1. Introduction
A challenge exists in the automotive industry to develop new, hot-rolled, high-strength low-alloy (HSLA) steels offering a balance of high tensile strength and superior stretch-flange formability to reduce vehicle weight without compromising safety, performance, or manufacturability [
1,
2]. The steel industry has responded by developing ferritic steels strengthened with extensive nano-sized precipitation [
2,
3,
4]. The single-phase ferritic matrix eliminates hard constituents and imparts superior stretch-flange formability, while its high yield and tensile strengths are derived from nano-sized precipitates. Titanium (Ti)-, niobium (Nb)-, or vanadium (V)-based microalloy systems are typically used for such HSLA steels, and molybdenum (Mo) is often added to strongly retard the precipitate coarsening rate [
4,
5]. Substitutional manganese (Mn) additions are also made to such HSLA steels to compensate for the low carbon levels and lower the Ar
3 transformation temperature for better refinement of the microalloy precipitation sizes [
6]. However, the Mo and Mn additions can enhance the hardenability of the steel, resulting in slower austenite decomposition kinetics and hard secondary phase constituents (e.g., bainite and/or martensite) in the final microstructures. Hard constituents are undesirable because stretch-flange formability is markedly reduced due to the nucleation of voids at the interfaces between the relatively hard and soft phases [
7,
8]. Therefore, the ability to obtain both fine microalloy precipitates and a single-phase ferritic matrix in the final microstructure requires attention to thermomechanical processing, due to its effect on the austenite decomposition behavior [
9].
Hot strip mills (HSM) are a key part of HSLA steel production. A typical, semi-continuous HSM consists of the following units: reheat furnace, roughing stand, transfer table, coilbox, finishing mill, runout table, and coiler [
10]. The reheat furnace heats the slab to a suitable temperature to start hot rolling. The roughing stand is used for major reductions in slab thickness. The transfer table carries the slab, now called the transfer bar, from the roughing stand to the finishing mill. A coilbox is sometimes used to wrap the transfer bar into a coil to obtain a more uniform temperature profile [
10]. The transfer bar is then delivered to the finishing mill, which is used for more precise gauge reductions. The strip is water-cooled from coolant headers along the runout table to control the final microstructure and then wrapped into a coil, which slowly cools to room temperature. Important processing steps during HSM processing to develop the desired, ferritic microstructures are controlled rolling in the finishing mill and accelerated cooling on the runout table (lowers the coiling temperature) [
9].
Three types of austenite recrystallization behaviors are typical in HSM processing: static recrystallization, austenite pancaking (i.e., avoidance of recrystallization), and dynamic/metadynamic recrystallization [
11]. The long interpass times and high temperatures, above the non-recrystallization temperature (T
nr), during rough rolling allow for nearly complete static recrystallization (SRX) to take place between rolling passes. Fine, equiaxed austenite grains are produced with negligible strain accumulation through SRX. Finish rolling is typically close to or below the T
nr, thus encouraging rapid strain-induced precipitation (SIP) of microalloy carbonitrides (e.g., Nb [
12], Ti [
13], and V [
14] containing HSLA steels). These precipitates retard or even prevent the SRX of austenite grains; leading to pancaking, greater strain accumulation, and the generation of defects like dislocations and deformation bands. One of the major differences between HSM and plate mill rolling schedules is the interpass times, which are much shorter in the HSM and range from 0.2–5 s [
11,
15,
16]. These short interpass times encourage strain accumulation and dynamic recrystallization (DRX) [
11,
17]. Fine, equiaxed austenite grains with negligible strain accumulation are also produced through DRX [
18].
