3.1 Effect of post-weld heat treatment on weld microstructure
The microstructures of BM and fusion zone (FZ) of weld are shown in Fig. 2. It can be seen that the BM mainly consists of white martensite with black carbides in Fig. 2(a), and grains are distributed in a band shape along the rolling direction as shown in Fig. 2(b). In Fig. 2(c), the prior austenite grains in FZ grow in a manner of columnar, and the martensitic laths present a parallel alignment in the grain interior. The δ-ferrite is distributed at the prior austenite grain boundaries (PAGB) and present two forms, the worm-shaped ferrite parallel to the growth direction of prior austenite and the non-parallel lathy ferrite. The mean content of δ-ferrite in FZ was evaluated by image processing software to be about 6 vol.%. It has been reported that the 5-15 vol.% δ-ferrite content in stainless steel weld can significantly suppress the solidification cracking of welding [1, 10]. Fig. 2(d) shows the local morphology at high magnification of FZ. It can be seen that a small quantity of nano-sized carbide particles is distributed in the martensite matrix. The X-ray diffraction patterns for BM and FZ of weld are shown in Fig. 3. It can be found that there is a small amount of retained austenite in BM. The absence of austenite diffraction peaks in FZ confirms that the retained austenite has completely transformed into martensite after welding. The volume fraction of austenite at room temperature was calculated according to the following method by measuring the integrated intensities of (110)α’ and (111)γ diffraction peaks [11]
![](https://myfiles.space/user_files/58890_add8f4303ffe25fa/58890_custom_files/img1599556242.png)
Where Vα’ and Vγ are the volume fraction of γ austenite and α martensite, respectively. Iα’ and Iγ are the integrated intensities of (110)α’ and (111)γ diffraction peaks, respectively. Using the above equations, the volume fraction of retained austenite of BM is determined as 8.4 vol.%.
Fig. 4 gives the SEM images of microstructural evolution in FZ under all the heat treated conditions. With the increase of aging time, the amount of carbides precipitated along PAGB and martensitic lath boundaries increase, and δ-ferrite is still visible at aging temperatures of 300 ºC and 400 ºC. At the aging temperature of 550 ºC, the δ-ferrite gradually spheroidizes, and eventually disappears as the aging time increases. Simultaneously, the characteristic of columnar grains in FZ also disappears. When aging at 600 ºC and 650 ºC, the processes of spheroidization and disappearance of δ-ferrite accelerated. Furthermore, at the higher temperature of 900 ºC (Ac3 temperature of FV520B steel is about 900 ºC [12]) austenitization occurs. With the increase of holding time, the martensitic laths become coarse after cooling.
Fig. 5 shows the SEM images of FZ under some particular heat-treatment parameters at the higher magnification of 15000×. As can be seen in Fig. 5(a), at 400 ºC for 9 min many fine granular-shaped and needle-shaped carbides appear within the size range from tens to hundreds of nanometers. According to the study of Lu et al. [13], the precipitates with the same morphology were also found in the Cr13 martensitic stainless steels tempered at 300 ºC, and this carbide was identified as M3C type (M is Fe or Cr) by transmission electron microscopy (TEM). As show in Fig. 5(b), the size of carbides increases, while the tempered martensite still retains the lathy feature. During laser welding the high cooling rate of the weld causes that there is no sufficient time for the precipitation of alloy elements, thus leading to the formation of the unstable supersaturated solid solution. The alloy elements in supersaturated solid solution tend to precipitate in the form of carbides, and the size of carbides also increases with aging time. The coarsening of prior austenite grain boundaries (PAGBs) may be related to the precipitation of carbides at grain boundaries. The precipitation of carbides results in the continuous rejection of carbon atoms from the martensite which causes a decrease of the tetragonality of the martensite matrix [14], thereby forming tempered martensite. Fig. 5(c) and 5(d) show the microstructures aged at 550 ºC for 2 h and 12 h, respectively. In Fig. 5(c) the lathy feature of the tempered martensite can still be observed, while after aging for 12 h the nucleation of new phases occurs at the PAGBs and these phases grow inside the prior austenite grains along the boundaries of martensite laths. Such phases can also be seen in the weld aged at 600 ºC for 1 h (Fig. 5(e)). It was reported that the reversion from martensite to austenite will occur when maraging steel is heated close to the Acl temperature [8, 15] (Ac1 temperature of FV520B is about 580 ºC [16]). During cooling, the austenite is expected to two types of transformation. The first is the forming of martensite, and the other is the forming of the more stable reversed austenite due to the high content of austenite stabilizing elements (Ni, Mn, etc.). Thus, it can be speculated that the relevant phases in Figs. 5(d) and 5(e) may be the reversed austenite which formed by reversed transformation and retain their morphology at the aging temperature after cooling. Its composition may be martensite, reversed austenite, or a mixed structure of martensite and reversed austenite. After aging at 600 ºC for 12 h, the strip-shaped and granular phases are distributed in the tempered martensite matrix. It is supposed that the phases in Fig. 5(f) are formed by the growth of the above-mentioned phases. As shown in Fig. 5(g) and 5(h), at aging temperature of 650 ºC the strip-shaped and granular grains are formed after only 1 h and obviously coarsened after 12 h.
