Kink band induced β→α phase transformation in lamellar α+β Ti alloys

Generally, deformation in Ti and Ti alloys are dominated by dislocations and deformation twins. In this work, a novel deformation band (DB) has been observed in the fully lamellar Ti-6Al-4V alloy. Through transmission electron microscopy (TEM) investigation, the DB is determined as a kink band (KB) resulting from the pile-up of  dislocations on { $ 10\bar{1}0 $ 101¯0} prismatic planes. More intriguingly, the KB can pass through the β phase and then induce the phase transformation from β to α phase, and the accompanying composition change may have been accommodated by diffusion driven by the high deformation energy under severe deformation conditions. GRAPHICAL ABSTRACT IMPACT STATEMENT First observing the existence of kink band and unveiling the kink band induced β to α phase transformation mechanism in α+β Ti alloys.


Introduction
Dual-phase α+β Ti alloys composed of hexagonalclose-packed (HCP) α phase and body-centered-cubic (BCC) β phase have received widespread attention owing to their high specific strength, low density and remarkable corrosion resistance [1][2][3].Compared with bimodal or fully-equiaxed α+β Ti alloys, the α+β Ti alloys with fully lamellar structure have better fracture toughness, anti-fatigue crack propagation performance, namely good damage tolerance, and are used in the aviation industry [1,4,5].It is because the fully lamellar α+β Ti alloys are composed of a large number of α/β interfaces, which can not only hinder dislocation movement, but also act as a dislocation source to promote dislocation generation [6][7][8][9][10].Actually, during the deformation process, not only dislocations, but also deformation twins, can also be stimulated to accommodate the plastic deformation in fully lamellar α+β Ti alloys [11][12][13][14].In this way, the interaction between dislocation/twin and β phase is inevitable, which will significantly affect the properties of alloys.Therefore, it is particularly crucial to explore the deformation modes and their interaction mechanism with the β phase in the fully lamellar α+β Ti alloys.
In the fully lamellar α+β Ti alloy, the volume fraction of β phase with BCC structure is only about 10%, and the deformation is still dominated by α phase with HCP structure.However, owing to the low symmetry of HCP structured metal, the von Mises criterion is not usually satisfied by the dislocation slip systems [15].Consequently, deformation twinning, slip banding and shear banding are activated to accommodate the plastic deformation [16][17][18].For example, our previous research has indicated that {10 12} twin can continuously propagate via sequential kinking in β phase, which effectively toughened the dual-phase lamellar Ti-6Al-4V alloy [11].Moreover, the interaction between the {11 21} twin and β phase can induce the nucleation of sequential {10 12} deformation twins in lamellar Ti-6Al-4V alloys, which plays a role of strain releasing [12].In addition to deformation twins, a large number of deformation bands (DB), such as slip bands and shear bands, can be produced in the fully lamellar Ti-6Al-4V alloys [6,14,[19][20][21].It is reported that a large number of slip bands will be generated in the fully lamellar Ti-6Al-4V alloy after fatigue deformation, which can run through the entire α cluster and continuously cross multiple α/β interfaces [20,21].Shear bands can also form in the severely deformed Ti-6Al-4V alloys, which destroy the local structure, as well as the mechanical properties, eventually lead to a failure [6].In materials with the HCP structure, besides slip bands and shear bands, a large number of kink bands (KB) can also nucleate during deformation [22,23].For example, in the Ti/Ti 3 Si 2 C composite, KB has even become the dominant deformation mode providing crystal plasticity at room temperature [17].However, in pure Ti and dual-phase Ti alloys, the formation mechanism of kink band is still unclear.Specially, the interaction behavior between KB and β phase has predictably huge effect on the deformation and failure of the fully lamellar Ti alloys.
In the present work, we observed a new type of microband in as-deformed Ti-6Al-4V alloy.With the help of TEM, the new type of micro-band can be determined as a KB which is induced by the pile-up of < a > type dislocations on the {10 10} prismatic plane.Besides, atomic crystal structure and elemental analysis show that the interaction between KB and β phase results in the transformation from β phase to α phase under sever plastic deformation.

Materials and method
An α+β titanium alloy Ti-6Al-4V with chemical compositions (in wt.%) of Al 6.05, V 4.10, O 0.06, Fe 0.05 and Ti balance is used in this work.The β transus temperature of this alloy is 970 ± 5°C.To produce a fully lamellar microstructure, this alloy was held at 1000°C for 1 h and then furnace cooled to room temperature.The microstructure of the as-annealed Ti-6Al-4V is shown in Figure 1(a).The rod tensile samples with a gauge length of 60 mm and a diameter of 12.5 mm were strained on the SANS-CMT 5205 with a strain rate of 10 −4 at room temperature.
After tensile test, typical deformation characters of lamellar Ti-6Al-4V were characterized by a FEI Tecnai G 2 F30 and a FEI Titan G 2 60-300 TEM, and the TEM samples were taken from the region within 2 mm from the tensile fracture and then the sample was cut to 2mm × 1mm × 0.5 mm and subsequently ground to about 40 μm thickness, followed by dimpling and ion milling by using a Gatan PIPS.

