Jia Xiao, Ming Li, Jian-ping Liang, Li Jiang, De-jun Wang, Xiang-xi Ye, Ze-zhong Chen, Na-xiu Wang, and Zhi-jun Li, Fine structure characterization of an explosively-welded GH3535/316H bimetallic plate interface, Int. J. Miner. Metall. Mater., 28(2021), No. 11, pp.1811-1820. https://doi.org/10.1007/s12613-020-2128-7
Cite this article as:
Jia Xiao, Ming Li, Jian-ping Liang, Li Jiang, De-jun Wang, Xiang-xi Ye, Ze-zhong Chen, Na-xiu Wang, and Zhi-jun Li, Fine structure characterization of an explosively-welded GH3535/316H bimetallic plate interface, Int. J. Miner. Metall. Mater., 28(2021), No. 11, pp.1811-1820. https://doi.org/10.1007/s12613-020-2128-7
Jia Xiao, Ming Li, Jian-ping Liang, Li Jiang, De-jun Wang, Xiang-xi Ye, Ze-zhong Chen, Na-xiu Wang, and Zhi-jun Li, Fine structure characterization of an explosively-welded GH3535/316H bimetallic plate interface, Int. J. Miner. Metall. Mater., 28(2021), No. 11, pp.1811-1820. https://doi.org/10.1007/s12613-020-2128-7
Cite this article as:
Jia Xiao, Ming Li, Jian-ping Liang, Li Jiang, De-jun Wang, Xiang-xi Ye, Ze-zhong Chen, Na-xiu Wang, and Zhi-jun Li, Fine structure characterization of an explosively-welded GH3535/316H bimetallic plate interface, Int. J. Miner. Metall. Mater., 28(2021), No. 11, pp.1811-1820. https://doi.org/10.1007/s12613-020-2128-7
An explosion-welded technology was induced to manufacture the GH3535/316H bimetallic plates to provide a more cost-effective structural material for ultrahigh temperature, molten salt thermal storage systems. The microstructure of the bonding interfaces were extensively investigated by scanning electron microscopy, energy dispersive spectrometry, and an electron probe microanalyzer. The bonding interface possessed a periodic, wavy morphology and was adorned by peninsula- or island-like transition zones. At higher magnification, a matrix recrystallization region, fine grain region, columnar grain region, equiaxed grain region, and shrinkage porosity were observed in the transition zones and surrounding area. Electron backscattered diffraction demonstrated that the strain in the recrystallization region of the GH3535 matrix and transition zone was less than the substrate. Strain concentration occurred at the interface and the solidification defects in the transition zone. The dislocation substructure in 316H near the interface was characterized by electron channeling contrast imaging. A dislocation network was formed in the grains of 316H. The microhardness decreased as the distance from the welding interface increased and the lowest hardness was inside the transition zone.
Ultrahigh temperature molten salt (UHTMS) thermal storage systems usually use chlorine salt and fluorine salt as the medium for heat transfer and storage. These systems typically operate at temperatures greater than 700°C [1]. Despite the high thermal efficiency, the UHTMS thermal storage systems put extremely high demands on their structural materials, including excellent high-temperature strength and matching with the molten salt. GH3535, a solid-solution-strengthened nickel-based alloy, has been regarded as a promising candidate because of its good high-temperature strength and molten salt anticorrosion [2–5]. However, this alloy is too expensive to be widely used in commercial UHTMS thermal storage systems. Bimetallic plates could be a more cost-effective structural material for UHTMS thermal storage systems and molten salt reactors. Bimetallic plates are usually composed of one inexpensive substrate with strong mechanical properties and one anticorrosion cladding layer, combining acceptable strength, excellent corrosion resistance, and cost effectiveness [6–8]. In this study, 316H stainless steel and a GH3535 alloy were selected as the substrate material and the anticorrosion cladding layer of bimetallic plates, respectively.
