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Effect of a brief post-weld heat treatment on the microstructure evolution and pitting corrosion of laser beam welded UNS S31803 duplex stainless steel
Corrosion Science 65 (2012) 472–480
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Effect of a brief post-weld heat treatment on the microstructure evolution and pitting corrosion of laser beam welded UNS S31803 duplex stainless steel Yanze Yang a, Zhiyu Wang b, Hua Tan a, Jufeng Hong a, Yiming Jiang a, Laizhu Jiang b, Jin Li a,⇑ a b
Department of Materials Science, Fudan University, Shanghai 200433, PR China Research and Development Center, Baosteel Co. Ltd., Shanghai 201900, PR China
a r t i c l e
i n f o
Article history: Received 30 April 2012 Accepted 20 August 2012 Available online 30 August 2012 Keywords: A. Stainless steel B. TEM C. Welding C. Pitting corrosion
a b s t r a c t The effect of laser-beam welding and subsequent short-time post-weld heat treatment at different temperatures and holding time on microstructure evolution and pitting corrosion behavior of UNS S31803 duplex stainless steel was investigated. The results showed the as-welded joint displayed impaired pitting corrosion resistance and that pitting preferentially occurred at ferrite grain in the fusion zone. Shorttime heat treatment enhanced the pitting corrosion resistance of welded joint. Optimal post-weld heat treatment of 3 min heat treatment at 1080 °C was identified at which the joint restored the pitting corrosion resistance lost during welding process. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction Duplex stainless steels (DSSs), which contain approximately equal amounts of ferrite (a) and austenite (c) phases, offer a unique combination of high corrosion resistance and mechanical properties [1–4]. Thus, DSSs are increasingly being used for various applications. It is well established that the good properties of DSS heavily depend on the balanced two-phase microstructure and the absence of deleterious secondary phases such as r-phase, v-phase, carbides and nitrides etc. [5–9]. DSSs generally show impaired corrosion resistance and toughness when welded, especially after laser-beam welding (LBW). LBW offers many advantages over conventional arc welding process such as low heat input, short cycle time and good cosmetic welds which make LBW especially suitable to high volume applications. However, the very fast cooling rate of LBW produces welds with excessive a-phase content. Since the solubility of nitrogen (N) in a-phase is much lower than that in c-phase, an N supersaturation in a-phase is prone to occur. Thus, the LBW process increases the a-phase volume fraction [10–13] and brings on the precipitation of chromium nitrides [14–20] due to the excessive nitrogen content in the ferrite phase as well as the lack of time for a ? c transformation [21,22]. Ramirez et al. [23] suggested that chromium nitrides precipitated at the interior of the ferrite grains of DSS after quenching from 1350 °C. As reported by Perren et al. [24], chromium nitrides lowered the corrosion resistance of DSS by enhancing critical current densities and passivation potential. ⇑ Corresponding author. Tel./fax: +86 21 65643648. E-mail address: [email protected] (J. Li). 0010-938X/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2012.08.054
Since LBW upsets the balanced a/c duplex microstructure and promotes the deleterious phases, it is necessary to eliminate the problems associated with excessive ferrite content and intermetallic phases. Post-weld heat treatment (PWHT) is recommended in practice [25]. PWHT is especially important for chemical and fossil energy industries that require the applications of materials with high corrosion resistance according to the American and European standards (ASTM A928/A928M and NORSOK MDS D42) [26]. Our group has studied the effect of PWHT temperature (1020– 1120 °C) on the pitting corrosion resistance of the plasma-arc welded (PAW) DSS 2304 [27]. The weak region of pitting was found at the heat-affected zone (HAZ) and the corrosion resistance of PAW welded joints was restored after 4 min PWHT at 1080 °C. Investigation performed by Young et al. [28] suggested that short-time PWHT at 1050 °C was able to effectively raise the austenite content and impact the toughness of the DSS welds. In summary, literature review has shown that there would be a strong relationship between the PWHT parameters, microstructure and pitting corrosion behavior of welded joints. However, little work has been carried out regarding the effects of short-time PWHT parameters, in particular annealing temperature and holding time, on the pitting corrosion behavior of the laser-beam welded DSS joints, even though such knowledge is important for improving the efficiency of the LBW process. In the present paper, duplex stainless steel UNS S31803 (DSS 2205) was joined by autogenous laser beam welding (LBW) without filer metal and subsequently annealed at different temperatures ranging from 1020–1100 °C for 1 min and 3 min. Potentiostatic critical pitting temperature technique (CPT) was employed to evaluate the corresponding pitting corrosion resistance, which
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has been proven to be more sensitive and reliable than other pitting characterizing approaches such as pitting potential [29,30]. The microstructure and the precipitation of chromium nitrides were investigated by scanning electron microscope (SEM), atomic force microscopy (AFM) and transmission electron microscope (TEM). The main objectives of this study are to: 1. Reveal the influence of the laser beam welding on the microstructure and pitting corrosion resistance of the DSS 2205 joint. 2. Assess the influence of short-time PWHT parameters (annealing temperature and holding time) on the microstructure evolution and the pitting corrosion behavior of laser beam welded DSS 2205 joint and to explain the mechanism of such behavior. 3. Determine the optimum short-time PWHT temperature and holding time to restore the pitting corrosion resistance lost during welding process. 2. Experimental 2.1. Material and sample preparation DSS 2205 (chemical composition in wt.% C: 0.021, P: 0.008, Ni: 5.72, Cr: 22.36, Mn: 1.52, Si: 0.46, Mo: 3.02, Co: 0.069, S: 0.001, N: 0.17, Fe: bal.) plates of 1.5 mm in thickness were obtained from Baosteel Group Corporation. Welding was performed using autogenous LBW without filler metal and the corresponding parameters were listed in Table 1. Considering the efficiency of practical application, the PWHT holding time should be as short as possible. According to industrial experiences, the holding time of PWHT for 1 mm thick sheet (or pipe) is about 1 min. Therefore, 1 min and 3 min annealing treatment was adopted for the 1.5 mm DSS 2205 samples in this study to evaluate the effectiveness of shorttime PWHT. Short-time PWHT was performed at 1020 °C, 1050 °C, 1080 °C and 1100 °C, respectively in a muffle furnace in atmosphere. Firstly, the muffle furnace was heated to the setting temperature and then the welded joint was put into the furnace and held for 3 min or 1 min at the setting temperature. After that, the welded joint was taken out and quenched in water. The specimens for electrochemical tests and microscopy examination, with dimension of 12 mm 12 mm 1.5 mm were cut from the joint including the fusion zone and the heat-affected zone. 2.2. Electrochemical measurements The pitting corrosion resistance of the specimens were evaluated by the CPT technique in 1 mol/L NaCl solution [31]. All the electrochemical measurements were carried out using a PARSTAT2273 potentiostat in a classical three-electrode cell. In this cell, a platinum foil served as the counter electrode and a saturated calomel electrode (SCE) served as the reference electrode. All potentials quoted in this paper referred to this reference electrode. The specimen acting as working electrode was mounted in epoxy resin. In order to avoid crevices, the area of the specimen near the interface between the resin and the specimen was sealed with
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silica gel sealant [32]. The exposed working electrode surface area was 100 mm2. Prior to each CPT measurement, the working electrode was ground with successive grade silicon carbide sand paper up to 1000 grit, degreased with ethanol, rinsed with distiller water and dried in air. The test solution, 1 mol/L NaCl solution, was made up from analytical grade reagents and distilled water according to the ASTM standard: G150–99 ‘‘Standard Test Method for Electrochemical Critical Pitting Temperature Testing of Stainless steels’’. The CPT measurement procedure is described as follows: (1) the surface of the working electrode was initially polarized cathodically at 0.9 VSCE for 120 s to reduce the oxide film formed in air. (2) working electrode was then stabilized at open circuit potential for 30 min in test solution at 2 °C; (3) a static potential of 0.75 VSCE was applied to working electrode and the electrolyte temperature was increased continuously at a rate of 1 °C/min controlled by a water bath. The corresponding current was recorded simultaneously with the increasing temperature. The experiment was terminated when the current density increased to 100 lA/cm2. The CPT is defined as the temperature at which the current density equals to 100 lA/cm2. The test solution was kept bubbled with pure N2 gas to remove the dissolved O2 throughout the test. All electrochemical tests were repeated at least three times. 2.3. Optical, SEM, TEM and AFM analyses Specimens for optical, SEM and AFM characterization were ground with successive grade abrasive papers up to 1500 grit and then polishing using 1.5 lm diamond paste. Before the observation, the specimens were electrolytically etched by 30 wt.% KOH solution at 2.0 VSCE for 15 s. Quantitative metallography was carried out using light optical microscopy (LOM) and Image-Pro Plus 6.0 software. The volume fractions of the two phases were measured in ten different locations in the same zone and the average value was used in this paper. The chemical compositions of the c-phase and a-phase were analyzed using an energy dispersive X-ray spectroscopy (EDS) linked to a SEM system (SCE Philips Xl 30FEG). Since the concentration of alloying elements may vary at different locations in the same zone, ten measurements were carried out and the average value was used. Two methods were employed to reveal the presence of chromium nitride in the DSS 2205 weld: (1) direct observation by using TEM; (2) observation by using SEM and AFM after electrolytic etching by 10 wt.% oxalic acid solution at 7 VSCE for 30 s [33]. Before TEM observation, thin foils of specimens were prepared following conventional TEM sample preparation methods. The samples were first mechanically polished to approximately thickness of about 0.1 mm followed by polishing at 25 VSCE in a solution of 10% perchloric acid in 90% methyl alcohol with a twin-jet unit. After that the samples were further thinned, the analyses were performed in a transmission electron microscopy (JEM-2010, JEOL) operated at 200 kV. A commercial atomic force microscopy (Dimension Icon, Bruker Corporation) was used to obtain images and topographical information of the studied specimens. The measurements were performed in the contact mode (using SNL cantilevers with a spring constant k = 0.12 N/m), and the AFM images were analyzed by NanoScope Analysis version 1.4 software package.
