2017 Volume 57 Issue 6 Pages 1138-1140
In this work, the effects of heat treatment parameters were investigated in order to introduce a novel third-generation advanced high-strength steel (AHSS) in the FeCrNiBSi alloy system. For this purpose, hot rolling was carried out on rapidly solidified specimens at a 950–1100°C temperature range with a reduction of 60%, followed by 50% cold-rolling reduction. The heat treatment was carried out at a 700–1100°C temperature range for 20–120 minutes. The results showed that a novel third-generation AHSS was obtained by heat treatment at 1100°C for 120 minutes. The additional second-step heat treatment at 700°C for 20 minutes resulted in a better combination of mechanical properties, with a formability index of 39.6 GPa%. These mechanical properties were attributed to microbands formation and δ-ferrite-to-austenite transformation-induced plasticity mechanisms.
Higher crash-worthiness, lower cost, lower fuel consumption, and lower gas emission are the current trends in the automotive industry. Advanced high-strength steels (AHSSs) with superior mechanical properties are strong enough to allow down-gauging while maintaining sufficient strength, which could result in vehicle weight reduction. While strengthening mechanisms (e.g., grain refinement, micro alloying, severe plastic deformation) continue to represent a pragmatic approach to enhancing the strength of existing AHSSs, there is also significant interest in the development of new alloys.1) Increasing the stability of austenite to delay the onset of necking using strain-induced transformation of austenite is a key to obtaining high ductility and high strength during alloy development.
Recently, novel FeCrNiBSi AHSSs (AM2B®) were developed and introduced by the current authors based on grain size, solid solution, and precipitation-strengthening mechanisms using three different casting cooling rates (slow, accelerated and fast) and in as-cast, hot-rolled and heat-treated conditions.2,3,4) The results on the as-cast specimens showed that increasing the casting cooling rate from slow to accelerated and fast leads to a transition in the matrix phases from two-phase austenite–δ-ferrite to a single austenitic matrix.2) As an index of formability and absorbed energy of materials, product of ultimate tensile strength (UTS) and total elongation percentage (El.%) has been applied to tailor steels roughly for automobile application.5,6) The heat-treated cast fast-cooled specimen showed a formability index of 24 GPa% which is the best among cast slow- and accelerated-cooled specimens.2) The investigation on the heat-treated hot-rolled slow-, accelerated- and fast-cooled specimens showed that the heat-treated hot-rolled fast-cooled specimen displayed the best mechanical properties (formability index=37.8 GPa%) which is attributed to strain-induced precipitation, strain-induced phase transformation, microband formation, and austenite grain rotation.3)
In the current work, the effects of heat treatment processes were investigated on the cold-rolled fast-cooled specimens; with the aim of introducing a novel third-generation AHSS in the FeCrNiBSi alloy system with a novel δ-ferrite-to-austenite induced plasticity mechanism.
Cast ingots of AM2B® (FeCrNiBSi alloy system, Fe59.0Cr15.0Ni11.5B9.2Si5.3 (at.%)) AHSS were prepared using an induction furnace under 10−4-mbar vacuum. Specimens (50×50×5 mm3) were produced using a water-cooled Cu-mold technique (fast-cooled) in a pressure vacuum caster machine. The as-cast specimens were heated to 1100°C for 120 minutes and then hot rolled at a 950–1100°C temperature range with a total reduction of 60% (5 mm-thick specimen to plate thickness of 2 mm), followed by a 50% reduction (2 mm to 1 mm plate thickness) with cold rolling. The as-cold-rolled specimens were subjected to three different heat treatment processes. In the first heat treatment process, the as-cold-rolled specimens were isothermally heated to 1100°C and held for 120 minutes (1100 specimens). In the second heat treatment process, some of the heat-treated specimens were subjected to an additional isothermal heat treatment at 700°C for 20 minutes (1100/700 specimens). In the third heat treatment process, the as-cold-rolled specimens were heated to 700°C and held for 20 minutes (700 specimens). The microstructures were observed using field emission gun scanning electron microscopy (FESEM) after etching with Kalling’s No. 2 etchant. A Philips X’Pert PRO X-Ray diffractometer (XRD) using Co-Kα (λ=0.179 nm) was used to verify the structures of the specimens. Hardness was measured using the Vickers method. The room-temperature uniaxial tensile test was conducted at a cross-head speed of 0.6 mm/min with dog-bone shaped specimens (11.4 mm gauge length and 3.0 mm width) using a Santam tensile testing machine. Yield stress (Y.S.) was determined using the 0.2% offset plastic strain method. The tensile specimens were machined from the plates parallel to the cold-rolling direction. Three specimens were hardness- and uniaxial tensile-tested, and the mean values were reported.
