*3.3.1. Influence of PWHT on microhardness*

obtained platelets at high temperatures (**Figure 8**). Starting from 750°C (FA1), a secondary α phase at prior β grain boundaries designated as grain boundary α [25] starts to appear (**Figure 8(b)** and **(c)**). Furthermore, the thickness of grain boundary α gradually increases with increasing PWHT temperature. After DA, the colonies of parallel α plates are formed upon recrystallization processes. In the case of the martensitic structure, fine α' needles are nearly orthogonal, and α colonies are not so pronounced (see **Figure 4(b)**). As we can see in **Fig‐ ure 8(c)**, in DA condition, the α colony size is clearly visible and is approximately 20–30 µm. These α colonies should not be confused with prior β grains. The prior β grain size was not altered after PWHT; it was an order of magnitude larger than the α colony size and cannot be

Because the HAZ has a bimodal microstructure and consists of equiaxed primary α grains in transformed β matrix, the effect of PWHT on the HAZ microstructure is somehow the combination of the above‐mentioned effects on the BM and FZ, that is, coarsening of primary

The distribution of microhardness across the laser beam welded Ti‐6AL‐4V butt joint in the as‐welded condition is presented in **Figure 9**. No significant difference in microhardness profiles among three testing positions was found; that is, no thickness gradient was observed in the present study. The average microhardness value of the base material was found to be 336 ± 8 HV 0.5. The FZ exhibited the highest average microhardness, approximately 396 ± 10 HV 0.5 (roughly 18% greater than that in the BM), and it decreased abruptly as the distance from the FZ line increased. The microhardness distribution within the FZ was quite uniform without significant deviations from the average value. This result is related to the use of Ti Grade 5 as a filler wire material. In the case of LBW with Ti Grade 2 (commercially pure Ti) filler material, a significant decrease in microhardness in the centre of the FZ was observed [27].

The increase in microhardness from the BM via the heat‐affected zone to the FZ centre is correlated with local changes in microstructure, which were activated during LBW and subsequent cooling. The occurrence of maximum hardness in the FZ is related to the formation of a strong martensitic structure due to high cooling rates upon solidification. Acicular α' phase produced by the diffusionless transformation from the high‐temperature β phase field exhibits

α grains and transformation of fine martensitic structure into coarse lamellar.

seen in **Figure 8**.

124 Study of Grain Boundary Character

**3.3. Microhardness**

**Figure 8.** Influence of PWHT on the microstructure of the FZ.

The variations of average microhardness values for the BM and FZ in the as‐welded condition and after PWHT under different conditions are shown in **Figure 10**. The hardness in the FZ was generally higher than that in the BM, but their difference depended strongly on the annealing temperature. The changes in microhardness after conducting PWHT are strictly related to the microstructural changes that occurred during the PWHT. Because the effect of PWHT depends on the initial microstructure, the BM and FZ underwent different transfor‐ mations during PWHT and will be discussed separately.

Annealing at 540°C for 4 h slightly increased the average microhardness of the BM from 336 ± 8 HV 0.5 in the as‐welded condition to 351 ± 10 HV 0.5 after annealing. This unexpected hardening effect upon low‐temperature annealing was not evident from OM and SEM observations and can be attributed to precipitation hardening of the α phase by coherent Ti3Al particles [1, 28]. During annealing, significant alloy element partitioning takes place; that is, α phase is enriched with α stabilizing elements (Al), and β phase is enriched with β stabilizing elements (V, Fe) owing to diffusion processes. This fact was proved by EDX point analysis. **Figure 11** shows the average content of alloying elements in α and β phases of the BM after PWHT. As shown in **Figure 11**, Al content in the α phase and V and Fe content in the β phase increase with increasing temperature. Coherent α2 particles can then be precipitated in the alpha phase by ageing owing to increased Al content. This age‐hardening effect of the α phase in Ti‐6Al‐4V by Ti3Al particles was well documented by Lutjering et al. [28]. In the Ti‐6Al‐4V alloy, the Ti3Al solvus temperature is approximately 550°C. Annealing at temperatures lower than 550°C will precipitate α2 particles, whereas a heat treatment at 600°C will be only a stress‐ relieving treatment [29]. The latter was in a good agreement with our findings. Heat treatment at 650°C for 1 h led to average microhardness in the BM of 331 ± 10 HV 0.5, which is slightly lower than that in the starting condition. Evidence of this ageing phenomenon is not apparent from optical microscopy images and EBSD results because the size of the α2 particles is approximately several nanometres as reported in Ref. [28]. The application of transmission electron microscopy (TEM) techniques to study this complex phenomenon should improve identification of the α2 phase.

