Journal of Materials Science 【-逻*辑*与-】amp; Technology, 2020, 49(0): 211-223 doi: 10.1016/j.jmst.2020.02.032

Research Article

Phase transformation and structural evolution in a Ti-5at.% Al alloy induced by cold-rolling

Bingqiang Weia, Song Nia, Yong Liua, Xiaozhou Liaob, Min Song,a,*

a State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China

b School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia

Corresponding authors: * E-mail address:msong@csu.edu.cn(M. Song).

Received: 2019-11-25   Revised: 2020-01-2   Accepted: 2020-01-9   Online: 2020-07-15

Abstract

The phase transformation, deformation mechanism and their correlation in a cold-rolled Ti-5at.%Al alloy were investigated. Two types of phase transformations from a hexagonal close-packed (HCP) structure to a face-centered cubic (FCC) structure were observed: the basal-type (B-type) with an orientation relationship of $<1\bar{2}\text{10}{{\text{}}_{\text{HCP}}}<1\bar{1}\text{0}{{\text{}}_{\text{FCC}}}$ and {0001}HCP//{111}FCC, and the prismatic-type (P-type) with an orientation relationship of $<1\bar{2}\text{10}{{\text{}}_{\text{HCP}}}<1\bar{1}\text{0}{{\text{}}_{\text{FCC}}}$ and ${{\text{ }\!\!\{\!\!\text{ 10}\bar{1}\text{0 }\!\!\}\!\!\text{ }}_{\text{HCP}}}\text{// }\!\!\{\!\!\text{ 110}{{\text{ }\!\!\}\!\!\text{ }}_{\text{FCC}}}$. The two types of transformations both accommodate the strain along the < c> axis of the HCP matrix. With the proceeding of deformation, different deformation mechanisms were activated in the FCC and the HCP structures, respectively, which led to a faster grain refinement rate in the FCC structure than in the HCP matrix. Deformation twins with zero macroscopic strain were prevalent in the FCC domains produced by the B-type transformation, while deformation twins with macroscopic strain were normally observed in the FCC domains produced by the P-type transformation. This is in accordance with the lattice mismatches produced during phase transformation. The easy occurrence of deformation twinning in the FCC structure contributed significantly to the grain refinement process. In addition, the interaction between neighboring FCC domains produced by the two types of phase transformations also accelerated the grain refinement process.

Keywords: Ti-Al alloy ; Deformation mechanism ; Phase transformation ; Grain refinement ; Deformation twinning

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Bingqiang Wei, Song Ni, Yong Liu, Xiaozhou Liao, Min Song. Phase transformation and structural evolution in a Ti-5at.% Al alloy induced by cold-rolling. Journal of Materials Science & Technology[J], 2020, 49(0): 211-223 doi:10.1016/j.jmst.2020.02.032

1. Introduction

Deformation mechanisms of metallic materials have been widely studied, since they have significant effects on the microstructural evolution and mechanical properties [[1], [2], [3], [4], [5]]. Dislocation slip and deformation twinning are two main deformation mechanisms for metallic materials. For face-centered cubic (FCC) metals, dislocation activities including slip are prevalent because of the abundant slip systems. Deformation twinning can be activated in FCC metals with medium or low stacking fault energy, especially at high strain rates and/or low temperatures when the critical resolved shear stress for twinning is lower than that for slip [6]. For hexagonal close-packed (HCP) metals, in addition to dislocation slip (including basal slip, prismatic slip and pyramidal slip [2,7,8]), deformation twinning is an important plastic deformation mode. Because the Burgers vectors of most dislocations in HCP materials do not have a c-axis component, the plastic strain with a c-axis component produced by twinning becomes critical to accommodate plastic deformation in HCP materials [2,3].

In addition to dislocation slip and deformation twinning, deformation-induced phase transformation can also be activated to accommodate strain, which could be an additional deformation mode for metallic materials. Among different kinds of phase transformations, stress-induced HCP to FCC phase transformation has been widely observed in metals and alloys, including Hf [9], Zr [10], Co [11,12], the Co-32 %Ni alloy [13,14], Ti [[15], [16], [17]], Ti-based alloys [[18], [19], [20]], stainless steels [21,22], and InAs nanowires [23]. In our previous investigation [10], the basal-type (B-type) and the prismatic-type (P-type) HCP to FCC phase transformations were observed, leading to two orientation relationships between the the HCP and the FCC phases. In the B-type transformation, the longitudinal boundary between the two phases is parallel to the (0001) basal plane of the HCP phase, while in the P-type transformation, the longitudinal boundary between the two phases is parallel to a $\text{(10}\bar{1}\text{0)}$ prismatic plane of the HCP phase. Different mechanisms were proposed for the two types of phase transformations [10]. Note that the formation of the FCC phase in the HCP matrix affects significantly the subsequent plastic deformation by activating various deformation mechanisms in different phase domains, which eventually affects the mechanical properties of the materials. Previous study on a dual-phase (HCP and FCC) cobalt processed by surface mechanical attrition revealed that different deformation mechanisms were initiated simultaneously and played various roles on the microstructural evolution [11]. However, the role of HCP to FCC phase transformation on the subsequent deformation of the materials and the correlation among various deformation mechanisms have not been clearly clarified.

In this study, the occurrence and correlation among different deformation mechanisms of a cold-rolled Ti-5at.%Al alloy were investigated systematically using transmission electron microscopy (TEM). Two types of HCP to FCC phase transformations occurred during deformation. As the deformation continued, the FCC structure and the HCP matrix exhibited different grain refinement rates, attributed to the different deformation mechanisms that were activated in different structures. The correlation between the different deformation mechanisms and their contributions to plastic deformation were discussed.