The effects of the heat treatment processes on the mechanical properties of the as-cold-rolled specimens are shown in Fig. 1. The engineering curves (Fig. 1(a)) show two distinct flow behaviors. The first type of flow curve shows a linear elastic region followed by rapidly increasing stress levels caused by increasing plastic strain levels (curves of the as-cold-rolled and 700 heat-treated specimens, Fig. 1(a)). The second type shows a gradually increasing plastic flow with increasing plastic strains after a linear elastic region (flow curves of the 1100 and 1100/700 heat-treated specimens, Fig. 1(a)). As indicated in Fig. 1(b), after each heat treatment process, Y.S. and El.% showed an overall increasing tendency, while UTS showed a decreasing trend. The specimens began to yield at 200–377.4 MPa, followed by different degrees of El.% prior to failure, in a range of 2.7–45.2%, and UTS range of 875.5–1259.1 MPa. It is obvious that the specimens heated at 1100°C for 120 minutes showed an improvement in El.%, such that the 1100 and 1100/700 heat-treated specimens showed a major increase in El.% to 40% and 45.2%, respectively. In terms of hardness values (Fig. 1(c)), maximum work hardening was displayed in the 1100/700 heat-treated specimen. The formability indexes in the current study were 3.4, 5.6, 36.2, and 39.6 GPa% for the as-rolled and 700, 1100, and 1100/700 heat-treated specimens, respectively. Therefore, using rapid solidification and applying appropriate heat treatment improved the formability index of these novel AHSSs (1100 and 1100/700 specimens) to the extent that are in the third-generation AHSS envelope, with a formability index of 25–50 GPa%.5,6) It is worth mentioning that the as-cold-rolled and 700 heat-treated specimens are competitive with martensitic AHSS with a formability index of (3–18) GPa%.5,6)
Mechanical properties of the as-cold-rolled and heat-treated cold-rolled specimens: (a, b) engineering stresses and total elongation percentage; (c) Vickers hardness.
Figure 2 shows secondary electron FESEM images at low and high magnifications from gripper (Fig. 2(a)) and gauge (Fig. 2(b)) sections of the 1100/700 heat-treated tensile specimen after fracture. XRD patterns of the gripper and gauge sections of the as-cold-rolled and 700, 1100, and 1100/700 heat-treated specimens are presented in Fig. 2(c). The identified phases using XRD analysis are also shown in Table 1. The microstructure of the gripper section of the 1100/700 heat-treated specimen contains recrystallized austenite grains (6.3±3 μm length, 5±2 μm width) with semi-oval to round M2B (M=Fe, Cr, using energy dispersive spectroscopy) precipitates (2±0.5 μm length, 1.4±0.4 μm width) at the grain boundaries. While the XRD pattern shows δ-ferrite in the gripper section of the 1100/700 heat-treated specimen, it could not be distinguished by FESEM. There are also some annealing twins in the microstructure of the gripper section of the 1100/700 heat-treated specimen. The gauge section of the 1100/700 heat-treated specimen contains severely deformed and elongated austenite grains, with and without microbands, as shown in Fig. 2(b). The existence of microbands in some austenite grains implies a strong dependence of microbands formation to crystallographic orientation of grains during tensile loading. A high density of ultrafine and nanometer precipitates is also seen in both the gripper and gauge sections of the 1100/700 heat-treated specimen. Using XRD patterns (Table 1), CrSi2, FeSi, and FeB are identified as ultrafine and nanometer precipitates in the gauge and gripper sections of the 1100/700 heat-treated specimen. The XRD patterns indicate that δ-ferrite-to-austenite transformation-induced plasticity also occurred in the 1100/700 heat-treated specimen, as shown in Fig. 2(c). The same phenomenon also occurred in the 1100 heat-treated specimens. These XRD results are repeatable, as the XRD test was repeated three times for each of the heat-treated specimens. Therefore, heat treatment at 1100°C for 120 minutes resulted in the formation of δ-ferrite in the 1100 and 1100/700 heat-treated specimens; there is no evidence of δ-ferrite in the as-cold-rolled and 700 heat-treated specimens (Fig. 2(c)). The fast-cooled AM2B® AHSS specimens have been investigated in the as-cast, as-hot-rolled and as-cold-rolled as well as heat-treated conditions by the authors.2,3) The same phenomenon of δ-ferrite-to-austenite transformation-induced plasticity has been also observed in the heat-treated cast fast-cooled specimens2) while the heat-treated hot-rolled fast-cooled specimens showed austenite-to-martensite transformation-induced plasticity and austenite grain rotation.3) More specific information on deformation mechanisms of the heat-treated cast and hot-rolled AM2B® AHSS specimens can be found elsewhere.2,3) It is reported that the initial microstructure may affect the heat treatment response and the mechanical properties as well.7) Therefore, the heterogeneous deformation of the fast-cooled specimen during cold rolling (strain incompatibility of austenite and M2B phases) may influence the response of the cold-rolled specimen with respect to the hot-rolled specimen to the same heat treatment cycle (1100/700 heat treatment cycle) as observed in our previous and current papers.3) This different initial microstructure of 1100/700 heat-treated cold-rolled specimen (in terms of nature of matrix grains, matrix grains and M2B dimensions and volume fractions and local chemical compositions) can lead to different deformation mechanisms as reported in the current work.
