**3. Results**

Figure 2a presents the microstructures of the solution-annealed hot-rolled 2205 steels. These consisted of elongated austenite islands in the ferrite matrix and no apparent intermetallic phases were observed. During aging at 850 ◦C for 10 min, the χ phase preferentially nucleated at the boundary of ferrite and grew through the adjacent ferrite, as presented in Figure 2b. This was discussed in the authors' previous work [29]. Furthermore, it could be observed form Figure 2c,d, that, as the aging time increased to 3 and 4 h, the σ phase originating from the transformation of α → γ2 + σ would gradually appear and being coarse, where the γ2 was the secondary austenite. Therefore, it could be considered that high amounts of the σ phase existed in the matrix during aging for increased durations. Besides, a low amount of χ phase was also detected and distributed at the grains of the σ phase. Figure 2e,f presents the EDS line-scan profile of the hot-rolled 4 h aged samples, where the σ phase with a relatively low Cr and Mo contents could cause the generation of a chromium depleted region.

Figure 3 present the metallographic structure of 2205 duplex stainless steel with cold deformation subsequently to solid-solution and aging treatments. Figure 3a presents the microstructure of the solution-annealed cold-deformed sample under the optical microscope, and it can be seen that the austenitic phase with a lighter color was distributed within the ferrite phase, but the microstructures of the cold-rolled samples became more elongated along the cold rolling direction, as compared to the hot-rolled samples. Moreover, the austenite grains of the cold-deformed samples were fine and non-uniform, with a narrow strip shape, as a result of different local deformation. Following aging for 10 min at 850 ◦C, the bright white χ phase could also be observed in Figure 3b. As the aged time increased to 3 h and 4 h, as presented in Figure 3c,d, the σ phase precipitates began to appear at the boundary of the ferrite and austenite phases, but the corresponding grain size was lower when compared to the hot rolled samples. Figure 3e,f presents the EDS line-scan profiles of the cold-rolled sample aged for 4 h, suggesting that the precipitated χ and σ phases were enriched in Cr and Mo. This led to the uneven distribution of the alloying elements, such as Cr and Mo, within the matrix.μm

**Figure 2.** Microstructure of cold-rolled 2205 duplex stainless steel: (**a**) optical microscope (OM) morphology of hot-rolled specimen; (**b**) scanning electron microscopy (SEM)-scattered electron imaging (BSE) morphology of 10 min aged; (**c**) 3 h aged and (**d**) 4 h aged specimens at 850 ◦C; and, (**e**,**f**) electron dispersive X-ray spectroscopy (EDS) line-scan profile of hot-rolled 4 h aged specimen.

In Figure 4, the σ phase volume fractions of the cold-rolled and hot-rolled 2205 steels were plotted as a function of aging time. As the aging time increased, the amount of σ phase in the cold-rolled samples became gradually higher when compared to the hot-rolled samples at the same aging time. In particular, the precipitations of the cold-rolled samples reached approximately 38.2% for the specimen that was aged at 850 ◦C for 4 h, corresponding to 23.9% in the hot-rolled samples. The XRD diffraction spectra and local magnifications of the cold-rolled and hot-rolled samples are presented in Figure 5. It could be observed that the peak intensity of α(110), relative to the γ(111) decreased with aging time. Also, the sample that was aged for 4 h exhibited a low-sized ferrite peak,

indicating that a major fraction of the ferrite was transformed into the σ phase. The diffraction peak intensity of the σ phase in the cold-rolled sample was significantly higher when compared to the hot rolled sample following aging for 4 h, which suggested that an additional amount of the σ phase was precipitated in the cold-rolled sample. It could also be observed from Figure 5 that the peak intensity of the α(200) and α(211) in the cold-rolled solution-annealed samples was significantly higher when compared to the hot-rolled samples.