2. Experimental

The material used in this study was a commercial Ti-5at.%Al plate, purchased from Zhongnuo New Materials Technologies Inc. (Beijing, China) [24]. The as-received sample before deformation exhibits equiaxed grains with an average grain size of ~20 μm. Small bars with dimensions of 40 mm × 10 mm × 3 mm (L × W×H) were cut from the plate by spark machining. These small bars were then rolled at room temperature multiple times with a thickness reduction of 0.3 mm per pass to obtain the deformed samples with thickness reductions of 20 %, 40 % and 60 %, marked as R20, R40 and R60, respectively. The TEM specimens cut from the transverse plane were firstly grounded to 30 μm, glued onto the copper rings with a diameter of 3 mm and then Ar+ ion-milled to perforation using a Gatan 691 PIPS system with a voltage of 3 kV. For the initial and final milling, the ion incidence angle was set at 7° and 3°, respectively. Conventional TEM and high-resolution TEM (HRTEM) investigations were performed using an FEI Titan G2 60-300 microscope with an image Cs corrector operated at 300 kV.

3. Results and discussion

3.1. Microstructural evolution

3.1.1. Microstructures of the un-deformed material and R20 sample

Fig.1(a) shows a typical bright-field TEM image of the un-deformed material. No lamella and only a few dislocations were observed. A selected area electron diffraction (SAED) pattern inset at the upper right corner of Fig.1(a) indicates a pure HCP structure with the zone axis of $\left[ 1\bar{2}10 \right]$. Fig.1(b) shows a bright-field TEM image of the R20 sample, in which lamellas with widths of several tens of nanometers were observed. Fig.1(c) shows an SAED pattern of the area marked by a red dotted circle in Fig.1(b), indicating the existence of an FCC structure in the lamellas. In addition, twins were also formed inside the FCC structure, which will be discussed in detail later. The orientation relationship (OR) between the HCP matrix and the FCC lamella as presented in Fig.1(c) is ${{\text{ }\!\![\!\!\text{ 1}\bar{2}\text{10 }\!\!]\!\!\text{ }}_{\text{HCP}}}\text{// }\!\![\!\!\text{ 1}\bar{1}\text{0}{{\text{ }\!\!]\!\!\text{ }}_{\text{FCC}}}$ and (0001)HCP//(111)FCC, which was introduced via the B-type HCP to FCC phase transformation [10]. Fig.1(d) presents an HRTEM image viewed along [1$\bar{2}$10]HCP and [1 $\bar{1}$0]FCC, showing the interface between the HCP matrix and an FCC lamella, further confirming the OR. The inset in the upper left corner of Fig.1(d) is an enlarged Fourier filtered HRTEM image of the area outlined by the white square in Fig.1(d), showing the atomic arrangement around the interface of the two phases. It has been proved by previous investigations [9,10,23] that the B-type HCP to FCC transformation was accomplished by a/3<10$\bar{1}$0>HCP partial dislocations gliding on every other basal plane [25,26], resulting in the change of the stacking sequence of close-packed planes from …ABAB… to…ABCABC…

Fig. 1.

Fig. 1.   (a) A bright field TEM image of the un-deformed Ti-5at.%Al alloy, with the inset an SAED pattern under the [1 $\bar{2}$10] zone axis of the HCP matrix. (b) A bright-field TEM image of the R20 sample with FCC lamellas inside. (c) An SAED pattern taken from the area marked by the red dotted circle in (b), showing the HCP to FCC phase transformation with the OR of <1 $\bar{2}$10>HCP/​​​​​​​/<1 $\bar{1}$0>FCC and {0001}HCP//{111}FCC. (d) An HRTEM image of an interface between the HCP and FCC phases, with the inset being a Fourier filtered image taken from the area marked by the square in Fig.1(d).


Except for the B-type phase transformation, the deformation-induced P-type phase transformation was also observed in the R20 sample. Fig.2(a) shows a bright-field TEM image containing a P-type FCC lamella. The OR identified from a corresponding SAED pattern is [1 $\bar{2}$10]HCP//[1 $\bar{1}$0]FCC and (101-0)HCP//(110)FCC or (0001)HCP//(001)FCC, as shown in Fig.2(b). This OR is identical to the [0001]HCP//[001]FCC and (101-0)HCP//(110)FCC reported in previous investigations [10,17], which was observed from different directions for the same type of transformation [27]. Fig.2(c) presents an HRTEM image of the area marked by the red square in Fig.2(a) and the inset is an enlarged Fourier filtered HRTEM image of the area outlined by the white square, showing the atomic arrangement around the interface of the two phases. Previous work elucidated that the P-type transformation is accomplished via pure-shuffle and shear-shuffle mechanisms [10,17].

Fig. 2.

Fig. 2.   (a) A bright-field TEM image of the R20 sample with another type of FCC lamella inside. (b) The SAED pattern taken from the interface of the HCP and FCC phases in (a), showing the HCP to FCC phase transformation with the OR of [0001]HCP//[001]FCC and (10$\bar{1}$0)HCP//(110)FCC. (c) An HRTEM image of the interface of the HCP and FCC phases, with the inset being a Fourier filtered image taken from the area marked by the square. (d) A bright-filed TEM image of a {101-2}<$\bar{1}$011> deformation twin, and the insets in the upper left and lower right corners are an SAED pattern and an HRTEM image of the {101-2}<$\bar{1}$011> twin, respectively.


