Effect of “475 °C embrittlement” on duplex stainless steels localized corrosion resistance

Effect of “475 °C embrittlement” on duplex stainless steels localized corrosion resistance

Corrosion Science 47 (2005) 909–922 www.elsevier.com/locate/corsci Effect of ‘‘475 C embrittlement’’ on duplex stainless steels localized corrosion r...

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Corrosion Science 47 (2005) 909–922 www.elsevier.com/locate/corsci

Effect of ‘‘475 C embrittlement’’ on duplex stainless steels localized corrosion resistance F. Iacoviello b

a,*

, F. Casari b, S. Gialanella

b

a Universita` di Cassino, via G. Di Biasio 43, 03043, Cassino (FR), Italy Universita` di Trento, Dip. dei Materiali e Tecnologie Industriali, via Mesiano 77, Trento, Italy

Received 10 November 2003; accepted 11 June 2004 Available online 3 September 2004

Abstract The influence of the 475 C ageing treatment on the localized corrosion resistance of austenitic–ferritic (duplex) stainless steels has been investigated by means of double loop electrochemical potentiodynamic reactivation (DL-EPR) and potentiostatic tests. The effect of different ferrite/austenite (a/c) volume fractions has been also considered. For this purpose, two different 22 Cr 5 Ni duplex stainless steels, with different a/c ratios (respectively equal to 1 and 1.5) have been investigated. The ageing treatment at 475 C was conducted up to 1000 h. The resulting microstructural modifications were analyzed with transmission electron microscopy observation. The microstructure resulting from solid state transformations, like spinodal decomposition and G-phase precipitation, were characterized and the relevant mechanisms identified. An evident influence of the ferrite and austenite volume fraction on the Gphase formation and spinodal decomposition kinetics were found. The effects on the localized and selective corrosion attack susceptibility were also investigated.  2004 Elsevier Ltd. All rights reserved. Keywords: A. Stainless steel; B. TEM; C. Selective oxidation

*

Corresponding author. Tel.: +39 0776 2993681; fax: +39 0776 2993695. E-mail address: [email protected] (F. Iacoviello).

0010-938X/$ - see front matter  2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2004.06.012

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1. Introduction Due to their good mechanical properties and their excellent corrosion resistance, in many environments and operating conditions, like chloride induced stress corrosion, both under generalized and localized, duplex stainless steels are successfully used in chemical, petrochemical, nuclear, fertilizer and food industries [1–3]. Depending on their chemical composition, these steels are prone to age hardening and embrittlement over a wide temperature range. This is mainly due to precipitation phenomena that may occur inside ferrite grains and at ferrite–austenite grain boundaries [4]. Three different critical temperatures ranges are present: • Above 1050 C. Duplex stainless steels, that have a fully ferritic solidification structure, upon cooling, partly transform into austenite. This transformation is reversible: therefore, any temperature increase above 1050 C implies a ferrite volume fraction increase and a decrease in the partition coefficients of the alloying elements [5]. • Between 1050 and 600 C. This critical temperature range is characterized by the formation, both inside austenite and ferrite, of a variety of secondary phases that may precipitate with incubation times that are strongly affected by the chemical composition [6]: r phases, nitrides (Cr2N, p), secondary austenite, v and R phases, carbides (M7C3, M23C6). The precipitation of these carbides, nitrides and secondary phases strongly influences mechanical properties and corrosion resistance of duplex stainless steels [7–9]. • Between 600 and 300 C. This temperature range is characterized by the spinodal decomposition of ferrite into Cr-poor a and Cr rich a 0 domains. Other precipitation processes would also occur. Among them, the main one is the Ni, Si, Mo-rich G phase precipitation [10–13]. These particles are very small (usually from 1 to 10 nm, occasionally up to 50 nm) and they precipitate, more or less uniformly, within the ferrite grains, depending on the actual chemical composition of the steel (e.g. Mo-bearing steels show a more uniform precipitation than Mo-free steels). However, these particles are shown to form preferentially on dislocations and at a–c interfaces. Their composition depends not only on the steel composition, but also on the ageing conditions. For instance, the overall concentration in G-forming elements increases from 40% to 60% if tempered at 350 C respectively for 1000 and 30,000 h. Solubilised duplex stainless steels display a good corrosion resistance in many environments such as sulphuric acid, hydrochloridric acid or nitric acid [14]. On the other hand, they are rather susceptible to several localized corrosion attacks, such as pitting, crevice or intergranular corrosion. Nonetheless, duplex steels afford resistance to generalized or localized corrosion attacks that is still higher than that of many austenitic grades. Duplex stainless steels are also characterized by a lower susceptibility to the intergranular corrosion, related to the high temperature of the sensitization phenomenon, than austenitic or ferritic grades [13,15]. Over the last years,

