HEAT TREATMENT OF HIGH-STRENGTH COLD-ROLLED STEEL STRIP

Abstract

A heat treatment of a high-strength, cold-rolled steel strip or band, comprising the steps of: a) soaking a cold-rolled steel strip or band, b) cooling the soaked steel strip or band, c) heat-treating the cooled strip or band; d) cooling the heat-treated steel strip or band to the ambient temperature range; such that the steel strip or band has a microstructure comprising various ferrites, retained austenite, and martensite. The main components of the steel composition comprise carbon, manganese, silicon, and aluminum in addition to iron.

Description

HEAT TREATMENT OF HIGH-STRENGTH COLD-ROLLED STEEL STRIP OR BANDS FIELD OF INVENTION The present invention relates to a method of heat treatment of a high-strength cold-rolled steel strip or band. Several types of cold-rolled steels and manufacturing processes have been proposed in the art to meet the requirements of automotive applications. For example, extra-low carbon steel is used in automotive steel strips due to its formability. This type of steel exhibits a tensile strength in the range of 280–380 MPa. HSLA (high-strength low-alloy) steels contain microalloying elements. They are hardened through a combination of precipitation hardening and grain refinement. Advanced high-strength steels (AHSS), such as dual-phase (DP) steels and transformation-induced plasticity (TRIP) steels, are currently typical of the high-ductility, high-strength steels used in the automotive manufacturing industry. In DP steels, the presence of martensite within a ferrite matrix allows for tensile strengths exceeding 450 MPa combined with good cold formability. BACKGROUND OF THE INVENTION To simultaneously achieve a high yield strength-to-tensile strength ratio and even higher tensile strength, i.e., above 800 MPa, steels with complex microstructures (CP) have been developed, including ferrite, bainite, martensite, and / or retained austenite. However, due to the difference in deformation capabilities between ferrite, bainite, or martensite structures and retained austenite structures, these steels generally have inferior flange or wing formability under stretch. Therefore, their use is limited to automotive parts that do not require high formability. TRIP-type tempered martensitic steel (Q&P steel, produced by quenching and parting), consisting of hardened martensite as the matrix phase and residual austenite, and TRIP-type bainitic ferrite steel (TBF steel, produced by austempering or bainitic quenching), consisting of bainitic ferrite as the matrix phase and residual austenite, offer advantages such as high strength due to their hardened bainitic ferrite and / or martensite structure, and outstanding elongation because the matrix is ​​carbide-free, and fine residual austenite grains can easily form at the bainitic ferrite boundary in a lattice-like structure. Therefore, carbide-free bainitic ferrite or tempered martensitic steels are expected to achieve good flanging or draw-flanging capabilities due to their uniform lattice or fine mesh structure.The hardness heterogeneities due to the presence of only a small amount of martensite in these microstructures will allow these types of steel to achieve good drawing or deep drawing capabilities. However, due to the limitations of current continuous production lines, the expected beneficial combination of strength and ductility properties could not be achieved with currently available steel recipes. These limitations include, among other things, that the The QCRC Ln / Lznz / E / YILI reheating furnaces of current continuous annealing (CA) and continuous galvanizing (CG) lines are often only suitable for subjecting steel strips or bands to intercritical or recrystallization heat treatment. For example, in some current annealing lines, the maximum annealing temperature is limited to 890°C. Furthermore, the cooling rates in current CA / CG lines are limited within a fixed range. In addition, the overaging or excess time available for many CA / CG lines is limited; for example, this time frame is less than approximately 160 seconds, which places significant time constraints on the completion of any desired transformation during overaging. For example, document WO2013 / 144373A1 has disclosed a cold-rolled TRIP steel with a polygonal ferrite matrix having a specific composition comprising chromium and a particular microstructure, and possessing a tensile strength of at least 780 MPa. This steel is said to enable its production on a conventional industrial annealing line that includes an overaging / austempering section. Specifically, at a relatively high overaging / austempering temperature, the austempering or quenching time can be less than 200 seconds. Documents EP2831296B1 and EP2831299 describe TBF steels with a tensile strength of at least 980 MPa that could also be produced on a conventional production line. However, the preferred overaging / austempering or quenching times of 280–320 seconds are too long to allow production on a large number of conventional production lines. In other words, the bainitic transformation kinetics are too slow to complete the bainitic transformation within the limited time frame in the overaging section to obtain the required microstructure on a conventional production line. BRIEF DESCRIPTION OF THE INVENTION An object of the invention is to provide a cold-rolled steel strip or band having a desired combination of high tensile strength (TS) and high total elongation (TE) properties with a reasonable yield strength (YS), such as TS > 850 MPa; TE > 14% and / or YS > 500 MPa, in particular a steel strip or band for use in automotive applications, or a suitable alternative. Another object of the invention is to provide a method for the heat treatment of a cold-rolled steel strip or band to obtain the desired combination of properties as mentioned above, in particular a heat treatment that can be carried out using existing production lines, or a suitable alternative. Another object of the invention is to provide a high-silicon cold-rolled steel strip having a desired combination of properties, which