Journal of The Electrochemical Society, 152 共9兲 B327-B333 共2005兲

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0013-4651/2005/152共9兲/B327/7/$7.00 © The Electrochemical Society, Inc.

Effect of Annealing Temperature on Anodic Activation of Rolled AA8006 Aluminum Alloy by Trace Element Lead Yingda Yu,a Øystein Sævik,a Jan Halvor Nordlien,b and Kemal Nisancioglua,* a

Department of Materials Technology, Norwegian University of Science and Technology, N-7491 Trondheim, Norway Hydro Aluminium Research and Development Materials Technology, N-4256 Håvik, Norway

b

Certain aluminum alloys exhibit electrochemical activation in chloride media as a result of annealing at temperatures above 350°C, attributed to the enrichment of trace element Pb at the oxide-metal interface. Activation is manifested by a significant negative shift in the corrosion potential and a significant increase in the anodic current when polarized above the corrosion potential in aqueous chloride solution. The degree of activation normally increases with increasing annealing temperature and time. This work presents electrochemical and surface analytical characterization of alloy AA8006 共nominal composition in wt % 1.5 Fe, 0.4 Mn, 0.2 Si, 0.02 Mg兲 as a function of annealing temperature in the range 350-600°C. For fixed annealing time, maximum activation and Pb enrichment were obtained at an annealing temperature of 450°C independent of the original surface condition. This was attributed to an enhancement of Pb enrichment related to the presence of Mg in the alloy. Reduced activation with further increase in the heat treatment temperature was attributed to the formation of a dense thermal oxide consisting of the crystalline MgAl2O4 spinel embedded in amorphous aluminum oxide. The results furthermore show that the nanocrystalline grain refined surface layer 共GRSL兲 on this alloy, originally formed by the hot-rolling process, oxidizes probably before becoming recrystallized during heat treatment. Stable nanocrystals remain after heat treatment at temperatures as high as 600°C. © 2005 The Electrochemical Society. 关DOI: 10.1149/1.1972728兴 All rights reserved. Manuscript submitted February 7, 2005; revised manuscript received April 7, 2005. Available electronically July 21, 2005.

Thermomechanical processing has a significant role in the determination of the surface properties of aluminum alloys. Electrochemical activation of various commercial aluminum alloys resulting from high-temperature heat treatment has been a subject of attention for the past several years because of its importance in galvanic and filiform corrosion.1-9 The phenomenon is observed by heat treatment at temperatures above 350°C. The anodic activation was characterized by deep corrosion potential transients in slightly acidified chloride solutions, starting from highly negative potentials in relation to the usual pitting or corrosion potential of about −0.75 VSCE.10 In addition, anodic polarization in neutral chloride solutions gave a high anodic current output with oxidation peaks corresponding to the potential arrests, as also confirmed by potentiostatic data.5-7,11 Recent work on commercial alloys AA8006 and AA3102 has shown the cause of activation to be the enrichment of the material surface by lead, which was present in the material as a trace element only at the parts per million 共ppm兲 level.5,9 Further work on model-binary AlPb alloys verified the role of Pb in activation.11,12 By heat treatment at 600°C, surface concentration of enriched Pb reached the order of 1 wt % independent of the bulk Pb content in the test range of 5 to 50 ppm. Mechanical polishing or caustic etching removed most of the enriched lead from the surface. However, the etched surfaces were enriched again with Pb as a result of reannealing in agreement with the work on commercial alloys.1-9 Transmission electron microscopy 共TEM兲 work11 indicated enrichment of Pb at the surface and at the grain boundaries near the surface. In addition, segregated metallic Pb particles on the order of 10 nm in size were detected at the oxide-metal interface, and many of these particles were in contact with the metal surface.12 It was suggested that the electrochemical activation was related to reduced passivity of the overlying oxide by Pb enrichment as a result of heat treatment and further segregation of metallic lead as a result of selective oxidation of the more active aluminum matrix during anodic polarization in aqueous chloride environment.12 An important difference between the commercial alloys 3005 and 8006 and the commercial alloy 3102 and the model AlPb binary alloys is that the annealing temperature needed to activate the first group was significantly lower than needed for the second group. While annealing for 1 h at 350°C was sufficient to activate the first group,8 1 h annealing at 600°C was needed for the second group for a comparable level of activation,9 judged by the current output at a

