Localized Corrosion of Carbon Steel Weldments

CORROSION ENGINEERING SECTION

Localized Corrosion of Carbon Steel Weldments M. Sephton and P.C. Pistorius*

ABSTRACT Dominant mechanisms in the reduced localized corrosion resistance of carbon steel weldments were investigated with the aid of experimental sulfur-rich (0.1 wt% S) and low-sulfur (0.01 wt% S) carbon steels, exposed to 3.5% sodium chloride (NaCl). Experimental techniques included mapping current density in the solution above corroding samples (by using a scanning vibrating electrode), zero-resistance ammetry, exposure of autogenously welded samples to flowing solution, and examination of microstructure by optical and scanning electron microscopy. It was found that galvanic corrosion between different microstructures in the weldment is probably not the main cause for the localized attack, and it is proposed that, for the conditions investigated, grooving is caused primarily by the unique combination of active sulfides and networks of small sulfides on original austenite grain boundaries in the fusion line area. KEY WORDS: autogenous welding, carbon steel weldments, grooving, localized corrosion, microstructures, scanning vibrating electrode technique, sulfides, weld thermal cycle

INTRODUCTION Many failures of welded carbon steel products have been attributed to the selective corrosion of the weld. This corrosion is found in fusion and electrical resistance-welded products, occurs in varied environments, and is not confined to one region of the weldment.1-6 This type of corrosion is characterized by grooves parallel to the weld. Submitted for publication May 1999; in revised form, June 2000. * Department of Materials Science and Metallurgical Engineering, University of Pretoria, Pretoria 0002, South Africa.

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Extensive research on the corrosion of weldments has been performed to formulate corrosion prevention strategies. However, the mechanisms involved in the localized corrosion of the weldments are not fully clear. During the weld thermal cycle, a relatively narrow region of the material is subjected to a wide range of peak temperatures and cooling rates, resulting in a variety of microstructures and changes in properties of the affected area. This microstructural variety complicates investigation into possible corrosion mechanisms. Dominant mechanisms influencing the corrosion behavior of the weld appear to be the formation of microstructures susceptible to corrosion or the redistribution of sulfides during the weld thermal cycle.7-8 The present work was an attempt to elucidate the relative importance of these corrosion mechanisms. The scope of this investigation was limited to carbon steels with high manganese content and two variations in sulfur content, which were autogenously welded.

MATERIAL A cast of carbon steel, high in sulfur, was prepared. The high sulfur content was decided upon to accelerate and intensify any effect of sulfur or sulfides. A stoichiometric amount of manganese was added to ensure the formation of essentially manganese-rich sulfides, as manganese sulfide (MnS) is considered to have a detrimental effect on corrosion

0010-9312/00/000255/$5.00+$0.50/0 © 2000, NACE International

CORROSION–DECEMBER 2000


TABLE 1 Chemical Composition of Material Investigated (wt%) Fe

C

Mn

S

Al

Cu

Cr

Ni

P

High sulfur

Bal.

0.165

2.34

0.101

≈ 0.05

≈ 0.006

≈ 0.02

≈ 0.019

≈ 0.003

Low sulfur

Bal.

0.11

2.10

0.009

0.032

≈ 0.006

≈ 0.02

≈ 0.019

0.003

TABLE 2 Heat Treatment for Different Microstructures and Resulting Hardness Microstructure

Heat Treatment

Vickers Hardness (kg/mm2)

Ferrite/pearlite

Austenitizing at 930°C for 15 min, annealing

188.8 ± 3.4

Bainite

Austenitizing at 930°C for 15 min, quenching in molten lead at 475°C for 30 min, followed by quenching in water

229.4 ± 4.6

Fresh martensite

Austenitizing at 930°C for 15 min, quenching in brine

465.9 ± 8.4

Tempered martensite

Austenitizing at 930°C for 15 min, quenching in brine, followed by tempering at 400°C for 1 h

