Intergranular Corrosion of Copper-Containing AA6x x x AlMgSi Aluminum Alloys

Transcription

1 Intergranular Corrosion of Copper-Containing AA6x x x AlMgSi Aluminum Alloys Magnus Hurlen Larsen, John Charles Walmsley, Otto Lunder, Ragnvald H. Mathiesen and Kemal Nisancioglu J. Electrochem. Soc. 2008, Volume 155, Issue 11, Pages C550-C556. doi: / alerting service Receive free alerts when new articles cite this article - sign up in the box at the top right corner of the article or click here To subscribe to Journal of The Electrochemical Society go to: ECS - The Electrochemical Society

2 C550 Journal of The Electrochemical Society, C550-C /2008/ /C550/7/$23.00 The Electrochemical Society Intergranular Corrosion of Copper-Containing AA6xxx AlMgSi Aluminum Alloys Magnus Hurlen Larsen, a, *,d John Charles Walmsley, b,c Otto Lunder, c, ** Ragnvald H. Mathiesen, b and Kemal Nisancioglu a, **,z a Department of Materials Science and Engineering and b Department of Physics, Norwegian University of Science and Technology, N-7491 Trondheim, Norway c SINTEF Materials and Chemistry. N-7465 Trondheim, Norway AlMgSi AA6xxx-series aluminum alloys are generally resistant to intergranular corrosion IGC. However, copper may introduce susceptibility to IGC; its role was investigated by using model alloys with 0.02, 0.18, and 0.7 wt % Cu. The lowest coppercontaining alloy was resistant to IGC in accelerated corrosion testing. The 0.18 wt % copper alloy showed superficial etching in the naturally aged condition and was highly susceptible to IGC in the underaged temper, but was only slightly susceptible in the peak aged or overaged condition. High-resolution field emission scanning electron microscopy imaging showed no visible grain boundary precipitation in the T4 and underaged tempers, whereas the T6 and overaged tempers had grain boundaries decorated with Cu-containing precipitates. Field emission transmission electron microscopy investigation of the underaged material showed a copper-enriched grain boundary layer and an adjacent copper-depleted zone. The reduced susceptibility to IGC upon extended artificial aging was attributed to the consumption of the copper-rich grain boundary film by the growth of grain boundary precipitates The Electrochemical Society. DOI: / All rights reserved. Manuscript submitted April 17, 2008; revised manuscript received June 18, Published September 22, Aluminum alloys in the AA6xxx-series AlMgSi Cu are generally considered to be corrosion resistant although they can be susceptible to intergranular corrosion IGC, and it has been proposed that this could be caused by Si and Cu depletion, 1-3 precipitation of elemental Si or Cu-containing phases, 2-7 or the anodic dissolution of the intermetallic phase Mg 2 Si 4,8 along grain boundaries. Shi 9 reported that model analogs of alloy AA6111, containing wt % Cu, were all susceptible to IGC, the susceptibility increasing with increasing Cu content. While investigating filiform corrosion on a painted AA6111 sheet, which was subjected to rectification by grinding and a simulated paint bake heat-treatment, Liu et al. 10 observed IGC in the deformed surface layer, formed as a result of high shear surface deformation of the alloy They found Q-phase Al 4 Mg 8 Si 7 Cu 2 precipitates along the grain boundaries of the layer, while no such precipitates were present in the bulk grain boundaries. Solute depletion was not detected along the grain boundaries. Moreover, the IGC did not affect these grain boundary precipitates, and IGC was attributed to microgalvanic coupling between the cathodic precipitates and the surrounding matrix. In a series of recent papers, Svenningsen et al showed the importance of heat-treatment on the IGC resistance of a model AA6xxx alloy extrusions containing nominally wt % 0.5 Mg, 0.6 Si, 0.2 Fe, 0.2 Mn, and about 0.02 or 0.2 wt % Cu. The alloy with lower Cu was not susceptible to IGC, while the alloy with the higher Cu content was susceptible in the air-cooled and naturally aged condition following extrusion. This was attributed to the formation of a continuous copper-rich film along the grain boundaries together with coarse Q-phase particles. However, the expected corresponding solute depletion of adjacent grain was not detected. The alloy became resistant after aging to the T6 condition, even though additional Q-phase precipitated along the grain boundaries. This was attributed to the copper-rich film by becoming discontinuous and incorporated into the precipitates. It was also observed that water-quenched material became susceptible to IGC at early stages of age hardening, before grain boundary precipitation became noticeable. Overaging caused an increased susceptibility to pitting corrosion, resulting from the coarsening of the Q-phase precipitates both along the grain boundaries and within grain interiors and the corresponding depletion of the surrounding matrix. * Electrochemical Society Student Member. ** Electrochemical Society Active Member. d Present address: Nexans Norway AS, P.O. Box 42, N-1751 