Development of Low-Alloyed Low-Carbon Multiphase Steels under Conditions Similar to Those Used in Continuous Annealing and Galvanizing Lines

Abstract

In the present work, a Cr+Mo+Si low-alloyed low-carbon steel was fabricated at laboratory scale and processed to produce multiphase advanced high-strength steels (AHSS), under thermal cycles similar to those used in a continuous annealing and galvanizing process. Cold-rolled steel samples with a microstructure constituted of pearlite, bainite, and martensite in a matrix ferrite, were subjected to an intercritical annealing (817.5 °C, 15 s) and further isothermal bainitic treatment (IBT) to investigate the effects of time (30 s, 60 s, and 120 s) and temperature (425 °C, 450 °C, and 475 °C) on the resulting microstructure and mechanical properties. Results of an in situ phase transformation analysis show that annealing in the two-phase region leads to a microstructure of ferrite + austenite; the latter transforms, on cooling to IBT, to pro-eutectoid ferrite and bainite, and the austenite-to-bainite transformation advanced during IBT holding. On final cooling to room temperature, austenite transforms to martensite, but a small amount is also retained in the microstructure. Samples with the lowest temperature and largest IBT time resulted in the highest ultimate tensile strength/ductility ratio (1230.6 MPa-16.0%), which allows to classify the steel within the third generation of AHSS. The results were related to the presence of retained austenite with appropriate stability against mechanically induced martensitic transformation.

1. Introduction

One of the current demands of the automotive industry is to reduce the weight of vehicles, which is considered an efficient way to reduce both fuel consumption and CO₂ emissions [1]. Reducing the thickness of steel sheets and increasing the strength-to-ductility ratio is often a contradictory task in conventional steels [2,3]. Therefore, to assure passengers safety, modifications in the design of vehicles and in the selection of materials to fabricate their components are required [4]. Advanced high strength steels (AHSS) are promising candidates to satisfy these needs, since they allow significant reduction of the weight of several vehicles’ components while presenting extraordinary mechanical properties at the same time [5].

Unlike conventional steels, in which strength increases, while elongation to fracture decreases, AHSS can exhibit higher strength-to-ductility ratio due to the combination of multiphase structures [6]. Recently, the development of the “3rd Generation” of advanced high-strength steels has attracted increasing attention [7,8,9]. This generation of AHSS seeks to provide a better combination between high strength and ductility compared to the first generation, but without the joining problems and high costs associated with the second generation [10,11,12].

Computer simulations of intercritical continuous cooling transformation (CCT) diagrams, calculated as a function of chemical composition, have been reported as an approach to design new chemistries to obtain multiphase steels [2,13]. However, computational programs have still some limitations regarding the changes in the chemical composition that may occur during intercritical annealing and further isothermal bainitic treatments, which may change the austenite decomposition behavior during the thermal cycles, making it difficult to estimate properly the transformation products resulting from intercritical annealing followed by isothermal bainitic treatment. The development of low-alloyed multiphase AHSS under similar conditions to those used in a CAG process is a current challenge [2,14].

Some authors studied the effects of the partial substitution of Si (0.82–1.53%) by Al for the fabrication of C-Mn-Si-Al multiphase steels [15]. Processing involved homogenizing, hot rolling, and cold rolling in two steps, intercritical annealing and a final isothermal bainitic treatment [15]. Multiphase steels containing ferrite, bainite, and austenite were obtained by intercritical annealing at 750 °C for 420 s followed by isothermal bainitic treatment at 390 °C for 300 s [15]. However, such microstructure would be difficult to obtain under conditions simulating a continuous annealing and galvanizing process considering that the melting point of Zn is about 420 °C.

Other researchers have developed third-generation advanced high-strength steels with C (0.12–0.20%), Mn (4.98–10.02%), Si (3.11–3.17 %), Al (3.05–3.19%) [16]. In this case, hot-rolled plate samples (6 mm in thickness) were intercritically annealed at an elevated temperature for 15 min and water quenched [16]. Although mechanical properties were classified within the third generation of advanced high strength steel, the high contents of the alloying elements led to a high cost of production, and this becomes a deterrent to extensive use of the steel [16].

The separated effects of Mo and Cr on the on the microstructure and mechanical properties of NbV-microalloyed steels have been also investigated [17]. They found that additions of either 0.2 wt.% Mo or 0.2 wt.% Cr resulted in formation of a microstructure consisting mainly of bainite and martensite. Mo facilitates the bainite transformation, and solid solution strengthening; it also decreases the rate of dynamic recrystallization of austenite, leading to grain refinement [17]. Cr also facilitates the bainite transformation and delays Fe₃C precipitation in low carbon steels, however, the solid solution strengthening effect of Cr is about 6 times weaker than the one of Mo [17]. Si stabilizes ferrite and suppresses the formation of cementite during the IBT, stabilizing the retained austenite [17].

Multi-element microalloying with minor additions can be an efficient way to optimize steel composition and mechanical properties. Amongst the published literature, the combined effects of Mo, Cr, and Si in multi-low-alloyed steels has been scarcely investigated to produce multiphase advanced high strength steels under conditions similar to those used in a continuous annealing and galvanizing process. In the present work a Cr+Mo+Si low-alloyed low-carbon steel was fabricated at laboratory scale to further investigate the effects of time and temperature of the isothermal bainitic treatment on the resulting microstructures and mechanical properties. Investigated variables were set considering conditions similar to those used in CAG lines for thin steel sheets.