378.0 ± 8.3

resistance. A second cast, with a sulfur content closer to actual casts used in industry and an order of magnitude lower than that of the high-sulfur steel, also was prepared. The manganese content was kept similar to that of the high-sulfur steel to change only one parameter. (A change in manganese content would have changed the hardenability of the steel, and hence the microstructures in the weldment.) The chemical composition of the materials is listed in Table 1. Both casts were made in a vacuum induction furnace with aluminum as the deoxidizer. Following casting, the ingots were soaked at 1,200°C for 2 h and subsequently rolled down to a 10-mm (0.4-in.) thickness with a finishing temperature of 900°C, followed by air cooling. A normalizing heat treatment was given, soaking the ingots at 900°C for 30 min and cooling in air. In addition, a further normalizing heat treatment (1 h at 900°C, followed by cooling in air) was given to specimens cut from the rolled sheet before welding, to ensure a homogeneous microstructure. The specimens were 150 mm by 10 mm by 7 mm (6.0 in. by 0.4 in. by 0.3 in.) and were placed flat on the largest side during air cooling. After this treatment, the highand low-sulfur steel consisted of fine ferrite grains and pearlite colonies and had a Vickers hardness of 198 kg/mm2 and 190 kg/mm2, respectively.

EXPERIMENTAL PROCEDURES AND RESULTS

Microstructures Susceptible to Corrosion It has been suggested that certain microstructures formed during the weld thermal cycle, specifi-

CORROSION–Vol. 56, No. 12

cally bainite, are less corrosion resistant than others.7 It has also been proposed that a galvanic couple develops between the different microstructures in the weldment, promoting the selective corrosion of certain parts. The types of microstructures commonly found in welds of low-carbon steels are martensite, bainite, ferrite, and pearlite. Some autotempering of the martensite may occur. Products such as residual austenite also may be present, but are mainly found in higher alloyed steels. This investigation concentrated on ferrite/pearlite, bainite, fresh martensite, and tempered martensite. These microstructures were prepared by heattreating the sulfur-rich material (0.1% S, 2.3% Mn, 0.14% C). The specimens were austenitized at 930°C for 15 min, and quenched either in brine or in a lead bath (T = 475°C), or annealed. Quenching in brine led to the formation of fresh martensite. To produce the tempered martensite sample, the fresh martensite was heat-treated at 400°C for 1 h. A summary of the heat treatments, as well as the resulting hardness of the various microstructures, are given in Table 2. Scanning electron microscope (SEM) images of the microstructures resulting from the various cooling rates are presented in Figure 1. Some sulfide inclusions could be distinguished and are indicated by arrowheads. To obtain an indication of the corrosion behavior of the prepared microstructures, polarization diagrams were constructed for each separate microstructure by potentiodynamic scanning in 3.5%

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(a)

(b)

(c) FIGURE 1. SEM of microstructures investigated in galvanic couples: (a) ferrite/pearlite, (b) bainite, and (c) fresh martensite. (Some of the sulfide inclusions are indicated by arrowheads).

FIGURE 2. Polarization behavior of different microstructures in 3.5% NaCl, indicating significant differences in corrosion potential (vs AgAgCI reference electrode). The curves were mathematically corrected for IR drop. B = bainite, F/P = ferrite/pearlite, FM = fresh martensite, and TM = tempered martensite.

were allowed to corrode for 1 h under open-circuit conditions before the scans. Potential scanning was performed in the positive direction from the corrosion potential at a rate of 2 mV/s. The solution resistance (RΩ) of the set-up was determined with the aid of impedance measurements and was ≈ 12.8 Ω. Measured potentials were mathematically corrected for the IR drop after the scans. Results are presented in Figure 2. Differences in corrosion potential existed, indicating possible galvanic effects between the different microstructures. However, these polarization tests are short term, nearly instantaneous, and prolonged corrosion will likely change the corrosion behavior of the microstructures. With this in mind, longer-term polarization resistance tests also were performed. In these tests, separate microstructures were allowed to corrode under open-circuit conditions for an extended period, while oxygen was bubbled through the cell, saturating the solution with dissolved oxygen, and, at the same time, stirring it. During the exposure period, polarization resistance (Rp) measurements were periodically performed. This is a useful method for rapidly measuring relative corrosion rates or changes in corrosion.9 Rp was determined by calculating the gradient of the potential/current density scans where i = 0 (i.e., at the corrosion potential):10 1 / R p = δi / δE |i = 0

(1)

and this can be converted to the corrosion rate (icorr) through the following relationship:10 sodium chloride (NaCl) at room temperature and in a corrosion cell open to the atmosphere. Samples were ground to a 1,200-grit finish and an area of 1.23 cm2 (0.19 in.2) was exposed to the electrolyte. Specimens

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R p = B / i corr

(2)

where B = 0.02 V.