Halden, Norway. z kemaln@material.ntnu.no The purpose of the present work was to provide further detailed analysis of the grain boundary microstructure of IGC-susceptible tempers without extensive grain boundary precipitation to elucidate a better understanding of the IGC mechanism. Accelerated corrosion testing was combined with scanning electron microscopy SEM and scanning transmission electron microscopy STEM. Experimental Materials. Two of the alloys prepared, containing 0.02 and 0.18 wt % Cu alloys A and B, respectively, were similar to those reported by Svenningsen et al A third alloy, with a higher Cu concentration of 0.7 wt %, denoted alloy C, was used for additional investigation of the grain boundary microstructure of the IGCsusceptible tempers by using the higher copper content to facilitate the analytical STEM study. The properties of alloy C were otherwise not of prime interest in this work, and it was therefore not investigated to the same extent as alloys A and B. The chemical compositions of the alloys are given in Table I. Alloys A and B were extruded in a laboratory press, whereas alloy C was supplied as a rolled plate. All materials were given a solution heat-treatment 30 min at 540 C in a molten salt bath, followed by water quenching to room temperature. The quench delay was less than 2 s. In addition to samples in the as-quenched and naturally aged T4 temper, samples of all alloys were aged at 185 C for 42 min underaged, 5h peak aged T6, and 24 h overaged. The underaged temper was similar to the paint bake treatment used for painted sheets and extrusions in automotive applications. 19 Corrosion testing. The accelerated corrosion test BS-ISO method B, 20 was used to rank the susceptibility of samples to IGC. The test involved degreasing, alkaline etching, and desmutting in concentrated nitric acid, followed by 24 h immersion in 30 g NaCl/l + 1 vol % HCl solution. The morphology and degree of attack were assessed by examination of metallographic cross sections. Table I. Chemical composition of the alloys used (in weight percent). The compositions were determined by use of optical emission spectroscopy. Mg Si Cu Fe Mn Cr Alloy A Alloy B Alloy C

3 Journal of The Electrochemical Society, C550-C C551 SEM and TEM characterization. SEM samples were prepared by mechanical grinding and polishing to 1 m diamond paste finish. These samples were then electropolished in cold 32 to 37 C 1 part HNO parts methanol solution for 2 min at an applied voltage of 12 V to show grain boundaries. This method was suggested and used by Svenningsen for visualization of the grain boundary precipitates. 15 It is not useful for analytical characterization because the chemical treatment alters the composition and microstructure of the intermetallic phases. Transmission electron microscopy TEM samples were prepared by grinding and polishing specimens to about 100 m in thickness, followed by electropolishing in the same solution as above, maintained at 30 C, at an applied voltage of 20 V. The foils were ion-beam milled at low voltage and angle 2.5 kv and for min prior to TEM examination to remove electropolishing artifacts. A JEOL 2010F field emission-tem operated at 200 kv was used for STEM imaging and microanalysis. Energy dispersive spectroscopy EDS X-ray analysis was performed at a nominal probe size of 0.7 or 1.0 nm using an Oxford Inca EDS analysis system equipped with automatic drift compensation. Element mapping was used to analyze enrichment and depletion as this introduces less beam damage as compared to point and line analyses and gives a good local impression of compositional variation Care was taken to align the grain boundaries parallel to the electron beam before mapping started. Diffraction patterns were used to confirm that the analyzed boundaries were of high angle type. The EDS maps were further processed statistically by geometrical filtering of the images followed by integration of the line scans along selected grain boundary segments to improve the signal-to-noise ratio. In the case of the Cu maps, a background correction was obtained by interpolating the background signals on either side of the Cu K peak and subtracting this from the Cu signal map. The standard deviation of the remaining data scatter was then calculated. This was not possible for the Mg and Si K peaks due to the nonlinear background signal in that part of the EDS spectra and their partial overlap with the intense Al K peak and L-peaks of heavier elements Fe, Cr, Mn. In this case, the maps were filtered geometrically by using a combination of nearest-neighbor two-dimensional 2D median and smoothing algorithms. The grain boundary geometry was approximated in the form of straight line segments, and the image pixels were grouped by their distance to the grain boundary line. This allowed the calculation of integrated composition profiles as a function of distance from the grain boundary line segment. Grain boundary precipitates were treated as excluded image regions, such that the resulting