Although a great deal of effort has been put into forming metal matrix composites (MMCs) using laser additive manufacturing techniques, the behavior of metal matrices and reinforcements during heating and cooling remains unclear. Most studies rely on thermodynamic models to implicitly elucidate phase evolution and microstructure formation, and often overlook the dynamics involved.
Most studies on 316L+TiC composites have focused on the effects of processing parameters and reinforcement volume fraction on microstructure and properties. There is a lack of clear studies on the phase evolution kinetics under rapid heating and cooling conditions for laser additive manufacturing, and heating/cooling conditions and process kinetics play a key role in determining the final microstructure. Therefore, revealing phase evolution dynamics is crucial for understanding the microstructure formation mechanism and microstructure control.
However, due to the rapid and local nature of the laser melting process, it is very difficult to study the phase evolution kinetics during laser melting. To address this challenge, recent advances in synchrotron X-ray technology have made it possible to study the kinetics of the laser melting process. However, when it comes to in-situ characterization of MMC laser additive manufacturing, the focus is primarily on the effects of stable nanoparticles on spatter, porosity, melt pool, and crack dynamics. To date, no studies have reported direct observation of melting, solidification, dissolution, and precipitation kinetics during laser additive manufacturing of reaction-based MMCs.
In this study, Professor Lianyi Chen's team at the University of Wisconsin-Madison performed in-situ X-ray diffraction experiments to characterize the phase evolution kinetics of a 316L + 10 vol.%TiC system during laser melting. Complex phase evolution behaviors, including incomplete TiC dissolution and three-step TiC precipitation, are elucidated. Three different types of precipitates generate unique graded TiC micro- and nanostructures, increasing yield strength by 71% from 513 MPa to 877 MPa, tensile strength from 628 MPa to 1054 MPa, a 68% increase, and Young's modulus from 193 GPa to 221 GPa by 14%. The results of this study provide a theoretical basis for the unique microstructure and advanced MMC material design by laser additive manufacturing.
Link to article
Hatps://doi.org/10.1016/j.Aktamat.2024.119875
Main graphics
Figure 1.Sample preparation and in-situ laser melting X-ray diffraction experiments. (a) Schematic diagram of the sample collection point for tensile testing and in-situ X-ray diffraction experiments. (b) Schematic diagram of in-situ laser melting X-ray diffraction experiment.
Figure 2. In-situ X-ray diffraction characterization of melting and solidification kinetics during 316L laser melting. (a) XRD intensity plot of 0 sec to 4 sec during 316L laser melting. (b) XRD intensity plot of 0 sec to 0.5 sec during the melting of the 316L laser. (c) XRD image of 316L before and after laser melting.
Figure 3. In-situ X-ray diffraction characterization of melting, solidification, dissolution and precipitation kinetics during melting by 316L+10vol.%TiC laser at a scanning speed of 0.05 m/s. (a) XRD intensity plot of 0 s to 4 s during 316L+10vol.%TiC laser melting. (b) XRD intensity plot of 0 s to 0.5 s during melting of 316L+10vol.%TiC laser. (c) Evolution of peak intensities of γ-Fe 200 and TiC 111 during in-situ laser melting experiments. (d) Evolution of peak intensities of γ-Fe 200 and TiC 111 during cooling during in-situ laser melting experiments. (e) XRD patterns of 316L+10vol.%TiC laser before and after melting. (f) Evolution of the Q value of TiC 111 during the cooling process of the in-situ laser melting experiment.
Figure 4. SEM microscopic characterization of 316L (a) and 316L +10vol.%TiC (b-d).
Figure 5. Inverse polarity diagram (IPF) images of formed 316L (a) and formed 316L +10vol.%TiC (b).
Fig. 6.Thermodynamic calculation of Ti and C concentrations in a 316L liquid under equilibrium conditions.
Figure 7. Temperature evolution and distribution during laser melting. (a) Temperature evolution during heating during laser melting. (b) Temperature distribution of the liquid along the direction of the X-ray beam.
Figure 8. Dissolution thermodynamics and kinetic calculations of TiC laser melting process with initial particle sizes of 800 nm (a) and 5 μm (b).
Figure 9. Dynamic precipitation of TiC during laser melting of 316L+10vol.%TiC. (a) Phase diagram of 316L-TiC: stage I, TiC precipitation before 316L solidification, stage 2, TiC precipitation during 316L solidification, and stage 3, TiC precipitation after 316L solidification. (b) Thermohydrodynamic simulations of the temperature evolution during cooling of the laser melting process. (c) In phase I, thermodynamics and kinetics calculate the integral number of the TiC precipitated object as a function of temperature. (d) Thermodynamically and dynamically-calculated TiC precipitate as a function of temperature in Phase III.
Key conclusions:
(1) Based on the in-situ X-ray diffraction results, the evolution of the 316L phase and the TiC phase during the melting process of 316L + 10 vol.%TiC laser is revealed. The TiC phase undergoes a partial dissolution and a three-step precipitation process. Three different types of TiC precipitates were observed: primary TiC particles, eutectic TiC networks, and fine TiC nanoprecipitates. Together, these precipitates contribute to the formation of a unique layered TiC structure in 316L.
(2) Under equilibrium conditions, 1771 K is required to completely dissolve 10 vol.% TiC in 316L. When considering kinetics, a higher temperature (2494 K) is required to completely dissolve TiC. The thermohydrodynamic simulation results show that all the liquid metals along the path of the X-ray beam reach more than 1771 K during the laser melting process, while only 41% of the liquid metals reach 2494 K.
(3) The high cooling rate causes high supercooling, which leads to the deviation of the residual liquid composition in the primary precipitation process of TiC to the subeutectic region, resulting in the solidification of the γ-Fe phase before the eutectic reaction. Due to the significantly lower diffusion coefficient of Ti in γ-Fe, the precipitation rate of TiC in solidified 316L is very slow. Therefore, the precipitation of TiC in solidified 316L cannot reach the equilibrium state.
(4) The addition of TiC to 316L can significantly increase the yield strength (513 MPa to 877 MPa), the tensile strength from 628 MPa to 1054 MPa, and the Young's modulus from 193 GPa to 221 GPa.
About the Author
Lianyi Chen, an associate professor in mechanical engineering at the University of Wisconsin-Madison, received her Ph.D. from Zhejiang University in 2009, with research interests in additive manufacturing, nanoelement-based metal design, in-situ characterization, advanced materials and fabrication technologies for fusion energy.
Personal Data Source: https://directory.engr.wisc.edu/me/Faculty/Chen_Lianyi/
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