2022 Volume 62 Issue 3 Pages 568-576
A Gleeble thermo-mechanical simulator was used to simulate the welds of duplex stainless steel S32101. A micro-precipitate of chromium-nitride and its surroundings in the simulated welds was analyzed in detail by SEM/AES and EC-AFM. The AES analysis revealed that a chromium-depleted area is present in the γ phase near the chromium-nitrides precipitated at the α/γ grain boundary. The EC-AFM observation revealed that the duplex stainless steel preferentially dissolves from the chromium-depleted area.
Stainless steels are widely used as construction materials in seawater environments. In the past, austenitic stainless steel was preferred due to its mechanical and anti-corrosive properties. Especially in a severely corrosive environment (e.g., high-chloride seawater), high-corrosion-resistant alloys such as nickel-based alloys are used. For example, UNS S31603 grade is mainly used as piping material in reverse-osmosis desalination plants. Recently, duplex and super-duplex stainless steels have been replacing S31603 because such plants require a longer lifetime.1) They are cheaper than nickel-based alloys and have higher strength and higher corrosion resistance than that of S31603; thus, it is suitable for desalination plants, which are more highly corrosive environments than ordinary seawater. However, duplex and super-duplex stainless steels have higher nickel and chromium contents compared to S31603, which increases the cost of materials for plant construction: the cost of pipes, valves, and other parts. Despite use of such expensive materials at desalination plants, water leakage may occur due to pitting or crevice corrosion (depending on the means of manufacturing the parts that corroded). As for countermeasures against such corrosion, the causes of corrosion at desalination plants have been investigated.2,3)
When duplex stainless steel is used for piping and various other apparatus in seawater environments, high-quality welding is indispensable for fabricating joints. Duplex stainless steel has ferrite (α) phase and austenite (γ) phase. However, previous research has shown that certain metallic compounds and other phases (σ, χ phases, etc.) may precipitate in duplex stainless steel during heat treatment.4,5) One such heat-affected problem is precipitation at the grain boundary, which sometimes causes low corrosion resistance. ISO and ASTM standards for mechanical and corrosion properties regulate such precipitation by phase ratio and precipitation ratio at the grain boundary.6,7,8) In general, the precipitation is controlled by heat treatment after welding.9,10) However, it is preferable to control it by optimizing maximum welding temperature and cooling rate, although the welding conditions are quite complex. A Gleeble thermo-mechanical simulator was applied to studies the effects of heat on corrosion of duplex stainless steel.11,12,13,14) These studies showed that to improve corrosion resistance of the heat-affected zone (HAZ), σ phases, secondary γ phase,15) and chromium compounds in duplex stainless steel must be minimized.
Most previous researches11,16,17,18,19,20,21,22,23,24,25) focused on the chromium precipitates such as chromium-nitrides, which precipitate in the α phase and grain boundaries of the duplex stainless steel.24) Previous studies showed that Cr2N precipitate deteriorates pitting corrosion resistance of duplex stainless steels.4,5,13,20,21) Sathirachinda et al. measured the Volta potential differences between the Cr2N precipitates and matrix (α and γ phases) of a super-duplex stainless steel (UNS S32750) using a Scanning Kelvin Probe Force Microscope (SKPFM) to investigate the effect of Cr2N on the corrosion resistance.4,21) They concluded that fine Cr2N precipitates (<100 nm) formed in the water-quenched super-duplex stainless steel from 1125°C may not affect the pitting properties, and the larger precipitates formed by holding at 800°C for 3-min may detrimentally affect the corrosion resistance. Bettini et al. reported the similar result for UNS S32205 stainless steel.20) Some researchers reported that pitting corrosion may initiate in the chromium-depleted area around the Cr2N precipitate.20,22) However, it has not been experimentally confirmed that there is a chromium-depleted area around Cr2N in duplex stainless steels. In addition, few studies have in-situ observations of preferential dissolution processes in micro area such as the chromium-depleted area. Therefore, in the present study, in order to clarify the preferential corrosion mechanism of duplex stainless steel, we analyzed the boundary between chromium-nitride and the matrix in detail by Scanning Electron Microscope/Auger Electron Spectroscopy (SEM/AES). Furthermore, we observed the preferential dissolution behavior of the chromium-nitride and its surroundings in synthesized seawater solution in-situ by Electrochemical Atomic Force Microscopy (EC-AFM).
