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Effect of Mn addition on the precipitation and corrosion behaviour of 22% Cr economical duplex stainless steel after isothermal aging at 800 °C
Journal of Alloys and Compounds 699 (2017) 334e344
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Effect of Mn addition on the precipitation and corrosion behaviour of 22% Cr economical duplex stainless steel after isothermal aging at 800 C Zaiqiang Feng a, Yinhui Yang b, *, Jia Wang c a b c
School of Mechanical Engineering, North China University of Water Resources and Electric Power, Zhengzhou 450045, PR China School of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, PR China School of Materials Science and Engineering, Tongji University, Shanghai 200092, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 February 2016 Received in revised form 22 December 2016 Accepted 3 January 2017 Available online 4 January 2017
The microstructural and corrosion behaviour of Fe-22Cr-1.9Ni-2.3Mo-0.2N-xMn low nickel type duplex stainless steel (DSS) with Mn contents ranging from 4.3 to 9.7 wt% were studied after aging treatment at 800 C lasting from 30 to 930 min. The corrosion behaviour was investigated in 3.5 wt% NaCl solution by measurements of potentiodynamic polarisation and electrochemical impedance spectroscopy. The microstructural analysis shows that the s precipitates mainly nucleated along g/d phase boundaries and grew into the adjacent d-ferrite by prolonging the aging time, having the same eutectoid structures as commercial high-Ni DSS. At high Mn contents, the growth rate of s precipitates was initially high with aging times from 30 to 150 min, but decreased gradually up to 930 min with more Mn participation. A Mn content of 4.3 wt% brought only about a 150 mV pitting potential (Eb) drop for the aged specimens, while further addition of Mn significantly lowered Eb at the early stage of aging with more aging-induced s precipitates. The metastable pit nucleation resistance decreased with Mn content from 4.3 to 6.9 wt% and was suppressed by a higher Mn addition of 9.7 wt% for aging treatments from 150 to 930 min. The pits mainly initiated along d/g boundaries and within d ferrite, then propagated around s precipitates with increasing aging time, and more Mn gives rise to both the number and size of pits in the early stage of aging treatment. The charge transfer resistance (Rt) depends strongly on the existence of defects caused by the s precipitates, while increasing Mn content slightly decreases the passive film resistance (Rfilm) in specimens aged from 60 to 930 min. © 2017 Elsevier B.V. All rights reserved.
Keywords: Metals and alloys Intermetallics Microstructure Precipitation Corrosion
1. Introduction Duplex stainless steels (DSSs) are widely employed in many fields, such as the oil, petrochemical, nuclear and desalination industries, with a two-phase structure of ferrite and austenite, which takes advantage of their attractive combination of good corrosion resistance and mechanical properties under harsh corrosive environments [1,2]. However, the increasing price of Ni has made the cost of high-Ni-content DSS prohibitive for many industrial applications. Therefore, it is urgent to find proper alternatives for Ni. For instance, Mn-N DSSs are being actively developed by replacing Ni in high-Ni DSSs with cheaper Mn and N. Both Mn and N are known to
* Corresponding author. E-mail address: [email protected] (Y. Yang). http://dx.doi.org/10.1016/j.jallcom.2017.01.031 0925-8388/© 2017 Elsevier B.V. All rights reserved.
be good austenite stabilisers, which enhance the strengthening effect and improve the corrosion resistance of steels [3]. Mn, which is almost an order of magnitude cheaper than Ni at an equivalent weight [4,5], can further help increase the solubility of N in stainless steels [6], whereas the low solubility of N in liquid steel under atmospheric pressure has been problematic for production. DSSs are frequently exposed to high temperatures during hot working processes, which easily induce intermetallic phase precipitation. For example, Cr- and Mo-rich phases, such as Cr23C6 carbide and c and s phases, are easily formed when DSSs are subjected to elevated temperatures of 450e1000 C [7,8]. These precipitates often cause a loss of toughness and corrosion resistance. Moreover, previous studies noted that the detrimental s phase forms most rapidly at 800 C [9,10]. Johnson et al. [11] also reported that the initial nose of s phase formation in the timetemperature-transformation diagram of DSS occurs earliest at
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800 C. Thus, it is effective to investigate precipitation behaviour with different Mn additions under severe conditions of an 800 C aging temperature. The Cr-depleted zones in the vicinity of s precipitates are susceptible to pitting corrosion due to weakened passivation. Zhang et al. [12] investigated the effect of isothermal aging on the pitting corrosion resistance of UNS S82441, and found pitting transfers from the austenite phase into Cr-depleted zones with increasing aging time. It has been demonstrated that adding Cr and Mo can accelerate the precipitation of the s phase at high temperatures [13], while the partition of these alloy elements in two phases can be altered in the presence of Mn. Therefore, during aging treatment of 800 C, different Mn contents can lead to various degrees of precipitate transformations, which take place within the ferrite matrix and along grain boundaries due to the much faster diffusion rates of the alloying elements within ferrite compared to austenite [14,15]. It is known that the detrimental effect of Mn on the pitting corrosion resistance is associated with nonmetallic MnS inclusions in high-S alloys. The reported effects of Mn on the pitting corrosion
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resistance are controversial. Toor reported that the resistance to pitting and metastable pitting corrosion in a novel high-Mn Ni-free DSS decreased with increasing Mn content since a great number of (Mn, Cr) oxides acted as preferential sites of pitting [5]. Jang et al. believed that the tensile properties of CD4MCU0-cast DSS were determined by the volume fraction of hard ferritic phase and the shape of austenitic phase, and the resistance to pitting corrosion and to stress corrosion cracking in 3.5% NaCl þ5% H2SO4 aqueous solution was recovered with Mn contents from 0.8 to 2% [4]. However, there is little research available regarding the influence of Mn on the corrosion behaviour of DSSs caused by intermetallic phase precipitation during high temperature aging treatment. This study selects Mn content as the main parameter and investigates its effect on the microstructure evolution and corrosion behaviour of low nickel DSS by isothermal aging ranging from 15 to 930 min at 800 C. The relationships between the aginginduced precipitation phases under various Mn contents and pitting corrosion behaviour of the newly-developed DSS are discussed in detail.
Table 1 Chemical composition of hot rolled 19% Cr DSS (wt.%). Elements
C
Si
Mn
S
P
Cr
Ni
Mo
Cu
N
Other
Alloy 1 Alloy 2 Alloy 3
0.01 0.01 0.01
0.12 0.11 0.14
4.28 6.90 9.71
0.005 0.004 0.005
0.007 0.006 0.006
22.47 22.49 22.31
1.94 1.93 1.88
2.33 2.25 2.25
0.21 0.22 0.23
0.21 0.22 0.23
Bal.
Fig. 1. Optical images of the specimens aged at 800 C for: (a) 4.3% Mn, 15 min; (b) 6.9% Mn, 15 min; (c) 9.7% Mn, 15 min; (d) 4.3% Mn, 150 min; (e) 6.9% Mn, 150 min; (f) 9.7% Mn, 150 min; (g) 4.3% Mn, 930 min; (h) 6.9% Mn, 930 min; (i) 9.7% Mn, 930 min.
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2. Experimental 2.1. Materials and heat treatment The experimental steels were prepared by induction melting in a 25 kg vacuum induction furnace, followed by casting into ingots. The cast ingots were hot forged into 30 mm plates, then, the asforged samples were hot rolled into 12 mm plates at temperatures ranging from 1050 to 1180 C. The chemical compositions of the rolling plates, designated as Alloys 1, 2 and 3, are shown in Table 1, of which the Mn contents are 4.3, 6.9 and 9.7 wt%, respectively. The as-rolled plates were machined to the size of 12 12 50 mm and solution-treated at 1040 C for 30 min. In order to study the precipitation behaviour under different Mn contents, all the solution-treated specimens were then subject to aging treatment at 800 C for times varying from 15 to 930 min, followed by water quenching.
amount of austenite phases increased marginally with Mn content increasing from 4.3% to 9.7%. At the early aging stage of 15 min (Fig. 1aec), no obvious secondary phase appeared in all the specimens. However, with a longer aging time of 150 min (Fig. 1def), some dark granular precipitates nucleated at g/d interfaces in specimens with 4.3% and 6.9% Mn, while more precipitates formed
2.2. Microstructural characterisation The microstructure of the aged specimens was revealed by electrochemically etching in 40 wt% KOH solution for 10e15 s at room temperature. The chemical compositions of the aginginduced precipitates and g and d phases were measured by energy dispersive X-ray spectroscopy (EDS) attached to a scanning electron microscope (SEM, FEI Quanta 200) with a Robinson backscattered electron detector. In addition, X-ray diffraction (XRD) was used for identification of the precipitates. The volume fractions of the precipitates were measured using the method of manual point count according to ASTM E 562 [16]. 2.3. Electrochemical measurements A Bio-Logic VMP-300 potentiostat was used to perform all the electrochemical measurements in a three-electrode cell. A platinum sheet and a saturated calomel electrode (SCE) were used as the counter and reference electrodes, respectively. Electrochemical potentiodynamic polarisation was performed in a deaerated 3.5 wt % NaCl solution at a temperature around 25 C (±1 C). The specimens, embedded in epoxy resin with an exposure area of 100 mm2, acted as working electrodes. Prior to each experiment, the specimens were ground mechanically up to 3000 grit, rinsed with distilled water and dried in hot air. All potentials are given against the SCE. Then, the corrosion morphology was observed by SEM. The polarisation curves were recorded potentiodynamically with a scan rate of 0.1 mV s1, starting from the free corrosion potential to transpassive potential (vs. SCE). When the passive film breaks down with a rapid rise in the current density, Eb was determined as the current density reached a value of 100 mA/cm2. EIS measurements were executed at open circuit potential (OCP) in a deaerated 3.5 wt% NaCl solution. The frequency was swept from 100 kHz down to 10 mHz at 10 data cycles/decade, with an amplitude of sinusoidal wave of 10 mV. The impedance at the OCP was carried out after immersion at OCP for 30 min. The data were interpreted on the basis of equivalent electrical circuits using the Zview 2.70 program to fit the experimental data. 3. Results and discussion 3.1. Microstructural analysis The microstructural evolution of specimens with different Mn contents during isothermal aging at 800 C is shown in Fig. 1. The bright etched austenite (g) islands were formed of a lath-like morphology within the grey etched ferrite (d) matrix, and the
Fig. 2. XRD results of the specimens aged at 800 C for: (a) 4.3% Mn, 930 min; (b) 6.9% Mn, 930 min; (c) 9.7% Mn, 930 min.