composition profiles would not be affected by their presence. The reported compositions are expressed as X-ray counts per pixel, and these results should be regarded as semiquantitative, showing the trends for enrichment and depletion at and adjacent to the grain boundaries. Results Corrosion test. Figure 1 shows that alloy A 0.02 wt % Cu was not susceptible to intergranular corrosion in the tempers investigated. Figure 2a shows that alloy B 0.18% Cu was corroded in the naturally aged condition T4, but the attack was not intergranular. The underaged alloy B was highly susceptible to IGC to a depth of more than 500 m over the entire surface Fig. 2b. In the peakaged T6 and overaged tempers, the material was still slightly susceptible to IGC, but with only a localized, shallow attack Fig. 2c. The results for alloys A and B confirm the similar results obtained by Svenningsen et al The high Cu alloy alloy C was highly susceptible to IGC in the underaged temper, as shown in Fig. 3. Corrosion of this alloy was not investigated in other tempers because the purpose of including alloy C in this paper was to throw indirect light to the grain boundary microstructure of alloy B, especially the grain boundary depletion of Cu, as will be shown below, which was not easily observable in the low Cu content alloy B. Figure 1. Cross-sectional micrographs of alloy A in a underaged temper and b overaged temper after accelerated corrosion test according to BS-ISO method B. SEM imaging. SEM investigation of the naturally aged and underaged alloys did not show any grain boundary precipitation, as shown in Fig. 4a for alloy B. Peak aged and overaged materials showed grain boundaries decorated with precipitates, as shown in Fig. 4b for alloy B. The particles were characterized earlier as either the Q-phase Al 4 Mg 8 Si 7 Cu 2 or a precursor phase. Grain boundary precipitation of the cathodic Q-phase clearly did not correlate with the observed susceptibility to IGC. TEM characterization. Observation of a number of grain boundaries during examination of thin foils of several replicate samples confirmed that the grain boundaries of the naturally aged T4 temper of all alloys were free from precipitates. Only scattered, discrete dispersoids of AlFeMnSi-type were detected, as shown for alloy B in Fig. 5a. The dispersoids showed dark contrast in the bright-field images as they scatter electrons more strongly than the adjacent matrix. However, small precipitates, which were too small to be visible on electropolished samples in the SEM, were imaged on the grain boundaries of the underaged temper of all alloys as shown for alloy C in Fig. 5b. Matrix precipitation of the Q -phase was still not detectable in these tempers. The grain boundaries in the overaged temper were decorated with precipitates, as shown for alloy B in Fig. 5c. Extensive matrix precipitation was also evident. Alloy C was selected in Fig. 5b instead of alloy B because the same area of alloy C in Fig. 5b was investigated again later in the paper Fig. 10 for Cu depletion of the grain boundaries. Underaged alloy B is investigated in more detail in Fig STEM imaging of the underaged tempers of alloys B and C showed precipitates along the grain boundaries, with typical size in the range nm Fig. 6. Qualitative EDS analysis showed that

4 C552 Journal of The Electrochemical Society, C550-C Figure 3. Cross-sectional micrograph of alloy C in the underaged temper after accelerated corrosion test. To examine grain boundary depletion profiles, a larger area for analysis was chosen, with a slightly bigger spot size and a higher probe current. The elemental maps and the corresponding integrated line scan are shown in Fig. 7 and 8, respectively, for the underaged alloy B showing the clear Si depletion profile although no Cu depletion profile was found. A slight indication of Mg depletion of the grain boundary zone could be inferred from Fig. 7d. EDS mapping of grain boundaries in the overaged material alloy B showed extensive precipitation of Q-phase discrete precipitates along the grain boundaries, as shown in Fig. 9. In addition, matrix precipitation of the Q -phase could be seen. 24 Figure 2. Cross-sectional micrographs of alloy B in a T4, b underaged, and c T6 tempers after accelerated corrosion test. Corrosion morphology of the overaged variant was similar to that of T6. the precipitates were a quaternary phase containing Al, Mg, Si, and Cu. These particles were recently identified as the Q -phase in an alloy with similar composition and thermal history to alloy B. 24 The most significant observation was the continuous, copper-enriched film, which was only a few nanometers thick, as shown in Fig. 6, together with EDS elemental mapping, for alloy B. This observation is consistent with earlier findings from studies of a similar alloy susceptible to IGC. 16 The analysis showed the layer contained little or no Mg or Si. Figure 4. SEM micrographs of a underaged and b overaged tempers of alloy B. The samples were electropolished to enhance the image contrast.

