3.1 Weld formation
Figure 4 shows the macroscopic shape of the weld seam and the macroscopic shape of the joint cross-section, the cover layer weld is oxidized slightly, the weld surface is faint yellow. The filler material was well-spread in the groove and fully fused with the base material from both sides. There are only a few porosity defects in the joint cross-section, and the diameter of the porosity does not exceed 0.2 mm. The fusion between the layers and the side walls of the joint is well, and the fusion width of the weld is within the range of 7-7.5 mm, the height of the cover layer is about 3.5mm, in addition, the height of other filling welds is about 1.8-2.0mm.
Comparing the original model with the deformation data, the joint deformation cloud is depicted in Fig. 5, disregard of the displacement deviation caused by the unfilled position of the bevel and extracting the deformation data only for the edge of the test plate. The deformation results show that there is a slight warpage of the test plate after welding, and the maximum deformation of the test plate edge is only 0.25 mm, and the maximum warpage angle is calculated to be about 0.19°. Therefore, the symmetrical U-shaped groove was used and alternating welding sequence can suppress the deformation after welding, effectively.
3.2 Microstructure examination
The microstructure of the interlaminar weld zone is intercepted and studied in three different areas of the upper, middle and lower parts of the joint, respectively. The interlayer microstructure characteristics of different areas of the joint are analyzed, and the variability of the tissue morphology of the weld zone between different areas of the layers is quantitatively compared. Figure 6 demonstrates the interlayer microstructure of the cover layer. As seen in Fig. 6 (d) and (e), there are distinct upper boundary of the remelting zone and fusion line in the upper interlayer region of the joint. The measured statistics show that the width of the remelting zone is in the range of 900 ~ 1000 µm. Figure 6 (b) and (c) shows the microstructure near the interlayer fusion line, and it can be seen that from the figure that a short needle-like α' phase appears near the fusion line. Above it, i.e., the part of the cover layer weld, the widmanstatten structure consisting of long and wide clusters of needle-like martensite regularly arranged, among where the white color is the needle-like α' phase and the black border between the phases is the β phase; while below the fusion line, the basketweave structure consisting of cross-arrangement of needle-like martensite appears inside the coarse β columnar crystal. In addition, the second welding heat cycle in the upper weld, the fusion line below the part of the original β columnar grain boundary gradually tends to be "blurred".
Figure 7 shows the interlayer microstructure morphology of the middle and lower part of the joint. Unlike the interlayer organization of the upper part of the joint, the upper boundary of the remelting zone in the middle and lower part disappears completely due to the occurrence of multiple solid-state phase transformations. Therefore, in this part, the interlayer microstructure near the fusion line of the two regions was intercepted for analysis. Figure 7 (c) and (d) shows that short needle-like martensite tissue also appears near the fusion line in the middle of the joint. As shown in Fig. 7 (d), the statistical measurements of the martensite length dimensions of ten groups in the field of view are randomly performed, and it can be seen that the average martensite length near the fusion line is around 12 µm. Figure 7 (f) and (g) show that although the width dimension of the columnar crystals in the lower part of the joint is smaller, the widmanstatten structure consisting of coarse needle-like martensite clusters still appears above the fusion line. In addition, as seen in Fig. 7 (g), the original β columnar grain boundary below the fusion line in the lower part of the joint has nearly disappeared.
The formation mechanism of the short needle-like α' phase near the upper boundary of the remelting zone and the fusion line is studied, as shown in Fig. 8. Ti-6Al-4V titanium alloy laser welding is a rapid heating and cooling process, in the process of rapid cooling of the high-temperature β phase, solid solution in which the alloying elements are too late to diffuse. At this time, the occurrence of martensitic phase transformation belongs to the non-diffusion type phase transformation. Many studies have shown that, similar to the liquid-solid phase transformation, martensite phase transformation is also a nucleation and growth process [8]. Therefore, this paper analyzed the formation mechanism of the short needle-like α' phase from the perspective of nucleation.
As shown in Fig. 8 (a), during the narrow gap laser welding of Ti-6Al-4V titanium alloy, the liquid molten pool can be divided into two parts: the lower weld, which is secondarily melted by the laser heat, and the new weld, which is formed by filling in the molten wire. While the formation of the upper weld, the surface area of the lower weld absorbs the laser energy and remelts. Subsequently, the wire melts and transitions into the molten pool under the action of the laser beam. During the transition of the molten droplet, a strong stirring effect on the formed molten pool is inevitable, as shown in Fig. 8 (b). Therefore, near the upper boundary of the remelting zone, there are obvious energy, structural as well as compositional undulations, which provide the nucleation conditions for the subsequent martensite phase transformation. At the same time, after solidification of the weld, a large number of defects such as interstitial atoms, dislocations, and laminations are retained near the boundary of the remelting zone due to the stirring behavior of the melt pool. When the high-temperature β phase cooled to the phase transition temperature, the high-density defects will provide a large number of nucleation conditions for the formation of the α' phase. Near the boundary, the α' phase first nucleates and grows at the location of point defects, line defects, etc. As a result, a large number of short needle-like α' phases are formed near the boundary of the remelting zone.
