Strong metals are greatly desired for lightweight and energy-efficient industrial design. The strengthening of metals is traditionally accomplished by the additive contributions of various obstacle ...families (e.g., solid solutions, particles, and grain boundaries) and dislocation self-interactions that impede dislocation motion. In the present work, unlike a traditional additive understanding, a distinctive nonadditive strengthening mixture rule for obstacles and dislocations were validated based on experimental and modeling analyses in numerous cold-worked steels and aluminum alloys. Concretely, numerous well-annealed body-centered cubic steels and face-centered cubic aluminum alloys were prepared, in which the hierarchical strength levels of solid solutions, grain boundaries, and/or particles were estimated. The above specimens were then cold rolled to various strain levels. Dislocation densities were quantified by utilizing X-ray diffraction line-profile-analysis, and the dislocation density was found to increase faster with an increasing strain level when a high strength was presented in well-annealed specimens. When plotting the yield stresses as a function of the square roots of the dislocation densities in massive distorted samples based on the Taylor hardening law, it is interesting to note that an approximate single linearity was obtained. Individual dislocation strengthening was found to respond to the total strength in the deformed specimens, which indicated that the full nonadditive strengthening mixture rule was employed between the obstacle families and the dislocation contributions. The mechanisms of the observed nonadditive strengthening were also discussed by implementing additional experiments and transmission electron microscopy observations. Then, the modified constitutive models based on both one and two internal parameters’ Knocks-Mecking models were developed respectively, which excellently captured the effects of the various obstacle families on dislocation storage processes. The developed models also rationalized the observed nonadditive strengthening mixture rule.
Display omitted
•The dislocation densities accumulated faster in the cold-worked steels and Al alloys with higher primary strength.•Approximate single linearity was obtained by plotting the yield stresses as a function of square roots of the dislocation densities.•The full nonadditive strengthening mixture rule between the contributions of various obstacle families and dislocation junctions were observed.•The modified constitutive models based on both one and two internal parameters' Knocks-Mecking models were developed respectively.•The observed nonadditive strengthening mixture rule was captured well by the developed models.
The reverse transformation behavior during heating in Fe-10%Mn-0.1%C (mass%) martensitic alloy consisting of α’-martensite, ε-martensite and retained austenite was investigated using the in-situ ...neutron diffraction. When the temperature was elevated with a heating rate of 10 K/s, the ε→γ reverse transformation occurred first at the temperature range of 535–712 K, where Fe and Mn hardly diffused. In the temperature range where the ε→γ reverse transformation occurred, the full width at half maximum of the 200γ peak increased, indicating that the austenite reversed from ε-martensite contains high-density dislocations. In addition, the transformation temperature hardly depends on the heating rate and the crystal orientation of the reversed austenite was identical to that of the prior austenite (austenite memory), which suggests that the ε→γ reverse transformation would proceed through the displacive mechanism. After completion of the ε→γ transformation, the α’→γ reverse transformation occurred at the temperature range of 842–950 K. When the heating rate is low (<10 K/s), the reverse transformation start temperature significantly depends on the heating rate. It could be because the diffusional reverse transformation accompanying the repartitioning of Mn occurs. On the other hand, a higher heating rate (≥10 K/s) resulted in the disappearance of the heating rate dependence. This was probably due to the change in the transformation mechanism to the massive-type transformation, which is diffusional transformation without repartitioning of Mn.
Effect of ferrite grain size on dislocation strengthening was investigated in low carbon steels (0.006%C-0.15%C) with various grain sizes from 1 to 100 µm. In specimens with slight deformation, ...dislocation density increases in proportion to the inverse of ferrite grain size. In the dislocation density range below 2×1014/m2, dislocation density increases linearly against deformation strain but it tends to level off due to the dynamic recovery of dislocations when dislocation density has exceeded it. On the other hand, tensile tests revealed that yield stress follows the Hall-Petch relation for as-annealed specimens but follows the Bailey-Hirsch relation for cold rolled specimens. This means that flow stress depends on only the dislocation density regardless of grain size. As a result, it was concluded that the introduction of dislocations has been promoted with decreasing ferrite grain size and this results in the increase of flow stress in the uniform deformation region.