5 Journal of The Electrochemical Society, C550-C C553 Figure 11a shows a bright-field image and compositional profiles across a grain boundary area of alloy B in the T4 temper naturally aged after solution heat-treatment. This alloy temper combination was not susceptible to IGC, although a low level of Cu segregation along with the corresponding curve for the standard deviation of data and a weak but clear segregation of Si at the grain boundaries are indicated in the X-ray EDS line scans shown in Fig. 11b. In contrast, Si segregation was not detected in the precipitate-free grain boundary segments of the artificially aged tempers. The standard deviation of the Cu signal indicates that the grain boundary Cu enrichment is real within at least a 95% confidence level. Table II summarizes the results from corrosion testing and microstructure analysis. The table indicates that MgSi-type phases were not detected at the grain boundaries in the Cu-containing phases in the conditions listed in the table. In all these cases, the grain boundary phases, which could be characterized, were Q for the underaged 24 and Q for T6 and overaged tempers. Recall that the present specimens were quenched in water after solution heattreatment. MgSi-type phases were detected in earlier work at the grain boundaries of specimens similar to alloy B, which were air cooled after extrusion or solution heat-treatment. Figure 5. Typical S TEM images of a naturally aged T4 alloy B brightfield image, b underaged alloy C dark-field image, and c overaged alloy B bright-field image. The gray precipitates decorating the grain boundaries in b and c are the Q and Q phases, respectively. The large black spots at the grain boundary in a and at the bottom right of c are AlMnSiFe dispersoids. Due to the low level of Cu in solid solution 0.2 wt %, grain boundary depletion of this element could not be easily studied due to the sensitivity limits of the EDS technique. Alloy C 0.7 wt % Cu was therefore used to investigate Cu segregation and depletion at grain boundaries at a level suitable for the sensitivity of the STEM analysis. TEM imaging of underaged alloy C revealed Q -phase particles along the grain boundaries. A similar result was obtained for alloy B, although not shown here. Results for grain boundary precipitation in alloy B were also published earlier. 16,24 Figure 10a shows an integrated line scan analysis across the same grain boundary of alloy C as in Fig. 5b. As grain boundary precipitates were excluded from the analysis, the line scan represents only the Cu in solid solution near the grain boundary and the Cu-rich film at the grain boundary. The Cu profile shows the presence of Cudepleted zones on both sides of the grain boundary as compared to the bulk level attained approximately 50 nm away from the grain boundary. Figures 10b and c show, respectively, 2D maps of Cu and Si for an analyzed area that includes Q dispersoids. Figure 6. Color online Bright-field image of underaged alloy B left and corresponding EDS elemental mapping of marked grain boundary area. Discussion The present results support the earlier findings 16 that the Cuenriched film along the grain boundaries of the susceptible underaged temper is the main cause of the occurrence of IGC in the accelerated corrosion test. Copper segregation also causes a depletion of copper from the adjacent matrix from Cu. However, the Cu-depleted zone was detected only in the high Cu alloy alloy C. The microgalvanic couple created by the noble Cu film and the adjacent copper and silicon-depleted zone, causes the intense IGC corrosion attack observed. Given the IGC resistance of the T6 and overaged tempers, this type of IGC morphology cannot be explained by the presence of discretely spaced Q -particles alone, although both the film and the particles act as the local cathodes in the propagation of IGC. The Cu-free sample alloy A is not susceptible to IGC, although a solute Si -depleted grain boundary zone may form also in this alloy. Hence, the Si-depleted zone alone, without the continuous Cu-rich film along the grain boundary, does not provide the conditions necessary for IGC. Thermodynamic calculations 16 indicate that elemental Cu and an AlCu phase cannot form at the temperatures used for heat-treatment. Based on this, it was argued 16 that the grain boundary film had to be a precursor to the Q-phase, although Mg and Si could not be detected. The structure of this grain boundary film could not be clarified in the present study. However, the results show that the film does not contain appreciable amounts of Mg or Si, and hence it probably