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Fig. 3. SEM micrographs of specimens aged at 800 C for: (a) 4.3% Mn, 630 min; (b) 6.9% Mn, 630 min; (c) 9.7% Mn, 630 min (d) EDX spectrum of s-phase corresponding to (b).
along g/d interfaces and grew into the adjacent d-ferrite regions. These dark secondary precipitates can be identified as the s phase due to its heterogeneous precipitation with the formation of bright secondary austenite (g2) during high temperature aging of 700e900 C [17,18]. As the aging reached 930 min, it is noted that the coarser dark s phase nucleated at g/d interfaces in the specimen with 4.3% Mn (Fig. 1g), whereas the s precipitates increased both in number and size with higher Mn content (Fig. 1hei), which is attributed to the enhanced diffusion of solute elements with higher Mn content. The XRD results of the specimens are presented in Fig. 2. In addition to the major peaks from d-ferrite and austenite, the minor diffraction peaks from the s phase were identified in specimens of different Mn contents, proving the presence of the s phase in the aged specimens. It is also observed that the intensity of s (411) and s (112) peaks increased with adding Mn from 4.3% to 6.9% and up to 9.7%, corresponding to the reduction of the d (1 1 0) peak intensity as a result of the d / s þ g2 transformation [19]. This suggested that the precipitation was accelerated by adding more Mn. Fig. 3 presents the SEM images for specimens aged after 630 min with different Mn contents. For the specimen with 4.3% Mn content, some s particles precipitated at d/g interfaces and penetrated into the d-ferrite matrix. However, when the Mn content was 6.9% and 9.7% (Fig. 3b and c), the lamellar structure, consisting of s and g2 phases, can be clearly observed, the same eutectoid structures were also obtained in commercial DSSs aged at high temperature in spite of very low Mn contents [20]. As shown in Fig. 3d, these s phases were further confirmed through EDX analysis, featuring high Cr and low Mo contents [15]. In addition, it is detected that the s precipitate was partitioned with a higher Mn content than dferrite matrix in the specimen with 6.9% Mn, suggesting that the s
precipitates have a closer affinity to Mn [21,22]. The variations of s phase precipitation at aging of 800 C with different Mn contents are shown in Fig. 4. It is observed that the volume fraction of s phase increased with increasing Mn content during aging time from 15 to 930 min, indicating quicker formation of s precipitates at high Mn content. At early stages of aging time from 30 to 150 min, it is inferred from the slope of the curve that the rate of s phase precipitation in 9.7 wt% Mn content samples was
Fig. 4. Variation of s phase volume fraction as a function of aging time at 800 C in samples with different Mn contents.
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higher than that in the 6.9 and 4.3 wt% Mn specimens. This can be attributed to higher diffusivity of Mn, which is similar to the effect of Mo on the nucleation of the s phase in DSS [23]. However, with the aging time finally increased to 930 min, the amount of s phase in 6.9 wt% Mn Alloy 2 rapidly increased and nearly became saturated, while in the 9.7 wt% Mn Alloy 3, it increased rather slowly due to complete decomposition of d-ferrite. As a consequence, the
Table 2 Pitting and passivation values of the materials tested in 3.5% NaCl solution at different Mn contents. Heat treatment condition
Eb (mV vs. SCE)
ip (mA cm2)
icorr (mA cm2)
Ecorr (mV vs. SCE)
Aging at 800 C 4.3 wt% Mn Aging at 800 C 4.3 wt% Mn Aging at 800 C 4.3 wt% Mn Aging at 800 C 4.3 wt% Mn Aging at 800 C 4.3 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 9.7 wt% Mn Aging at 800 C 9.7 wt% Mn Aging at 800 C 9.7 wt% Mn Aging at 800 C 9.7 wt% Mn Aging at 800 C 9.7 wt% Mn
for 15 min,
591
2.04
1.90
434
for 60 min,
571
2.29
1.95
633
for 150 min,
513
2.69
1.92
644
for 450 min,
445
1.66
1.39
495
for 930 min,
435
1.67
1.45
539
for 15 min,
521
2.29
0.96
606
for 60 min,
429
2.14
1.81
529
for 150 min,
269
2.15
1.84
532
for 450 min,
267
2.04
1.71
549
for 930 min,
234
2.14
1.84
607
for 15 min,
439
2.69
2.62
654
for 60 min,
252
2.19
2.61
656
for 150 min,
154
2.63
1.40
712
for 450 min,
216
2.18
1.54
654
for 930 min,
169
3.47
3.64
771
driving force for the precipitation of s phases tends to be strong with higher Mn contents. 3.2. Potentiodynamic measurements and pitting corrosion resistance analysis Fig. 5 shows the potentiodynamic polarisation curves of all specimens obtained for aging times of 15e930 min in 3.5 wt% NaCl solution at 25 C. As the sweep potential increased to the pitting potential (Eb), the sharp increase in current density was associated with an occurrence of stable pits, indicating the breakdown of the
Fig. 5. Potentiodynamic polarisation curves of samples with different Mn contents in 3.5 wt% NaCl solution aged at 800 C from 15 to 930 min: (a) Alloy1 with 4.3 wt% Mn; (b) Alloy2 with 6.9 wt% Mn; (c) Alloy3 with 9.7 wt% Mn.
Fig. 6. Difference of Eb and Ecorr for different Mn addition samples aged at 800 C tested in 3.5 wt% NaCl solution.
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Fig. 7. SEM-BSE images of specimens aged at 700 C after potentiodynamic polarisation testing for: (a) 4.3 wt% Mn, aged for 60 min; (b) 6.9 wt% Mn, aged for 60 min; (c) 9.7 wt% Mn, aged for 60 min; (d) 4.3 wt% Mn, aged for 930 min; (e) 6.9 wt% Mn, aged for 930 min; (f) 9.7 wt% Mn, aged for 930 min.
passive film. Alloy 1 with 4.3 wt% Mn displays similar polarisation curves throughout all the aging stages, showing a wide potential domain of passivity in the potential range between 0.4 and 0.45 VSCE (Fig. 5a). The Eb value decreased slightly with increasing aging time from 15 to 930 min, owing to the formation of more Crdepleted zones caused by s precipitates with longer aging time [24,25]. Moreover, at high Mn content, in addition to negative shifting of Eb, the difference of Eb between specimens aging for 15 and 930 min was enlarged (Fig. 5b and c), which can be attributed to the ripening of Cr-depleted zones with rapid coalescence of s phases [26]. This was consistent with the above mentioned microstructural observations and discussion. The results of the polarisation measurements are summarised in Table 2. For specimens with 4.3 wt% Mn, the corrosion potentials (Ecorr) decreased with increasing aging time from 15 to 60 and 150 min, while it showed small changes with further aging time to
930 min, this can be explained by the nucleation of precipitates in the initial aging stages. The variation in Ecorr of most samples with 6.9 and 9.7 wt% was small, with their values being around 600 mVSCE. However, Ecorr decreased to 771 mVSCE in the 9.7 wt% Mn specimen with 930 min aging due to the formation of more s phases, and the corresponding corrosion current density, icorr, increased by about 2 mA cm2 compared with lower Mn addition specimens at this aging time. This suggests that the addition of Mn has less of an effect on (Ecorr) SCE of aged specimens than on the low Mn content commercial DSS reported earlier [27]. The passivation current densities, ip, are similar regardless of Mn content and aging conditions, indicating good passivation stability for different aged samples in the presence of Cl. However, the values of Eb exhibited significant differences among samples of different Mn contents. For the specimen with 4.3 wt% Mn, Eb decreased gradually by about 150 mV with increasing aging time from 15 to 930 min, showing
Table 3 Mean compositions (wt.%) and PRENs of g, d and g2 in samples with different Mn contents subject to aging at 800 C for 630 min. Material
Phase
Cr
Mn
Mo
Ni
Na
PRENb
4.3 wt% Mn, aging for 630 min
Austenite (g) Ferrite (d) Secondary austenite (g2) s phase Austenite (g) Ferrite Secondary austenite (g2) s phase Austenite (g) Ferrite (d) Secondary austenite (g2) s phase
21.4 23.1 19.5 25.7 21.2 23.5 20.1 27.1 19.6 22.4 19.6 25.8
5.0 3.9 5.1 5.3 7.9 5.4 7.6 7.5 10.2 8.4 9.5 9.9
1.9 2.5 1.2 2.0 1.8 2.6 1.8 2.2 1.9 2.9 1.8 2.6
2.5 1.5 2.9 1.3 2.4 1.2 2.2 1.1 2.5 1.6 1.5 1.5
0.37 0.05 0.10
33.77 28.95 21.36
0.39 0.05 0.10
30.94 28.18 21.44
0.41 0.05 0.10
27.97 25.07 19.04
6.9 wt% Mn, aging for 630 min
9.7 wt% Mn, aging for 630 min
a b
Nitrogen in ferrite is taken as the saturation value z 0.05%, the rest partitions to austenite [35]. PRENs were calculated with the following equation: PREN ¼ %Cr þ 3.3%Mo þ 30%N1 %Mn.
