Materials Science and Engineering A 541 (2012) 110–119
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Superplastic deformation mechanism and mechanical behavior of a laser-welded Ti–6Al–4V alloy joint Shuhai Chen a , Jihua Huang a,∗ , Donghai Cheng a,b , Hua Zhang a , Xingke Zhao a a b
School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing, 100083, PR China School of Aeronautical Manufacturing Engineering, Nanchang Hangkong University, Nanchang, 330063, PR China
a r t i c l e
i n f o
Article history: Received 16 September 2011 Received in revised form 15 January 2012 Accepted 3 February 2012 Available online 11 February 2012 Keywords: Superplasticity Laser welding Phase transformation Titanium alloys
a b s t r a c t The mechanical behavior and superplastic deformation mechanism of a laser-welded Ti–6Al–4V alloy joint were investigated. Uniaxial tensile tests were performed on welded specimens at 870–920 ◦ C temperature and 10−3 to 10−1 s−1 strain rate. The microstructural evolution of the weld zone was observed under the strains of 43%, 143%, 229%, and 387%. The laser-welded joint was found to have good superplasticity under a suitable strain rate; the highest joint elongation reached 397%. Superplastic deformation in the weld zone accompanied the globularization of the as-welded microstructure. Continuous globularization ensured a good superplasticity of the laser-welded joint. As a major cause of the globularization of lamellar structures in the weld zone, the stress activated the diffusion of Al atoms by changing the chemical potential at the boundaries of the ␣ phase. Consequently, ␣ →  phase transformation occurred. The globularization of the as-weld microstructures was considered to be governed by this transformation and the development of grain boundary sliding. © 2012 Elsevier B.V. All rights reserved.
1. Introduction The two-phase ␣/ titanium (Ti) alloy Ti–6Al–4V is a widely used advanced structural material in aeronautics. This alloy has the advantages of low density, high strength, toughness, good hightemperature properties, and plasticity. Diffusion bonding (DB) combined with superplastic forming (SPF) is an advanced hot working method for Ti–6Al–4V alloys. Given their excellent design flexibility and considerably light weight, DB/SPF integral structures of Ti alloy are successfully used in the aviation industry, such as in hatch doors, airframes, wings, among others [1–3]. Recently, aircrafts have been required to possess more integrated and larger structures. This requirement is difficultly fulfilled by the DB/SPF technology. Firstly, DB requires a heavy pressure on the components, which restricts the design flexibility of complex components and the mold life. Secondly, to obtain reliable joining, DB should be performed at high temperatures and long durations. However, under these conditions, production efficiency is decreased and leads to significant grain growth, which is dramatically detrimental to the mechanical properties of the components. Therefore, studies on the combination of SPF with other welding methods under certain circumstances are being conducted.