Explosive welding is generally selected to fabricate bimetallic plates [9–11], including steel/stainless steel [12], steel/Cu [13], steel/Ti [14], steel/Al [15], Al/Ti [16], Al/Cu [17], and Al/Ni [18] because of the ability to weld dissimilar metal materials with different thermal expansion coefficients or mechanical properties. The structure and performance of the interface directly affect the quality of the explosive welding plates, which has gained the attention of researchers. Zhang et al. [19] examined the heterogeneous microstructure and mechanical properties of explosively-welded 2205 stainless steel/X65 pipe steel bimetallic sheets. The microstructure investigation illustrated two kinds of bonds at the wavy interface of 2205/X65: metal–metal and metal–solidified melt. Different structures, such as pores, localized melted zones, peninsulas, and island morphology, were obtained close to the interface. Kaya et al. [20] performed an investigation of explosive welding of Grade A ship steel-duplex stainless steel composites. The completely straight joining interface was transformed into a wavy interface near the explosion zone. With increasing distance from the explosion zone, the wavelength and amplitude gradually increased due to increasing cold deformation. Akbari Mousavi and Farhadi Sartangi [21] defined the weldability field of cp-titanium/AISI 304 stainless steel and studied the effect of the explosive load on the bonding interface. A transition region from a smooth interface to a wavy interface was obtained with an increase in explosive load. The interface was highlighted by a special sharp transition between two materials; however, local melting also occurred on the front slope of the wavy interface under high explosive loads. In contrast, fine characterization of the explosively-welded interface between a nickel-based alloy and stainless steel has rarely been reported. The structure, elemental composition, and strain distribution of the explosively-welded interface reflect the rationality of the welding procedure and the post-weld heat treatment of the bimetallic plates. Therefore, the fine structure characterizations of the interface are essential for the explosive welding of the GH3535/316H bimetallic plates.
In this study, the explosion-welded interface of a GH3535/316H bimetallic plate was finely characterized using scanning electron microscopy (SEM), energy dispersive spectrometry (EDS), electron probe microanalyzer (EPMA), and electron backscattered diffraction (EBSD), which revealed the macro–micro morphology, elemental composition, and strain distribution of the interface of explosively-welded GH3535/316H.
2.
Experimental
In this investigation, 316H stainless steel was chosen as the substrate to ensure high-temperature mechanical properties, and GH3535 with excellent anticorrosion was used as the cladding plate. The chemical composition of the GH3535/316H bimetallic plates are listed in Table 1. In the fabrication of the GH3535/316H composite, a flyer plate (316H) with a dimension of 400 mm × 500 mm × 5 mm and the base plate (GH3535) with a dimension of 380 mm × 480 mm × 25 mm were arranged in a classical parallel scheme. The explosive material was rock expanded ammonium nitrate explosive, and the detonation velocity was 2200 m/s. The welding assembly was placed on a sand foundation and all explosions were performed on soil ground. After explosive welding, the final thicknesses of 316H and GH3535 were 25 mm and 5 mm, respectively, and the heat-treatment scheme of plates was 650°C for 1 h to eliminate the stress at a heating rate of 166°C/h, and the samples were air cooled.
Table
1.
Chemical compositions of GH3535 and 316H wt%
Rectangular samples with dimensions of 15 mm × 15 mm × 2 mm were cut from the interior of the bimetallic plate, and away from the explosion starting point by spark cutting. The samples were successively ground by SiC papers to 2000 grit and polished with a 0.05 μm Al2O3 suspension. The macro–micro morphology of the bimetallic plate interface was examined by SEM (Carl-Zeiss, Berlin, Germany). The elemental composition and distribution of the bimetallic plate interface were detected by EDS (Oxford instruments, Oxford, United Kingdom) and EPMA (Shimadzu Corporation, Kyoto, Japan). The microstructures and strain distribution of the bimetallic plate interface were characterized by EBSD (Oxford instruments, Oxford, United Kingdom). Before EBSD examination, mechanical polishing, and vibration-assisted final polishing were conducted to obtain a suitable surface for EBSD. Microhardness tests were performed using a Vickers microindenter (ZwickRoell GmbH & Co. KG, Ulm, Germany) testing machine under a load of 0.98 N with a dwell time of 15 s. Measurements were performed on the GH3535, 316H, and interface areas, and the hardness was the average of three indentations.
3.
Results and discussion
3.1
Macromorphology of the bimetallic plate interface
The macromorphology and several typical structures of the bonding interface of bimetallic plates are displayed in Fig. 1. A metallurgical bond interface was formed between GH3535 and 316H during explosive welding. A wavy morphology was observed at the bond interface of the GH3535/316H bimetallic plates, whose wavelength and amplitude was approximately 310 and 120 µm, respectively. The wavy morphology can be ascribed to variations in the velocity distribution at the collision point and periodic disturbances of the materials [21]. The welding interface can be changed from straight to wavy with increasing explosive loading [22]. Therefore, a wavy interface can be obtained.
Fig.
1.
