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Microstructure, mechanical and corrosion properties of friction stir welded high nitrogen nickel-free austenitic stainless steel
Microstructure, mechanical and corrosion properties of friction stir welded high nitrogen nickel-free austenitic stainless steel H.B. Li, Z.H. Jiang, H. Feng, S.C. Zhang, L. Li, P.D. Han, R.D.K. Misra, J.Z. Li PII: DOI: Reference:
S0264-1275(15)00453-0 doi: 10.1016/j.matdes.2015.06.103 JMADE 148
To appear in: Received date: Revised date: Accepted date:
8 March 2015 26 May 2015 10 June 2015
Please cite this article as: H.B. Li, Z.H. Jiang, H. Feng, S.C. Zhang, L. Li, P.D. Han, R.D.K. Misra, J.Z. Li, Microstructure, mechanical and corrosion properties of friction stir welded high nitrogen nickel-free austenitic stainless steel, (2015), doi: 10.1016/j.matdes.2015.06.103
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ACCEPTED MANUSCRIPT Microstructure, mechanical and corrosion properties of friction stir welded high nitrogen nickel-free austenitic stainless steel
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College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China;
Laboratory for Excellence in Advanced Steel Research, Center for Structural and Functional Materials Research
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School of Materials and Metallurgy, Northeastern University, Shenyang 110819, China
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H.B. Lia*, Z.H. Jianga, H. Feng a, S.C. Zhanga, L. Lia, P.D. Hanb, R.D.K. Misrac, J.Z. Lid
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and Innovation, and Department of Metallurgical and Materials Engineering, University of Texas at El Paso, 500 W. University Avenue, El Paso, TX 79968, USA China Friction Stir Welding Center, Beijing Aeronautical Manufacturing Technology Research Institute, Beijing
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* Corresponding author: H.B. Li
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100024, China
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*E-mail address: [email protected] (H.B. Li) Tel.: +86 24 83686453; fax: +86 24 23890559 Address: School of Materials and Metallurgy, Northeastern University, Shenyang110819, China
ABSTRACT
Friction stir welding (FSW) was applied to a 2.4 mm thick high nitrogen nickel-free austenitic stainless steel plate using tungsten-rhenium (W-Re) tool. The high-quality weld was successfully produced at a tool rotation speed of 400 rpm and a traveling speed of 100 mm/min. The microstructure, mechanical and corrosion properties of the weld were studied. The nitrogen content of the weld was almost identical to that of base metal (BM). FSW refined grains in the stir zone (SZ) through dynamic recrystallization and led to increase in hardness and tensile strength within the SZ, while the ductility was slightly decreased. The failure of tensile specimens occurred in the BM.
ACCEPTED MANUSCRIPT TEM results revealed precipitates of Cr23C6 of size ~1µm in the SZ, although their content was small. The precipitation of Cr23C6 and increase in δ-ferrite in the SZ led to small decrease in both
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pitting and intergranular corrosion resistance. Keywords: Friction stir welding; High nitrogen austenitic stainless steel; Microstructure;
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Mechanical properties; Corrosion resistance
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1. Introduction
High nitrogen austenitic steels (HNAS) are attractive candidates for the replacement of
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conventional Fe-Cr-Ni austenitic steels [1]. Nitrogen in steels has a significant effect in stabilizing
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austenitic, enabling partial or complete replacement of expensive Ni used to stabilize the austenitic
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phase [2]. Additionally, the synergetic effect of nitrogen and other alloying elements (Cr, Mo, V, etc.) simultaneously improves the mechanical properties and corrosion resistance of steels [3,4]. With the
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development of pressurized smelting equipment, HNAS with excellent comprehensive properties containing more than 1.0 wt.% of nitrogen are being widely considered for transportation, medical applications, power generation, and cryogenic industries, etc. [5-7]. However, the wide application of HNAS for structural applications is largely dependent on its welding characteristics. Conventional fusion welding processes, such as GTAW, TIG, GMAW and laser welding; of HNAS exhibit a number of limitations: (a) loss of nitrogen by desorption of nitrogen or formation of nitrogen associated pore in the weld zone, (b) precipitation of Cr-nitrides and carbides in weld zone and heat-affected zone (HAZ), (c) solidification cracks in weld zone and liquation cracks in the HAZ [8-10]. When welding of HNAS with nitrogen content exceeding the nitrogen solubility under atmospheric pressure is carried out using fusion welding methods, the weld defects are prominent, and lead to significant deterioration in mechanical properties and
ACCEPTED MANUSCRIPT corrosion resistance of the weld. To alleviate and even avoid problems associated with the conventional fusion welding processes
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of HNAS, solid-state joining technique is considered appropriate. Friction Stir Welding (FSW) was invented as an innovative solid-state welding technology by TWI in 1991[11]. In the FSW process,
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a rotating tool containing a pin and a shoulder is plunged into the joint between two workpieces,
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generating heat by friction. Once the material is heated to a visco-plastic state, the tool is translated
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along the weld line. Plasticized base material passes around the tool, where it is consolidated because of the force applied by the shoulder of the tool. Finally, it leaves a solid phase bond
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between the two pieces. Initially, FSW was successfully applied to metals with moderate melting
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point, such as Al, Cu and Mg alloys [12-14]. The ability to use FSW for joining of high
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softening-temperature materials has recently been brought to fruition because of the development of specialized tool material [15,16]. The feasibility of FSW for steels, Ti alloys, and Ni-base alloys has
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been reported to exhibit sound weld joint [17-20]. In recent years, several studies on the FSW of HNAS have been carried out [21-23]. Park et al. [21] conducted a feasibility work and demonstrated that the defects associated with melting/solidification phenomena in fusion welding process of the HNAS joint can be eliminated. Miyano et al. [22] also performed FSW of 2-mm thick Fe-23Cr-0Ni-1Mo-1N austenitic steel using different parameters. Although optimum welding parameters were determined in relation to the mechanical properties of the joint, the microstructural characteristics were not explored. D. Wang et al. [23] carried out a detailed examination of the microstructure and the mechanical properties of the friction stir welded Fe-18.4Cr-15.8Mn-2.1Mo-0.69N-0.04C austenitic stainless steel. The authors reported the FSW significantly refined the austenitic grains and increased hardness and strength in the stir zone (SZ).
ACCEPTED MANUSCRIPT In the present study, a high nitrogen nickel-free austenitic stainless steel with 0.96 wt.% nitrogen was processed with FSW. The microstructural evolution, mechanical, and corrosion properties were
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studied to reveal the microstructure and properties of friction stir welded HNAS.
2.1 Material preparation and friction stir welding process
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2. Experimental procedure
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The HNAS material used in this study was made using vacuum induction furnace and electroslag
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remelting furnace in nitrogen atmosphere [24]. The chemical composition of the steel is listed in Table 1. The experimental cold rolled plate was of dimensions 150 mm×60 mm×2.4 mm. The plates
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were solution treated at 1100 °C for 90 min, followed by water quenching and then ground to
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remove the oxide and contamination of both the top surface and joint surface. In order to assist tool
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penetration as well as reduce wear of the tool, a pilot hole whose size was similar to the tool probe was drilled at the start position of the plate.