∗ Corresponding author. Tel.: +86 010 62334859; fax: +86 010 62334859. E-mail address:
[email protected] (J. Huang). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2012.02.011
Friction stir welding (FSW) is a novel solid-phase joining technique for light alloys. In FSW, the relatively fine-grained microstructure of a parent material is retained. The FSW of aluminum (Al) and other light alloys has been utilized for several years. The combination of FSW with SPF may ideally replace DB/SPF. Hence, the application of FSW/SPF on Al and magnesium alloys warrants more attention [4–13]. Recently, a preliminary investigation on the FSW/SPF of Ti alloys has been conducted by Ramulu and coworkers [14–16]. Future studies focused on improving the FSW/SPF of Ti alloys are expected. Besides the abovementioned solid-phase joining methods of DB and FSW, other welding methods that can be combined with SPF include gas tungsten arc welding (GTAW), plasma arc welding (PAW), electron beam welding (EBW), and laser beam welding (LBW). Homer et al. [17] have investigated the superplastic behavior of GTAW-welded specimens of the Ti–6Al–4V alloy. Specimen elongation is found to be less than 150%. Kruglov et al. [18] have reported that the weld bead by GTAW almost has no plastic deformation during the bulge-forming of the Ti alloy tailored blank. Hence, the superplasticity of a GTAW-welded joint is not ideal. Zhang et al. [19] have studied the superplastic deformation and microstructures of a Ti-weld bead by PAW. The weld bead by PAW has better superplasticity than that by GTAW. The microstructures of the weld have mixed microstructures of globular ␣ and long strip ␣ during bulge forming.Joints welded by high-energy beam welding such as EBW and LBW exhibit possibly better superplasticity. High-energy beam welding is rapid, has a narrow heat-affected
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zone (HAZ), and finer-grained weld beads [20]. The EBW technique needs a vacuum environment, which extremely limits its application range. In contrast, LBW with SPF can result in more integrated and larger structures. Will et al. [21,22] have investigated the superplastic deformation of multi-sheet structures with Ti alloy by LW/SPF. However, the relationship between the microstructure and mechanical property of the laser-welded bead is not successfully established. Our previous works [23,24] have indicated that the laser-welded plate of the Ti–6Al–4V alloy has good superplasticity under suitable tensile conditions. The acicular structure of the weld bead is transformed into equiaxed grains during the tensile process, which has great potential applications. However, two issues on the superplastic deformation of laser-welded Ti–6Al–4V alloy joints need to be addressed. First, the mechanical behavior of a laser-welded plate, which contains a partial parent metal, does not adequately reflect that of the weld zone (WZ). To control the formation of structural parts during superplastic deformation, the mechanical behavior of the WZ needs to be evaluated. Second, the mechanism of the superplastic deformation of a WZ is unclear. As to the mechanisms of superplastic deformation for ideal microstructures, the well-accepted phenomenon of grain boundary sliding (GBS) is mainly responsible for large joint elongations [25–29]. However, as-welded microstructures are not ideal for superplastic deformation due to the non-equilibrium thermal process during laser welding. The globularizationglobularization mechanisms of lamellar microstructures in Ti alloys during thermomechanical processing may provide useful information. The mechanisms of transformation from lamellar microstructures into equiaxed grains can be summarized as follow. Firstly, both low- and high-angle subboundaries across ␣ plates are formed by the generation and recovery of dislocations during hot deformation. The  phase penetrates into the ␣ plates along these sub-boundaries, which separate the ␣ lamella into shorter segments [30–34]. Secondly, platelet coarsening or globularizationglobularization is caused by diffusion across a chemical potential/solute concentration gradient between the sharp edge of a platelet and the flat interphase boundary [35,36]. Thirdly, the globularizationglobularization of lamellar microstructures may be related to an ␣ →  phase transformation [37,38], although this has not been confirmed. Non-equilibrium microstructures in a laser-welded zone are different from equilibrium lamellar structures because of the high solidification speed during laser welding. Hence, ␣ →  phase transformation may play an important role in the globularizationglobularization of welded microstructures. In addition, the abovementioned mechanisms are obtained by compressive deformation. The influence of tensile deformation on microstructures is different from that of compression. In the present study, the effects of deformation conditions such as temperature and strain rate on the mechanical behavior of a laser-welded zone were systematically described. Based on observations and analyses of the effect of strain on microstructural evolution, the transformation mechanism from as-welded into equiaxed structure was clarified.
Fig. 1. Dimensions of the hot tensile specimen.
