Microstructure evolution and mechanical properties of ...

Materials and Design 118 (2017) 207–217

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Microstructure evolution and mechanical properties of friction stir welding super-austenitic stainless steel S32654 Huabing Li a,⁎, Shouxing Yang a, Shucai Zhang a,⁎, Binbin Zhang a, Zhouhua Jiang a, Hao Feng a, Peide Han b, Jizhong Li c a b c

School of Metallurgy, Northeastern University, Shenyang 110819, China College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China Guangdong Welding Institute (China-Ukraine E. O. Paton Institute of Welding), Guangzhou 510650, China

H I G H L I G H T S

G R A P H I C A L

A B S T R A C T

• Friction stir welding was successfully applied to super-austenitic stainless steel S32654. • Grain structure evolution in stir zone is dominated by continuous dynamic recrystallization. • Strain rate plays a dominated effect on grain refinement. • Grain refinement, high density dislocations and substructures together improve hardness and strength, but reduce elongation. • The more suitable welding parameters are determined as 300 rpm and 100 mm/min for this steel.

a r t i c l e

i n f o

Article history: Received 6 October 2016 Received in revised form 31 December 2016 Accepted 11 January 2017 Available online 13 January 2017 Keywords: Friction stir welding Super-austenitic stainless steel S32654 Microstructure evolution Recrystallization Mechanical properties

a b s t r a c t Super-austenitic stainless steel S32654 sheets with 2.4 mm thickness were successfully welded by friction stir welding (FSW) at the rotational speeds of 300 and 400 rpm with a constant traverse speed of 100 mm/min using W-Re tool. The sound joints with almost no nitrogen loss were successfully produced. The microstructure evolution was characterized by optical digital microscope (ODM), scanning electron microscopy with energy-dispersive spectroscopy (SEM-EDS), electron backscatter diffraction (EBSD) and transmission electron spectroscopy (TEM). The results suggest that the grain structure evolution in stir zone (SZ) is dominated by continuous dynamic recrystallization (CDRX). The strain rate plays a dominated effect on obvious grain refinement. The band structures containing W and Re are generated due to the wear between tool probe and steel in SZ. Furthermore, the microhardness measurements and transverse tensile tests indicate that the grain refinement combining with high density dislocations and substructures improves the hardness and strength, but greatly reduces the plastic deformation capacity of joints. The more suitable welding parameters are determined as 300 rpm and 100 mm/min for this steel. © 2017 Elsevier Ltd. All rights reserved.

⁎ Corresponding authors. E-mail addresses: [email protected] (H. Li), [email protected] (S. Zhang).

http://dx.doi.org/10.1016/j.matdes.2017.01.034 0264-1275/© 2017 Elsevier Ltd. All rights reserved.

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H. Li et al. / Materials and Design 118 (2017) 207–217

1. Introduction As one of the highest alloyed austenitic stainless steel grades, superaustenitic stainless steel (SASS) S32654 with pitting resistance equivalent number (PREN) (PREN = wt.% Cr + 3.3 wt.% Mo + 16 wt.% N) value higher than 56 exhibits outstanding corrosion resistance and mechanical properties [1–3]. This steel is believed to have a great potential for application in extremely harsh service environments, such as seawater handling systems, chemical processing equipment, waste incineration systems and plate heat exchangers [4,5]. In general, the conventional fusion welding methods, such as GTAW, TIG and laser welding, of high alloyed stainless steels have some drawbacks like grain growth and welding defects (cracks, holes or undercuts) [6,7]. For S32654, it is more sensitive to precipitation of secondary phases (σ, χ and Laves phase) due to its higher alloy contents [8–10]. Taking into account the higher welding heat input in fusion welding, it is more likely to induce segregation and precipitation for this steel. Therefore, it is a great technical challenge to choose an appropriate welding method for S32654. To avoid the problems associated with traditional fusion welding process, friction stir welding (FSW) invited as an innovative welding technique by TWI in 1991 is considered appropriate [11]. As a solidstate welding technology, FSW exhibits lower peak temperature, which leads to short thermal cycle and excellent weldability. In addition, the lower heat input for FSW process is benefit to suppress segregation, precipitation and the formation of defects [12,13]. Recently, FSW has been widely applied to the materials with moderate melting point like Al, Mg and Cu alloys [14–16]. With the development of specialized tool material, this technology has been promoted for applications of higher melting point alloys like carbon steels [17,18], stainless steels [19–24] and Ni-base alloys [25,26]. So far, the exploration of FSW SASS is only confined to the researches on the feasibility, microstructure and mechanical properties of lower alloyed SASS [27–30]. Mehranfar et al. [27] suggested that a nanostructured layer of about 91 μm thick could be produced on the specimen surface of super-austenitic steel using FSW technology. Mironov et al. [28] found that S31254 was successfully welded by FSW method and the microstructure of the stir zone (SZ) was primarily governed by discontinuous recrystallization. It also can be confirmed by Sato et al. [29] that the low heat input FSW was an effective method to produce a joint with relatively mechanical properties for NSSC 270. As can be revealed by the research of Klingensmith et al. [30], the AL-6XN plates have been successfully joined via a double-sided FSW process and SZ had a fine equiaxed grain structure and the higher hardness. However, there is very little information reported on microstructure evolution and mechanical properties of FSW SASS with higher alloy content (such as S32654). The primary purpose of this paper is an attempt to evaluate the feasibility of FSW SASS S32654 with high quality joints, and to explore a suitable FSW process. At the same time, the microstructure evaluation and mechanical properties of joints were investigated. The mechanisms of dynamic recrystallization (DRX) and grain refinement were discussed in detail. It is expected that the present research would provide a new method for welding higher alloy content of stainless steels and alloys. 2. Experimental procedure The SASS S32654 used in the present study was smelt using an induction furnace under nitrogen atmosphere. The chemical composition of the ingot is shown in Table 1. The ingot was hot-forged and hot-rolled within the temperature range from 1100 to 1250 °C into a 3 mm thick plate. The hot-rolled plate was solution treated at 1200 °C for 1 h, followed by water quenching to ensure all secondary phases dissolving back into the matrix. Then the plate was ground to remove the oxide and cut into the dimensions of 200 mm × 75 mm × 2.4 mm. A pilot