2. Experimental Procedure

2.1. Design of Chemistry for the Experimental Steel

The chemical composition proposed for the experimental steel was set based on the procedures reported by Navarrete, 2021 to obtain advanced high strength steels [2]. To this end, simulations with JMatPro 8.0 about the behavior of continuous cooling transformation (CCT) and time-temperature-transformation (TTT) diagrams were conducted as a function of chemical composition to intentionally move the transformation curves and be able to obtain multiphase steels under conditions similar to those used in a CAG line. Calculations were performed using a licensed module for General Steels (database ver. 8.0). The effects of temperature and time on grain size were considered. In addition, after obtaining the approximate volume fraction of austenite by dilatometry, carbon content in austenite (Cγ) was determined dividing the carbon content in steel (Cst) by the volume fraction of austenite (Vγ) formed during the intercritical annealing according to Cγ = Cst/Vγ [18], (being 0.254%). This concentration was used to recalculate the CCT and TTT diagrams. However, considering that the presence of carbide-forming elements such as V, Nb, and Ti changes the effective carbon content in the γ-phase, and other factors such as the presence of pre-existing phase boundaries and redistribution of Mn and Si, which are not considered in the models used in the software, and which may affect the decomposition of intercritical austenite, interpretation of results obtained based merely on these diagrams is still difficult.

2.2. Fabrication and Processing

Once obtaining theoretical CCT and TTT diagrams that allow obtaining multiphase steels, the experimental steel was produced at laboratory scale. Fabrication was done by conventional fusion and casting techniques. Ingots with dimensions of 10 cm width × 10 cm length × 15 cm height were subjected to homogenization heat treatment at 1100 °C for 1 h. For further processing, ingots were sectioned to obtain samples of 2.5 cm width × 2.5 cm length × 10 cm height. Samples were reheated at 1100 °C for 10 min and processed by hot rolling at temperatures between 1100 °C and 900 °C (T > Ac₃) to obtain hot-rolled steel strips of 2.3 mm thickness, which were air-cooled. Hot-rolled samples were pickled and cold-rolled to obtain thin steel sheets of 1.0 mm thickness.

2.3. Heat Treatments

Cold-rolled samples were then subjected to a final heat treatment simulating the conditions of a CAG process (where annealing is conducted at temperatures between 800–875 °C, and galvanizing stage is performed at temperatures between 425–475 °C, both during short processing times). Thermal cycles were carried out in a LINSEIS L78 quenching dilatometer. The temperature was measured and controlled by a type K thermocouple (0.38 mm in diameter), welded to the sample surface. The length change was measured by an LVDT sensor through a quartz push-rod. Experiments were performed under vacuum. Heating was performed using an inductive coil and cooling was conducted with He. Dilatometric measurements performed as a function of temperature allowed the determination of the critical transformation temperatures on continuous heating (Ac₁, Ac₃) and cooling (M_s, M_f). Ac₁ and Ac₃, represent the temperature for the start and completion of ferrite to austenite transformation, while M_s and M_f are the starting and ending temperatures for the austenite transformation to martensite. Once the critical temperatures were determined, cold-rolled samples were subjected to thermal cycles considering the next information: (i) annealing stage at temperatures within the intercritical region (Ac₁ < T < Ac₃), and (ii) isothermal bainitic treatment (IBT) at temperatures above the start temperature for martensite (T > M_s). These thermal treatments were carried out to modify the microstructure of the cold-rolled material to produce multiphase steels. Intercritical annealing was intended to produce ferrite + austenite, aiming that the later phase transformed into bainite during IBT and to martensite during final cooling to achieve microstructures consisting of a mixture of ferrite + bainite + martensite, or ferrite + bainite + austenite, depending on the carbon enrichment achieved during IBT.

Firstly, a dilatometric analysis was conducted to determine the temperature needed to obtain a proportion of about 50% ferrite + 50% austenite during heating. From the results obtained, the conditions of the intercritical annealing were set constant (817.5 °C for 15 s). Heating rate (20 °C/s), cooling rate from annealing to IBT (15 °C/s), and cooling rate from IBT to room temperature (2 °C/s) were also kept constant to further investigate the effects of IBT. To this end, samples were subjected to different times (30, 60, 120) and temperatures (425 °C, 450 °C and 475 °C), considered similar conditions to those used in a CAG process. Figure 1 shows a schematic representation of the thermal cycles and the processing variables used to produce multiphase steels.

2.4. Characterization

The contents of sulfur and carbon in the experimental steel were determined in a simultaneous carbon/sulfur LECO CS230 analyzer by infrared absorption spectroscopy according to the procedures of ASTM E-1019. The concentration of other elements was determined by optical emission spectrometry using a Spectro spectrometer model LabS according to standard ASTM E-415.

Table 1. Chemical composition of the experimental steel (in wt.%).

Element	Fe	C	Mn	Cr	Mo	Si	Nb	S	P
Composition	Balance	0.14	1.66	0.49	0.476	0.725	0.04	0.004	0.002

gives the chemical composition of the experimental steel.

Heat-treated samples were characterized by electron backscattered diffraction (EBSD). This technique was used as a tool for identification and quantification of phases, using image quality (IQ) values and the multi-peak model. IQ values are highly sensitive to experimental factors, thus its application to identify and quantify phases requires the normalization of such values in order to reduce these errors and increase the comparability of the measured values [19,20,21]. Two IQ-EBSD maps (75 µm × 75 µm) were obtained for each condition of IBT.