comparison. Fig. 1 shows the dimensions of the specimens. The welded specimens were cut along the test direction parallel to the weld line. Tensile tests were performed at 870–920 ◦ C temperature and 10−3 to 10−1 s−1 strain rate. The temperature was controlled by a three-zone furnace and thermocouples located around the sample. During the hot tensile tests, the specimens were heated at 1 ◦ C/s. Tensile tests were performed for 5 min after reaching the target temperature, as shown in Fig. 2. To analyze the microstructural evolution process of superplastic deformation in the WZ, the microstructures were observed under strains of 43%, 143%, 229%, and 387%. In the present study, standard grinding and polishing sample preparation procedures were used. Solutions composed of 2% HF, 10% HNO3 , and 88% H2 O were used to etch the samples. The microstructures of the specimens before and after tensile tests were observed by optical and scanning electron microscopes. The crystal phases of the laser welding seam were identified by a microbeam X-ray diffractometer. Transmission electron microscopy (TEM) was used to characterize the microstructures in detail. The specimens for TEM were prepared by cutting, grinding, and ion milling to electron transparency. 3. Results 3.1. Microstructure of the joint Fig. 3a shows the macro-section of a typical laser-welded joint. The joint had a width of 1.2 mm, and was composed of a WZ and a HAZ. The width of the HAZ was only 0.4 mm. Coarse grains appeared in the WZ, but fine grains were found in the HAZ. Fig. 3b–e exhibits the microstructure morphology of the parent metal, WZ, and HAZs near the WZ and parent metal. The microstructures of the parent metal consisted of equiaxed ␣ grains and  phases around the grain boundaries (Fig. 3b). A mass of acicular structure with different crystallographic orientations appeared in the WZ, as shown in Fig. 3c. To identify the crystal phases of the seam, X-ray diffraction (XRD) analysis was performed at WZ, as shown in Fig. 4. Besides martensite ␣ , some  phases were also found in the WZ. Considering the rapid cooling of the welded metal during laser welding; the as-welded metal transformed into a mass of acicular martensite ␣ and residual  phases.
2. Experimental materials and procedures A 0.8 mm-thick Ti–6Al–4V alloy annealed sheet was used in the experiment. The balanced chemical composition was 6.2Al–4V–0.08C–0.25Fe–0.05N–0.015H–Ti. The microstructure consisted of equiaxed ␣ grains and  phases around the grain boundaries. All joints were welded using a 4 kW CO2 laser at 1300 W power and 3 m/min welding speed. Hot tensile tests were carried out on both welded and unwelded specimens for
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Fig. 2. Heating process during a hot tensile test.
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Fig. 3. Microstructures of the joint: (a) macro-section, (b) parent metal, (c) weld zone (WZ), (d) heat-affected zone (HAZ) near the WZ and (e) HAZ near the parent metal.
The martensite ␣ was supersaturated solid solutions. Fig. 5 shows the microstructure morphology of the WZ under TEM. High-density dislocations appeared inside the weld. The dislocations may have evolved from the boundary or sub-boundary by their recovery and rearrangement during the hot tensile test. This test positively influenced the globularizationglobularization of the as-welded microstructures. The microstructures consisted of an alternately distributed ␣ and residual  phases, as shown in Fig. 5a. There were martensite colonies that consisted of acicular subgrains in more or less the same orientation. The residual  phases existed in the sub-boundaries between adjacent subgrains, as shown in Fig. 5b (indicated by the arrows). A little rhombic martensite ␣ with compound twin structures was formed in the WZ, as shown in Fig. 5b (indicated by the ellipse).
The microstructure of the HAZ was composed of ␣ +  + ␣ phases. Acicular martensite ␣ in the HAZ near the WZ was greater in quantity and denser than those near the parent metal, as shown in Fig. 5d and e. The microstructure of the Ti–6Al–4V joint by LBW is described in more detail elsewhere [39]. 3.2. Mechanical behavior of superplastic deformation The mechanical behavior of the superplastic deformation of the welded specimens has been investigated in our previous works [23,24]. However, the results do not exhibit complete mechanical behavior of the WZ because the welded specimens partially include the parent metal in addition to the WZ and HAZ. Knowledge of the mechanical behavior of the superplastic deformation
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Fig. 4. XRD patterns at the weld zone.