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some researches have been carried out about the possibility for duplex stainless steels to suffer from a selective or a localized corrosion attack also at lower temperatures, corresponding to the ‘‘475 C embrittlement’’ temperature range [16,17]. It turns out that duplex stainless steels show a localized corrosion attack only after a long term ageing heat at low temperatures. The responsibility for this degradation phenomenon is attributed to the spinodal decomposition and the consequent Cr-rich a 0 precipitation and the resulting Cr depletion in the alloy surrounding matrix. This Cr depletion is more evident in the neighborhood of the G-phase particles [18]. Probably, the localized corrosion resistance is also affected by G-phase precipitation. The aim of this work is the analysis of the localized corrosion resistance in two 22 Cr 5 Ni austenitic–ferritic stainless steels characterized by different a/c ratios. The investigation was performed after several ageing heat treatments at 475 C (up to 1000 h). This resistance was evaluated by means of electrochemical potentiokinetic (EPR tests) and potentiostatic tests. The results of the corrosion tests were related to the evolution of the mechanical properties of the steels [7,9].

2. Materials and experimental procedures Two different 22 Cr 5 Ni austenitic ferritic stainless steels were considered for the investigation (Tables 1 and 2). They were characterized by different a/c ratios, respectively equal to 1 and 1.5. These steels were both investigated after solubilisation at 1050 C for 1 h, followed by a water quenching. Tempering heat treatments were conducted at 475 C for times up to 1000 h, namely 1, 10, 100 and 1000 h. Microstructure analyses were performed using a transmission electron microscope (TEM). Thin foils for TEM observations were prepared by a preliminary mechanical thinning, down to 100 lm thickness, followed by twin-jet polishing in a solution of 2butoxiethanol (90%) and perchloric acid (10%) at 3 C, using a voltage of 3 V and a current of 3.6 mA. Observations were performed with an analytical instrument operated at 120 kV, using a double-tilt sample holder for a more efficient acquisition of selected area electron diffraction (SAED) patterns. The microscope was equipped

Table 1 Chemical composition for the 22 Cr 5 Ni duplex stainless steel with an a/c ratio equal to 1 C 0.019

Si 0.39

Mn 1.51

P 0.022

S 0.002

Cr 22.45

Ni 5.50

Mo 3.12

N 0.169

Table 2 Chemical composition for the 22 Cr 5 Ni duplex stainless steel with an a/c ratio equal to 1.5 C 0.025

Si 0.385

Mn 1.428

S 0.011

P 0.028

Ni 5.639

Cr 22.78

Mo 2.941

Cu 0.148

Co 0.160

N 0.129

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with an energy dispersive X-ray spectrometer (EDXS). Concentrations of the main elements, but carbon, were evaluated using a standardless analysis program. The 475 C ageing influence on the localized corrosion susceptibility was assessed by the ‘‘double loop electrochemical potentiokinetic reactivation’’ test (DL-EPR test). This test is quite effective in investigating quantitatively the degree of sensitization of steel alloys, and the resulting corrosion susceptibility [19]. Although this test is actually standardized for austenitic stainless steels, it is widely applied also to stainless steels with different microstructures, such as austenitic–ferritic ones [20,21]. In this work an aqueous solution 0.5M H2SO4 + 0.01 KSCN was considered, with KSCN added as a depassivating agent [22], with an exposed area of the sample of 1 cm2. After establishing of Ecorr, the specimen was polarized from the initial potential (500 mV/SCE), in the cathodic region, to a version potential of +200 mV/SCE, with a sweep rate of 50 mV/min. As soon as this potential was reached, the scanning direction was reversed and, always considering the same sweep rate, the potential was decreased to the initial potential. Tests were performed at 25 C and they were repeated five times, in order to control the results reproducibility. Working electrodes consisted of cubic samples insulated in epoxy resin, polished with a 1200 silicon carbide paper, rinsed in alcohol and dried before each measurement. A platinum counter electrode and a saturated calomel reference electrode were used. Potentiostatic tests were performed (same solution of DL-EPR tests; potential pause equal to 0 mV/SCE for 5 min; test duration between 3 and 15 min) at potential values selected on the basis of the DL-EPR tests results. Specimens surfaces were analyzed by means of a scanning electron microscope (SEM), in order to identify the main attack morphologies. The results of the DL-EPR and potentiostatic corrosion tests were compared with the evolution of the mechanical properties of the steels. In fact, both steels were formerly investigated as concerns the influence of 475 C ageing on the slow strain rate tensile properties (a/c ratio equal to 1.5 [7]) and on the fatigue crack propagation resistance (a/c ratio equal to 1 [9]). An evident influence of 475 C ageing time on the evolution of the mechanical properties was always identified (Figs. 1 and 2), although this influence was less evident than the modifications obtained for higher tempering temperatures (about 800 C) for the same or shorter tempering times.