SEM images of the welding interface between GH3535 and 316H: (a) wavy interface; (b) high-resolution image (HRI) of B in (a); (c–f) HRI of rectangles corresponding to insets in (b).
As demonstrated in Fig. 1(a), wide and narrow transition zones were observed at the front and back slopes of the waves, respectively. Other locations on the waves indicate a clear demarcation between the two plates. Fig. 1(b) shows a characteristic wave at the bonding interface as marked by rectangle B in Fig. 1(a), in which four different structures (marked by rectangles C, D, E, and F) were observed. As demonstrated in rectangle C, one peninsula-like transition zone was observed to extend into the 316H matrix. In rectangle D, a portion of the 316H matrix was an island-shaped inlaid in the transition zone. The peninsula morphology is produced by the combined effect of the explosive force and metal vortex flow. When the explosive force and metal vortex flow are significantly intense, an island morphology can be easily formed [19]. The transition zone with porosity and a defect-free zone were observed in rectangle E and rectangle F, respectively. Generally, the occurrence of a transition zone is related to the explosive load. Because the waviness increases with explosive loading, a severe wave deformation may lead to wave crest separation and drifting in some cases, and vortices are formed at the interface with excessive explosive loads [21].
EPMA was utilized in order to obtain different element distributions at the interface. Fig. 2 displays the EPMA mappings of Fe, Ni, Mo, and Cr. Element diffusion through the interface was not evident, except for in the transition zones. The concentration of each element in the transition zone corresponds to the average value of 316H and GH3535. Additionally, EPMA mapping of the elements also suggests that there was the formation of an interlayer (3–10 μm) along the waves, with a chemical composition similar to the transition zone.
Fig.
2.
Area distribution of elements at the explosively-welded GH3535/316H bimetallic plate interface: (a) target area of EPMA mapping; (b–e) EPMA mappings of Fe, Ni, Mo, and Cr of the typical wave in (a).
EDS line scanning was also selected to further analyze the element distribution at different positions along the wavy interface (Fig. 3). The line distributions of the elements of the narrow and wide transition zones at the back and front slopes of the waves are presented in Figs. 3(b) and 3(f), respectively. The step-like element distributions indicate that the transition zones are not element diffusion regions. The content of each element in the transition zone was between that of 316H and GH3535, which is consistent with the EPMA results (Fig. 2). The width of the element diffusion layer between the transition zone and 316H matrix or GH3535 alloy was approximately 2–5 μm (Figs. 3(b) and 3(f)). As shown in Fig. 3(b), the peaks of Mo and Si were also observed at the scanning line b between 30 and 60 μm from the GH3535, which indicates the existence of carbides with rich Mo and Si in the matrix of the GH3535 alloy [5]. Figs. 3(d) and 3(g) present the line distributions of the elements at the peak and trough of the wavy interface without transition zones. The width of the element diffusion layer between the substrate plate and cladding plate was approximately 2–5 μm, and the step distribution of the elements was not observed. Therefore, there are two types of interface bonding modes in the bimetallic plate, matrix–matrix and matrix–transition zone–matrix. Figs. 3(c) and 3(e) illustrate that the element distributions of the island and peninsula structures at the interface were a staggered morphology. The transition zone and substrates were mixed due to the strong explosive force and metal vortex flow.
Fig.
3.
Line distributions of the elemental composition of explosively-welded GH3535/316H bimetallic plate interface: (a) target area of the EDS line distributions; (b–g) mass percentage of the elements on lines b–g in (a).
3.2
Micromorphology of the bimetallic plate interface
As demonstrated in Fig. 4, the transition zone and two plates near the interface were further characterized by backscattered electron imaging. The regions marked by the rectangles B, C, D, E, F, G, H, I, and J were observed at higher magnification. During explosive welding, high-pressure shock waves were generated when two plates collided, and strong plastic deformation occurred in the contact area. The matrix grains on both sides of the interface showed filamentary plastic flow and the plastic flow appears as helical near the transition zone, as demonstrated in rectangles H and I. Rectangles B–D show the grain morphology inside the wide transition zones at the front slope of the wave. Fine-equiaxed grains were observed in the wide transition zones in contact with GH3535 (rectangle B), and columnar grains (rectangle C) and equiaxed grains (rectangle D) were observed in the interior of the wide transition zones. Rectangle J presents the shrinkage porosity in the wide transition zones. Rectangle F shows the fine columnar grains and equiaxed grains in the narrow transition zones at the back slope of the wave. Rectangle G shows that the two materials were directly in contact with each other with a clear dividing line at the interface without transition zones. Rectangle E illustrates that the fine-equiaxed grains were observed at the vortex tip where the transition zone was contacted with GH3535.