Table 1 Chemical composition of the S32654 used in this study (wt.%). C

Cr

Ni

Mo

N

Mn

Si

Cu

P

S

Fe

0.012

24.45

22.58

7.38

0.54

2.92

0.39

0.46

0.005

0.002

Bal.

hole whose size was similar to that of the tool probe was drilled at the start position of the plate to assist tool penetration. The FSW was performed using a gantry welding machine named FSW-LM-AL25. The process ran along the centerline of plates aligned in the rolling direction at a constant welding speed of 100 mm/min (ν) and different rotational speeds (ω) of 300 and 400 rpm. The welding tool was made of tungsten-rhenium (W-Re) with a convex shoulder diameter of 20 mm and a tapered probe. The probe was tapered from 8 mm at the upper diameter to 4.5 mm at the lower diameter with a length of 2.3 mm. The plates were fixed onto a tungsten alloy backing plate to prevent any displacement. It is well accepted that both increasing rotational speed and decreasing welding speed can reduce the download force. Combined with the values reported in literatures [25, 31] and our previous research [7], the download force was controlled at 20 kN. The tilt angle was 0°. In order to minimize surface oxidation of joints and tool, argon gas shield was employed at the volume flow rate of 21 L/min during FSW process. After FSW, the joints were observed according to the visual inspection to evaluate the surface quality. Then the nitrogen content analysis, microstructure observation, Vickers microhardness and tensile tests of joints were performed, and the schematic illustration of welding and sampling is shown in Fig. 1. The principal coordinate of FSW was designed as transverse direction (TD), welding direction (WD) and normal direction (ND), respectively. The nitrogen content of the base metal (BM) and the stir zone (SZ) was detected by a LECO TC-500 nitrogen/oxygen analyzer to evaluate whether the nitrogen desorption occurred during FSW process. The metallographic samples were cut perpendicular to WD and ground with SiC embedded papers of up to 2000 grit and polished with 2.5 μm diamond polishing paste. In order to display the transverse cross section microstructure, the samples were electrolytically etched using 5 g oxalic acid + 100 mL concentrated hydrochloric acid at 5 V for 3– 5 s, and then observed by Olympus DSX 510 optical digital microscope (ODM) and Care Zeiss Ultra Plus scanning electron microscope (SEM). The samples for electron backscatter diffraction (EBSD) analysis were electrolytically polished at 20 V for 20–25 s in the solution of 12.5 vol.% HClO4 + 87.5 vol.% CH2H5OH. The EBSD regions were schematically shown as red-colored rectangles in Fig. 1 and the EBSD datas were acquired using the indexed patterns of Kikuchi bands and analyzed by the HKL Channel 5 software. The Tecnai G 2 20 transmission electron microscopy (TEM) operated at 200 kV was used to observe dislocations, substructures and precipitates in SZ. The standard disks of TEM samples with the diameter of 3 mm were ground to almost 50 μm thickness, and then thinned to electron transparency in the solution of 8 vol.% HClO4 + 92 vol.% CH2H5OH at 40 V and − 20 °C for 30 s using a Struers TenuPol-5 twin-jet electropolisher. To evaluate the mechanical properties of FSW joints at different welding parameters, the Vickers microhardness on the transverse cross section was performed by Future-Tech FM-700 microhardness tester. The hardness profiles were measured at 1 mm deep from the joints surface with regular intervals of 0.5 mm using a load of 2 N for 15 s. In addition, the transverse tensile tests were carried out by SANS-CMT 5105 computer control electronic universal tensile testing machine with a constant cross head speed of 1 mm/min at room temperature. A configuration of the representative tensile samples was shown in Fig. 1. Before each test, the tensile sample was ground to remove the surface ripples and avoid notch effect.