Table 1 Laser welding parameters used in the experiment. Laser beam configuration
Unit
Value
Laser power Travel speed Spot size on surface Focal position Deviation distance to focus spot Shielding gas flow rate (Ar) Backing gas flow rate (Ar)
W Mm/min Mm2 – Mm L/min L/min
1500 2000 1.5 1.5 At specimen surface 0 15 15
3. Results and discussion 3.1. Influence of LBW on microstructure and pitting corrosion resistance Laser beam welding is a high energy density and low heat input process, resulting in a very small heat affect zone (HAZ) and a deep and narrow fusion zone (FZ) with very little distortion [34].
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3.2. Microstructure Fig. 1 shows the morphologies of different zones of DSS 2205 LBW as-welded joint. Fig. 1a illustrates the typical duplex microstructure of the DSS 2205 base metal. The c-phase was shown as the lighter region while the dark region was the a-phase. The austenite was evenly distributed in the ferrite matrix and no apparent hazard phase was observed. Fig. 1b presents the microstructure of a typical laser beam welded joint. The width of fusion zone (FZ) was about 1 mm and the heat affected zone (HAZ) could hardly be identified. Therefore, special attention was paid to the FZ in this work. Fig. 1c shows the enlarged center part of FZ, 50–100 lm equiaxial ferrite grains were observed. While at the edge of FZ, as shown in Fig. 1d, 200 lm 30 lm columnar ferrite grains were observed grown towards the center of FZ driven by temperature gradient. The effect of the grain size/orientation on the pitting corrosion resistance of FZ was not considered in this paper. It is assumed that there is no obvious difference in the pitting corrosion resistance between different parts of FZ since the pitting corrosion resistance of FZ largely depends on the ferrite phase formed during the very fast cooling stage of the LBW, as will be discussed below. Generally, three types of austenite are formed during the apostsolidification transformation at the cooling stage of welding [35]: (1) austenite formed as allotriomorphs at the prior-ferrite grain boundaries (GBA); (2) austenite formed as Widmanstätten sideplates grown into ferrite grains (WA); (3) austenite formed as intragranular precipitates in ferrite grains (IGA). Fig. 1c shows that the GBA (light) located at the boundaries of ferrite grains due to the nucleation and growth of austenite at the boundaries. The amount of WA and IGA was very small. The average a-phase volume fraction was 93% in FZ (austenite phase volume fraction: 7%). Gilath et al. [36] estimated the magnitude of the cooling rate of LBW FZ and suggested it to be 103 K/s. Due to the very high cooling rate of LBW, the transformation from a-phase to c-phase as well as the diffusion of alloying elements were severely suppressed. In addition to the excessive a-phase, chromium nitride played an
important role in corrosion resistance properties of DSS. Chromium nitride introduced chromium depleted regions that were susceptible to corrosion attack [37]. It was reported that chromium nitrides formed in HAZ and FZ due to the influence of welding thermal cycle [10,12,38]. Fig. 2a shows the morphologies of the as-welded joint etched by oxalic acid. Small pits occurred within ferrite phase, indicating the presence of Cr2N precipitation. AFM depth profile and three dimensional AFM topography (Fig. 2b and Fig. 2c, respectively) reveal that the width of pits was about 0.2–2 lm and the depth of pits was about 100–150 nm. TEM image and selected area diffraction pattern (SADP), as shown in Fig. 2d and Fig. 2e, respectively, provide direct experimental evidence of the hexagonal Cr2N. No obvious intermetallic compound such as r-phase, v-phase and carbides was detected. The length of the needle-like Cr2N was estimated to be 200–400 nm. The observed accumulations and alignment of the Cr2N within ferrite grains (Fig. 2d) corresponded to the high dislocation density of those sites and could be explained through the diffusion impediment of N at these dislocations during rapid cooling. Thus, the pits observed by SEM and AFM characterization were correlated to Cr2N that precipitated in a-phase of FZ during welding process [39]. 3.2.1. Pitting corrosion resistance Fig. 3 presents the typical curves of the current density against solution temperature for the as-received DSS 2205 base metal and the as-welded specimens without PWHT. The CPT decreased from 56 °C to 42 °C after LBW process, indicating a severe degradation of the pitting corrosion resistance. Pit morphologies were characterized by optical/scanning electron microscope, as shown in Fig. 4a and Fig. 4b, respectively. Three stable pits were observed in the FZ of the as-welded specimen. Pit 1# clearly occurred in the aphase; pit 2# occurred in a-phase and was restrained by the c/a boundary; it 3# spanned multiple a/c domains with the c-phase mostly remained and the a-phase severely attacked. Thus, it was found that the mechanism of pitting corrosion was the selective
Fig. 1. Optical micrographs of as-welded DSS 2205 joints etched by 30 wt.% KOH solution showing (a) base metal (BM), (b) fusion zone (FZ), (c) enlarged centre of the fusion zone and (d) edge of fusion zone. Austenite (c) is light and ferrite (a) is dark. The ferrite phase volume content is more than 90% in the fusion zone.
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Fig. 2. SEM, AFM and TEM micrographs indicating Cr2N precipitated in the ferrite phase of the FZ; (a) The surface morphology of as-welded joint observed by SEM after etched by 10% oxalic acid, (b) two-dimensional depth profile of the pits formed on specimen surface (16 lm 16 lm), (c) the corresponding three-dimensional AFM topography and (d) the morphology of Cr2N observed by TEM and the diffraction pattern of Cr2N.
Fig. 3. Typical current density as a function of temperature for the DSS 2205 aswelded joints (LBW) and the DSS 2205 as-received base metal (BM).
dissolution of the a-phase in the FZ. Pits first initiated inside the aphase and then propagated into the a-phase until they reached the c-phase, indicating that pit growth was restrained by the c-phase. The occurrence of pits in FZ after CPT measurement manifested that the FZ with excessive a content with Cr2N precipitation was less resistant. Therefore, it is important to improve the corrosion resistance of FZ with appropriate PWHT process.