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The welding process was carried out using an FSW machine (FSW-LM-AL25) developed by CFSWC at a traveling speed of 100mm/min and a rotational speed of 400 rpm. The tungsten-rhenium (W-Re) tool used during FSW process had a convex shoulder of 15 mm in diameter and a tapered probe. The pin was tapered from 8 mm at the shoulder to 4.5 mm with a length of 2.3 mm. During the welding process, the tilt angle was 0° and downward force was controlled at 20KN. The backing plate material was tungsten alloy, which has high thermal conductivity. An argon shielding gas at the flow rate of 21L/min was employed to avoid surface oxidation. 2.2 Macro- and microstructural evolution Following FSW, the joint was first evaluated on the basis of visual and dye penetration inspection to reveal the weld defects. Then, specimens were extracted from various locations of the weld, as
ACCEPTED MANUSCRIPT shown in Fig. 1 and from the base metal plate. The nitrogen content of the SZ and the BM was detected by Nitrogen/oxygen analyzer (LECO TC 500) to determine whether nitrogen desorption
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occurred in the weld. The metallographic specimens were prepared perpendicular to the weld direction and
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mechanically ground with SiC paper to 2000 grade and polished with 2.5 µm diamond. The samples
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were etched using 10 wt. % oxalic acid and then examined by optical microscope (OM, Olympus
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DSX 500) and scanning electron microscope (SEM, Carl Zeiss Ultra Plus). The grain size was measured with the mean liner intercept (MLI) technique. The specimens from the SZ for electron
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back-scattered diffraction (EBSD) in SEM were prepared by electrolytic polishing in a solution of
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10 vol. % perchloric acid in ethanol. Thin disks of diameter of 3 mm were cut from the SZ parallel
at 200 kV.
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2.3 Mechanical testing
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to the weld, and jet polished for transmission electron microscopy (TEM, TECNCI G2 20) studies
The mechanical properties were studied using combination of Vickers hardness (FUTURE-TECH, FM-700) and transverse tensile test (SANS-CMT5105). The hardness profile was measured at 1mm deep from the surface with regular intervals of 0.2 mm using a load of 0.5 N for 10 s across the cross-section of the weld. The configuration of a representative tensile specimen is shown in Fig. 2. Tensile specimens were ground to remove the surface ripples and flash to avoid notch effect. Three tests were conducted at a constant cross head speed of 3.0 mm/s at room temperature. The fracture surface specimen was observed using SEM. 2.4 Electrochemical measurements For corrosion tests, specimens of 8.1 mm×8.1 mm were machined from the weld joint and the base material. The passivation treatment of the specimens was performed in 20-30 wt.% nitric acid
ACCEPTED MANUSCRIPT solution at 50℃ for 1 h to prevent the occurrence of crevice corrosion. The electrodes were mounted in epoxy resin with an exposed area of 0.64 cm2 and polished with SiC paper from 200 to
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2000 grits. All the electrochemical measurements were carried out using a PARSTAT 2273 that comprised of three electrodes. The BM and weld specimens were the working electrodes, and a
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platinum foil (Pt) and a saturated calomel electrode (SCE) were used as the counter and reference
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electrodes, respectively. The potentiodynamic polarization curves were obtained in 3.5 wt% NaCl
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solution at 25 °C ±0.5 °C from -0.2 V below the OCP at a scan rate of 20 mV min-1 to the anodic direction.
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Meanwhile, intergranular corrosion (IGC) resistance was examined through the double loop
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electrochemical potentiokinetic reactivation (DL-EPR). The testing was carried out at a scanning
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rate of 1.6667 V/s at 25℃±0.5℃ in 2 M H2SO4+1 M NaCl+0.01 M KSCN solution. The solution was deaerated by purging with N2 for 30 min before the test, and a nitrogen atmosphere was
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maintained during the test. The IGC susceptibility of the steel was evaluated by comparing the ratio of the peak reactivation current density Ir and the peak activation current density Ia. All the measurements were repeated at least three times. 3. Results and discussion
3.1 Joint appearance and inspection result Fig. 3 shows a typical joint appearance processed at welding speed of 100 mm/min and tool speed rotating of 400 rpm from the top surface. No groove-like defects [25] were observed along the weld line, indicating that the heat input was sufficient to ensure the material flow during FSW process [26,27]. Dye penetrant inspection was conducted to reveal surface cracks and the inspection results confirmed the absence of surface cracks in the weld. It should be pointed out that the dyed parts correspond to fin and the weld center line exceeded the welding range. Considering that only
ACCEPTED MANUSCRIPT surface-breaking defects can be detected with dye penetrant inspection, the possibility of internal defects was studied using cross-sectional microscopy and mechanical properties tests.
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The nitrogen content in SZ detected by Nitrogen/oxygen analyzer was 0.94 wt.%, which is similar to that of BM (0.96 wt.%), this implied that the SZ of the FSW processed HNAS did not
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exhibit nitrogen loss. Several authors have reported similar results [21,23]. Compared with other
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welding methods, FSW is an effective joining process to prevent nitrogen desorption.
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3.2 Microstructural evolution during FSW
The cross-sectional macrograph of the joint was consistent with the size of pin as shown in Fig.
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4(a). The joint had no internal defects, such as cracks, pores or cavities, which were often observed
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in other FSW studies [21,28] and fusion welds [8,9]. The microstructure was characterized by the
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stir zone (SZ), heat-affected zone (HAZ), thermo-mechanically affected zone (TMAZ) and base metal (BM). The TMAZ can be classified into TMAZ-AS and TMAZ-RS according to the
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advancing side and retreating side of the rotating tool, which are related to the left and right sides of the weld center.
Fig. 4(b) and Fig. 4(c) are the optical micrographs of BM and HAZ, respectively, which consist of austenitic and twins inside austenitic grains. The average grain size of BM was ~24.7 µm. HAZ had a microstructure similar to that of BM and the parent grains did not experience significant growth. This can be explained by the lower peak welding temperature and high cooling rate [22,29]. The cooling speed was very high, not only because the plate was thin, but also because of the fast-flow of argon shielding gas, which took away heat. On the other hand, the SZ consisted of very fine equiaxed grains, as shown in Fig. 4(d), and the average grain size measured at the middle of the SZ was ~6.3 µm. The fine equiaxed grains were produced by dynamic recrystallization induced by frictional heat and plastic deformation [21,30]. The average grain size of SZ was significantly
ACCEPTED MANUSCRIPT smaller than the BM due to dynamic recrystallization. Similar to previous studies, the grains in the SZ contained few twin boundaries caused by the stir function of the tool [23,31].
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Based on earlier studies [25,32], the TMAZ/SZ boundary on the advancing side of the tool was clearly sharper than that on the retreating side. In the present study there was no apparent difference
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between the two boundaries in the macrograph (Fig. 4(a)). However, as regards the microstructure
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of TMAZ-RS and TMAZ-AS (Fig. 4(e) and (f)), the asymmetry of FSW was clearly observed. The
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grains in the TMAZ-AS were severely-elongated and re-oriented perpendicular (arrowed) to the taper direction of the tool. The boundary between TMAZ-RS/SZ was not distinct because of weaker
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mechanical deformation compared to the TMAZ-AS [33].
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Fig. 5 is the SEM micrograph near the top surface and bottom of the SZ. It can be clearly seen
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that there was decrease in grain size from top to the bottom direction. This was attributed to the inhomogeneous distribution of the peak temperature and different cooling speed in the SZ, as
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previously reported [34]. The higher temperature and therefore the higher cooling speed contributed to remarkable grain growth at the top of the SZ [26,35]. TEM studies were carried out to study the microstructure in detail and are presented in Fig. 6. The equiaxed grains in the SZ were relatively small, and there was obvious difference in the dislocation density. Fine equiaxed grains with very low dislocation density were observed as shown in Fig. 6(a). This indicated that dynamic recrystallization occurred in the SZ, which was induced by the intense plastic deformation and frictional heating during the FSW progress. The positive effect of Mo and lower stacking fault energy due to nitrogen addition can promote dynamic recrystallization [36,37]. Moreover, the SZ had a grain structure with a relatively high density of sub-boundaries and dislocation pile up and tangles (Fig. 6(b), Fig. 6(c)), which indicated incomplete dynamic recrystallization in the SZ. The results confirmed that the density of dislocation of grains
ACCEPTED MANUSCRIPT in the SZ was inhomogeneous, and in general, the dislocation density was higher. The results in the present study were consistent with the typical microstructure in the FSW of other steels [23,31].