The image quality (IQ) is proportional to the sharpness of the Kikuchi Pattern obtained by EBSD, which is associated with the presence of crystalline defects. An elastically distorted lattice results in smeared Kikuchi Patterns and low image quality values, but a highly distorted lattice will have low IQ values [22]. Normalization of IQ values minimizes these misleading effects, such as assuming that the operating factors have an equal effect on each constituent in the microstructure. Normalization of IQ values can be calculated as [22]:

$$I{Q}_{Normalized}=\frac{I{Q}_{Initial}-I{Q}_{Min}}{I{Q}_{Max}-I{Q}_{Min}} x 100$$

(1)

where $I{Q}_{Max}$and $I{Q}_{Min}$, represent the maximum and the minimum IQ values in the scanning set, respectively. This calculation allows comparing microstructures with the same components, but different volume fractions thereof, without the need of information of a standard specimen [22].

$$N={\displaystyle \sum }_{i=1}^k{n}_i$$

(2)

$$IQ\cong {\displaystyle \sum }_{i=1}^{k}ND \left(n_i,{\mu }_i,{\sigma }_i\right)$$

(3)

$$Min \left(k\right)$$

(4)

$$\left|IQ-{\displaystyle \sum }_{i=1}^{k}ND \left(n_i,{\mu }_i,{\sigma }_i\right) \right| \le \epsilon$$

(5)

where $N$ is the number of the total scan points; $k$ is the number of normal distributions; ε is the minimum acceptable error; $ND \left(n_i,{\mu }_i,{\sigma }_i\right)$ is the ith normal distribution with the total data number of $n_i$, the group mean value of ${\mu }_i,$ and the standard deviation of ${\sigma }_i$ [22]. Adjusting every normal distribution and satisfying the Equations (3) and (5) produces a minimum difference between the initial IQ distribution and the sum of all single normal distributions, as shown in the Equation (5) [22]. In this procedure, each of these normal distributions represents one component in the microstructure [22].

Samples for EBSD analysis were cut using a Buehler IsoMet low speed precision saw, and grinded using sandpapers Struers No. 240, 340, 500, 800, 1200, and 2400. Samples were then polished in four steps starting with diamond paste with abrasive particle sizes of 3 µm, 1 μm, and 0.25 μm, and finishing with colloidal silica with abrasive particle size of 0.05 µm. The confidence index (CI) considered to proceed with the EBSD analysis was equal or superior to 0.85, which represents an image with at least 85% of successfully indexed data. Samples with image quality (IQ) values less than 0.85% were polished again until a desired value was achieved (≥0.85). Voltage used for EBSD analysis was 20 kV.

The microstructure of both cold-rolled material and heat-treated samples was characterized by optical microscopy (OM) using an Olympus GX51 microscope. Images were acquired using the software QCapture Pro. Characterization by scanning electron microscopy (SEM) was conducted in a Jeol 6610LV microscope and in a high-resolution focused ion beam (FEI-Scios dual-beam) scanning electron microscope. Samples for microscopic observation were prepared by conventional metallographic techniques (including grinding and polishing), to achieve a mirror-like surface, and etched using 2% Nital. Mechanical properties (yield strength, YS; ultimate tensile strength, UTS; elongation to fracture, E_f) were determined by uniaxial tensile tests according to standard ASTM E-8. Miniature specimens, machined at scale from standard specimens, were used for the following reasons: (a) to conduct thermal cycles under controlled conditions (heating and cooling rates, temperature, time), (b) to perform an in situ phase transformation analysis and (c) to determine mechanical properties, which together allows determining the correlation between phase transformations that occur during thermal cycles and the resulting microstructure and mechanical properties, as a function of time and temperature of IBT.

3. Results and Discussion

3.1. Effects of Time and Temperature of Isothermal Bainitic Treatment on Phase Transformations and Microstructure

Before evaluating the microstructural changes caused by thermal cycles, it is important to know the microstructural characteristics of the cold-rolled material. Figure 2 shows the microstructure of the cold-rolled steel sheet before being subjected to heat treatments. In general, the microstructure of the cold-rolled steel is characterized by a uniform distribution of phases along the steel thickness (Figure 2a), which according to Figure 2b, correspond to pearlite (P), bainite (α_B), and martensite (α’) in a ferrite (α) matrix.

Figure 3 shows the CCT and TTT diagrams were calculated with the chemical composition of the experimental steel (Table 1), and recalculated, considering the amount of carbon in austenite during intercritical annealing. CCT diagrams show the effect of cooling rates on the austenite decomposition during continuous cooling conditions (Figure 3a). Cooling curves represented by the black lines coming from 805°C (temperature at the proportion of 50% ferrite + 50% austenite), from left to right of the diagram, correspond to cooling rates of 100 °C/s, 10 °C/s, 1 °C/s, 0.1 °C/s, and 0.01 °C/s, respectively. Purple, magenta, and brown lines correspond to start, 50% and 90% of the martensite transformation. Blue-colored line indicates 1% of bainite transformation, while light green and dark green lines correspond to 1% of pearlite transformation and remaining austenite, respectively. The horizontal black lines indicate the transformation temperatures on heating, according to which, A₁ = 740.9 °C and A₃ = 853.7 °C. Hardness values (HRC) obtained from decomposition of austenite are also included in the diagram. For cooling rates between 100 °C/s and 1 °C/s, intercritical austenite will transform mainly into martensite with a hardness of about 47 HRC. For a cooling rate of 1 °C/s, a decrease in hardness is observed due to the formation of bainite, and this decrease is more significant for slower cooling rates due to the formation of pearlite.