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small pitting corrosion resistance reduction in the aging time range. With higher Mn content (6.9 wt%), Eb decreased quickly after 60 min aging time and flattened out at around 250 mV until 930 min. In addition, Eb was shifted to a more negative value in the 9.7 wt% Mn sample at the equivalent aging time than that for the 6.9 wt% Mn sample, indicating that higher Mn contents enhance the susceptibility of the aged specimens to pitting corrosion. It is noted that Mn promoted the formation of s precipitates to some degree in aged specimens from the above microstructural analysis. Meanwhile, the s precipitates play an important role in decreasing localised corrosion resistance because of the relatively largevolume Cr-depleted zones surrounding the particles [26]. Thus, a close relationship is observed between the decrease in pitting corrosion resistance of aged samples and the increase of the amount of s phase, indicating that the amount of s precipitates influenced by varying the Mn content plays a critical role in the propagation of the initial pits [28]. The difference between Eb and Ecorr is a measure of the extent of the passive region on the polarisation curve and provides an indication of the susceptibility to pitting and the metastable pits occurring in this region [29e31]. The resistance of the material to metastable pit nucleation increases when this difference increases. The values of (Eb - Ecorr)SCE from samples of different Mn contents at 800 C aging are shown in Fig. 6. The values were similar in all samples at the initial aging of 15 min, then decreased with Mn content from 4.3 to 6.9 wt% and up to 9.7 wt% as the aging time was further prolonged to 930 min, indicating that the metastable pit nucleation resistance decreased with higher Mn content for aging treatment at this temperature. However, as the Mn content increased from 6.9 to 9.7 wt%, the values of (Eb - Ecorr)SCE increased instead with aging time from 150 to 930 min, meaning the metastable pit nucleation was suppressed to some extent. Fig. 7 shows SEM-BSE images of the corroded surfaces of the aged specimens at 800 C for 60 and 330 min after potentiodynamic polarisation tests performed in 3.5 wt% NaCl solution at 25 C. As shown in the SEM surface morphology from 60 min-aged specimens with 4.3 and 6.9 wt% Mn (Fig. 7a and b), the pits initiated not only along d/g phase boundaries and d/d grain boundaries, but also within the d ferrite matrix. This could be explained by the difference in the chemical composition, namely the pitting resistance equivalent (PREN) of the two phases [32]. With a high Mn content of 9.7 wt% (Fig. 7c), an incomplete eutectoid decomposition of the d ferrite (d / s þ g2) induced more pits around the s phases. This indicates that Mn may reduce pitting corrosion resistance in the initial stage of aging. With longer aging treatment up to 930 min, the back-scattered images (Fig. 7def) revealed the areas attacked by the pits are active black regions around s precipitates, which have lower Cr contents in g2 than that in primary g and provide preferential sites for pitting attacks [12]. These black pit regions were propagated with increasing Mn addition from 4.3 to 6.9 wt% and up to 9.7 wt% due to the formation of coarser s phases. These s precipitates made their adjacent regions depleted in Cr and thus incoherent with the matrix, triggering and lowering the pitting potential (Epit)SCE in the Cl-based medium [33,34]. The mean compositions and PRENs of g, d and g2 in specimens with different Mn contents subjected to aging at 800 C for 630 min are listed in Table 3. It shows that the PREN of ferrite was lower than that of the austenite phase in all the specimens, meaning that the primary sites for pit nucleation were in ferrite. Meanwhile, the secondary austenite (g2) has a much lower value of PREN than both the austenite and ferrite phases, indicating that the formation of more g2 accompanied by the nucleation of s and the decomposition of d ferrite (d / s þ g2), inducing more pits around s phases, in
agreement with the post-potentiodynamic polarisation corrosion morphology analysis. It is also concluded from the table that the PREN of d-ferrite decreased with adding Mn from 4.3 to 9.7 wt% due to the effect of Mn on the pitting corrosion [36], indicating that the ferrite was more susceptible to pit attacks in the presence of high Mn content at the early stage of aging. Furthermore, it is noted that the s phase precipitation was partitioned with Mn by EDS analysis, suggesting the addition of Mn accelerated the formation of s precipitates after aging for 90 min, which mostly triggered pitting corrosion in d ferrite.
Fig. 8. Nyquist plots for aged specimens tested in 3.5 wt% NaCl solution with different Mn additions: (a) Alloy1 with 4.3 wt% Mn addition; (b) Alloy2 with 6.9 wt% Mn addition; (c) Alloy3 with 9.7 wt% Mn addition.
Z. Feng et al. / Journal of Alloys and Compounds 699 (2017) 334e344
3.3. EIS measurements EIS measurements were performed to determine the effects of Mn on the electrochemical properties of the passive film aged at 800 C. Fig. 8 presents the Nyquist spectra obtained at the OCP of different aging times with different Mn contents. Impedance spectra are normally displayed in the form of a Nyquist diagram, where the opposite of the imaginary part of impedance is plotted vs. the real part. The diagrams display a capacitive arc that does not end at low frequency, and the radius of the semi-circular arc is related to the polarisation resistance of the passive film [37]. The radius of the capacitive arc decreases with increasing aging time, indicating that the corrosion
341
resistance was lowered with weakening of the passive ability due to s phase precipitation. With the Mn content increased from 4.3 to 9.7 wt% (Fig. 8aec), it is shown that the impedance value decreases at the initial aging stage of 15 min, then flattens out by prolonging the aging time to 930 min, revealing a high sensitivity to aging precipitation. The impedance responses for differently aged specimens are presented in Bode format in Fig. 9. They show time constants at high and low frequencies. The shape of the curves is usually associated with the response of an inhomogeneous film [38]. For the specimen with 4.3 wt% Mn (Fig. 9a), the phase angle q is almost invariable around 70 for 15 min, where the time constants from charge transfer process and passive film overlap together as
Fig. 9. Bode plots for aged specimens with (a) 4.3 wt% Mn, (b) 6.9 wt% Mn, (c) 9.7 wt% Mn contents in 3.5 wt% NaCl solution.
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Fig. 10. Equivalent electric circuit tested to model the experimental EIS data. Rs is solution resistance, Rt is charge transfer resistance, QPEt is the capacitive at the interfaces, Rfilm is oxide film resistance and QPEfilm is the capacitive of the oxide film.
a horizontal region at low and middle frequencies. This can be attributed to a compact layer spontaneous passive film growth in the early aging period. However, the frequency region of the horizontal part becomes narrower by increasing the aging time to 930 min. These indicate that passive film defects may be generated around s precipitates to weaken the passive state [39]. When the Mn content increases from 6.9 to 9.7 wt% (Fig. 9bec), no obvious changes are observed for the phase angle and frequency region of the horizontal part with aging treatment from 60 to 930 min, indicating that the capacitance feature of the passive film was retained, despite the accelerating effect of s precipitates at higher Mn contents. Meanwhile, the frequency region of the horizontal part decreased slowly with aging time from 15 to 930 min in the 9.7 wt% Mn content specimen, indicating that the passive ability was weakened, which is associated with the formation of more Cr-rich s precipitates caused by d-ferrite decomposition. Different models were proposed to interpret impedance spectra on passive metal surfaces [40e42]. Two regions corresponding to two time constants can be distinguished in Fig. 9, thus the electrode system stainless steel/NaCl solution for different aged specimens can be schematically illustrated in Fig. 10. In this circuit, Rs is the electrolyte resistance, Rt is the charge transfer resistance and QPEt is the capacitance of the corrosion product layer, including the defects that were caused by the formation of ionically conducting paths across the s precipitates. Rfilm is the resistance of the passive film and QPEfilm is the capacitance of the passive layer within d/g phases. All the samples with different Mn contents are spontaneously passive at the OCP, thus, this equivalent circuit is suitable for revealing the corrosion mechanism. This model assumes that the passive film does not totally recover the metal and can be treated as a defective layer [40,41]. The inhomogeneous degree on metallic substrates was related to precipitation initiation and growth rate with prolonging aging time. Table 4 lists the best fitting electrical parameters based on the
Fig. 11. Rp resistance from the fitting procedure for samples with different Mn contents aged at 800 C tested in 3.5 wt% NaCl solution.
circuit in Fig. 10. The Rt of the aged specimens is small and decreases with increasing Mn content from 4.3 to 6.9 wt% and 9.7 wt %, which indicate that more defects in the passive film appear as more s precipitates are formed with higher Mn content. The resistance of Rt depends strongly on the existence of defects, as more Cr-depleted zone around s-precipitates increases the number of defects in the d-ferrite matrix and along d/g boundaries. For different specimens, the oxide film resistance (Rfilm) decreased gradually with prolonging aging time from 15 to 930 min, while increasing Mn content slightly decreased Rfilm in specimens aged from 60 to 930 min. This indicates that the electrical conductivity of the films increases as more s precipitates are formed. Polarisation resistance, Rp (Rp ¼ Rt þ Rfilm), is commonly used as a measure of the resistance of a metal to corrosion damage [39]. The Rp values with different aging times and Mn contents are shown in Fig. 11. At the initial stage of 15 min, the Rp value for 4.3 wt% Mn is obviously higher than those for 6.9 and 9.7 wt% Mn, showing better corrosion resistance. As the aging time increased to 60 min, Rp decreases similarly regardless of Mn content, which flattens out by prolonging the aging time to 930 min in spite of a slightly higher Rp in the 4.3 wt% Mn sample. This suggested the corrosion resistance of the studied alloys is less likely to be affected after s phase precipitation by long aging time treatment with different Mn additions.