A thermomechanical control process (TMCP) was performed for two types of medium-Mn martensitic steels to investigate the effect of microstructure control on fracture behavior. The TMCPed 5%Mn-0.1%C ...steel (mass%) exhibited a slight improvement in toughness, and the partially occurring intergranular fracture was suppressed completely owing to the elongated prior austenite grains. On the other hand, the TMCPed 10%Mn-0.1%C steel exhibited a remarkably enhanced low-temperature toughness with separation, in which a large number of sub-cracks were observed parallel to the impact direction on the main fracture surface. The separation induced by the elongated prior austenite grains reduced the triaxial stress at the crack tip, which in turn contributed to improved toughness. Furthermore, the ultra-refined microstructure obtained by the combination of the γ → ε → α’ martensitic transformation and TMCP without recrystallization improved the toughness of 10%Mn-0.1%C steel.
Display omitted
The effects of retained austenite upon softening during low-temperature tempering at 373 K were investigated using martensitic carbon steels with and without retained austenite. To increase the ...amount of retained austenite, 10 mass% Ni was added to the base carbon steel (Fe-0.3C alloy). During tempering, the hardness decreased more rapidly in the Ni-added steel containing 6 vol.% retained austenite than in the base steel without retained austenite. Analyses of the microstructure and the carbon content in the solid solution (i.e., the solute carbon concentration) revealed that the retained austenite tended to suppress carbide precipitation and significantly reduced the solute carbon concentration in the martensitic matrix. We demonstrated that retained austenite acts as an effective absorption site for solute carbon in the martensitic matrix; however, the partitioned carbon is unevenly localized near the martensite/austenite interface, owing to the poor diffusivity at 373 K.
The relationship between hardness and solute carbon concentration estimated via electrical resistivity measurement was investigated in as-quenched and tempered martensitic steels containing carbon of ...0.3–0.6 mass%. As a result of corelating the amount of hardening due to carbon in solid solution with the solute carbon concentration, by the calculation to subtract precipitation strengthening, grain refinement strengthening, dislocation strengthening, and softening due to retained austenite from the total strengthening, we derived an equation of solid solution strengthening, where the hardening increases proportionally to the 1/2 or 2/3 power of the solute carbon concentration. It was confirmed that the effects of the factors other than solid solution strengthening due to carbon on hardness are relatively small in tempered specimens when the tempering temperature is less than 673 K; therefore, the change in hardness in tempered martensitic steels can be mostly explained by solute carbon concentration regardless of carbon content.
Microstructures of both as-quenched and intercritical-annealed medium-Mn steel were investigated by transmission electron microscopy to acquire a better understanding of γ–ɛ and γ–α′ martensitic ...transformations. The orientation relationships among each phases were determined as (1¯10)α′//(0001)ε,111α′//112¯0ε for the case of as-quenched sample, and (1¯10)α′//(1¯11)γ//(0001)ε,111α′//110γ//112¯0ε for the case of intercritical-annealed sample, respectively. In addition, the presences of core-shell type microstructures with the gradient of Mn concentration were confirmed with ɛ–martensite being surrounded by retained γ. Mn concentration of each phase was found dependent on the growth of reversed γ controlled by Mn diffusion during intercritical annealing. It was strongly suggested that the gradient of Mn concentration occurred during intercritical annealing affected the phase stability, which resulted in the formation of ε/γ core-shell microstructures.
Display omitted
The electrical resistivity of low-carbon martensitic steels was measured to estimate the carbon concentration in the solid solution. Since electrical resistivity is influenced not only by solute ...carbon but also by substitutional elements, lattice defects, and second phase, the effects of these factors need to be subtracted from the total electrical resistivity to obtain an accurate solute carbon concentration via this method. Consequently, the effects of dislocations and grain boundaries were much smaller than those of solute elements, representing approximately 1–2% of the total electrical resistivity in martensitic steel. However, substitutional elements and retained austenite were found to be significantly effective. By subtracting these effects from the measured value, the change in electrical resistivity owing to solute carbon (Δρsol.C) could be formulated as a function of the carbon concentration in the solid solution of martensite (Csol) as follows:Δρsol.CmΩmm = 0.25 × Csolmass%The estimated solute carbon concentration was confirmed to correspond to the directly measured value by atom probe tomography.
PDF