Results and discussion
Firstly, the microstructure of as-deformed lamellar Ti-6Al-4V alloy is illustrated in Figure 1.Due to the severe plastic deformation and complex stress conditions around the fully lamellar Ti-6Al-4V tensile fracture, a large number of DB generated near the fracture.Figure 1b and c presents bright-field (BF) and dark field (DF) TEM images of the DB, respectively.It can be found that DB has a distinct interface with the matrix.The angle between [1 210] of the DB and of the matrix is about 5°−15°m easured by TEM, so the DB and the matrix forms a low angle grain boundary (LAGB).Normally, the LAGB is generated by the continuous arrangement of a series of edge dislocations or screw dislocations at the interface.Figure 1d is the high-resolution TEM (HRTEM) image of the interface between the α matrix and the DB, viewed along the [1 210] of the matrix.It can be seen that the interface between the matrix and the DB is near their (10 10) plane.In addition, at the α/DB interface, the Burgers vector of the interface dislocation can hardly be obtained.The reason may be that the dislocation Burgers vector component is parallel to the [1 210] direction at the interface.
To further analyze the type of DB, we performed diffraction analysis on the DB and the matrix.Figure 2a  and b show the BF and DF TEM images of the DB, respectively, where the matrix is viewed along its [1 210].The c and d indicate positions for selected area electron diffraction (SAED) of the matrix and the DB, as shown in Figure 2c and d, respectively.Compared with the diffraction pattern of [1 210] zone axis in Figure 2c, only the 0002 diffraction spots in Figure 2d are prominent, indicating that the DB has rotated around the [0001] axis relatively to the matrix.After that, the Kikuchi patterns of the matrix and the DB were collected, as displayed in Figure 2e and f, respectively.It can be found that the [1 210] pole of the Kikuchi pattern of the DB has moved along the 0002 Kikuchi line and is far away from the [1 210] pole of the matrix, which indicates that the DB is relatively rotated around the [0002] direction (rotation in the (0002) plane), compared with the matrix.In addition, according to the corresponding relationship between the Kikuchi line and the matrix crystallography, the 0002  S2(c-d).In this way, based on our experiment results shown in Figure 2(a-b), the boundary can be determined as tilt boundary and we can confirm that the DB in the fully lamellar Ti-6Al-4V alloy is a KB formed around the [0002] axis.
With the nucleation of KB in α phase, the interaction between KB and β phase is inevitable.As shown in Figure 3a, KB can continuously pass through the β phase.To further understand the interaction mechanism between KB and the β phase, HRTEM analyses of the interaction region between KB and the β phase in Figure 3a was performed, as shown in Figure 3b.Based on fast Fourier transformation (FFT) on the region c (β phase without transmission of KB, Figure 3c) and the region d (β phase with transmission of KB, Figure 3d), respectively, the structure transformation in different area can be determined.It can be found that the FFT pattern corresponding to the c region illustrates that it still maintains the crystal structure of the β phase, while the FFT pattern corresponding to the d region shows that the crystal orientation is along the [2 11 0] direction of the α phase, which proves that the crystal structure of the d region has changed from the BCC β phase to the HCP α phase.Furthermore, we characterized the element distribution of the interaction region between the KB and the β phase to further confirm the phase transformation from β to α phase.In Ti-6Al-4V alloy, V and Fe are stable elements of β phase.As a result, compared with α phase, V and Fe elements are rich in β phase, while Ti and Al elements are short.Hence, we  can use the content of Ti, Al, V, Fe elements in different regions to determine whether the β phase is transformed.Figure 3e is a high-angle annular dark field scanning-TEM (HAADF-STEM) image of the interaction region, and energy dispersive spectrometer (EDS) analysis was carried out on the red box area, as shown in Figure 3f.It can be found that in the interaction area between the β phase and the KB, the content of elements such as V and Fe is lower than that of the β phase region without interaction, and at the same time the content of elements such as Ti and Al are higher than that of the β phase region without interaction, which proves that the interaction between KB and β phase results in the transformation from β to α phase.Moreover, due to the shear deformation induced by kink band is not so obvious along [1 210] direction, we tilt the sample away from [1 210] direction and can see the obvious shear deformation of β phase induced by kink band in Fig. S3.
As mentioned above, the nucleation of KB can be attributed to the accumulation of < a > type dislocation on the {10 10} plane.The sketch of the nucleation process of KB is illustrated in Figure 4a.With the continuous slipping and accumulation of the < a > dislocations on the prismatic plane, a KB will nucleate with a rotation relative to the matrix around [0001] axis.Meanwhile, the schematic image of the atomic-scale interface structure is also presented in Figure 4b.The sketch obviously indicates the regular distribution of the full dislocations with a Burgers vector of 1/3a < 11 20 > on the KB/matrix interface.The regular arrangement of these 1/3 a < 11 20 > full dislocations result in the formation of KB which has a rotation angle 5°−15°relative to the matrix.With the nucleation of KB in matrix, the deformation can be accommodated and then contribute to the plasticity of lamellar Ti-6Al-4V, which is the same as the effect of KB on toughness reported in magnesium alloys [22].In the fully lamellar Ti-6Al-4V alloy, it can be judged that the nano KB observed in our experiment becomes an additional way to accommodate the deformation apart from slipping and deformation twinning.Furthermore, Figure 4c-e shows the atomic models of the KB/matrix interface in three different directions.Due to the overlap between matrix and KB, full dislocations with a Burgers vector of 1/3 < 11 20 > cannot be observed at the interface.That is why we cannot see the distribution of dislocations at KB/matrix interface in Figure 1c along the [1 210] direction.
Moreover, as mentioned above, with the interaction between KB and β phase, the β phase will finally transform to the α phase.So how does the β phase convert to the α phase?As early as 1934, Burgers proposed the kinetic mechanism of β phase to α phase transition under loading [24].As shown in Figure 5, under certain stress conditions, when the β phase is subjected to shear stress along the {112} < 111 > direction, the atoms on different {112} planes will move, so that the original direction on the same {110} plane (the angle is 70°32') is transformed into the < 11 20 > direction of the hexagonal structure (the angle is 60°).After that, with the increasing of shear stress, the atoms on each {110} plane of the β phase with the BCC structure move along the { 1100} direction of the α phase, which eventually induces the β phase with the BCC structure completely transforming into the α phase with the HCP structure.The observed change in composition accompanying the transformation (Figure 3f) can be attributed to diffusion driven by the high deformation energy, similar to what occurs in severe plastic deformation process [25,26].