Table 4 Best fitting parameters for impedance spectra after stabilisation at the OCP. Aging time and Mn content
Rs (U cm2)
QPEt (U1 cm2 sn)
n1
Rt (U cm2)
QPEfilm (U1 cm2 sn)
n2
Rfilm (U cm2)
15 min, 4.3 wt% Mn 60 min, 4.3 wt% Mn 150 min, 4.3 wt% Mn 450 min, 4.3 wt% Mn 930 min, 4.3 wt% Mn 15 min, 6.9 wt% Mn 60 min, 6.9 wt% Mn 150 min, 6.9 wt% Mn 450 min, 6.9 wt% Mn 930 min, 6.9 wt% Mn 15 min, 9.7 wt% Mn 60 min, 9.7 wt% Mn 150 min, 9.7 wt% Mn 450 min, 9.7 wt% Mn 930 min, 9.7 wt% Mn
9.94 10.42 9.559 9.058 9.499 9.906 6.735 9.257 10.32 10.22 10.41 9.691 11.07 9.126 8.997
2.0655E-6 1.4708E-5 1.6285E-5 1.5012E-5 1.4126E-5 1.3678E-5 1.1866E-5 1.4675E-5 1.7982E-5 2.005E-5 7.6361E-6 9.6485E-6 1.5993E-5 1.6984E-5 2.2031E-5
0.81315 0.90643 0.89617 1.027 1.06 1.048 1.022 1.053 1.055 1.035 0.95033 0.96959 1.058 1.025 0.92285
336.3 50.65 26.46 7.027 8.512 8.254 5.598 8.291 10.01 11.26 17.74 9.719 10.65 11.5 1.6413
1.4706E-7 5.4361E-6 6.6551E-6 6.9093E-6 5.9946E-6 7.7887E-6 6.5796E-6 5.9275E-6 5.5146E-6 5.6385E-6 1.4748E-6 3.0281E-6 5.6474E-6 5.3235E-6 2.376E-11
0.78205 1.112 1.099 0.85444 0.85367 0.86778 0.83991 0.84406 0.84309 0.83036 0.79981 0.83539 0.84878 0.82978 0.46624
1412200 123450 92291 109760 149390 126720 105430 94420 88993 71886 238790 140730 115580 111940 77582
Z. Feng et al. / Journal of Alloys and Compounds 699 (2017) 334e344
4. Conclusions The influence of Mn addition on the microstructural and corrosion behaviour of aged 22 Cr economical DSSs can be concluded as follows: (1) The s precipitates nucleated at g/d interfaces and grew into the adjacent d-ferrite with increasing aging time from 15 to 930 min. By adding Mn in concentrations from 4.3 wt% to 9.7 wt%, s precipitates were formed more rapidly in the early stage of aging time from 30 to 150 min, but decreased gradually up to 930 min. The acceleration of s phase precipitation with increasing Mn can be attributed to the partition of the Mn element. (2) For the 4.3 wt% Mn content specimen, only about a 150 mV Eb drop occurred with increasing aging time from 15 to 930 min in 3.5 wt% NaCl solution, but Eb dropped significantly at the early aging stage with higher Mn contents of 6.9 and 9.7 wt% due to the formation of more aging-induced s precipitates. The metastable pit nucleation resistance decreased with Mn content from 4.3 to 6.9 wt% and was suppressed by a higher Mn addition of 9.7 wt% for aging treatment from 150 to 930 min. (3) In the initial aging stage of 60 min, the pit initiation sites occurred mainly along d/g boundaries and within d-ferrite due to a low PREN value of ferrite, which gradually occurred around s precipitates up to 930 min due to lower PREN value of the g2 product phase. Higher Mn content increased both the number and size of pits in the early stage of aging treatment due to the effect of Mn in increasing s precipitates. (4) EIS measurements showed that the charge transfer resistance, Rt, depends strongly on the existence of defects caused by s precipitates. Both Mn content and aging time were observed to affect the parameters associated with the passive film. In the initial aging stage of 15 min, the value of Rfilm in the sample of 4.3 wt% Mn is significantly higher than that of 6.9 and 9.7 wt% Mn, whereas higher Mn contents slightly decreased the value of Rfilm in the specimens aged from 60 to 930 min. Acknowledgements This work was mainly supported by the Fundamental Research Funds for the National Natural Science Foundation of China (No. 51261010). References [1] K.M. Lee, H.S. Cho, D.C. Choi, Effect of isothermal treatment of SAF 2205 duplex stainless steel on migration of delta/gamma interface boundary and growth of austenite, J. Alloy. Compd. 285 (1999) 156e161. [2] Y. Yang, B. Yan, The microstructure and flow behavior of 2205 duplex stainless steels during high temperature compression deformation, Mater. Sci. Eng. A 579 (2013) 194e201. [3] J. Li, Z. Ma, X. Xiao, J. Zhao, L. Jiang, On the behavior of nitrogen in a low-Ni high-Mn super duplex stainless steel, Mater. Des. 32 (2011) 2199e2205. [4] Y.H. Jang, S.S. Kim, J.H. Lee, Effect of different Mn contents on tensile and corrosion behavior of CD4MCU cast duplex stainless steels, Mater. Sci. Eng. A 396 (2005) 302e310. [5] I.-U.-H. Toor, P.J. Hyun, H.S. Kwon, Development of high MneN duplex stainless steel for automobile structural components, Corros. Sci. 50 (2008) 404e410. [6] R.F.A. Jargelius-Pettersson, Application of the pitting resistance equivalent concept to some highly alloyed austenitic stainless steels, Corrosion 54 (1998) 162e168. [7] C.J. Park, H.S. Kwon, Effects of aging at 475 C on corrosion properties of tungsten-containing duplex stainless steels, Corros. Sci. 44 (2002) 2817e2830. [8] E. Angelini, B. Benedetti, F. Rosalbino, Microstructural evolution and localized corrosion resistance of an aged superduplex stainless steel, Corros. Sci. 46 (2004) 1351e1367.
