Microstructure of dual-phase TiMo alloy
The typical microstructure and atomic structure of the dual-phase TiMo alloy are presented in Fig. 1(a)–(c). The two phases had distinct contrasts and different shapes. The atomic structures of the two phases shown in the high-angle annular dark-field (HAADF) image illustrated that the dark plates were associated with the HCP structure and the bright plates with the BCC structure, corresponding to the α phase and β phase, respectively. The clear atomic arrangement was acquired from zone axis \(\left[\stackrel{-}{1}11\right]\) of β and \(\left[2\stackrel{-}{1}\stackrel{-}{1}0\right]\)of α at a 2 nm scale. Furthermore, the two phases were distributed homogeneously and interlaced. Specifically, the α phase was in the form of a submicron-sized plate and partitioned the bright β phase into comparative-size plates and blocks. Notably, the plate-shaped phase exhibited a short side with a width of about 80 nm and a longitudinal direction with a length of up to several micrometers. As such, α-β interfaces and α-α intersections were considered the main barriers to the transportation of dislocation plasticity in the dual-phase TiMo alloys. The atomic structures of α-α intersections and α-β interfaces were characterized and displayed in Fig. 1(d) and 1(e), respectively. It was shown that two α grains at the α-α interface were symmetric and well fitted without any stress concentration. Meanwhile, the α-β phase boundary demonstrated coherent character with a classical Burgers orientation relationship of <\(11\stackrel{-}{2}0\)>α//<\(1\stackrel{-}{1}1\)>β (Fig. 1(e))[14]. A dislocation with a Burgers vector of 1/2<110> (marked by “⊥”) appeared close to the boundary instead of lying at the boundary. It was noted that the coherent interface between nanoparticles and matrix with low mismatch and little lattice distortion could effectively facilitate the release of stress concentrations, generating considerable plastic strain[15]. The high coherency of α-β interfaces and α-α intersections implied limited resistance for dislocation transferring.
Strain transportation in dual-phase architecture
We performed in situ TEM straining tests to observe the dynamic dislocation behaviors within the two phases and directly revealed the transportation of dislocation plasticity within this dual-phase structure. According to the observation over a wide range of strains, it was determined that dislocations were preferentially activated in the α phase during the early stage of deformation, as shown in the serial images captured from Supplementary Movie 1 in Fig. 2. These pictures were acquired at the [\(11\stackrel{-}{2}0\)] zone axis with the [\(01\stackrel{-}{1}1\)] g vector. The fast movement of dislocations in α plates dominated the deformation process at the initial stage, which was attributed to the densely packed atoms in the α phase[16] and the lower critical shear stress for dislocation slip in the HCP structure[12]. Interestingly, the dislocations activated in the α phase tended to slip along the longitudinal direction of α plates in most cases. As shown in Fig. 2(a), dislocations glided along the longitudinal directions in both α plates, although the two distinct plates oriented orthogonally. In the horizontal plate, dislocations (marked with red dashed lines) moved from left to right. Moreover, the dislocations (marked by dashed lines with other colors) moved from top to bottom in the vertical plate. Overall, the dislocations in the α phase glided along the longitudinal direction of the plates. As α plates oriented in multiple directions and intersected with other α plates, dislocations activities occurred in multiple directions, resulting in homogeneous deformation comparative to only one or limited preferential slips. Generally, the Schmid factor determines which grain activates the dislocation first and which slip system is activated[17]. In this case, the dislocations preferentially glided along each plate’s longitudinal direction, especially in the α phase. It was also noted that the width and length of each alpha plate reside at about 80 nm and several micrometers, respectively. The shape characteristics could result in anisotropic properties in the two directions, including Young’s Modulus, yield strain, and tensile strength[18]. Therefore, it was expected that the anisotropic shape could result in an anisotropic Hall-Petch effect, affecting the movement of dislocations in the α phase. And the result was consistent with our observation.