3. Results and discussion 3.1. 475 C embrittled duplex stainless steels microstructure TEM analysis The typical solution annealed (1050 C) microstructure of both investigated stainless steels is characterized by precipitate free grains, that are elongated along the rolling direction (Fig. 3, duplex stainless steel with a/c = 1). Optical microscope or SEM observations do not display any microstructural modifications following the 475 C tempering heat treatment. In this respect, TEM observations proved more successful. Considering the duplex stainless steel tempered for 1000 h at 475 C with an a/c ratio equal to 1, in

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-6

da/dN [m/cycle]

10

Annealed 375°C-1000h 475°C-1000h 800°C-1h

-7

10

-8

10

-9

10

-10

10

10

3

1/2

50

∆K [MPa m ] Fig. 1. Influence of the duplex stainless steel tempering temperature on the fatigue crack growth resistance (a/c ratio equal to 1 [9]).

1400

UTS [MPa]

1200

1000

800

600

475°C 800°C

1

10

100

1000

10000

Time [hours] Fig. 2. Influence on the low strain rate UTS (a/c ratio equal to 1.5 [7]) of the duplex stainless steel tempering temperature and time.

Fig. 4 a ferritic grain, as observed along the h111i direction, with the corresponding selected area diffraction pattern is shown. Dislocations lines are visible inside the grains and piled-up along grain boundaries. As to the diffraction pattern, no other reflections than those of the ferritic phase are visible. Similar observations were conducted along the h100i and h110i directions. As for the h111i direction, the diffraction patterns of the other two considered orientations of the ferritic structure, h100i and h110i, did not show any other reflection than those of the body centered cubic cell. This would suggest that no other phase precipitated in the ferritic matrix. However, the presence of the G-phase, reported to form in similar steels as the present one, cannot be ruled out. Nucleation of this phase is highly favored by the extraenergy associated with dislocations, particularly, and other microstructural defects possibly present in the material. Dislocations structures have been systematically observed in the examined material. Some dislocation lines, when observed at higher magnification, display a pinned structure as the one shown in Fig. 5. This may

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Fig. 3. SEM micrograph of the investigated a/c = 1 duplex stainless steel in solution annealed condition (after electrochemical etching in oxalic acid, 3 V—60 s).

Fig. 4. TEM micrograph of a ferritic region. Bright field image and selected area diffraction pattern for h111i zone axis. 22 Cr 5 Ni duplex stainless steel tempered for 1000 h at 475 C (a/c ratio equal to 1).

indicate that precipitation has actually started, but it is still at too an early stage to produce a detectable contribution to diffraction patterns. Still considering the higher magnification image (Fig. 5), it is possible to see a slightly modulated contrast, that could be due to the early stages of the spinodal decomposition of ferrite into a and a 0 phases [13]. They both have a body centered cubic (bcc) structure, with a very similar parameter and thus retain coherent interfaces. Considering the duplex stainless steel with an a/c ratio equal to 1.5 and tempered for 1000 h at 475 C, some differences are evident. In the ferritic matrix a number of elongated precipitates, having dimensions of 50–100 nm, are visible (Fig. 6) and the electron diffraction pattern, although too scarce to provide precise crystallographic information, still indicates the precipitates have their own (means, different from the matrix) crystalline structure. EDXS analyses (Table 3) have been conducted in different regions. Precipitates are richer in silicon, nickel and molybdenum as com-

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Fig. 5. TEM bright field image of a dislocation in a ferritic region imaged along h110i direction. 22 Cr 5 Ni duplex stainless steel tempered for 1000 h at 475 C (a/c ratio equal to 1).