TTT diagram, also called as isothermal transformation diagram, has a very important application like austempering, which is commonly employed in the industry for achieving specific microstructures and properties in steel. The diagram relates to the austenite evolution during isothermal transformation schedules for a given chemistry [23]. It is understood if steel cools down from austenitizing temperature to a transformation temperature and held constant during the completion of transformation. Black, green, and blue lines with markers in the TTT diagram of Figure 3b represent the start transformation curves for ferrite, pearlite, and bainite, respectively. Gray and black continuous lines represent the fraction of austenite that transforms into bainite. These diagrams were used to define the chemical composition of the experimental steel and have a better understanding of the phase transformations that can occur during heat treatments of steel.

It can be expected from these diagrams that, during annealing at temperatures between 740.9 °C and 853.7 °C, intercritical ferrite and austenite may coexist. During isothermal treatment at temperatures above M_s (about 300 °C), part of the intercritical austenite formed during annealing may transform into bainite. Finally, on final cooling, austenite may transform into martensite allowing the formation of multiphase steels containing ferrite + bainite + martensite. Depending on the carbon enrichment obtained during IBT, austenite can also be retained at room temperature, favoring the development of multiphase microstructures constituted of ferrite + bainite + martensite + austenite.

Some authors have demonstrated that, at the early stage of the ferrite to austenite phase transformation (during 30 s of annealing at 800 °C), substitutional alloying elements, such as Mn and Si, are hardly partitioned [24]; therefore, it is considered that the formation of austenite occurs mainly only with the partitioning of carbon. However, considering the limitations of JMatPro regarding the carbon enrichment or local variations in chemical composition, it becomes of vital importance to investigate the effects of time and temperature of IBT on the resulting microstructure and mechanical properties. Next, the effects of continuous heating and cooling on the variation in the dilation curves are presented. This to determine the critical transformation temperatures on heating and cooling for the experimental steel to further set the thermal cycles parameters under conditions similar to those used in a CAG process.

Figure 4a shows the dilation curve obtained during continuous heating of the cold-rolled steel. Two changes are observed, the first between 665 °C and 725 °C (stage I) and the second between 748 °C and 900 °C (stage II). Above 900 °C, the dilation curve presents a linear behavior with temperature, which suggests that austenite is the stable phase above such temperature. Some authors have investigated the effect of heating rate on the behavior of recrystallization and austenitization of cold-rolled steels. They observed that during recrystallization of steel, there was a change in the dilation curve, characterized by a linear behavior between ΔL and temperature [25]. Temperature required for such process increase with the increase in the heating rate, which was attributed to less time available for the nucleation process. These researchers reported that for a low carbon steel, recrystallization occurred between 650 °C and 725 °C (for a heating rate of 50 °C/s) prior to the ferrite– austenite phase transformation [25]. It is known that recrystallization temperature depends on several factors, such as chemical composition, grain size, heating rate, and plastic deformation level [25]. In the present work, considering that experimental steel was subjected to cold rolling before heating, and according to the findings of other works [25], the first change observed in the dilation curve at temperatures between 665°C and 725 °C (Figure 4a) can be attributed to recrystallization of the cold-rolled steel. Furthermore, considering that, in the present work, heating rate was relatively fast (20 °C/s) and recrystallization occurred in ΔT = 60 °C, it can be concluded that the recrystallization process happened in a very short time (3 s). The second change observed in the dilation curve on heating of steel, which was observed at temperatures between 748 °C and 900 °C, is characterized by a slight contraction on heating (stage II in Figure 4a). Some authors reported a similar behavior for a hypo-eutectoid steel at temperatures between 760 °C and 840 °C [26]. Such change in the dilation curve was attributed to the atomic rearrangement that occurs as a result of the allotropic change in steel [26]. The volumetric change during the ferrite– austenite phase transformation can be calculated from the crystallographic data using a = 4r/√3 and a = √8r, for body-centered cubic (BCC) and face-centered cubic (FCC) systems, respectively [27], with a being a side of the unit cell and r the atomic radius of iron. The volume in a unit cell for BCC iron before transformation is 0.023467 nm³, which is occupied by two atoms of Fe, since there are two atoms for the unit cell in a BCC crystalline structure [27]. The volume in FCC iron is 0.046307 nm³, but this volume is occupied by four atoms of iron. Therefore, two BCC cells should be compared (with a volume of 0.046934 nm³) with each FCC cell. The volume percentage change during the transformation is given by ((0.046307 nm³ – 0.046934 nm³)/0.046934 nm³) × 100, which is equal to −1.34% [27]. This indicated that iron contracts when heating 1 – 0.0134 = 0.9866 cm³ after phase transformation [27], which can explain the contraction observed in the dilation curve (stage II), and can be attributed to the ferrite –austenite phase transformation that occurs on continuous heating. As can be seen in Figure 4a, the transformation of austenite from ferrite (F) + second phases (SP) does not occur uniformly; it is firstly slow, and then speeds up around 790 °C. Some authors have reported that the first stage consists on the formation of austenite from the dissolution of carbides between the temperatures Ac_1s (dissolution starts) and Ac_1f (dissolution finishes), and the second stage consists on the transformation of ferrite into austenite between the temperatures Ac_1f and Ac₃. The zone between Ac_1s and Ac₃ is what is known as the intercritical zone related to the Fe–Fe₃C phase diagram [28].

