A alloy above the eutectic temperature resulted in constitutional liquation of large θ-Al2Cu particles. Size of the θ particles ranged from 5 to15 µm. As results from computational investigation of constitutional liquation in Al–Cu alloys [16], this size precipitates can react with the matrix by eutectic-type reaction during heating at higher rate than 2 × 10− 2 and 2 × 10− 3 °C/s, respectively. The applied heating rate was
Fig. 5 – Intergranular corrosion of the heat treated 2017A alloy after marine atmosphere exposure for 2 years: (a) general view, and (b) magnified detail (SEM). Longitudinal section.
Fig. 6 – Al–Cu- and Al–Cu–Si-containing eutectic particles left in dissolved grain boundaries of corroded 2017A alloy. Longitudinal section (SEM).
higher (order of 10− 1 °C/s), so all the Al2Cu particles visible at the light microscope magnifications were subject to non-equilibrium eutectic melting. The equilibrium particles of the hardening phase became apparent in the form of rod-shaped dispersoids (see Fig. 4), characteristic for overaged condition. This is comprehensible, considering the fact that in the naturally aged condition the hardening phase particles are in the form of coherent G–P zones, which are subject to dissolving before nucleation of the equilibrium θ phase. So, it can be assumed that, during heating before reaching the complete solid-state dissolution temperature, precipitation and coagulation of θ phase take place. Under rapid heating condition, dissolution temperature increases with increasing heating rate and with increasing initial particle size [17]. This means that, in the case when in the matrix co-exist phases with particles differing by at least one size order, at a temperature above the eutectic point, incipient melting occurs at the interface of large particles, while the dispersed hardening particles are subject to dissolution in the α-matrix. The time required for the incipient melting is decidedly shorter than that required for dissolving the particles in the solid solution. It is interesting that although, as can be seen in Fig. 1, the grain boundaries were not especially privileged with regard to θ phase arrangement, after the applied heating and cooling cycle the fraction of the eutectic and divorced eutectic particles solidified along grain boundaries (Fig. 3) is clearly larger than it could be expected. This phenomenon well explains the grain boundary penetration mechanism wherein liquid film from the liquating particles infiltrates and migrates along the grain boundary regions. As the metastable intergranular liquid produced by constitutional liquation of θ phase reacts with the adjacent solid grains through back-diffusion of solute atoms across the solid–liquid interface, the solid–liquid interfacial energy is considerably low. For this reason, the liquid can effectively wet the grain boundaries and spread along intergranular regions. Since the liquid phase was formed by melting the Al2Cu phase particles with the surrounding α matrix, its chemical composition was close to the eutectic composition. Therefore, during cooling, the liquid phase solidified by epitaxial growing
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from the pre-existing α particles and, finally, solidified as lamellar eutectic or divorced eutectic (Figs. 4 and 5). A similar solidifying course of droplets in contact with the α matrix was observed in [13–15,18]. According to the authors of [13], the critical concentration depends on efficiency of the nucleation site: the α/liquid interface or the partially melted θ-Al2Cu particles. In their opinion, featureless θ particles in the matrix can be only parts of the original Al2Cu left in the liquid phase. But, as can be seen in Fig. 4a, the presence of α-phase regions with no dispersive precipitates solidified around the featureless particles and the particles with the dendritic-like internal structure, as well as around the particles with mixed structure (Fig. 4c), can speak in favour of massive solidification of Al2Cu [18]. According to this mechanism, undercooled liquid in contact with α phase should solidify as supersaturated solid solution with a gradient concentration of Cu, and in sufficient undercooling situation should yield a massive θ phase or θ phase with interdendritic segregation. In the light of the above results, grain boundary corrosion of the present material is evidently due to the presence of noble Al2Cu phase along the grain boundaries, either in normal form or as elongated divorced eutectic. Electrochemically based mechanisms for intergranular corrosion of Al-Cu and Al–Cu–Mg alloys are based on the existence of a preferential anodic path along the grain boundaries. This active region along the grain boundaries is thought to be caused by either solute depleted zones