Compared to the un-deformed sample, abundant dislocations were observed in the HCP matrix of the R20 sample, as shown in both Fig.1(b) and Fig.2(a). Besides, a few deformation twins in the HCP matrix were produced in the R20 sample, as shown in Fig.2(d). The insets in the upper left corner and lower right corner in Fig.2(d) are an SAED pattern and an HRTEM image of the HCP twin, respectively. The twinning plane was {10$\bar{1}$2}, a typical {10$\bar{1}$2}<$\bar{1}$011> tension twin activated by c-axis tension [28]. Therefore, dislocation slip, phase transformation and twinning occurred in the plastic deformation process of the Ti-5at.%Al alloy to accommodate strain. Extensive TEM observations showed that the number of the FCC lamella formed via the P-type phase transformation was less than that of the B-type transformation, indicating that the B-type transformation was preferred. In addition, the morphologies of the FCC phase zones formed by the B-type and the P-type transformation are both of elliptical shape. The aspect ratio of length to width for the B-type FCC phase is relative smaller than that of the P-type one, while the B-type FCC phase is more densely distributed.

3.1.2. Microstructures of the R40 sample

Fig.3(a) shows a bright-field TEM image of the R40 sample. Different structural features were observed in different areas in Fig.3(a), dividing the whole area into three domains, with the domain boundaries marked by dotted white lines. SAED was used for phase identification at different areas. Fig.3(b) shows an SAED pattern taken from area b in Fig.3(a), indicating that this area is of an FCC structure with deformation twins. Fig.3(c) presents a dark-field image taken using the reflection marked by the white arrow in Fig.3(b), showing that the widths of the FCC twins are in the nano-scale range. Fig.3(d) shows an SAED pattern taken from area d in Fig.3(a), from which both HCP and FCC patterns were detected, indicating that this area contained an interface of the two phases. The OR presented in the SAED pattern indicates the occurrence of B-type HCP to FCC phase transformation. The experimental results showed that the FCC area expanded with increasing deformation strain, and a high density of deformation twins formed in the FCC domain. Fig.3(e) presents an SAED pattern obtained from area e in Fig.3(a), which shows an HCP structure with the zone axis of [01$\bar{1}$1]. In addition, the diffuse arc-shaped spots in Fig.3(e) indicate the existence of low-angle grain boundaries. In the HCP domain, dislocations generated by plastic deformation accumulated, interacted and tangled with each other, and rearranged spatially to form parallel microbands [5,29,30]. The boundaries of the microbands are indicated by white arrows in Fig. 3(a) and the widths of these microbands are about hundreds of nanometers. The microbands tended to breakdown into elongated segment structures as the strain increased [5,31,32].

Fig. 3.

Fig. 3.   (a) A bright-field TEM image of the R40 sample. (b) An SAED pattern taken from the FCC area b in (a). (c) A dark-field TEM image using the reflection marked by the white arrow in (b). (d) An SAED pattern taken from the interface between the HCP and FCC area d in (a). (e) An SAED pattern taken from the HCP matrix area e in (a).


Fig.4(a) shows a TEM image of another area in the R40 sample, from which microbands with widths of ~100 - 300 nm were readily observed. Figs.4(b) and (c) show SAED patterns with a zone axis of [1$\bar{2}$13] for the HCP structure, which were taken from areas b and c in Fig.4(a), respectively. The diffuse arc-shaped diffraction spots in Fig.4(c) indicate the existence of small-angle misorientations among the microbands. In addition, the microbands display a tendency to split and breakdown with increasing shear strain, as shown by the white arrows in Fig.4(a). There are two ways for the subdivision of the microbands [5,[31], [32], [33]]: (1) the microband boundaries tended to curve under the localized shear deformation and some laths extruded out locally to form a bulge, which can accelerate the longitudinal splitting of the thin laths; (2) a high density of dislocations gathered together to cut off the microbands, which resulted in the transverse splitting of the microbands and finally led to the breakdown of microbands into elongated segments. Note that the B-type phase transformation was quite popular but the P-type transformation was hardly observed in the R40 sample.

Fig. 4.

Fig. 4.   (a) A bright-field TEM image of the HCP matrix in the R40 sample. (b, c) SAED patterns taken from the areas b and c marked by red dotted circles in (a), respectively.


3.1.3. Microstructures of the R60 sample

When the material was cold rolled to a thickness reduction of 60 %, a large number of FCC domains was observed, as presented in Supplementary Fig.S1. Both the FCC and HCP crystals were further refined. Figs.5(a) and (b) show a bright-field TEM image and a dark-filed TEM image, respectively, of a typical HCP domain in the R60 sample. The diffraction spot used for dark-field imaging is marked by a red circle in the SAED pattern as inset in Fig.5a. Both TEM images and the diffraction pattern indicate that the microbands were further subdivided into elongated segments as the strain increased. Fig.5(c) shows a bright-field TEM image and its corresponding SAED pattern of a typical FCC domain in the R60 sample. The image and the distinctive continuous rings in the inset SAED pattern indicate that the FCC domains were completely refined to nanograins. The distribution of the grain size summarized from 200 grains is presented in Fig.5(d) and the average grain size is estimated to be ~23 nm. The average grain size agrees well with the widths of the nanotwins formed previously in the FCC domains, and it is therefore proposed that the nanograins at this stage were evolved from nanotwins via twin-twin and twin-dislocation interactions [[34], [35], [36], [37]].

Fig. 5.