* Electrochemical Society Active Member.

selected potential. At the same time, the negative shift in the corrosion potential obtained by heat treatment of the first group at 350°C was in general larger than that obtained by heat treatment of the second group at 600°C. The present work was therefore undertaken to obtain a more detailed chemical and microstructural information on hot-rolled commercial alloy AA8006 for exploring the cause of these differences in the activation behavior. In particular, it is of interest to investigate the reason why certain commercial alloys, such as AA8006, activate more easily and at lower annealing temperatures than, e.g., binary model AlPb alloys fabricated from the pure components. It is further desirable to obtain additional evidence of Pb segregation resulting from heat treatment of a commercial alloy, specifically the actual location and state of the segregated Pb, the effect of heat treatment in general on the near-surface chemistry and microstructure of the alloy, and the effect of these factors on the electrochemical and corrosion behavior. Experimental The commercially rolled AA8006 material was supplied in the form of fully annealed O temper 0.8 mm thick sheets. The chemical composition of the material determined by spectrographic analysis is given in Table I. Samples were heat treated at different selected temperatures in an air circulation furnace for 2 h, followed by quenching in distilled water at room temperature and degreasing in acetone. For both as-received and heat-treated samples, alkaline etching was carried out in 10 wt % NaOH at 60°C for 10 s, removing approximately 1 ␮m material from the sample surface, and the samples were then desmutted for 1 min in concentrated nitric acid at 25°C. Electrochemical polarization measurements were performed at a potential sweep rate of 6 mV/min in a neutral 5 wt % NaCl electrolyte exposed to ambient air at 25 ± 1°C, and stirred by a mechanical stirrer at a fixed rate. The samples were mounted in a polyetherether-ketone 共PEEK兲 electrode holder, giving an exposed area of 1.33 cm.13 The cell geometry, including the container size, the stirrer geometry, and location were identical in all electrochemical runs to obtain identical hydrodynamic conditions in the test solution. In all cases, polarization curves were recorded at least twice in order to ensure reproducibility of the plotted results. In preparing TEM cross-sectional specimens, four strips of sample material, 4 mm wide, were first glued together with epoxy and then mechanically ground into a 3 mm diameter rod to keep the

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Table I. Chemical composition of AA8006, measured by spectrographic analysis (wt %). Mn