Fig.
4.
(a) Backscattered electron image of the transition zones at the welding interface and two plates near the interface; (b–j) HRI of the corresponding rectangles B–J in (a).
To determine the characteristics of the interface between the joined plates, the parameters of the welding along the whole length of the specimen were measured and the equivalent thickness of the melted area (ETR) was calculated. The ETR coefficient (ηETR) is determined according to Eq. (1) [23–24]:
ηETR=∑ni=1PiL
(1)
where P is areas of the melted zone (μm2), i is the number of melted zone, and L is the interface line length (μm).
The value of ηETR, which should not exceed 10, is 0.08. Theoretically, ηETR can be considered an indicator of the bond’s quality, which also presents the number of intermetallic layers created between the joined materials [25]. The higher the ηETR value, the lower the mechanical properties of the material.
EDS mappings of the transition zone at the welding interface and two plates near the interface are presented in Fig. 5. Fig. 5(a) shows the EDS mapping (i.e., the position of the wide transition zone shown in Fig. 4) of the wavy interface area. The EDS mappings of Fe, Ni, Mo, Cr, and C in the area are shown in Figs. 5(b)–5(f), respectively. The element distribution of the transition zone and two plates near the interface were uniform. Additionally, the same elemental composition was discovered in the fine-equiaxed grains, columnar grains, and equiaxed grains in the transition zone. In rectangle E in Fig. 4, the elemental composition of the fine-equiaxed grains at the vortex tip was the same as that of GH3535. Therefore, the fine-equiaxed grains should be formed by recrystallization of GH3535. Because of the disappearance of oxygen and carbon aggregation in the EDS mappings, the precipitate or oxide was not observed in the transition zone at the interface.
Fig.
5.
EDS mappings of the transition zone of the bimetallic plate interface and two plates near the interface: (a) target area of EDS mapping; (b–f) EDS mappings of Fe, Ni, Mo, Cr, and C for (a).
As shown in Fig. 6, EBSD was used to investigate the grain morphology and size of the transition zone. According to the grain morphology, the transition zone was divided into four different regions: the matrix recrystallization region (A), the fine grain region (B), the columnar grain region (C), and the equiaxed grain region (D). Fig. 6(b) presents the average grain size of each region. Generally, the solidification structure contains surface fine grains, columnar grains, and central equiaxed grains. In this study, the transition zone also showed similar morphological characteristics. The density difference or adiabatic heating of gases compressed between the plates causes localized melting of the vortex of a wave, which is due to the trapped jet adiabatic heating inside vortices at the front and back slopes of the waves [21]. The melting zone surrounded by relatively cold metal is subjected to a very high cooling rate and the estimated cooling rate is about 105–107 K/s [21]. The rapid solidification leads to fine-equiaxed and columnar grains in the local melted region (transition zone). Because of the fastest cooling rate of the area (B) in the transition zone that directly contacts with two plates, and the substrates can be used as the base of nonuniform nucleation, fine-equiaxed grains are formed in this area. Due to the fastest heat dissipation perpendicular to the interface, the columnar grains grow in the direction perpendicular to the interface. Fig. 6(a) also shows that the grain morphologies were different between nearby GH3535 and 316H in the transition zone. The grains in the transition zone in contact with GH3535 were mostly fine-equiaxed (B1, B2), and the grains inside the transition zone near GH3535 were mostly equiaxed (D). Fewer fine-equiaxed grains were produced in the transition zone in contact with 316H (B3), and most of the grains near 316H in the transition zone were columnar (C1). In the columnar grain region C2, the shape of the grains near 316H was longer than that of the columnar grains near GH3535. The grains were more likely to form as columnar in the transition zone near 316H, which grows faster than the columnar grains near GH3535 due to faster heat dissipation of 316H. In the transition zone near GH3535, equiaxed grains were easily formed due to the slower heat dissipation of GH3535.
Fig.
6.
Fine structure characterization of the wide transition zone (local melted region) in Fig. 5(a): (a) typical area division of the transition zone; (b) average grain size of each typical region.