Fig. 5.   (a) A bright-field TEM image of the R60 sample for the HCP matrix, with the inset being an SAED pattern obtained from the area outlined by a red dotted circle. (b) A dark-field TEM image of (a) using the reflection marked by the red circle in the inset of (a). (c) A bright-field TEM image of the R60 sample for the FCC structure and the inset is a corresponding SAED pattern. (d) The grain size distribution of the FCC structure in the R60 sample.


The above experimental observations showed clearly that the FCC area expanded with increasing deformation strain and at the same time different deformation mechanisms were activated in the HCP and FCC domains. Dislocation slip predominated in the HCP domains, while deformation twinning was preferred in the FCC domains, resulting in different grain refinement rates. The grain refinement process for HCP domains can be described as follows: (1) formation of microbands via the generation, accumulation and tangling of dislocations, (2) subdivision of the microbands to form elongated segments and fine grains. At the same time, a few {10$\bar{1}$2}<$\bar{1}$011> tension twins were observed, which also participated in the microstructural evolution of the HCP matrix. The evolution process of the FCC structure can be described as: (1) formation of FCC lamellas via the HCP to FCC phase transformation, (2) formation of deformation twins, and (3) formation of nanograins via twin-twin and dislocation-twin interactions. Note that the grain refinement process for the FCC structure was faster than that of the HCP matrix, revealing that the HCP to FCC phase transformation accelerated the grain refinement process of the whole material.

3.2. Deformation twins inside the FCC structure

3.2.1. Zero-macroscopic strain deformation twins

Fig.6(a) presents typical twins with a 9R structure observed in the FCC phase of the R20 sample, and the inset is an Fourier filtered image taken from the area marked by the white square in Fig.6(a), where the 9R structure could be readily identified by the repeated stacking with a period of 3 times the interplanar spacing of {111} planes [25,38,39]. It can be seen that the 9R phase presented as a step of a coherent twin boundary (CTB). Previous experiments and molecular simulations [25,38,40] elucidated that the Σ3{112} incoherent twin boundaries (ITBs) were produced by gliding three different partials with a repeatable sequence b2:b1:b3 on successive {111} planes, where b1 is a pure edge partial dislocation with the Burgers vector of 1/6<11$\bar{2}$>. b2 and b3 are mixed-type dislocations with opposite screw components and the Burgers vectors of b2 and b3 are 1/6<$\bar{2}$11> and 1/6<1$\bar{2}$1>, respectively. The sum of the Burgers vectors of the three partial dislocations equals to zero, which could lead to zero net macroscopic strain during a twinning process [26,41]. Therefore, twins formed in this way produce zero net macroscopic strain. In the absence of stress, the initial compact Σ3{112} ITBs can dissociate spontaneously into two tilt boundaries, bounding a 9R phase [42,43], to decrease the elastic energy by emitting Shockley partials on every three {111} planes. This results in a repeated pattern with a period of 3 times the interplanar spacing of {111} planes, as revealed in the inset of Fig.6(a). Under shear stress, the partial b1 would move to one side, and the paired partials b2 and b3 to the other side. The partial b1 is initially more mobile than the other two partials (b2 and b3) [40], giving rise to a non-equilibrium state of the Σ3{112} ITBs and the propagation of the 9R structure. The resulting stacking faults and the interaction between the three partials will pull the b2, b3 partials to approach b1 until reaching another equilibrium state, which eventually promote the migration of the ITBs and subsequently the growth of the deformation twin [25,40]. The deformation process contains a repeated process between a non-equilibrium state and the equilibrium state, and the twinning via migration of the ITBs is an important twinning process [40,44]. Fig.6(b) shows an HRTEM image of another 9R structure which is under propagation. The inset is a fast Fourier transformation (FFT) pattern taken from the area marked by a square in Fig.6(b) and the extra diffraction spots correspond to a periodicity of three times the interplanar of {111} planes, indicating the formation of the 9R structure. It is seen that many dislocations reside in the step of the 9R structures and CTBs, which played a role in destroying the coherency of CTBs and impeding the propagation of the ITB [45,46]. Moreover, there was a pinning effect when the dislocations interact with the ITBs, which could strengthen the materials since a higher stress is needed to reactive the motion of the ITB [47]. A longer 9R after propagation was presented in the Supplementary Fig.S2. Two more HRTEM images of FCC twins with ITBs inside are shown in Supplementary Fig.S3.

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Fig. 6.   (a) An HRTEM image of zero macroscopic strain twins with the 9R structure. The inset is a Fourier filtered image taken from the area marked by the square. (b) Another HRTEM image of the 9R structure, with the inset being an FFT pattern taken from the area marked by the square. (c) An HRTEM image observed in a large area of the FCC structure in the R40 sample, and (d) an enlarged HRTEM image corresponding to the area marked by the square in (c).