209

Fig. 1. Schematic illustration of friction stir welding and transverse tensile test specimen (A—tensile sample, B—hardness sample, C—EBSD sample, D—metallographic sample, E—nitrogen analysis sample). (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)

3. Results and discussion

3.2. Joint cross-section appearance

3.1. Joint surface appearance and nitrogen content

Fig. 4 shows the ODM overviews of the transverse cross sections of FSW samples at 300 and 400 rpm. Obviously, the sound joints were successfully produced at both welding parameters. The weldment profiles appear “basin-shape” which is consistent with the size of the tool probe. The left- and right-hand sides of the weldment center correspond to the advancing side (AS) and retreating side (RS) of rotation tool, respectively (Fig. 4). In both cases, the joints can be split into several distinct regions: base metal (BM), stir zone (SZ) and the narrow transition region-commonly called thermo-mechanically affected zone (TMAZ). According to AS and RS of rotation probe, the TMAZ can be divided into TMAZ-AS and TMAZ-RS. No obvious heat affected zone (HAZ) was found in either sample, indicating that the growth of parent grains could be effectively inhibited during FSW process. This can be explained by the lower peak temperature and higher cooling rate compared with traditional welding methods [7,34]. The sheets were fully penetration welded without internal welding defects like cracks, pores or cavities. This is mainly because the sufficient heat input elevated the

Fig. 2 exhibits the typical joint surface appearances at a constant welding speed of 100 mm/min and the rotational speeds of 300 and 400 rpm. Under the above two welding parameters, the samples were perfectly welded without any surface-breaking defects, such as pores, groove-like defects and surface galling. The appearances of weldments are flat and glossy, indicating that the heat input was sufficient to ensure good material fluidity during FSW process at both rotational speeds [32,33]. The nitrogen contents of SZ at 300 and 400 rpm detected by nitrogen/oxygen tester analyzer were 0.5387 and 0.5401 wt.% respectively, which are similar to that of the BM (0.5409 wt.%) as shown in Fig. 3. Similar results about FSW high nitrogen austenitic stainless steels have also been reported [7,34]. Therefore, it can be concluded that the FSW is an effective welding method for S32654 containing high nitrogen content to prevent nitrogen desorption.

Fig. 2. Appearance of FSW zone at different rotational speeds: (a) front side of 300 rpm, (b) reverse side of 300 rpm, (c) front side of 400 rpm, (d) reverse side of 400 rpm.

210


Fig. 3. Nitrogen contents in BM and SZ at different rotational speeds.

temperature and effectively promoted the plastic flow of materials around the tool probe [6]. In addition, as a solid-state welding technology without going through the process of melting and solidification, FSW could effectively prevent from nitrogen emission and the formation of nitrogen pores for this steel. To observe the microstructure of each region more clearly, high magnification optical images were obtained as shown in Fig. 5. The corresponding microstructure characteristics of BM, HAZ and SZ at 300 and 400 rpm are very similar. The BM exhibits a typical austenitic microstructure consisting of coarse austenite grains and a large number of annealing twins (Fig. 5a and f). The HAZ (Fig. 5b and g) has a similar microstructure to BM. The SZ contains extremely fine equiaxed grains due to DRX induced by frictional heat generated and severe plastic deformation [21,35], as shown in Fig. 5c and h. The morphologies of TMAZ-AS and TMAZ-RS at 300 rpm are relatively symmetrical, and there is no obviously sharp boundary or band structures (BSs) forming in this joint. However, compared with the microstructure of TMAZ at 300 rpm (Fig. 5d and e), the shape and width of TMAZ at 400 rpm (Fig. 5i and j) change greatly. For example, it is clearly observed that the microstructure of TMAZ-AS and TMAZ-RS is asymmetric. The grains in the TMAZ-AS are significantly elongated (indicated by the arrow in Fig. 5i), and their distinctive feature is a streamline distribution relative to the SZ borderline. A sharp boundary can be observed between SZ and TMAZ-AS (Fig. 5i). In addition, it is noteworthy that a large number of BSs appear in the AS of SZ at 400 rpm (Fig. 5i), while this phenomenon is not obvious at 300 rpm. These BSs with the feature of streamline or lamellar structure are parallel to the SZ borderline. This distribution of BSs is greatly different from many other researches, and their BSs distributed in the center or bottom of SZ were mainly attributed to the intermetallic particles precipitation or foreign inclusion formation during FSW process [21,28].

