2021 Volume 61 Issue 7 Pages 2150-2158
Phosphorus (P) addition is expected to simultaneously increase the strength and corrosion resistance of weathering steels. However, P causes solidification cracking in the fusion welding process and reduces the toughness of the steel. To avoid these problems, the P content of weldable SMA490AW weathering steel is currently limited to below 0.035 mass%. High P steels which are impossible to be joined by the fusion welding process, can be joined by a solid-state joining process, i.e., friction stir welding (FSW). Because the stir zone obtained by FSW contained very fine grains, its toughness was expected to improve. This study applies FSW to high-P weathering steels and examines the weldability of the product. The microstructural evolution and mechanical properties of the stir zone were investigated at different welding temperatures. The macroscopic cross-sectional observations of the FSW joints revealed crack-free structures even in steel containing 0.3 mass% P. Moreover, FSW significantly refined the grain structure in the stir zone. Consequently, the ductile-to-brittle transition temperature of the stir zone was approximately 150°C lower in the steel containing 0.3 mass% P and welded below A1 (average grain size = 2.5 µm) than in the base material (average grain size = 23 µm). It appears that the grain refinement by FSW overcomes the embrittlement caused by excessive P content.
Weathering steels containing alloying elements, such as Cu, P, Cr and Ni, exhibit excellent corrosion resistance under atmospheric conditions by forming a protective rust layer on their surface. These steels provide effective long-term anticorrosion performance without the need of painting and thus reduce the life-cycle cost of structures. Hence, they are widely used, mainly for bridges.
In weathering steels, P is a suitable alloying element because it improves weatherability and strength at minimal cost. Previous studies have reported that the weatherability of steel improves as the P content increases to at least 0.5 mass%.1,2,3,4,5,6) Moreover, the addition of a small amount of P effectively increases strength.7,8) However, P reduces the toughness of the base material (BM)9,10,11,12,13,14,15,16) as well as its weldability.17,18,19,20,21)
The negative effect of P on weldability is caused by solidification cracking, which is a form of high-temperature embrittlement that occurs due to lowering of the solidification temperature.21) Therefore, SMA490AW,22) a commercial weathering steel, contains less than 0.035 mass% P to ensure that it can be fusion-welded. SPAH23) is another example of weathering steel, which contains 0.1 mass% P; it exhibits excellent weatherability but poor performance in fusion welding.
To overcome the problems related to weldability, friction stir welding (FSW) can be employed. FSW is a solid-state joining process, in which materials are joined at a temperature below their melting point. This process is primarily used for aluminum alloys.24,25,26,27) Many studies on practical application of FSW for steels have been conducted.28,29,30,31,32,33,34) FSW does not involve solidification; hence, joining of steels with large amounts of P is possible.
Regarding the effect of P on toughness, many studies have reported that P segregation at the grain boundaries reduces toughness.9,10,11,12,13,14,15) Yamaguchi15) explained that an increase in P segregation at the grain boundary reduces the cohesive energy of the grain boundary, thereby reducing the fracture stress at the boundary. In contrast, Suzuki et al.9) showed that in steels with approximately ≧0.01 mass% C, C first segregates at the grain boundaries, resulting in site competition between C and P; consequently, the segregation of P at the grain boundary is inhibited, which leads to suppression of grain boundary fractures. Weathering steels contain approximately 0.1 mass% C; thus, grain boundary fracture is unlikely to occur if P content is suitably controlled and an appropriate fabrication process is employed. Nevertheless, the solute P within the grains generally hinders the gliding of dislocations due to a significant solid solution strengthening effect; thus, an increase in P content may causes cleavage fracture and reduce the toughness.
Grain refinement is effective for improving the toughness.35,36,37,38,39) For example, in ferritic steel, the ductile-to-brittle transition temperature (DBTT) decrease to 282°C when the grain diameter is refined from 20 μm to 1 μm.35) For grain refinement, in addition to conventional rolling technology,36,37) combinations of severe plastic deformation, such as equal-channel angular pressing (ECAP)38) and accumulative roll-bonding (ARB),14,39) and heat treatment have been proposed. During FSW, an extremely large shear strain is introduced together with heat generation, which induces dynamic recrystallization and leads to a significantly fine microstructure.31,32) Therefore, in high-P steels, FSW is expected to improve joint toughness through grain refinement.