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Additionally, the precipitation of Cr23C6 of size of ~1 µm was occasionally observed in the SZ, as shown in Fig. 6(d). To clearly understand the variation in the distribution of phases in equilibrium
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with temperature and phase transformation, the phase diagram of the steel was calculated with
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Thermo-calc software. The vertical line in Fig. 7 denotes the chemical composition of the studied
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steel. It should be noted that the precipitation of intermetallic phases, such as Cr2N and σ phase were not observed in the SZ. This may be attributed to the short thermal cycle and different crystal
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structure of precipitations. During the FSW process, argon gas shielding device as oxidation
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protection and cooling method moved with FSW tool. The cooling intensity was higher at the high
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temperature range (800 ℃ ~900 ℃ ), which shortened the residence time, thus prevented precipitation of the Cr2N and σ phase. In the lower temperature range below 800 ℃, the cooling
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speed was relatively lower and the residence time was prolonged due to the decrement of argon cooling intensity. The Cr2N,σ and Cr23C6 phases at the temperature range below 800℃ could be prone to nucleate and preferentially precipitate than Cr2N,σ phase at temperature range from 800℃ to 900 ℃. In addition, the crystal structure of Cr23C6 phase is fcc, and Cr2N and σ phases are hexagonal and tetragonal structure, respectively [38]. Due to the same crystal structure with austenitic matrix, the Cr23C6 phase more preferentially precipitated. Although the introduction of high strain and dynamic recrystallization can accelerate mutual diffusion of the alloying elements [39], the short thermal cycle during FSW process allowed insufficient time for precipitation to nucleate and grow, so only less amount and small size Cr23C6 was found. Fig. 8(a) and Fig. 8(b) are the EBSD maps of BM and SZ, respectively. The white spots distributed in the colored matrix (austenitic grains) represent ferrite. The ferrite phase can be clearly
ACCEPTED MANUSCRIPT observed before and after FSW process, and furthermore increases after FSW process. The remaining ferrite in the BM was mainly attributed to element segregation such as N, Cr and Mo [40].
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Some δ-ferrite was locally formed in HNAS with high nitrogen content of 0.96 wt.%, when produced by electroslag remelting under nitrogen atmosphere. Moreover, it is noteworthy that the
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SZ contains more ferrite than the BM. This suggests that δ-ferrite in the SZ not only originated from
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the remaining ferrite in the BM, but also from the transformation of the austenitic matrix during the
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FSW process.
The local nitrogen content can be less than 0.96 wt.% as a result of element segregation, which
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resulted in a deviation from the equilibrium transformation towards lower nitrogen. At the same
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time, the temperature also should be taken into consideration during phase transformation process.
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Miyano et al. [22] measured the thermal cycle at the back of 2 mm thick Fe-23Cr-0Ni-1Mo-1N stainless steel during FSW at identical welding parameters, as used in the present study (400 rpm,
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100 mm/min). It was found that the temperature of the bottom surface in the SZ could reach ~1020 °C and the upper zone of SZ was estimated to be at slightly at higher temperature. Considering the plate used in this study is slightly thicker than pervious studies, the SZ should experience a higher temperature. Thus, considering the deviation from the equilibrium transformation, the temperature in the SZ during the FSW process presumably reached the temperature of γ- austenitic and δ-ferrite coexistence region, as shown in red shadow area in Fig. 7. The phase transformation to form new δ-phase took place during FSW process, such that the amount of δ-phase was increased. Previous studies also suggested that δ-ferrite can form in the SZ of austenitic steels [23,31]. 3.3 Mechanical properties The hardness distribution in the welded sample as a function of distance from the weld center is
ACCEPTED MANUSCRIPT presented in Fig. 9. The hardness distribution was roughly symmetrical. The hardness increased from the as-received BM (ranged from 360 to 370HV) region toward the TMAZ, reaching a
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maxima at SZ (ranged from 390 to 400HV) with a width of ~8 mm. There is not an obvious fluctuation in both the BM and the SZ, possibly due to the uniform distribution of grain sizes in the
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measured direction. The increase in the hardness was apparently caused by the refinement of grains
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during dynamic recrystallization in the SZ according to the well-known Hall-Petch relationship. In
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addition, the high density dislocations formed in the severely deformed SZ was also responsible for the increment of hardness [31]. Even though the phase change, i.e. the austenitic - ferrite
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transformation due to the element segregation took place, the hardness value may not be obviously
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affected by the formation of the discontinuous ferrite because of their extremely small amount and
of the material.
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grain size. Thus, the grain refinement and high density dislocations were the major hardening factor
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The stress-strain curves and test results obtained from the tensile test of both BM and FSW joint are plotted in Fig.10(a) and Fig.10(b). Similar to the hardness results, transverse 0.2% offset yield (770MPa) and ultimate tensile strengths (1100MPa) of the FSW joint was overmatched relative to the BM. It is confirmed that no small defects were formed in the welds due to the sufficient heat input. However, the ductility of the FSW joint was lower compared with BM. The SZ zone exhibited a smaller amount of deformation compared with the BM during test, which meant lower percentage of elongation. This can be explained by the following reasons. Firstly, it is usually right and proper that grain refinement can simultaneously improve the strength and toughness for most of metal material. But lots of literatures reported the ductility of high nitrogen steels and nitrogen containing austenitic stainless steel decreased with grain refinement [41,42], which was contrary to the traditional fine grain strengthening theory. So for HNS FSW joint, the grain refinement in SZ
ACCEPTED MANUSCRIPT may result in the decrement of toughness. High density dislocations in the SZ generated by mechanical deformation were the key factor decreasing the percentage of elongation [43]. In
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addition, the presence of less amount and small size of Cr23C6 carbides exhibits a certain extent effect on the decrement of percentage of elongation [44].
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Moreover, the tensile samples fractured in the BM near the TMAZ (inserted figure in Fig. 10(b)).
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Examination of the fracture surface BM (Fig. 10(c)) and FSW joint (Fig. 10(d)) by SEM exhibited
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dimples with a feature microvoid coalescence, which indicated characteristics of ductile fracture. But there were some differences between them, the size distribution of dimple for FSW joint was
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uneven than that of BM, and the average size of the dimple for BM was larger than that of SZ. The
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results of morphology of fracture surfaces represented that the BM exhibits higher toughness than
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FSW joint. Additionally, the unevenness of fracture near TMAZ for FSW joint indicated the microstructure near TMAZ was inhomogeneous.