Calculated corrosion rates are given in Figure 3. From these corrosion rates, it seems that there is very little difference in the corrosion rates of the different microstructures after longer immersion periods. This observation was supported by the measured corrosion potential (of all the microstructures as a function of time), which tended to have a value of ≈ 0.63 V vs silver-silver chloride (Ag-AgCl) in all instances. These results indicate that little or no galvanic interaction between the different microstructures is likely in the long term. This conclusion was tested by constructing special galvanic couples. A schematic illustration of a galvanic couple between two microstructures is shown in Figure 4. Experiments were performed in an aerated 3.5% NaCl solution at room temperature and lasted 22 h. Couples were allowed to corrode freely while the current passing between them was measured by a zero resistance ammeter (ZRA) in the circuit. Oxygen was introduced into the cell via a diffuser situated in the lower quarter of the corrosion cell. The oxygen bubbles had the added effect of stirring the solution, thereby preventing stagnant conditions. The average galvanic currents are given in Figure 5. Galvanic currents were taken as the average of the last 2 h of the tests. This period was decided upon since it fell well within the range of stabilized current flow and corresponded to the time when the uncoupled corrosion rate stabilized. On the graph, the different points for each microstructure combination refer to repeat experiments. The sign convention is such that a net cathodic current on the microstructure that is listed first gives a positive galvanic current. There appeared to be little difference in the galvanic behavior of the different microstructures, and there was no consistent tendency of any microstructure, in any given couple, toward predominantly anodic or cathodic behavior.

FIGURE 3. Corrosion rate of different microstructures as a function of time, indicating little difference in behavior after an extended period of immersion. ◆: fresh martensite; 䡺: bainite; : ferrite/pearlite; 䢇: tempered martensite.

FIGURE 4. Sample used for galvanic couple. The exposed surface was always parallel to the rolling direction. Surfaces were wet-ground to a 1,200-grit finish immediately before exposure.

Effect of Welding To establish the effect of welding on the microstructure and corrosion behavior of the weld area, welded samples of the high- and low-sulfur steels underwent tests involving long-term exposure to flowing 3.5% NaCl solution. Autogenous tungsten inert gas (TIG) welding was used as the welding procedure. This has the major advantage of producing a thermal history comparable to that of fusion welding with a filler material, while not introducing the added complexity of alloying elements. Welding resulted in a weld metal and heataffected zone (HAZ) consisting of martensite in both steels, as indicated by the Vickers microhardness graph for the high-sulfur steel presented in Figure 6. Welds were exposed to flowing 3.5% NaCl (at 30°C) for 5 weeks in a flow channel open to the


FIGURE 5. Average galvanic currents (with standard deviations) for different combinations of microstructures in aerated and agitated 3.5% NaCl (an area of ≈ 0.9 cm2 of both microstructures in the couples was exposed to the solution), showing an almost random distribution.

atmosphere. After exposure, grooves formed in the fusion line area of the high- and low-sulfur steels. The grooves seemed to consist of a multitude of pits of various depths. In fact, it seemed that the grooves formed by growth of pits (areas of discrete localized attack) that eventually coalesced, with the deepest part of the groove on or near the fusion line. No pits could be observed in the HAZ.

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FIGURE 6. Typical microhardness profile over a weldment of the high sulfur steel, indicating the high hardness of the weld metal and HAZ. BM = base metal and WM = weld metal.

sulfur steels) were followed by using a scanning vibrating electrode. The scanning vibrating probe is a sensitive tool for directly measuring the potential differences resulting from current variations as corrosion occurs. It converts the potential gradient into an alternating current voltage by mechanical motion of the sensing electrode. Amplification of the alternating current potentials is far more effective than the amplification of direct current potentials, resulting in a more sensitive technique.11 With Ohm’s law, the measured potential can be converted to the component of current density at the point in the direction the electrode vibrates (z-direction):

i Z = K ( ∆V / ∆z )

(a)

(b) FIGURE 7. (a) Typical result of the vibrating electrode scans over the high-sulfur steel weldment, and (b) profilometer result over the same weldment after 5 weeks’ exposure to flowing 3.5% NaCl, comparing the two sets of results. The broken lines in (b) indicate the transition zones between base metal and HAZ, and HAZ and weld metal, respecitvely.