In this study, weathering steels with 0.1–0.3 mass% P, which cannot be fusion-welded, were produced, and the weldability of these steels using FSW was investigated. Furthermore, the microstructures and mechanical properties of the BMs and the stir zones (SZs) of FSW were evaluated, and the effects of P content and grain refinement on the mechanical properties of the steels were investigated.
Table 1 lists the chemical compositions of the steels used in this study. SMA490AW is a commercial weathering steel with only 0.01 mass% P and can be fusion-welded. The samples of 0.1P and 0.3P are the high P weathering steels fabricated in the laboratory. These samples were prepared using the following method. The ingots were produced using electrolytic iron (99.9%), pure carbon (99.9%), pure silicon (99.9%), electrolytic manganese (99.9%), ferrophosphorus (Fe with 20 mass% P), pure copper (99.99%), and electrolytic chromium (99%). These metals were placed in an alumina crucible and melted in a high-frequency induction melting furnace. This melting process was performed in a 0.5-atm argon atmosphere. The final solidified sections of the ingots were cut and shaped into φ30 × 40 mm pieces. Subsequently, hot rolling was performed nine times at 1000°C until the plate thickness became 3.5–4.0 mm. This was followed by heat treatment at 1000°C for 10 min and then air cooling. The hot rolling temperature of 1000°C was higher than Ae3 point of the 0.1P and 0.3P samples, which implies that the hot rolling was conducted in the austenite phase. Finally, the plates were ground into a 3 mm thickness to remove the surface oxide layer.
| Material | Chemical composition (mass%) | ||||||||
|---|---|---|---|---|---|---|---|---|---|
| C | Si | Mn | P | S | Cu | Ni | Cr | Al | |
| SMA490AW | 0.12 | 0.20 | 1.14 | 0.01 | 0.002 | 0.32 | 0.10 | 0.48 | – |
| 0.1P | 0.10 | 0.51 | 0.09 | 0.10 | <0.001 | 0.49 | 0.01 | 0.39 | 0.03 |
| 0.3P | 0.08 | 0.50 | 0.09 | 0.27 | <0.001 | 0.49 | 0.01 | 0.40 | 0.02 |
FSW was conducted on the fabricated plates. The welding tool was made of a WC-based material and equipped with a columnar probe without a thread. The shoulder diameter, probe diameter, and probe length were 15.0 mm, 6.0 mm and, 2.9 mm, respectively. Two types of welding conditions were employed: welding at a temperature above A3 (rotation and welding speeds were 400 rpm and 150 mm/min, respectively) and below A1 (rotation and welding speeds were 80 rpm and 150 mm/min, respectively). K-type thermocouples were placed on the bottom surface and at the centerline of the materials to measure the changes in temperature during FSW.
Microstructure observations and Vickers hardness tests were performed on the cross-section of the joints perpendicular to the welding direction. The microstructures were observed using a field-emission scanning electron microscope (FE–SEM) equipped with an electron backscattered diffractometer (EBSD). The specimens for the SEM observations were mechanically polished and etched using a solution containing 4% nitric acid and 96% ethanol. The samples for the EBSD measurement were mechanically polished and electropolished at 15 V using a solution containing 10% perchloric acid and 90% glacial acetic acid. The Vickers hardness profiles of the joints were measured at a load of 2.94 N for 15 s. Tensile and impact tests were performed to investigate the mechanical properties of the BMs and the SZs. The tensile tests were performed at a strain rate of 5 × 10–3 s–1. The length, width, and thickness of the gauge section were 5, 2, and 2 mm, respectively. The direction of the gauge section was perpendicular to the welding direction, as shown in Fig. 1(a). During the tensile test, the strain was measured using the digital image correlation (DIC) method, and the nominal strain was evaluated with an initial gauge length of 4.5 mm. The impact test was performed using a miniaturized Charpy impact machine (Tanaka, MI20T-D05KJ type) for a specimen with a length of 20 mm, width of 1 mm, and thickness of 1 mm (Fig. 1(b)). This specimen had a V-shaped notch with a depth of 0.2 mm and a root radius of 0.08 mm. The specimens were picked up such that the notch was at the joint center for the SZ, and at a position of at least 10 mm away from the joint center for the BM. The DBTT in this study was defined as the temperature at the boundary between the completely brittle and brittle–ductile transition because the influence of specimen size was negligible.