Fig. 4. Micrographs of pits in the FZ of the as-welded DSS 2205 specimen after CPT measurements. Austenite (c) is light and ferrite (a) is dark. (a) Stable pits observed by optical microscopy, (b) enlarged same site observed by SEM. pit 1# occurred in the ferrite grain; pit 2# was restrained by the c/a boundary; pit 3# spanned a/c domains with c-phase mostly remained and a-phase severely attacked. (BM is base metal and FZ is fusion zone).
3.3. Influence of PWHT LBW process upsets the balanced two-phase microstructure and the pitting corrosion resistance of DSS 2205 welded specimens. In order to eliminate the harmful effects, short-time PWHT was performed to restore the corresponding properties of welded joints.
3.3.1. Effect of PWHT on microstructure evolution Fig. 5 and Fig. 6 show the microstructure change of DSS 2205 LBW joints after 1 min and 3 min PWHT at different temperatures, respectively. Compared with the as-welded condition shown by Fig. 1b, c and d, more and coarser GBA, WA and IGA can be
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Fig. 5. Typical microstructure of the LBW welded fusion zone after 1 min PWHT at (a) 1020 °C, (b) 1050 °C, (c) 1080 °C, (d) 1100 °C. The formed austenite is presented as GBA (Grain Boundary Austenite), WA (Widmanstätten Austenite) and IGA (Intragranular Austenite).
Fig. 6. Typical microstructure of the LBW welded fusion zone after 3 min PWHT at (a) 1020 °C, (b) 1050 °C, (c) 1080 °C and (d) 1100 °C.
observed. The relationship between a/c volume fraction and the corresponding PWHT parameters was depicted by Fig. 7. For 1 min PWHT, a-phase volume fraction decreased gradually but
was never lower than 50%. While for specimens subjected to 3 min PWHT, a-phase volume fraction decreased from 1030 °C to 1080 °C and then increased slightly from 1080 °C to 1100 °C. The
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short-time PWHT. Fig. 9 presents the equilibrium phase diagram of DSS 2205 obtained by Thermo-Cal software, in which important temperatures could be drawn, such as the phase-balance temperature (Tb), about 1140 °C, at which the volume fraction of a-phase is equal to that of c-phase, Above Tb,a-phase volume fraction increases sharply. Thus, heat treatment above 1140 °C should be avoided in order to improve the c volume fraction in FZ. Whereas below 1000 °C, harmful phases such as sigma phase, chi-phase and chromium nitrides will precipitate in the matrix and the corrosion resistance of welded joint is prone to be compromised. Therefore, it is essential to choose an appropriate annealing temperature between 1020 °C and 1100 °C.
Fig. 7. Relationship between ferrite phase volume fraction of DSS 2205 joint and the short-time PWHT parameters (holding time and heat treatment temperature).
minimum amount of ferrite content, 46%, was obtained at 1080 °C. It is well known that the higher the solution annealing temperature, the more the a-phase. The observed deviation of a-phase volume change from 1020 °C to 1080 °C could be attributed to the extremely unbalanced as-welded microstructure. During the PWHT, the suppressed solid transformation of a ? c was active due to the diffusion of austenite stabilizer (mainly Ni and N) with longer holding time and higher annealing temperatures. It can be seen a balanced state of a/c transformation was reached after 3 min PWHT at 1080 °C, which is related with a balance between a ? c transformation suppressed by fast cooling after welding and activation by PWTH and the thermodynamic c ? a transformation from 1020 °C to 1100 °C. From 1080 °C to 1100 °C, the effect of ascending annealing temperatures became dominant and ferrite content increased as expected. Fig. 8a shows the morphology of DSS 2205 LBW joint after 3 min PHWT at 1080 °C, etched by oxalic acid. No obvious Cr2N related pit was found compared with Fig. 2a. Fig. 8b shows typical a/ c duplex microstructure of DSS 2205 recorded by TEM. Fig. 8c and Fig. 8d, the interpretation of selected area diffraction patterns (SADP), illustrate diffraction patterns of a-phase and c-phase, respectively. Again, there was no obvious Cr2N precipitation compared with Fig. 2d. Thus, SEM combined with TEM characterization suggested the Cr2N formed during welding was dissolved into the matrix after short-time PWHT. Although both the welding and the short-time PWHT are not equilibrium procedures, equilibrium phase diagram is still helpful to understand the microstructure change observed before/after
3.3.2. Effect of PWHT on pitting corrosion resistance We determined the CPT values of LBW welded DSS 2205 specimens after 3 min and 1 min PWHT. The relationship between CPT and PWHT parameters (heat treatment temperature and holding time) was plotted in Fig. 10. Compared with that of as-welded specimens, all heat treated specimens showed enhanced CPT values, indicating the beneficial effect of short-time PWHT. The highest CPT, 55 °C, obtained after 3 min PWHT at 1080 °C was close to the CPT value for the as-received base metal (as shown in Fig. 3). This suggested that the welded joint annealed 3 min at 1080 °C regained the lost pitting corrosion resistance during welding. For 1 min PWHT, CPT values increased gradually with annealing temperature, but the beneficial effect of PWHT was less pronounced. Therefore, it was concluded that (1) short-time PWHT was necessary to improve the pitting corrosion resistance of the LBW welded joint; (2) 3 min PWHT was more effective than 1 min PWHT; (3) the optimal short-time PWHT temperature was 1080 °C. Fig. 11 shows pit morphologies of LBW welded specimens annealed at 1080 °C for 3 min after CPT measurements. Metastable pits were observed occurred both in ferrite (Fig. 11a) and austenite domains (Fig. 11b) indicating that similar corrosion resistance was achieved between ferrite and austenite phases. Besides, no obvious morphology change was found in the base metal after the CPT experiments. 3.3.3. Relationship between pitting resistance and pitting equivalent number of single phase The pitting corrosion behavior of duplex stainless steels was closely related to local chemical composition of alloying elements (such as Cr, Mo and N), namely the pitting resistance equivalent number (PREN) [40]. In addition, it has been acknowledged that the pitting corrosion resistance of DSSs depends on the PREN of the weaker (less resistant) phase. The best pitting corrosion resistance of DSS is reached when a-phase and c-phase exhibit equal or similar PREN values [33,40]. For further understanding the effect of
Fig. 8. Surface morphologies of DSS 2205 LBW welded joint after 3 min PWHT at 1080 °C, (a) SEM image shows no Cr2N related pits, etched by 10% oxalic acid, (b) a and c domains characterized by TEM (c) diffraction pattern of a-phase and (d) diffraction pattern of c-phase.
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Fig. 9. Equilibrium phase diagram of DSS 2205 obtained by Thermo-Cal software. Database: ECFE4.
a-phase and c-phase at 1000–1100 °C (about 10 5 and 10 7 m2s 1, respectively), for 3 min PWHT, the N levels in a-phase were
Fig. 10. Relationship between critical pitting temperature (CPT) and short-time PWHT parameters (annealing temperature and holding time). Measured in 1.0 mol L 1 NaCl solution.