The volume fraction of transformed austenite from second phases (SP) or ferrite phase can be determined from the dilation curves and the lever rule method, as shown in Figure 4b. To this end, two tangent lines need to be traced parallel to the linear part of the dilation curve, followed by tracing vertical straight lines to intersect such tangents; the total length of the vertical lines will be defined by the segment AC (Figure 4b). The intersection with the dilation curve will define segments AB and BC, which, according to the lever rule method, will allow determining the transformed austenite fraction and remaining ferrite + second phases or single ferrite through AB/AC*100 and BC/AC*100, respectively, as shown in Figure 4b. Above 900 °C, a linear behavior between ΔL and temperature is observed, which suggests that the austenite phase transformation has been completed and therefore austenite will be the stable phase above this temperature.

Figure 5a shows the evolution of ferrite (F), second phases (SP), and austenite (A), as determined from the dilation curve and the lever rule method during the continuous heating of the cold-rolled steel. Ac₁ and Ac₃ critical transformation temperatures are about 748 °C and 900 °C, respectively, which suggests a variation between 7.1 °C and 46.3 °C, with respect to the values calculated with the software. The proportion of ferrite and second phases decrease, and the austenite increases with the increase in temperature. This behavior is associated with the ferrite/second phases–austenite transformation that occurs on continuous heating.

Figure 5b shows the dilation curve obtained on continuous cooling from the austenite phase field to room temperature. A slope change characterized by an expansion on cooling is observed between 403 °C and 257 °C. This behavior is related to the austenite– martensite transformation [29]. It is known that the volume of martensite is higher than the one of austenite, therefore, during the progress of the martensitic transformation, at temperatures below M_s, an expansion is observed with the decrease in temperature, which is associated with a larger amount of martensite [29]. At temperatures below M_f, a linear behavior is observed in the dilation curve ΔL vs. temperature, which indicates that the steel does not undergo any additional microstructural change. This result opens the possibility to promote the bainite transformation at temperatures similar than the ones used in CAG lines, which in addition to the possibility of forming ferrite + austenite during annealing at T ≤ 890 °C, could favor the development of multiphase steels under similar conditions to the ones used in an industrial process. From the results of Figure 1, Figure 2, Figure 3 and Figure 4, the intercritical annealing was set at 817.5 °C for 15 s to further investigate the effects of time (30 s, 60 s, 120 s) and temperature (425 °C, 450 °C, 475 °C) of the isothermal bainitic treatment (above M_f, Figure 5), and to evaluate the feasibility to obtain multiphase steels under the conditions mentioned before.

Figure 6, Figure 7 and Figure 8 show the dilation curves obtained during intercritical annealing, isothermal bainitic treatment and final cooling. All of them were obtained after 120 s of IBT for temperatures of 425 °C, 450 °C, and 475 °C.

The changes in the dilation curves during the intercritical annealing are shown in Figure 6. As can be seen, during isothermal annealing at 817.5 °C, there is a contraction of steel. As mentioned before, this contraction is associated with the ferrite to austenite transformation. Two changes in the slope of the dilation curve are observed during isothermal annealing, the first one is characterized by a rapid reduction of ΔL during the first three or four seconds of holding, but from that time on, the magnitude of change in ΔL with time is less significant. It appears then that the formation of austenite during intercritical annealing occurs mainly during the first seconds of annealing, but then the rate slows down making its formation difficult.

Some authors investigated the effect of heating rate on the austenite formation in low-carbon, high-strength steels annealed in the intercritical region [30]. They also report the volume fraction of austenite formed during the continuous heating to intercritical temperatures, and the effect of the heating rate on the volume fraction of austenite formed during isothermal holding as a function of intercritical annealing temperature and holding time [30]. The chemical composition of the investigated steel was C: 0.08%, Mn:1.9%, Mo+Cr+Si: 0.06%, Nb:0.010%, Al:0.045%, N:0.006%. The temperature required to achieve a 50% ferrite + 50% austenite on continuous heating was about 817 °C and 827 °C for heating rates of 50°/s and 10 °C/s, respectively [30]. The former is similar to the one obtained in the present work for a heating rate of 20 °C/s. The amount of austenite formed isothermally at 820 °C after 15 s of holding (similar to the time used in the present work) was about 4.93% and 6.25% for heating rates of 50 °C/s and 10°C/s, respectively [30]. As observed in Figure 6a, the isothermal transformation of austenite at 817.5 °C is very slow, especially after the first three or four seconds, which also suggests that the evolution of austenite is slow. The amount of transformed austenite changes from 50% to about 52.5% and 57.5% when the temperature of annealing is increased from 817.5 °C to 820 °C and 825 °C, respectively. Therefore, considering that: (a) the effect of time is less significant than the one caused by temperature, (b) the time used for annealing is short, and (c) other authors obtained amounts of austenite around 5% for similar conditions of IBT, the amount of austenite formed during intercritical annealing is expected to be about 5%, which could be reasonable considering the results reported elsewhere [30].