Fig.6(c) presents another HRTEM image of a large FCC area in the R40 sample. This area can be divided into two domains based on the microstructural features, the boundary of which was marked by white dotted lines. The {111} planes within and outside the dotted area were indicated by yellow and green solid lines, respectively. There exists a misorientation angle of 55° between these two domains, which agrees well with the angle difference between two FCC phase domains formed via B-type phase transformation and P-type phase transformation, respectively, in a crystal. The 9R structure was observed in both domains, as indicated by white arrows. Note that most of the 9R structure within the dotted area can extend across the whole grain. The white squared area is further magnified in Fig.6(d), in which a 9R phase terminated in the grain was observed (marked by the green dotted ellipse), indicating a state of the 9R structure under propagation. Ma et al. [48] showed that the ITBs can readily nucleate at the boundary of a nano-scaled FCC grain and split to produce a 9R structure across the whole grain. It is therefore likely that the 9R structure in Fig.6(c) originated from a grain boundary, i.e., the partials forming the 9R structure were emitted from the grain boundary. In addition, the smooth grain boundary with little kinks in Fig.6(c) implies the existence of net zero macroscopic strain. Wu et al. [26] proposed that the zero-strain twins could make it easy for grains to rotate and slide during further deformation, which promotes the grain refinement process. It is worth mentioning that twins with zero macrostrain were more prevalent in the FCC domains formed via the B-type phase transformation than in the FCC domains formed via the P-type phase transformation. The reason is discussed below.

Fig.7(a) shows schematics of the b1, b2, b3 partials for HCP and FCC structures viewed along the [0001]HCP and [111]FCC directions. The formation of 9R structure in the FCC domains consists of three steps as presented by a schematic diagram in Fig.7(b): (1) HCP to FCC phase transformation, (2) nucleation of Σ3{112} ITBs at the HCP-FCC phase boundary and (3) split of the Σ3{112} ITBs under external stress. Considering the source of the b1, b2 and b3 partials, the nucleation of the ITBs might be closely related to the B type phase transformation process. It is known that the B type phase transformation was accomplished by three different a/3<10$\bar{1}$0> partial dislocations (b1, b2 and b3) gliding on the basal plane with the angle of 120° between each other, which also produced a near zero net macroscopic strain (presented in Supplementary Fig.S4) [9,10,23]. The b1, b2 and b3 partials that produce Σ3{112} ITBs in the FCC phase present the same direction to those dominate the phase transformation in the HCP matrix, as shown in Fig.7(a). Therefore, the Σ3{112} ITBs could be readily nucleated at the phase boundary after the B-type phase transformation. Furthermore, the interface between the HCP and the FCC structures was presented in a non-equilibrium state as the HCP matrix would progressively transform into the FCC structure during plastic deformation, and the moving interface between the two phases could thus be activated as the source of partials for the nucleation of Σ3{112} ITBs. Previous study [44] proposed that the Σ3{112} ITBs could nucleate via the synchronized emission of three (b1:b2:b3) partials as “zonal” twinning dislocations from a grain boundary (here, the phase interface) under high stress and a proper stress orientation, which could be readily provided by the cold rolling process. In addition, the ITBs can also nucleate at the boundary of the nano-scaled FCC structure, as shown in Fig.6(c) and previous work [48]. This opens up more possibilities for the breakdown of ITB equilibrium when the grain size is reduced to the nano scale, and forms the 9R structure across the whole grain (Fig.6c).

Fig. 7.

Fig. 7.   Schematics for the formation of ITBs and the 9R structure. (a) Schematics of the b1, b2, b3 partials for the HCP and FCC structures viewed from [0001]HCP and [111]FCC directions, respectively. (b) Schematic for the HCP to FCC transformation and formation of ITBs and the 9R structures under the action of b1, b2 and b3 partials.


3.2.2. Deformation twins with macroscopic strain

Deformation twins with macroscopic strain were also observed, normally in FCC lamellas formed via the P-type phase transformation. Fig.8(a) shows a bright-field TEM image of an area containing an FCC lamella formed via the P-type phase transformation in the R20 sample. It is seen that several deformation twins formed in the lamella, which were indicated by white arrows. Fig.8(b) shows an HRTEM image of the area outlined by a white square in Fig.8a, from which the orientation relationship for the P-type phase transformation and the deformation twins with CTBs is readily identified. This type of twinning usually occurred via the gliding of the same twinning partial dislocations on successive {111} planes, which produced a net macroscopic strain and sheared the grain into a new shape (often called a monotonic twinning process [41]). The morphology of the phase boundary between the HCP and the FCC structures was thus changed by twinning induced shear strain, as presented by the yellow dotted lines in Fig.8(b) and the shear direction of the twins was indicated by the white arrows. It is obviously seen that twinning partials were emitted from HCP/FCC phase boundaries, which acted as a similar role to grain boundaries in nanocrystalline materials since the width of the FCC lamella is about tens of nanometers. Fig.8(c) shows another HRTEM image with a P-type FCC lamella inside, and deformation twins were produced in the left part of the lamella in the R20 sample. The phase boundary was outlined by the yellow dotted lines and its direction was also changed by the shear strain produced by the twinning process. There was a grain boundary at the lower left part of the FCC lamella, as marked by a white dotted line. The 55° misorientation angle between the two FCC domains indicates that the FCC domain at the lower left part formed via the B-type HCP to FCC phase transformation. Both phase boundaries and grain boundaries could act as sources for twinning partials [6,49]. At the right side of the lamella with a relatively large width, no twin was observed and some misfit dislocations were detected, as shown in Fig.8(c). Experimental results showed a clear grain size effect on the propensity of deformation twinning in the FCC phase, i.e., twins were hardly seen in relatively large FCC areas, which is consistent with the literature reports on grain size effect on deformation twinning [41,50]. From the above observations, it can be seen that partial dislocations emitted from phase boundaries play an important role in the nucleation and growth of deformation twins.

Fig. 8.

Fig. 8.   (a) A bright-field TEM image of the P-type phase transformation and (b) an HRTEM image taken from the area marked by the square in (a). (c) An HRTEM image of deformation twins with macroscopic strain observed in the FCC structure corresponding to the P-type phase transformation.