By making close observation of the origin of dislocations in the α phase, it was found that the dislocation activities were generated from both α-α junctions and α-β interfaces, and they glided along the longitudinal direction of α plates. For the α-α junction, most dislocations glided along the longitudinal direction of the α strip (Fig. 3(a) and Supplementary Movie 2). The dislocations first piled up at the junction (marked by an orange arrow in Fig. 3(a)), causing stress concentrations. The Burgers vector of the dislocations in the α phase was 1/3<\(11\stackrel{-}{2}0\)>, and the slip plane was {0001} as characterized by Burgers circuit analysis, as shown in Fig. 3(b). It was also noted that a few dislocations activated from α-α junctions glided in multiple directions (Fig. 3(c)). These dislocations moved toward the α-β phase boundary and were impeded by the phase boundary. As the dislocations interacted with the α-β phase boundary, dislocation activities were initiated in the neighboring β phase (pointed out by red arrows). In this case, the activation of dislocations in the relatively hard β phase might result from the coherency of the α-β phase boundary, which could provide channels for dislocations to slip through. Consequently, both phases were plastically deformed and contributed to homogeneous deformation. For the α-β phase boundary, as the applied stress increased, it was observed that dense dislocations were located at the α-β phase boundary and could be de-pinned under the applied stress, as marked by the red dashed line in Fig. 3(d). For instance, from t + 8s to t+14s, the high density of the newly excited dislocations generated from the α-β phase boundary moved along α plates (indicated by the red arrow). These pictures were acquired at the [\(11\stackrel{-}{2}0\)] zone axis with the [\(01\stackrel{-}{1}0\)] g vector. The systematic dislocation analysis demonstrated that the type of mobile dislocations in vertical α plate was <a> type and the slip planes were on prismatic planes. Post-mortem characterizations of the dislocation structure of the deformed sample (shown in Fig. 3(e)) also illustrated that the dislocations regularly lined up along the α stripe at low strain. The picture was acquired at the [0001] zone axis with [0–110] g vector under the TEM model. The dislocation plasticity primarily transports in the α phase along the longitudinal direction, regardless of where the dislocation originated.
Micropillar compression test
It has been well-accepted that strain transportation within the dual-phase structure is critical in accommodating deformation and strengthening. It was considered that the dislocation activities transmitted through the α-β phase boundary in the dual-phase alloy since it was the key interface with the largest area. While in this case, it was found that α-α junctions played an important role in transporting dislocation plasticity at the early stage of plastic deformation. The longitudinal direction of the α phase could be regarded as the main channel for dislocation slip, and the transportation of deformation was accomplished by the dislocation transfer between various α plates through the α-α junctions. Therefore, it was speculated that the validity of α-α junctions should be a significant factor affecting the mechanical properties of materials. To quantitatively analyze the contribution of such structures to the mechanical properties, micropillars with a diameter of 0.5, 1, and 3 µm were prepared and compressed (see Methods). Four engineering stress-strain curves for each sized micropillar from the in situ SEM compression tests were plotted in Fig. 4(a). The steady serrated curves signified that all dual-phase TiMo pillars were deformed continuously without strain burst, indicating superior plastic stability. In addition, all pillars displayed uniform deformation since all pillars deformed into a bulged shape after compression (Fig. 4(b)). The high deformation stability of the dual-phase pillars resulted from the architecture of the two phases. Interestingly, the different-sized pillars displayed yield strengths in a similar range. However, data was more scattered at smaller sizes, distinct from the traditional size effect of single-phase metals, i.e., demonstrating that the reduction of sizes dramatically enhances strength[19, 20]. The 3 µm pillars exhibited about equal yield strength at 1050 MPa, and the 1 µm pillars displayed variations in yield strength from 1000 to 1200 MPa, while the yield strength of 0.5 µm pillars fluctuated from 900 to 1300 MPa. Note that the dislocations were first activated in the α phase during the plastic deformation of the dual-phase alloy. The size of the α plates, especially the length in the longitudinal direction, determined the yield strength. Using FIB to extract cross-sections from the deformed pillars, it was observed that the length in the longitudinal direction of α plates and the number of α-α intersections contained in the micropillars varied specifically at the smallest tested size (Fig. 4(c)). In large pillars, the number of α plates was significant. Therefore, the mechanical data was closer to the statistical behavior. However, as the sample size decreased, although the yielding of the pillar was still related to the dislocation activities within α plates, the mechanical data varied since the distinct shape and orientation of each α plate and how different α plates intersected had a great influence on the transportation of dislocation plasticity. Thus, the yield strength in the smaller pillar showed considerable fluctuation.