Schematic of (a) sampling position of tensile test specimen and impact test specimen and (b) shape of impact test specimen.
Figure 2(a) shows the heat cycles during FSW under each welding condition, and Fig. 2(b) shows the pseudo-binary phase diagram of the Fe(0.1C-0.5Si-0.1Mn-0.5Cu-0.4Cr)-P steel. The maximum temperatures during the welding above the A3 and below the A1 temperatures were 1018°C (higher than the Ae3 temperatures of the 0.1P (907°C) and 0.3P (985°C) samples) and 695°C (less than the Ae1 temperatures of the 0.1P (739°C) and 0.3P (744°C) samples), respectively. These values indicate that each condition has be applied as intended.
(a) Heat cycles obtained by thermocouples located on the bottom surface at the centerline for SMA490AW; (b) Fe(0.1C-0.5Si-0.1Mn-0.5Cu-0.4Cr)-P pseudo-binary phase diagram evaluated by thermodynamic calculation, using the Thermo-Calc. TCFE9.
Figure 3 shows cross-sectional photographs of the friction-stir-welded joints for all samples. All samples exhibited defect-free joints.
Photographs showing the cross section of FSW joints for 0.1P and 0.3P samples under the two different FSW conditions.
Figure 4 shows the SEM microstructures of the BMs and SZs. The BM for each sample consisted of ferrite and pearlite structures. The ferrite grain diameters of the 0.1P and 0.3P samples are 49 μm and 23 μm, respectively. The SZ obtained at a welding temperature above A3 (SZ above A3) consisted of low-temperature transformation products, such as bainite and martensite, as well as ferrite and pearlite. These products were formed during the transformation from austenite. The amount of low-temperature transformation products in the SZ decreased with increasing the P content. Furthermore, after FSW, the grain diameters of the 0.1P and 0.3P samples decreased to 12 and 7.3 μm after FSW, respectively. The SZ obtained at a welding temperature below A1 (SZ below A1) consisted of ferrite and fine spherical cementite. The ferrite grain diameters of the 0.1P and 0.3P samples were 1.4 and 2.5 μm, respectively. Contrary to the expectation, the 0.3P sample had a larger ferrite diameter. The SZs of both materials were extremely refined via FSW below the A1 temperature. Moreover, the sizes of fine cementite for the SZs obtained below A1 were approximately 0.5 μm or less.
SEM microstructures of base material and stir zones under above A3 and below A1 conditions for 0.1P and 0.3P samples.
Figure 5 shows ferrite microstructures consisting of high-angle grain boundaries with a misorientation of >15°; these microstructures were obtained through inverse pole figure (IPF) maps acquired using the EBSD. The grains of the SZs of all samples were finer than those of the BMs due to dynamic recrystallization during the FSW. In particular, the grains of the SZ obtained below A1 were considerably finer than those of the SZ obtained above A3. The SZ obtained below A1 for the 0.1P sample showed relatively homogeneous fine grains with an average grain diameter of approximately 1.4 μm; in contrast, one of the 0.3P samples showed a bimodal structure with an average grain diameter of approximately 2.6 μm.
EBSD map showing grain boundaries of base materials and stir zones under above A3 and below A1 conditions for 0.1P and 0.3P samples.
