*3.2.3. Heat affected zone*

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

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,

**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.

(*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*

120 Study of Grain Boundary Character

8] joints.

HAZ displays the transition region between the acicular morphology in the FZ and globular structure in the BM. It is usually divided into two subregions based on the β transus temper‐ ature. In the HAZ adjacent to the FZ (near‐HAZ), the temperatures exceed the β transus during LBW. Consequently, this region consists mostly of the transformed acicular microstructure. Because the temperatures in the HAZ adjacent to the BM (far‐HAZ) were lower than the β transus, its microstructure is very similar to that of the BM. Microstructure analysis of both the near‐HAZ and far‐HAZ is shown in **Figure 5**.

**Figure 5.** Microstructure of the HAZ in the as‐welded condition. Orientation maps, inverse pole figures and pole fig‐ ures of the HAZ adjacent to the BM (a) and the HAZ adjacent to the FZ (b).

The microstructure of the HAZ zone adjacent to the BM (1.4 mm from the weld centre) remained nearly the same after thermal cycles were imposed by the LBW process but was characterized by finer globular grains having an average grain size of 2.1 ± 0.7 µm and wider distribution of crystal directions between <0 0 0 1>//ND and <1 0 ‐1 0>//ND compared with that of the base material (**Figure 5(a)**). The main component is still (0 0 0 1)[2 ‐1 ‐1 0] ] (12.1 mrd <  *f* < 12.8 mrd); however, (0 1 ‐1 0)[2 ‐1 ‐1 0] was superimposed by (0 8 ‐8 1)[2 ‐1 ‐1 0] (*f* = 6.5 mrd), leading to misorientations between 3.6° and 4.4°. The BM component (0 1 ‐1 0)[2 ‐1 ‐1 4] has been dissolved, and new components such as (0 6 ‐6 5)[5 ‐5 0 6] (*f* = 6.4 mrd) and (0 3 ‐3 ‐4)[14 3 ‐17 15] (*f* = 9.4 mrd) have emerged within the orientation band (*φ*2 = 0°, *φ*1 = 60°, 0°< *φ* < 80°). The misorientations of these components are 26.2° and 37.3° with regard to (0 1 ‐1 0)[2 ‐1 ‐1 4]. The orientation band (*φ*2 = 0°, *φ*1 = 120°, 0°< *φ* < 180°) showed the occurrence of further com‐ ponents such as (0 5 ‐5 1)[‐3 1 2 5] (*f* = 5.6 mrd) and (0 6 ‐6 5)[‐8 9 ‐1 12] (*f* = 4.6 mrd), resulting in misorientations of 7.7° and 24.7° with regard to (0 1 ‐1 0)[‐2 1 1 4] (see pole figure in **Figure 5(a)**). The misorientations between the base material and heat‐affected zone indicate microstructural distortion of the laser beam welded Ti‐6Al‐4V butt joint due to heat input and subsequent rapid cooling during the joining process.

A significant change in microstructure was observed for the near‐HAZ at a distance of 1.05 mm from the FZ (**Figure 5(b)**). The equiaxed initial microstructure of the base material was transformed into an acicular morphology with a small amount of embedded globular grains. The transformation was connected with a grain refinement and a weakening of the (0 0 0 1)[2 ‐1 ‐1 0] component dominating the BM and HAZ adjacent to the BM. Some of the crystals rotated around the angles between 4.6° and 4.8°, which led to formation of (0 0 0 1)[4 ‐7 3 0] (5.9 mrd < f < 6.5 mrd). A further portion of crystals rotated around the <2 ‐1 ‐1 0>//RD crystal direction and tilted at an angle of 10.3°, resulting in the formation of the (0 1 ‐1 ‐1)[2 ‐1 ‐1 0] component (f = 7.3 mrd). Furthermore, the microstructure transformation led to the formation of new components such as (0 3 ‐3 2)[14 ‐12 ‐2 15] (f = 4.5 mrd), (0 1 ‐1 ‐1)[1 0 ‐1 1] (f = 2.1 mrd), (0 1 ‐1 ‐1)[‐6 ‐1 7 8] (f = 3.9 mrd) and (0 1 ‐1 ‐1)[‐17 19 ‐2 21] (f = 3.0 mrd) and to dissolution of (0 6 ‐6 5)[5 ‐5 0 6]. The (1 1 ‐2 0) pole figure shows a tendency towards the formation of <1 1 ‐ 2 0>//RD fibre texture present in the FZ also (**Figure 4(d)**).