the short-time PWHT on pitting corrosion behavior of LBW welded joint, the local concentrations of alloying elements in FZ and BM after 3 min PWHT at 1050 °C 1080 °C and 1100 °C were measured using a SEM–EDX system and the results were listed in Table 2. Because of the insensitivity of EDX techniques to nitrogen, an approximate calculation was employed to determine the nitrogen content of single phase. Since the very high diffusivity of nitrogen in
assumed to be the saturation value, 0.05 wt.%. And the nitrogen levels in the c-phase were thereby the rest which could be calculated based on the content of nitrogen in the whole alloy and the phase volume fraction [41,42]. Table 2 shows that Cr and Mo were comparatively enriched in a-phase while Ni and N enriched in c-phase during the short-time PWHT [42]. More importantly, the content of alloying elements was influenced by variations of the two phases volume fraction. In the FZ, as the volume fraction of the c-phase increased and that of the a-phase decreased from 1050 °C to 1080 °C, Cr and Mo, the a stabilizers, were enriched in the a-phase. Meanwhile, the concentration of N decreased in the c-phase. In the BM, a dilution of the Cr and Mo in ferrite phase and slightly increase of nitrogen concentration in austenite phase was observed from 1020 °C to 1100 °C. This is in agreement with previous literature [33,42]. Fig. 12 shows the PREN value of ferrite and austenite phases of BM and FZ and the corresponding critical pitting temperature vs. annealing temperatures for 3 min PWHT. The PREN values were calculated by the following equation: PREN = wt.% Cr + 3.3 wt.% Mo + k wt.% N (k = 16–30) [2,37,42]. Given k = 23, as the annealing temperature increased from 1050 °C to 1080 °C, the PREN values of the ferrite phase in FZ, PREN-FZ(a) increased and that of austenite phase PREN-FZ (c) decreased slightly due to the increase of Cr and Mo content in ferrite phase and the decrease of nitrogen in austenite phase. Consequently, the PREN difference between the two phases in FZ decreased. From 1080 °C to 1100 °C, the PREN values difference between the two phases in FZ enlarged since Cr and Mo
Fig. 11. Metastable pits morphologies observed on the surface of the FZ after 3 min PWHT at 1080 °C. (a) Metastable pits occurred in the ferrite phase of the FZ and (b) Metastable pits occurred in the austenite phase of the FZ.
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Y. Yang et al. / Corrosion Science 65 (2012) 472–480 Table 2 Alloying elements concentration of ferrite and austenite phases in FZ (Fusion Zone), BM (Base Metal) after 3 min PWHT at 1050 °C, 1080 °C and 1100 °C. Annealing temperature (°C)
Positions
Phase
Volume fraction (%)
Cr
Mo
Ni
N
1050 °C
FZ
c a c a
47 53 49 51
21.79 22.87 21.50 23.08
2.46 3.53 2.28 3.75
6.02 5.64 6.46 5.01
0.305 0.050 0.290 0.050
c a c a
54 46 48 52
21.99 23.45 21.75 22.89
2.48 3.56 2.32 3.72
6.34 4.96 6.11 5.27
0.271 0.050 0.302 0.050
c a c a
47 53 46 54
21.81 22.80 21.83 22.77
2.44 3.49 2.45 3.51
6.92 5.11 5.91 5.57
0.305 0.050 0.311 0.050
BM 1080 °C
FZ BM
1100 °C
FZ BM
Elements (wt.%)
stainless steel was investigated in 1 mol/L NaCl solution and following conclusions can be drawn:
Fig. 12. PREN-BM(a), PREN-BM (c), PREN-FZ(a), PREN-FZ (c) and CPT as function of different annealing temperatures of 3 min PWHT. (PREN = wt.% Cr + 3.3 wt.% Mo + 23 wt.% N), BM = Base Metal, FZ = Fusion Zone.
diluted in ferrite phase. The PREN-FZ(a) and PREN-FZ(c) were very close at 1080 °C. Fig. 12 also shows that PREN values of ferrite phase in the BM, PREN-BM(a), decreased continually with the increasing temperature. This phenomenon can be explained by a dilution of the ferrite-stabilizing elements (Cr and Mo) with an increase of the ferrite phase volume in BM. At the same time, a slight increase of nitrogen concentration in austenite phase lead to the increasing PREN-BM(c) values. This is in agreement with previous literature [33,42]. Thus, very similar PREN-BM(c), PREN-BM(a), PREN-FZ(c) and PREN-FZ(a) were obtained at 1080 °C, which is consistent with the highest CPT at this temperature. It should be noted that the pitting corrosion resistance of the BM is higher than that of FZ as shown in Fig. 12. Fig. 11a and b show that the metastable pits were observed in both a and c phases FZ, indication a similar pitting corrosion resistance was reached for the two phases in FZ. In summary, the improvement of corrosion resistance of DSS 2205 LBW welded joints can be attributed to the change of PREN(a) and PREN(c) of BM and FZ and the elimination of deleterious intermetallic phase (Cr2N) by short-time PWHT. 4. Conclusions In this work, the effect of short-time PWHT temperature and holding time on microstructure evolution and pitting corrosion behavior of laser beam welded UNS S31803 (DSS 2205) duplex
1. Autogenous LBW resulted in non-equilibrium microstructure in FZ of DSS 2205 with a-phase volume fraction more than 90%. Cr2N precipitated in the a-phase of FZ which has been confirmed by TEM and AFM. 2. The as-welded DSS 2205 samples exhibited severely impaired pitting corrosion resistance with CPT value decreasing from 56 °C (base metal) to 42 °C.Stable pits occurred preferentially at a-phase in the FZ due to the unbalanced microstructure and the Cr depleted zone induced by Cr2N. 3. Short-time PWTH improved the corrosion resistance of the welded joint and more austenite was formed in FZ compared with the as-welded state. For the 3 min PWHT, as the annealing temperature increased in the range of 1020–1100 °C, the ferrite volume fraction in FZ reached its lowest value at 1080 °C. As the annealing temperature increased from 1020 °C to 1080 °C, the CPT value increased to 55 °C, but further increasing the annealing temperature decreased the CPT value. 4. To restore the pitting corrosion resistance of the welded joint, the optimum short-time PWHT parameters were determined: 3 min heat treatment at 1080 °C. This result can be explained by variations of PREN (a) and PREN (c) values using an N factor of 23. The values of PREN (a) and PREN (c) were very close at 1080 °C in both BM and FZ at which the best corrosion resistance of welded joint was obtained.
Acknowledgements The authors would like to thank helpful collaboration of Baosteel. This work is supported by the National Science Foundation of China (Grant Nos. 50871031, 51131008, 51134010 and 51071049) and Industrialization Project of Shanghai New and High technology. References [1] J. Charles, S. Bernhardson, Proc. Conf Duplex Stainless Steels’91, vol. 2, Beaune, France, (1991) 3–48. [2] J.O. Nilsson, Overview of super duplex stainless steel, Mater. Sci. Technol. 8 (1992) 685–700. [3] Y.S. Sato, H. Kokawa, Preferential precipitation site of sigma phase in duplex stainless steel weld metal, Scripta. Mater. 40 (1999) 659–663. [4] V. Muthupandi, P. Bala Srinivasan, S.K. Seshadri, S. Sundaresan, Effect of weld metal chemistry and heat input on the structure and properties of duplex stainless steel welds, Mater. Sci. Eng., A. 358 (2003) 9–16. [5] Z. Cvijovic´, G. Radenkovic´, Microstructure and pitting corrosion resistance of annealed duplex stainless steel, Corro. Sci. 48 (2006) 3887–3906.