consists of nonequilibrium supersaturated Cu in solid solution or a nonequilibrium AlCu compound. Segregation of Mg and Si, along with Cu, to the grain boundaries is evident due to the formation of the Q -phase particles. The absence of Mg and Si in the film can be explained by the higher mobility of these elements at the grain boundaries, thus readily becoming incorporated in particles rather than remaining in the film. The assumed lower mobility of copper at the grain boundaries inhibits the incorporation of Cu in the grain boundary precipitates. This preserves the continuous grain boundary film between the grain boundary particles in the underaged condition. Although the reported near-grain boundary concentration profiles are semiquantitative, the background corrected line scan in Fig. 10a indicates that the Cu concentration approaches zero close to the boundary. The Si depletion Fig. 7 and 8 could not be corrected for the background signal, suggesting that the solute depletion could be considerably larger than it appears in the reported data. The depletion of Si also appears to extend farther into the grain than the corresponding Cu depletion, probably owing to the higher diffusivity of Si. In the peak-aged and overaged tempers, the susceptibility to IGC was reduced, but not completely eliminated. The elimination of the

6 C554 Journal of The Electrochemical Society, C550-C Figure 7. Bright-field S TEM image of the analyzed grain boundary region of a underaged alloy B with corresponding b Cu, c Si, and d Mg maps. grain boundary film as a result of aging appeared to reduce the susceptibility. However, some limited Cu-enriched film was present at some grain boundaries in the overaged material, which could account for the localized and superficial IGC observed. IGC probably occurred locally at the grain boundaries with a copper-enriched film. Extensive precipitation of Q -phase particles also occurred within the grains. This process lowered the Cu content in solid so- Figure 8. Color online Integrated line scans of Si and Cu, showing depletion of Si along the grain boundary and grain boundary enrichment of Cu. The line scans were extracted from the maps in Fig. 7 using the procedure described in the text. Figure 9. Color online Low-magnification STEM analysis of an overaged sample alloy B, showing the bright-field TEM image top along with Cu, Si, and Mg EDS maps of the area outlined in the bright-field image.

7 Journal of The Electrochemical Society, C550-C C555 Figure 11. Bright-field STEM image of the grain boundary area of a naturally aged alloy B sample, and corresponding integrated line scans of Si upper curve and Cu middle curve. The lowest curve is the standard deviation of the Cu signal. The profiles were extracted from unfiltered data. Shi 9 as Swiss cheese in appearance. The grain boundary segregation in the naturally aged T4 temper Fig. 11 is statistically reliable. However, we do not know whether the result can be generalized to all boundaries, and whether the observed segregation is continuous Figure 10. Color online a Copper distribution across a particle-free segment of a grain boundary of underaged alloy C 0.7 wt % Cu; see Fig. 5b, obtained by integration of line scan data along the grain boundary, as described earlier in the paper. The calculated local standard deviation showing the smoothed-out error in the integrated profile is also plotted. 2D maps of b Cu and c Si concentrations of the analyzed area, in which the lightest areas exhibit the highest concentrations. lution and thus equalized the potential difference between the grain boundary areas and the bulk grains, thereby reducing the risk of IGC. In contrast to the aged tempers, no significant grain body precipitation was present in the underaged temper. Hence, the grain bodies in this case are expected to exhibit the corrosion potential of the AlMgSiCu solid solution rather than that of a solute-depleted matrix. Because the corrosion potential of the solid solution matrix would be expected to be nobler than the depleted grain boundary zone, the electrochemical driving force for IGC would be further enhanced. The naturally aged material T4 was not susceptible to IGC. However, the Cu-containing alloys showed superficial etching during the corrosion test. It is believed to be caused by selective corrosion of aluminum, autocatalytically driven by the simultaneous enrichment of solid solution Cu at the surface, which acts as the cathode. This corrosion mode is probably similar to that reported by Table II. Summary of corrosion results and electron optical characterization. Corrosion type Grain boundary precipitates Grain boundary depletion Alloy A T4 NC a ND b ND Underaged NC ND ND T6 NC MgSi type ND Overaged NC MgSi type ND Alloy B T4 Superficially ND; slight Si and ND etched; no IGC Cu enrichment Underaged Extensive IGC Q dispersoids Si and slight Mg connected with continuous Cu film T6 Slight, localized Discrete Q-phase ND Overaged IGC Slight, localized IGC particles Discrete Q-phase particles Alloy C Underaged c Extensive IGC Q dispersoids connected with continuous Cu film a NC: No corrosion. b ND: Not detectable. c Investigated only in the underaged condition. ND Si and Cu