**Figure 3.** Microstructure of cold-rolled 2205 duplex stainless steel: (**a**) OM morphology of cold-rolled specimen; (**b**) SEM-BSE morphology of 10 min aged; (**c**) 3 h aged and (**d**) 4 h aged specimens at 850 ◦C; and, (**e**,**f**) EDS line-scan profile of cold-rolled 4 h aged specimen.

Figure 6 presents the potentiodynamic polarization curves comparison between the cold-rolled and hot-rolled 2205 duplex stainless steels subsequently to different aging treatments with a 3.5 wt % NaCl solution. Moreover, Tables 2 and 3 present the *E*pit results in comparison from the potentiodynamic polarization curves for the hot-rolled and cold-rolled 2205 steels, respectively. It can be observed that the polarization curves of the hot-rolled and cold-rolled samples without aging were similar, while both of the pitting potentials (*E*pit) exceeded 1000 mV, indicating that the pitting resistance was highly consistent for the solid-solution annealed cold-rolled and hot-rolled steels. Adversely, for the cold-rolled

sample, the *E*pit presented in Figure 6b apparently decreased following aging for 10 min only, while no current density at the end of the polarization curve existed, suggesting that the sample surface was not repassivated. As the aging time increased to 20 min, the *E*pit was reduced to 643 mV, which was lower than the *E*pit of the hot-rolled samples aged for 3 h. Through the aging time further increase to 40 min, the *E*pit decreased to 548 mV, where the value was even lower than the hot-rolled samples aged for 4 h. When the aging time of the cold-rolled sample increased to 1 h or beyond, the *E*pit further decreased, and remained it below 500 mV.

**Figure 4.** Changes in σ phase volume fraction of cold-rolled and hot-rolled of 2205 duplex stainless steel with prolonging aging time.

**Figure 5.** Comparison of X-ray diffraction patterns between cold-rolled and hot-rolled 2205 duplex stainless steel without aging treatment, and after aging for 10 min, 4 h at 850 ◦C, (**a**) overall patterns; and, (**b**) amplification patterns of cold-rolled and hot-rolled 4 h aged specimens.

**Table 2.** *E*pit resulted from potentiodynamic polarization curves for hot-rolled 2205.


**Table 3.** *E*pit resulted from potentiodynamic polarization curves for cold-rolled 2205.

**Figure 6.** Comparison of potentiodynamic polarization curves between cold-rolled and hot-rolled 2205 duplex stainless steel after different aging treatment, (**a**) cold-rolled specimens, and (**b**) hot-rolled specimens.

Figure S1 presents the real impedance vs. the imaginary impedance plot at each frequency for the 2205 duplex stainless steels with different aging times with the 3.5% NaCl solution. It could be observed that the Nyquist diagrams of the cold-rolled and hot-rolled samples exhibited a depressed semicircle with a capacitive arc. Moreover, the diameter of the capacitive semicircle in the cold-rolled sample was lower as compared to the hot-rolled samples, indicating that the passive film stability of the cold-rolled samples was worse when compared to the hot-rolled samples. The equivalent circuit presented in Figure S2 was proposed for the EIS data fitting to quantify the electrochemical parameters. In this equivalent circuit, Rs is the solution resistance and Rt stands for the charge-transfer resistance. CPE1 represents the capacitance of the double electrical layers. CPE2 symbolizes the capacitance of the passive film on the metal surface. Rf is the passive film resistance. The electrochemical impedance parameters of the cold-rolled and hot-rolled samples that were obtained from the fitting of the EIS diagrams are presented in Tables S1 and S2, respectively. The passive film resistance (Rf) of the cold-rolled samples exceeded the surface charge transfer resistance (Rt), which occurred similarly for the hot-rolled samples. The passive film of the latter played a major role in the corrosion resistance. Furthermore, the Rf of the cold-rolled samples was significantly lower as compared to the hot-rolled samples, indicating that the passive films of the samples treated by cold rolling were significantly weaker compared to the hot-rolled samples. Moreover, the Rf change in the cold-rolled samples aged for a short time was not apparent, but the values of Rf for the samples following aging for 1 h or longer times significantly decreased, indicating that the passivation film became quite weaker subsequently to aging for a long time. The Rt also decreased as the aging time increased, which demonstrated that, the migration of the charged particles in the double layer between the electrode and the electrolyte solution gradually became easy.