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[25] H. Liu, X. Lu, X. Jin, Phase transformation and mechanical properties in laser continuous heat treatment welds, Mater. Des. 32 (2011) 2269–2276. [26] P. Ferro, A. Tiziani, Influence of induction and furnace postweld heat treatment on corrosion properties of SAF 2205 (UNS 31803), Welding J. 87 (2008) 298– 306. [27] H. Tan, Z. Wang, Y. Jiang, Y. Yang, B. Deng, H. Song, J. Li, Influence of welding thermal cycles on microstructure and pitting corrosion resistance of 2304 duplex stainless steels, Corros. Sci. 55 (2012) 368–377. [28] M.C. Young, L.W. Tsay, C.S. Shin, S.L.I. Chan, The effect of short time post-weld heat treatment on the fatigue crack growth of 2205 duplex stainless steel welds, Int. J. Fatigue 29 (2007) 2155–2162. [29] C. Zhong, X. Tang, Y.F. Cheng, Corrosion of steel under the defected coating studied by localized electrochemical impedance spectroscopy, Electrochim. Acta 53 (2008) 4740–4747. [30] C. Zhong, M.F. He, L. Liu, Y.T. Wu, Y.J. Cheng, Y.D. Deng, B. Shen, W.B. Hu, Lower temperature fabrication of continuous intermetallic coatings on AZ91D magnesium alloy in molten salts, J. Alloys Compd. 504 (2010) 377–381. [31] ASTM G150–99(2010), ‘‘Test Method for Electrochemcial Critical Pitting Temperature Testing for Stainless steels’’, West Conshohocken, PA: ASTM, 2010. [32] B. Deng, Y. Jiang, J. Liao, Y. Hao, C. Zhong, J. Li, Dependence of critical pitting temperature on the concentration of sulphate ion in chloride-containing solutions, Appl. Surf. Sci. 253 (2007) 7369–7375. [33] H. Tan, Y. Jiang, B. Deng, T. Sun, J. Xu, J. Li, Effect of annealing temperature on the pitting corrosion resistance of super duplex stainless steel UNS S32750, Mater. Charact. 60 (2009) 1049–1054. [34] S.S.M. Tavares, C. Scandian, J.M. Pardal, T.S. Luz, F.J. da Silva, Failure analysis of duplex stainless steel weld used in flexible pipes in off shore oil production, Eng. Fail. Anal. 17 (2010) 1500–1506. [35] Y. 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) 3756–3763. [36] I. Gilath, J.M. Signamarcheix, P. Bensussan, A comparison of methods for estimating the weld-metal cooling rate in laser welds, J. Mater. Sci. 29 (1994) 3358–3362. [37] S.-T. Kim, S.-H. Jang, I.-S. Lee, Y.-S. Park, Effects of solution heat-treatment and nitrogen in shielding gas on the resistance to pitting corrosion of hyper duplex stainless steel welds, Corros. Sci. 53 (2011) 1939–1947. [38] R. Cervo, P. Ferro, A. Tiziani, F. Zucchi, Annealing temperature effects on superduplex stainless steel UNS S32750 welded joints. II: pitting corrosion resistance evaluation, J. Mater. Sci. 45 (2010) 4378–4389. [39] A.-M. El-Batahgy, Effect of laser welding parameters on fusion zone shape and solidification structure of austenitic stainless steels, Mater. Lett. 32 (1997) 155–163. [40] L. Weber, P.J. Uggowitzer, Partitioning of chromium and molybdenum in super duplex stainless steels with respect to nitrogen and nickel content, Mater. Sci. Eng., A. 242 (1998) 222–229. [41] L.F. Garfias-Mesias, J.M. Sykes, C.D.S. Tuck, The effect of phase compositions on the pitting corrosion of 25 Cr duplex stainless steel in chloride solutions, Corros. Sci. 38 (1996) 1319–1330. [42] L. Zhang, W. Zhang, Y. Jiang, B. Deng, D. Sun, J. Li, Influence of annealing treatment on the corrosion resistance of lean duplex stainless steel 2101, Electrochim. Acta. 54 (2009) 5387–5392.
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Effect of a brief post-weld heat treatment on the microstructure evolution and pitting corrosion of laser beam welded UNS S31803 duplex stainless steel Yanze Yang a, Zhiyu Wang b, Hua Tan a, Jufeng Hong a, Yiming Jiang a, Laizhu Jiang b, Jin Li a,⇑ a b
Department of Materials Science, Fudan University, Shanghai 200433, PR China Research and Development Center, Baosteel Co. Ltd., Shanghai 201900, PR China
a r t i c l e
i n f o
Article history: Received 30 April 2012 Accepted 20 August 2012 Available online 30 August 2012 Keywords: A. Stainless steel B. TEM C. Welding C. Pitting corrosion
a b s t r a c t The effect of laser-beam welding and subsequent short-time post-weld heat treatment at different temperatures and holding time on microstructure evolution and pitting corrosion behavior of UNS S31803 duplex stainless steel was investigated. The results showed the as-welded joint displayed impaired pitting corrosion resistance and that pitting preferentially occurred at ferrite grain in the fusion zone. Shorttime heat treatment enhanced the pitting corrosion resistance of welded joint. Optimal post-weld heat treatment of 3 min heat treatment at 1080 °C was identified at which the joint restored the pitting corrosion resistance lost during welding process. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction Duplex stainless steels (DSSs), which contain approximately equal amounts of ferrite (a) and austenite (c) phases, offer a unique combination of high corrosion resistance and mechanical properties [1–4]. Thus, DSSs are increasingly being used for various applications. It is well established that the good properties of DSS heavily depend on the balanced two-phase microstructure and the absence of deleterious secondary phases such as r-phase, v-phase, carbides and nitrides etc. [5–9]. DSSs generally show impaired corrosion resistance and toughness when welded, especially after laser-beam welding (LBW). LBW offers many advantages over conventional arc welding process such as low heat input, short cycle time and good cosmetic welds which make LBW especially suitable to high volume applications. However, the very fast cooling rate of LBW produces welds with excessive a-phase content. Since the solubility of nitrogen (N) in a-phase is much lower than that in c-phase, an N supersaturation in a-phase is prone to occur. Thus, the LBW process increases the a-phase volume fraction [10–13] and brings on the precipitation of chromium nitrides [14–20] due to the excessive nitrogen content in the ferrite phase as well as the lack of time for a ? c transformation [21,22]. Ramirez et al. [23] suggested that chromium nitrides precipitated at the interior of the ferrite grains of DSS after quenching from 1350 °C. As reported by Perren et al. [24], chromium nitrides lowered the corrosion resistance of DSS by enhancing critical current densities and passivation potential. ⇑ Corresponding author. Tel./fax: +86 21 65643648. E-mail address: [email protected] (J. Li). 0010-938X/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2012.08.054
Since LBW upsets the balanced a/c duplex microstructure and promotes the deleterious phases, it is necessary to eliminate the problems associated with excessive ferrite content and intermetallic phases. Post-weld heat treatment (PWHT) is recommended in practice [25]. PWHT is especially important for chemical and fossil energy industries that require the applications of materials with high corrosion resistance according to the American and European standards (ASTM A928/A928M and NORSOK MDS D42) [26]. Our group has studied the effect of PWHT temperature (1020– 1120 °C) on the pitting corrosion resistance of the plasma-arc welded (PAW) DSS 2304 [27]. The weak region of pitting was found at the heat-affected zone (HAZ) and the corrosion resistance of PAW welded joints was restored after 4 min PWHT at 1080 °C. Investigation performed by Young et al. [28] suggested that short-time PWHT at 1050 °C was able to effectively raise the austenite content and impact the toughness of the DSS welds. In summary, literature review has shown that there would be a strong relationship between the PWHT parameters, microstructure and pitting corrosion behavior of welded joints. However, little work has been carried out regarding the effects of short-time PWHT parameters, in particular annealing temperature and holding time, on the pitting corrosion behavior of the laser-beam welded DSS joints, even though such knowledge is important for improving the efficiency of the LBW process. In the present paper, duplex stainless steel UNS S31803 (DSS 2205) was joined by autogenous laser beam welding (LBW) without filer metal and subsequently annealed at different temperatures ranging from 1020–1100 °C for 1 min and 3 min. Potentiostatic critical pitting temperature technique (CPT) was employed to evaluate the corresponding pitting corrosion resistance, which
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has been proven to be more sensitive and reliable than other pitting characterizing approaches such as pitting potential [29,30]. The microstructure and the precipitation of chromium nitrides were investigated by scanning electron microscope (SEM), atomic force microscopy (AFM) and transmission electron microscope (TEM). The main objectives of this study are to: 1. Reveal the influence of the laser beam welding on the microstructure and pitting corrosion resistance of the DSS 2205 joint. 2. Assess the influence of short-time PWHT parameters (annealing temperature and holding time) on the microstructure evolution and the pitting corrosion behavior of laser beam welded DSS 2205 joint and to explain the mechanism of such behavior. 3. Determine the optimum short-time PWHT temperature and holding time to restore the pitting corrosion resistance lost during welding process. 2. Experimental 2.1. Material and sample preparation DSS 2205 (chemical composition in wt.% C: 0.021, P: 0.008, Ni: 5.72, Cr: 22.36, Mn: 1.52, Si: 0.46, Mo: 3.02, Co: 0.069, S: 0.001, N: 0.17, Fe: bal.) plates of 1.5 mm in thickness were obtained from Baosteel Group Corporation. Welding was performed using autogenous LBW without filler metal and the corresponding parameters were listed in Table 1. Considering the efficiency of practical application, the PWHT holding time should be as short as possible. According to industrial experiences, the holding time of PWHT for 1 mm thick sheet (or pipe) is about 1 min. Therefore, 1 min and 3 min annealing treatment was adopted for the 1.5 mm DSS 2205 samples in this study to evaluate the effectiveness of shorttime PWHT. Short-time PWHT was performed at 1020 °C, 1050 °C, 1080 °C and 1100 °C, respectively in a muffle furnace in atmosphere. Firstly, the muffle furnace was heated to the setting temperature and then the welded joint was put into the furnace and held for 3 min or 1 min at the setting temperature. After that, the welded joint was taken out and quenched in water. The specimens for electrochemical tests and microscopy examination, with dimension of 12 mm 12 mm 1.5 mm were cut from the joint including the fusion zone and the heat-affected zone. 2.2. Electrochemical measurements The pitting corrosion resistance of the specimens were evaluated by the CPT technique in 1 mol/L NaCl solution [31]. All the electrochemical measurements were carried out using a PARSTAT2273 potentiostat in a classical three-electrode cell. In this cell, a platinum foil served as the counter electrode and a saturated calomel electrode (SCE) served as the reference electrode. All potentials quoted in this paper referred to this reference electrode. The specimen acting as working electrode was mounted in epoxy resin. In order to avoid crevices, the area of the specimen near the interface between the resin and the specimen was sealed with