of the WZ and HAZ can provide substantial insight into the design of LBW/SPF. In the present paper, the stress of the WZ was obtained by deducting the load of the parent metal. At high temperatures, the microstructures of the HAZ and parent metal had little difference [40]. Therefore, high-temperature mechanical properties of the HAZ and parent metal were assumed to be the same. The stress of the WZ was calculated as: w =
sum × Ssum − m × Sm Sw
(1)
where sum is the stress of the welded specimen (in MPa), Ssum is the load area of the welded specimen (in mm2 ), m is the stress of the parent metal (in MPa), Sm is the load area of the parent metal in the welded specimen (in mm2 ), w is the stress of the WZ (in MPa), and Sw is the load area of the WZ (in mm2 ). Based on the cross-section of the WZ, the Ssum , Sm , and Sw areas were 4.0, 3.2, and 0.8 mm2 , respectively. sum and m were obtained by the tensile tests of the welded specimens and parent metal. Fig. 6 shows the calculated stress–strain curves of the WZ under strain rates of 10−1 , 10−2 , and 10−3 s−1 . The stress in the early stage of deformation sharply increased with increased strain because of strain hardening. In contrast, the increase in the stress rate gradually subsided with increased strain because of the strengthened flow softening behavior. Hence, the flow stress peaked at a critical strain, and then decreased with further straining under the strain rates of 10−1 and 10−2 s−1 , as shown in Fig. 6a and b, respectively. However, after peaking, the flow stress became constant (even slightly increased) with increased strain of 10−3 s−1 (Fig. 6c). Dynamic equilibrium was achieved between strain hardening and flow softening. Therefore, the laser-welded joint had good superplasticity at low strain rates. The flow stress also decreased with increased temperature rise and decreased strain rate. The temperature rise induced decreased yield strength of the matrix; the decreased strain rate increased the deformation time such that dynamic recovery completely proceeded. All these phenomena contributed to the decreased flow stress. Peak stress is an important parameter for analysing mechanical behavior during superplastic deformation. Fig. 7 shows the influence of the temperature and strain rate on the peak stress of the WZ and parent metal. The peak stress of the WZ and parent metal clearly decreased with increased temperature and decreased strain rate. The peak stress of the WZ was clearly higher than that of the parent metal because of the acicular microstructure of the WZ. Fig. 8 shows the effects of the temperature and strain rate on the elongation of the WZ and parent metal. The elongation of the WZ was about one half of the parent metal. However, the high
Fig. 5. TEM image of the weld zone: (a) alternately distribution of ␣ and residual  phases and (b) residual  phases between adjacent subgrains.
elongation (233%–397%) of the WZ was still within the superplastic range. These features further indicated that the WZ had good superplasticity. At high strain rates (10−1 and 10−2 s−1 ), the elongation of the WZ and parent metal increased with increased temperature. At a low
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Fig. 7. (a) Peak stress of the weld zone. (b) Peak stress of parent metal.
3.3. Microstructure evolution
Fig. 6. Stress–strain curves of the welded specimens under different strain rates: (a) 10−1 s−1 , (b) 10−2 s−1 and 10−3 s−1 .
strain rate of 10−3 s−1 , the elongation initially increased, and then decreased with increased temperature. Subsequently, the elongation reached a maximum value at 900 ◦ C. Fig. 9 shows the photos of specimens at strain rate of 10−3 s−1 before and after tensile. This result was due to the insignificant growth of the grain at high strains because of the short test duration. In contrast, growth was significant at low strains because of the long test duration. Coarse grains also induced the decreased plastic deformation ability.