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3.4 Corrosion resistance
The electrochemical polarization curves of BM and SZ in 3.5% NaCl solution are presented in Fig. 11, where the plots have similar shape. The plots had a wide range of passive region, and similar results were obtained in our previous research [45]. The corrosion potential (Ecorr) and the corrosion current densities (icorr) were similar. However, the pitting potential (Eb,10) value of SZ was marginally lower, which implied lower pitting resistance. In general, localized corrosion resistance of welds is affected by microstructural changes according to previous research. On the one hand, the precipitation of secondary phase, such as sigma, Cr2N, Cr23C6, would lead to generation of Cr-depleted zone, which is sensitive to corrosion [1,4]. Ha et al. [1] reported that the pitting corrosion resistance of HNAS was deteriorated by lamellar Cr2N precipitate formed along grain boundaries. On the other hand, severe plastic
ACCEPTED MANUSCRIPT deformation and dynamic recrystallization contributed to the formation of fine grains in SZ, which can enhance the corrosion resistance [46-48]. Sarlak et al. [46] reported that corrosion performance
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of the stir zones of the lean duplex steel were not deteriorated, and even exhibited superior corrosion resistance than the base metal.
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In the present study, although precipitation of Cr23C6 occurred, the amount of precipitation was
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less, may exhibit little influence the corrosion resistance. Thus, considerable amount of δ-ferrite
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must be taken into account for the decrement of the pitting corrosion resistance. Fig. 12 represents SEM morphologies of δ-ferrite in BM and SZ and SEM-EDS line profile through δ-ferrite phase.
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The SEM-EDS results showed that the Cr and Mo contents were higher than that of austenite, and
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Mn and N were lower. The formation of Cr-depletion zone was found at the interface between the
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austenitic and δ-ferrite as shown Fig. 12(c). Several authors have reported that the formation of a Cr-depleted zone promoted pitting corrosion in the SZ [49,50]. Thus, the decrease of Eb,10 of the SZ
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compared to the BM was attributed to the formation of a Cr-depleted zone due to δ-ferrite formation, regardless of grain refinement.
The EPR results of the BM and SZ are presented in Fig. 12. The activation peaks of both the samples were nearly same, and the reactivation peaks could hardly be observed. Since the solution treated BM appeared desensitized, the reactivation current density was much lower than that of the activation. From the local amplification of reactivation peaks in inserted graph, the reactivation peak of the SZ was slightly higher than that of BM, and Ir/Ia ratio for the BM and the SZ were 0.596 and 1.089, respectively. This indicated that the intergranular corrosion of the SZ was more sensitive than the BM. It is well-known that the sensitivity to intergranular corrosion of stainless steel is closely related to the depletion of chromium along the grain boundarieds [51,52]. In present study, the intergranular
ACCEPTED MANUSCRIPT corrosion was attributed to the formation of δ-phase and a small quantity of Cr23C6, which led to the formation of Cr-depleted zone at the interface between the austenitic and δ-ferrite as shown
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Fig.12(c). Thus, the formation of δ-ferrite led to a lower degree of intergranular corrosion resistance of the SZ [50,53].
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4. Conclusions
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In this study, the FSW of a high-nitrogen nickel-free austenitic stainless steel plate was performed
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at a tool rotation speed of 400rpm and welding speed of 100mm/min. Microstructural characteristics, mechanical and corrosion properties were explored. The results of the study can be summarized as
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follows:
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(1) A sound FSW joint was acquired without any defects. The SZ did not show nitrogen loss, and
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was characterized by fine-grained microstructure. (2) Based on TEM and EBSD studies, a small amount of Cr23C6 was observed and amount of
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δ-ferrite increased in SZ during FSW process. (3) The hardness of SZ was slightly higher than the BM. The tensile strength and yield strength of the joint were higher than the BM, while the elongation was slightly lower. (4) Based on measurements of pitting and intergranular corrosion resistance, the corrosion resistance of SZ was slightly decreased in comparison to the BM.
Acknowledgments The present research was financially supported by National Natural Science Foundation of China (No.51304041, 51434004 and U1435205) and Program for New Century Excellent Talents in University (No. N130502001). RDKM gratefully acknowledges support from University of Texas at El Paso, USA. References
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ACCEPTED MANUSCRIPT residual stress in friction stir welding of 304L stainless steel. Mater. Des. 2014;55:812–20. [34] Schmidt H, Hattel J, Wert J. An analytical model for the heat generation in friction stir welding.
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ACCEPTED MANUSCRIPT [44] Zheng LG, Hu XQ, Kang XH, Li DZ. Precipitation of M23C6 and its effect on tensile properties of 0.3C–20Cr–11Mn–1Mo–0.35N steel. Mater. Des. 2015;78:42–50.
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stainless steel. Mater. Des. 2015;66:209–16.
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Figure and Table Captions
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Table 1 Chemical composition of base material (wt.%)
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Fig. 2. Configuration of transverse tensile specimen (mm)
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Fig. 1. Scheme of extraction of samples (A-EBSD specimen, B-Metallographic specimen, C-Tensile specimen, D-Nitrogen-oxygen analysis specimen, E-TEM specimen, F-Electrochemical corrosion testing specimen, G- Hardness specimen; mm)
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Fig. 3. (a) Appearance of FSW zone and (b) dye penetration inspection result
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Fig. 4. (a) Optical macrograph of transverse section of the joint and microstructure indicated in the macrograph, (b)BM, (c) HAZ, (d) SZ, (e)TMAZ-AS, (f)TMAZ-RS Fig. 5. SEM micrographs of the SZ (a) near the top surface and (b) at the bottom
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Fig. 7. The equilibrium phase diagram calculated with Thermo-calc software Fig. 8. Phase maps of the weld: (a)BM, (b)SZ Fig. 9. Microhardness profile of FSW HNAS joint Fig. 10. (a) stress-strain curves and macrograph of the fractured FSW joint (inserted figure), (b) Results of tensile test, (c) SEM morphologies of fracture surfaces (BM), (d) SEM morphologies of fracture surfaces (FSW joint) Fig. 11. Potentiodynamic polarization curves of BM and SZ in 3.5wt% NaCl solution Fig. 12. SEM morphologies of δ-ferrite in BM (a) and SZ(b), (c) SEM-EDS line profile through δ-ferrite phase Fig. 13. EPR curves of the BM and SZ in 2 M H2SO4+1 M NaCl+0.01 M KSCN
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Advancing side
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Rolling and welding direction
Tool rotation direction
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Fig. 1. Scheme of extraction of samples (A-EBSD specimen, B-Metallographic specimen, C-Tensile specimen, D-Nitrogen-oxygen analysis specimen, E-TEM specimen, F-Electrochemical corrosion testing specimen, GHardness specimen; mm)
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Fig. 2. Configuration of transverse tensile specimen (mm)
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Retreating side
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HAZ
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Advancing side
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Fig. 4. (a) Optical macrograph of transverse section of the joint and microstructure indicated in the macrograph, (b)BM, (c) HAZ, (d) SZ, (e)TMAZ-AS, (f)TMAZ-RS
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Fig. 5. SEM micrographs of the SZ (a) near the top surface and (b) at the bottom
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Fig. 6. TEM images of SZ: (a) Fine grains with very low dislocation density, (b) high density dislocation network structure, (c) grain structure with relatively high density of subboundaries and dislocations, (d) precipitation of Cr23C6
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Fig. 7. The equilibrium phase diagram calculated with Thermo-calc software
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(a)
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Fig. 8. Phase maps of the weld: (a)BM, (b)SZ
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Fig. 9. Microhardness profile of FSW HNAS joint
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(a)
Fracture location(BM)
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The joint
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Fig. 10. (a) stress-strain curves and macrograph of the fractured FSW joint (inserted figure), (a)Results of tensile test, (c) SEM morphologies of fracture surfaces (BM), (d) SEM morphologies of fracture surfaces (FSW joint)
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Fig. 11. Potentiodynamic polarization curves of BM and SZ in 3.5wt% NaCl solution
δ-ferrite
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δ-ferrite
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Fig. 12. SEM morphologies of δ-ferrite in BM (a) and SZ(b), (c) SEM-EDS line profile through δ-ferrite phase
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Fig. 13. EPR curves of the BM and SZ in 2 M H2SO4+1 M NaCl+0.01 M KSCN
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Graphical abstract
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Si
Cr
Mn
Mo
Ni
HNAS 0.058 0.19 19.55 19.5 2.26
--
S
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P
Al
O
N
0.003 0.015 0.04 0.0048 0.96
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Steel
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Table 1 Chemical composition of base material (wt.%)
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Highlights
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1. Friction stir welding (FSW) was successfully applied to a high nitrogen nickel-free austenite steel with 0.96 wt.% nitrogen. 2. The small and little amount Cr23C6 was found. 3. The increment of the amount of δ-ferrite in the SZ was discussed. 4. The corrosion results revealed the SZ exhibited excellent corrosion resistance.