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silica gel sealant [32]. The exposed working electrode surface area was 100 mm2. Prior to each CPT measurement, the working electrode was ground with successive grade silicon carbide sand paper up to 1000 grit, degreased with ethanol, rinsed with distiller water and dried in air. The test solution, 1 mol/L NaCl solution, was made up from analytical grade reagents and distilled water according to the ASTM standard: G150–99 ‘‘Standard Test Method for Electrochemical Critical Pitting Temperature Testing of Stainless steels’’. The CPT measurement procedure is described as follows: (1) the surface of the working electrode was initially polarized cathodically at 0.9 VSCE for 120 s to reduce the oxide film formed in air. (2) working electrode was then stabilized at open circuit potential for 30 min in test solution at 2 °C; (3) a static potential of 0.75 VSCE was applied to working electrode and the electrolyte temperature was increased continuously at a rate of 1 °C/min controlled by a water bath. The corresponding current was recorded simultaneously with the increasing temperature. The experiment was terminated when the current density increased to 100 lA/cm2. The CPT is defined as the temperature at which the current density equals to 100 lA/cm2. The test solution was kept bubbled with pure N2 gas to remove the dissolved O2 throughout the test. All electrochemical tests were repeated at least three times. 2.3. Optical, SEM, TEM and AFM analyses Specimens for optical, SEM and AFM characterization were ground with successive grade abrasive papers up to 1500 grit and then polishing using 1.5 lm diamond paste. Before the observation, the specimens were electrolytically etched by 30 wt.% KOH solution at 2.0 VSCE for 15 s. Quantitative metallography was carried out using light optical microscopy (LOM) and Image-Pro Plus 6.0 software. The volume fractions of the two phases were measured in ten different locations in the same zone and the average value was used in this paper. The chemical compositions of the c-phase and a-phase were analyzed using an energy dispersive X-ray spectroscopy (EDS) linked to a SEM system (SCE Philips Xl 30FEG). Since the concentration of alloying elements may vary at different locations in the same zone, ten measurements were carried out and the average value was used. Two methods were employed to reveal the presence of chromium nitride in the DSS 2205 weld: (1) direct observation by using TEM; (2) observation by using SEM and AFM after electrolytic etching by 10 wt.% oxalic acid solution at 7 VSCE for 30 s [33]. Before TEM observation, thin foils of specimens were prepared following conventional TEM sample preparation methods. The samples were first mechanically polished to approximately thickness of about 0.1 mm followed by polishing at 25 VSCE in a solution of 10% perchloric acid in 90% methyl alcohol with a twin-jet unit. After that the samples were further thinned, the analyses were performed in a transmission electron microscopy (JEM-2010, JEOL) operated at 200 kV. A commercial atomic force microscopy (Dimension Icon, Bruker Corporation) was used to obtain images and topographical information of the studied specimens. The measurements were performed in the contact mode (using SNL cantilevers with a spring constant k = 0.12 N/m), and the AFM images were analyzed by NanoScope Analysis version 1.4 software package.
Table 1 Laser welding parameters used in the experiment. Laser beam configuration
Unit
Value
Laser power Travel speed Spot size on surface Focal position Deviation distance to focus spot Shielding gas flow rate (Ar) Backing gas flow rate (Ar)
W Mm/min Mm2 – Mm L/min L/min
1500 2000 1.5 1.5 At specimen surface 0 15 15
3. Results and discussion 3.1. Influence of LBW on microstructure and pitting corrosion resistance Laser beam welding is a high energy density and low heat input process, resulting in a very small heat affect zone (HAZ) and a deep and narrow fusion zone (FZ) with very little distortion [34].
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3.2. Microstructure Fig. 1 shows the morphologies of different zones of DSS 2205 LBW as-welded joint. Fig. 1a illustrates the typical duplex microstructure of the DSS 2205 base metal. The c-phase was shown as the lighter region while the dark region was the a-phase. The austenite was evenly distributed in the ferrite matrix and no apparent hazard phase was observed. Fig. 1b presents the microstructure of a typical laser beam welded joint. The width of fusion zone (FZ) was about 1 mm and the heat affected zone (HAZ) could hardly be identified. Therefore, special attention was paid to the FZ in this work. Fig. 1c shows the enlarged center part of FZ, 50–100 lm equiaxial ferrite grains were observed. While at the edge of FZ, as shown in Fig. 1d, 200 lm 30 lm columnar ferrite grains were observed grown towards the center of FZ driven by temperature gradient. The effect of the grain size/orientation on the pitting corrosion resistance of FZ was not considered in this paper. It is assumed that there is no obvious difference in the pitting corrosion resistance between different parts of FZ since the pitting corrosion resistance of FZ largely depends on the ferrite phase formed during the very fast cooling stage of the LBW, as will be discussed below. Generally, three types of austenite are formed during the apostsolidification transformation at the cooling stage of welding [35]: (1) austenite formed as allotriomorphs at the prior-ferrite grain boundaries (GBA); (2) austenite formed as Widmanstätten sideplates grown into ferrite grains (WA); (3) austenite formed as intragranular precipitates in ferrite grains (IGA). Fig. 1c shows that the GBA (light) located at the boundaries of ferrite grains due to the nucleation and growth of austenite at the boundaries. The amount of WA and IGA was very small. The average a-phase volume fraction was 93% in FZ (austenite phase volume fraction: 7%). Gilath et al. [36] estimated the magnitude of the cooling rate of LBW FZ and suggested it to be 103 K/s. Due to the very high cooling rate of LBW, the transformation from a-phase to c-phase as well as the diffusion of alloying elements were severely suppressed. In addition to the excessive a-phase, chromium nitride played an
important role in corrosion resistance properties of DSS. Chromium nitride introduced chromium depleted regions that were susceptible to corrosion attack [37]. It was reported that chromium nitrides formed in HAZ and FZ due to the influence of welding thermal cycle [10,12,38]. Fig. 2a shows the morphologies of the as-welded joint etched by oxalic acid. Small pits occurred within ferrite phase, indicating the presence of Cr2N precipitation. AFM depth profile and three dimensional AFM topography (Fig. 2b and Fig. 2c, respectively) reveal that the width of pits was about 0.2–2 lm and the depth of pits was about 100–150 nm. TEM image and selected area diffraction pattern (SADP), as shown in Fig. 2d and Fig. 2e, respectively, provide direct experimental evidence of the hexagonal Cr2N. No obvious intermetallic compound such as r-phase, v-phase and carbides was detected. The length of the needle-like Cr2N was estimated to be 200–400 nm. The observed accumulations and alignment of the Cr2N within ferrite grains (Fig. 2d) corresponded to the high dislocation density of those sites and could be explained through the diffusion impediment of N at these dislocations during rapid cooling. Thus, the pits observed by SEM and AFM characterization were correlated to Cr2N that precipitated in a-phase of FZ during welding process [39]. 3.2.1. Pitting corrosion resistance Fig. 3 presents the typical curves of the current density against solution temperature for the as-received DSS 2205 base metal and the as-welded specimens without PWHT. The CPT decreased from 56 °C to 42 °C after LBW process, indicating a severe degradation of the pitting corrosion resistance. Pit morphologies were characterized by optical/scanning electron microscope, as shown in Fig. 4a and Fig. 4b, respectively. Three stable pits were observed in the FZ of the as-welded specimen. Pit 1# clearly occurred in the aphase; pit 2# occurred in a-phase and was restrained by the c/a boundary; it 3# spanned multiple a/c domains with the c-phase mostly remained and the a-phase severely attacked. Thus, it was found that the mechanism of pitting corrosion was the selective
Fig. 1. Optical micrographs of as-welded DSS 2205 joints etched by 30 wt.% KOH solution showing (a) base metal (BM), (b) fusion zone (FZ), (c) enlarged centre of the fusion zone and (d) edge of fusion zone. Austenite (c) is light and ferrite (a) is dark. The ferrite phase volume content is more than 90% in the fusion zone.