3.3. Lattice misfit in phase boundaries

The lattice misfits on the HCP-FCC interfaces formed by the two types of HCP to FCC phase transformations were analyzed from HRTEM images. Figs.9(a) and (b) show schematic diagrams for the two types of phase transformations. The interplanar spacing with error bars was measured from the HRTEM images. For the HCP matrix the lattice parameters obtained from the interplanar spacing are: aHCP = 3.03 ± 0.06 Å, cHCP = 4.79 ± 0.06 Å, and for the FCC structure it is aFCC = 4.36 ± 0.06 Å. The parameters are very close to that obtained by molecular dynamics simulations [51] and XRD experimental results [52] in Ti.

Fig. 9.

Fig. 9.   (a,b) Atomic schematics of the lattice parameters and ORs for the two types of HCP to FCC phase transformations. (c,d) Fourier filtered HRTEM images of the interfaces of the HCP and the FCC phases for two types of phase transformations.


For the B-type transformation, as shown in Fig.9(a), lattice expansions of +1.3 % and +5.0 % exist along the [10$\bar{1}$0] and [0001] directions, respectively, indicating that lattice expansion is mainly along the < c> axis of the HCP matrix. Previous work showed that tension along the < c> axis not only triggered the B-type phase transformation [9], but also stabilize the FCC phase under a hydrostatic expansion stress [53,54]. Therefore the lattice expansion can effectively accommodate the tensile-strain along the < c> axis, which is comparable to the {10$\bar{1}$2}<$\bar{1}$011> tension twins in the HCP structure. For the P-type phase transformation, as shown in Fig.9(b), there exists a large lattice expansion of +17.2 % along the [10$\bar{1}$0] direction and a lattice compression of -8.7 % along the [0001] direction. This suggests that the HCP is subjected a tensile stress along the [10$\bar{1}$0] axis and a compressive stress along the [0001] direction, when the P-type phase transformation occurs. The stress is consistent with the shear direction of the deformation twin in the FCC area formed via the P-type phase transformation as shown in Figs.8(b) and (c), indicating that both the phase transformation and deformation twinning are important ways to accommodate deformation during cold-rolling. In addition, the compression twinning modes of {10$\bar{1}$1}<10$\bar{1}$$\bar{2}$> and {11$\bar{2}$2}<11$\bar{2}$$\bar{3}$> in the HCP structure can normally be activated under an effective compressive stress along the c-axis, resulting a compressive strain normal to the basal plane and a tensile strain normal to the prismatic plane [17,55]. This presents the same stress state with the P-type phase transformation, and thus implies that the P-type phase transformation can accommodate macrostrain comparable to compression twins. The misfit strain resulted from the lattice misfit was produced during phase transformation, which led to the formation of misfit dislocations at phase boundaries for both the B-type and the P-type phase transformations, as shown in Figs.9(c) and (d), respectively. The Burgers vector of the misfit dislocations for the B-type phase transformation was b= <11$\bar{2}$0>HCP, which presented along the [10$\bar{1}$0]HCP direction on the (0001)HCP plane. The Burgers vector of the misfit dislocations for the P-type phase transformation was b= a/2<01$\bar{1}$>FCC, which presented along the [11$\bar{2}$]FCC direction on the (111)FCC plane. Therefore, there existed two components of the Burgers vector in the [001]FCC and [110]FCC directions, which released the lattice misfit strain along the [001]FCC//[0001]HCP and [110]FCC// [10$\bar{1}$0]HCP, respectively. From the above analysis, it is concluded that the misfit strain in the P-type phase transformation is larger than that of the B-type counterpart, implying that the driving force for the P-type phase transformation is also larger than that for the B-type one. This explains our observations that the number of the P-type transformation was much less than that of the B-type transformation.

3.4. Interaction between two types of phase transformation

Fig.10(a) presents a bright-field TEM image taken from the R20 sample, in which some triangular bulges were clearly observed, as indicated by white arrows. The inset in Fig.10(a) is an SAED pattern taken from the bulge marked by the white square, showing the co-existence of the B-type and P-type phase transformations. The orientation relationships between the HCP phase and FCC phase in the two types of phase transformations are: ${{\text{ }\!\![\!\!\text{ 1}\bar{2}\text{10 }\!\!]\!\!\text{ }}_{\text{HCP}}}\text{// }\!\![\!\!\text{ 1}\bar{1}\text{0}{{\text{ }\!\!]\!\!\text{ }}_{\text{B-FCC}}}\text{// }\!\![\!\!\text{ 1}\bar{1}\text{0}{{\text{ }\!\!]\!\!\text{ }}_{\text{P-FCC}}}$, (0001)HCP//(111)B-FCC//(001)P-FCC, and there is a misorientation angle of 55° around the [1$\bar{1}$0]FCC between two neighboring FCC domains formed by the two types of phase transformations. Fig.10(b) shows an HRTEM image of part of the bulged area indicated by the white square in Fig.10(a). The {111} planes for the FCC structure induced by the P-type and the B-type phase transformation were indicated by yellow and green solid lines, respectively. The misorientation between them is 55°, which agrees well with the theoretical angle difference between FCC areas formed via the B-type and the P-type phase transformations in a crystal. Interestingly, the bulge with the FCC structure formed via the B-type phase transformation “grows” into the FCC lamella formed via the P-type phase transformation, implying that there was interaction between the two types of phase transformation and the B-type one could consume the P-type one. Fig.10(c) presents an HRTEM image of another area in the R20 sample where both the P-type and B-type phase transformations were observed. It’s interesting that the FCC structure induced by the P-type phase transformation (as marked by dotted lines) was surrounded by the HCP structure and the FCC structure induced by the B-type phase transformation. Fig.10(d) shows an HRTEM image of an FCC area in the R40 sample. Based on the structural features, the area can be divided into three domains, with the domain boundaries marked by dotted lines. By measuring the misorientation angles among them, the domain (FCCP 1) in the middle part of the image was with the FCC structure evolved from the P-type phase transformation, while the other two domains (FCCB 2 and FCCB 3) were with the FCC structure evolved from the B-type phase transformation. Moreover, the 9R structure was produced in the middle FCC domain as marked by the white square in Fig.10(d) and the relatively smooth grain boundary implies the existence of net zero macroscopic strain.