S0264-1275(15)00453-0 doi: 10.1016/j.matdes.2015.06.103 JMADE 148
To appear in: Received date: Revised date: Accepted date:
8 March 2015 26 May 2015 10 June 2015
Please cite this article as: H.B. Li, Z.H. Jiang, H. Feng, S.C. Zhang, L. Li, P.D. Han, R.D.K. Misra, J.Z. Li, Microstructure, mechanical and corrosion properties of friction stir welded high nitrogen nickel-free austenitic stainless steel, (2015), doi: 10.1016/j.matdes.2015.06.103
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ACCEPTED MANUSCRIPT Microstructure, mechanical and corrosion properties of friction stir welded high nitrogen nickel-free austenitic stainless steel
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College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China;
Laboratory for Excellence in Advanced Steel Research, Center for Structural and Functional Materials Research
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School of Materials and Metallurgy, Northeastern University, Shenyang 110819, China
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a
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H.B. Lia*, Z.H. Jianga, H. Feng a, S.C. Zhanga, L. Lia, P.D. Hanb, R.D.K. Misrac, J.Z. Lid
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and Innovation, and Department of Metallurgical and Materials Engineering, University of Texas at El Paso, 500 W. University Avenue, El Paso, TX 79968, USA China Friction Stir Welding Center, Beijing Aeronautical Manufacturing Technology Research Institute, Beijing
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* Corresponding author: H.B. Li
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100024, China
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*E-mail address: [email protected] (H.B. Li) Tel.: +86 24 83686453; fax: +86 24 23890559 Address: School of Materials and Metallurgy, Northeastern University, Shenyang110819, China
ABSTRACT
Friction stir welding (FSW) was applied to a 2.4 mm thick high nitrogen nickel-free austenitic stainless steel plate using tungsten-rhenium (W-Re) tool. The high-quality weld was successfully produced at a tool rotation speed of 400 rpm and a traveling speed of 100 mm/min. The microstructure, mechanical and corrosion properties of the weld were studied. The nitrogen content of the weld was almost identical to that of base metal (BM). FSW refined grains in the stir zone (SZ) through dynamic recrystallization and led to increase in hardness and tensile strength within the SZ, while the ductility was slightly decreased. The failure of tensile specimens occurred in the BM.
ACCEPTED MANUSCRIPT TEM results revealed precipitates of Cr23C6 of size ~1µm in the SZ, although their content was small. The precipitation of Cr23C6 and increase in δ-ferrite in the SZ led to small decrease in both
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pitting and intergranular corrosion resistance. Keywords: Friction stir welding; High nitrogen austenitic stainless steel; Microstructure;
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Mechanical properties; Corrosion resistance
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1. Introduction
High nitrogen austenitic steels (HNAS) are attractive candidates for the replacement of
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conventional Fe-Cr-Ni austenitic steels [1]. Nitrogen in steels has a significant effect in stabilizing
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austenitic, enabling partial or complete replacement of expensive Ni used to stabilize the austenitic
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phase [2]. Additionally, the synergetic effect of nitrogen and other alloying elements (Cr, Mo, V, etc.) simultaneously improves the mechanical properties and corrosion resistance of steels [3,4]. With the
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development of pressurized smelting equipment, HNAS with excellent comprehensive properties containing more than 1.0 wt.% of nitrogen are being widely considered for transportation, medical applications, power generation, and cryogenic industries, etc. [5-7]. However, the wide application of HNAS for structural applications is largely dependent on its welding characteristics. Conventional fusion welding processes, such as GTAW, TIG, GMAW and laser welding; of HNAS exhibit a number of limitations: (a) loss of nitrogen by desorption of nitrogen or formation of nitrogen associated pore in the weld zone, (b) precipitation of Cr-nitrides and carbides in weld zone and heat-affected zone (HAZ), (c) solidification cracks in weld zone and liquation cracks in the HAZ [8-10]. When welding of HNAS with nitrogen content exceeding the nitrogen solubility under atmospheric pressure is carried out using fusion welding methods, the weld defects are prominent, and lead to significant deterioration in mechanical properties and
ACCEPTED MANUSCRIPT corrosion resistance of the weld. To alleviate and even avoid problems associated with the conventional fusion welding processes
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of HNAS, solid-state joining technique is considered appropriate. Friction Stir Welding (FSW) was invented as an innovative solid-state welding technology by TWI in 1991[11]. In the FSW process,
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a rotating tool containing a pin and a shoulder is plunged into the joint between two workpieces,
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generating heat by friction. Once the material is heated to a visco-plastic state, the tool is translated
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along the weld line. Plasticized base material passes around the tool, where it is consolidated because of the force applied by the shoulder of the tool. Finally, it leaves a solid phase bond
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between the two pieces. Initially, FSW was successfully applied to metals with moderate melting
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point, such as Al, Cu and Mg alloys [12-14]. The ability to use FSW for joining of high
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softening-temperature materials has recently been brought to fruition because of the development of specialized tool material [15,16]. The feasibility of FSW for steels, Ti alloys, and Ni-base alloys has
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been reported to exhibit sound weld joint [17-20]. In recent years, several studies on the FSW of HNAS have been carried out [21-23]. Park et al. [21] conducted a feasibility work and demonstrated that the defects associated with melting/solidification phenomena in fusion welding process of the HNAS joint can be eliminated. Miyano et al. [22] also performed FSW of 2-mm thick Fe-23Cr-0Ni-1Mo-1N austenitic steel using different parameters. Although optimum welding parameters were determined in relation to the mechanical properties of the joint, the microstructural characteristics were not explored. D. Wang et al. [23] carried out a detailed examination of the microstructure and the mechanical properties of the friction stir welded Fe-18.4Cr-15.8Mn-2.1Mo-0.69N-0.04C austenitic stainless steel. The authors reported the FSW significantly refined the austenitic grains and increased hardness and strength in the stir zone (SZ).
ACCEPTED MANUSCRIPT In the present study, a high nitrogen nickel-free austenitic stainless steel with 0.96 wt.% nitrogen was processed with FSW. The microstructural evolution, mechanical, and corrosion properties were
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studied to reveal the microstructure and properties of friction stir welded HNAS.
2.1 Material preparation and friction stir welding process
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2. Experimental procedure
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The HNAS material used in this study was made using vacuum induction furnace and electroslag
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remelting furnace in nitrogen atmosphere [24]. The chemical composition of the steel is listed in Table 1. The experimental cold rolled plate was of dimensions 150 mm×60 mm×2.4 mm. The plates
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were solution treated at 1100 °C for 90 min, followed by water quenching and then ground to
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remove the oxide and contamination of both the top surface and joint surface. In order to assist tool
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penetration as well as reduce wear of the tool, a pilot hole whose size was similar to the tool probe was drilled at the start position of the plate.