Figure 6 shows the grain diameter distributions in the SZs obtained below A1 for the 0.1P and 0.3P samples. The dashed lines represent the average grain diameter. For the 0.1P sample, all grains had a diameter of <4 μm. Contrarily, for the 0.3P sample, several grains had a diameter of >4 μm, and a second peak existed at a grain diameter of approximately 4 μm (arrow in Fig. 6(b)). This result indicates a bimodal-like grain size distribution.
Area fraction distribution of the grain sizes in the stir zone under below A1 condition for (a) 0.1P and (b) 0.3P samples.
Figure 7 shows the engineering stress–strain curves of all samples. The curve for the BM of SMA490AW is also shown for comparison. In each sample, the tensile strength increased in the order of BM < SZ above A3 < SZ below A1. For both BMs and SZs, the 0.3P sample had a tensile strength higher than that of the 0.1P sample. In the stress–strain curves of the SZs obtained below A1, large yield point elongations were observed regardless of the P content. Furthermore, the 0.1P sample with relatively uniform fine ferrite and cementite exhibited an excellent balance between tensile strength and total elongation.
Nominal stress-nominal strain curves of base materials and stir zones under above A3 and below A1 conditions for (a) 0.1P and (b) 0.3P samples, and the base material for SMA490AW for comparison.
Figure 8 shows the temperature dependence of the absorbed impact energy for all samples. For the 0.1P sample (Fig. 8(a)), the DBTTs decreased in the order of SZ below A1 < SZ above A3 < BM. In particular, the SZ obtained below A1 exhibited ductile fracture even at a boiling point of liquid nitrogen, as well as high absorbed energy. Similar to those for the 0.1P sample, the DBTTs for the 0.3P sample (Fig. 8(b)) also decreased in the order of SZ below A1 < SZ above A3 < BM. However, compared to those for the 0.1P sample, the DBTTs for the 0.3P sample increased due to the increase in P content in both the BM and the SZs. The SZ obtained below A1 for the 0.3P sample exhibited brittle fracture at about −180°C, which deviated from the behavior of the 0.1P sample.
Temperature dependence of absorbed impact energy of base materials and stir zones under above A3 and below A1 conditions for (a) 0.1P and (b) 0.3P samples, and the base material for SMA490AW for comparison.
Figure 9 shows the fracture surfaces of an entirely brittle-fractured specimen (arrows in Fig. 8) for all samples. The BMs in each sample exhibited a cleavage fracture surface, where a river pattern was observed. Furthermore, the SZs obtained above A3 exhibited a cleavage fracture surface regardless of the P content. However, the unit size of the fracture surface of the SZ obtained above A3 was smaller than that of the BM. In contrast, the appearance of the fracture surface of the SZ obtained below A1 for the 0.3P sample was significantly different from that of the SZ obtained above A3. Intergranular fracture surfaces with a size of approximately 1 μm (arrow in Fig. 9(e)), corresponding to the ferrite grain size, were observed in some areas. No brittle fracture surface was observed on the SZ obtained below A1 for the 0.1P sample, because ductile fracture was exhibited at all temperatures.
SEM images showing the brittle fracture surfaces of base materials and stir zones under above A3 and below A1 conditions for 0.1P and 0.3P samples.
The cross-sectional macro-photographs (Fig. 3) show that the joints were defect-free for all samples under all conditions. FSW can join weathering steels containing 0.1 mass% C even with at least 0.3 mass% P.
P impairs fusion-weldability primarily because it promotes solidification cracking. The solid–liquid coexistence region during solidification has extremely low ductility; such a coexistence region is called the brittleness temperature range (BTR). Therefore, when the strain due to solidification shrinkage exceeds the critical strain, solidification cracking occurs.40) The P accumulates in the liquid phase during solidification and expands the BTR to cover lower temperatures. This expansion increases the risk of solidification cracking. The phase diagram in Fig. 2(b) shows that as the temperature drops from the stable temperature range of the liquid phase, the δ phase is formed first, followed by the γ phase. Subsequently, the three-phase (δ+γ+L) coexistence, and finally, the two-phase (δ+γ) coexistence is reached. However, at the final solidification stage, the concentration of P increases remarkably. As shown in Fig. 2(b), the increase in P concentration expands the temperature range of the δ+γ+L coexistence state to cover lower temperatures, resulting in high-temperature cracking. In contrast, the solid-state joining method FSW can join steels without the solidification process, resulting in defect-free joints.