Fig. 10.

Fig. 10.   (a) A bright-field TEM image containing both the B-type and the P-type phase transformations, with the inset being an SAED pattern taken from the area outlined by the white box. (b) An HRTEM image corresponding to the area outlined by the box in (a). (c) Another HRTEM image containing the two types of phase transformations. (d) Another HRTEM image in a large area of the FCC structure in the R40 sample.


It should be noted that it is not just coincidence for observing areas that contain both the two types of phase transformations. According to the orientation relationship of the two types of transformations observed experimentally, a 3-dimentional structure of the material can be built, and then the crystallographic projection and the diffraction patterns under different zone axes can be deduced [27]. Fig.11(a) shows a schematic diagram of the diffraction pattern for the B-type transformation. Figs.11(b) and (c) show schematic diagrams for diffraction patterns of the HCP and FCC phases viewed along two other directions. Fig.12(a) shows a schematic diagram of the diffraction pattern for the P-type transformation. Figs.12(b) and (c) show schematic diagrams for diffraction patterns of the HCP and FCC phases viewed along two other directions. All the orientation relationships for the B-type and P-type phase transformation were summarized in Supplementary Tables S1 and S2, respectively. It can be seen that the two types of phase transformation can be seen simultaneously in [1$\bar{2}$10]HCP//[1$\bar{1}$0]FCC HRTEM images. In addition, the longitudinal boundaries for the two types of transformation were perpendicular to each other along the [1$\bar{2}$10]HCP//[1$\bar{1}$0]FCC zone axis, implying that the two types of FCC lamellas would meet during the deformation process, as presented in Supplementary Fig.S5. Figs.10(a) and (b) show that the B-type transformation was predominant during the interaction of the two types of transformation since it could “grow” into the P-type one. If there pre-exists a P-type FCC lamella in the HCP matrix, the B-type FCC lamella may propagate and transmit across the P-type one during the subsequent deformation process, separating the P-type FCC lamella into several blocks via interaction between domains formed by the two types of phase transformation. As the B-type FCC structure expanded, the P-type FCC structure was embedded as small blocks in the large area of a B-type FCC domain, as shown in Fig.6(c) and Fig.10(d). That is one of the reasons why the P-type transformation was hardly observed in the R40 sample.

Fig. 11.

Fig. 11.   (a) A schematic of the experimental observed OR for the B-type phase trans-formation. (b,c) Schematics of the OR viewed along two other directions of the B typephase transformation.


Fig. 12.

Fig. 12.   (a) A schematic of the experimental observed OR for the P-type phase trans-formation. (b,c) Schematics of the OR viewed along two other directions of the P typephase transformation.


Fig.13(a) shows an HRTEM image of another area with the FCC structure in the R40 sample. Areas 1, 2 and 3 marked in Fig.13(a) were further magnified in Figs.13(b) - (d), respectively. These areas contained FCC domains evolved from both B-type and P-type phase transformations. The domain boundaries were indicated by white dotted lines. In addition, deformation twins were observed in the FCC domain formed by the B-type phase transformation. The {111} planes that formed the twin relationship were indicated by white sold lines in Fig.13. Interestingly, the twins stopped and deflected by the FCC domains evolved from the P-type phase transformation, as indicated by the white arrows. It’s therefore proposed that there is an interaction between P-type and B-type phase transformation-induced FCC crystals, which effectively promotes the grain refinement process of the FCC structure during the deformation.

Fig. 13.

Fig. 13.   (a) An HRTEM image observed in a large area of the FCC structure. (b,c,d) are enlarged HRTEM images corresponding to the areas indicated by the numbers 1, 2, and 3, respectively.


Three different deformation modes, namely, deformation-induced HCP to FCC phase transformation, dislocation slip and deformation twinning occurred in the cold-rolled Ti-5at.%Al alloy to accommodate strain and eventually led to grain refinement. The occurrence and correlation among different mechanisms were summarized as follows. Firstly, two types of deformation induced HCP to FCC transformation (B-type and P-type) occurred, both accommodated macroscopic strain along the < c> axis of the HCP matrix. Analysis of the lattice mismatch between the two phases indicates that the B-type transformation can accommodate about 5% tensile-strain along the < c> axis. The P-type transformation can accommodate about 8.7 % compressive-strain along the < c> axis. Secondly, the HCP to FCC transformation continued progressively with deformation strain, which accelerated the grain refinement process of the material. The FCC structure produced by phase transformation divided the HCP matrix into several refined platelets, as shown in Supplementary Fig.S6. The expansion of the FCC area with increasing deformation strain is in agreement with the previous study on Hf [56] in that the HCP Hf powder could be totally transformed into the FCC Hf under high-energy ball milling for a long time. Thirdly, different deformation mechanisms were activated in the HCP and FCC structures, and the FCC structured domains exhibited a faster rate of grain refinement than the HCP domains. Dislocation slip predominated in the HCP domain while deformation twinning was preferred in the FCC domain. The HCP to FCC phase transformation promoted twinning of the material since twinning can easily occur in the FCC structure, which played an important role in accelerating the grain refinement process of the FCC structure. The zero macroscopic strain twins were prevalent in the FCC domains formed via the B-type phase transformation, which is related to the zero macroscopic strain HCP to FCC transformation mechanism. In addition, the twins with zero macroscopic strain can also be produced by emitting partials from the interface between the FCC domains formed via the B-type and the P-type phase transformations (Figs.6c and 10 d). The deformation twins with macroscopic strain were normally observed in the P-type ones, which could be ascribed to the relatively larger lattice misfit in the P-type transformation. Fourthly, the interaction between the two types of phase transformations resulted in the P-type FCC structure being embedded as small blocks in the large areas of the B-type FCC domains. The interface between the two phases could thus be activated as the source of partials for the nucleation of twinning dislocations, which also accelerated the grain refinement process.