Mg

Fe

Si

Cu

Zn

0.4

0.02

1.5

0.16

0.02

0.02

surface regions of interest in the rod center. These rods were then cut into 0.5 mm thick slices perpendicular to the glued surfaces, mechanically ground to about 120 ␮m thickness, and thinned further by using a dimple grinder, yielding a sample thickness of about 30 ␮m. The electron transparency was finally achieved by ion milling at an angle of 3.5°. Some TEM cross-sectional specimens were generated also by ultramicrotomy, using well-documented procedures.14 Stripped surface oxide specimen was prepared by immersing an annealed sample into a mercuric chloride-methanol solution to remove the thermal oxide film from the aluminum substrate, followed by cleaning in distilled water and retrieving onto TEM sample grid. TEM observations were carried out using a Philips CM30 microscope with a point-to-point resolution of 0.23 nm. The microscope was equipped with an energy dispersive X-ray spectroscopic 共EDS兲 detector, operated at an accelerating voltage of 300 kV. Qualitative elemental depth profiling was performed by using an LECO GDS750A optical emission spectrometer 共GD-OES兲 in the direct current mode. Results The results supplement an ample amount of electrochemical, microstructural, and surface-analytical data published on alloy 8006 in earlier work. In summary, the material exhibited densely distributed intermetallic phases of type ␣-Al共Mn,Fe兲Si and Al6共Mn,Fe兲.1-4 These were of the order of 1 ␮m in the bulk and up to 2 orders of magnitude smaller very near the surface of the as-received material, which was covered by a nanocrystalline grain refined surface layer 共GRSL兲 of varying thickness, with maximum thickness at about 1 ␮m.4 The nanograin boundaries were decorated with Mg-rich oxide particles entrained during hot rolling, and the nanosized intermetallics which precipitate as a result of the thermomechanical processing, including hot rolling and subsequent annealing steps.6,7 The nature of corrosion potential and anodic polarization data obtained is already summarized in the introduction of this paper. The corrosion morphology resulting from natural exposure to slightly acidified chloride solution and by anodic polarization in neutral chloride solutions at selected potentials has also been described in detail.5 Moreover, qualitative GD-OES analysis indicated enrichment of the trace element Pb at the surface, in which the nanocrystalline layer acted as a reservoir for the enriched Pb.8 Pb was enriched only at the metal-oxide interface as a result of heat treatment of the deep caustic-etched surface, which did not contain the nanocrystalline surface. Electrochemical activation of the surface was attributed to the presence of enriched Pb. Electrochemistry.— Figure 1 shows the effect of 2 h air annealing in the range 350-600°C on the anodic behavior of as-received AA8006 in 5 wt % NaCl solution, in comparison to the behavior of an etched sample. The material was already active in the as-received condition as a result of hot rolling and annealing during fabrication, i.e., already exposed to temperatures as high as 450°C. Therefore, annealing at the lower temperature of 350°C did not give any further Pb enrichment, and the polarization curve for the sample annealed at 350°C was nearly identical to that of the as-received condition.6 Significant negative shift of the entire polarization curve as compared with the etched condition could be noted as discussed in detail in earlier papers.5,8 With increasing annealing temperature, the anodic curves shifted in the positive potential direction while

Figure 1. Effect of annealing temperature on the polarization behavior of as received alloy 8006 in 5% NaCl solution. The samples were annealed for 2 h followed by quenching in water.

remaining significantly active relative to the etched condition in all cases. However, the passive zone extended in the positive direction together with the corrosion potential. If the surface was first etched and then annealed, the anodic behavior obtained was to some extent similar to that for the asreceived samples, as shown in Fig. 2. Activation was reduced for samples annealed at temperatures lower than 450°C. Maximum activation, as judged by the general anodic current level, appeared to be attained by annealing at 450°C. Annealing at higher temperatures gave reduced activation similar to the behavior of the as-received surfaces. Increasing the annealing periods beyond 2 h did not significantly affect the shape of the polarization curves in the annealing range 350 to 600°C. It should be reiterated that these results differ from those for the binary model AlPb alloys11,12 in the fact that the activation of alloy 8006 occurred much easier at lower temperatures of annealing. This was independent of whether or not a GRSL was present. GD-OES.— The depth profiles shown in Fig. 3 and 4 are plotted in terms of arbitrary concentration units because of uncertainty about the Pb calibration in the measurements. The numbers given in the ordinate roughly correspond to weight percent. However, these

Figure 2. Effect of annealing temperatures on the polarization behavior of etched alloy 8006 in 5% NaCl solution. The samples were annealed for 2 h followed by quenching in water.

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Figure 3. GD-OES elemental depth profiles for alloy 8006 in the following conditions: 共a兲 as-received, 共b兲 annealed for 2 h at 450°C, and 共c兲 annealed for 2 h at 600°C. Heat-treated specimens were quenched in water.