or anodic precipitates [1,3]. The IGC susceptibility of these alloys is a function of the aging time, where the resistance to IGC increases with overaging [3,9–12]. The authors of [3] proved that intergranular corrosion of Al–Cu alloys is not due to a difference in potentials between grain boundaries and grain bodies but to a difference in their breakdown potentials. As this mechanism conditions corrosion occurrence to the presence of Cu-depleted zones along grain boundaries with the breakdown potential lower than that of the grain bodies (in the underaged condition with high concentration of Cu in the matrix), it does not explain the investigated case of grain boundaries corrosion in the 2017A alloy. In the marine atmosphere (containing Cl− anions capable to breakdown passivity of the Al), the 2017A alloy in naturally aged condition (underaged condition), proved resistance to both intergranular and pitting corrosion. This means that, after hyperquenching in water and natural ageing, the alloy behaves as a monophase alloy and is cathodically protected. This confirms the conclusions resulting from the research presented in [9–12] that, as opposed to air cooling, hyperquenching in water ensures IGC resistance to Cu-rich alloys. Rapid heating of the alloy to a temperature above the eutectic point with immediate air cooling, resulting in constitutional liquation, generated the Cu concentration gradient in the grain boundary regions different than in the case of a proper heat treatment. Considering size of precipitates inside the grains and low hardness of the alloy, Cu content in the matrix was reduced to 0.1–0.2% [3,11]. The nonequilibrium heterogenic solidification of α phase ended by eutectic or massive solidification, following non-equilibrium partial melting of grain boundaries, had to result in increasing Cu concentration from 0.1–0.2% to the eutectic concentration
of 33.2% or Al2Cu content of 54% [15,17]. Qualitative EDX maps of Cu distribution along a grain boundary (Fig. 7a, b) in some fragments proved the concentration of Cu on grain boundaries and no depleted zone along the grain boundaries. This does not exclude the occurrence of a nanoscale thick Cu-rich film, whose detection requires applying FE-(S) TEM instruments [10]. This would mean low Cu content in the matrix between dispersive precipitates, increasing at the boundaries, and high concentration in the eutectic and θ particles arranged along
Fig. 7 – EDX elemental maps of grain boundary region: (a) SEM image, (b) Al distribution, and (c) Cu distribution.
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the grain boundaries. So, non-equilibrium melting of grain boundaries of the 2017A alloy results in the conditions, in which the grain boundary corrosion must be due to differences in potentials between Cu-rich grain boundaries and Cudepleted grain bodies. This is evidenced by the wide corroded zone along grain boundaries (Fig. 6) and the progression with time dissolving of the grain matrix (Fig. 5b).
5.
Conclusions
In the microstructure of the extruded commercial alloy 2017A in naturally aged condition, a significant portion of the hardening Al2Cu phase exists in the form of coherent precipitates and the remaining part in the form of large particles arranged on grain boundaries and inside the grains. In the circumstances of rapid heating to a temperature above the eutectic point, large particles are subject to constitutional liquation. Non-equilibrium melting of particles on grain boundaries results in spreading the Cu-rich liquid along grain boundaries. At the same time overaging of the alloy occurs by precipitating of the hardening phase inside the grains and, as a consequence, depletion of the matrix in Cu. During cooling, the liquid phase solidifies like hypoeutectoid alloy by forming α phase on the interface α-matrix/liquid and ended by forming Al2Cu phase by eutectic or massive transformation. As a result, high Cu concentration in the areas adjacent to the grain boundary, compared to low concentration of the depleted grain interior, creates the condition under which a high-cathodic area is in contact with a high-anodic area. Thus, the main factor producing severe intergranular corrosion of the examined alloy in marine atmosphere is the difference in electrode potentials, rather than the difference in pitting potentials. Susceptibility of Cu-containing aluminium alloys to constitutional liquation and, as a consequence, severe reduction of their resistance to intergranular corrosion should be taken into account in the cases of applying the processes requiring heating to high temperatures, like homogenizing treatment (when heat-treating equipment, which provides high heating rates, is employed), welding or laser treatment.
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