Hot torsion testing has been employed in numerous hot rolling simulations and recrystallization studies of HSLA steels for its ability to impose large amounts of strain while accurately controlling temperature, interpass time, and strain rate [
9,
15,
18,
19,
20,
21]. The conditioning of austenite into pancaked grains during HSM processing is important for final microstructural development. Whitley et al. [
19] considered the evolution of the austenite grain morphology during hot torsion testing (shear deformation), and
Figure 1 [
19] shows a schematic overview of the expected austenite morphology at several stages throughout testing. The torsional axis is vertical to the page in
Figure 1. Grains undergo one or both of the following processes at a given time: (i) shear deformation, thus becoming more elongated in morphology; and (ii) recrystallization, thus becoming refined and equiaxed in nature. Austenite grains are assumed to be initially equiaxed after soaking at high temperatures (
Figure 1a), and they become elongated and rotated by the application of shear strain (
Figure 1b) when viewed normal to the torsional axis of a cylindrical specimen [
19]. If deformation occurs above the T
nr, SRX of the austenite is expected given sufficient interpass time (
Figure 1c). However, if deformation occurs below the T
nr, pancaking of the austenite is expected and results in rotated grains with higher aspect ratios (
Figure 1d). Austenite grains accommodate strain in this manner until there is sufficient driving force for recrystallization in the form of stored strain energy [
19]. Strain-free grains can form in the regions of highest stored strain energy (e.g., grain boundaries and deformation substructure) (
Figure 1e).
A metallographic technique was developed by Whitley et al. [
19] to quantify the shear strain accumulation within (prior) austenite microstructures produced via hot torsion testing. Tested samples are sectioned parallel to the torsional axis and metallographically prepared to observe the prior austenitic microstructures. In this “tangential” plane cross-section, the inclination angles (
) of prior austenite grains and other microstructural features can be measured with respect to the torsional axis and used to estimate the amount of shear strain that accumulated during testing (γ
acc) with the relationship
Multiple austenite deformation-recrystallization cycles during multi-pass torsion testing can result in mixed (prior) austenite microstructures displaying a distribution of inclination angles [
18]. These distributions reflect the local variation in strain accumulation that can result during thermomechanical processing.
Figure 2 [
19] shows an example of this technique applied to a 1045 steel microalloyed with V that underwent industrial bar rolling simulations via hot torsion testing. Various microstructural features are highlighted with their corresponding inclination angles with respect to the torsional axis, which indicated:
- (A)
Manganese sulfide (MnS) inclusion, elongated and initially oriented parallel to the rolling direction. Since MnS inclusions do not recrystallize during thermomechanical processing, γacc represents the total shear strain imposed;
- (B)
Highly elongated prior austenite grain with the same inclination angle as MnS, suggesting no recrystallization during thermomechanical processing;
- (C)
Elongated prior austenite grain with an inclination angle less than (A) and (B), indicating some degree of recrystallization during thermomechanical processing, followed by subsequent deformation. The measured inclination angle represents the amount of shear strain that accumulated after the last recrystallization event, assuming an equiaxed grain morphology after recrystallization;
- (D)
Fine, equiaxed prior austenite grains that indicate recrystallization without shear strain accumulation.
The main aim of this work was to investigate differences in austenite strain accumulation before decomposition and the associated influence on (prior) austenite morphology and microstructural development after isothermal transformation for a low-carbon, Ti-Mo microalloyed steel.
2. Materials and Methods
A low-carbon, Ti-Mo microalloyed steel was investigated, and its chemical composition is shown in
Table 1. The experimental alloy was received as 16 mm thick, hot-rolled steel from Baoshan Iron & Steel Co. (Shanghai, China).
Table 2 provides estimates of particular critical transformation temperatures using empirical equations found in the literature. These critical transformation temperatures were used to guide the thermomechanical processing of the experimental alloy. The following empirical equations were used: Andrews for Ac
1 and Ac
3 [
22], Schacht for Ar
1 [
23], Pickering for Ar
3 [
24], Lee #2 for B
s [
25], and Borrato for T
nr [
26]. The M
s temperature was determined experimentally using dilatometry [
27].
Solutionizing temperatures for relevant compounds in austenite were determined based on solubility expressions [
28,
29]. The calculations showed that titanium nitride (TiN) remains undissolved during solid-state processing, thus all nitrogen (N) was assumed to be removed from solid solution. The evolution of equilibrium phases as a function of temperature was predicted with Thermo-Calc
® (Thermo-Calc Software, Solna, Sweden, Version 2019) using the TCFE9 database (assuming all N was already incorporated into TiN precipitates), and the results are shown in
Figure 3. The MC equilibrium phase represents a mixed microalloy carbide exhibiting the NaCl (B1) crystal structure without the incorporation of N. From the determined equilibrium dissolution temperature of MC, a soaking temperature of 1250 °C was selected.