Fig. 6. TEM micrograph displaying a detail of the precipitates present in the ferritic matrix with SAED: precipitate spots are visible (a/c ratio equal to 1.5).

Table 3 EDXS quantitative analyses (wt.%) of different regions of the sample Reference

Si

Cr

Fe

Ni

Mo

Precipitate (Fig. 6) Matrix close to precipitate (Fig. 6) Matrix far from precipitate (Fig. 6) Austenitic grain (Fig. 9) Ferritic grain with spinodal precipitates (Fig. 9)

2 1 0 1 1

24 29 29 22 33

55 62 59 67 58

7 5 6 8 5

12 3 7 2 3

pared to the surrounding matrix, which results to be depleted of the above elements, as can be seen from a comparison between the concentrations of matrix regions close and far away from the precipitates. Obviously, this depletion is more pronounced in the neighborhood of the precipitates than in remote regions of the matrix.

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Fig. 7. TEM micrograph displaying precipitation along grain boundaries and dislocations (a/c ratio equal to 1.5).

Considering different tilting angles, it comes out that precipitates are mostly localized along dislocations and grain boundaries (Fig. 7), where diffusion, that drives precipitate ripening, occurs more rapidly. Another feature that can be clearly seen in the TEM micrographs is a fine dispersion of nanometric precipitates (5–10 nm). This is shown in greater detail in Fig. 8. The image, taken along the [001] zone axis of the bcc lattice of the ferritic matrix, displays a microstructure that may come from a spinodal decomposition, a transformation which is to be expected in this kind of steels. No extra-spots, in addition to those of the ferrite phase, have been observed in the diffraction pattern, indicating a complete coherency of the precipitates with respect to the matrix. This would also agree with the proposed spinodal-like structure, in which no sharp interface separates the newly forming from the parent phase. As an additional note, it is to be said that, in agreement with the spinodal origin of the clusters, they can be imaged under different sample orientations, as their contrast is mostly due to a different concentration as compared to the parent phase. The fine

Fig. 8. TEM micrograph of the fine precipitation present in ferritic grains with SAED (a/c ratio equal to 1.5).

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Fig. 9. TEM micrograph displaying two grains (austenitic on the left and ferritic with fine precipitation on the right) and corresponding SAED (a/c ratio equal to 1.5).

dispersion has been observed only in the ferritic grains, whereas the austenite phase displays quite a clear contrast, with annealing twin and rows of dislocation loops (Fig. 9). SAED analysis shows that austenite structure has not any specific orientation relationship with ferrite. EDXS measurements (Table 3) confirm the austenite and ferrite phases identification, according to the Ni partition phenomenon. TEM analysis, performed on the two investigated duplex stainless steels after a tempering treatment for 1000 h, shows a strong difference in the spinodal decomposition and G-phase precipitation that is probably due to the different a/c ratio. In fact, considering that a/a and a/c grain boundaries are preferential precipitation sites, the higher the a/c ratio the more evident should result the microstructural modifications. 3.2. Corrosion tests results DL-EPR tests results are characterized by a good reproducibility. For all the investigated conditions, only one curve is sufficient to characterize the applied potential–current density behavior. Both the investigated stainless steels in solutioned conditions do not show any reactivation maxima. The activation curves depend on the a/c ratio (Fig. 10).

2

Current density [mA/cm ]

10

α / γ = 1.5

1 α /γ = 1

0.1 -0.5

-0.4

-0.3

-0.2

-0.1

0.0

0.1

0.2

Potential [V/SCE] Fig. 10. DL-EPR test results for the investigated duplex stainless steels (after solutioning treatments).