It should be noted that the addition of the Al element also affects the activation of the deformation mechanism since Al effectively reduces the stacking fault energy (SFE) of the (0001) basal plane of Ti. The SFE for Ti-5at.%Al is about 204 mJ/m2, which is much lower than that in pure Ti (300 mJ/m2) [57,58]. Basal stacking faults are very common in the Ti-5at.%Al alloy, as presented in Supplementary Fig.S7. Lower SFE in the Ti-5at.%Al alloy indicates that the nucleation and propagation of the B-type HCP to FCC transformation (through partial gliding on every other basal planes) is easier in the alloy than in pure Ti. In addition, the B-type transformation is more widely observed in TEM, indicating that the B-type transformation should be easier than the P-type transformation in the Ti-5at.%Al alloy.

Thermodynamically, Ti has two kinds of stable crystal structures: HCP (α phase at low temperatures) and body centered cubic (BCC, the β phase at high temperatures). The FCC phase is an unstable one. However, the FCC structure can be produced under severe plastic deformation to release high stress concentration. The transformation process was accomplished instantaneously due to the high-speed motion of partial dislocations. Once the FCC phase was produced, the elastic constraint produced by the lattice misfit between the FCC phase and the HCP matrix helps stabilize the FCC structure. Theoretically, the BCC structure can also be produced due to the local temperature raising during a severe deformation process in pure Ti, Zr and Hf. However, it is difficult to form the BCC structure by cold deformation in the Ti-5at.%Al alloy since Al is a strong stabilizer of the α phase in Ti alloys [59]. The HCP to FCC transformation is mainly activated by shear stress, while the transformation between HCP and BCC is mainly based on thermal activation. The mechanisms for the two types of transformations should be the same for Ti [17,27], Zr [10] and Hf [9]. While, the probability of generation and size of the two types of FCC phases may be different considering the different c/a ratios and physical properties of Ti, Zr and Hf.

4. Conclusions

The occurrence and correlation among different deformation mechanisms in a Ti-5at.%Al alloy during cold-rolling were investigated systematically. The following conclusions are drawn:

(1)Two types of HCP to FCC phase transformations occurred during cold rolling: the B-type transformation with an orientation relationship of ${{\text{ }\!\![\!\!\text{ 1}\bar{2}\text{10 }\!\!]\!\!\text{ }}_{\text{HCP}}}\text{// }\!\![\!\!\text{ 1}\bar{1}\text{0}{{\text{ }\!\!]\!\!\text{ }}_{\text{FCC}}}$and (0001)HCP//(111)FCC, and the P-type transformation with an orientation relationship of ${{\text{ }\!\![\!\!\text{ 1}\bar{2}\text{10 }\!\!]\!\!\text{ }}_{\text{HCP}}}\text{// }\!\![\!\!\text{ 1}\bar{1}\text{0}{{\text{ }\!\!]\!\!\text{ }}_{\text{FCC}}}$, ${{\text{(10}\bar{1}\text{0)}}_{\text{HCP}}}\text{//(110}{{\text{)}}_{\text{FCC}}}$ and (0001)HCP//(001)FCC. The HCP matrix transformed progressively into the FCC structure with increasing deformation strain.

(2)Different deformation mechanisms were activated in the HCP and FCC structures as the deformation proceeded. The FCC domains exhibited a faster grain refinement rate and smaller final grain sizes than the HCP domains. Dislocation slip predominated in the deformation of the HCP domains while deformation twins were preferred in the FCC domains. Twins with zero macroscopic strain were prevalent in the FCC domains produced by the B-type phase transformation, while deformation twins with macroscopic strain were normally observed in the FCC domains formed by the P-type transformation. Besides, twins with zero macroscopic strain can also be produced by partial dislocation emissions from the interface of two FCC domains formed via the B-type and the P-type phase transformations, respectively. Deformation twinning played an important role in accelerating the grain refinement process of the FCC structure.

(3)The interaction between the two types of phase transformations resulted in a microstructure in which P-type FCC domains were embedded as small blocks in B-type FCC domains. The interface between the two domains could act as sources of partials for twinning, which also accelerated the grain refinement process.

Acknowledgments

The financial supports from National Natural Science Foundation of China (51828102) and Australian Research Council (DP190102243) are appreciated.

Appendix A. Supplementary data

Supplementary material related to this article can be found, inthe online version, at doi:https://doi.org/10.1016/j.jmst.2020.02.032.

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