results are at best semiquantitative, and they should be regarded as order of magnitude estimates rather than accurate concentration measurements. The depth calibration, based on cross-sectional TEM micrographs, is reliable. Elemental depth profiles obtained from the as-received and airannealed AA8006 alloy verified earlier data8 insofar as the enrichment of Pb together with significant Mg enrichment at the surface, as shown in Fig. 3. In addition, expected increasing enrichment of Mg with increasing annealing temperature was documented. The data suggested that increasing annealing temperature did not increase the peak Pb concentration, which appeared to have attained a saturation value, in agreement with the quantitative GD-OES data for the model AlPb alloys.11,12 The maximum, as deduced by inspection of the Al and O profiles, occurred at the metal-oxide interface. The magnitude of lead maximum was estimated from the semiquantitative analysis of the depth profiles to be of the order 1 wt %, also in agreement with the work on AlPb and alloy AA3102.9 The maximum appears to be largest for the sample annealed at 400°C. With increasing annealing temperature, the peaks became broader, possibly indicating that an increasingly thicker surface layer became enriched by Pb. The amount of Pb enriched in the near-surface region, obtained by integrating the profiles with respect to depth, shown in Table II, increased with increasing annealing temperature. Broadening of the depth profiles with increasing annealing temperature could result partly from roughening of the surface by oxide growth and recrystallization as a result of annealing. However, the information obtained about the total Pb enriched near the surface, as determined from the area under the curve, should be realistic.12 The profiles for the etched surface, shown in Fig. 4, indicated that Pb was removed from the metal surface. A small amount 共fraction of a percent兲 of Pb existed in the oxide, as judged again by inspection of the Al and O profiles shown in Fig. 4a. Annealing of the etched surface at 450°C resulted in a significant enrichment of Pb at the surface, much higher than that obtained by similar anneal-

ing of the as-received variant as shown in Fig. 4b. By inspection of the Al and O profiles, the Pb peak was observed to be located again at the metal-oxide interface. However, there was not as much Pb at a distance deeper than 0.05 ␮m from the surface, i.e., most of the metal was enriched very close to the metal-oxide interface, as indicated by the shift of the peak to the left relative to its position in the corresponding as-received specimen. The peak corresponded to about 2-3 wt % of Pb at the interface. Annealing at the higher temperature 550°C decreased the peak concentration of Pb at the metaloxide interface to the more common level of about 1 wt %. At the same time the peak became broader. However, it was still closer to the metal surface than the peak for the corresponding as-received variant. The peak integrals 共Table II兲 indicated that the amount of Pb enriched at the near-surface region increased with increasing annealing temperature also for the case of the etched variants. TEM.— TEM characterization of cross-sectional specimens verified the earlier results1-5 that a grain refined surface layer 共GRSL兲 existed on the as-received surface, as shown in Fig. 5a. The thickness of GRSL in the figure was about 0.5 ␮m. However, the layer thickness was not uniform on the basis of observations on several cross-sectional TEM foils. After annealing for 2 h at 600°C, the remnants of GRSL were still present in the form of finely dispersed Al nanograins, as indicated by arrows, in the TEM cross-sectional micrograph, Fig. 5b, which was prepared by ultramicrotomy. The rest of the GRSL was probably oxidized to form the thick oxide layer visible in the micrograph, rather than becoming recrystallized. The metallic aluminum microstructure of the nanograins was confirmed by TEM microdiffraction analysis. The diffraction pattern example inserted in Fig. 5b was indexed as Al 关321兴 orientation of the double-arrow-marked grain in the figure. Figure 6a shows the TEM cross-sectional morphology after alkaline etching. The wavy surface morphology indicated nonuniform etching, which removed the top surface oxide and most of GRSL.

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Figure 4. GD-OES elemental depth profiles for alloy 8006 in the following conditions: 共a兲 etched for 10 s in 5% NaOH and desmutted for 60 s in concentrated HNO3; 共b兲 treated as in 共a兲 and annealed for 2 h at 450°C; 共c兲 treated as in 共a兲 and annealed for 2 h at 550°C. Heat treated specimens were quenched in water.

However, at some locations as marked by arrows in Fig. 6a, remnants of the GRSL survived etching. Figure 6b shows the crosssectional micrograph of an etched sample after reannealing for 2 h at 550°C followed by water quenching. The only features observable were a large primary particle and the new thermally grown oxide characterized by dense crystals with flat bottoms grown into the metal. Significant effort to locate evidence of Pb, either in segregated form or in solid solution with the aluminum matrix, gave limited results. Specifically, a few segregated Pb-containing particles could be detected on a foil sectioned from a sample which was annealed for 2 h at 450°C, as shown in Fig. 7, with the corresponding EDS spot analysis in Fig. 8. As indicated by the high-resolution TEM 共HRTEM兲 image shown in Fig. 7b, this particle was located at the metal-oxide interface and in contact with the Al metal surface. The inserted HRTEM micrograph in Fig. 7b, enlarged from the arrowmarked area in the figure, exhibited perpendicular lattice fringes

Table II. Total Pb enriched and/or segregated at the surface of alloy 8006 as a result of annealing for 2 h, as determined by integrating the GD-OES depth profiles in Fig. 3 and 4 between 0 and 0.5 ␮m with arbitrary units, but roughly corresponding to mg/m2.