At high temperatures, elemental diffusion was very spontaneous. Consequently, the reaction ␣ → ␣ +  proceeded within 5 min before the load was applied. The acicular martensite ␣ then transformed into a lamellar ␣ phase, which played an important role in microstructure evolution. Fig. 10 shows the evolution of the lamellar structure under the strains of 43%, 143%, 229%, and 387%, as well as deformation temperature of 920 ◦ C. At 43% strain, the lamellar ␣ phase clearly decreased compared with the as-welded acicular martensite microstructures, as shown in Fig. 10a. The long direction of the lamellar ␣ phases tended to be parallel with the tensile direction. The plastic deformation of the lamellar ␣ phases was not significant. However, some studies have reported severe deformation during compressive deformation [30–36]. In the present study, the mechanism of tensile deformation was found to be different from compressive deformation. Fig. 10b shows a magnified image of the microstructure at the zone of ␣ colonies. These colonies had the characteristics of lamellas, but also presented splitting features. The morphological characteristics of the cleavage of some lamellar ␣ phases along the long direction were observed. The  phases markedly thickened between these splitting lamellas. The ␣ →  phase transformation occurred between adjacent lamellas of the colonies. With increased strain, the aspect ratio (i.e., length:width ratio) of the ␣ phase clearly decreased, and more globular grains appeared in the WZ. However, some lamellar ␣ phases still remained in
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The specimens were broken when the strain reached 387%. Fig. 10g and h exhibit the microstructures of the WZ after the specimens were broken. The transformation from as-welded to equiaxed structure was clearly achieved outright in the WZ. 3.4. Effects of different strains and temperatures on the globularizationglobularization mechanism
Fig. 8. (a) Elongation of the weld zone. (b) Elongation of parent metal.
certain zones, as shown in Fig. 10c. At 229% strain, the lamellar grains almost disappeared, and globular fractions remarkably increased, as shown in Fig. 10e. Nevertheless, under higher magnification, string grains were found to form from the splitting lamella ␣ phase, and some grains still had slightly high length:width ratios, as shown in Fig. 10f. The ␣/ phase boundaries had irregular morphologies, which may have been caused by the ␣ →  phase transformation during the deformation.
Superplastic deformation in the WZ accompanied the globularization of lamellar structures. To investigate quantitatively the globularizationglobularization process during superplastic deformation, the ␣ phase with a length:width ratio lower than 2 was defined as the equiaxed grain. The equiaxed grain rate in the WZ of 0.5 mm × 0.2 mm was defined as the globularizationglobularization fraction. Fig. 11 shows the influences of different strains and temperatures on the globularizationglobularization fraction. Globularization fraction increased with increased strain. The deformation temperature affected, to some extent, the globularizationglobularization of the lamellar ␣ phase. The influence of strain was clearly higher than that of the temperature. Therefore, strain was a major cause of the globularizationglobularization of the lamellar structure in the WZ. As shown in Fig. 11, temperature had little influence on the globularizationglobularization at certain strains. To clarify the influence of the annealing temperature on the globularization without hot deformation, the as-welded specimens were held at 920 ◦ C for 5 s, 5 min, and 50 min. Their microstructures are shown in Fig. 12. The fine acicular martensite ␣ transformed into the coarse lamellar ␣ phase, and the length as well as thickness of the ␣ phase increased with increased holding time. This result indicated that the globularization of the as-welded microstructures did not spontaneously appear during the annealing. Annealing alone did not result in the globularization of the as-welded microstructures. Rather, globularization was due to the high anisotropy of the interface energy of the dual-phase Ti alloy microstructures. Consequently, the lamellar structure was conferred a higher stability than the spherical shaped, based on a thermodynamics perspective [41,42]. When fine-grained globular microstructures are annealed, globular ␣ phase particles may become more like plates [37]. Based on the microstructure in Fig. 12, this phenomenon was also established from the fact that the length and thickness of the ␣ phase increased with increased holding time. The lamellar structures can also be transformed into equiaxed structures during annealing by increasing the interface energy, such as in a pre-deformation procedure [34]. The aforementioned results indicated that the strain was essential to the promotion of the globularization of the as-welded microstructures with lamellar structures. On one hand, the globularization of the lamellar structures enhanced the ability of the plastic deformation of the seam. On the other hand, the continuous plastic deformation of the seam promoted the globularization of the lamella structures. These two phenomena were mutually reinforcing. During plastic deformation, the effect of the strain on the microstructure evolution was shown by the stress. Therefore, the effect of the stress on the globularization was discussed in detail. 4. Discussion
Fig. 9. Photos of specimens at strain rate of 10−3 s−1 before and after tensile.