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If this value is equal to 1, only one maximum is evident in the activation curve, corresponding to an applied potential of 240 mV/SCE, with a current density of about 1 mA/cm2. If a/c ratio is equal to 1.5, the higher ferrite volume fraction implies the presence of two maxima, with higher density current values: the first corresponds to about 280 mV/SCE and the second corresponds to about 180 mV/SCE. This maximum is more evident and it is characterized by a current density of about 4 mA/ cm2 (up to four times higher than the value obtained considering the a/c ratio equal to 1). For the solutioned duplex stainless steels, differences in the DL-EPR results are due to the different austenite content that is nobler than ferrite: the higher the austenite volume fraction, the lower the activation maximum density current. The influence of the 475 C annealing on the DL-EPR results depends on the a/c ratio. If a/c ratio is equal to 1 (Fig. 11), both the activation and the reactivation curves do not differ from the curves obtained considering the stainless steel after solubilisation for all the investigated treatment durations: the activation maximum corresponds to the same potential (about 250 mV/SCE) with almost the same current density (about 1 mA/cm2) and no reactivation maxima are observed. This behavior should be compared with the influence of the tempering heat treatment on the fatigue crack propagation resistance (Fig. 1). TEM observations do not show an evident G formation. The very early stages of a possible ferrite spinodal decomposition and G-phase precipitation, mainly at dislocation lines, can be inferred. A limited dislocations pinning, due to the presence of G particles, may be able to decrease the steel fatigue crack propagation resistance. However, these transformations, if actually underway, would not imply a significant matrix depletion of G-forming elements. Therefore, differences in localized corrosion resistance between the solutioned and the 475 C embrittled steel are not observed. If a/c ratio is equal to 1.5 (Fig. 12), an evolution of the DL-EPR test results is evident for the 475 C tempering temperature, for example treated for both 100 and 1000 h. After 100 h at 475 C, the maximum current density value corresponding to 180 mV/SCE is almost constant, but the second lower maximum almost disap-

2

Current density [mA/cm ]

10

1 475°C-1000h Solubilized

0.1 -0.5

-0.4

-0.3 -0.2 -0.1 0.0 Potential [V/SCE]

0.1

0.2

Fig. 11. a/c = 1 duplex stainless steel EPR test results (after solutioning and after ageing at 475 C for 1000 h).

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2

Current density [mA/cm ]

10

475°C-1000h activation reactivation

919

475°C - 100h

Solubilized

1

0.1 -0.5

-0.4

-0.3

-0.2

-0.1

0.0

0.1

0.2

Potential [V/SCE] Fig. 12. a/c = 1.5 duplex stainless steel EPR test results (after solubilisation and after ageing at 475 C for 100 and 1000 h).

pears. Furthermore, a very low current density reactivation maximum started to appear. As a consequence of a longer tempering treatment, an increase of the activation maximum current density is to be expected (from about 4 up to about 10 mA/cm2). On the other hand a lower activation maximum at 280 mV/SCE has been detected, probably due to the increase of the first maximum. Furthermore, two different reactivation maxima are evident. The first corresponds to 280 mV/SCE, with a current density of about 3.5 mA/cm2 (this value is comparable with the activation maximum obtained with the solubilised steel). The second one is less evident at 480 mV/SCE, with a current density of about 0.5 mA/cm2. Scanning electron microscope surface analysis performed on 475 C–1000 h tempered duplex stainless steels specimens after potentiostatic tests shows a morphology attack that depends both on the a/c ratio and on the applied potential. For the a/c ratio equal to 1 (higher austenite content) no localized corrosion attack morphologies are observed for all the investigated potentials. For the a/c ratio equal to 1.5 (lower austenite content) localized attack morphologies are observed corresponding to both the reactivation maxima. Considering the higher maximum (corresponding to about 280 mV/SCE) a selective attack of ferrite grain boundaries is evident (Fig. 13). Considering the lower maximum (corresponding to about 480 mV/SCE) a selective attack morphology on ferrite grains is more evident (Fig. 14). Since grain boundaries are the preferential, but not unique, G-phase precipitation sites, and spinodal decomposition usually takes place more uniformly in ferrite grains, these two reactivation maxima are connected to two different phenomena: the spinodal decomposition (480 mV/SCE) and the G-phase precipitation (280 mV/SCE). Comparing the corrosion resistance and the mechanical behavior evolution, as a function of the tempering duration (Fig. 2), in the solution treated and the 475 C embrittled steel, and considering TEM analysis results (Figs. 7–9), it is evident that more advanced stage of the microstructural transformations imply both mechanical properties and localized corrosion resistance modifications. Dislocations pinning, due to the G particles presence and to the spinodal decomposition, is more evident

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Fig. 13. Duplex stainless steel with an a/c ratio equal to 1.5: SEM surface analysis after a potentiostatic test performed at 280 mV/SCE for 5 min.

Fig. 14. Duplex stainless steel with an a/c ratio equal to 1.5: SEM surface analysis after a potentiostatic test performed at 480 mV/SCE for 15 min.

and the matrix local depletion of G-forming elements implies the 475 C embrittled stainless steel to be more susceptible to a localized corrosion attack.