The formation mechanism of short needle-like α' phases near the fusion line is similar to the above-mentioned case. As shown in Fig. 8 (c), compared with the upper boundary of the remelting zone, the stirring behavior of the molten pool near the fusion line is relatively weakened. However, after liquid solidification, the original solid-liquid interface also retains a certain number of nucleation prone positions such as point defects, dislocations and stacking faults. In the meantime, because it is close to the base metal, the undercooling near the fusion line is higher. Many literatures show that when the undercooling increases, the critical nucleation decreases successfully, which means that the number of nucleation increases. Therefore, a large number of α' The phase nucleates and grows near the original solid-liquid interface, forming a short needle shape α'. The fusion zone formed by aggregation is shown in Fig. 8 (c).
Meanwhile, some of the β columnar crystals below the fusion line showed the fuzzy or even disappearance of grain boundaries, and this phenomenon appeared in several interlayer weld regions such as the upper, middle and lower parts of the joint, as shown in Fig. 9 (a) and (b). To explain the above phenomenon, this study is carried out to investigate the microstructure evolution behavior of the heat-affected zone below the fusion line. As seen in Fig. 9 (a) and (b), the grain boundary disappearance region is basically located within 200 µm below the fusion line, and this region is subject to the same strong thermal action during the forming of the upper layer weld, although remelting does not occur. Due to the different cooling conditions, the microstructure before and after the solid-state phase transformation may differ, at this moment the original β columnar grain boundary morphology is weakened.
In addition, the grain boundary is a collection of a large number of point defects and line defects (dislocations). First, the local area below the fusion line experiences a sudden temperature rise and increased atomic activity due to the thermal influence of the upper weld. As a result, point defects concentrated at grain boundaries start to migrate under the action of driving force, and some vacancy defects migrate to dislocations or disappear in migration by combining with interstitial atoms. At the same time, due to the increased atomic activity, in addition to point defects, dislocations in the role of internal stresses in the weld is activated and slip, while the slip process of dissimilar dislocations occur in combination and offset. Therefore, the grain boundary near the point defects, line defects density will be further reduced. In summary, due to the tissue evolution of the secondary solid state phase transformation and the influence of defect migration, the fusion line below part of the original β columnar grain boundaries in the upper layer of the weld forming process under the influence of heat gradually tends to blur or even disappear.
3.3 The weld joint mechanical properties
Figure 10 shows the path selection schematic for the micro-Vickers hardness test, and a total of four paths were selected. The hardness test results are shown in Fig. 11. The results show that the joint of heat-affected zone and weld hardness are much higher than the base material after the martensitic phase transformation is completed. The weld area below the cover layer shows an obvious "softening" phenomenon. The average hardness of the cover layer weld is 379.4 HV, while the average hardness level of the other layers of the weld is only 339.1 HV. Analysis shows that the upper layer of the weld forming process of multiple welding thermal cycle-induced tissue evolution and defect migration behavior is the main reason for the "softening" of the lower layer of the weld.
In order to further investigate the influence of interlayer tissue variability in different areas of the joint on the mechanical properties, the microstructure of specimen #5 was analyzed. The failure of the specimens occurred in the parent material area, and the microscopic morphology near the fracture surface of specimen #5 showing in Fig. 12. The fracture surface penetrates a large number of equiaxed grains in the base material area. Therefore, the fracture mode of the tensile specimens can be tentatively judged as grain penetration ductile fracture.
Figure 13 (a) shows the displacement-stress variation curves of different tensile specimens. During the tensile process, with the displacement is increased, the specimen first enters the elastic deformation stage in where the stress level increases linearly. Subsequently, when the specimen deformation exceeds the critical value, the tensile specimen enters the plastic deformation stage in where the deformation is irreversible. Finally, with the increasing loading force, the microcracks inside the specimen continued to expand to the critical length and the specimen failed to fracture. Comparing the test results of tensile specimens sampled from different areas of the joint, it can be seen that the tensile strength of the specimens ranged from 920 to 990 MPa, and the tensile strength of different specimens did not very much, and did not show an obvious pattern of change.
The post-break length L and post-break cross-sectional area A were measured for the failed specimens, and the results are shown in Fig. 13 (b). The measurement results showed that the specimen #1 had the highest shrinkage at break, reaching 20%, while the specimen #5, which was sampled from the upper part of the joint, had a shrinkage at break of only 8.96% as the sampling position increased. Similarly, the elongation after break showed the same trend, with specimens #1 and #2, which were sampled from the lower part of the joint, showing an elongation after break of about 4%, while the elongation after break of specimen #5 decreased to 2.55%. Tensile specimens can be divided into three parts: base material, heat-affected zone and weld zone, and all specimens have the same mechanical properties base material, so different specimens in the tensile test show plasticity variability associated with the weld, heat-affected zone.
The fracture morphology observation of the tensile specimen was carried out, as shown in Fig. 14. From the figure, the fractures are fibrous and show a grayish color. The fracture areas are all equiaxed tough nests, which can be judged that the fracture modes of the tensile specimens all belong to the microporous aggregation type of toughness fracture. In addition, the composition scan of the local area of the fracture shows that the mass fraction of V element in the fracture section has increased, while the content of Al element has decreased.