GBS and its accommodation process are well accepted as the major mechanisms of superplastic deformation. Even under nonideal conditions (e.g., low temperatures, non-ideal microstructures, etc.) where the mechanisms are complex, GBS is still an important mechanism of superplastic deformation. First of all, despite the GBS is difficult to no-equiaxed grain, the GBS of non-ideal microstructures has been observed [34,37,38]. Secondly, the GBS plays more important role with globularization development in WZ.
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Fig. 10. Effects of different strains on the microstructures, (a) and (b) 43%. Effects of different strains on the microstructures, (c) and (d) 143%. Effects of different strains on the microstructures, (e) and (f) 229%. Effects of different strains on the microstructures, (g) and (h) 387%.
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Fig. 11. Effects of strain and temperature on the globularization fraction.
In fact, the continuous globularization of the as-welded lamellar structure was a key factor in maintaining the good superplastic deformation of the WZ, according to the above observation of the as-welded microstructure evolution. Therefore, the superplastic deformation mechanism for the as-welded microstructures can be clarified by their globularization mechanisms. During the superplastic deformation of ideal microstructures, an increase in  volume fraction was observed [43,44]. This result indicated the existence of an ␣ →  phase transformation. For lamellar microstructures during compressive deformation, an ␣ →  phase transformation significantly influences globularization [37,38]. This influence may increase for as-welded non-equilibrium microstructures. At the initial stage of deformation, a mass of ␣ phase colonies that were unfavorably oriented with respect to the applied stress may have rotated, facilitating further shear. During rotation, these colonies suffered shear stress along the long direction of the lamellar ␣ phase. Based on the observed microstructures in the WZ, there was a thin residual  phase between adjacent lamella ␣ phases inside the colonies. Therefore, the GBS between the residual  and lamellar ␣ phases was initially activated by the shear stress because the sliding resistance of ␣/ GBS was far less than that of ␣/␣ GBS [28,29]. On one hand, the GBS causes a stress gradient along the tilted grain boundaries to the tensile axis [44]. On the other hand, the long direction of the ␣ lamella tended to be parallel with the tensile direction of the specimens. Consequently, the ␣ lamella roughly suffered compressive and tensile stresses in the transverse and long directions, respectively. The stress gradient between these boundaries induced the changes in the chemical potential of the vacancies. Consequently, a diffusional flow of atoms from the compressive to the tensile boundaries occurred. Al is the most diffusionally mobile of the alloying elements in Ti alloys [38]. The diffusion of Al atoms along this region with the lattice defects induced the ␣ →  phase transformation because Al is an ␣-stabilizer. Therefore, the  phase was formed along the long direction of the lamella. Fig. 13 shows the microstructure morphology induced by the ␣ →  phase change and ␣/ GBS under 143% strain, 0.1 s−1 strain rate, and 900 ◦ C temperature. The ␣ →  phase change created favorable conditions for the globularization of the lamellar structures. The roughly parallel lamellar ␣ phases in some zones suffered shear stress in the transverse direction. Unlike compressive deformation, which could significantly deform the lamellar ␣ phase, the tensile deformation level by the shear stress was small. However, tensile deformation continuously generated a number of dislocations depending on the grain matrix deformation inside the
Fig. 12. Effects of different holding times on the as-welded microstructures, (a) 5 s, (b) 5 min, and (c) 50 min.
lamellar ␣-phase. At the same time, recovery processes resulted in the rearrangement of the dislocation structures of the ␣-phase, and transverse low-energy semi-coherent boundaries were then formed. The stress ensured the continuous introduction of lattice dislocations into the boundaries, resulting in increased energy. Therefore, the stress further induced the conversion of the semicoherent boundaries into higher-energy non-coherent boundaries inside the lamellar ␣-phase. Finally, unstable dihedral angles between the interphase ␣/ and intraphase ␣/␣ boundaries were formed. The equilibrium between the various surface tension forces at the interphase/intraphase triple junctions led to increased chemical potential along intraphase ␣/␣ boundaries. Al atoms diffused into the interphase ␣/ boundaries, resulting in ␣ →  phase transformation and the formation of a groove along the ␣/␣ boundary.
