*3.2.1. Base material*

**Figure 2.** Transverse cross section macrograph of the laser beam welded Ti‐6Al‐4V butt joint.

118 Study of Grain Boundary Character

terms of the weld profile imperfections for the aerospace industry.

The most frequently observed defects in laser weldments of titanium alloys are underfills and porosity [3, 5, 16, 17]. These imperfections are particularly undesirable for structures subjected to cyclic loading because they lead to stress concentration and consequently premature crack formation. The evaporation and expulsion of the molten material from the weld pool and liquid metal flow around the keyhole are dominant processes affecting the formation of underfills [2, 18], which are always present in cases of autogenous laser welding. The use of an additional filler wire in the present work allowed overfilling to be produced and geometrical weld imperfections such as underfills to be partially eliminated. However, filler wire resulted in weld reinforcements from both face and root sides. Abrupt change in the thickness due to weld reinforcement leads to stress concentration at the weld toes and roots and consequently reduces the fatigue strength of the joints. Although the macrograph presented in **Figure 2** reveals almost no underfills, a single weld cross section cannot guarantee the uniformity of the weld profile over the whole length of the seam. This problem can usually be solved by extracting more than one specimen for metallurgical examination. In the present work, the maximum measured underfill depth was approximately 70 µm, which is less than 3% of the specimen thickness. The maximum observed reinforcement was approximately 350 µm. The geometric profile imperfections of laser weldments used in the aerospace industry are strictly limited by several standards: AWS D17.1 [19] and EN 4678 [20]. In terms of underfills and weld reinforcements, EN 4678 is more stringent; the maximum allowed underfill depth for butt joints is 5% of the total thickness and maximum reinforcement is 490 µm for the specimens of the 2.6‐mm‐thick material. All welds in our work confidently passed the acceptance criteria in

The spherical shape of most pores observed in the present study indicates gas‐type porosity. A number of researchers have investigated the main causes of porosity when laser welding titanium alloys [18, 21, 22]. Potential sources for porosity formation are mainly from the presence of excessive hydrogen in the FZ, which is rejected upon solidification, and keyhole instability leading to the entrapment of shielding gases. The investigation of the influence of The as‐received BM microstructure of the 2.6‐mm‐thick Ti‐6Al‐4V sheet consists of globular α grains with average grain size of 3.1 ± 0.8 µm and an intergranular retained β, as shown in **Figure 3**. This corresponds to what is known as a mill‐annealed microstructure for a hot worked plate that is not fully recrystallized [1]. **Figure 3(a)** shows incompletely recrystallized regions with lamellar morphology, which are not fully transformed to equiaxed α grains upon mill annealing after a hot rolling process. The bright regions in **Figure 3(a)** are the equiaxed or lamellar α grains, and the dark regions are the intergranular β grains distributed at α grain boundaries. Texture analysis is represented by pole figures and inverse pole figures in the cross‐sectional plane. It should be kept in mind that the normal to the cross section coincides with the transverse direction of the sheet, and the S direction stands for the thickness direction (see **Figure 2**). The colour in the crystal orientation map (**Figure 3(b)**) is based on a colour‐ coded inverse pole figure, in which different colours represent different crystallographic orientations. Large red regions in the orientation map correspond to not fully recrystallized lamellar regions. They have the same colour owing to their near‐equal crystallographic orientation and were not considered for the calculation of the average grain size.

**Figure 3.** Microstructure of the BM in the as‐received condition. (a) OM image, (b) orientation map in the cross section, (c) inverse pole figure, and (d) (0 0 0 1) and (1 1 –2 0) pole figures in the cross‐sectional plane.