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The welding process was carried out using an FSW machine (FSW-LM-AL25) developed by CFSWC at a traveling speed of 100mm/min and a rotational speed of 400 rpm. The tungsten-rhenium (W-Re) tool used during FSW process had a convex shoulder of 15 mm in diameter and a tapered probe. The pin was tapered from 8 mm at the shoulder to 4.5 mm with a length of 2.3 mm. During the welding process, the tilt angle was 0° and downward force was controlled at 20KN. The backing plate material was tungsten alloy, which has high thermal conductivity. An argon shielding gas at the flow rate of 21L/min was employed to avoid surface oxidation. 2.2 Macro- and microstructural evolution Following FSW, the joint was first evaluated on the basis of visual and dye penetration inspection to reveal the weld defects. Then, specimens were extracted from various locations of the weld, as
ACCEPTED MANUSCRIPT shown in Fig. 1 and from the base metal plate. The nitrogen content of the SZ and the BM was detected by Nitrogen/oxygen analyzer (LECO TC 500) to determine whether nitrogen desorption
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occurred in the weld. The metallographic specimens were prepared perpendicular to the weld direction and
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mechanically ground with SiC paper to 2000 grade and polished with 2.5 µm diamond. The samples
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were etched using 10 wt. % oxalic acid and then examined by optical microscope (OM, Olympus
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DSX 500) and scanning electron microscope (SEM, Carl Zeiss Ultra Plus). The grain size was measured with the mean liner intercept (MLI) technique. The specimens from the SZ for electron
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back-scattered diffraction (EBSD) in SEM were prepared by electrolytic polishing in a solution of
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10 vol. % perchloric acid in ethanol. Thin disks of diameter of 3 mm were cut from the SZ parallel
at 200 kV.
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2.3 Mechanical testing
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to the weld, and jet polished for transmission electron microscopy (TEM, TECNCI G2 20) studies
The mechanical properties were studied using combination of Vickers hardness (FUTURE-TECH, FM-700) and transverse tensile test (SANS-CMT5105). The hardness profile was measured at 1mm deep from the surface with regular intervals of 0.2 mm using a load of 0.5 N for 10 s across the cross-section of the weld. The configuration of a representative tensile specimen is shown in Fig. 2. Tensile specimens were ground to remove the surface ripples and flash to avoid notch effect. Three tests were conducted at a constant cross head speed of 3.0 mm/s at room temperature. The fracture surface specimen was observed using SEM. 2.4 Electrochemical measurements For corrosion tests, specimens of 8.1 mm×8.1 mm were machined from the weld joint and the base material. The passivation treatment of the specimens was performed in 20-30 wt.% nitric acid
ACCEPTED MANUSCRIPT solution at 50℃ for 1 h to prevent the occurrence of crevice corrosion. The electrodes were mounted in epoxy resin with an exposed area of 0.64 cm2 and polished with SiC paper from 200 to
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2000 grits. All the electrochemical measurements were carried out using a PARSTAT 2273 that comprised of three electrodes. The BM and weld specimens were the working electrodes, and a
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platinum foil (Pt) and a saturated calomel electrode (SCE) were used as the counter and reference
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electrodes, respectively. The potentiodynamic polarization curves were obtained in 3.5 wt% NaCl
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solution at 25 °C ±0.5 °C from -0.2 V below the OCP at a scan rate of 20 mV min-1 to the anodic direction.
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Meanwhile, intergranular corrosion (IGC) resistance was examined through the double loop
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electrochemical potentiokinetic reactivation (DL-EPR). The testing was carried out at a scanning
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rate of 1.6667 V/s at 25℃±0.5℃ in 2 M H2SO4+1 M NaCl+0.01 M KSCN solution. The solution was deaerated by purging with N2 for 30 min before the test, and a nitrogen atmosphere was
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maintained during the test. The IGC susceptibility of the steel was evaluated by comparing the ratio of the peak reactivation current density Ir and the peak activation current density Ia. All the measurements were repeated at least three times. 3. Results and discussion
3.1 Joint appearance and inspection result Fig. 3 shows a typical joint appearance processed at welding speed of 100 mm/min and tool speed rotating of 400 rpm from the top surface. No groove-like defects [25] were observed along the weld line, indicating that the heat input was sufficient to ensure the material flow during FSW process [26,27]. Dye penetrant inspection was conducted to reveal surface cracks and the inspection results confirmed the absence of surface cracks in the weld. It should be pointed out that the dyed parts correspond to fin and the weld center line exceeded the welding range. Considering that only
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The nitrogen content in SZ detected by Nitrogen/oxygen analyzer was 0.94 wt.%, which is similar to that of BM (0.96 wt.%), this implied that the SZ of the FSW processed HNAS did not
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exhibit nitrogen loss. Several authors have reported similar results [21,23]. Compared with other
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welding methods, FSW is an effective joining process to prevent nitrogen desorption.
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3.2 Microstructural evolution during FSW
The cross-sectional macrograph of the joint was consistent with the size of pin as shown in Fig.
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in other FSW studies [21,28] and fusion welds [8,9]. The microstructure was characterized by the
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stir zone (SZ), heat-affected zone (HAZ), thermo-mechanically affected zone (TMAZ) and base metal (BM). The TMAZ can be classified into TMAZ-AS and TMAZ-RS according to the
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advancing side and retreating side of the rotating tool, which are related to the left and right sides of the weld center.
Fig. 4(b) and Fig. 4(c) are the optical micrographs of BM and HAZ, respectively, which consist of austenitic and twins inside austenitic grains. The average grain size of BM was ~24.7 µm. HAZ had a microstructure similar to that of BM and the parent grains did not experience significant growth. This can be explained by the lower peak welding temperature and high cooling rate [22,29]. The cooling speed was very high, not only because the plate was thin, but also because of the fast-flow of argon shielding gas, which took away heat. On the other hand, the SZ consisted of very fine equiaxed grains, as shown in Fig. 4(d), and the average grain size measured at the middle of the SZ was ~6.3 µm. The fine equiaxed grains were produced by dynamic recrystallization induced by frictional heat and plastic deformation [21,30]. The average grain size of SZ was significantly
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Based on earlier studies [25,32], the TMAZ/SZ boundary on the advancing side of the tool was clearly sharper than that on the retreating side. In the present study there was no apparent difference
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between the two boundaries in the macrograph (Fig. 4(a)). However, as regards the microstructure
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of TMAZ-RS and TMAZ-AS (Fig. 4(e) and (f)), the asymmetry of FSW was clearly observed. The
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grains in the TMAZ-AS were severely-elongated and re-oriented perpendicular (arrowed) to the taper direction of the tool. The boundary between TMAZ-RS/SZ was not distinct because of weaker
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mechanical deformation compared to the TMAZ-AS [33].
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Fig. 5 is the SEM micrograph near the top surface and bottom of the SZ. It can be clearly seen
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that there was decrease in grain size from top to the bottom direction. This was attributed to the inhomogeneous distribution of the peak temperature and different cooling speed in the SZ, as
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previously reported [34]. The higher temperature and therefore the higher cooling speed contributed to remarkable grain growth at the top of the SZ [26,35]. TEM studies were carried out to study the microstructure in detail and are presented in Fig. 6. The equiaxed grains in the SZ were relatively small, and there was obvious difference in the dislocation density. Fine equiaxed grains with very low dislocation density were observed as shown in Fig. 6(a). This indicated that dynamic recrystallization occurred in the SZ, which was induced by the intense plastic deformation and frictional heating during the FSW progress. The positive effect of Mo and lower stacking fault energy due to nitrogen addition can promote dynamic recrystallization [36,37]. Moreover, the SZ had a grain structure with a relatively high density of sub-boundaries and dislocation pile up and tangles (Fig. 6(b), Fig. 6(c)), which indicated incomplete dynamic recrystallization in the SZ. The results confirmed that the density of dislocation of grains
ACCEPTED MANUSCRIPT in the SZ was inhomogeneous, and in general, the dislocation density was higher. The results in the present study were consistent with the typical microstructure in the FSW of other steels [23,31].