To further explore the reasons for the formation and distribution of BSs at 400 rpm, SEM observations and EDS analyses were performed (marked a, b, c and d included in box), as shown in Fig. 6. The BSs formed not only in the AS of SZ (Fig. 6a, c and d) but also at the upper surface of SZ (Fig. 6b). Moreover, the mass fraction of white granular substances were found inside the BSs. EDS results show that the granular substances contain a certain amount of tungsten (W) and rhenium (Re) elements (marked A and C in Fig. 6d and e). The contents of W and Re in position A are 14.77 wt.% and 5.51 wt.%, and that in position C are 29.38 wt.% and 9.22 wt.% (Table 2). The dense regions in the vicinity of granular substances (marked B in Fig. 6c) also include 5.45 wt.% W and 1.62 wt.% Re, respectively (Table 2). However, both W and Re are not detected in the matrix between these BSs (marked D in Fig. 6d and Table 2). It can be considered that the BSs formed in SZ at 400 rpm are owing to higher friction load between weldments and tool probe. At 400 rpm, the strain and the frictional heat are much higher; which results in the wear of the tool. The worn tool particles (containing W and Re) were entrained into SZ by the flow materials during FSW process. The alternating distribution characteristic of BSs in SZ may be attributed to the special material transport phenomenon and stress state described as follows [36,37]. In the stirring process, a very thin layer of material surrounding the tool probe suffered from rotational extrusion. This led to the formation of large strain and strain rate gradients around the tool probe. When the relative distance between worn particles and tool probe appreciably exceeded the strain rate range, the particles were swept away and tended to move to the low strain rate region, causing the formation of clear alternate bands. In addition, the materials closing to the tool probe in AS were exposed to a higher shear stress than that in RS. It was easier to understand why the BSs formed in the AS of SZ. Besides, the BSs formed at the upper surface of SZ can be as a result of the wear between shoulder and samples. 3.3. Microstructure evolution 3.3.1. Dynamic recrystallization mechanism During the FSW process, the grains of SZ experienced intense plastic deformation and frictional heating, which led to a high degree of grain refinement because of DRX. EBSD analyses were carried out on BM, SZ and TMAZ-AS to clarify the DRX mechanism, as shown in Figs. 7–9. Image quality maps, grain boundary maps and misorientation angle distributions for BM and SZ at 300 and 400 rpm are illustrated in Fig. 7, where low-angle boundaries (LABs) having misorientation angles of 2–6°, 6–15° are depicted as pink and blue lines, and high-angle boundaries (HABs) with misorientation angles higher than 15° are marked by black lines. The corresponding proportions of LABs and HABs are summarized in Fig. 8. The BM consists of 97.3% HABs and only 2.7% LABs (Fig. 8), indicating that the microstructure is almost full of recrystallized structure [25]. After FSW at 300 and 400 rpm, the misorientation angles distribution of SZ changes greatly. The proportions of HABs are

Fig. 4. Optical low magnification overviews of the transverse cross sections of FSW specimens according to rotational speeds: (a) 300 rpm and (b) 400 rpm.


211

Fig. 5. Optical high magnification images of various microstructural zones indicated in Fig. 4: the BM (a) (f), HAZ (b) (g), SZ (c) (h), TMAZ-AS (d) (i) and TMAZ-RS (e) (j) of 300 and 400 rpm, respectively.

decreased to 88.2% and 85.2%, and the LABs proportions are increased to 11.8% and 14.8%, respectively. These changes demonstrate that the DRX of SZ at both samples is insufficient [25]. From the grain boundary maps of the SZ at 300 and 400 rpm (Fig. 7b), it can be seen that three types of grain boundaries including LABs with misorientation angles of 2°6°, LABs with misorientation angles of 6–15° and HABs were connected in turn (as indicated by red circle in Fig. 7b). A few HABs were not arranged as a clear grain structure, on the contrary, they were described by isolated HABs segments distributing in LABs with misorientation angles of 2–6° (red arrowed in Fig. 7b). It seems that the misorientation angles of LABs are progressively increased and eventually transformed into HABs during grain evolution process [38–40]. The changes of microstructure appear to fit the

definition of continuous dynamic recrystallization (CDRX) [41,42]. Moreover, according to the image quality maps (Fig. 9a) and grain boundary maps (Fig. 9b) of TMAZ-AS at 400 rpm, the grain size and grain boundary distribution did not appear obvious gradient transition characteristic. But a sharp borderline between BM and SZ could be observed clearly. Meanwhile, a series of deformation bands formed in the coarse grain approaching SZ (Fig. 9a), and these deformation bands could be considered to be subgrain boundaries (misorientation angle b 2o) [39]. It is interesting to note that some LAB segments with misorientation angles of 2–6° were preferentially generated from these deformation bands (as indicated by blue arrows in Fig. 9b). This reveals that subgrain boundaries continuously accumulated and gradually transformed into LABs, which is a typical characteristic of CDRX [43,

Fig. 6. Optical macrostructure of transverse section of the joint at rotational speed of 400 rpm (a) and SEM microstructure corresponding to different regions of the macrostructure marked b, c, d and e included in Fig. 6(a).