8 C556 Journal of The Electrochemical Society, C550-C or localized along the grain boundaries. Answers to these questions require a more comprehensive TEM study of the T4 temper. It should be reiterated, however, that segregation of Cu was at least an order of magnitude smaller than the grain boundary segregation found in the underaged temper. Moreover, Si segregation was in the form of enrichment rather than depletion. Si is known to form a highly passivating layer of SiO 2, 25,26 and in this form it does not participate in the corrosion process. These conditions are not sufficient to cause IGC. Conclusions 1. AlMgSiCu alloys were susceptible to IGC in the underaged temper as a result of a nearly continuous, Cu-enriched film along the grain boundaries and an adjacent zone depleted of Cu and Si. 2. The IGC susceptibility of the underaged temper was due to microgalvanic coupling between the noble Cu film and the adjacent solute-depleted active zone. 3. The reduced susceptibility observed after artificial aging to the maximum hardness T6 condition was attributed to the loss of the Cu-rich film due to discrete Q-phase precipitation. 4. Precipitation of the hardening Q -phase in the bulk grains further reduced the microgalvanic driving force causing IGC. 5. High IGC resistance and good mechanical properties can be achieved simultaneously by optimizing the artificial aging process to obtain favorable precipitation both at the grain boundary and grain interior, and a near-optimal property combination is obtained with the T6 temper. Acknowledgments Fabian Eckermann ETH Zürich provided the high Cu model alloy. This work was part of a Norwegian national research program entitled Light Metal Surface Science, supported by The Norwegian Research Council and Hydro Aluminium. Norwegian University of Science and Technology assisted in meeting the publication costs of this article. References 1. L. F. Mondolfo, Aluminium Alloys: Structure and Properties, Butterworths, London D. O. Sprowls and R. H. Brown, in Proceedings of Conference on Fundamental Aspects of Stress Corrosion Cracking, R. W. Staehle, Editor, p. 466, National Association of Corrosion Engineers, Houston, TX V. Guillaumin and G. Mankowski, Corrosion (Houston), 56, H. P. Godard, W. B. Jepson, M. R. Bothwell, and R. L. Kane, The Corrosion of Light Metals, p. 70, John Wiley & Sons, New York V. Guillaumin and G. Mankowski, Corros. Sci., 42, A. K. Bhattamishra and K. Lal, Mater. Des., 18, Aluminium-Properties and Physical Metallurgy, J. E. Hatch, Editor, Chap. 7, American Society of Metals, Metals Park, OH T. Onda, Y. Hirano, and T. Doko, in Proceedings of the 6th Automotive Corrosion and Prevention Conference, Society of Automotive Engineers, Inc., p A. Shi, Ph.D. Thesis, Pennsylvania State University, University Park, PA Y. Liu, X. Zhou, G. E. Thompson, T. Hashimoto, G. M. Scamans, and A. Afseth, Acta Mater., 55, M. Fishkis and J. C. Lin, Wear, 206, H. Leth-Olsen and K. Nisancioglu, Corros. Sci., 40, H. Leth-Olsen, A. Afseth, and K. Nisancioglu, Corros. Sci., 40, H. Leth-Olsen, J. H. Nordlien, and K. Nisancioglu, Corros. Sci., 40, G. Svenningsen, J. E. Lein, A. Bjørgum, J. H. Nordlien, Y. Yu, and K. Nisancioglu, Corros. Sci., 48, G. Svenningsen, M. H. Larsen, J. C. Walmsley, J. H. Nordlien, and K. Nisancioglu, Corros. Sci., 48, G. Svenningsen, M. H. Larsen, J. H. Nordlien, and K. Nisancioglu, Corros. Sci., 48, G. Svenningsen, M. H. Larsen, J. H. Nordlien, and K. Nisancioglu, Corros. Sci., 48, D. J. Lloyd, D. R. Evans, and A. K. Gupta, Can. Metall. Q., 39, British Standard BS 11846: D. B. Williams, A. J. Papworth, and M. Watanabe, J. Electron Microsc., 51, S V. J. Keast and D. B. Williams, J. Microsc., 199, D. T. Carpenter, M. Watanabe, K. Barmak, and D. B. Williams, Microsc. Microanal., 5, H. S. Hasting, J. C. Walmsley, A. T. J. van Helvoort, C. D. Marioara, S. J. Andersen, and R. Holmestad, Philos. Mag. Lett., 86, K. Mizuno, A. Nylund, and I. Olefjord, Corros. Sci., 43, K. Nisancioglu, J. Electrochem. Soc., 137,