EDS results of the positions denoted by the arrows in Figs. 5(e) and 5(f) are shown in Table 2. In Fig. 5(e), the Cu content at the position 1 is higher than the average content of the base metal. It indicates that the segregation of Cu has occurred here. It has been reported that Cu bearing martensitic precipitation hardened stainless steel will precipitate nano-scale Cu-rich particles during aging treatment [6, 17]. The Ni content at positions 2 and 3 is higher than the average content of the base metal. It implies that the segregation of Ni has occurred at these positions. The reversed austenite may be formed in Ni-rich regions after aging. However, it is difficult to use SEM to observe the morphology of reversed austenite in martensitic matrix due to its small size.
Fig. 6 shows the XRD patterns of the FZ under different heat treatment conditions. At room temperature the volume fraction of reversed austenite can be calculated using Eq. (1) and (2), and its variation with aging time is plotted in Fig. 7. At aging temperature of 550 ºC the reversed austenite diffraction peaks cannot be found at 2 h, while after a longer aging time of 12 h, the amount of reversed austenite reaches 7.0 vol.%. When the aging temperature increases to 600 ºC, there is no reversed austenite under the holding times below 1 h. As the aging time increases to 1 h, 2h and 12 h, the volume fraction of reversed austenite reaches 2.9 vol.%, 9.0 vol.% and 15.6 vol.% respectively. When the aging temperature further rises to 650 ºC, the reversed austenite is formed after aging for only 9 min. Additionally, it is noted that the amount of austenite increases first and then decreases with increasing time. In particular, after aging for 12 h the FZ no longer contains the austenite at room temperature. The amount of reversed austenite is related to the content and stability of the austenite formed during aging. Furthermore, the amount of austenite during aging is associated with the temperature and the holding time. The higher temperature and the longer time are desirable for the formation of austenite. The stability of austenite at the aging temperature is related not only to the content of austenite-stabilizing elements such as Ni and Mn in austenite, but also to the concentration of vacancies [8, 14, 15]. When the aging temperature is at 550 ºC, the reversed transformation occurs in a diffusion manner, then it requires a long time to obtain austenite. When heated to 600 ºC, the formation rate of reversed austenite becomes faster, and so the amount of reversed austenite at room temperature also increases. The reason why the reversed austenite disappears at the aging condition of 650 ºC and 12 h can be explained as follows. The limited Ni concentration and the increasing reversed austenite at the aging temperature give rise to a decrease of the average content of Ni in austenite, resulting in a decrease in the stability of austenite and an increase in the Ms point. Therefore, all austenite is transformed into martensite during cooling [18]. Wang et al. [6] found that as aging time was prolonged, the segregation of Ni, Mn, Si and Nb occurred and the co-precipitates were formed accompanied with precipitation of Cu. This is also one reason of a decrease of Ni content in reversed austenite.
3.2 Effect of post-weld heat treatment on Vickers hardness and impact toughness of weld
Fig. 8 shows the Vickers hardness of the weld in as-welded and different heat treated conditions. For the as-welded specimen, the hardness of FZ is higher than that of the heat affected zone (HAZ) but slightly lower than that of the BM. At aging temperatures of 300 ºC and 400 ºC, the average hardness of FZ and HAZ increases with aging time, and the hardness difference between the three regions narrowed. In the aging temperature range of 550-650 ºC, the hardness curves of welds present the maximum values at the aging time of 9 min, and then decrease with increasing aging time.