It is well-known that the corrosion resistance has always been considered in regard to the microstructures of materials. For the hot-rolled 2205 duplex stainless steel, the precipitation of σ phase gradually increased as the aging time increased, while the corrosion resistance of the samples significantly decreased. In order to explain the correlation between the microstructure and corrosion resistance, the corrosion morphology of the hot-rolled 2205 duplex stainless steels aging for 10 min and 4 h, following potentiodynamic polarization, was characterized through SEM-BSE with an EDS system. In the results, each residual phase was confirmed. As presented in Figure 7a, in the hot-rolled 2205 aged for shorter aging times, pitting nucleation preferably occurred on the grain boundaries or on the ferrite/austenite interfaces. Furthermore, a severe pitting corrosion of the samples aged for 4 h was observed, as presented in Figure 7b. Tables 4 and 5 confirmed the chemical compositions of the phases presented in Figure 7.

**Figure 7.** Corrosion morphology of hot-rolled 2205 duplex stainless steel after potentiodynamic polarization by SEM-BSE, (**a**) 10 min aged; and, (**b**) 4 h aged.


**Table 4.** Chemical composition of phases showed in Figure 7a (wt %).

**Table 5.** Chemical composition of phases showed in Figure 7b (wt %).


The corrosion morphology comparisons demonstrated that the corrosion of cold-rolled samples had apparent selectivity, in which the order of corrosion of each phase apparently differed from the hot rolled samples. Figure 8 present the corrosion morphology of the cold-rolled 2205 duplex stainless steel aged for 10 min, 1 h and 4 h; subsequently, to potentiodynamic polarization, through SEM-BSE. The pitting corrosion of the cold rolled samples was more evenly distributed on the sample surfaces when compared to the hot rolled samples, and it was easy to be concentrated in certain areas. Figure 8a presents that the phase boundary of cold-rolled 2205 preferred to be corroded in the 3.5 wt % NaCl solution, whereas it was also worth being noted that stripe patterns appeared in the corrosion morphology of the sample that was aged for 10 min. The corresponding EDS analysis is presented in Table 6. The chemical composition of the stripe was the same as the surrounding austenite, which might be α -martensite. Furthermore, as presented in Figure 8b–d, the surfaces of the cold-rolled 2205 samples still remained a high amount of shallow white σ phases following corrosion. Therefore, differently from the hot-rolled 2205 sample, the cold-rolled 2205 DSS, subsequent to aging treatment, might be preferentially corroded from the phase boundary and the α -martensite in the 3.5 wt % NaCl solution. Also, the precipitates were basically not subjected to corrosion in the initial process.

**Figure 8.** Corrosion morphology of cold-rolled 2205 duplex stainless steel after potentiodynamic polarization by SEM-BSE, (**a**) 10 min aged; (**b**,**c**) 1 h aged; and, (**d**) 4 h aged.

**Table 6.** Chemical composition of strain-induced martensite and γ phase showed in Figure 8a (wt %).


To investigate in detail the corrosion behavior of the hot-rolled and cold-rolled 2205 DSS, the σ phase was selectively dissolved for the sample fabrication without the σ phase through the electrochemical method. The potentiodynamic polarization curves among the original specimens and the specimens without an σ phase are presented in Figure 9. For the hot-rolled materials, the Epit of the sample without σ phase would increase to the solution-annealed values. For the cold-rolled 2205 DSS, the precipitation content of the σ phase increased, whereas the precipitation speed of the precipitates also increased. This appeared to undermine the corrosion resistance of the cold-rolled 2205 DSS, as compared to the hot-rolled materials. However, even if the σ phase was eliminated from the cold-rolled samples, the Epit increased from 443 mV to 630 mV, being quite lower when compared to the non-aging sample (1002 mV). This suggested a more complicated influence factor on the corrosion resistance of the cold-rolled samples. The passive current of the samples without σ phases was higher when compared to the solid solution samples, which might be caused by the formation of a compact structure between the steel and the epoxy resin.