2.1. Hot Torsion Testing
Hot torsion testing was accomplished using a Gleeble
® 3500 system equipped with the Hot Torsion Mobile Conversion Unit (Dynamic Systems Inc., Poestenkill, NY, USA). Sub-sized torsion samples were machined from the as-received material according to the schematic illustration in
Figure 4, where the rolling direction was parallel to their lengths.
Figure 5 shows a photograph inside the Hot Torsion Mobile Conversion Unit chamber, highlighting the setup used during testing. Both ends of the sample were restrained to keep the reduced gauge length fixed during torsion testing, and the Gleeble
® 3500 was programmed to minimize axial stresses by adjusting the stroke arm. Axial stresses did not exceed ±10 Mpa during testing. Helium (He) gas was used as the quenchant for all tests and was directed from quench heads both in front and behind the sample. The torsion motor coupler within the Hot Torsion Mobile Conversion Unit was set for 20° free rotation, which allowed rapid acceleration of the torsion motor during deformation as well as a rapid reduction in torque on the sample during interpass times. Hot torsion testing was performed under the protective environment of argon (Ar) gas to minimize oxidation and decarburization near the surface of the sample. The oxygen partial pressure within the chamber was maintained under 30 ppm during testing and monitored using a PurgEye
® 200 oxygen sensor (Huntingdon Fusion Techniques, Burry Port, UK).
The temperature of each sample was monitored at the mid-length of the reduced gauge section using a Metis Model MQ11 optical pyrometer (Process Sensors Corporation, Milford, MA, USA) that was calibrated at 1100 °C prior to testing. The pyrometer is not reliable below ~700 °C, so it was used to control temperature during heating, soaking, and deformation. An alternative method was required to control temperature during the accelerated cooling and isothermal holding steps due to the relatively low temperatures employed. Attaching a thermocouple to the fixed shoulder of the sample, where limited deformation occurs, and accounting for the temperature difference between the shoulder and mid-length of the reduced gauge section, proved to be an efficient method for controlling temperatures below ~700 °C. Therefore, a Type K thermocouple was spot welded to the surface of the sample roughly 0.5 mm away from the fixed shoulder (as shown in
Figure 5) for testing that included isothermal holding. Each thermocouple wire was insulated with a small section of ceramic tubing to prevent short-circuiting. An isothermal holding temperature of 650 °C was planned, and preliminary testing showed that an offset value of approximately 18 °C (i.e., shoulder temperature of 632 °C) was appropriate to account for the temperature difference between the shoulder and reduced gauge section.
Microstructural gradients within samples undergoing torsion testing require the selection of a specific radial position to determine deformation parameters since imposed shear strain (γ) varies with the radius of the reduced gauge section (
r), according to
where
is the angle of twist (in radians),
L is the reduced gauge length of the sample, and
is the expected inclination angle with respect to the torsional axis corresponding to the imposed shear strain. The radial position used in this work to determine deformation parameters during HSM processing simulations was the “effective radius”, which is positioned at 72.4 pct of the radial distance from the central axis [
30]. Barraclough et al. showed that this location represents the bulk behavior for materials having a wide variety of strain rate sensitivities and/or strain hardening behaviors [
30].
The following equations were used to convert the pass-by-pass shear strains and strain rates into appropriate angles of twist and twisting times for simulation purposes, as well as to convert the resulting torque to shear stress. The angle of twist for each pass was calculated according to
and the twisting time (
t) for each pass was calculated according to
where
is the shear strain rate. Barraclough et al. [
30] developed an equation to convert torque (Γ) to shear stress (τ) considering the effective radius, which assumes pure torsion and uniform shear strain along the length of the reduced gauge section
where
is the outer surface radius and
is the inner surface radius. Finally, the shear stress and strain values were converted to equivalent true stress (
) and strain (
) values by applying the Von Mises criterion
and used to determine the MFS for each pass according to
where
a and
b are the initial and final equivalent true strains per pass, respectively. The integrals were solved using analytical solutions of logarithmic regressions of the data according to
where
a1 and
b1 are constants. This approach typically assumes uniform constitutive mechanical properties of the material through the cross-section.