212


Table 2 The results of the EDS analysis of BS and matrix at different regions. EDS points

A B C D

Chemical composition (wt.%) Cr

Ni

Mo

Mn

Fe

W

Re

23.69 23.99 17.11 25.55

10.80 21.25 9.78 22.37

13.33 6.07 7.40 6.85

1.98 2.83 1.50 3.39

29.91 38.79 24.61 40.99

14.77 5.45 29.38 0

5.51 1.62 10.22 0

44]. The TEM analyses of SZ after FSW at 300 and 400 rpm (Fig. 10) provided a fundamental insight into the DRX process. Some apparent dislocation walls or subgrains were found in both the FSW samples (marked as white arrows in Fig. 10a and b), and their numbers increase with the rotational speed, which could be another evidence for the occurrence of CDRX. Based on the above EBSD and TEM analyses (Figs. 7–10), it is concluded that the DRX mechanism of FSW S32654 is dominated by CDRX. However, it is noting that the DRX mechanism of FSW S32654 is inconsistent with the that of FSW conventional austenitic stainless steels like AISI 316L or super austenitic stainless steels with relatively lower alloy content like S31254, whose DRX mechanisms are primarily dominated by DDRX [19,28]. The difference in DRX mechanisms could be attributed to the distinct metallurgical characteristics of the materials, especially stacking fault energy (SFE) [45]. According to these literatures, both AISI 316L and S31254 are usually classified as low SFE materials. In such materials, partial dislocations are too far separated to easily recombine and climb, so cross slip and dynamic recovery occurs slowly under hot deformation process. That is, dislocations can hardly rearrange themselves into subboundaries and thus pile up inside the grains. The accumulated dislocations could gradually transform into the crystal nuclei, subsequently the crystal nuclei grow into new grains via a process of DDRX. Up to now, there is little information available about the SFE of S32654. According to the two appropriate empirical equations

Fig. 8. The summary of the grain size and the different types of grain boundary proportion during FSW of S32654.

proposed by Dai et al. [46] and Scharam and Reed [47], the SFE value of S32654 could be evaluated respectively. 0:5 SFE ¼ γ0 þ 1:59Ni‐1:34Mn þ 0:06Mn2 ‐1:75Cr þ 0:01Cr2 þ 15:21Mo‐5:59Si‐60:69 ðC þ 1:2NÞ þ26:27ðC þ 1:2NÞ ðCr þ Mn þ MoÞ0:5 þ 0:61½Ni ðCr þ MnÞ0:5 mJ=m2

ð1Þ where γ0 is the SFE value of pure austenitic iron at room temperature, and its value is about 36–42 mJ/m2 [46]. SFE ¼ −53 þ 6:2Ni þ 0:7Cr þ 3:2Mn þ 9:3Mo mJ=m2

ð2Þ

in the Eqs. (1) and (2), the symbol for the alloying element represents its weight percentage.

Fig. 7. (a) EBSD image quality maps, (b) grain boundary maps and (c) misorientation angle distributions for the BM (top) and the SZ of FSW at rotational speeds of 300 rpm (middle) and 400 rpm (bottom). LABs are depicted as pink and blue lines, HABs are marked by black lines on the grain boundary maps, respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)


213

Fig. 9. (a) EBSD image quality maps and (b) grain boundary maps for the TMAZ-AS of FSW at rotational speeds of 400 rpm. (For interpretation of the references to color in this figure, the reader is referred to the web version of this article.)