Sample

Annealing temperature 共°C兲

Pb at surface 共mg/m2兲

As received As received As received Etched Etched Etched

No annealing 450 600 No annealing 450 550

4.0 4.9 6.3 1.4 3.1 3.5

giving a lattice distance of 0.287 nm in both directions, as determined by using the bulk Al 共111兲 lattice 0.234 nm as the internal reference.12 Even though this value is nearly identical to the lattice distance 0.286 nm of Pb 共111兲,15 the particle was probably not pure Pb because the corresponding lattice fringes for Pb 共111兲 should lie at a 70.5° configuration with respect to one another. HRTEM data of the observed Pb-rich particle did not reveal further information about the composition and structure. The elements other than Pb analyzed by EDS 共Fig. 8兲, such as Mg, Mn, and Fe, may be pick-up from the surrounding matrix. Figure 9a and b shows the TEM micrographs of surface oxides stripped from sample surfaces annealed for 2 h at 500 and 600°C, respectively. TEM selected area electron diffraction 共SAED兲 pattern of the dark contrast crystalline phases in both figures, which corresponded to the flat-bottomed morphology in Fig. 6, revealed the spinel structure MgAl2O4.16 Only a tiny fraction of the oxide crystals was characterized as ␥-Al2O3 crystals. Both the spinel and ␥-Al2O3 crystals were embedded in an amorphous aluminum oxide matrix, which corresponded to the well-known oxide structure for Mg-containing aluminum alloys discussed in the literature.16 A comparison of Fig. 9a and b demonstrated the significant increasing size of the spinel crystalline phase with increasing annealing temperature from 500 to 600°C. Discussion If the segregation of the trace element Pb is a normal activationcontrolled diffusion process, increase of both the surface enrichment of Pb and anodic activation of the surface would be expected by increasing annealing temperature. However, the effect of annealing temperature on alloy 8006 appeared to be more complex. Alloy 8006 became electrochemically activated at a lower annealing temperature than the model binary alloy AlPb12 for a fixed time of annealing. Furthermore, the highest surface concentration of Pb was

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Figure 6. TEM cross-sectional micrographs of the etched surface showing 共a兲 remnants of GRSL marked by arrows and 共b兲 spinel oxide crystals 共arrows兲 formed by 2 h annealing at 550°C.

Figure 5. TEM cross-sectional micrographs of 共a兲 as-received sample and 共b兲 as-received sample annealed for 2 h at 600°C, followed with water quenching. Nanocrystals of aluminum remaining from the original GRSL are indicated by arrows on the annealed sample.

obtained by annealing at 450°C rather than at 600°C for alloy 8006, whereas the higher annealing temperature gave the highest Pb concentration at the surface of the binary model AlPb alloy. However, the total amount of Pb enriched at the near-surface region, obtained by integrating the Pb depth profiles, increased with increasing annealing temperature, as was also the case for the model binary AlPb alloy. The extent of activation, as measured by the negative shift in the corrosion potential and the anodic current output above the corrosion potential, was apparently determined by the Pb concentration at the metal surface, rather than the integrated amount of Pb present near the surface. Among the commercial alloys investigated earlier, alloys 3005,7 5754,6 and an “Alloy X”3 with the nominal composition 0.4% Fe, 0.2% Si, 0.2% Mn, 0.1% Mn exhibited behavior similar to alloy 8006, while alloys, 1050,3 1080,3 1200,3 and 31029 were similar to the binary AlPb. One difference which could give this distinction between the two alloy groups is the fact that the first group contained Mg, and the second group did not. A direct role of Mg in the activation mechanism was ruled out because it existed in a stably oxidized state at the surface. However, magnesium is known to dif-