Table 3 summarizes the hot torsion testing schedule applied. The testing parameters were developed after consideration of the literature [
9,
11,
15,
18] and industrial processing [
16,
31]. Samples were heated at 5 °C/s to a soaking temperature of 1250 °C and held for 5 min to dissolve microalloy carbides. Rough rolling simulations consisted of four identical passes between 1240 and 1150 °C, each with relatively long interpass times to promote SRX of the austenite. Finish rolling simulations consisted of seven passes: either between 1150 and 1000 °C (designated as High T Finish) or between 1050 and 900 °C (designated as Low T Finish). Note that the designations of either High T Finish or Low T Finish simulations include the identical rough rolling simulations. These temperature ranges were selected to be above or mostly below the estimated T
nr of approximately 1000 °C to develop (prior) austenite microstructures with drastically different strain accumulation prior to decomposition. The roughing-to-finishing delay was 30 and 100 s for the High T Finish and Low T Finish simulations, respectively. Short interpass times are typical during HSM processing, which promote austenite pancaking and possibly DRX of the austenite later during finish rolling [
11,
15,
16]. Overall, the amounts of true strain imparted during rough and finish rolling simulations were about 1.60 and 2.40, respectively, totaling about 4.00. The cooling rate between all passes was kept constant at 5 °C/s to ensure accurate temperature control. Additionally, relatively low target shear strain rates were utilized to ensure accuracy of the imparted shear strains. After the last finishing pass (F7), samples were either: (i) quenched as rapidly as possible (~43 °C/s) to room temperature to investigate the prior austenite grain (PAG) size and morphology, or (ii) accelerated cooled at ~30 °C/s to an isothermal holding temperature of 650 °C, held for 30 min, and finally quenched as rapidly as possible to room temperature to investigate polygonal ferrite characteristics and the presence of any secondary phase constituents.
Previous research [
19] has shown that a tangential orientation is best for investigating PAG morphologies and quantifying the amount of shear strain that accumulated within the microstructure for samples tested via hot torsion. This orientation is presented schematically in
Figure 6. Tested samples were prepared in the tangential orientation according to the following procedure. First, the reduced gauge section was cut free from its ends on both sides of the sample. Next, the gauge section was cut in half (perpendicular to the torsional axis) to reveal the “thermal plane”. Note that the thermal plane corresponds to the mid-length of the reduced gauge section, approximately where the optical pyrometer was aligned prior to testing. Then, each piece was cut in half (parallel to the torsional axis). Finally, these quartered sections were mounted in Bakelite and precision ground to the radial position of interest using measurements of chord length to reveal the tangential plane.
In addition to the effective radius, two other radial positions were selected to further investigate how strain influences austenite conditioning and the final microstructures during HSM processing. Recall that shear strain varies with the radius of the reduced gauge section according to Equation (2). The target shear strain for the HSM processing simulations was 8.00, where the effective radius (0.724 radial position) was used to determine the deformation parameters. Radial positions of 0.50 and 0.90 were also selected to represent an extensive range of possible shear strain accumulation. The approximate shear strain for the 0.50 radial position is 5.52, which might represent processing of thicker gauge material, and the approximate shear strain for the 0.90 radial position is 9.94, which might represent processing of thinner gauge material. All three radial positions were investigated using the tangential orientation for each condition produced via hot torsion testing.
2.2. Microstructural Characterization
Samples for microscopic evaluation were sectioned, mounted in Bakelite, and prepared using standard metallographic procedures. Samples were etched with either a 1 pct nital or a modified Béchet–Beaujard reagent. The 1 pct nital reagent was used to reveal ferrite grain boundaries and secondary phase constituents. The modified Béchet-Beaujard reagent was used to reveal PAG boundaries and consisted of 200 cm3 of deionized water, 2.6 g of picric acid solids, 8 cm3 of Teepol (wetting agent), and 2 cm3 of hydrochloric acid. This reagent was heated to 65 °C on a temperature-controlled hot plate equipped with a thermocouple feedback and stirred with a magnetic stir rod throughout the etching process. After each interval of etching with the modified Béchet-Beaujard reagent, samples were immersed in a methanol bath, ultrasonicated, and dried with a heat gun. PAG boundaries were highlighted in black to enhance their clarity within the provided micrographs.
