*3.5.2. Influence of PWHT on fatigue*

The influence of PWHT on the fatigue performance of the laser beam welded Ti‐6Al‐4V butt joints is shown in **Figure 16**. It should be kept in mind that all data given in **Figure 16** represent the flush milled condition. If the annealing temperature is lower than 750°C, the results of the tests lie in the scattering range of the as‐welded S‐N curve. No significant influence of low‐ temperature annealing (*T* < 750°C) on fatigue performance was found. Conditions SRA1 and SRA2 had approximately the same fatigue limit of approximately 500 MPa, considering the inherent scatter of fatigue experiments.

Annealing at temperatures above 750°C leads to the slight improvement of fatigue properties. The specimens annealed at 850°C (FA2) and subjected to DA showed the highest value of fatigue limit in the present work: 550 MPa. For comparison, as‐welded specimens heat treated at lower temperatures endured less than 500,000 cycles at that level of stress. The fatigue limit for the FA1 condition was not achieved in the current study. The last two FA1 specimens exhibited very low fatigue life compared with the position of other points for this condition. The examination of fracture surfaces revealed the presence of relatively large clusters of pores with the size of approximately 300 µm. These specimens were probably extracted from the region of the plate with worse quality of the welding (run‐in or run‐outs). Even considering the scatter of the results, the general trend, that annealing at high temperatures (>750°C) increases the fatigue strength of the joint and shifts the S‐N curve towards higher values of stresses, can be clearly seen in **Figure 16**.

#### **3.6. Microfractography**

Fracture surfaces were studied using OM and SEM to identify the locations of crack initiation, region of stable crack propagation and overload region morphology. After careful fracture surface examination of the broken S‐N specimens with an optical microscope, it was concluded that almost 100% of these failures started from internal welding defects, that is, pores with average diameters of approximately 10–100 µm. These pores play the role of structural discontinuities and stress concentrators (notches). The existence of such high stresses in the specimen leads to the initiation of the microcracks in very early periods of fatigue life [37]. **Figure 17** shows the fracture surface overall views of the heat‐untreated machined specimen (a) and the specimen after PWHT at 850°C for 1 h (b). Crack initiation sites are clearly seen. The typical distance from the surface to the crack nucleation point is approximately 300–700 µm; however, less frequently, pores appear closer to the surface. Cracks initiated from a single pore are very rare and occur only if the size of the pore is approximately 100 µm. More frequently, the cracks start from the subsurface clusters of pores (usually two or three pores), as shown in **Figure 17(a)**. The distance between cracks in subsurface clusters is typically less than the pore diameter and does not exceed 10 µm. According to AWS D17.1 [19], two or more discontinuities in the welding zone should be treated as one when the spacing between them is less than the dimension of the larger discontinuity. This allows us to consider subsurface clusters of pores as single defects with an approximate size of 150–250 µm.

The area around the crack initiation site is slightly brighter than the region of stable crack growth. This white circle resembles a "fish eye" fracture, common for steels in the ultra‐long‐ life regime, when fatigue fracture origins are mostly at non‐metallic inclusions in the subsur‐ face [37]. The "fish eye" in the present study always had a radius equal to the distance from the surface to the crack nucleation site. The difference in the colours is probably caused by cyclic contact of the fracture surfaces in the absence of atmospheric effects within the "fish eye" and by cyclic contact of the fracture surfaces in the presence of atmospheric gases outside the "fish eye" [37].

*3.5.2. Influence of PWHT on fatigue*

132 Study of Grain Boundary Character

inherent scatter of fatigue experiments.

stresses, can be clearly seen in **Figure 16**.

**3.6. Microfractography**

The influence of PWHT on the fatigue performance of the laser beam welded Ti‐6Al‐4V butt joints is shown in **Figure 16**. It should be kept in mind that all data given in **Figure 16** represent the flush milled condition. If the annealing temperature is lower than 750°C, the results of the tests lie in the scattering range of the as‐welded S‐N curve. No significant influence of low‐ temperature annealing (*T* < 750°C) on fatigue performance was found. Conditions SRA1 and SRA2 had approximately the same fatigue limit of approximately 500 MPa, considering the

Annealing at temperatures above 750°C leads to the slight improvement of fatigue properties. The specimens annealed at 850°C (FA2) and subjected to DA showed the highest value of fatigue limit in the present work: 550 MPa. For comparison, as‐welded specimens heat treated at lower temperatures endured less than 500,000 cycles at that level of stress. The fatigue limit for the FA1 condition was not achieved in the current study. The last two FA1 specimens exhibited very low fatigue life compared with the position of other points for this condition. The examination of fracture surfaces revealed the presence of relatively large clusters of pores with the size of approximately 300 µm. These specimens were probably extracted from the region of the plate with worse quality of the welding (run‐in or run‐outs). Even considering the scatter of the results, the general trend, that annealing at high temperatures (>750°C) increases the fatigue strength of the joint and shifts the S‐N curve towards higher values of

Fracture surfaces were studied using OM and SEM to identify the locations of crack initiation, region of stable crack propagation and overload region morphology. After careful fracture surface examination of the broken S‐N specimens with an optical microscope, it was concluded that almost 100% of these failures started from internal welding defects, that is, pores with average diameters of approximately 10–100 µm. These pores play the role of structural discontinuities and stress concentrators (notches). The existence of such high stresses in the specimen leads to the initiation of the microcracks in very early periods of fatigue life [37]. **Figure 17** shows the fracture surface overall views of the heat‐untreated machined specimen (a) and the specimen after PWHT at 850°C for 1 h (b). Crack initiation sites are clearly seen. The typical distance from the surface to the crack nucleation point is approximately 300–700 µm; however, less frequently, pores appear closer to the surface. Cracks initiated from a single pore are very rare and occur only if the size of the pore is approximately 100 µm. More frequently, the cracks start from the subsurface clusters of pores (usually two or three pores), as shown in **Figure 17(a)**. The distance between cracks in subsurface clusters is typically less than the pore diameter and does not exceed 10 µm. According to AWS D17.1 [19], two or more discontinuities in the welding zone should be treated as one when the spacing between them is less than the dimension of the larger discontinuity. This allows us to consider subsurface

clusters of pores as single defects with an approximate size of 150–250 µm.

The area around the crack initiation site is slightly brighter than the region of stable crack growth. This white circle resembles a "fish eye" fracture, common for steels in the ultra‐long‐

**Figure 17.** OM images of fracture surfaces. (a) As‐welded condition, 525 MPa, 3,488,300 cycles and (b) PWHT (FA2), 650 MPa, 1,763,300 cycles.

The SEM images with higher magnification of the specimen shown in **Figure 17(b)** revealing the topography of the fracture surface in different zones of crack growth are shown in **Figure 18**. The zone adjacent to the pore (**Figure 18(a)**) is characterized by low values of stress intensity factor and shows fibrous morphology. A comparison of the fracture face with the microstructure (see **Figure 8**) suggests that the elongated fracture features correspond to individual α laths. The crack propagated radially from the pore and was dominated by a transgranular mode of cracking. At high stress intensity factors (**Figure 18(b)**), the fracture topography was mainly characterized by typical fatigue striations and secondary cracks. The overload region exhibited small, shallow dimples, which are indicative of ductile fracture due to microvoid coalescence (**Figure 18(c)**).

**Figure 18.** SEM images of fatigue fracture surfaces. PWHT (FA2), 650 MPa, 1,763,300 cycles. (a) Region close to the pore, (b) region of stable crack growth, and (c) overload region, final fracture.