fuse readily to the surface by heat treatment,4-9,11,12,16,17 probably along the grain boundaries. Metal oxidizes once it comes in contact with oxygen close to the alloy surface. Mg and Pb have a certain affinity to one another in terms of solid solution solubility and formation of intermetallic compounds. Furthermore, the presence of Mg in solid solution depresses the melting point of Pb.18 The anomaly observed in the microstructure determination of Pb-rich particles by the present HRTEM investigation may be attributed to the presence of Mg in the particles. The solid solution concentration of Pb in aluminum metal is very small and reported to be about 0.023 wt % at 450°C.19 Our results so far indicated a maximum concentration of about 1 wt % at the surface, which is much higher than the reported bulk solubility and an apparent supersaturation level at the surface.11,12 Even higher surface concentration observed by GD-OES depth profiling of etched, 450°C-annealed sample must be due to Pb enrichment or segregation in other forms such as metallic particles or other Pb-rich compounds. HRTEM analysis of cross-sectional foils seemed to support this hypothesis. The exact nature of the surface Pb, in particular the type of association between Pb and Mg at the surface, if any, still needs to be determined, although this is a difficult task given the small concentrations averaged over area or volume which can be quantitatively characterized by the available analytical techniques. Work is in progress by use of ternary model AlMgPb alloys prepared from the pure components. The decrease observed in the GD-OES Pb peak by annealing at 600°C relative to 450°C was accompanied by broadening of the peak. A similar phenomenon was reported also with increasing time of annealing at 600°C on AlPb model alloys.12 This type of data can be attributed to increasing roughness of the surface or propensity for nonuniform sputtering with increasing annealing temperature.12 Roughness can result from oxidation of the surface and Pb particle

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Figure 7. 共a兲 TEM cross-sectional micrograph of sample annealed for 2 h at 450°C, showing a Pb-rich particle 共arrow兲, and 共b兲 HRTEM image of the Pb-rich particle.

segregation, while nonuniform sputtering may be due to varying physical properties of crystallographic planes of larger grains obtained by increasing the annealing temperature. Large surface grains in turn result from recrystallization of the GRSL or exposure of bulk grains as a result of oxidation of the overlying GRSL. The present TEM data indicated also an increased probability of detecting segregated Pb particles by annealing at 450°C relative to annealing at

Figure 8. EDS spot analysis spectrum of the Pb-rich particle in Fig. 7.

Figure 9. TEM micrographs of the thermal oxides stripped from samples annealed for 2 h at 共a兲 500°C and 共b兲 600°C.

600°C, followed by water quenching. Pb particle segregation is known to become enhanced also by slow cooling of samples after annealing.12 These results support the theory of Pb enrichment proposed by Gundersen,11 involving rapid diffusion of Pb to the surface during heating the sample to the final annealing temperature. The fact that it was difficult to observe segregated particles on samples annealed at 600°C and water quenched indicated that Pb was in solid solution with the aluminum matrix at the surface. The reduced activation resulting from annealing at temperatures higher than 450°C is difficult to explain based on the view that the Pb in solid solution with the aluminum matrix is responsible for the activation, since the solid solution concentration should increase with increasing temperature. However, this study could not ascertain the type of Pb-containing phases formed by annealing at 450°C, which may be responsible for the activation, and their stability with increasing temperature. Furthermore, formation of a dense spinel oxide with increasing annealing temperature may contribute to the passivity of the surface.16 In addition, evaporation of the surface Pb with increasing temperature may become possible. However, all available evidence so far suggests increasing amount of Pb in the near-surface region with increasing annealing temperature. Any particulate Pb segregating at the surface at lower temperatures, e.g., around 450°C during heating may become incorporated in solid