Figure 7 shows the part of the dilation curves obtained during cooling from intercritical annealing to isothermal bainitic treatment at 425 °C (Figure 7a), 450 °C (Figure 7b), and 475 °C (Figure 7c). A slope change is observed in the three ΔL vs. time curves, which according to the ΔL vs. temperature curves, occurs at temperatures about 710 °C and 588 °C, respectively. Some authors reported the dilation curves obtained as a function of cooling rate (0.1 °C/s, 10 °C/s, and 100 °C/s) after intercritical annealing at 800 °C in a steel containing C: 0.15%, Mn:1.906%, Si:0.26%, Cr:0.413%, Ti:0.044%, and B:0.0010% [31]. The dilation curve reported from intercritical annealing during continuous cooling at 10 °C/s (similar to the one used in the present work from annealing to IBT) exhibited a similar change in the slope at about 674 °C and 576 °C, which was attributed to the formation of ferrite (Fs) and bainite (Bs) during cooling, respectively [31]. Therefore, the result obtained in the present work suggest that part of the intercritical austenite may transform into ferrite at about 710 °C (Fs) and bainite at around 588 °C (Bs); according to Figure 7, the time available for the formation of ferrite on cooling is about 4 s. Other authors reported a comparative study between theoretical diagrams calculated with JMatPro, and experimental CCT diagrams constructed from the results of dilatometric experiments [32]. They found that the experimental bainite transformation curves were presented at higher temperatures than the ones predicted by the software [32]. The difference in the austenite transformation following intercritical heat treatments was related to the nucleation mechanism of pro-eutectoid (new) ferrite at the pre-existing phase boundaries. Pre-existing interfacial surfaces such as ferrite–austenite (α–γ) anticipate the formation of bainite [32]. These results are consistent with the differences observed in the present work between results obtained by dilatometry and the CCT diagrams calculated with the software.

Figure 8 shows the variation in the magnitude of the dilation during IBT after 120 s, at 425 °C, 450 °C, and 475 °C. It is clear that the expansion due to the bainite transformation is more significant at 425 °C, but it decreases with the increase in IBT temperature to 450 °C (Figure 8a,b). Similar results have been obtained in other works [33]. For instance, some authors conducted the isothermal bainitic treatment at temperatures of 365 °C, 380 °C, 400 °C, 420 °C, and 430 °C during times up to 1800 s [33]. They found that the increase in the IBT temperature caused a decrease in the magnitude of the dilation [33]; the higher dilation was observed at 365 °C, which was attributed to a higher formation of bainite. The amount of bainite increases, and the kinetics of bainite transformation is accelerated with a decrease of holding temperature [33]. Carbon enrichment determines the difference in Gibbs free energy, if the difference between initial and final carbon enrichment is lower, the lower is difference between Gibss free energy and, hence, the smaller amount of bainite would be created. This behavior is observed in Figure 8 when IBT temperature is increased from 425 °C to 450 °C. The increment of bainite amount with a decrease of holding temperature can be explained by the $T_o^{\prime }$ theory [34], which follows that the maximum amount of bainite that can be obtained at any temperature is limited by the fact that the carbon content of the residual austenite must not exceed the $T_o^{\prime }$ curve of the phase diagram [35]. The driving force for the formation of new plates decreases as the carbon concentration in the untransformed austenite approaches the $T_o^{\prime }$ composition, at which the free energy of bainite and austenite phases become identical. An opposite behavior is observed when IBT temperature is increased up to 475 °C, in this case, it is observed a contraction in the dilation curve, the absence of expansion suggests that the transformation of bainite does not occur during the isothermal holding (Figure 8c), since even a small fraction of bainite could lead to an expansion in the dilation curve. Although bainite can be formed during cooling from intercritical annealing to IBT as shown in Figure 8c, the results suggest that during isothermal treatment at 475 °C, the Gibbs energy between austenite and bainite is in equilibrium or near-equilibrium conditions, and as a result, the transformation stops. The contraction observed at 475 °C is not easy to explain since at this temperature, the formation of pearlite or ferrite becomes difficult. Other authors investigated the effects of temperature (350–470 °C) and time (50–600 s) of IBT on the bainite transformation by dilatometry [36]. They found that, when IBT was conducted at 470 °C, the bainite transformation rate was significantly reduced, and observed that the dilatation diminishes to zero as the temperature is raised toward the BS temperature [36]. This information supports the result obtained in the present work regarding the absence of the bainite transformation at 475 °C, however, the contraction observed in Figure 8c is finally explained with the temperature vs time plot experimentally obtained, which is shown in Figure 9. As can be seen in this figure, there is a gradual decrease in the temperature during the beginning of the isothermal bainitic treatment, which is observed for the three temperatures investigated. The higher the temperature of IBT, the higher the difference between the target and the experimental temperature, and the longer the time needed to achieve the desired temperature. Therefore, apparently, the contraction observed in Figure 8c is influenced by both the absence of isothermal bainite transformation and the decrease in the steel temperature to achieve the desired IBT temperature. This observation is supported by the expansion observed in samples with IBT at 425 °C and 450 °C, even with the decrease in the temperature observed in Figure 7c, which was related to the occurrence of the bainite transformation.

Figure 10 shows the dilation curves obtained on final cooling after 120 s of isothermal bainitic treatment. It is observed that steel experiments another change in the dilation curve during cooling from IBT to room temperature, which is attributed to the austenite to martensite phase transformation. The time required for this transformation does not show any significant change since this transformation occurs by a displacive mechanism [37]. The expansion seems to be slightly higher in samples subjected to 475 °C, which could be related to a higher amount of austenite after IBT.