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Fig. 2. SEM, AFM and TEM micrographs indicating Cr2N precipitated in the ferrite phase of the FZ; (a) The surface morphology of as-welded joint observed by SEM after etched by 10% oxalic acid, (b) two-dimensional depth profile of the pits formed on specimen surface (16 lm 16 lm), (c) the corresponding three-dimensional AFM topography and (d) the morphology of Cr2N observed by TEM and the diffraction pattern of Cr2N.
Fig. 3. Typical current density as a function of temperature for the DSS 2205 aswelded joints (LBW) and the DSS 2205 as-received base metal (BM).
dissolution of the a-phase in the FZ. Pits first initiated inside the aphase and then propagated into the a-phase until they reached the c-phase, indicating that pit growth was restrained by the c-phase. The occurrence of pits in FZ after CPT measurement manifested that the FZ with excessive a content with Cr2N precipitation was less resistant. Therefore, it is important to improve the corrosion resistance of FZ with appropriate PWHT process.
Fig. 4. Micrographs of pits in the FZ of the as-welded DSS 2205 specimen after CPT measurements. Austenite (c) is light and ferrite (a) is dark. (a) Stable pits observed by optical microscopy, (b) enlarged same site observed by SEM. pit 1# occurred in the ferrite grain; pit 2# was restrained by the c/a boundary; pit 3# spanned a/c domains with c-phase mostly remained and a-phase severely attacked. (BM is base metal and FZ is fusion zone).
3.3. Influence of PWHT LBW process upsets the balanced two-phase microstructure and the pitting corrosion resistance of DSS 2205 welded specimens. In order to eliminate the harmful effects, short-time PWHT was performed to restore the corresponding properties of welded joints.
3.3.1. Effect of PWHT on microstructure evolution Fig. 5 and Fig. 6 show the microstructure change of DSS 2205 LBW joints after 1 min and 3 min PWHT at different temperatures, respectively. Compared with the as-welded condition shown by Fig. 1b, c and d, more and coarser GBA, WA and IGA can be
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Fig. 5. Typical microstructure of the LBW welded fusion zone after 1 min PWHT at (a) 1020 °C, (b) 1050 °C, (c) 1080 °C, (d) 1100 °C. The formed austenite is presented as GBA (Grain Boundary Austenite), WA (Widmanstätten Austenite) and IGA (Intragranular Austenite).
Fig. 6. Typical microstructure of the LBW welded fusion zone after 3 min PWHT at (a) 1020 °C, (b) 1050 °C, (c) 1080 °C and (d) 1100 °C.
observed. The relationship between a/c volume fraction and the corresponding PWHT parameters was depicted by Fig. 7. For 1 min PWHT, a-phase volume fraction decreased gradually but
was never lower than 50%. While for specimens subjected to 3 min PWHT, a-phase volume fraction decreased from 1030 °C to 1080 °C and then increased slightly from 1080 °C to 1100 °C. The
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short-time PWHT. Fig. 9 presents the equilibrium phase diagram of DSS 2205 obtained by Thermo-Cal software, in which important temperatures could be drawn, such as the phase-balance temperature (Tb), about 1140 °C, at which the volume fraction of a-phase is equal to that of c-phase, Above Tb,a-phase volume fraction increases sharply. Thus, heat treatment above 1140 °C should be avoided in order to improve the c volume fraction in FZ. Whereas below 1000 °C, harmful phases such as sigma phase, chi-phase and chromium nitrides will precipitate in the matrix and the corrosion resistance of welded joint is prone to be compromised. Therefore, it is essential to choose an appropriate annealing temperature between 1020 °C and 1100 °C.
Fig. 7. Relationship between ferrite phase volume fraction of DSS 2205 joint and the short-time PWHT parameters (holding time and heat treatment temperature).
minimum amount of ferrite content, 46%, was obtained at 1080 °C. It is well known that the higher the solution annealing temperature, the more the a-phase. The observed deviation of a-phase volume change from 1020 °C to 1080 °C could be attributed to the extremely unbalanced as-welded microstructure. During the PWHT, the suppressed solid transformation of a ? c was active due to the diffusion of austenite stabilizer (mainly Ni and N) with longer holding time and higher annealing temperatures. It can be seen a balanced state of a/c transformation was reached after 3 min PWHT at 1080 °C, which is related with a balance between a ? c transformation suppressed by fast cooling after welding and activation by PWTH and the thermodynamic c ? a transformation from 1020 °C to 1100 °C. From 1080 °C to 1100 °C, the effect of ascending annealing temperatures became dominant and ferrite content increased as expected. Fig. 8a shows the morphology of DSS 2205 LBW joint after 3 min PHWT at 1080 °C, etched by oxalic acid. No obvious Cr2N related pit was found compared with Fig. 2a. Fig. 8b shows typical a/ c duplex microstructure of DSS 2205 recorded by TEM. Fig. 8c and Fig. 8d, the interpretation of selected area diffraction patterns (SADP), illustrate diffraction patterns of a-phase and c-phase, respectively. Again, there was no obvious Cr2N precipitation compared with Fig. 2d. Thus, SEM combined with TEM characterization suggested the Cr2N formed during welding was dissolved into the matrix after short-time PWHT. Although both the welding and the short-time PWHT are not equilibrium procedures, equilibrium phase diagram is still helpful to understand the microstructure change observed before/after
3.3.2. Effect of PWHT on pitting corrosion resistance We determined the CPT values of LBW welded DSS 2205 specimens after 3 min and 1 min PWHT. The relationship between CPT and PWHT parameters (heat treatment temperature and holding time) was plotted in Fig. 10. Compared with that of as-welded specimens, all heat treated specimens showed enhanced CPT values, indicating the beneficial effect of short-time PWHT. The highest CPT, 55 °C, obtained after 3 min PWHT at 1080 °C was close to the CPT value for the as-received base metal (as shown in Fig. 3). This suggested that the welded joint annealed 3 min at 1080 °C regained the lost pitting corrosion resistance during welding. For 1 min PWHT, CPT values increased gradually with annealing temperature, but the beneficial effect of PWHT was less pronounced. Therefore, it was concluded that (1) short-time PWHT was necessary to improve the pitting corrosion resistance of the LBW welded joint; (2) 3 min PWHT was more effective than 1 min PWHT; (3) the optimal short-time PWHT temperature was 1080 °C. Fig. 11 shows pit morphologies of LBW welded specimens annealed at 1080 °C for 3 min after CPT measurements. Metastable pits were observed occurred both in ferrite (Fig. 11a) and austenite domains (Fig. 11b) indicating that similar corrosion resistance was achieved between ferrite and austenite phases. Besides, no obvious morphology change was found in the base metal after the CPT experiments. 3.3.3. Relationship between pitting resistance and pitting equivalent number of single phase The pitting corrosion behavior of duplex stainless steels was closely related to local chemical composition of alloying elements (such as Cr, Mo and N), namely the pitting resistance equivalent number (PREN) [40]. In addition, it has been acknowledged that the pitting corrosion resistance of DSSs depends on the PREN of the weaker (less resistant) phase. The best pitting corrosion resistance of DSS is reached when a-phase and c-phase exhibit equal or similar PREN values [33,40]. For further understanding the effect of
Fig. 8. Surface morphologies of DSS 2205 LBW welded joint after 3 min PWHT at 1080 °C, (a) SEM image shows no Cr2N related pits, etched by 10% oxalic acid, (b) a and c domains characterized by TEM (c) diffraction pattern of a-phase and (d) diffraction pattern of c-phase.