Conclusion
In summary, we found that a large quantity of nanoscale DB appeared near the fracture of the fully lamellar Ti-6Al-4V alloy after tensile deformation.With the help of HRTEM and Kikuchi pattern, the DB is determined as KB formed by the accumulation of 1/3 a < 11 20 > dislocations on {10 10} prismatic plane in the α phase, which induce a rotation angle 5-15°around [0001] direction relative to the matrix.In addition, we also observed that the KB can pass through the β phase.After the interaction between KB and β phase, the β phase will transform to the α phase, in accordance with the Burgers orientation relationship under severe plastic deformation.All these discoveries provide further insights into the deformation behavior in lamellar α+β Ti alloy.

Disclosure statement
No potential conflict of interest was reported by the author(s).

Figure 1 .
Figure 1.The microstructure of the as-annealed fully lamellar Ti-6Al-4V alloy and typical morphology of DB after tensile deformation: (a) the morphology of the as-annealed fully lamellar Ti-6Al-4V alloy with alternatively distributed α phase and β phase.(b) BF TEM image of DB; (c) DF TEM image of DB; (d) HRTEM image of the interface between matrix and DB.

Figure 2 .
Figure 2. The typical morphology, SAED pattern and Kikuchi pattern of matrix and DB under the direction of [1 210] M : (a) BF TEM image of DB; (b) DF TEM image of DB; (c) SAED pattern of matrix; (d) SAED pattern of DB; (e) Kikuchi pattern of matrix; (f) Kikuchi pattern of DB.

Figure 3 .
Figure 3.The morphology for the interaction area between β phase and KB: (a) Typical BF-TEM image of the KB; (b) a HRTEM image of the interaction area between β and KB indicated in (a); (c, d) The corresponding FFT of the area shown in the white and yellow square indicated in (b), respectively; (e) Typical HADDF-STEM image of the interaction area between KB and β phase, and the enlarged morphology of shear deformation in β phase at the right corner; (f) the EDS analysis of the interaction area shown in (e), showing preponderance of Ti, Al and lack of V, Fe in the β phase sheared by KB.

Figure 5 .
Figure 5. 3D schematic lattice correspondence between the BCC β phase and the HCP α phase during β to α transformation.(a) The original 3D lattice structure of β; (b) the transformation from β phase to α phase; (c) the final 3D lattice structure of α phase.
CONTACTShijian Zheng sjzheng@hebut.edu.cnTianjin Key Laboratory of Materials Laminating Fabrication and Interface Control Technology, School of Materials Science and Engineering, Hebei University of Technology, 300401, Tianjin, People's Republic of China; Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 110016, Shenyang, People's Republic of China * These authors contribute equally to this work Supplemental data for this article can be accessed online at https://doi.org/10.1080/21663831.2024.2313076.