Lean-duplex stainless steel specimens of UNS S32101 were used as base-alloy sheets. The chemical composition of the specimens is listed in Table 1. The specimens were cold-rolled to the desired thickness (2 mm). In the specimens, a HAZ was simulated by using a Gleeble 3180 unit for rapid heating and cooling. The Gleeble unit was previously applied as a weld HAZ simulator.11,12,13,14) In this study, in accordance with the actual welding process, the base-alloy sheets were rapidly heated to 1493 K (1220°C) and cooled to 423 K (150°C) at a rate of 80 K/s.
| Grade | C | N | Ni | Cr | Mo | Mn |
|---|---|---|---|---|---|---|
| S32101 | 0.02 | 0.21 | 1.67 | 21.6 | 0.32 | 0.52 |
Anodic polarization curve measurements were acquired by measuring the specimens immersed in aerated synthesized seawater by using an ordinary three-electrode cell. The chemical composition of the synthesized seawater is represented in Table 2. The salinity and pH were 3.5 mass% and 8.2. A platinum wire (φ1.0 mm/150 mm long) and saturated KCl Ag/AgCl (SSE) were used as the counter electrode and reference electrode, respectively. Anodic polarization curve measurement was conducted at a potential scan rate of 0.33 mV∙s−1 from an open circuit potential (OCP) to 1.5 V vs. SSE using a potentiostat (HZ-7000, Hokuto Denko Co. Ltd.). Working electrode was the base-alloy sheets and the simulated-HAZ specimens (12 mm × 100 mm × 2 mm). The specimen surface was polished with SiC abrasive paper to a 1200-grit finish and then cleaned ultrasonically in acetone. It was covered with an electrodeposition coating, leaving 100 mm2 as an effective electrode surface area. The polarization curve measurements were measured multiple times to confirm the reproducibility.
| Component | MgCl2∙6H2O | CaCl2∙2H2O | SrCl2∙6H2O | KCl | NaHCO3 |
| mass% | 1.11 | 0.15 | 0.004 | 0.069 | 0.02 |
| Component | KBr | NaF | NaCl | Na2SO4 | H3BO3 |
| mass% | 0.01 | 0.0003 | 2.45 | 0.41 | 0.41 |
The EC-AFM images were acquired in the synthesized seawater (Table 2) by using an E-sweep nano-navi station (Hitachi High Technologies Co. Ltd.) with a platinum-wire counter electrode and a special leak-less SSE (eDAQ Pty Ltd.) reference electrode. The pH of the synthesized seawater was adjusted to 4 by adding HCl to enhance the pitting corrosion. It took 0.25-h to acquire one image. The simulated-HAZ specimen (12 mm × 8 mm × 1 mm) was used as the working electrode. It was polished with SiC abrasive papers and alumina paste to a mirror surface, and then covered with insulating tape, leaving approximately 28 mm2 (ϕ6 mm) as the electrode surface area. The timeline of the EC-AFM measurement is shown in Fig. 1. The HAZ electrode was first immersed at OCP for 4-h in the synthesized seawater, secondly polarized at constant current of 35.4 μA∙cm−2 (as galvanostatic polarization) for 1-h, and finally kept at OCP for about 2-h. The first AFM image was recorded after 2.25-h immersion ((a) in Fig. 1). The second image was taken just before the end of the first 4-h immersion (b). The third image was recorded at 0.25-h galvanostatic polarization (c) after the first 4-h immersion. The fourth and fifth images were taken at 0.5-h (d) and 1.5-h (e) of the second immersion after 1-h galvanostatic polarization, respectively. All observations were carried out at room temperature (22 ± 3°C).
Timeline of EC-AFM measurement. The electrode was held in three consecutive stages: 4-h immersion, 1-h galvanostatic polarization, and 2-h immersion. The galvanostatic polarization was carried out at constant current density of 35.4 μA∙cm−2. AFM images were recorded in (a)–(e).
In the galvanostatic polarization, when the HAZ electrode was polarized to a higher potential than the pitting potential (Epit), it could be difficult to observe the growth of pit by EC-AFM because the growth rate is extremely high. Therefore, in this study, EC-AFM observation was performed in the passive potential region just below Epit where metastable pits form. The metastable pit grows relative slowly and repassivated. The constant current of 35.4 μA∙cm−2 was determined so that the HAZ electrode was polarized in the passive potential region where metastable pits formed.
2.3. Surface AnalysisA SEM (Hitachi SU6600, Hitachi High Technologies Co. Ltd.) and An AES (PHI700Xi, ULVAC-PHI, Inc.) were utilized for the surface analysis before and after the EC-AFM measurements.