**Figure 10.** PWHT influence on the average microhardness in the FZ and BM.

As presented in **Figure 10**, annealing in the temperature range of 650–850°C does not affect the average microhardness in the BM significantly, and It gradually decreased with increasing temperature. A substantial decrease in microhardness was observed after recrystallization annealing at 920°C for 45 min. These results are related to grain coarsening in accordance with the Hall‐Petch mechanism and are consistent with microstructural observations discussed in the previous section. Because the second step of duplex annealing is equal to SRA1 heat treatment, this led to hardening of the coarsened structure obtained after RA due to precipi‐ tation of α2 phase as discussed above.

**Figure 11.** Alloying element partitioning in the BM due to PWHT. (a) Al in α phase and (b) V and Fe in β phase.

The effect of various PWHT on the average microhardness in the FZ has a generally very similar trend as discussed for the BM; however, the mechanisms leading to these results were not the same because the initial microstructures were completely different. The hardening effect observed in the FZ after SRA1 annealing can be attributed to tempering of the martensitic structure. This effect at relatively low temperatures has already been observed by a number of researchers [6, 30, 31]. Because metastable α' martensite is supersaturated in β stabilizers owing to the diffusionless transformation β → α', upon annealing, it decomposes into α + β by precipitation of incoherent β particles at dislocations or β phase layers at plate boundaries [28]. Chesnutt et al. [32] investigated ageing of β‐quenched Ti‐6Al‐4V and used TEM to show microprecipitation of β phase particles in tempered martensite. Precipitation hardening takes place only after SRA1 annealing, whereas heat treatment at 650°C for 2 h leads to partial decomposition of martensite with attendant reduction of microhardness. Starting from the temperature of 750°C, grain coarsening of an acicular microstructure in the FZ is responsible for the gradual reduction of microhardness with increasing temperature (see **Figure 10**). After recrystallization, the annealing microstructure in the FZ was completely transformed into equilibrium coarse lamellar α + β morphology with the lowest hardness values nearly equal to that of the BM in the as‐received condition. The hardening mechanism that took place after ageing in DA was apparently the same as described above for the BM because after recrystal‐ lization, both the BM and FZ consisted of equilibrium α + β phases.

#### **3.4. Residual stress analysis**

particles [1, 28]. During annealing, significant alloy element partitioning takes place; that is, α phase is enriched with α stabilizing elements (Al), and β phase is enriched with β stabilizing elements (V, Fe) owing to diffusion processes. This fact was proved by EDX point analysis. **Figure 11** shows the average content of alloying elements in α and β phases of the BM after PWHT. As shown in **Figure 11**, Al content in the α phase and V and Fe content in the β phase increase with increasing temperature. Coherent α2 particles can then be precipitated in the alpha phase by ageing owing to increased Al content. This age‐hardening effect of the α phase in Ti‐6Al‐4V by Ti3Al particles was well documented by Lutjering et al. [28]. In the Ti‐6Al‐4V alloy, the Ti3Al solvus temperature is approximately 550°C. Annealing at temperatures lower than 550°C will precipitate α2 particles, whereas a heat treatment at 600°C will be only a stress‐ relieving treatment [29]. The latter was in a good agreement with our findings. Heat treatment at 650°C for 1 h led to average microhardness in the BM of 331 ± 10 HV 0.5, which is slightly lower than that in the starting condition. Evidence of this ageing phenomenon is not apparent from optical microscopy images and EBSD results because the size of the α2 particles is approximately several nanometres as reported in Ref. [28]. The application of transmission electron microscopy (TEM) techniques to study this complex phenomenon should improve

identification of the α2 phase.