P influences the microstructure and reduces the joint toughness in the fusion welding process,41,42,43) for example, by delaying the precipitation of cementite from the untransformed austenite and by promoting the formation of the martensite–austenite (M–A) component, which reduces toughness.41) In contrast, during FSW at the maximum temperature below A1, the transformation from ferrite to austenite does not occur, which prevents M–A formation. Moreover, as discussed in Section 4.2.3, significant grain refinement could improve the toughness of the SZ owing to FSW. Therefore, FSW of high-P steel prevents defects in the joints, as well as improves the mechanical properties of the joints.
4.2. Mechanism of Microstructure FormationFerrite grains were refined due to the increase in the P content in the BM and SZ obtained above A3 (Figs. 4, 5). This result could be attributed to the solute drag effect of P. P segregates at both the austenite and the ferrite grain boundaries.9,10,11,12,13,14,15,44) The solid solubility of P in austenite is lower than that in ferrite; hence, the grain boundary segregation is assumed to be substantial in austenite. Therefore, the austenite grains formed during the fabrication process of the BM and during the FSW above A3 are presumably refined as the P content increases, owing to the solute drag effect. Accordingly, the grain structures obtained through subsequent cooling also tend to be refined as the P content increases.
In the SZ obtained above A3, the amount of low-temperature transformation products, such as bainite and martensite, decreased due to the increase in the P content (Fig. 4). In general, P is a ferrite-forming element, so which raises the A3 temperature and promotes ferrite formation.45,46) Furthermore, as already mentioned, P leads to austenite grain refinement during FSW, promoting ferrite transformation. However, P also enhances hardenability.47) When the P content was increased in this study, the increase in the A3 temperature and the refinement of the austenite grains were presumably more significant than the increase in the hardenability. Consequently, the amount of low-temperature transformation products in the SZ was reduced with an increase in P content.
By contrast, in the SZ obtained below A1, microstructures composed of fine ferrite and fine cementite were obtained regardless of the P content (Figs. 4 and 5). During FSW below A1, the ferrite was refined through dynamic recrystallization, and the pearlite present in the BM was crushed and finely dispersed. However, the ferrite in the 0.3P sample was larger than that in the 0.1P sample and exhibited a bimodal structure. As already discussed, the amount of low-temperature transformation products in the SZ obtained above A3 for the 0.3P sample was less than that for the 0.1P sample. The same can be presumed for the amount of pearlite in the BM. The 0.3P sample had a smaller amount of pearlite than the 0.1P sample did (Figs. 4(a) and 4(d)). When the BM composed of ferrite and pearlite was friction-stir-welded at a temperature below A1, the nucleation frequency of the dynamically recrystallized grains in the region containing pearlite was presumably higher than that in the ferrite region; in addition, more significant grain refinement during FSW occurred in the region of prior pearlite in the BM. Therefore, the ferrite size in the 0.3P sample after FSW could be larger than that in the 0.1P sample, because the 0.3P sample had a smaller amount of pearlite in the BM. Moreover, the amount of pearlite in the BM of the 0.3P sample was small, and the distribution of the pearlite was considerably heterogeneous; hence, the 0.3P sample was assumed to have a prominent bimodal structure.
4.3. Mechanical Properties 4.3.1. Strength and ElongationIn the nominal stress–strain curves (Fig. 7) of the SZ obtained above A3, the 0.1P and 0.3P samples behaved differently; the 0.1P sample showed continuous yielding, whereas the 0.3P sample showed discontinuous yielding. This result could be due to the difference in the amount of low-temperature transformation products, such as bainite and martensite, between the 0.1P and 0.3P samples. Steels consisting of ferrite and low-temperature transformation products exhibit continuous yielding owing to the presence of mobile dislocations at the interface between soft and hard structures.48) Therefore, the continuous yielding of the 0.1P sample could be due to the larger amount of low-temperature transformation products. In contrast, the discontinuous yielding of the 0.3P sample could be because of the considerably small amount of martensite. Moreover, the large amounts of bainite and martensite resulted in poor elongation of the SZ obtained above A3 for the 0.1P sample compared with that for the 0.3P sample.