Anodic polarization curve measurements of the base metal (BM) and the simulated HAZ of the S32101 specimen in synthesized seawater are shown in Fig. 2. The passive region of the BM appears from the OCP (−0.1 V) to about 0.4 V. The passive current is in the order of 10−8 ~ 10−7 A∙cm−2, although it oscillates. No pits were found on the surface of the electrode after potentiostatic polarization at 0.0 V, which was performed separately. Therefore, the current oscillation of 1 nA∙cm−2 or less is considered to be due to external noise. The rapid increase in current from about 0.4 V is due to the onset of pitting corrosion because growth-type pits were observed on the electrode surface after the anodic polarization curve measurements. On the contrary, the anodic behavior of the HAZ is more complex; that is, the HAZ shows almost the same passivation current as the BM below 0.0 V. However, the passive current increases rapidly from 0.05 V. The second passive region appears in the potential range of 0.1 to 0.3 V. This increase in passive current is considered to be due to the increase in the week spots of the passive film introduced by the Gleeble simulator. The current oscillation for the HAZ was observed in the second passive region (in 0.1 to 0.3 V) which may be due to the onset of the metastable pits. The steep increase in current due to formation of growth-type pits starts from 0.3 V. From these anodic polarization characteristics, it can be concluded that the number of defects of the passive film are significantly increased and pitting potential Epit is shifted more negatively by the heat treatment using the Gleeble simulator.
Anodic polarization curve measurements of BM and HAZ of S32101 in synthetic seawater. (Online version in color.)
The microstructures of the as-polished BM and HAZ of the S32101 specimen before the corrosion test are shown in Figs. 3(a) and 3(b), respectively. The phase ratios of γ and α phase and the chemical composition of each phase determined by Energy Dispersive X-ray Spectroscopy (EDX) are listed in Table 3. The phase ratios of the BM and HAZ are not so different and fall within the provisions of ISO15156-1.6) As for the composition determined by SEM/EDX, the BM and HAZ show no significant difference. These results indicate that it is difficult to explain the difference in the anodic behaviors of the BM and HAZ shown in Fig. 2 in terms of the phase ratio and the chemical composition in macro-sized observations.
SEM images of as-polished BM (a) and HAZ (b) of S32101 before corrosion test.
| Grade | Part | Phase | Phase ratio | Si | Cr | Mn | Fe | Ni |
|---|---|---|---|---|---|---|---|---|
| S32101 | BM | α | 48.6 | 0.52 | 22.56 | 4.56 | 71.44 | 0.93 |
| γ | 51.4 | 0.52 | 20.50 | 5.41 | 71.82 | 1.76 | ||
| HAZ | α | 46.8 | 0.44 | 22.62 | 4.88 | 70.90 | 1.17 | |
| γ | 53.2 | 0.45 | 20.34 | 5.36 | 71.94 | 1.91 |
Microstructure of as-polished BM and HAZ before the corrosion tests were observed by SEM/AES. The BM has no precipitate as shown in Fig. 4. The chromium precipitates detected in the α phase of HAZ are shown in Fig. 5. For HAZ, the precipitates were present in the α phase and at the grain boundary between the α/γ phases; however, they were not detected in the γ phase.26,27,28,29,30) From the SEM image, it can be seen that multiple precipitates with size of 1 μm or less are present in the α phase. From the element distribution of nitrogen and chromium, it can be seen that four chromium-nitrides and other chromium precipitates were detected. The result of line AES analysis of the precipitates is shown in Fig. 6. It shows that two precipitates are chromium-nitrides, and the chromium concentration in the vicinity of the nitride decreases slightly, but the chromium-depleted area in the α phase could not be clearly confirmed from the surface AES analysis and line AES analysis.
SEM image of as-polished BM (a) and element distributions of N (b), Cr (c) and Fe (d) before corrosion test. (Online version in color.)
SEM image of as-polished HAZ and element distributions of N (b), Cr (c) and Fe (d) before corrosion test. (Online version in color.)
SEM image of the precipitates in ferrite phase of as-polished HAZ (a) and AES line-scan profile across the precipitates (b) before corrosion test. The electron beam was scanned along a red dotted arrow. (Online version in color.)
The precipitates on a grain boundary of the α/γ phases are shown in Fig. 7. From the distribution of chromium and nitrogen, it can be seen that three chromium-nitrides and one other chromium precipitate are present near the grain boundary. Among them, the two precipitates are present on the α phase side, and the remaining two precipitates are just on the grain boundary. It should be noted that the chromium-depleted area is clearly observed in the vicinity of the multiple precipitates in Fig. 7(c). It is possible that because the multiple chromium precipitates are precipitated nearby, the chromium-depleted area is clearly visible between those precipitates. In case of the precipitates in the α phase (Fig. 5), each precipitate is present at a certain interval, so the chromium-depleted area was not clearly confirmed.