The base material is characterized by preferred crystal directions such as <0 0 0 1>//ND and <1 0 ‐1 0>//ND, as shown by the inverse pole figure in **Figure 3(c)**. The microtexture components have been determined using the orientation distribution function (ODF) at sections of *φ*2 = 0° and *φ*2 = 30°, where *φ*<sup>2</sup> is an Euler angle. The results analysis showed that the Ti‐6Al‐4V base material contained the components (0 0 0 1)[2 ‐1 ‐1 0] (*f* = 12.7 mrd (multiple of a random distribution), (0 1 ‐1 0)[2 ‐1 ‐1 4] (4.3 mrd < *f* < 9.4 mrd) and (0 1 ‐1 0)[2 ‐1 ‐1 0] (*f* = 1.6 mrd). The (0 0 0 1) pole figure shows that basal planes are aligned in both the rolling direction (RD) and transverse direction (TD) (see **Figure 3(d)**). According to the investigation of Salem [23], this type of texture indicates that the Ti‐6Al‐4V sheet was cross‐rolled. The pole density of the (0 0 0 1) pole figure at RD shows an opening angle of 90° and corresponds to orientation bands at (*φ*2 = 0°, *φ*1 = 60° and *φ*1 = 120°, 0°≤ *φ* ≤ 180°) and (*φ*2 = 30°, *φ*1 = 90°, 0° ≤ *φ* ≤ 180°).

#### *3.2.2. Fusion zone*

The microstructure of the FZ is characterized by columnar prior β grains that grow from the HAZ in the direction opposite that of the heat flow and impinge at the weld centreline after solidification (see **Figures 2** and **4**). The FZ prior β grain size depends primarily on the weld energy input, with a higher energy input promoting a larger grain size [2, 18, 24]. In the present study, the average prior β grain size in the as‐welded condition was approximately 200–300 µm. Within the prior β grains, the FZ predominantly consists of an acicular α' martensitic structure, resulting from the diffusionless β → α' transformation upon high cooling rates encountered in the LBW process. Ahmed et al. [25] investigated the effect of different cooling rates on microstructural reactions in Ti‐6Al‐4V and found out that fully α' martensitic transformations take place at cooling rates above 410°C/s. The martensitic microstructure was characterized by long orthogonally oriented thin plates having acicular morphology. **Figure 4(a)** and **(b)** show a similar microstructure in the FZ of laser beam welded Ti‐6Al‐4V butt joints. Ahmed et al. observed preferential grain boundary formation of secondary α morphology at cooling rates in the range between 410 and 20°C/s. Because no secondary α at prior β grain boundaries was observed in the current work, according to the results of Ahmed et al., we can conclude that the cooling rate in the welding zone was high enough to provide β → α' diffusionless transformation during LBW. The observed micro‐ structure is typical for fusion zones of laser beam welded [2, 5] and electron beam welded [7, 8] joints.

**Figure 4.** Microstructure of the FZ. (a) OM image of the microstructure, (b) orientation map in the cross section, (c) inverse pole figure, and (d) (0 0 0 1) and (1 1 –2 0) pole figures in the cross section plane.

Local melting and subsequent solidification of Ti‐6Al‐4V has changed the initial microtexture significantly. The main microtexture components in the FZ are (0 6 ‐6 1)[8 ‐3 ‐5 12] (*f* = 3.7 mrd), (0 6 ‐6 1)[‐14 5 9 24] (*f* = 2.2 mrd), (0 1 ‐1 1)[‐4 ‐1 5 6] (*f* = 3.3 mrd), (0 2 2 ‐3)[2 ‐1 ‐1 0] (*f* = 3.2 mrd) and (0 1 ‐1 1)[2 ‐1 ‐1 0] (*f* = 2.8 mrd), which were examined at 100× magnification to obtain higher statistical weight. The (0 0 0 1)[2 ‐1 ‐1 0] component (*f* = 2.5 mrd) is compa‐ ratively much less pronounced (**Figure 4(c)** and **(d)**) than it was in the BM (**Figure 3(d)**). The preferred direction of crystal growth was the <1 1 ‐2 0>//RD crystal direction during solidi‐ fication. Furthermore, the presence of <1 1 ‐2 0>//RD fibre texture is visible.