4. Conclusions In the present work two 22 Cr 5 Ni austenitic–ferritic (duplex) stainless steels, featuring different ferrite–austenite ratios (e.g. a/c equal to 1 and 1.5), were investigated as concerns microstructural transformations that can take place during heat treatments at 475 C for times up to 1000 h. TEM investigations were performed and the results were compared to the evolution of the mechanical properties and the localized corrosion resistance analysis. This aspect was investigated by means of double loop electrochemical potentiodynamic reactivation (DL-EPR) and potentiostatic tests.

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On the basis of the experimental results, the following conclusions can be drawn: • TEM observations suggest that the kinetics of spinodal decomposition and Gphase precipitation is influenced by the ferrite/austenite ratio, in agreement with previous studies [14]. The 475 C-1000 h tempered duplex stainless steel with the higher ferrite content (a/c = 1.5) is characterized by a more advanced spinodal decomposition and G-phase precipitation, as compared to the steel with the lower ferrite content (a/c = 1). • A reduction in the resistance of the fatigue crack growth has been observed since the very early stage of spinodal decomposition, i.e., in the steel having a a/c = 1 with very small G-phase particles (Fig. 1). The effect is even more pronounced in the a/c = 1.5 material, and can be ascribed to dislocations interaction with more or less developed G-phase particles. Fluctuations in the chemical composition of ferrite grains, with a well developed spinodal structure and coarse G-phase precipitates, induced a larger susceptibility to localized corrosion in the steel with a higher ferrite content. Acknowledgement Prof. A. Molinari, University of Trento, is warmly acknowledged for useful discussions. References [1] P. Lacombe, B. Baroux, G. Beranger, Les aciers inoxydables, Les e`ditions de physique, Les Ulis Cedex A, France, 1990, p. 663. [2] F. Dupoiron, J.P. Aoudouard, Scandinavian Journal of Materials 25 (1996) 95. [3] J. Charles, Duplex 2000 6th World, Venice, Italy, October 2000, p. 1. [4] J. Charles, Duplex stainless steel Õ91, Beaune, France, October 1991, p. 3. [5] A.J. Strutt, G.N. Lorimer, C.V. Roscoe, K.J. Gradwell, Duplex stainless steel Õ86, The Hague, Netherlands, October 1986, p. 310. [6] J. Charles, P. Pugeault, F. Dupoiron, D. Catelin, Bulletin du Cercle des Me´taux XV (1988) 24. [7] F. Iacoviello, M. Habashi, M. Cavallini, Materials Science and Engineering A 224 (1997) 116. [8] F. Iacoviello, J. Galland, M. Habashi, Corrosion Science 260, 40, 8 (1998) 1281. [9] F. Iacoviello, M. Boniardi, G.M. La Vecchia, International Journal of Fatigue 21 (October) (1999) 957. [10] M. Guttmann, Duplex stainless steel Õ91, Beaune, France, October 1991, p. 79. [11] A. Mateo, L. Llanes, M. Anglada, A. Redjaimia, G. Metauer, Journal of Materials Science (1997) 4533. [12] F. Danoix, P. Auger, Materials Characterisation 44 (2000) 177. [13] C.J. Park, H.S. Kwon, Corrosion Science 44 (2002) 2817. [14] R.N. Gunn, Duplex stainless steels—Microstructure properties and applications, Abington Publishing, Cambridge, UK, 1997. [15] F. Iacoviello XXIX AIM Meeting, 13–15 November 2002, Modena, Italy, p. 153. [16] K. Ravindranhath, S.N. Malhortra, Corrosion Science 37 (1995) 121. [17] J.E. May, C.A.C. de Sousa, S.E. Kuri, Corrosion Science 45 (2003) 1395. [18] F. Danoix, S. Chambreland, J.P. Massoud, P. Auger, Duplex stainless steel Õ91, Beaune, France, October 1991, p. 111.

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[19] ASTM G 108-92 Standard test method for electrochemical reactivation (EPR) for detecting sensitization of AISI 304 and 304L stainless steel, Section 3, 03.02. [20] N. Lopez, M. Cid, M. Puiggali, I. Azkarate, A. Pelayo, Materials Science and Engineering 229 (1997) 123. [21] V. Cihal, R. Stefec, Electrochemical Acta 46 (2001) 3867. [22] T.-F. Wu, T.-P. Cheng, W.-T. Tsai, Journal of Nuclear Materials 295 (2001) 233.