3.1 Overall morphology of joints
Figure 3(a) shows the overall morphologies of as-welded joints under different forging pressures. It can be seen that the joints are well-formed without cracks and pores, illustrating the welding parameters selected in this research are reasonable. However, when the forging pressure is 29 MPa, the flash near the 42CrMo side is not completely extruded. It reveals that the size of the flash will decrease with decreasing forging pressure. As depicted in Fig. 3(b), when the forging pressure is decreased from 40 to 29 MPa, the width of the weld zone (WZ) is increased from 6.7mm to 8.2mm on the contrary. The friction heat is mainly affected by the friction pressure at the constant rotational speed. Under friction pressure, the high temperature will soften the material generated on the contact surface. Then the forging pressures result in more soft materials being extruded out to form the flash, and the less soft materials form the WZ. Therefore, under the higher forging pressure, the as-welded joint has a narrower WZ.
3.2 Effect of tempering on microstructure
3.2.1 Microstructure of the 42CrMo carbon steel side
Figure 4 reveals the microstructure characteristics of the 42CrMo carbon steel side joint with and without tempered under 35 MPa forging pressure and 60MPa friction pressure. Figures 4a-c show the microstructure of the 42CrMo steel along the cross-section without tempered, and tempered microstructures are shown in Figs. 4e-f. Martensite was deformed in WZ and carbides are decorating grain boundaries and inter-block boundaries, as shown in Fig. 4a. After tempering, the microstructure exhibits a dense lamellar structure which is tempering sorbite (Fig. 4d). The thermal-mechanically affected zone (TMAZ) is characterized by highly deformed grains, which are coarser than those of WZ but finer than those of BM due to the effect of and friction heating and stress[25]. As shown in Fig. 4b, tempering martensite pearlite and carbides do not appear in the TMAZ during the cool process. Heavy plastic deformation occurred in the TMAZ, resulting in the streamlined microstructure. On the contrary, due to the tempering, the microstructure of TMAZ (Fig. 4e) and HAZ (Fig. 4f) is tempered sorbite and decorated with some ferrites. Considering the relatively low temperature and fast heat dissipation from the WZ to HAZ, a higher volume of ferrite was observed in the HAZ than that in the WZ and TMAZ. Hence, regardless of tempering or not, there are obvious differences in the microstructure of TMAZ and HAZ.
From Figs. 4e-f, it can be seen that TMAZ has a large torsional deformation, but the temperature is slightly lower than that of the WZ. The lack of driving force leads to insufficient dynamic recrystallization, and the deformation streamline is retained. This zone shows bending deformation grains and some fine recrystallized grains. The temperature in HAZ drops below the dynamic recrystallization temperature, which is not enough to drive the metal to deform plastically[26]. Therefore, the microstructure in this zone only changes slightly under the action of high temperature and continues to maintain the BM morphology.
3.2.2 Microstructure of the 36Mn2V steel side
The microstructure characteristics along the radial direction of the 36Mn2V steel side of the joint are shown in Figs. 5a-f. Along the friction interface, the microstructure without tempered is martensite and carbides (Fig. 5a). In TMAZ, it can be found that the microstructure mainly contains martensite and some ferrite. As for HAZ, it consists of coarse base material and precipitated carbides, as shown in Fig. 5c. After tempering, the WZ is transformed into acicular ferrite, which increases the concentration of solid dissolved alloy elements, leading to reducing the segregation, refining uniform grains and increasing the fracture resistance of the weld.
3.3 Effect of forging pressure on microstructure of joints
This section reveals the effect of forging pressure on the microstructure of the joint. the microstructure of the 42CrMo side from WZ to HAZ after tempering under forging pressures of 40MPa and 29Mpa was selected as the analysis object, as shown in Fig. 6. It can be seen that compared with 29MPa joint (Fig. 6b), the microstructure of WZ is smaller under the forging pressure of 40MPa (Fig. 6a), which is also greatly deformed caused by the larger deformation force. When the forging pressure increases, the microstructure does not change significantly but the grain refinement is obvious. Comparing with Figs. 6b and e, the microstructure undergoes greater torsional deformation and sufficient dynamic recrystallization, resulting in more deformed microstructure and finer grains. In the HAZ, (Figs. 6c and f), the grain refinement effect is not significant with increase of forging pressure. This is attributed to that the microstructure in this zone is minimally affected and frictional heat with a slight effect of stress, undergoing only slight changes compared to base metals.
3.4 Effect of welding parameters on mechanical properties
3.4.1 Tensile tests
Figure 7 shows the tensile test results of joints with different welding parameters after heat treatment. Regardless of the process parameters, heat treatment or not, both the elastic deformation stage and the yield stage can be seen in the stress-strain curve. There are two types of stress-strain curves. One type is in which the stress reaches its maximum at relatively small strains, while another type is in which the stress reaches its maximum just before the specimen fracture. It can be seen that the first type happens when the joints are tempered (Fig. 7a) and Fig. 7b shows the second type. Different types of curves can reflect the fracture characteristics of different materials. The first type of curve is often associated with the fracture of ductile materials, while the other type tends to represent brittle materials, where the maximum stress value typically occurs just before the material fractures.