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Fig. 9. Equilibrium phase diagram of DSS 2205 obtained by Thermo-Cal software. Database: ECFE4.
a-phase and c-phase at 1000–1100 °C (about 10 5 and 10 7 m2s 1, respectively), for 3 min PWHT, the N levels in a-phase were
Fig. 10. Relationship between critical pitting temperature (CPT) and short-time PWHT parameters (annealing temperature and holding time). Measured in 1.0 mol L 1 NaCl solution.
the short-time PWHT on pitting corrosion behavior of LBW welded joint, the local concentrations of alloying elements in FZ and BM after 3 min PWHT at 1050 °C 1080 °C and 1100 °C were measured using a SEM–EDX system and the results were listed in Table 2. Because of the insensitivity of EDX techniques to nitrogen, an approximate calculation was employed to determine the nitrogen content of single phase. Since the very high diffusivity of nitrogen in
assumed to be the saturation value, 0.05 wt.%. And the nitrogen levels in the c-phase were thereby the rest which could be calculated based on the content of nitrogen in the whole alloy and the phase volume fraction [41,42]. Table 2 shows that Cr and Mo were comparatively enriched in a-phase while Ni and N enriched in c-phase during the short-time PWHT [42]. More importantly, the content of alloying elements was influenced by variations of the two phases volume fraction. In the FZ, as the volume fraction of the c-phase increased and that of the a-phase decreased from 1050 °C to 1080 °C, Cr and Mo, the a stabilizers, were enriched in the a-phase. Meanwhile, the concentration of N decreased in the c-phase. In the BM, a dilution of the Cr and Mo in ferrite phase and slightly increase of nitrogen concentration in austenite phase was observed from 1020 °C to 1100 °C. This is in agreement with previous literature [33,42]. Fig. 12 shows the PREN value of ferrite and austenite phases of BM and FZ and the corresponding critical pitting temperature vs. annealing temperatures for 3 min PWHT. The PREN values were calculated by the following equation: PREN = wt.% Cr + 3.3 wt.% Mo + k wt.% N (k = 16–30) [2,37,42]. Given k = 23, as the annealing temperature increased from 1050 °C to 1080 °C, the PREN values of the ferrite phase in FZ, PREN-FZ(a) increased and that of austenite phase PREN-FZ (c) decreased slightly due to the increase of Cr and Mo content in ferrite phase and the decrease of nitrogen in austenite phase. Consequently, the PREN difference between the two phases in FZ decreased. From 1080 °C to 1100 °C, the PREN values difference between the two phases in FZ enlarged since Cr and Mo
Fig. 11. Metastable pits morphologies observed on the surface of the FZ after 3 min PWHT at 1080 °C. (a) Metastable pits occurred in the ferrite phase of the FZ and (b) Metastable pits occurred in the austenite phase of the FZ.
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Y. Yang et al. / Corrosion Science 65 (2012) 472–480 Table 2 Alloying elements concentration of ferrite and austenite phases in FZ (Fusion Zone), BM (Base Metal) after 3 min PWHT at 1050 °C, 1080 °C and 1100 °C. Annealing temperature (°C)
Positions
Phase
Volume fraction (%)
Cr
Mo
Ni
N
1050 °C
FZ
c a c a
47 53 49 51
21.79 22.87 21.50 23.08
2.46 3.53 2.28 3.75
6.02 5.64 6.46 5.01
0.305 0.050 0.290 0.050
c a c a
54 46 48 52
21.99 23.45 21.75 22.89
2.48 3.56 2.32 3.72
6.34 4.96 6.11 5.27
0.271 0.050 0.302 0.050
c a c a
47 53 46 54
21.81 22.80 21.83 22.77
2.44 3.49 2.45 3.51
6.92 5.11 5.91 5.57
0.305 0.050 0.311 0.050
BM 1080 °C
FZ BM
1100 °C
FZ BM
Elements (wt.%)
stainless steel was investigated in 1 mol/L NaCl solution and following conclusions can be drawn:
Fig. 12. PREN-BM(a), PREN-BM (c), PREN-FZ(a), PREN-FZ (c) and CPT as function of different annealing temperatures of 3 min PWHT. (PREN = wt.% Cr + 3.3 wt.% Mo + 23 wt.% N), BM = Base Metal, FZ = Fusion Zone.
diluted in ferrite phase. The PREN-FZ(a) and PREN-FZ(c) were very close at 1080 °C. Fig. 12 also shows that PREN values of ferrite phase in the BM, PREN-BM(a), decreased continually with the increasing temperature. This phenomenon can be explained by a dilution of the ferrite-stabilizing elements (Cr and Mo) with an increase of the ferrite phase volume in BM. At the same time, a slight increase of nitrogen concentration in austenite phase lead to the increasing PREN-BM(c) values. This is in agreement with previous literature [33,42]. Thus, very similar PREN-BM(c), PREN-BM(a), PREN-FZ(c) and PREN-FZ(a) were obtained at 1080 °C, which is consistent with the highest CPT at this temperature. It should be noted that the pitting corrosion resistance of the BM is higher than that of FZ as shown in Fig. 12. Fig. 11a and b show that the metastable pits were observed in both a and c phases FZ, indication a similar pitting corrosion resistance was reached for the two phases in FZ. In summary, the improvement of corrosion resistance of DSS 2205 LBW welded joints can be attributed to the change of PREN(a) and PREN(c) of BM and FZ and the elimination of deleterious intermetallic phase (Cr2N) by short-time PWHT. 4. Conclusions In this work, the effect of short-time PWHT temperature and holding time on microstructure evolution and pitting corrosion behavior of laser beam welded UNS S31803 (DSS 2205) duplex
1. Autogenous LBW resulted in non-equilibrium microstructure in FZ of DSS 2205 with a-phase volume fraction more than 90%. Cr2N precipitated in the a-phase of FZ which has been confirmed by TEM and AFM. 2. The as-welded DSS 2205 samples exhibited severely impaired pitting corrosion resistance with CPT value decreasing from 56 °C (base metal) to 42 °C.Stable pits occurred preferentially at a-phase in the FZ due to the unbalanced microstructure and the Cr depleted zone induced by Cr2N. 3. Short-time PWTH improved the corrosion resistance of the welded joint and more austenite was formed in FZ compared with the as-welded state. For the 3 min PWHT, as the annealing temperature increased in the range of 1020–1100 °C, the ferrite volume fraction in FZ reached its lowest value at 1080 °C. As the annealing temperature increased from 1020 °C to 1080 °C, the CPT value increased to 55 °C, but further increasing the annealing temperature decreased the CPT value. 4. To restore the pitting corrosion resistance of the welded joint, the optimum short-time PWHT parameters were determined: 3 min heat treatment at 1080 °C. This result can be explained by variations of PREN (a) and PREN (c) values using an N factor of 23. The values of PREN (a) and PREN (c) were very close at 1080 °C in both BM and FZ at which the best corrosion resistance of welded joint was obtained.
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