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Additionally, the precipitation of Cr23C6 of size of ~1 µm was occasionally observed in the SZ, as shown in Fig. 6(d). To clearly understand the variation in the distribution of phases in equilibrium
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with temperature and phase transformation, the phase diagram of the steel was calculated with
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Thermo-calc software. The vertical line in Fig. 7 denotes the chemical composition of the studied
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steel. It should be noted that the precipitation of intermetallic phases, such as Cr2N and σ phase were not observed in the SZ. This may be attributed to the short thermal cycle and different crystal
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structure of precipitations. During the FSW process, argon gas shielding device as oxidation
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protection and cooling method moved with FSW tool. The cooling intensity was higher at the high
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temperature range (800 ℃ ~900 ℃ ), which shortened the residence time, thus prevented precipitation of the Cr2N and σ phase. In the lower temperature range below 800 ℃, the cooling
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speed was relatively lower and the residence time was prolonged due to the decrement of argon cooling intensity. The Cr2N,σ and Cr23C6 phases at the temperature range below 800℃ could be prone to nucleate and preferentially precipitate than Cr2N,σ phase at temperature range from 800℃ to 900 ℃. In addition, the crystal structure of Cr23C6 phase is fcc, and Cr2N and σ phases are hexagonal and tetragonal structure, respectively [38]. Due to the same crystal structure with austenitic matrix, the Cr23C6 phase more preferentially precipitated. Although the introduction of high strain and dynamic recrystallization can accelerate mutual diffusion of the alloying elements [39], the short thermal cycle during FSW process allowed insufficient time for precipitation to nucleate and grow, so only less amount and small size Cr23C6 was found. Fig. 8(a) and Fig. 8(b) are the EBSD maps of BM and SZ, respectively. The white spots distributed in the colored matrix (austenitic grains) represent ferrite. The ferrite phase can be clearly
ACCEPTED MANUSCRIPT observed before and after FSW process, and furthermore increases after FSW process. The remaining ferrite in the BM was mainly attributed to element segregation such as N, Cr and Mo [40].
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Some δ-ferrite was locally formed in HNAS with high nitrogen content of 0.96 wt.%, when produced by electroslag remelting under nitrogen atmosphere. Moreover, it is noteworthy that the
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SZ contains more ferrite than the BM. This suggests that δ-ferrite in the SZ not only originated from
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the remaining ferrite in the BM, but also from the transformation of the austenitic matrix during the
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FSW process.
The local nitrogen content can be less than 0.96 wt.% as a result of element segregation, which
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resulted in a deviation from the equilibrium transformation towards lower nitrogen. At the same
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time, the temperature also should be taken into consideration during phase transformation process.
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Miyano et al. [22] measured the thermal cycle at the back of 2 mm thick Fe-23Cr-0Ni-1Mo-1N stainless steel during FSW at identical welding parameters, as used in the present study (400 rpm,
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100 mm/min). It was found that the temperature of the bottom surface in the SZ could reach ~1020 °C and the upper zone of SZ was estimated to be at slightly at higher temperature. Considering the plate used in this study is slightly thicker than pervious studies, the SZ should experience a higher temperature. Thus, considering the deviation from the equilibrium transformation, the temperature in the SZ during the FSW process presumably reached the temperature of γ- austenitic and δ-ferrite coexistence region, as shown in red shadow area in Fig. 7. The phase transformation to form new δ-phase took place during FSW process, such that the amount of δ-phase was increased. Previous studies also suggested that δ-ferrite can form in the SZ of austenitic steels [23,31]. 3.3 Mechanical properties The hardness distribution in the welded sample as a function of distance from the weld center is
ACCEPTED MANUSCRIPT presented in Fig. 9. The hardness distribution was roughly symmetrical. The hardness increased from the as-received BM (ranged from 360 to 370HV) region toward the TMAZ, reaching a
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maxima at SZ (ranged from 390 to 400HV) with a width of ~8 mm. There is not an obvious fluctuation in both the BM and the SZ, possibly due to the uniform distribution of grain sizes in the
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measured direction. The increase in the hardness was apparently caused by the refinement of grains
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during dynamic recrystallization in the SZ according to the well-known Hall-Petch relationship. In
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addition, the high density dislocations formed in the severely deformed SZ was also responsible for the increment of hardness [31]. Even though the phase change, i.e. the austenitic - ferrite
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transformation due to the element segregation took place, the hardness value may not be obviously
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affected by the formation of the discontinuous ferrite because of their extremely small amount and
of the material.
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grain size. Thus, the grain refinement and high density dislocations were the major hardening factor
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The stress-strain curves and test results obtained from the tensile test of both BM and FSW joint are plotted in Fig.10(a) and Fig.10(b). Similar to the hardness results, transverse 0.2% offset yield (770MPa) and ultimate tensile strengths (1100MPa) of the FSW joint was overmatched relative to the BM. It is confirmed that no small defects were formed in the welds due to the sufficient heat input. However, the ductility of the FSW joint was lower compared with BM. The SZ zone exhibited a smaller amount of deformation compared with the BM during test, which meant lower percentage of elongation. This can be explained by the following reasons. Firstly, it is usually right and proper that grain refinement can simultaneously improve the strength and toughness for most of metal material. But lots of literatures reported the ductility of high nitrogen steels and nitrogen containing austenitic stainless steel decreased with grain refinement [41,42], which was contrary to the traditional fine grain strengthening theory. So for HNS FSW joint, the grain refinement in SZ
ACCEPTED MANUSCRIPT may result in the decrement of toughness. High density dislocations in the SZ generated by mechanical deformation were the key factor decreasing the percentage of elongation [43]. In
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addition, the presence of less amount and small size of Cr23C6 carbides exhibits a certain extent effect on the decrement of percentage of elongation [44].
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Moreover, the tensile samples fractured in the BM near the TMAZ (inserted figure in Fig. 10(b)).
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Examination of the fracture surface BM (Fig. 10(c)) and FSW joint (Fig. 10(d)) by SEM exhibited
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dimples with a feature microvoid coalescence, which indicated characteristics of ductile fracture. But there were some differences between them, the size distribution of dimple for FSW joint was
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uneven than that of BM, and the average size of the dimple for BM was larger than that of SZ. The
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results of morphology of fracture surfaces represented that the BM exhibits higher toughness than
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FSW joint. Additionally, the unevenness of fracture near TMAZ for FSW joint indicated the microstructure near TMAZ was inhomogeneous.