SEM image of the precipitate at α/γ grain boundary of as-polished HAZ (a) and element distributions of N (b), Cr (c) and Fe (d) before corrosion test. (Online version in color.)
It can be seen in Fig. 7(c) that the chromium-depleted areas around these chromium precipitates are present in both the α and γ phases. The presence of the chromium-depleted area near the chromium nitride marked by “×” in Fig. 7(a) (which is precipitated on the grain boundary) is not clear from the result of surface AES analysis shown in Fig. 7(c). Accordingly, that chromium-nitride was subjected to line AES analysis, the result of which is shown in Fig. 8. This figure confirms that this chromium-nitride has a thin chromium-depleted area with width of 0.5 μm. The depletion area exists in both the α and γ phases, although it is more prominent in the γ phase because the γ phase has a lower diffusion rate of chromium.31) From the results of AES point analysis of this precipitate, the chemical composition of the precipitate was determined to be 14.9at% C, 13.1at% N, 31.2at% Cr, and 40.8at% Fe. In other words, it was found that the chromium-nitride is actually a composite nitride of chromium and iron. Regarding carbon, it is not possible to determine whether it is present as a chromium or iron carbide or a contaminant on the surface. It was previously reported that chromium-carbide and σ phase precipitated in duplex stainless steel.32) However, the cooling rate was considerably lower than that used in the present study, in which a similar cooling rate to that used in the actual welding process was used. Previous and present results indicated the risk of chromium-depletion occurred in actual engineering process of lean duplex stainless steel.
As mentioned in Section 3.3, precipitates of chromium-nitride and the presence of the chromium-depleted area were found in the vicinity of the α/γ grain boundary of the HAZ. To confirm whether these sites preferentially dissolve, the α/γ grain boundary of the HAZ was subjected to in-situ observation by EC-AFM in synthesized seawater (pH4). EC-AFM images were recorded under immersion (OCP) and galvanostatic polarization. The obtained EC-AFM images are shown in Fig. 9. The HAZ electrode was set so that the α/γ grain boundary where the chromium-nitrides are in the center of the images. It took about 2-h to set up the EC-AFM system. An AFM image (a) in Fig. 9 was taken to start immediately after the setting was completed. At this point, the electrode had already been immersed in the solution for about 2.25-h. After the electrode was immersed for another 2-h, galvanostatic polarization at constant current density of 35.4 μA∙cm−2 was started. An AFM image measured immediately before the galvanostatic polarization, that is, after immersion for a total of about 4-h, is shown in Fig. 9(b). AFM images taken at 0.25-h and 1-h after the start of the galvanostatic polarization are shown in Figs. 9(c) and 9(d), respectively. As for image (d), the observation failed immediately after the completion of galvanostatic polarization because of gas evolution. Thus, the observation was recorded at 0.5-h after completion of the galvanostatic polarization. The EC-AFM images (a) to (d) in Fig. 9 were recorded at the point (a) to (d) in Fig. 1, respectively. The height of the γ phase is lower than that of the α phase in all images. This difference is not because the γ phase dissolves faster than the α phase, because similar height differences can be seen for the sample immediately after polishing.
EC–AFM images of HAZ in synthetic seawater (pH4). The AFM images of (a) to (d) were recorded at points (a) to (d) in Fig. 1, respectively. “Cr-N” and “Cr-dep” indicate the precipitation sites of chromium-nitride and the chromium-depleted area, respectively. (Online version in color.)
An SEM image and element distribution of nitrogen, chromium, and iron recorded by AES after the EC-AFM experiment are shown in Fig. 10. The figure confirms that three chromium-nitrides exist at the α/γ grain boundary. By comparing the EC-AFM image in Fig. 9 with the SEM/AES image in Fig. 8, the precipitation sites of the chromium-nitrides in the AFM images are identified. They are shown as “Cr–N” in Fig. 9.
SEM image of the precipitates at α/γ grain boundary of HAZ and element distributions of N (b), Cr (c) and Fe (d) after EC-AFM observation. (Online version in color.)