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Wu, Precipitation behavior of s phase in 19Cre9Nie2Mn and 18Cre0.75 Si stainless steels hot-rolled at 800 C with various reduction ratios, Mater. Sci. Eng. A 467 (2007) 181e189. [14] I. Zucato, M.C. Moreira, I.F. Machado, S.M.G. Lebrao, Microstructural characterization and the effect of phase transformations on toughness of the UNS S31803 duplex stainless steel aged treated at 850 C, Mater. Res. 5 (2002) 385e389. [15] T.H. Chen, J.R. Yang, Effects of solution treatment and continuous cooling on s-phase precipitation in a 2205 duplex stainless steel, Mater. Sci. Eng. A 311 (2001) 28e41. [16] ASTME 562 standard practice for Determining Volume Fraction by Systematic Manual Point Count. [17] T.H. Chen, K.L. Weng, J.R. Yang, The effect of high-temperature exposure on the microstructural stability and toughness property in a 2205 duplex stainless steel, Mater. Sci. Eng. A 338 (2002) 259e270. [18] C.S. Huang, C.C. Shih, Effects of nitrogen and high temperature aging on s phase precipitation of duplex stainless steel, Mater. Sci. Eng. A 402 (2005) 66e75. [19] Y.H. Yang, B. Yan, J. Wang, J.L. Yin, The influence of solution treatment temperature on microstructure and corrosion behavior of high temperature ageing in 25% Cr duplex stainless steel, J. Alloy. Compd. 509 (2011) 8870e8879. [20] Y. Maehara, Y. Ohmori, J. Murayama, N. Fujino, T. Kunitake, Effects of alloying elements on s- phase precipitation in d-g duplex phase stainless steels, Metal. Sci. 17 (1983) 541e547. €m, Sigma phase precipitation in duplex stainless steel [21] H. Sieurin, R. Sandstro 2205, Mater. Sci. Eng. A 444 (2007) 271e276. [22] H.L. Yakel, Atom distributions in sigma phases. II. Estimations of average siteoccupation parameters in a sigma phase containing Fe, Cr, Ni, Mo and Mn, Acta Crystallogr. 39 (1983) 28e33. [23] S.-B. Kim, K.-W. Paik, Y.-G. Kim, Effect of Mo substitution by W on high temperature embrittlement characteristics in duplex stainless steels, Mater. Sci. Eng. A 247 (1998) 67e74. [24] N. Lopez, M. Cid, M. Puiggali, Influence of O-phase on mechanical properties and corrosion resistance of duplex stainless steels, Corros. Sci. 41 (1999) 1615e1631. [25] M. Martins, L.C. Casteletti, Sigma phase morphologies in cast and aged super duplex stainless steel, Mater. Charact. 60 (2009) 792e795. [26] Z. Zhang, H. Zhang, D. Han, L. He, Y. Jiang, J. Li, Precipitation evolution in duplex stainless steel during isothermal aging at 700 C, Mater. Sci. Technol. 30 (2014) 451e457. [27] M.A. Domínguez-Aguilar, R.C. Newman, Detection of deleterious phases in duplex stainless steel by weak galvanostatic polarization in alkaline solution, Corros. Sci. 48 (2006) 2560e2576. [28] K.M. Adhe, V. Kain, K. Madangopal, H.S. Gadiyar, Influence of sigma-phase formation on the localized corrosion behavior of a duplex stainless steel, J. Mater. Eng. Perform. 5 (1996) 500e506. [29] M. Kaneko, H.S. Isaacs, Pitting of stainless steel in bromide, chloride and bromide/chloride solutions, Corros. Sci. 42 (2000) 67e78. [30] J. Hu, K. Tian, W.Y. Chu, Electrochemical corrosion behavior of Al18B4O33w/Al composite, J. Mater. Sci. 40 (2005) 5147e5151. [31] S.M. Alvarez, A. Bautista, F. Velasco, Influence of process parameters on the corrosion resistance of corrugated austenitic and duplex stainless steels, Mater. Tech. 47 (2013) 317e321. [32] J.O. Nilsson, P. Kangas, T. Karlsson, A. Wilson, Mechanical properties, microstructural stability and kinetics of s-phase formation in 29Cre6Nie2Moe0.38N superduplex stainless steel, Metall. Mater. Trans. A 31A (2000) 35e45. [33] Y.H. Yang, B. Yan, J. Li, J. Wang, The effect of large heat input on the microstructure and corrosion behaviour of simulated heat affected zone in 2205 duplex stainless steel, Corros. Sci. 53 (2011) 3756e3763. [34] V. Vignal, N. Mary, R. Oltra, J. Peultier, A mechanicaleelectrochemical approach for the determination of precursor sites for pitting corrosion at the microscale, J. Electrochem. Soc. 153 (2006) B352eB357. [35] L. Zhang, Y. Jiang, B. Deng, W. Zhang, J. Xu, J. Li, Effect of aging on the corrosion resistance of 2101 lean duplex stainless steel, Mater. Charact. 60 (2009) 1522e1528. [36] G. Rondelli, B. Vicentini, A. Cigada, Influence of nitrogen and manganese on localized corrosion behaviour of stainless steels in chloride environments, Mater. Corros. 46 (1995) 628e632. [37] H. Luo, C.F. Dong, X.G. Li, K. Xiao, The electrochemical behaviour of 2205 duplex stainless steel in alkaline solutions with different pH in the presence of chloride, Electrochim. Acta 64 (2012) 211e220.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of Mn addition on the precipitation and corrosion behaviour of 22% Cr economical duplex stainless steel after isothermal aging at 800 C Zaiqiang Feng a, Yinhui Yang b, *, Jia Wang c a b c
School of Mechanical Engineering, North China University of Water Resources and Electric Power, Zhengzhou 450045, PR China School of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, PR China School of Materials Science and Engineering, Tongji University, Shanghai 200092, PR China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 February 2016 Received in revised form 22 December 2016 Accepted 3 January 2017 Available online 4 January 2017
The microstructural and corrosion behaviour of Fe-22Cr-1.9Ni-2.3Mo-0.2N-xMn low nickel type duplex stainless steel (DSS) with Mn contents ranging from 4.3 to 9.7 wt% were studied after aging treatment at 800 C lasting from 30 to 930 min. The corrosion behaviour was investigated in 3.5 wt% NaCl solution by measurements of potentiodynamic polarisation and electrochemical impedance spectroscopy. The microstructural analysis shows that the s precipitates mainly nucleated along g/d phase boundaries and grew into the adjacent d-ferrite by prolonging the aging time, having the same eutectoid structures as commercial high-Ni DSS. At high Mn contents, the growth rate of s precipitates was initially high with aging times from 30 to 150 min, but decreased gradually up to 930 min with more Mn participation. A Mn content of 4.3 wt% brought only about a 150 mV pitting potential (Eb) drop for the aged specimens, while further addition of Mn significantly lowered Eb at the early stage of aging with more aging-induced s precipitates. The metastable pit nucleation resistance decreased with Mn content from 4.3 to 6.9 wt% and was suppressed by a higher Mn addition of 9.7 wt% for aging treatments from 150 to 930 min. The pits mainly initiated along d/g boundaries and within d ferrite, then propagated around s precipitates with increasing aging time, and more Mn gives rise to both the number and size of pits in the early stage of aging treatment. The charge transfer resistance (Rt) depends strongly on the existence of defects caused by the s precipitates, while increasing Mn content slightly decreases the passive film resistance (Rfilm) in specimens aged from 60 to 930 min. © 2017 Elsevier B.V. All rights reserved.
Keywords: Metals and alloys Intermetallics Microstructure Precipitation Corrosion
1. Introduction Duplex stainless steels (DSSs) are widely employed in many fields, such as the oil, petrochemical, nuclear and desalination industries, with a two-phase structure of ferrite and austenite, which takes advantage of their attractive combination of good corrosion resistance and mechanical properties under harsh corrosive environments [1,2]. However, the increasing price of Ni has made the cost of high-Ni-content DSS prohibitive for many industrial applications. Therefore, it is urgent to find proper alternatives for Ni. For instance, Mn-N DSSs are being actively developed by replacing Ni in high-Ni DSSs with cheaper Mn and N. Both Mn and N are known to
* Corresponding author. E-mail address: [email protected] (Y. Yang). http://dx.doi.org/10.1016/j.jallcom.2017.01.031 0925-8388/© 2017 Elsevier B.V. All rights reserved.
be good austenite stabilisers, which enhance the strengthening effect and improve the corrosion resistance of steels [3]. Mn, which is almost an order of magnitude cheaper than Ni at an equivalent weight [4,5], can further help increase the solubility of N in stainless steels [6], whereas the low solubility of N in liquid steel under atmospheric pressure has been problematic for production. DSSs are frequently exposed to high temperatures during hot working processes, which easily induce intermetallic phase precipitation. For example, Cr- and Mo-rich phases, such as Cr23C6 carbide and c and s phases, are easily formed when DSSs are subjected to elevated temperatures of 450e1000 C [7,8]. These precipitates often cause a loss of toughness and corrosion resistance. Moreover, previous studies noted that the detrimental s phase forms most rapidly at 800 C [9,10]. Johnson et al. [11] also reported that the initial nose of s phase formation in the timetemperature-transformation diagram of DSS occurs earliest at
Z. Feng et al. / Journal of Alloys and Compounds 699 (2017) 334e344
800 C. Thus, it is effective to investigate precipitation behaviour with different Mn additions under severe conditions of an 800 C aging temperature. The Cr-depleted zones in the vicinity of s precipitates are susceptible to pitting corrosion due to weakened passivation. Zhang et al. [12] investigated the effect of isothermal aging on the pitting corrosion resistance of UNS S82441, and found pitting transfers from the austenite phase into Cr-depleted zones with increasing aging time. It has been demonstrated that adding Cr and Mo can accelerate the precipitation of the s phase at high temperatures [13], while the partition of these alloy elements in two phases can be altered in the presence of Mn. Therefore, during aging treatment of 800 C, different Mn contents can lead to various degrees of precipitate transformations, which take place within the ferrite matrix and along grain boundaries due to the much faster diffusion rates of the alloying elements within ferrite compared to austenite [14,15]. It is known that the detrimental effect of Mn on the pitting corrosion resistance is associated with nonmetallic MnS inclusions in high-S alloys. The reported effects of Mn on the pitting corrosion
335
resistance are controversial. Toor reported that the resistance to pitting and metastable pitting corrosion in a novel high-Mn Ni-free DSS decreased with increasing Mn content since a great number of (Mn, Cr) oxides acted as preferential sites of pitting [5]. Jang et al. believed that the tensile properties of CD4MCU0-cast DSS were determined by the volume fraction of hard ferritic phase and the shape of austenitic phase, and the resistance to pitting corrosion and to stress corrosion cracking in 3.5% NaCl þ5% H2SO4 aqueous solution was recovered with Mn contents from 0.8 to 2% [4]. However, there is little research available regarding the influence of Mn on the corrosion behaviour of DSSs caused by intermetallic phase precipitation during high temperature aging treatment. This study selects Mn content as the main parameter and investigates its effect on the microstructure evolution and corrosion behaviour of low nickel DSS by isothermal aging ranging from 15 to 930 min at 800 C. The relationships between the aginginduced precipitation phases under various Mn contents and pitting corrosion behaviour of the newly-developed DSS are discussed in detail.
Table 1 Chemical composition of hot rolled 19% Cr DSS (wt.%). Elements
C
Si
Mn
S
P
Cr
Ni
Mo
Cu
N
Other
Alloy 1 Alloy 2 Alloy 3
0.01 0.01 0.01
0.12 0.11 0.14
4.28 6.90 9.71
0.005 0.004 0.005
0.007 0.006 0.006
22.47 22.49 22.31
1.94 1.93 1.88
2.33 2.25 2.25
0.21 0.22 0.23
0.21 0.22 0.23
Bal.
Fig. 1. Optical images of the specimens aged at 800 C for: (a) 4.3% Mn, 15 min; (b) 6.9% Mn, 15 min; (c) 9.7% Mn, 15 min; (d) 4.3% Mn, 150 min; (e) 6.9% Mn, 150 min; (f) 9.7% Mn, 150 min; (g) 4.3% Mn, 930 min; (h) 6.9% Mn, 930 min; (i) 9.7% Mn, 930 min.