According to results of dilatometry, it can be concluded that, during heating of steel above the critical transformation temperature Ac_1s, ferrite + second phases transform into austenite causing a contraction under isothermal conditions during annealing (first phase transformation). This allows the coexistence of ferrite and austenite during annealing. Austenite transforms into pro-eutectoid ferrite and bainite during cooling from intercritical annealing to IBT, and the bainite transformation progresses during IBT, except for the IBT temperature of 475 °C, where the transformation stops. During final cooling, austenite transforms to martensite, but a certain amount can be also retained as reported elsewhere [38].

Figure 11 shows the microstructure of heat-treated samples obtained by scanning electron microscopy as a function of IBT temperature after 120 s. As can be seen, for all the conditions investigated, a fine-grained microstructure is observed. All of them show intercritical ferrite (α_Int), bainite (α_B), pro-eutectoid (new) ferrite (α_N), retained austenite (RA), and martensite (α’). This later has been reported to have certain amount of austenite, and is sometimes identified as a mixture or martensite/austenite [39]. Unlike the ferrite morphology in the cold-rolled material, the heat-treated samples show ferrite grains with equiaxed morphology, confirming that recrystallization of the cold-rolled material occurred during heating of the steel, as shown previously in the dilatometric analysis (Figure 4a). In the simplest terms, bainite can be defined as a non-lamellar aggregate of lath- or plate-shaped ferrite and carbide [39]. The microstructural characteristics vary with composition and temperature of transformation. When the cementite formation is prevented, the microstructure mainly consists of lath-like ferrite and carbon-enriched residual austenite or martensite (M/A) constituents on the lath boundaries [39]. When diffusion of carbon becomes slower, some the carbon is precipitated as fine carbide particles inside the ferrite plates. The remaining carbon escapes into the austenite and may precipitate as interplate carbide. It is possible to suppress the carbide component of bainite in steels with sufficient concentration of alloying elements, such as Si or Al [39]. The microstructure and other features of this carbide-free variety are then very similar to the one with carbides, but it consists of an aggregate of ferrite plates and untransformed austenite [39].

Martensite in ferrous alloys exhibits various morphologies, such as chiefly lath, lenticular, and thin plate, depending on chemical compositions and Ms temperature [37]. Lath and lenticular are the two major morphologies of α′ martensite. Lath martensite is formed in Fe-C (<0.6% C), Fe-Ni (<28% Ni), and Fe-Mn (<10% Mn) alloys, and most heat-treatable commercial steels, and has overwhelming industrial significance because it is a basic structure of high strength steels. Lenticular martensite appears in Fe-high C (0.8–1.8% C) and Fe-high Ni (29–33% Ni) alloys. The other three α′ martensites are not common in ferrous alloys [37]. Lath martensites have a tendency to become aligned parallel to one another. The current view held is that the austenite grain is divided into packets (a group of parallel laths with the same habit plane) and that each packet is further subdivided into blocks (a group of laths of the same orientation, i.e., the same variant of the K-S orientation relationship) [37].

Phases shown in Figure 11 have been identified based on the information described above and considering their similarities with the microstructural characteristics reported in other works [39,40,41,42].

Figure 12 shows EBSD-IQ maps of samples subjected to 120 s of IBT at three different temperatures: 425 °C, 450 °C, and 475 °C. The corresponding analysis by the multi-peak method using normalized IQ values is also presented. As can be seen, all samples present four individual distributions related to the presence of: ferrite (distribution with higher IQ values), bainite (distribution with intermediate IQ values), martensite (distribution with lower IQ values), and retained austenite (with IQ values between martensite and bainite, as reported elsewhere). A similar behavior was observed for the other IBT conditions. The phases quantification is presented in

Table 2. Percentage of phases and microconstituents as a function of IBT time/temperature (vol. %).

IBT Conditions		Phases/Microconstituents
Temperature (°C)	Time (s)	Ferrite (vol. %)	Bainite (vol. %)	Martensite (vol. %)	Austenite (vol. %)
425	30	46.6	20.40	30.83	2.17
	60	51.4	26.15	18.89	3.56
	120	48.5	31.30	15.51	4.69
450	30	47.0	17.10	34.00	1.90
	60	51.7	21.20	25.24	1.86
	120	48.2	23.8	24.44	3.56
475	30	45.6	9.10	43.62	1.68
	60	48.2	10.4	39.64	1.76
	120	53.3	9.75	35.24	1.71

. In general, it is observed that, for a specific IBT temperature, the amount of bainite increases with increasing IBT time. For a specific IBT time, the amount of bainite is higher in the sample with the lowest IBT temperature (425 °C). Increasing the IBT temperature causes a decrease in the amount of bainite obtained, being the lowest values obtained in samples with higher IBT temperature, which, as explained before, is related to the reduction in the Gibbs free energy difference between austenite and bainite. The low amount of bainite in samples with IBT at 475 °C can be mainly attributed to the austenite– bainite phase transformation that occurs during cooling from intercritical annealing to IBT. The presence of retained austenite can be related to carbon enrichment that occurs due to the formation of pro-eutectoid ferrite and bainite. The retained austenite fraction could be obtained directly from the phase fraction either by EBSD or X-ray diffraction [43]. Regarding EBSD-IQ analysis, some authors reported the presence of ferrite, bainite, retained austenite, and martensite after annealing at 770 °C followed by cooling to 450 °C at 15 °C/s and air cooling to room temperature, after a holding time of 120 s at 450 °C. They quantified the above-mentioned phases and microconstituents by EBSD-IQ. To this end, they normalized the IQ values and used the multi-peak model [44]. EBSD-IQ values for retained austenite were reported to be between martensite and bainite [44] (see the figure below). It has been also reported that both bainitic ferrite and its associated retained austenite can have high dislocation densities [45]; the latter can only be attributed to plastic relaxation effects [45]. According to these observations, phases in micro-constituents were quantified by EBSD-IQ as shown in Figure 12.