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3.4 Corrosion resistance
The electrochemical polarization curves of BM and SZ in 3.5% NaCl solution are presented in Fig. 11, where the plots have similar shape. The plots had a wide range of passive region, and similar results were obtained in our previous research [45]. The corrosion potential (Ecorr) and the corrosion current densities (icorr) were similar. However, the pitting potential (Eb,10) value of SZ was marginally lower, which implied lower pitting resistance. In general, localized corrosion resistance of welds is affected by microstructural changes according to previous research. On the one hand, the precipitation of secondary phase, such as sigma, Cr2N, Cr23C6, would lead to generation of Cr-depleted zone, which is sensitive to corrosion [1,4]. Ha et al. [1] reported that the pitting corrosion resistance of HNAS was deteriorated by lamellar Cr2N precipitate formed along grain boundaries. On the other hand, severe plastic
ACCEPTED MANUSCRIPT deformation and dynamic recrystallization contributed to the formation of fine grains in SZ, which can enhance the corrosion resistance [46-48]. Sarlak et al. [46] reported that corrosion performance
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of the stir zones of the lean duplex steel were not deteriorated, and even exhibited superior corrosion resistance than the base metal.
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In the present study, although precipitation of Cr23C6 occurred, the amount of precipitation was
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less, may exhibit little influence the corrosion resistance. Thus, considerable amount of δ-ferrite
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must be taken into account for the decrement of the pitting corrosion resistance. Fig. 12 represents SEM morphologies of δ-ferrite in BM and SZ and SEM-EDS line profile through δ-ferrite phase.
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The SEM-EDS results showed that the Cr and Mo contents were higher than that of austenite, and
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Mn and N were lower. The formation of Cr-depletion zone was found at the interface between the
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austenitic and δ-ferrite as shown Fig. 12(c). Several authors have reported that the formation of a Cr-depleted zone promoted pitting corrosion in the SZ [49,50]. Thus, the decrease of Eb,10 of the SZ
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compared to the BM was attributed to the formation of a Cr-depleted zone due to δ-ferrite formation, regardless of grain refinement.
The EPR results of the BM and SZ are presented in Fig. 12. The activation peaks of both the samples were nearly same, and the reactivation peaks could hardly be observed. Since the solution treated BM appeared desensitized, the reactivation current density was much lower than that of the activation. From the local amplification of reactivation peaks in inserted graph, the reactivation peak of the SZ was slightly higher than that of BM, and Ir/Ia ratio for the BM and the SZ were 0.596 and 1.089, respectively. This indicated that the intergranular corrosion of the SZ was more sensitive than the BM. It is well-known that the sensitivity to intergranular corrosion of stainless steel is closely related to the depletion of chromium along the grain boundarieds [51,52]. In present study, the intergranular
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Fig.12(c). Thus, the formation of δ-ferrite led to a lower degree of intergranular corrosion resistance of the SZ [50,53].
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4. Conclusions
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In this study, the FSW of a high-nitrogen nickel-free austenitic stainless steel plate was performed
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at a tool rotation speed of 400rpm and welding speed of 100mm/min. Microstructural characteristics, mechanical and corrosion properties were explored. The results of the study can be summarized as
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follows:
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(1) A sound FSW joint was acquired without any defects. The SZ did not show nitrogen loss, and
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was characterized by fine-grained microstructure. (2) Based on TEM and EBSD studies, a small amount of Cr23C6 was observed and amount of
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δ-ferrite increased in SZ during FSW process. (3) The hardness of SZ was slightly higher than the BM. The tensile strength and yield strength of the joint were higher than the BM, while the elongation was slightly lower. (4) Based on measurements of pitting and intergranular corrosion resistance, the corrosion resistance of SZ was slightly decreased in comparison to the BM.
Acknowledgments The present research was financially supported by National Natural Science Foundation of China (No.51304041, 51434004 and U1435205) and Program for New Century Excellent Talents in University (No. N130502001). RDKM gratefully acknowledges support from University of Texas at El Paso, USA. References
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Figure and Table Captions
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Table 1 Chemical composition of base material (wt.%)
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Fig. 2. Configuration of transverse tensile specimen (mm)
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Fig. 1. Scheme of extraction of samples (A-EBSD specimen, B-Metallographic specimen, C-Tensile specimen, D-Nitrogen-oxygen analysis specimen, E-TEM specimen, F-Electrochemical corrosion testing specimen, G- Hardness specimen; mm)
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Fig. 3. (a) Appearance of FSW zone and (b) dye penetration inspection result
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Fig. 4. (a) Optical macrograph of transverse section of the joint and microstructure indicated in the macrograph, (b)BM, (c) HAZ, (d) SZ, (e)TMAZ-AS, (f)TMAZ-RS Fig. 5. SEM micrographs of the SZ (a) near the top surface and (b) at the bottom
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Fig. 7. The equilibrium phase diagram calculated with Thermo-calc software Fig. 8. Phase maps of the weld: (a)BM, (b)SZ Fig. 9. Microhardness profile of FSW HNAS joint Fig. 10. (a) stress-strain curves and macrograph of the fractured FSW joint (inserted figure), (b) Results of tensile test, (c) SEM morphologies of fracture surfaces (BM), (d) SEM morphologies of fracture surfaces (FSW joint) Fig. 11. Potentiodynamic polarization curves of BM and SZ in 3.5wt% NaCl solution Fig. 12. SEM morphologies of δ-ferrite in BM (a) and SZ(b), (c) SEM-EDS line profile through δ-ferrite phase Fig. 13. EPR curves of the BM and SZ in 2 M H2SO4+1 M NaCl+0.01 M KSCN
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Advancing side
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Rolling and welding direction
Tool rotation direction
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Fig. 1. Scheme of extraction of samples (A-EBSD specimen, B-Metallographic specimen, C-Tensile specimen, D-Nitrogen-oxygen analysis specimen, E-TEM specimen, F-Electrochemical corrosion testing specimen, GHardness specimen; mm)
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Fig. 2. Configuration of transverse tensile specimen (mm)
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Retreating side
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3mm (d)
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Advancing side
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Fig. 4. (a) Optical macrograph of transverse section of the joint and microstructure indicated in the macrograph, (b)BM, (c) HAZ, (d) SZ, (e)TMAZ-AS, (f)TMAZ-RS
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Fig. 5. SEM micrographs of the SZ (a) near the top surface and (b) at the bottom
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Fig. 6. TEM images of SZ: (a) Fine grains with very low dislocation density, (b) high density dislocation network structure, (c) grain structure with relatively high density of subboundaries and dislocations, (d) precipitation of Cr23C6
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Fig. 7. The equilibrium phase diagram calculated with Thermo-calc software
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Fig. 8. Phase maps of the weld: (a)BM, (b)SZ
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Fig. 9. Microhardness profile of FSW HNAS joint
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The joint
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Fig. 10. (a) stress-strain curves and macrograph of the fractured FSW joint (inserted figure), (a)Results of tensile test, (c) SEM morphologies of fracture surfaces (BM), (d) SEM morphologies of fracture surfaces (FSW joint)
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Fig. 11. Potentiodynamic polarization curves of BM and SZ in 3.5wt% NaCl solution
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Fig. 12. SEM morphologies of δ-ferrite in BM (a) and SZ(b), (c) SEM-EDS line profile through δ-ferrite phase
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Fig. 13. EPR curves of the BM and SZ in 2 M H2SO4+1 M NaCl+0.01 M KSCN
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Graphical abstract
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Si
Cr
Mn
Mo
Ni
HNAS 0.058 0.19 19.55 19.5 2.26
--
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Al
O
N
0.003 0.015 0.04 0.0048 0.96
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Steel
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Table 1 Chemical composition of base material (wt.%)
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Highlights
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1. Friction stir welding (FSW) was successfully applied to a high nitrogen nickel-free austenite steel with 0.96 wt.% nitrogen. 2. The small and little amount Cr23C6 was found. 3. The increment of the amount of δ-ferrite in the SZ was discussed. 4. The corrosion results revealed the SZ exhibited excellent corrosion resistance.