A clearer picture of the extent of overall corrosion damage was obtained by following the contours of the surface with a sensitive profilometer. Profiles were measured at set intervals (≈ 0.5 cm) perpendicular to the weld to cover a wide area. An example of such a profile over the high-sulfur steel weldment is given in Figure 7(b). The profiles show that the weld metal corroded the most, followed by the HAZ, and then the base metal. The grooves that formed between the weld metal and HAZ also could be distinguished. The initiation and propagation of localized corrosion on the weldment (of the high- and low-

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(3)

where K is the conductivity of the solution, ∆V is the measured potential amplitude, and ∆z is the amplitude of vibration of the electrode. For these experiments, ∆z was 30 µm, the electrode tip diameter was ≈ 30 µm, and the frequency of vibration was 77.6 Hz. The welds were ground-flush and wet-ground to a 1,200-grit finish before exposure to the electrolyte. Measurements were performed in flowing 3.5% NaCl, open to the atmosphere, and at room temperature. Tests lasted for 211 h and 121 h for the high- and low-sulfur steel, respectively. During the tests, the corrosion products were removed every 24 h by rinsing the surface with fresh test solution. A typical result (after 117.5 h exposure to the solution and 30 min after the last rinsing) is shown in Figure 7(a). Each line represents a gradient profile over the weldment at a different position. Profiles were progressively displaced in the y-direction to distinguish among the different positions. The distance between the profiles therefore has no fundamental significance. However, the relative potentials on a single profile indicate the positions of relative anodic and cathodic areas. Figure 7(b) compares the corrosion profile (profilometer results) for the high-sulfur steel with the typical result in Figure 7(a), showing the similarity between the two sets of results. Figure 8(a) shows the progression of corrosion at one position on the high-sulfur steel sample, and Figure 8(b) shows this on the low-sulfur sample. In these graphs, the profiles are again progressively displaced in the y-direction, but in this instance to distinguish among different times of measurement. Negative peaks in the profiles indicate areas of anodic activity. The position of the fusion line can at best be given as a range of positions because of the varying width of the weld metal. In addition, the positioning of the sample and vibrating electrode had to be performed manually, limiting the position measurement accuracy.



Results of the scans over weldments show that, although the activity varied from time to time, the most pronounced attack seemed to be concentrated in the fusion line area for the high-sulfur steel and in the HAZ of the low-sulfur steel. SEM investigations of the fusion line area revealed that in the weld metal and adjacent hightemperature HAZ, sulfide inclusions were closely associated with the original austenite grain boundaries (Figure 9), at some places almost forming continuous networks. The sulfides in the low temperature HAZ did not seem to be associated with a specific region. Corrosion has been noted to initiate at, or in the vicinity of, sulfide inclusions.5,7-8,12-14 A possible reason for corrosion initiation at sulfides is that “active sulfides” exist with a reactive region around them. This region has been found to initiate corrosion.7,12 Active sulfides probably are formed when existing sulfides dissolve and re-precipitate during the weld thermal cycle. It is thought that, since the cooling rate is high, not all the sulfur transforms back to MnS; and a sulfur-rich layer is formed around the inclusions. The presence of such active sulfides can be detected with the aid of the so-called micro-corrosion test (MCT).15 This test is an etch method for carbon steel and involves immersion of a finely polished sample in 3% NaCl for 30 s. Active sulfides can be distinguished by optical microscopy (at 100X magnification) by the presence of a colored halo. Graphs showing the area density of active sulfides in various regions of the weldment are shown in Figures 10(a) and (b). In the high- and low-sulfur steel, such active sulfides could be detected, and the highest concentrations generally correspond with areas of increased corrosive attack (as shown by the scanning vibrating electrode results in Figure 8). Scanning vibrating electrode results also indicated differences in corrosion in the HAZ of both samples. The relative anodic and cathodic behavior of the high- and low-temperature regions in the HAZ was tested further by measuring the average galvanic currents flowing between simulated high- and lowtemperature HAZ microstructures (formed by austenitizing the high-sulfur steel in a weld simulator at 1,485°C and 900°C, respectively) during exposure to aerated 3.5% NaCl for 22 h at room temperature. Tests were performed in the same manner as described for the different microstructures in the previous section (Figure 5). Results are presented in Figure 11. Average galvanic currents indicate that the hightemperature HAZ is anodic with respect to the lowtemperature one, which might be attributable to the presence of active sulfides and the change in sulfide morphology, resulting from the weld thermal cycle.