The SFE value of S32654 was calculated to be about 208.5 and 182.1 mJ/m2, which is similar to that of the Al (about 200 mJ/m2) [48], ferritic stainless steel (about 200 mJ/m2) [49] and IF steels (about 200 mJ/m2) [50] wildly considered as high SFE materials. So S32654 could be inferred as a high SFE metal. During FSW process, the cross slip and climb of dislocations occurred rapidly and dislocations rearranged themselves into subgrain boundaries easily. Then the misorientation angles of the subgrain boundaries progressively increased through continuously absorbing dislocations, leading to subgrains rotation. Eventually, the subgrains transformed into a real new DRX grain. It is theoretically proved that the CDRX was the operative DRX mechanism for S32654 during FSW process. 3.3.2. Grain refinement mechanism It is commonly accepted that heat input is the primary factor affecting grain refinement during FSW process. The pseudo heat index (ω2/ν) was defined to describe the dependence of the heat input to the rotational speed (ω) and welding speed (ν). The power 2 for ω indicates that the contribution of rotational speed to heat input is far greater than that of welding speed [19]. Accordingly, at a constant welding speed in the FSW, the higher the rotational speeds, the more the heat input and the larger the grain size. In other words, the grain size in SZ is directly proportional to the rotational speed (heat input). This view has been confirmed by several reports about FSW super or ordinary austenitic stainless steels. For example, Sato et al. [29] found that the average grain size in SZ of SASS NSSC 270 increased from 2 μm to 10 μm when raising the rotational speed from 400 rpm to 800 rpm. Hajian et al. [19] showed that increasing rotational speed from 200 rpm to 315 rpm resulted in greater grain size during FSW 316L stainless steel. Similarly, Reynolds et al. [51] also observed the same law between average grain size and rotational speed during the process of FSW 304L stainless steel. From the EBSD maps (Figs. 7 and 8), it can be more clearly seen that the grain sizes of SZ at 300 and 400 rpm are 1.4 and 1.3 μm, respectively, which is much refiner than that (62.7 μm) of BM. However, the mean grain size does not increase but slightly reduce with increasing the rotation speed from 300 rpm to 400 rpm. This phenomenon is obviously different from other researches [19,29,51]. In fact, the grain size in SZ after FSW is not only decided by heat input (temperature), but also by strain

rate [52]. The higher strain rate is conducive to reducing grain size, which is mainly revealed in the following two aspects. On the one hand, increasing the rotational speed i.e. strain rate results in the dramatic enhancement of dislocation density (Fig. 10) and substructure (Figs. 7 and 10) in SZ, so the critical radius of subgrain would be smaller (nucleation would be easier), and the nucleation rate intensely increases [53]. On the other hand, under higher strain rate, there is not enough time for dislocations movement and grain boundaries migration, resulting that the DRX grains growth is restrained [54,55]. As a result, the higher nucleation and lower growth under high strain rate condition would promote the formation of much finer DRX grains. In conclusion, both the heat input (temperature) and strain rate will increase with the rotational speed, and have an opposite effect on the grain size. In the present research, the grain size at 400 rpm is a bit less than that at 300 rpm, hence it can be concluded that the strain rate is dominated for grain refinement of FSW S32654. Lots of researches [56,57] suggested the grain refinements were also associated with the Zener-Hollomon (Z) parameter, which reflected the _ as follow: peak temperature (T) compensated strain rate (ε) QA Z ¼ ε_ exp RT

ð3Þ

where QA is the hot deformation activation energy and R is gas constant. _ in FSW by the Eq. The earlier studies for estimating the strain rate (ε) (4): ε_ ¼ 2π

de Rm Le

ð4Þ

where de and Le represent the effective diameter and depth of the DRX zone. In these studies, the effective radius of DRX zone is about π/8 of shoulder diameter, the depth of DRX zone is about π/4 of the plunge depth. Rm is the average flow rate of material which is assumed to be about half of rotational speed. The peak temperature (T) is often evaluated by the following equation [12,58]: T¼K

α

ω2 4

ν 10

Tm

ð5Þ

where ω and ν are the tool rotational speed and traverse speed. Tm is the melting point of material and K and α are contents. According to the Eqs. (3), (4) and (5), the relationship between Z values and rotational speed (ω) is derived as follow: M Z ¼ 13:65 ω exp 2α ω

Fig. 10. TEM micrographs of FSW samples at rotational speeds of 300 (a) and 400 rpm (b).