Similar to the solidification structure, the shrinkage porosity can also be formed in the transition zone. When the melting region is formed at the vortices, the deformation of the two plates near the interface is converted to a circular movement, and the melting zone is stirred vigorously, resulting in miscibility of GH3535, 316H, and the molten zone. This miscibility promotes the formation of peninsula- and island-like structures, and the severe deformation of the matrix material also promotes the formation of a spiral plastic flow on the two sides of the interface. The severe plastic deformation in the matrix near the transition zone also results in the unindexed zone in the EBSD mapping. The generation of the fine-equiaxed grains (A) in GH3535 at the interface may be because the heat in the local melted region diffuses slowly in GH3535, which causes recrystallization of the GH3535 near the interface. Concurrently, high strains exist at the tip of the substrate mixed with the local melted region, which makes it easy for strain-induced recrystallization to occur.
3.3
Interface strain analysis
As shown in Fig. 7, EBSD was also selected to characterize the interface morphology and strain distribution of the bimetallic plates. The grain sizes of the GH3535 alloy and 316H steel were approximately 62 and 46 µm, respectively. There was no significant change in the grain size of the two plates away from the interface, while the grain size in the transition zone was much smaller than that of the matrix (Fig. 7(c)). Fig. 7(d) shows the mapping of the kernel average misorientation (KAM) of the interface and surrounding plates. KAM is the average misorientation between every pixel and its surrounding pixels according to EBSD, which can evaluate the local strain and response to the density of dislocation in crystalline materials [26–27]. The strain degree becomes stronger with the distance to the interface, which was intensely higher than that in the two plates. Because of melting and resolidification, the stress from explosive welding in the transition zone was released, and the transition zone revealed a higher EBSD indexed ratio and weaker strain concentration compared with the matrix of the GH3535 alloy and 316H steel (Fig. 7(d)). The high-pressure shock wave and vortex structure generated during the explosive welding led to severe plastic deformation of the interface, which caused high strain at the interface. The map of KAM shown in Fig. 7(d) had a low-resolution region (white area in Fig. 7(d)) with a width of about 100 µm near the transition zone. Based on the band contrast (BC) map (Fig. 7(b)), there were a large number of black lines in the low-resolution region. When the quality of the Kikuchi diffraction pattern is poor, it is displayed in black in the BC map. The black line in the BC map can reflect crystal lattice defects, such as slip lines [28]. A large number of dislocations caused by the vortex in the matrix near the interface produces a low resolution in that region. The strain in this low-resolution region should be greater than that of the nearby GH3535 and 316H. Electron channeling contrast imaging (ECCI) was used to further analyze the lattice distortion of this region.
Fig.
7.
Interface morphology and strain distribution of the bimetallic plates: (a) target area of the EBSD mapping; (b) band contrast (BC) map; (c) grain morphology from the two plates to the interface; (d) mapping of KAM from the two plates to the interface.
The strain distribution in the transition zone at the bonding interface of the bimetallic plates is presented in Fig. 8. The strain of the small columnar grains and equiaxed grains in the transition zone was significantly less than that of the surrounding matrix of GH3535 and 316H. Due to the serious wave-shaped deformation caused by the explosion load, a stress concentration occurred in the vortex structure. In contrast, because the two plates inside the vortex melted and resolidified to form a local melted region, the internal stress of the vortex was released, resulting in low strain in the local melted region. The strain was greater because of the severe plastic deformation at the interface between the local melted region and substrate. The strain near the shrinkage porosity caused by rapid solidification in the local melted region was also high. Materials for a thermal storage system, such as molten salt tanks and loops, can sustain higher thermal stresses during thermal cycling over its service life. For the interface between GH3535 and 316H, the interface between the transition zone and two matrices, as well as the shrinkage porosity in the transition zone may become preferential initiation and expansion sites of thermal fatigue cracks due to the concentrated high strain. The position shown by the yellow dotted line in Fig. 8 is the fine recrystallized grains in GH3535 adjacent to the transition zone. This position shows a deficient strain due to the recrystallization that releases the stress.
Fig.
8.
Grain boundary and mapping of KAM of the wide transition zone (local melted region).