Microstructure of the gripper (a) and gauge (b) sections of the 1100/700 specimens, as well as XRD patterns of all the specimens.
Previous works in the literature8) have shown that the formation of δ-ferrite in heat-treated specimens resulted in Ni-rich austenite–δ-ferrite interfaces. Furthermore, it has been reported that the generated crystal defects during deformation can increase the austenite nucleation sites.9) Eisenlohr et al.10) showed that the velocities of lattice dislocations can be high and in some cases may approach shear wave velocity, and that as a result of internal friction, they can lead to a high energy release rate. This associated deformation energy can result in local temperature increases of up to several hundred degrees Kelvin.11) The likely local temperature increase during tensile loading of the 1100 and 1100/700 heat-treated specimens may bring the local area to an austenite–δ-ferrite two-phase region, where austenite nucleates at the austenite–δ-ferrite interfaces with higher crystal defects and Ni concentration. Deformation-associated heat energy is transferred to the adjacent surroundings that caused a high cooling rate, and during cooling, δ-ferrite transforms to austenite. However, this supposition requires a great deal of followup research for further verification.
It has been reported that in planar glide materials,12) by applying strain to the material, Taylor lattice formation is the main strain energy relaxation mechanism at low strains. With increasing strain, a domain boundary will be formed, which is a single dislocation wall consisting of geometrically necessary dislocations to accommodate the misfit orientation between the Taylor lattice domains. With further increasing strain, a second dislocation wall parallel to the first formed domain wall will appear. The region between the first and second formed walls is usually called a microband.13) It has been reported that the length of these microbands extends to 50–100% of the grain diameter, as seen in the current work (Fig. 2(b)).
So, as it was shown in Figs. 1 and 2 the high formability index of the heat-treated cold-rolled specimens (1100 and 1100/700) could be related to mechanisms involved in releasing the applied strain energy during deformation which are δ-ferrite-to-austenite transformation-induced plasticity and microband formation.
The principal conclusions that can be drawn from the present work are as follows:
• The present 1100 and 1100/700 heat-treated cold-rolled specimens showed an excellent combination of strength and ductility which are UTS=950.2 MPa, Y.S.=377.4, and MPa, El.%=40% and UTS=875.5, Y.S.=303.3, and El.%=45.2%, respectively.
• As a result of the best combination of mechanical properties in terms of formability index, which ranged 36.2–39.6 GPa%, the 1100 and 1100/700 heat-treated cold-rolled specimens are located in the targeted third-generation AHSS envelope.
• The deformed microstructures of the 1100 and 1100/700 heat-treated cold-rolled specimens exhibited δ-ferrite-to-austenite- and microband-induced plasticity mechanisms.
• The as-cold-rolled and 700 heat-treated cold-rolled specimens are competitive with martensitic AHSSs.
This research was ﬁnancially supported by the PhD grants for PhD students of Tarbiat Modares University (9260702001) and funding of Nano Structured Advanced Materials Technologies Development Co. (NAMAD-1392) for advanced high-strength steel innovation project. The authors gratefully appreciate Dr. Ghaffari (NAMAD Co.) for the support of this project.