M¼

104α Q A ν α RK α T m

ð6Þ ð7Þ

214


Based on the above equations, α, QA, R, K and Tm are constants during FSW S32654. Under a constant traverse welding speed (v = 100 mm/min), M is also considered as a constant. In our research, it is very likely that the Z value slightly increases when the rotational speed increases. It is expected that the higher Z values in FSW would lead to finer grains and vice versa [56,57]. Thus, this could further provide an evidence for the slight decrement of grain size with the rotational speed increasing from 300 to 400 rpm. Cui et al. [6] and Su et al. [59] also found that the grain in SZ presents a slight refinement with increasing the rotation speed during FSW process. In addition, Abedi et al. [60] also found that increasing strain rate promoted grain refinement during the high temperature deformation process of low density steel. Simultaneously, raising true strain significantly refined the grain size [61]. 3.4. Mechanical properties 3.4.1. Microhardness tests Fig. 11 shows the microhardness profiles for the two welded samples at 300 and 400 rpm along transverse direction. Both the hardness distributions are roughly symmetrical. In the narrow TMAZ, the changes of hardness values depend on the transient features of microstructure. Based on the hardness distribution, it is difficult to distinguish the existence of softened HAZ in both as-welded joints. The hardness values of the SZ at 300 rpm (363.3 Hv) and 400 rpm (378.0 Hv) are notably higher than that of BM (290.7 Hv), and the value for 400 rpm is slightly higher. According to the famous Hall-Petch formula [25,58], the obvious increment of hardness values in the SZ at 300 and 400 rpm can be attributed to high degree of grain refinement. The hardness changes versus rotational speed could be due to the following two aspects. Firstly, the grain size of SZ at 400 rpm is a little smaller than that at 300 rpm, which results in a slight increase in hardness. Secondly, the dislocation density (Fig. 10) as well as substructure (Figs. 7 and 10) increases with rotational speed, which can also improve the hardness of the weldments [62]. Accordingly, the grain refinement, high density dislocations and substructures are the main factors controlling the SZ of FSW S32654 hardening. 3.4.2. Transverse tensile tests The transverse tensile tests of BM and both joints were carried out and the corresponding results are shown in Figs. 12 and 13. The macrographs of tensile samples (Fig. 12) reveal that all the samples are elongated significantly after tensile tests. It is observed that the BM sample was elongated uniformly and broke at its center (Fig. 12a). However, both the two FSW weldments were stretched preferentially at BM and finally failed in BM. It is noting that only a low degree of deformation occurred in SZ, indicating that the SZ exhibits higher strength and lower elongation than BM. The stress-strain curves, tensile strength and elongation of all the tensile samples are plotted in Fig. 13. During tensile tests, all the tensile samples underwent four typical stages including elastic deformation,

plastic deformation, work hardening and fracture (Fig. 13a). The stages of plastic deformation and work hardening are more pronounced. Compared with the BM sample, the work hardening phenomena of two joints are more obvious in the plastic deformation process. In addition, the higher the rotational speed, the more significant the work hardening phenomenon is. The BM sample has an ultimate tensile strength (UTS) and yield strength (YS) of 875 and 445 MPa with an elongation of 65.2%, indicating that S32654 exhibits superior strength and excellent ductility. After FSW at 300 rpm, the UTS and YS values are 890 and 505 MPa, respectively. For the sample at 400 rpm, the two values are 895 and 510 MPa. This reveals that raising rotational speed leads to a slight increase in YS and UTS. However, the elongations of tensile samples at 300 and 400 rpm are reduced from 65.2% (BM sample) to 41.7 and 37.2%, respectively, which suggest that the plastic deformation capacities of joints decrease. Moreover, the SEM fractography exhibits apparent necking (Fig. 13c) and dimples with a microvoid coalescence feature (Fig. 13d), which reveals a typical characteristic of ductile fracture after FSW. It is well accepted that the grain refinement is the key factor in improving the strength of FSW samples [19]. The decrease of the grain size will dramatically increase the amount of grain boundaries, which further impedes dislocation motion, resulting in the increment the strength of the material. In other words, after FSW, the grain refinement in SZ causes a higher strength than that in BM. So the tensile deformation mainly concentrated on BM. Unlike the tensile process of homogeneous materials, the resistance effect imposed by SZ on the BM has a considerable influence on the whole joints. When the strain localization takes place in BM, the stress state of this region transforms from the uniaxial state to the multiaxial state, raising the resistance to further deformation of the joints [62,63]. Therefore, the strength of the whole joints increases. Similar to the hardness results, the increment of the dislocation density and substructures with the rotational speed are also contributed to the improvement of strength in SZ, and then enhanced the strength of the whole joint. Different from the strength, the plastic deformation capacities of both the FSW joints were lower than that of BM sample. The traditional fine grain strengthening theory suggested that grain refinement can simultaneously improve the strength and ductility for most of materials [7]. However, numerous literatures reported that grain refinement would decrease the ductility of high nitrogen steels and nitrogen containing austenitic stainless steel [6,7]. As one of the higher nitrogen content of austenitic stainless steel, S32654 joint may also exhibit a lower ductility due to the grain refinement in SZ. High density dislocations and substructures in SZ caused by mechanical deformation are mainly responsible for the decrement in ductility of joints [64]. And the reduction in grain size of SZ is favorable to decrease the effective slip distance and further accelerate dislocations accumulation against grain boundaries [19]. So the plastic deformation capacities of the two FSW joints are lower than that of BM sample. Due to the higher density dislocations (Fig. 10) and substructures at 400 rpm (Figs. 7 and 10), its hardening in

Fig. 11. Microhardness profile of FSW S32654 at different rotational speeds: (a) as-welded joint and (b) the corresponding hardness values of BM and SZ.