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2. Experimental 2.1. Materials and heat treatment The experimental steels were prepared by induction melting in a 25 kg vacuum induction furnace, followed by casting into ingots. The cast ingots were hot forged into 30 mm plates, then, the asforged samples were hot rolled into 12 mm plates at temperatures ranging from 1050 to 1180 C. The chemical compositions of the rolling plates, designated as Alloys 1, 2 and 3, are shown in Table 1, of which the Mn contents are 4.3, 6.9 and 9.7 wt%, respectively. The as-rolled plates were machined to the size of 12 12 50 mm and solution-treated at 1040 C for 30 min. In order to study the precipitation behaviour under different Mn contents, all the solution-treated specimens were then subject to aging treatment at 800 C for times varying from 15 to 930 min, followed by water quenching.
amount of austenite phases increased marginally with Mn content increasing from 4.3% to 9.7%. At the early aging stage of 15 min (Fig. 1aec), no obvious secondary phase appeared in all the specimens. However, with a longer aging time of 150 min (Fig. 1def), some dark granular precipitates nucleated at g/d interfaces in specimens with 4.3% and 6.9% Mn, while more precipitates formed
2.2. Microstructural characterisation The microstructure of the aged specimens was revealed by electrochemically etching in 40 wt% KOH solution for 10e15 s at room temperature. The chemical compositions of the aginginduced precipitates and g and d phases were measured by energy dispersive X-ray spectroscopy (EDS) attached to a scanning electron microscope (SEM, FEI Quanta 200) with a Robinson backscattered electron detector. In addition, X-ray diffraction (XRD) was used for identification of the precipitates. The volume fractions of the precipitates were measured using the method of manual point count according to ASTM E 562 [16]. 2.3. Electrochemical measurements A Bio-Logic VMP-300 potentiostat was used to perform all the electrochemical measurements in a three-electrode cell. A platinum sheet and a saturated calomel electrode (SCE) were used as the counter and reference electrodes, respectively. Electrochemical potentiodynamic polarisation was performed in a deaerated 3.5 wt % NaCl solution at a temperature around 25 C (±1 C). The specimens, embedded in epoxy resin with an exposure area of 100 mm2, acted as working electrodes. Prior to each experiment, the specimens were ground mechanically up to 3000 grit, rinsed with distilled water and dried in hot air. All potentials are given against the SCE. Then, the corrosion morphology was observed by SEM. The polarisation curves were recorded potentiodynamically with a scan rate of 0.1 mV s1, starting from the free corrosion potential to transpassive potential (vs. SCE). When the passive film breaks down with a rapid rise in the current density, Eb was determined as the current density reached a value of 100 mA/cm2. EIS measurements were executed at open circuit potential (OCP) in a deaerated 3.5 wt% NaCl solution. The frequency was swept from 100 kHz down to 10 mHz at 10 data cycles/decade, with an amplitude of sinusoidal wave of 10 mV. The impedance at the OCP was carried out after immersion at OCP for 30 min. The data were interpreted on the basis of equivalent electrical circuits using the Zview 2.70 program to fit the experimental data. 3. Results and discussion 3.1. Microstructural analysis The microstructural evolution of specimens with different Mn contents during isothermal aging at 800 C is shown in Fig. 1. The bright etched austenite (g) islands were formed of a lath-like morphology within the grey etched ferrite (d) matrix, and the
Fig. 2. XRD results of the specimens aged at 800 C for: (a) 4.3% Mn, 930 min; (b) 6.9% Mn, 930 min; (c) 9.7% Mn, 930 min.
Z. Feng et al. / Journal of Alloys and Compounds 699 (2017) 334e344
337
Fig. 3. SEM micrographs of specimens aged at 800 C for: (a) 4.3% Mn, 630 min; (b) 6.9% Mn, 630 min; (c) 9.7% Mn, 630 min (d) EDX spectrum of s-phase corresponding to (b).
along g/d interfaces and grew into the adjacent d-ferrite regions. These dark secondary precipitates can be identified as the s phase due to its heterogeneous precipitation with the formation of bright secondary austenite (g2) during high temperature aging of 700e900 C [17,18]. As the aging reached 930 min, it is noted that the coarser dark s phase nucleated at g/d interfaces in the specimen with 4.3% Mn (Fig. 1g), whereas the s precipitates increased both in number and size with higher Mn content (Fig. 1hei), which is attributed to the enhanced diffusion of solute elements with higher Mn content. The XRD results of the specimens are presented in Fig. 2. In addition to the major peaks from d-ferrite and austenite, the minor diffraction peaks from the s phase were identified in specimens of different Mn contents, proving the presence of the s phase in the aged specimens. It is also observed that the intensity of s (411) and s (112) peaks increased with adding Mn from 4.3% to 6.9% and up to 9.7%, corresponding to the reduction of the d (1 1 0) peak intensity as a result of the d / s þ g2 transformation [19]. This suggested that the precipitation was accelerated by adding more Mn. Fig. 3 presents the SEM images for specimens aged after 630 min with different Mn contents. For the specimen with 4.3% Mn content, some s particles precipitated at d/g interfaces and penetrated into the d-ferrite matrix. However, when the Mn content was 6.9% and 9.7% (Fig. 3b and c), the lamellar structure, consisting of s and g2 phases, can be clearly observed, the same eutectoid structures were also obtained in commercial DSSs aged at high temperature in spite of very low Mn contents [20]. As shown in Fig. 3d, these s phases were further confirmed through EDX analysis, featuring high Cr and low Mo contents [15]. In addition, it is detected that the s precipitate was partitioned with a higher Mn content than dferrite matrix in the specimen with 6.9% Mn, suggesting that the s
precipitates have a closer affinity to Mn [21,22]. The variations of s phase precipitation at aging of 800 C with different Mn contents are shown in Fig. 4. It is observed that the volume fraction of s phase increased with increasing Mn content during aging time from 15 to 930 min, indicating quicker formation of s precipitates at high Mn content. At early stages of aging time from 30 to 150 min, it is inferred from the slope of the curve that the rate of s phase precipitation in 9.7 wt% Mn content samples was
Fig. 4. Variation of s phase volume fraction as a function of aging time at 800 C in samples with different Mn contents.
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higher than that in the 6.9 and 4.3 wt% Mn specimens. This can be attributed to higher diffusivity of Mn, which is similar to the effect of Mo on the nucleation of the s phase in DSS [23]. However, with the aging time finally increased to 930 min, the amount of s phase in 6.9 wt% Mn Alloy 2 rapidly increased and nearly became saturated, while in the 9.7 wt% Mn Alloy 3, it increased rather slowly due to complete decomposition of d-ferrite. As a consequence, the
Table 2 Pitting and passivation values of the materials tested in 3.5% NaCl solution at different Mn contents. Heat treatment condition
Eb (mV vs. SCE)
ip (mA cm2)
icorr (mA cm2)
Ecorr (mV vs. SCE)
Aging at 800 C 4.3 wt% Mn Aging at 800 C 4.3 wt% Mn Aging at 800 C 4.3 wt% Mn Aging at 800 C 4.3 wt% Mn Aging at 800 C 4.3 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 6.9 wt% Mn Aging at 800 C 9.7 wt% Mn Aging at 800 C 9.7 wt% Mn Aging at 800 C 9.7 wt% Mn Aging at 800 C 9.7 wt% Mn Aging at 800 C 9.7 wt% Mn
for 15 min,
591
2.04
1.90
434
for 60 min,
571
2.29
1.95
633
for 150 min,
513
2.69
1.92
644
for 450 min,
445
1.66
1.39
495
for 930 min,
435
1.67
1.45
539
for 15 min,
521
2.29
0.96
606
for 60 min,
429
2.14
1.81
529
for 150 min,
269
2.15
1.84
532
for 450 min,
267
2.04
1.71
549
for 930 min,
234
2.14
1.84
607
for 15 min,
439
2.69
2.62
654
for 60 min,
252
2.19
2.61
656
for 150 min,
154
2.63
1.40
712
for 450 min,
216
2.18
1.54
654
for 930 min,
169
3.47
3.64
771
driving force for the precipitation of s phases tends to be strong with higher Mn contents. 3.2. Potentiodynamic measurements and pitting corrosion resistance analysis Fig. 5 shows the potentiodynamic polarisation curves of all specimens obtained for aging times of 15e930 min in 3.5 wt% NaCl solution at 25 C. As the sweep potential increased to the pitting potential (Eb), the sharp increase in current density was associated with an occurrence of stable pits, indicating the breakdown of the
Fig. 5. Potentiodynamic polarisation curves of samples with different Mn contents in 3.5 wt% NaCl solution aged at 800 C from 15 to 930 min: (a) Alloy1 with 4.3 wt% Mn; (b) Alloy2 with 6.9 wt% Mn; (c) Alloy3 with 9.7 wt% Mn.
Fig. 6. Difference of Eb and Ecorr for different Mn addition samples aged at 800 C tested in 3.5 wt% NaCl solution.