126 Study of Grain Boundary Character

**Figure 10.** PWHT influence on the average microhardness in the FZ and BM.

As presented in **Figure 10**, annealing in the temperature range of 650–850°C does not affect the average microhardness in the BM significantly, and It gradually decreased with increasing temperature. A substantial decrease in microhardness was observed after recrystallization annealing at 920°C for 45 min. These results are related to grain coarsening in accordance with the Hall‐Petch mechanism and are consistent with microstructural observations discussed in the previous section. Because the second step of duplex annealing is equal to SRA1 heat

The HDM method gives the average stresses in the area where the material was drilled, that is, in the circle of 0.6 mm in diameter. Because no depth‐dependent gradient of residual stresses was observed in the present study, each point is represented by the mean value of residual stresses at the depth of 0.3 mm. The distribution of residual stresses longitudinal and transverse to the weld line across the welding seam is shown in **Figure 12(a)**. High tensile longitudinal stresses up to 650 MPa are produced in the weld itself and the immediately adjacent parent material during solidification. These high tensile residual stresses near the weld are balanced by compressive longitudinal stresses further from the weld line. Transverse residual stresses have a similar profile but are an order of magnitude lower than the longitudinal component with a maximum value not exceeding 50 MPa. Our results are in good agreement with that reported by Cao et al. for laser beam welded Ti‐6Al‐4V alloy [33].

**Figure 12.** Residual stress distribution in the vicinity of the laser beam welded Ti‐6Al‐4V butt joint. (a) Residual stress‐ es in the as‐welded condition and (b) influence of PWHT on longitudinal residual stress profile.

**Figure 13.** Residual stresses in the BM after milling.

The stress‐relieving effect at elevated temperatures due to PWHT is presented in **Figure 12(b)**. Because longitudinal stresses significantly prevail over transverse components, for simplicity, **Figure 12(b)** presents only the evolution of longitudinal stresses upon heat treatment. Annealing at 540°C results in the stress relieffrom 650 MPa in the centre of FZ to approximately 90 MPa (nearly 85% effect). After heat treatment at 750°C, almost full stress relief was achieved, and PWHT's at higher temperatures completely remove welding‐induced residual stresses in the welding seam. It should be kept in mind that the presented results correspond to different specimen geometry from that used in fatigue testing. The specimen size should be considered as an influential factor when evaluating the effect of residual stresses on the fatigue of welded joints [34, 35]. After extracting S–N specimens from the welded plate due to the relatively short width of the gauge length (8 mm), residual stresses are almost fully removed. The latter was confirmed by the HDM technique in the extracted fatigue specimens and does not allow us to make any conclusions about the influence ofresidual stresses on the fatigue properties of laser beam welded Ti‐6Al‐4V butt joints. However, in the situations in which the width of the tested specimen is sufficient to keep the residual stresses after cutting or the welding line is parallel to the external stresses, the results presented in this work would be quite useful. Moreover, the stress‐relieving effect investigated here is necessary for comprehensive analysis of the PWHT oflaser beam welded butt joints. Large‐scale specimens must be tested to separately investigate the influence of residual stresses on the fatigue of LBW joints.

Incremental HDM allowed us to investigate the residual stresses that arise in the surface layer after machining. These stresses are of great interest because they significantly affect the unnotched fatigue properties of the material and will be discussed in the last section. **Figure 13** shows a residual stress distribution in the BM after milling. As we can see, high compressive residual stresses up to 550 MPa are formed in the 0.2‐mm‐thick surface layer. These compres‐ sive stresses in the surface layer are balanced by tensile stresses in the bulk material.