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Fig. 12. Macrograph of transverse tensile specimens for FSW weldment before loading and after fracture: (a) BM, (b) 300 rpm, (c) 400 rpm and (d) before loading.

SZ is more obvious than that at 300 rpm. So the plastic deformation capacity of the joint at 400 rpm is lower. In addition, a few fine Cr- and Mo-rich σ phases were discovered in the SZ of 400 rpm while no precipitates were found at 300 rpm, as shown in Fig. 14. The σ phase is identified by selected area electron diffraction patterns with a tetragonal structure and parameters a = 0.879 nm and c = 0.452 nm. A previous investigation [29] found that the maximum temperature during FSW super austenitic stainless steel was above 1200 °C, exceeding the precipitation temperature of σ phase ranging of 650–1100 °C [65], so it is more likely that the σ phase formed during the cooling process of FSW. At higher rotational speed, the peak temperature is much higher, so the cooling time become even longer, which is more conducive to σ phase participation. Hence, σ phase was only found in the sample at 400 rpm. Numerous literatures reported that σ phase can increase the strength and decrease the ductility for metal materials [66,67]. Therefore, the presence of σ phase existed in SZ at 400 rpm exhibits a certain extent effect on the increment of strength and the decrement of plastic deformation capacity for the joint. In summary, the microstructure of TMAZ-AS and TMAZ-RS at 300 rpm was relatively symmetric compared with that of 400 rpm. No obvious BSs and σ phase formed in SZ at 300 rpm, which means more homogeneous microstructure. At the same time, lower rotational speed greatly decreased the wear degree of tool probe. Additionally, the strength of the whole joint at 300 rpm is essentially consistent

with that at 400 rpm, while the plastic deformation capacity of joint at 300 rpm is much higher. Therefore, the parameters of 300 rpm and 100 mm/min are the more suitable choice for FSW SASS S32654. 4. Conclusion In the present study, super-austenitic stainless steel S32654 plate was successfully friction stir welded at 300 and 400 rpm with 100 mm/min using W-Re tool. The microstructure evolution and mechanical properties of weldments were investigated. The main conclusions can be summarized as follows: (1) The sound joints without any defects were acquired at both rotational speeds. The SZ was characterized by fine-grained microstructure and almost no nitrogen content loses during FSW processes. (2) There exist obvious band structures containing W and Re in SZ at 400 rpm, which is attributed to the wear between tool probe and weldment. (3) The grain structure evolution is dominated by CDRX owing to the high SFE of S32654. The mean grain size slightly reduces with increasing the rotation speed from 300 rpm to 400 rpm, which reveals the strain rate plays a dominated effect on grain refinement of FSW S32654.

Fig. 13. Stress-strain curves (a), results of tensile test (b), SEM morphologies of fracture surface at low magnification (c) and high magnification (d) at 400 rpm.

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Fig. 14. (a) σ phase micrograph of the SZ at 400 rpm and (b) the corresponding energy spectrum.

(4) The combined action of grain refinement, high density dislocations and substructures improves the hardness and strength, but greatly reduces the plastic deformation capacity of joints. The more suitable welding processes are determined as 300 rpm and 100 mm/min.

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Corrigendum

Corrigendum to “Microstructure evolution and mechanical properties of friction stir welding super-austenitic stainless steel S32654” [Mater. Des. 118 (2017) 207-217] Huabing Li a,⁎, Shouxing Yang a, Shucai Zhang a,⁎, Binbin Zhang a, Zhouhua Jiang a, Hao Feng a, Peide Han b, Jizhong Li c a b c

School of Metallurgy, Northeastern University, Shenyang 110819, China College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China Guangdong Welding Institute (China-Ukraine E. O. Paton Institute of Welding), Guangzhou 510650, China

The authors regret to inform that Fig. 7 is not correct and same as Fig. 6. The correct form of Fig. 7 is as follows:

Fig. 7. (a) EBSD image quality maps, (b) grain boundary maps and (c) misorientation angle distributions for the BM (top) and the SZ of FSW at rotational speeds of 300 rpm (middle) and 400 rpm (bottom). LABs are depicted as pink and blue lines, HABs are marked by black lines on the grain boundary maps, respectively. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

Author would like to apologize for the inconvenience caused. DOI of original article: http://dx.doi.org/10.1016/j.matdes.2017.01.034. ⁎ Corresponding authors. E-mail addresses: [email protected] (H. Li), [email protected] (S. Zhang).

http://dx.doi.org/10.1016/j.matdes.2017.02.051 0264-1275/© 2017 Elsevier Ltd. All rights reserved.