Table 3 shows the yield strength and ultimate tensile of the CDFW joints with various forging pressures and after heat treatment. Table 3 demonstrates that the yield strength increases initially and then decreases with the increase of forging pressure. The joint tempered under 35MPa forging pressure has the highest yield strength, reaching up to 798 MPa, which is 7.7% higher than that of joints with 40MPa and 29MPa forging pressures. The yield strength of the joints increases generally with the tempering. When the forging pressure increases from 29MPa to 40MPa, the yield strength of the tempered joint improves by 10.3%, 12.7%, and 5.3% respectively compared to the untempered joint.
The results of tensile tests in Fig. 8 show the elongation and contraction of samples welded at different forging pressures and heat treatment. No matter the tempering or not and the different forging pressure, the fracture occurred at the 36Mn2V base metal. Comparing with Fig. 8a and Fig. 8b, it can be observed that under different forging pressures, the maximum elongation before tempering is 14.2% and the maximum reduction of area is 64.5%. After tempering, the maximum elongation increases to 18.11%, representing an improvement of 29.6%, and the maximum reduction of area increases to 66.3%, an improvement of 2.8%. The elongation and contraction of samples exhibit the same trend as yield strength and tensile strength, increasing initially and then decreasing as the forging pressure increases. The reason for this phenomenon is that on the one hand, under the effect of friction pressure, the material comes into contact and generates heat through rotational motion. As the interface reaches a certain temperature, the metals gradually are softened, and the viscoplastic metals gradually wrap around the entire friction interface. The dominant heat generation mechanism transforms from initial sliding friction to plastic deformation within the viscoplastic metals[16, 27]. Under the effects of heat and force, dynamic recrystallization occurs in viscoplastic metals, leading to grain refinement and joint performance improvement [28]. On the other hand, when the forging pressure is relatively high, more softened metals are extruded to form a flash, reducing the amount of viscoplastic metals available for metallurgical bonding, and resulting in the decrease of joint mechanical properties. Conversely, when the forging pressure is relatively low, the softened metals and oxides, inclusions, etc. cannot be completely extruded from the interface, resulting in the decrease of joint mechanical properties. Therefore, when the forging pressure is at 35 MPa, it can both satisfy the requirements for metallurgical bonding and extrusion of the interfacial impurities, resulting in a joint with good plasticity and strength. Table 3 and Fig. 8 contain the mechanical properties test results of the joints after tempering. The main reason for the increase in yield strength, elongation, and ductility of the joint after tempering is the precipitation of carbides, which can pin dislocations and provide precipitation strengthening[29, 30]. Ferrite has a low solubility for carbon, and the difference in diameter between solute atoms and solvent atoms forms a lattice distortion stress field around the solute atoms[31]. This stress field interacts with dislocations, thereby increasing the yield strength.
Table 3
Tensile test results for joints welded under different forging pressures
No.1
|
Rotational speed
/RPM
|
Friction pressure
/MPa
|
Forging pressure
/MPa
|
Yield strength
/MPa
|
Ultimate tensile strength
/MPa
|
Heat treatment
|
1
|
1451
|
60
|
40
|
741
|
875
|
tempered
|
704
|
845
|
untempered
|
2
|
35
|
798
|
860
|
tempered
|
708
|
815
|
untempered
|
3
|
29
|
750
|
850
|
tempered
|
680
|
760
|
untempered
|
3.4.2 Impact tests
The location of Charpy V-notch was placed at the central of the weld and perpendicular to welding interface. The results of impact test joints are presented in Fig. 9. The highest impact toughness was observed at a forging pressure of 35MPa with 71.02J, while the lowest impact toughness occurred for the 40MPa with 46.41J. For untempered samples, the impact toughness of the numerous lath martensite generated at the welding interface was inferior to that of the pearlite, bainite, tempered sorbite, and ferrite present in the samples prepared at tempered samples. In addition, the impact toughness values of the joints tempered after CDFW under all forging pressures were higher than those of the untempered joints. After tempering, the untempered martensite structure begins to transform into tempered sorbite, and the joint gradually becomes softened. Compared with the untempered joint, its plasticity is improved and its ability to resist crack expansion becomes stronger, contributing to the increase of plastic impact energy absorption.