**Figure 9.** Comparison of potentiodynamic polarization curves between original specimens and specimens without σ phase, (**a**) hot-rolled specimens; and, (**b**) cold-rolled specimens.

#### **4. Discussion**

The 2205 DSS, χ phase would preferentially precipitate at the ferrite-ferrite boundaries, following aging at 850 ◦C. As the aging time increased, the σ phase would precipitate at both ferrite boundaries and ferrite-austenite boundaries. As the χ phase was a meta-stable phase, it would dissolve and transform into the σ phase along with the aging time increase. The σ phase could be transformed by a eutectoid reaction from ferrite and the dissolution of the χ phase. The σ phase precipitation could be accelerated by the cold deformation in the subsequent aging, which might be due to the increased defect and distortion energy during cold-rolling, promoting the ferrite transformation into the σ phase. Besides, it was proved that the austenite phase in duplex stainless steels would be transformed into ε-martensite or α -martensite during cold deformation, also being directly transformed into the α -martensite when the stacking fault energy was high [28]. Moreover, it is well known that ferrite has a body-centered cubic structure and the martensite has a body-centered tetragonal structure. Since the diffraction peak of α -martensite was consistent with the ferrite peak, the peak intensity of ferrite increased when containing α -martensite. Therefore, the XRD diffraction pattern that is presented in Figure 5 indicated the formation of α -martensite following cold rolling. By contrast,

due to the non-diffusive phase transition and unchanged chemical composition of the α -martensitic transformation, the direct observation of α -martensitic phase appeared difficult.

It has been proved that pitting occurring in chromium-depleted areas are associated with the precipitated phase [30,31]. When combined with Figures 2f and 7a, the ferrite was almost occupied by the σ phase and the secondary austenite, which often has relatively low Cr content than the σ or γ phases. Therefore, it could be concluded that the original grain-boundaries were the more susceptible to pitting, thus inducing the growth of the σ precipitates. During the initial stage, the pitting nucleation was located at the boundary of α/γ and around the σ phase, due to low Cr content and stability of the passive film in this position. With the transformation of α into σ and γ2, the γ2 was also corroded, allowing for the pitting to further extend within the α. Therefore, the pitting corrosion of the hot-rolled 2205 steels preference to occur around the σ phase was inevitable, due to the relatively high Cr and Mo contents of the σ phase. Moreover, pitting tended to further increase with the precipitation of an additional amount of σ phase in the hot-rolled 2205.

Besides the effect of σ phase on the corrosion resistance of the cold-rolled 2205 Dss, the influence factors of its corrosion were complicated. On the one hand, the cold-deformation of 2205 DSS easily generated dislocations, deformation twinning, and dislocation-twin interactions, where the nucleation of the pitting attack was quite likely to occur. Moreover, a more defective passive film on the cold-rolled 2205 DSS was likely to be formed at the defects of the grains. On the other hand, the strain-induced martensite in the cold-rolled 2205 might be active than the other precipitate and matrix phases. Consequently, it was also easily corroded in the early stage, which was similar to the reported role of strain-induced martensite in cold-worked 304 steels [32]. Therefore, the precipitated σ phase was the main factor affecting the corrosion resistance of the hot-rolled samples, which might be ascribed to the formation of a Cr-depleted zone around the σ phase, also preferring to be corroded. When the σ phase was removed, the corrosion resistance of the samples would be apparently restored, which suggested the chromium redistribution in the matrix of the hot-rolled 2205. Adversely, the σ phase elimination in the cold-rolled 2205 could not restore the *E*pit to the level of non-aging. Further research is required to be conducted to clarify the influence factor of the corrosion resistance for the cold-rolled 2205 DSS.