The strength of the SZ obtained below A1 was higher than that of the BM and the SZ obtained above A3 in both the 0.1P and 0.3P samples. Furthermore, the 0.1P sample had an excellent balance between strength and ductility. Compared with those of the BM, the yield point strength of the SZ obtained below A1 for the 0.1P and 0.3P samples were improved by 394 MPa and 353 MPa respectively. The Hall–Petch relationship holds between the yield strength and the grain diameter;49,50,51) thus, the improvement in the yield strength in the SZ obtained below A1 was mainly due to ferrite grain refinement.
The SZs obtained below A1 for both the 0.1P and 0.3P samples showed remarkable yield point elongation. Yield point elongation develops and increases with grain refinement.52,53,54,55) Therefore, the fine grains in the SZ obtained below A1 must lead to yield point elongation. Furthermore, the yield point elongation of the 0.1P sample was larger than that of the 0.3P sample; this phenomenon can be attributed to the finer grains of the 0.1P sample. In addition, the local elongation of the 0.1P sample was superior to that of the 0.3P sample. Grain refinement is reported to improve the reduction of area in a tensile test.56) Therefore, the SZ obtained below A1 for the 0.1P sample with grains finer than those of the 0.3P sample was assumed to exhibit improved local elongation. In contrast, the relatively small local elongation of the 0.3P sample might be due to the larger grains as well as the bimodal structure. In the bimodal structure of the 0.3P sample, the strain should be concentrated in the weak part of the structure with larger grains, resulting in reduced local elongation.
4.3.2. ToughnessCompared with those of the BM, the DBTTs of the SZs decreased both for the 0.1P and 0.3P samples. In particular, the SZ obtained below A1 for the 0.1P sample did not cause brittle fracture even at a boiling point of liquid nitrogen. Furthermore, the SZ obtained below A1 for the 0.3P sample exhibited excellent toughness with a DBTT as low as −175°C. The excellent toughness in the SZ obtained below A1 is assumed to be primarily controlled by grain refinement. Figure 10 shows the relationship between (d)−1/2 and DBTT for all samples, where d is the grain diameter. For comparison, previously reported data for C–Mn,57) C–Mn–Nb,58) and IF steels14) are also included. The slopes of the lines for the 0.1P and 0.3P samples were assumed to be equal to that of the C–Mn steel (0.15C-1.5Mn [mass%]), which has approximately the same amount of C. For the 0.1P sample, the DBTT of the SZ obtained above A3 was higher than the line of the BM. This result could be due to the presence of coarse martensite (>1 μm) in the SZ obtained above A3 (Fig. 4). Meanwhile, for the 0.3P sample, the DBTT of the SZ obtained below A1 was slightly higher than that of the BM. This phenomenon could be due to grain boundary embrittlement because the sample showed grain boundary fracture (Fig. 9). DBTT was slightly affected by the microstructure and fracture morphology. However, for the 0.1P and 0.3P samples, FSW both above A3 and below A1 significantly improved the toughness owing to grain refinement. In particular, the SZ obtained below A1 for the 0.3P sample showed a DBTT lower than that of a commercial SMA490AW (Fig. 8) specimen in which the amount of P was reduced to 0.01 mass%. This result indicates that grain refinement can overcome embrittlement due to high P content.
Effect of grain diameter on ductile-to-brittle transition temperature of base materials and stir zones for 0.1P and 0.3P samples; BM: base material, SZ: stir zone, M: martensite, B: bainite, IGF: intergranular fracture; Data of C-Mn-steel,57) C–Mn–Nb steel58) and IF steel14) are superimposed.