According to the CCT diagram calculated for the experimental steel and considering the cooling rate from intercritical annealing to IBT (10 °C/s), it was expected that the microstructures obtained from the thermal cycles contained a higher amount of martensite and a lower amount of bainite than the one that was obtained experimentally. As mentioned before, some authors found similar results, and differences were attributed to carbon enrichment and pre-existing interfacial surfaces such as ferrite–austenite (α–γ), which anticipate the formation of bainite [32].

3.2. Effects of Time and Temperature of Isothermal Bainitic Treatment on Mechanical Properties

The mechanical properties of the cold-rolled steel prior to being subjected to thermal cycles were: ultimate tensile strength (UTS) = 1419.2 MPa, yield strength (YS) = 1349.9 MPa, and elongation to fracture E_f = 5.5%. As observed in Figure 2, the microstructure of the cold-rolled material consists of a mixture of perlite, bainite, and martensite in a ferrite matrix. Martensite is considered as a hard, brittle phase. Additionally, it is well known that during deformation of steel at room temperature, part of the mechanical energy used for the deformation process is lost as heat, but other amount remains in the steel as stored energy as lattice defects (for example: dislocations). The greater the number of dislocations, the greater the strength of the material and the lower the elongation at fracture. Therefore, the high mechanical strength and low elongation at fracture of the cold-rolled material is consistent with the microstructural characteristics in such condition.

Figure 13 shows the mechanical properties obtained as a function of time and temperatures of the isothermal bainitic treatment; the changes in mechanical properties are related to the changes in the microstructure caused by the heat treatment. In general, the tensile strength and yield strength of heat-treated samples are lower than those of cold-rolled material (Figure 13a,b), but elongation to fracture is higher than that observed in the cold-rolled material (Figure 13c). This behavior is related to the reduction in the number of dislocations as a result of recrystallization and to the variation in the proportion of second phases. The highest values for these properties were obtained in samples with the lowest IBT temperature (425 °C), while the lowest values were obtained in samples with the highest IBT temperatures (475 °C) (Figure 13). Changes in mechanical properties are related to changes in the microstructure caused by heat treatments. Samples with IBT = 475 °C have the lower elongation to fracture, which can be mainly attributed to a higher amount of martensite, and the lower amount of bainite and retained austenite, compared to samples thermally treated at 425 °C and 450 °C.

The best combination between ultimate tensile strength and elongation to fracture is obtained in samples with IBT at 425 °C, which indicates that the presence of a ferrite matrix along with a higher amount of bainite/retained austenite, and a lower amount of martensite (compared to samples thermally treated at 450 °C and 475 °C) enhances work hardening. This result is very similar to the one observed in DP (martensite + ferrite) against TRIP (ferrite + bainite + retained austenite) steels, where TRIP steels have better combination between work hardening and elongation to fracture when comparing with DP steels [46]. Bainite has excellent toughness, and retained austenite can provide good values of ductility along with high work hardening when it is presented with appropriate stability against mechanically induced martensitic transformation [38]. Martensite, in contrast, is known for being a strong, hard, and brittle phase [37].

4. Conclusions

Cr+Mo+Si low-alloyed low carbon multi-phase steels were produced under conditions similar to those used in continuous annealing and galvanizing lines, with microstructures consisting of a mixture of ferrite, bainite, martensite and retained austenite.

The microstructure of the deformed steel recrystallizes during heating of steel. Subsequent transformation to austenite occurs both during heating and during annealing in the two-phase region. On cooling from intercritical annealing to IBT, austenite transforms into pro-eutectoid ferrite and bainite. The austenite to bainite transformation progresses during IBT, and during final cooling, austenite transforms to martensite, but a certain amount is also retained at room temperature.

The austenite–bainite phase transformation that occurs during IBT, is more significant for the lowest IBT temperature (425 °C), but the rate of this transformation decreases with the increase in temperature to 450 °C, which is attributed to the reduction in the Gibbs free energy between austenite and bainite with the increase in temperature. During IBT at 475 °C, this transformation is inhibited, suggesting that under these conditions, the Gibbs free energy difference between austenite and bainite is in equilibrium or near equilibrium conditions.

The mechanical properties of the investigated steels strongly depend on the microstructural characteristics resulting from thermal treatments. Samples subjected to IBT = 425 °C had the highest work hardening and highest ductility, while samples with IBT = 475 °C had the lowest work hardening and lowest elongation to fracture. The best combination between strength and ductility is obtained in samples with the highest amount of bainite and retained austenite, and the lower amount of martensite. Apparently, carbon enrichment in austenite during IBT due to the occurrence of bainite increases austenite stability against mechanically induced martensitic transformation.