Journal of The Electrochemical Society, 152 共9兲 B327-B333 共2005兲 solution with increasing temperature in the annealing furnace. A combination of such factors, along with the specific role attributed to Mg in surface enrichment of Pb, may lead to the reported observations. Finally, the results demonstrate the stability of the nanocrystalline layer against recrystallization at temperatures as high as 600°C. Expressed differently, the layer had a higher affinity to oxidize than recrystallize with increasing annealing temperature. The presence of remnants of the layer in the form of random metallic aluminum nanocrystals scattered over the surface even after annealing at 600°C indicated continuing stability of the crystals against growth by Zener pinning, possibly the ones most contaminated by oxide around the grain boundaries. Conclusions The data presented for the rolled commercial AA8006 aluminum alloy and data available for various other commercial and model alloys in the literature indicated that surface enrichment of the trace element Pb as a result of heat treatment requires annealing at lower temperatures 共450°C兲 on Mg-containing aluminum alloys than on Mg-free alloys 共600°C兲 for fixed annealing time 共2 h兲. The presence of magnesium in the alloy enhanced enrichment and segregation of Pb at the surface and thereby increased anodic activation of the alloy in chloride environment. Although the bulk concentration was at the ppm level, the trace element Pb became enriched at the surface of alloy 8006, probably in solid solution with the aluminum matrix, at a supersaturated concentration of 1 wt %. In addition, segregation of Pb-rich nanoparticles occurred by heat treatment at 450°C. This temperature gave optimal combination of enriched and segregated Pb at the surface insofar as the anodic activation of alloy 8006. The nanocrystalline grain refined surface layer 共GRSL兲 on alloy 8006 was stable against recystallization at annealing temperatures up to 600°C as a result of Zener pinning caused by the presence of

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nanosize oxide and intermetallic particles around the grain boundaries. Oxidation of the layer occurred before recrystallization came into effect. Acknowledgment This work was supported by the Norwegian Research Council under contract no. 158545/431. Norwegian University of Science and Technology assisted in meeting the publication costs of this article.

References 1. H. Leth-Olsen, J. H. Nordlien, and K. Nisancioglu, J. Electrochem. Soc., 144, L196 共1997兲. 2. H. Leth-Olsen and K. Nisancioglu, Corros. Sci., 40, 1179 共1998兲. 3. H. Leth-Olsen, A. Afseth, and K. Nisancioglu, Corros. Sci., 40, 1195 共1998兲. 4. H. Leth-Olsen, J. H. Nordlien, and K. Nisancioglu, Corros. Sci., 40, 2051 共1998兲. 5. Y. W. Keuong, J. H. Nordlien, and K. Nisancioglu, J. Electrochem. Soc., 148, B497 共2001兲. 6. A. Afseth, J. H. Nordlien, G. M. Scamans, and K. Nisancioglu, Corros. Sci., 43, 2359 共2002兲. 7. A. Afseth, J. H. Nordlien, G. M. Scamans, and K. Nisancioglu, Corros. Sci., 44, 145 共2002兲. 8. Y. W. Keuong, S. Ono, J. H. Nordlien, and K. Nisancioglu, J. Electrochem. Soc., 150, B547 共2003兲. 9. J. T. B. Gundersen, A. Aytac, J. H. Nordlien, and K. Nisancioglu, Corros. Sci., 46, 265 共2004兲. 10. K. Nisancioglu and H. Holtan, Corros. Sci., 18, 835 共1978兲. 11. J. T. B. Gundersen, A. Aytac, J. H. Nordlien, and K. Nisancioglu, Corros. Sci., 46, 679 共2004兲. 12. Ø. Sævik, Y. Yu, J. H. Nordlien, and K. Nisancioglu, J. Electrochem. Soc., 152, 䊏 共2005兲. 13. A. Broli and H. Holtan, Corros. Sci., 17, 59 共1977兲. 14. R. Furneaux, G. E. Thompson, and G. C. Wood, Corros. Sci., 18, 853 共1978兲. 15. Powder Diffraction File 共PDF 38–1420兲, International Center for Diffraction Data, Swarthmore, PA 共1987兲. 16. D. J. Field, G. M. Scamans, and E. P. Butler, Metall. Trans. A, 18, 463 共1987兲. 17. C. Lea and M. P. Seah, Philos. Mag., 35, 213 共1977兲. 18. J. M. Eldridge, E. Miller, and K. L. Komarek, Trans. Metall. Soc. AIME, 233, 1303 共1965兲. 19. F. A. Shunk, Constitution of Binary Alloys Second Supplement, McGraw-Hill, New York 共1963兲.