(a)

(b) FIGURE 8. Progression of corrosion over: (a) the high-sulfur steel weldment at one position (results after [——] 2 h; [- - -] 72.5 h; [——] 165 h; [——] 211 h of exposure are shown) and (b) the lowsulfur steel sample (results after [——] 4 h; [- - -] 23 h; [——] 50 h; [—— ] 121 h of exposure are shown) profiles are progressively displaced (2 × 10–5 units per profile) in the y-direction to distinguish between different scan times. Shortest times are indicated by the bottom profile and longest times by the top profile.

FIGURE 9. SEM micrograph of high-temperature HAZ, showing inclusions (indicated by arrowheads) on grain boundaries (etched in 2% nital).

Both samples had martensitic microstructures, confirming the suggestion that the grooving activity is not controlled by the microstructure of the steel matrix itself, but rather by sulfur redistribution.

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FIGURE 10. Profiles showing the area density of active sulfides in various regions of the weldment of high-sulfur and low-sulfur steels. The band indicates the minimum and maximum values.

FIGURE 11. Average galvanic currents showing high-temperature HAZ (Ht) to be anodic to low-temperature HAZ (Lt). The high-sulfur material (0.1% S) was used to simulate the HAZ microstructures, and an area of 1 cm2 of both microstructures in the couple was exposed to the test solution.

DISCUSSION Results of the corrosion tests on the microstructures that possibly could exist in the weldment suggest that, initially, there are some differences in corrosion behavior of the different microstructures. For instance, the initial corrosion potentials (Figure 2) indicated that there can exist a potential difference of ~ 30 mV between the microstructures. This difference has been found by some investigators to be large enough for a galvanic cell to develop.1 Measured galvanic currents (0.11 mA/cm2 to 0.22 mA/cm2) were similar to the corrosion currents of the uncoupled microstructures (≈ 0.11 mA/cm2).

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Hence, it does not appear that the proximity of two microstructures leads to greatly increased corrosion. The possibility exists that furnace preparation of HAZ microstructures might not exactly simulate actual welded HAZ microstructures. However, this investigation sought to establish whether one or more of the low-temperature transformation products, found by some investigators to be susceptible to corrosion,2 behave in a predominantly anodic fashion and whether the close proximity of these structures leads to the development of micro-galvanic cells in the weldment. Grooves that formed on both sides of and parallel to the weld, seemingly in a fully martensitic area for the high-sulfur (0.1 wt%) and low-sulfur (0.01 wt%) carbon steel casts, after exposure to flowing 3.5% NaCl, appeared to substantiate the conclusion that it is not the difference in microstructure that causes grooving since some martensitic regions grooved while others did not. It was evident that the weld thermal cycle affected the distribution and morphology of existing sulfides. It was shown how sulfides much smaller than those in the stringers in the base metal formed in networks on the original austenite grain boundaries. This was particularly apparent in the weld metal, but, more importantly, it also occurred in the high-temperature HAZ next to the fusion line (Figure 9). It is thought that the regions of high-temperature HAZ and weld metal (next to the fusion line), immediately adjacent to one another, experience comparable thermal cycles, which affect the sulfide morphology and distribution in the same manner. These networks of sulfides on grain boundaries form a natural path for corrosion. Whether the corrosion propagates along this path through the dissolution of the sulfides or the disbondment of the sulfides from the matrix or any of the other proposed mechanisms,8 the corrosive effect remains the same. The corrosion along networks of sulfides seems to be confirmed by the appearance of the grooves, where it was noted that the groove actually consists of multiple pits that eventually coalesced. A second aspect, which is most important in the fusion line area, is that the presence of active sulfides could be detected. It has been shown that these sulfides are susceptible to corrosion,7,12 and in this case the highest concentrations of active sulfides corresponded with areas of increased corrosive attack. It is proposed that the combination of active sulfides, caused by rapid cooling, and the accompanying redistribution of small sulfides in networks along prior austenite grain boundaries make the fusion line and immediate surroundings susceptible to localized attack. It is possible that corrosion initiates readily in response to the presence of the active sulfides and then propagates along the sulfide networks.