Z. Feng et al. / Journal of Alloys and Compounds 699 (2017) 334e344
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Fig. 7. SEM-BSE images of specimens aged at 700 C after potentiodynamic polarisation testing for: (a) 4.3 wt% Mn, aged for 60 min; (b) 6.9 wt% Mn, aged for 60 min; (c) 9.7 wt% Mn, aged for 60 min; (d) 4.3 wt% Mn, aged for 930 min; (e) 6.9 wt% Mn, aged for 930 min; (f) 9.7 wt% Mn, aged for 930 min.
passive film. Alloy 1 with 4.3 wt% Mn displays similar polarisation curves throughout all the aging stages, showing a wide potential domain of passivity in the potential range between 0.4 and 0.45 VSCE (Fig. 5a). The Eb value decreased slightly with increasing aging time from 15 to 930 min, owing to the formation of more Crdepleted zones caused by s precipitates with longer aging time [24,25]. Moreover, at high Mn content, in addition to negative shifting of Eb, the difference of Eb between specimens aging for 15 and 930 min was enlarged (Fig. 5b and c), which can be attributed to the ripening of Cr-depleted zones with rapid coalescence of s phases [26]. This was consistent with the above mentioned microstructural observations and discussion. The results of the polarisation measurements are summarised in Table 2. For specimens with 4.3 wt% Mn, the corrosion potentials (Ecorr) decreased with increasing aging time from 15 to 60 and 150 min, while it showed small changes with further aging time to
930 min, this can be explained by the nucleation of precipitates in the initial aging stages. The variation in Ecorr of most samples with 6.9 and 9.7 wt% was small, with their values being around 600 mVSCE. However, Ecorr decreased to 771 mVSCE in the 9.7 wt% Mn specimen with 930 min aging due to the formation of more s phases, and the corresponding corrosion current density, icorr, increased by about 2 mA cm2 compared with lower Mn addition specimens at this aging time. This suggests that the addition of Mn has less of an effect on (Ecorr) SCE of aged specimens than on the low Mn content commercial DSS reported earlier [27]. The passivation current densities, ip, are similar regardless of Mn content and aging conditions, indicating good passivation stability for different aged samples in the presence of Cl. However, the values of Eb exhibited significant differences among samples of different Mn contents. For the specimen with 4.3 wt% Mn, Eb decreased gradually by about 150 mV with increasing aging time from 15 to 930 min, showing
Table 3 Mean compositions (wt.%) and PRENs of g, d and g2 in samples with different Mn contents subject to aging at 800 C for 630 min. Material
Phase
Cr
Mn
Mo
Ni
Na
PRENb
4.3 wt% Mn, aging for 630 min
Austenite (g) Ferrite (d) Secondary austenite (g2) s phase Austenite (g) Ferrite Secondary austenite (g2) s phase Austenite (g) Ferrite (d) Secondary austenite (g2) s phase
21.4 23.1 19.5 25.7 21.2 23.5 20.1 27.1 19.6 22.4 19.6 25.8
5.0 3.9 5.1 5.3 7.9 5.4 7.6 7.5 10.2 8.4 9.5 9.9
1.9 2.5 1.2 2.0 1.8 2.6 1.8 2.2 1.9 2.9 1.8 2.6
2.5 1.5 2.9 1.3 2.4 1.2 2.2 1.1 2.5 1.6 1.5 1.5
0.37 0.05 0.10
33.77 28.95 21.36
0.39 0.05 0.10
30.94 28.18 21.44
0.41 0.05 0.10
27.97 25.07 19.04
6.9 wt% Mn, aging for 630 min
9.7 wt% Mn, aging for 630 min
a b
Nitrogen in ferrite is taken as the saturation value z 0.05%, the rest partitions to austenite [35]. PRENs were calculated with the following equation: PREN ¼ %Cr þ 3.3%Mo þ 30%N1 %Mn.
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small pitting corrosion resistance reduction in the aging time range. With higher Mn content (6.9 wt%), Eb decreased quickly after 60 min aging time and flattened out at around 250 mV until 930 min. In addition, Eb was shifted to a more negative value in the 9.7 wt% Mn sample at the equivalent aging time than that for the 6.9 wt% Mn sample, indicating that higher Mn contents enhance the susceptibility of the aged specimens to pitting corrosion. It is noted that Mn promoted the formation of s precipitates to some degree in aged specimens from the above microstructural analysis. Meanwhile, the s precipitates play an important role in decreasing localised corrosion resistance because of the relatively largevolume Cr-depleted zones surrounding the particles [26]. Thus, a close relationship is observed between the decrease in pitting corrosion resistance of aged samples and the increase of the amount of s phase, indicating that the amount of s precipitates influenced by varying the Mn content plays a critical role in the propagation of the initial pits [28]. The difference between Eb and Ecorr is a measure of the extent of the passive region on the polarisation curve and provides an indication of the susceptibility to pitting and the metastable pits occurring in this region [29e31]. The resistance of the material to metastable pit nucleation increases when this difference increases. The values of (Eb - Ecorr)SCE from samples of different Mn contents at 800 C aging are shown in Fig. 6. The values were similar in all samples at the initial aging of 15 min, then decreased with Mn content from 4.3 to 6.9 wt% and up to 9.7 wt% as the aging time was further prolonged to 930 min, indicating that the metastable pit nucleation resistance decreased with higher Mn content for aging treatment at this temperature. However, as the Mn content increased from 6.9 to 9.7 wt%, the values of (Eb - Ecorr)SCE increased instead with aging time from 150 to 930 min, meaning the metastable pit nucleation was suppressed to some extent. Fig. 7 shows SEM-BSE images of the corroded surfaces of the aged specimens at 800 C for 60 and 330 min after potentiodynamic polarisation tests performed in 3.5 wt% NaCl solution at 25 C. As shown in the SEM surface morphology from 60 min-aged specimens with 4.3 and 6.9 wt% Mn (Fig. 7a and b), the pits initiated not only along d/g phase boundaries and d/d grain boundaries, but also within the d ferrite matrix. This could be explained by the difference in the chemical composition, namely the pitting resistance equivalent (PREN) of the two phases [32]. With a high Mn content of 9.7 wt% (Fig. 7c), an incomplete eutectoid decomposition of the d ferrite (d / s þ g2) induced more pits around the s phases. This indicates that Mn may reduce pitting corrosion resistance in the initial stage of aging. With longer aging treatment up to 930 min, the back-scattered images (Fig. 7def) revealed the areas attacked by the pits are active black regions around s precipitates, which have lower Cr contents in g2 than that in primary g and provide preferential sites for pitting attacks [12]. These black pit regions were propagated with increasing Mn addition from 4.3 to 6.9 wt% and up to 9.7 wt% due to the formation of coarser s phases. These s precipitates made their adjacent regions depleted in Cr and thus incoherent with the matrix, triggering and lowering the pitting potential (Epit)SCE in the Cl-based medium [33,34]. The mean compositions and PRENs of g, d and g2 in specimens with different Mn contents subjected to aging at 800 C for 630 min are listed in Table 3. It shows that the PREN of ferrite was lower than that of the austenite phase in all the specimens, meaning that the primary sites for pit nucleation were in ferrite. Meanwhile, the secondary austenite (g2) has a much lower value of PREN than both the austenite and ferrite phases, indicating that the formation of more g2 accompanied by the nucleation of s and the decomposition of d ferrite (d / s þ g2), inducing more pits around s phases, in
agreement with the post-potentiodynamic polarisation corrosion morphology analysis. It is also concluded from the table that the PREN of d-ferrite decreased with adding Mn from 4.3 to 9.7 wt% due to the effect of Mn on the pitting corrosion [36], indicating that the ferrite was more susceptible to pit attacks in the presence of high Mn content at the early stage of aging. Furthermore, it is noted that the s phase precipitation was partitioned with Mn by EDS analysis, suggesting the addition of Mn accelerated the formation of s precipitates after aging for 90 min, which mostly triggered pitting corrosion in d ferrite.
Fig. 8. Nyquist plots for aged specimens tested in 3.5 wt% NaCl solution with different Mn additions: (a) Alloy1 with 4.3 wt% Mn addition; (b) Alloy2 with 6.9 wt% Mn addition; (c) Alloy3 with 9.7 wt% Mn addition.
Z. Feng et al. / Journal of Alloys and Compounds 699 (2017) 334e344
3.3. EIS measurements EIS measurements were performed to determine the effects of Mn on the electrochemical properties of the passive film aged at 800 C. Fig. 8 presents the Nyquist spectra obtained at the OCP of different aging times with different Mn contents. Impedance spectra are normally displayed in the form of a Nyquist diagram, where the opposite of the imaginary part of impedance is plotted vs. the real part. The diagrams display a capacitive arc that does not end at low frequency, and the radius of the semi-circular arc is related to the polarisation resistance of the passive film [37]. The radius of the capacitive arc decreases with increasing aging time, indicating that the corrosion
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resistance was lowered with weakening of the passive ability due to s phase precipitation. With the Mn content increased from 4.3 to 9.7 wt% (Fig. 8aec), it is shown that the impedance value decreases at the initial aging stage of 15 min, then flattens out by prolonging the aging time to 930 min, revealing a high sensitivity to aging precipitation. The impedance responses for differently aged specimens are presented in Bode format in Fig. 9. They show time constants at high and low frequencies. The shape of the curves is usually associated with the response of an inhomogeneous film [38]. For the specimen with 4.3 wt% Mn (Fig. 9a), the phase angle q is almost invariable around 70 for 15 min, where the time constants from charge transfer process and passive film overlap together as
Fig. 9. Bode plots for aged specimens with (a) 4.3 wt% Mn, (b) 6.9 wt% Mn, (c) 9.7 wt% Mn contents in 3.5 wt% NaCl solution.