From the SEM image in Fig. 10, it can be seen that the α/γ grain boundary is periodically wavy and that three chromium-nitrides are discontinuously precipitated on the convex parts of the α phase. It is therefore concluded that a γ phase exists between the chromium-nitrides. The profile of the relative height to the chromium-nitride is shown in Fig. 11. These height profiles were acquired when the AFM was scanned along the grain boundary in the direction of white dotted arrows in Fig. 9. In this observation, three high regions and four low regions were observed between X = 0 and X = 8.0 alternately. An AES line profile scanned along the same line as the height profile is shown in Fig. 12. It was recorded after the EC-AFM observation. Comparing Figs. 11 and 12 shows that the high plateaus measured by AFM correspond to the chromium-nitrides, and the chromium content of the γ phase between those chromium-nitrides at the grain boundary is much lower than that in the grain. The minimum value of chromium content in the γ phase is 14.2 at%.
Change in the relative height of the chromium-depleted areas (“Cr-dep” in Fig. 11) is shown in Fig. 13. The corrosion depth of the chromium-depleted areas was determined from the height of the chromium-nitrides (Xd =0) under the assumption that the dissolution of the chromium-nitrides was negligible during immersion and galvanostatic polarization. The average height of the highest part of each of the three chromium-nitrides was used as a reference point. Four chromium-depleted areas are shown in Fig. 11. In Fig. 13, the average value of the maximum corrosion depth of each chromium-depleted area is represented by a closed circle, and the maximum and minimum values are indicated by an error bar. Measurement points (a)–(d) in the figure were obtained from Figs. 9(a)–9(d), respectively. The corrosion depth at 1.5-h immersion after the end of galvanostatic polarization, although not shown in Fig. 9, is also plotted as (e) in the figure. Dissolution of the chromium-depleted area progresses by about 11.8 nm by 2-h immersion before galvanostatic polarization. The dissolution rate corresponds to current density of about 4 μA∙cm−2. It can be seen that the dissolution proceeds at a certain rate even at the OCP. Meanwhile, the dissolution progresses by about 12.5 nm by 1-h galvanostatic polarization, which corresponds to dissolution rate of approximately 10 μA∙cm−2. The dissolution in the chromium-depleted areas is actually enhanced by the galvanostatic polarization, but it is less than expected. The dissolution rate of 10 μA∙cm−2 is smaller than the applied current density of 35.4 μA∙cm−2. Furthermore, it can be seen from Fig. 13 that when the sample is immersed for another 1-h (e) after (d), the dissolution hardly proceeds. This is because the dissolution of the chromium-depleted areas in Fig. 9 was completed to some extent by the galvanostatic polarization. This result suggests that dissolution at the chromium-depleted areas did not proceed continuously for a long time. This is because the galvanostatic polarization was performed at a potential where no growth-type were formed. Figure 14 shows the change in potential in galvanostatic polarization. The potential during the galvanostatic polarization is in the range of 0.15 to 0.25 V. It can be seen from the polarization curve measurement of HAZ in Fig. 2 that this potential certainly corresponds to the passive potential where metastable pits were formed. It was also confirmed by an optical microscope that many small rust spots over the electrode surface after the AFM measurement were observed, but growth-type pits were not present. The EC-AFM study revealed that the initiation site of the metastable pit is the chromium-depleted area around the chromium nitride. Accordingly, the growth-type pit is also expected to be initiated from the chromium-depleted area.
Change in potential under galvanostatic polarization at constant current density of 35. 4 μA∙cm−2.
A specimen simulating the welded part (HAZ) of lean-duplex stainless steel UNS S32101 was prepared by the Gleeble thermo-mechanical simulator. The precipitated chromium-nitrides and its surroundings on the surface of the simulated HAZ were analyzed in detail by SEM/AES, and the dissolution behavior was in-situ investigated in the synthesized seawater (pH4) by EC-AFM. The main conclusions that can be drawn are as follows.
(1) The simulated HAZ showed a larger passive current and a lower pitting potential than the solution-treated base metal S32101. The simulator can successfully reproduce the heat-affected area of the welding.
(2) SEM/AES revealed that chromium-nitrides are precipitated in the α phase and on the α/γ grain boundary and that chromium-depleted area exists in the vicinity of the precipitated chromium- nitrides. The depletion was particularly remarkable in the case of chromium-nitrides precipitated at the α/γ grain boundaries. The shape of the observed grain boundaries was periodically wavy, multiple chromium-nitrides were discontinuously precipitated on the boundary of convex parts of the α phase, and chromium-depleted areas were formed in the γ phase between the chromium-nitrides.
(3) The preferential dissolution from the chromium-depleted area can be successfully monitored by EC-AFM.