The difference between the DBTTs of the 0.1P and 0.3P samples with identical grain diameters d (Fig. 10) could be attributed to the effect of P on DBTT. Increasing the amount of P by 0.2 mass% increased the DBTT to 115°C. In contrast, Suzuki et al.10) showed that in Fe–C–P ternary alloy (0.04 mass% C), when the P content is increased from 0.11 to 0.34 mass%, the DBTT increases to 55°C. In this study, the increase in DBTT due to the increase in P content was larger than that reported by Suzuki et al. The amount of C and the size and fraction of the second phase also affect the DBTT; however, these factors will be investigated in future work.
The SZ obtained below A1 for the 0.3P sample exhibited grain boundary fracture at temperatures below the DBTT (Fig. 9(e)). Analysis of the fracture surface unit indicates that fracture occurred at the ferrite grain boundary. This implies that the fracture strength at the grain boundary was lower than that at the cleavage. In the 0.3P sample, the grain boundary segregation of P presumably increased because FSW had been performed below A1.
In general, dynamically recrystallized ferrites, such as those obtained via FSW below A1, have fresh ferrite grain boundaries. Therefore, no segregation of P at the grain boundary is expected immediately after the formation of dynamically recrystallized ferrite. Moreover, even if grain boundary segregation of P occurs during the cooling process after FSW, it should be difficult for P to segregate to the extent that grain boundary fracture occurs due to ordinary diffusion of P. However, during FSW below A1, dynamic grain boundary segregation of P possibly occurred due to severe plastic deformation. Several reports have shown that P dynamically segregates at grain boundaries due to deformation.59,60,61,62) The proposed mechanism is as follows: Deformation creates vacancies in the grains and forms supersaturated vacancy–solute complexes. A concentration gradient then appears between the grains and the grain boundaries, and the complexes easily move to the grain boundaries. Consequently, impurities rapidly segregate at the grain boundaries. Such non-equilibrium grain boundary segregation reportedly increases with increasing strain rate.62) Therefore, non-equilibrium grain boundary segregation occurs in FSW in which a large amount of strain is introduced. To prove the occurrence of dynamic segregation, grain boundary segregation must be measured; this constitutes a topic for future research.
The following findings were obtained by investigating the weldabilities, microstructure evolution, and mechanical properties of the FSW joints of high-P (0.1–0.3 mass% P) weathering steel, which cannot be joined through fusion welding.
(1) FSW provided defect-free joints at both welding temperatures below A1 and above A3. This occurred because FSW does not involve solidification, which causes defects in fusion welding. A weathering steel containing 0.1 mass% C can be joined via FSW even with at least 0.3 mass% P.
(2) The grains of the SZs were remarkably finer than those of the BMs, particularly during FSW at temperatures below A1. The DBTTs decreased with grain refinement. In the 0.3P samples, the DBTT of the SZ obtained below A1 (~2 μm) was reduced to approximately 150°C from that of the BM (~23 μm). This result indicates that grain refinement can overcome embrittlement caused by a high P content.
(3) As the P content increased, the grains in the BMs and SZs obtained above A3 were refined. This phenomenon could be due to the solute drag effect of P. Meanwhile, in the SZs obtained below A1, the 0.3P sample had large grains and showed a bimodal structure. This finding can be attributed to the small amount of pearlite in the BM of the sample, resulting in a low nucleation rate of the dynamically recrystallized ferrite during FSW and increased heterogeneity of the nucleation site.
(4) In the BMs and SZs obtained above A3, the fracture surface of an entirely brittle-fractured specimen showed a cleavage fracture. Conversely, in the SZ obtained below A1 for the 0.3P sample, intergranular fracture surfaces were observed. Non-equilibrium dynamic grain boundary segregation of P presumably occurred during FSW, leading to grain boundary fracture.
This work was supported by the JST-Mirai Program (Grant Number JPMJMI19E5), Research Group II of The Iron and Steel Institute of Japan, and a Grant-in-Aid for Science Research from the Japan Society for the Promotion of Science, 19H00826, Japan. The authors also wish to acknowledge Prof. M. Ohata of Osaka University for providing valuable advice during this study.