ECCI was used with SEM to characterize the dislocation substructure of the 316H closed to the interface, as shown in Fig. 9. When the orientation contrast of the grains is shown as black in ECCI, the substructures of dislocation and twin in the grains will be displayed in white and bright colors. Therefore, ECCI can characterize dislocations and twins from a wide range of perspectives [29]. For 316H near the interface, areas of dense dislocations were observed inside the grains, and the equiaxed grains were elongated in the direction of the deformation (Fig. 9(a)). Fig. 9(b) is the HRI of the position shown by the white dotted line in Fig. 9(a). A complex network formed by a large number of dislocations was distributed in the grains near the interface. Notably, a large amount of deformation caused by the vortex structure in the matrix near the interface made it difficult to identify the matrix grains in contact with the interface, and presented a curved fibrous structure along the interface as shown by the yellow arrow in Fig. 9(a). Additionally, severe lattice distortion at this location resulted in a low indexed region in the EBSD mapping (Fig. 7). Dislocations were denser in the fibrous structure than in the nearby 316H. Reportedly, the dislocation substructure and twin substructure are refined with an increase in strain [29], which indicates that the fibrous structure has a higher strain concentration and dislocation density than the nearby matrix.
Fig.
9.
ECCI images of (a) 316H near the interface and (b) HRI of the area identified in (a).
The nonuniform strain distribution was close to the explosively-welded interface of GH3535/316H. Fine columnar and equiaxed grains in the transition zone and the recrystallized fine grains showed lower strain, while the interface between GH3535 and 316H, the interface between the transition zone and the two plates, and the solidification defects in the transition zone indicated higher strain. Therefore, subsequent welding parameters and post-weld heat-treatment methods should be formulated to eliminate this uneven distribution of strain.
3.4
Microhardness
As shown in Fig. 10, the microhardness of GH3535 was slightly greater than that of 316H and the microhardness generally decreased as the distance from the welding interface increased. Similar results have been reported [30]. It is generally accepted that this trend could be attributed to the high plastic deformation in the welding zone and high lattice distortion relative to the area at a distance. Fig. 10 also shows that the microhardness did not change substantially in either metal if the distance from the interface was greater than 530 µm. The maximum Vickers hardness of 316H near the interface was approximately HV 488, which was approximately 61.6% greater than the HV 302 hardness of 316H away from the interfacial region. The maximum Vickers hardness of GH3535 near the interface was approximately HV 580, which was approximately 75.2% greater than the HV 331 of the GH3535 away from the interfacial region. Additionally, the lowest hardness value was inside the transition zone, which is attributed to the local melting and resolidification of the transition zone to release stress and eliminate lattice distortion caused by explosive welding.
Fig.
10.
Vickers microhardness variation across the interface using a load of 0.98 N and 15 s dwell time.
The microstructure and interface strain distribution of explosively-welded GH3535/316H bimetallic plates were investigated with emphasis on the fine structure characterization of the transition zone at the interface and surrounding matrix. The conclusions of the study are follows.
(1) The bonding interface of bimetallic plates showed a wavy morphology. Two bonding modes were discovered at the interface: matrix–matrix and matrix–transition zone–matrix. The transition zone, peninsula-, and island-like structures were observed at the interface.
(2) The transition zone was produced by rapid solidification of the vortex structure after local melting. Four typical regions were divided into the transition zone and the surrounding matrix: the matrix recrystallization region, the fine grain region, the columnar grain region, and the equiaxed grain region. The shrinkage porosity was also observed in the transition zone.
(3) The grains in the transition zone near GH3535 were mostly fine-equiaxed grains and equiaxed grains, while the grains in the transition zone near 316H were mostly columnar grains. The grain size in the transition zone was much smaller than that of the substrates.
(4) The strain increases with a decrease in distance from the interface. After melting and resolidification, the transition zone and the surrounding recrystallization region in the GH3535 alloy had a very low strain. The interface and solidification defects in the transition zone concentrated higher strain.
(5) A complex network of dislocations was distributed in the grains of the 316H near the interface, and the equiaxed grains elongated along the deformation direction. The matrix in contact with the interface presented a fibrous structure along the interface. The strain concentration and dislocation density of the fibrous structure was greater than that of the surrounding matrix, and the severe lattice distortion at this location promoted low resolution of the EBSD.
(6) The microhardness generally decreased as the distance from the welding interface increased. The lowest hardness occurred inside the transition zone.
Acknowledgements
This work was financially supported by the National Natural Science Foundation of China (Nos. U2032205, 51971238, and 52005492), the Shanghai Outstanding Academic Leaders Plan (21XD1404300), the Natural Science Foundation of Shanghai (Nos. 18ZR1448000, 19ZR1468200, 20ZR1468600, and 21XD1404300), the Shanghai Sailing Program (Grant No. 19YF1458300), and the Youth Innovation Promotion Association, Chinese Academy of Science (No. 2019264).
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