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Fig. 10. Equivalent electric circuit tested to model the experimental EIS data. Rs is solution resistance, Rt is charge transfer resistance, QPEt is the capacitive at the interfaces, Rfilm is oxide film resistance and QPEfilm is the capacitive of the oxide film.
a horizontal region at low and middle frequencies. This can be attributed to a compact layer spontaneous passive film growth in the early aging period. However, the frequency region of the horizontal part becomes narrower by increasing the aging time to 930 min. These indicate that passive film defects may be generated around s precipitates to weaken the passive state [39]. When the Mn content increases from 6.9 to 9.7 wt% (Fig. 9bec), no obvious changes are observed for the phase angle and frequency region of the horizontal part with aging treatment from 60 to 930 min, indicating that the capacitance feature of the passive film was retained, despite the accelerating effect of s precipitates at higher Mn contents. Meanwhile, the frequency region of the horizontal part decreased slowly with aging time from 15 to 930 min in the 9.7 wt% Mn content specimen, indicating that the passive ability was weakened, which is associated with the formation of more Cr-rich s precipitates caused by d-ferrite decomposition. Different models were proposed to interpret impedance spectra on passive metal surfaces [40e42]. Two regions corresponding to two time constants can be distinguished in Fig. 9, thus the electrode system stainless steel/NaCl solution for different aged specimens can be schematically illustrated in Fig. 10. In this circuit, Rs is the electrolyte resistance, Rt is the charge transfer resistance and QPEt is the capacitance of the corrosion product layer, including the defects that were caused by the formation of ionically conducting paths across the s precipitates. Rfilm is the resistance of the passive film and QPEfilm is the capacitance of the passive layer within d/g phases. All the samples with different Mn contents are spontaneously passive at the OCP, thus, this equivalent circuit is suitable for revealing the corrosion mechanism. This model assumes that the passive film does not totally recover the metal and can be treated as a defective layer [40,41]. The inhomogeneous degree on metallic substrates was related to precipitation initiation and growth rate with prolonging aging time. Table 4 lists the best fitting electrical parameters based on the
Fig. 11. Rp resistance from the fitting procedure for samples with different Mn contents aged at 800 C tested in 3.5 wt% NaCl solution.
circuit in Fig. 10. The Rt of the aged specimens is small and decreases with increasing Mn content from 4.3 to 6.9 wt% and 9.7 wt %, which indicate that more defects in the passive film appear as more s precipitates are formed with higher Mn content. The resistance of Rt depends strongly on the existence of defects, as more Cr-depleted zone around s-precipitates increases the number of defects in the d-ferrite matrix and along d/g boundaries. For different specimens, the oxide film resistance (Rfilm) decreased gradually with prolonging aging time from 15 to 930 min, while increasing Mn content slightly decreased Rfilm in specimens aged from 60 to 930 min. This indicates that the electrical conductivity of the films increases as more s precipitates are formed. Polarisation resistance, Rp (Rp ¼ Rt þ Rfilm), is commonly used as a measure of the resistance of a metal to corrosion damage [39]. The Rp values with different aging times and Mn contents are shown in Fig. 11. At the initial stage of 15 min, the Rp value for 4.3 wt% Mn is obviously higher than those for 6.9 and 9.7 wt% Mn, showing better corrosion resistance. As the aging time increased to 60 min, Rp decreases similarly regardless of Mn content, which flattens out by prolonging the aging time to 930 min in spite of a slightly higher Rp in the 4.3 wt% Mn sample. This suggested the corrosion resistance of the studied alloys is less likely to be affected after s phase precipitation by long aging time treatment with different Mn additions.
Table 4 Best fitting parameters for impedance spectra after stabilisation at the OCP. Aging time and Mn content
Rs (U cm2)
QPEt (U1 cm2 sn)
n1
Rt (U cm2)
QPEfilm (U1 cm2 sn)
n2
Rfilm (U cm2)
15 min, 4.3 wt% Mn 60 min, 4.3 wt% Mn 150 min, 4.3 wt% Mn 450 min, 4.3 wt% Mn 930 min, 4.3 wt% Mn 15 min, 6.9 wt% Mn 60 min, 6.9 wt% Mn 150 min, 6.9 wt% Mn 450 min, 6.9 wt% Mn 930 min, 6.9 wt% Mn 15 min, 9.7 wt% Mn 60 min, 9.7 wt% Mn 150 min, 9.7 wt% Mn 450 min, 9.7 wt% Mn 930 min, 9.7 wt% Mn
9.94 10.42 9.559 9.058 9.499 9.906 6.735 9.257 10.32 10.22 10.41 9.691 11.07 9.126 8.997
2.0655E-6 1.4708E-5 1.6285E-5 1.5012E-5 1.4126E-5 1.3678E-5 1.1866E-5 1.4675E-5 1.7982E-5 2.005E-5 7.6361E-6 9.6485E-6 1.5993E-5 1.6984E-5 2.2031E-5
0.81315 0.90643 0.89617 1.027 1.06 1.048 1.022 1.053 1.055 1.035 0.95033 0.96959 1.058 1.025 0.92285
336.3 50.65 26.46 7.027 8.512 8.254 5.598 8.291 10.01 11.26 17.74 9.719 10.65 11.5 1.6413
1.4706E-7 5.4361E-6 6.6551E-6 6.9093E-6 5.9946E-6 7.7887E-6 6.5796E-6 5.9275E-6 5.5146E-6 5.6385E-6 1.4748E-6 3.0281E-6 5.6474E-6 5.3235E-6 2.376E-11
0.78205 1.112 1.099 0.85444 0.85367 0.86778 0.83991 0.84406 0.84309 0.83036 0.79981 0.83539 0.84878 0.82978 0.46624
1412200 123450 92291 109760 149390 126720 105430 94420 88993 71886 238790 140730 115580 111940 77582
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4. Conclusions The influence of Mn addition on the microstructural and corrosion behaviour of aged 22 Cr economical DSSs can be concluded as follows: (1) The s precipitates nucleated at g/d interfaces and grew into the adjacent d-ferrite with increasing aging time from 15 to 930 min. By adding Mn in concentrations from 4.3 wt% to 9.7 wt%, s precipitates were formed more rapidly in the early stage of aging time from 30 to 150 min, but decreased gradually up to 930 min. The acceleration of s phase precipitation with increasing Mn can be attributed to the partition of the Mn element. (2) For the 4.3 wt% Mn content specimen, only about a 150 mV Eb drop occurred with increasing aging time from 15 to 930 min in 3.5 wt% NaCl solution, but Eb dropped significantly at the early aging stage with higher Mn contents of 6.9 and 9.7 wt% due to the formation of more aging-induced s precipitates. The metastable pit nucleation resistance decreased with Mn content from 4.3 to 6.9 wt% and was suppressed by a higher Mn addition of 9.7 wt% for aging treatment from 150 to 930 min. (3) In the initial aging stage of 60 min, the pit initiation sites occurred mainly along d/g boundaries and within d-ferrite due to a low PREN value of ferrite, which gradually occurred around s precipitates up to 930 min due to lower PREN value of the g2 product phase. Higher Mn content increased both the number and size of pits in the early stage of aging treatment due to the effect of Mn in increasing s precipitates. (4) EIS measurements showed that the charge transfer resistance, Rt, depends strongly on the existence of defects caused by s precipitates. Both Mn content and aging time were observed to affect the parameters associated with the passive film. In the initial aging stage of 15 min, the value of Rfilm in the sample of 4.3 wt% Mn is significantly higher than that of 6.9 and 9.7 wt% Mn, whereas higher Mn contents slightly decreased the value of Rfilm in the specimens aged from 60 to 930 min. Acknowledgements This work was mainly supported by the Fundamental Research Funds for the National Natural Science Foundation of China (No. 51261010). References [1] K.M. Lee, H.S. Cho, D.C. Choi, Effect of isothermal treatment of SAF 2205 duplex stainless steel on migration of delta/gamma interface boundary and growth of austenite, J. Alloy. Compd. 285 (1999) 156e161. [2] Y. Yang, B. Yan, The microstructure and flow behavior of 2205 duplex stainless steels during high temperature compression deformation, Mater. Sci. Eng. A 579 (2013) 194e201. [3] J. Li, Z. Ma, X. Xiao, J. Zhao, L. Jiang, On the behavior of nitrogen in a low-Ni high-Mn super duplex stainless steel, Mater. Des. 32 (2011) 2199e2205. [4] Y.H. Jang, S.S. Kim, J.H. Lee, Effect of different Mn contents on tensile and corrosion behavior of CD4MCU cast duplex stainless steels, Mater. Sci. Eng. A 396 (2005) 302e310. [5] I.-U.-H. Toor, P.J. Hyun, H.S. Kwon, Development of high MneN duplex stainless steel for automobile structural components, Corros. Sci. 50 (2008) 404e410. [6] R.F.A. Jargelius-Pettersson, Application of the pitting resistance equivalent concept to some highly alloyed austenitic stainless steels, Corrosion 54 (1998) 162e168. [7] C.J. Park, H.S. Kwon, Effects of aging at 475 C on corrosion properties of tungsten-containing duplex stainless steels, Corros. Sci. 44 (2002) 2817e2830. [8] E. Angelini, B. Benedetti, F. Rosalbino, Microstructural evolution and localized corrosion resistance of an aged superduplex stainless steel, Corros. Sci. 46 (2004) 1351e1367.
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