The diffusion bonding method is one of the most essential manufacturing technologies for Ti-steel composite plates. In this paper, the atomic diffusion behavior at the Fe-Ti interface during the bonding process of Ti-steel composite plates is studied using classical diffusion theory and molecular dynamics (MD) simulation. Henceforth, the diffusion mechanism of Fe and Ti atoms at the bonding interface is obtained at the atomic scale. The results show that Fe and Ti atoms diffused deeply into each other during the diffusion process. This behavior consequently increased the thickness of the diffusion layer. Moreover, the diffusion quantity of Fe atoms to the Ti side was much greater than that of Ti atoms to the Fe side. Large plastic deformation and shear strain occurred at the diffusion interface during diffusion. The crystal structure of the diffusion zone was damaged and defects were generated, which was beneficial to the diffusion behavior of the interface atoms. As the diffusion time and temperature increased, the shear strain of the atoms at the interface also increased. Furthermore, there is a relationship between the mutual diffusion coefficient and the temperature. Subsequently, after the diffusion temperature was raised, the mutual diffusion coefficient and atomic disorder (Fe atom and Ti atom) increased accordingly.

The law of microstructure evolution and transformation mechanism of the α′ martensite decomposition during 400–600 °C were studied by the isothermal dilatometry. The transformation process of α′ martensite was quantitatively characterized based on Johnson–Mehl–Avrami (JMA) model, and the microstructure evolution under different aging processes was observed and compared on Scanning Electron Microscopy (SEM) and Transmission Electron Microscopy (TEM). The results showed that α′ → α + β is the elemental diffusion transformation, the position and shape of the precipitate gradually change with the holding time and temperature. The decomposition rate of α′ martensite was positively correlated with the aging temperature. The whole transformation process was divided into two stages based on the value of the Avrami exponent n, and the corresponding average values of the transformation activation energies Q are 46.1 kJ/mol and 116.8 kJ/mol, respectively. The calculated model had good agreement with the experimental data, and the transformation curve of α′ martensite with time and temperature during the isothermal aging at 400–600 °C was drawn.

In order to understand the non-isothermal transformation behavior of Ti-6Al-4V titanium alloy in the continuous heating stage of solution treatment, thermal dilatometry tests with heating rates of 0.1~0.8 °C/s were designed. The conversion between the expansion amount and the transformed volume fraction was realized by the lever principle, and the transformation characteristics of α + β → β were quantified based on the Kissinger-Akahira-Sunose (KAS) theory and the modified Johnson–Mehl–Avrami (JMA) model. The results show that the phase transformation kinetics curves present typical “S” patterns, and the element diffusion transformation controls the nucleation and growth of new grains during the transformation of α + β → β. The phase transformation interval gradually moves to high temperature regions with the increase of heating rates, and the phase transformation activation energy is 445.5 kJ·mol⁻¹. The phase transformation process is divided into three stages according to the relationship between the Avrami exponent n and the transformed volume fraction f_T. These three stages correspond to different stages of grain nucleation or growth.

This work combined theoretical calculation with experimental characterization to methodically study the anisotropy mechanism and evolution of the plastic behavior of pure titanium. Initially, a constant-strain uniaxial tensile test was used to measure the anisotropy of the yield behavior along the rolling direction (RD) and transverse direction (TD). Subsequently, the information of crystal orientation both before and after deformation was statistically characterized using electron backscatter diffraction (EBSD). Ultimately, the main deformation mechanism was determined by combining Schmid law with an analysis of the variation of SF values of each deformation mode with the angular relationship between the loading axis and the grain’s c-axis. The findings demonstrate that, for each slip system, the variation trend and value of the SF are influenced by the angle formed by the loading axis and the grain’s c- and a-axes. The primary result of dislocation slip activation is the change of the tilt angle of the grain c-axis from ND to TD, but this has little effect on the tilt angle of the grain c-axis from ND to RD. Prismatic slip dominates the tensile deformation along the RD. Pyramidal slip and pyramidal slip will be activated during the subsequent hardening, whereas basal slip is difficult to activate. The prismatic slip in the soft-oriented grain will be preferentially activated during the tensile deformation along the TD, and the prismatic slip and pyramidal slip will become the dominant deformation modes during the subsequent hardening. Some soft-oriented grains could activate basal slip and pyramidal slip, but dislocation slip is restricted and coordinated by {10-12}ET.