The increase in hardness during aging is mainly related to the dispersion strengthening caused by the precipitation of Cu-rich phases in a nanoscale [7, 19, 20]. Below the temperature of 400 ºC, the precipitation rate of Cu-rich phase is lower. However, when aging temperature is not lower than 550 ºC, the high precipitation rate makes the number of Cu-rich precipitates increase sharply in a short time, resulting in a high level hardness. However, the Cu-rich precipitates gradually grow with increasing holding time, thus leading to a weakening of dispersion strengthening effect. Additionally, the increase of tempered martensite and the occurrence of reversed austenite with aging time will also give rise to a softening effect which can even overcome the dispersion strengthening, resulting in a lower hardness compared with the as-welded.
At the same parameter three parallel samples were prepared for the impact toughness test. The mean values of the absorbed impact energy of the welded joints under different heat treatment processes are presented in Fig. 9. At aging temperatures of 300 ºC and 400 ºC, the impact toughnesses of welds are lower, and have a similar variation trend which first increase and then decrease with holding time. At the same holding time, the absorbed impact energies of the welds under aging temperature of 400 ºC are higher than that under 300 ºC. At aging temperature of 550 ºC, the impact toughness first decreases and then increases with holding time, the absorbed impact energy reaches the highest value of 61 J at aging time 12 h. For Fe-Cr alloys the Cr-rich ferrites will be formed at the elevated temperatures from 475 ºC to 550 ºC [1], and this results in the decrease of impact toughness after aging treatment. However, at the longer aging time of 12 h the impact toughness of the weld is improved due to the formation of reversed austenite [15]. At aging temperature of 600 ºC, the absorbed impact energy gradually increased with aging time, reaching the maximum of 75 J at aging time 12 h. Under aging temperature of 650 ºC, the absorbed impact energy increases first and then decreases as the aging time increases, reaching the maximum at aging time 2 h. Combined with the XRD results, it can be found that the impact toughness improves as the content of reversed austenite increases. In addition, with the increase of aging temperature and time, the growth of non-coherent Cu-rich precipitates can also lead to a decrease in hardness and an increase in impact toughness [14].
Fig. 10 shows the impact fracture morphologies of the specimens at aging temperatures of 550 ºC and 650 ºC under different aging times. At temperature of 550 ºC and aging time below 2 h, as shown in Fig. 10(a), 10(c) and 10(e), the fractographs exhibit the quasi-cleavage characteristics, which correspond to the lower toughness. At aging temperature 550 ºC for 12 h, as shown in Fig. 10(g), the large number of dimples on the fracture surface imply an improvement in the toughness of weld. The EDS analyses of the short rod-shaped second phase (denoted by the arrow in Fig. 10(e)) shows the element content 8.83Cr-1.48Ni-29.37Mn-18.85S (wt. %). The higher contents of S and Mn indicate that the second phase may be MnS. At aging temperature of 650 ºC, the specimens under four different aging times all show the dimple-like characteristics corresponding to a good toughness. Fig. 10(b), 10(d) and 10(f) exhibit tear-type dimples, while Fig. 10(h) shows fine equiaxed dimples. At aging temperature of 650 ºC for 12 h, the finer dimples, which is likely to be related to the formation of reversed austenite, do not produce an extra effect on the impact toughness. It differs from the significant improvement in toughness at aging temperature 550 ºC for 12 h.
Generally, the Vickers hardness presents a strong positive correlation with yield strength [20], then the value of Vickers hardness is selected to reflect the change trend of yield strength in this work. The scatter diagram of Vickers hardness vs. impact energy under all heat treatment conditions is plotted in Fig. 11, and the point number in the diagram is consistent with the heat treatment parameters in Table 3. In order to compare the variation of properties under different heat treatment parameters with the as welded, the diagram is divided into four parts by the dotted lines across the as-welded point. As shown in Fig. 11, the scatter diagram can give a clear indication for the selection of desired properties of welds. The parameter points in the upper right corner area of the diagram correspond to the excellent combination of strength and toughness, such as the point numbers of 17 and 14.