General imaging of microstructures was accomplished using light optical microscopy (LOM), where the micrographs were used to determine average grain size, phase area fraction, etc. LOM was performed with an Olympus Model PMG3 inverted light microscope (LECO Corporation, St. Joseph, MI, USA) with a PAXcam Model PX-CM digital camera (MIS Inc., Villa Park, IL, USA) and PAX-it! Image analysis software (MIS Inc., Villa Park, IL, USA, Version 7.8). Electron backscatter diffraction (EBSD) analysis was performed on isothermally transformed microstructures with a JSM-7000F field emission-scanning electron microscope (JEOL USA Inc., Peabody, MA, USA) to investigate polygonal ferrite grain diameter distributions. Prior to EBSD analysis, samples were metallographically prepared and vibratory polished for at least 4 h using 0.02 μm colloidal silica solution. EBSD scans were performed at an accelerating voltage of 20 keV, calibrated EBSD camera working distance of 18 mm, and step size of 0.1 μm. EBSD data were collected with a Hikari Pro detector (EDAX Inc., Mahwah, NJ, USA) using the TEAMTM software ( EDAX Inc., Mahwah, NJ, USA, Version 4.5), and the datasets were analyzed with the Orientation Imaging Microscopy Analysis© software ( EDAX Inc., Mahwah, NJ, USA, Version 8.1). EBSD datasets were cleaned using the following functions: Neighbor Orientation Correction (Level 3, Tolerance 5.0, Minimum Confidence Index (CI) 0.10); Grain CI Standardization (Tolerance 5.0, Minimum Size 3, Multi Row 1); and Neighbor CI Correlation (Minimum CI 0.30, Single Iteration).
Polygonal ferrite and prior austenite grain sizes were determined with the concentric circle method utilizing the ImageJ software (open-source). The intercepts of the circles with the grain boundaries were counted, and the average intercept lengths were calculated and reported as the average grain sizes. A total of 1000 or more grain boundary intercepts were counted for each condition to determine a representative grain size. Note that prior austenitic twins were observed within some PAGs, but these were not considered in the PAG size measurements. Aspect ratios of the PAGs were determined by measuring the major and minor axes of individual grains (assuming an elliptical shape) with the ImageJ software and calculating their ratios. Inclination angles of the PAGs with respect to the torsional axis were measured with the ImageJ software and used to quantify the amount of shear strain that accumulated within the (prior) austenite microstructures using the previously described metallographic technique developed by Whitley et al. [
19]. A total of 100 or more PAGs were measured for each condition to determine a representative aspect ratio and inclination angle.
The ImageJ software was also used to determine the area fraction of secondary-phase constituents using an image thresholding procedure. This procedure was employed because of the distinct difference in the etching response of the polygonal ferrite (carbon depleted) and secondary phase constituents (carbon rich) with 1 pct nital. First, LOM micrographs were thresholded from grayscale images to black-and-white images using a grayscale value range that best captured the secondary phase constituents but without capturing polygonal ferrite grain boundaries. After thresholding, pixels in the desired grayscale value range were transformed to white pixels (representing the secondary phase constituents), while all other pixels were transformed to black pixels (representing polygonal ferrite).
Figure 7 provides an example of the thresholding process. Finally, the ratio of the number of white or black pixels to the total number of pixels was used to determine the secondary phase constituents or polygonal ferrite area fractions, respectively. This procedure was repeated at five or more different areas of the microstructure for each condition to determine representative phase area fractions.
Samples were metallographically prepared and etched with 1 pct nital prior to Vickers microhardness testing in order to relate microhardness to microstructural features. Vickers microhardness testing was performed with a LM110 hardness tester (LECO Corporation, St. Joseph, MI, USA) according to ASTM Standard E384-17 [
32]. A 5 × 5 array of indents with an indentation load of 100 g was used to determine the representative values of the polygonal ferrite in each isothermally transformed condition. Indentations immediately adjacent to secondary phase constituents (including large TiN precipitates) were disregarded. Very low indentation loads of 10 g were used to investigate the small secondary phase constituents. Note that indentation loads like 10 g can result in consistently higher microhardness values [
33], so values determined with this indentation load were used for comparison only.