#### *3.2.4. Influence of heat treatment on the microstructure*

EBSD analysis of heat‐treated specimens revealed no significant texture transformations upon PWHT. The main texture components remained approximately the same with slight deviations in numerical values of peaks. Thus, our further attention will be focused mainly on the microstructural characteristics, which can be clearly seen from the OM observations. **Figure 6** shows the influence of PWHT on the average grain size in the BM. Heat treatments at tem‐ peratures less than 750°C did not change the grain size in the BM significantly (**Figure 7(a)**, **(b)**). Recrystallization processes leading to coarsening of the microstructure were activated at higher temperatures starting from 800°C (FA2, DA). The maximum average grain size was achieved after DA and was 4.9 ± 1.5 µm. OM images of the BM microstructure after PWHT at temperatures higher than 750°C are shown in **Figure 7**. From this figure, it can be seen that α phase is fully recrystallized after duplex annealing, and almost all lamellar regions were transformed into equiaxed grains (**Figure 7(c)**). All other conditions of PWHT that are not shown in **Figure 7** led to almost the same microstructure as in the starting condition, and these images are omitted.

Effect of Microstructure Transformations on Fatigue Properties of Laser Beam–Welded Ti‐6Al‐4V Butt... http://dx.doi.org/10.5772/66178 123

**Figure 6.** Influence of PWHT on the average grain size in the BM and the width of lamellae in the FZ.

**Figure 7.** Influence of PWHT on the microstructure of the BM.

*f* < 12.8 mrd); however, (0 1 ‐1 0)[2 ‐1 ‐1 0] was superimposed by (0 8 ‐8 1)[2 ‐1 ‐1 0] (*f* = 6.5 mrd), leading to misorientations between 3.6° and 4.4°. The BM component (0 1 ‐1 0)[2 ‐1 ‐1 4] has been dissolved, and new components such as (0 6 ‐6 5)[5 ‐5 0 6] (*f* = 6.4 mrd) and (0 3 ‐3 ‐4)[14 3 ‐17 15] (*f* = 9.4 mrd) have emerged within the orientation band (*φ*2 = 0°, *φ*1 = 60°, 0°< *φ* < 80°). The misorientations of these components are 26.2° and 37.3° with regard to (0 1 ‐1 0)[2 ‐1 ‐1 4]. The orientation band (*φ*2 = 0°, *φ*1 = 120°, 0°< *φ* < 180°) showed the occurrence of further com‐ ponents such as (0 5 ‐5 1)[‐3 1 2 5] (*f* = 5.6 mrd) and (0 6 ‐6 5)[‐8 9 ‐1 12] (*f* = 4.6 mrd), resulting in misorientations of 7.7° and 24.7° with regard to (0 1 ‐1 0)[‐2 1 1 4] (see pole figure in **Figure 5(a)**). The misorientations between the base material and heat‐affected zone indicate microstructural distortion of the laser beam welded Ti‐6Al‐4V butt joint due to heat input and