3.4.3 Microhardness
The microhardness distributions of the dissimilar joints along the WZ welded under different forging pressures and heat treatment are shown in Fig. 10. The average microhardness values of 42CrMo alloy steel and 36Mn2V steel are 310HV and 250HV, respectively. From the welding interface towards the BM, the microhardness is decreased over the TMAZ and dropped further in the HAZ. Due to the formation of the martensitic phase, the maximum hardness of 661HV appears near the welding interface of the untempered joint of 42CrMo steel. The microhardness at the welding interface is higher than that of the TMAZ and HAZ on the 36Mn2V steel side, which is attributed to the refined microstructure and carbides. Comparing with the as-welded joints with different forging pressures before and after tempering, the hardness of the untempered joints is higher than that of tempered joints. This is because after tempered, the original dislocations begin to slip, merge, and disappear. The untempered martensite structure gradually decomposes, the solubility of the alloy elements in the ferrite matrix decreases, and carbides gradually precipitate to form tempered sorbite[32]. The disappearance of dislocations reduces the "pinning" effect on the grain boundaries[30]. Therefore, the microhardness in the WZ, TMAZ, and HAZ decreases after tempering. Comparing with the as-welded joints with different forging pressures before and after tempering, the hardness of the untempered joints is higher than that of tempered joints. This is because after tempered, the original dislocations begin to slip, merge, and disappear. The untempered martensite structure gradually decomposes, the solubility of the alloy elements in the ferrite matrix decreases, and carbides gradually precipitate to form tempered sorbite[32]. The disappearance of dislocations reduces the "pinning" effect on the grain boundaries[30]. Therefore, the microhardness in the WZ, TMAZ, and HAZ decreases after tempering.
After tempering, the hardness of WZ increases with the increase of forging pressure. Increasing the forging pressure usually enhances its microhardness. Higher forging pressure results in greater plastic deformation in WZ, promoting grain refinement and achieving a grain refinement strengthening effect. Additionally, larger forging pressure can also increase the dislocation density in WZ, which generally enhances its hardness. However, in this study, the hardness of the untempered WZ shows the opposite trend compared to the tempered joint. The main reason is as follows: under lower forging pressure, more untempered martensite structure is retained in the WZ. During microhardness testing, the greater presence of martensite structure exhibits higher hardness. A softened zone is only observed on the 36Mn2V steel side due to the ferrite content in the HAZ. Comparing with the hardness values of three HAZs with different forging pressures in Fig. 10, the lowest hardness of 260HV was observed for the 40Mpa joint on the 36Mn2V carbon steel side, which has tempered after welding. In the TMAZ, the highest hardness was observed for the 29MPa joint on the 42CrMo side due to faster heat dissipation.
3.4.4 Fracture analysis
Figure 11 presents the macro-microscopic morphologies of the specimens after the tensile test for as-welded joints with 35MPa forging pressure. Figures 11a-c show the fracture morphologies of the tempered specimens after tensile testing while Figs. 11e-f show the fracture morphologies of the untempered specimens. The different zones in the radial direction defined in (a) are shown in Figs. 11b-c. The central area was considered as a rough fracture surface, which indicates improved metallurgical bonding, as shown in Fig. 11a and d. In Fig. 11a, a large number of dimples and tear ridges can be observed, indicating that the specimen exhibits good ductility. The number of dimples in Fig. 11a is significantly higher than in Fig. 11e, and the dimples in Fig. 11a are more uniformly distributed and more regularly shaped. These uniformly distributed dimples suggest that the tempered material possesses good toughness, which is consistent with the impact toughness results discussed in section 3.4.2. Under tensile stress, micropores in the material form and coalesce, ultimately leading to the formation of dimples. These micropores can form around grain boundaries, inclusions, or second-phase particles. Therefore, examining the detailed SEM micrographs in Figs. 11a and e, spherical second-phase particles can be seen around the dimples. The distribution of these second-phase particles affects the formation of dimples. After tempering, there are fewer second-phase particles in the structure and no aggregation occurs (Fig. 11b). Conversely, a higher number of second-phase particles tend to aggregate. These aggregated second-phase particles create defect concentration zones in the microstructure, leading to a decrease in mechanical properties.
Figures 11c and f show the tensile fracture peripheral area. Few dimples can be observed, and their parabolic shape indicates that the microstructure in this area was subjected to shear stress along the tensile direction. This area also exhibits heterogeneous microstructures, cleavage steps, and a significant number of second-phase particles, indicating lower tensile strength in the peripheral area. Consequently, during fracture, cracks initially formed in the peripheral area, accompanied by cleavage fracture. As the tensile load increased, cracks also appeared in the central area. Eventually, the cracks intersected, leading to the fracture of the specimen.
Figure 12 shows the macro and microscopic morphologies of the tensile fracture surfaces after tempering at forging pressures of 40MPa and 29MPa. Figures 12a-c show the fracture morphologies of the central and peripheral area at a forging pressure of 40MPa, while Figs. 12e-f show the fracture morphologies with forging pressure of 29MPa. Compared to Fig. 11b, the dimples in the central area at 40MPa (Fig. 12b) and 29MPa (Fig. 12e) are larger and fewer, with more regular shapes. The peripheral area (Fig. 12c and f) exhibits more pronounced cleavage fracture characteristics compared to 40 MPa (Fig. 11c), with clear cleavage steps and river-like patterns.
The effect of the forging pressure is explained as follows. As the forging pressure increases, the plastic deformation between metals is enhanced, improving the interfacial bonding strength. This results in the formation of larger but fewer dimples during fracture. The increase in forging pressure also makes the peripheral area more prone to displaying cleavage fracture characteristics, such as the cleavage steps and river-like patterns observed in Fig. 12(c). This indicates that at higher forging pressures, the material in the peripheral area becomes more brittle, making cracks more likely to initiate and propagate along the grain boundaries in these areas, ultimately leading to fracture.