CONCLUSIONS ❖ Literature shows that localized corrosion of carbon steel weldments occurs for various environments, steel types, and weld procedures. Only some fundamental aspects of this diverse phenomenon were considered here. This study is a starting place for systematic investigation into the causes of this localized type of corrosion. It is recognized that other parameters such as different steel types, heat input during welding, preheat treatment, residual stresses, grain size, and filler metal may have significance. ❖ Results of the investigation into the corrosion behavior of the microstructures that possibly could form in the weldment during the weld thermal cycle indicated a significant initial difference in corrosion behavior of the different microstructures, but as stable conditions were reached, the behavior did not differ markedly. This suggested that none of these microstructures was predominantly anodic or predominantly cathodic relative to the others. ❖ The corrosion behavior of the various microstructures probably does not differ sufficiently to cause the selective corrosion involved in the grooving phenomenon. ❖ Under the conditions investigated, it appears that the fusion line area is particularly susceptible to localized corrosive attack. ❖ Investigation of the fusion line area revealed that a redistribution of sulfides occurred and that the sulfides were closely related to the prior austenite grain boundaries. It was found that the highest concentrations of active sulfides correspond with areas of increased corrosion attack.


❖ It is proposed that the susceptibility to localized corrosion is caused by the unique combination of active sulfides and networks of small sulfide inclusions on original austenite grain boundaries.

ACKNOWLEDGMENTS The authors acknowledge the Council for Minerals Technology (Mintek) and the Center for Corrosion Engineering (sponsored by Iscor, Columbus, Eskom, and the THRIP program of the Department of Trade and Industry) for their financial support, and the Laboratory for Microscopy and Microanalysis of the University of Pretoria for use of their facilities. REFERENCES 1. 2. 3. 4. 5. 6.

7. 8. 9. 10. 11. 12.

13. 14. 15.

D.D. Marsden, CHEMSA 10 (1978): p. 135. E. Räsänen, K. Relander, Scand. J. Metall. 7 (1978): p. 11. K.G. Mishra, C.R. Das, Brit. Corros. J. 22, 3 (1987): p. 195. C.A. Loto, M.A. Matanmi, Brit. Corros. J. 24, 1 (1989): p. 36. N. Rothwell, M.E.D. Turner, MP 29, 2 (1990): p. 55. S. Endo, M. Nagae, T. Wada, “Preferential Corrosion Properties in CO2 Containing Environment and Mechanical Properties of Welded Joints of Linepipes,” Proc. 3rd Int. Offshore and Polar Engineering Conf., held 6-11 June 1993 (Singapore, The International Society of Offshore and Polar Engineers), p. 321. M. Kowaka, Metal Corrosion Damage and Protection Technology (New York, NY: Allerton Press, 1990), p. 71-76. S.C. Srivastava, M.B. Ives, Corrosion 43, 11(1987): p. 687. M.G. Fontana, Corrosion Engineering, 3rd ed. (New York, NY: McGraw-Hill, 1987), p. 502-503. M. Stern, A.L. Geary, J. Electrochem. Soc. 104 (1957): p. 56. H.S. Isaacs, Corros. Sci. 29, 2/3 (1989): p. 313. G. Wranglén, “Active Sulfides and Pitting Corrosion,” in Localized Corrosion—NACE 3, eds. R.W. Staehle, B.F. Brown, J. Kruger, A. Agrawal (Houston, TX: NACE International, 1974), p. 462. C. Kato, Y. Otoguro, S. Kado. Y. Hisamatsu, Corros. Sci. 18 (1978): p. 61. C. Duran, E. Treiss, G. Herbsleb, MP 9 (1986): p. 41. G. Wranglén, Corros. Sci. 14 (1974): p. 331.

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