A significant change in microstructure was observed for the near‐HAZ at a distance of 1.05 mm from the FZ (**Figure 5(b)**). The equiaxed initial microstructure of the base material was transformed into an acicular morphology with a small amount of embedded globular grains. The transformation was connected with a grain refinement and a weakening of the (0 0 0 1)[2 ‐1 ‐1 0] component dominating the BM and HAZ adjacent to the BM. Some of the crystals rotated around the angles between 4.6° and 4.8°, which led to formation of (0 0 0 1)[4 ‐7 3 0] (5.9 mrd < f < 6.5 mrd). A further portion of crystals rotated around the <2 ‐1 ‐1 0>//RD crystal direction and tilted at an angle of 10.3°, resulting in the formation of the (0 1 ‐1 ‐1)[2 ‐1 ‐1 0] component (f = 7.3 mrd). Furthermore, the microstructure transformation led to the formation of new components such as (0 3 ‐3 2)[14 ‐12 ‐2 15] (f = 4.5 mrd), (0 1 ‐1 ‐1)[1 0 ‐1 1] (f = 2.1 mrd), (0 1 ‐1 ‐1)[‐6 ‐1 7 8] (f = 3.9 mrd) and (0 1 ‐1 ‐1)[‐17 19 ‐2 21] (f = 3.0 mrd) and to dissolution of (0 6 ‐6 5)[5 ‐5 0 6]. The (1 1 ‐2 0) pole figure shows a tendency towards the formation of <1 1 ‐

EBSD analysis of heat‐treated specimens revealed no significant texture transformations upon PWHT. The main texture components remained approximately the same with slight deviations in numerical values of peaks. Thus, our further attention will be focused mainly on the microstructural characteristics, which can be clearly seen from the OM observations. **Figure 6** shows the influence of PWHT on the average grain size in the BM. Heat treatments at tem‐ peratures less than 750°C did not change the grain size in the BM significantly (**Figure 7(a)**, **(b)**). Recrystallization processes leading to coarsening of the microstructure were activated at higher temperatures starting from 800°C (FA2, DA). The maximum average grain size was achieved after DA and was 4.9 ± 1.5 µm. OM images of the BM microstructure after PWHT at temperatures higher than 750°C are shown in **Figure 7**. From this figure, it can be seen that α phase is fully recrystallized after duplex annealing, and almost all lamellar regions were transformed into equiaxed grains (**Figure 7(c)**). All other conditions of PWHT that are not shown in **Figure 7** led to almost the same microstructure as in the starting condition, and these

subsequent rapid cooling during the joining process.

122 Study of Grain Boundary Character

2 0>//RD fibre texture present in the FZ also (**Figure 4(d)**).

*3.2.4. Influence of heat treatment on the microstructure*

images are omitted.

Light microscopic investigations yielded that heat treatments up to 650°C (SRA2) did not affect the microstructure of the FZ significantly because such low temperatures were insufficient for martensite decomposition into equilibrium α + β structure. Sallica‐Leva et al. [26] studied the effects of heat treatment on the mechanical properties of an acicular α' martensite obtained by selective laser melting. In their work, the microstructure of samples heat treated at 650°C was very similar to that of the starting condition, whereas significant grain coarsening was observed after heat treatment at 800°C. The precipitation of β phase and the gradual transformation of α' into α phase by the diffusion of excess vanadium from α' to β phase were proposed to be the main events of martensite decomposition. Their results are in a good agreement with our findings. **Figure 6** shows the influence of PWHT on the average thickness of α' (α) laths in the FZ. Significant grain coarsening of the microstructure was observed after PWHT at tempera‐ tures higher than 750°C. This result is a consequence of the transformation of fine martensitic morphology into an equilibrium lamellar α + β structure and diffusion‐controlled growth of 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 seen in **Figure 8**.

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

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

#### **3.3. Microhardness**

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 higher strength and lower ductility, which are attributable to the fine size of martensitic plates and high defect density [28]. The influence of microtexture plays a minor role because <0 0 0 1>//ND and <1 0 ‐1 0>//ND crystal directions were aligned parallel to the indentation direction in both the BM and HAZ despite decreasing axial intensity. HAZ is characterized by strong inhomogeneity and plays a role of a transition zone from acicular martensitic morphology within the FZ to equiaxed microstructure in the BM. The strong spatial variation of the microstructure, namely, the decrease in the martensitic content, leads to a high gradient of microhardness inside the HAZ. Squilace et al. [2] reported that the hardness gradient in the HAZ is inversely proportional to the heat input during LBW.

**Figure 9.** Microhardness profile of the laser beam welded Ti‐6Al‐4V butt joint measured in the as‐welded condition.