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Fig. 13. Microstructure morphology induced by ␣ →  phase change and ␣/ GBS under 143% strain, 10−1 s−1 strain rate, and of 900 ◦ C temperature.
The ␣ →  phase transformation continuously proceeded in the groove, which caused an irregular morphology of the ␣/ phase boundaries, as shown in Fig. 10f. The deepening of the grooves from both sides of an ␣ lamella eventually led to the separation of ␣ fragments. Lamellar ␣-phases with different orientations in some zones did not rotate along the shear stress direction due to their struts. The compressive stress was concentrated around the fulcrum of the lamellar ␣ phase, which generated a number of dislocations in these zones. The rearrangement of the dislocation structure led to the formation of sub-boundaries. Simultaneously, the compressive stress concentration increased the chemical potential in these zones, resulting in the diffusion of Al atoms, which stimulated the ␣ →  phase transformation. The cusp of a lamellar ␣ phase penetrated other lamellar ␣ phases along the sub-boundaries, depending on continuous ␣ →  phase transformation. Consequently, the aspect ratio of the ␣-phase further decreased, as shown in Fig. 14. The ␣-phases were also broken into small grains by extrusion with one another, as shown in Fig. 15. These broken grains had irregular morphologies with sharp edges or platelets. The gradient in the chemical potential/solute concentration between the sharp edge and flat interphase boundary drove solute diffusion. Generally, the concentration gradient is typically quantified using the Gibbs–Thompson equation. In such cases, Al atoms diffuse from a sharp edge to a flat surface, resulting in ␣ →  phase transformation in the sharp edge, and  → ␣ phase transformation in the flat
Fig. 15. Breaking of ␣ grains by extrusion with one another.
surface. Ultimately, the sharp edges of the irregular morphologies decreased or the platelets thickened. In summary, ␣ →  phase transformation is the dominant mechanism in the globularization of as-welded microstructures during superplastic deformation. During this process, GBS played an important role in promoting globularization. On one hand, the lamellar ␣ phases were separated by ␣/ and ␣/␣ GBS. On the other hand, GBS increased the interfacial energy such that the chemical potential gradient increased, which induced strengthened elemental diffusion. 5. Conclusions Based on investigation of the mechanical behavior and mechanism of superplastic deformation of a Ti–6Al–4V alloy joint by laser welding, the following conclusions were drawn. (1) The laser-welded joint had good superplasticity at low strain rates; the highest joint elongation reached 397%. The peak stress of the WZ decreased with increased temperature and decreased strain rate. The value of the peak stress was clearly higher than that of the parent metal. The elongation initially increased and then decreased with temperature rise at low strain rates (10−3 s−1 ). The elongation reached a maximum value at 900 ◦ C. (2) Superplastic deformation in the WZ accompanied the globularization of the long-stripped structures. Strain was a major cause of the globularization of the lamellar structures in the WZ. Temperature had little influence on the globularization at certain strains. However, annealing alone did not result in the globularization of the as-welded microstructures. (3) The continuous globularization of the as-welded lamellar structures was a key factor in maintaining the good superplastic deformation of the WZ. The globularization of the as-welded microstructures was governed by ␣ →  transformation and GBS development. References [1] [2] [3] [4] [5] [6]
Fig. 14. Microstructures by each other penetration of ␣ phases under 229% strain, 10−3 s−1 strain rate, and 870 ◦ C temperature.
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