can be manufactured on conventional industrial production lines. Another object of the invention is to provide a steel composition for a high-strength cold-rolled steel strip and its heat treatment that allows the completion of the bainitic transformation in a conventional production line in order to obtain a desired microstructure. QCRC Ln / Lznz / E / YILI In view of this, the invention provides a method for heat treating a cold-rolled steel strip, which method comprises the following stages or steps: a) soaking a cold-rolled steel strip or band above (Ac3 - 60) for a soaking time t2 of 1 - 150 seconds, to obtain a cold-rolled steel strip or band having at least a partially austenitic microstructure; b) cool the soaked steel strip or band resulting from step a) to a temperature T4 in the range of Bn - Ms; c) heat treat the cooled strip obtained in step b) in a temperature range of Bs - T4 for a time period t5 of 30 - 300 seconds; d) cool the heat-treated steel strip or band to room temperature; so that the steel strip or band has a microstructure (in % vol.) comprising QCRC Ln / ίΖΠΖ / Β / ΥΙΛΙ polygonal ferrite (PF) + acicular ferrite (AF) + higher bainitic ferrite (HBF): 20 - 55; where the polygonal ferrite (PF): 0 - 45: lower bainitic ferrite (LBF): 20 - 65; retained austenite (RA): 5 - 20; martensite (M): 0-20; where the steel band or strip has a composition (in % by mass) comprising C: 0.15-0.35; Mn: 1.50-3.00; Yes: 0.50-2.00; Al: 0.01 - 1.50; P: less than 0.050; S: less than 0.020; N: less than 0.0080; where the sum (Si + Al) is 0.60; and optionally one or more elements selected between < Cr < 0.35; 0 < Cu < 0.20; < Ni < 0.50; 0< Mo<0.30; 0< Nb<0.10; 0< V<0.10; 0 < Ti < 0.10; < B < 0.0030; 0 <Ca< 0.0050; < REM < 0.0100, where REM is one or more rare earth metals; and the rest is iron and unavoidable impurities. The method of the invention allows the production of a cold-rolled steel strip or band that has a specific composition and microstructure, and a combination of properties desirable for automotive parts that require high strength, formability, and weldability. The invention solves the problem of slow bainitic transformation kinetics by introducing an adequate amount of pro-eutectoid ferrite and controlling its morphology, obtaining fine grains of austenite by controlling the temperature and upper annealing time, and using a modified overaging process in a production line. This method according to the invention can be carried out using existing continuous annealing and galvanizing lines within the limitations regarding the maximum temperature in the annealing section, cooling rate ranges, and overaging time window at production rates that are typical of these production lines. The heat-treated, cold-rolled steel strip can be coated with zinc, for example, by hot-dip galvanizing or electrogalvanizing. A hot-dip galvanizing stage can be easily integrated into the heat treatment according to the invention. The terms used to describe the critical transformation temperatures of a steel are given below, as any expert in the field knows. Ae3: Equilibrium temperature for the transformation of ferrite into austenite and of austenite into ferrite. Ac3: Temperature at which, during heating, the transformation of ferrite into austenite ends. Ac3 is usually higher than Ae3, but tends towards Ae3 as the heating rate approaches zero. In this invention, Ac3 is measured at a heating rate of 3°C / s. Ar3: Temperature at which austenite begins to transform into ferrite during cooling. Bs: Temperature at which, during cooling, the transformation of austenite into bainite begins. Bn: nose temperature of the bainitic transformation on the time-temperature transformation (TTT) curve of a steel, in which the transformation of austenite to bainite has the fastest kinetics. Ms: Temperature at which, during cooling, the transformation of austenite into martensite begins. Mf: Temperature at which, during cooling, the transformation of austenite into martensite is completed. A practical problem with Mf is that the fraction of martensite formed during cooling approaches the maximum achievable amount only asymptotically, meaning that martensite formation continues at very low temperatures. For practical reasons and in the context of this invention, Mf is therefore taken to be the temperature at which 90% of the maximum achievable amount of martensite has formed. These critical phase transformation temperatures can be determined by dilatometer experiments. Alternatively, the Ac3, Bs, and Ms points of the steel according to the invention QCRC Ln / ίZΖΠZ / Β / YΙΛΙ can be calculated beforehand based on their composition, using commercial software such as JmatPro, or using the following empirical formulas: Ac3(°C) = 942-260C + 35Si - 35Mn + 125AI - 11 Cr - 14Cu Bs(°C) =839-86Mn-23Si-67Cr + 35^AI-270( 1 -exp (-1.330)) Ms (°C) = 539 - 423C - 30.4Mn - 7.5SI + 30AI In these formulas, component X of the steel composition is represented in % by weight. In this specification all temperatures are represented in degrees Celsius, all compositions are given in weight percent (% wt) and all microstructures are given in volume percent (% vol), except where explicitly stated otherwise. BRIEF DESCRIPTION OF THE FIGURES In the attached figures: Figures 1a and 1b are an EBSD map showing the microstructure features of bainitic ferrite from a low-temperature bainitic ferrite (Figure 1a), and a high-temperature bainitic ferrite (Figure 1b), respectively. Figure 2 is a histogram of the disorientation angle of a low-temperature bainitic ferrite and a high-temperature bainitic ferrite. Figure 3 is a diagram showing a generally applicable time versus temperature profile of one modality of the method according to the invention. DETAILED DESCRIPTION OF THE INVENTION The following is an explanation of the composition, the steps or stages of the method, and the microstructure according to the invention. Composition Carbon: 0.15 - 0.35% A sufficient amount of carbon is required to strengthen and stabilize the retained austenite, which provides the TRIP effect. Therefore, the carbon content is greater than 0.15%, preferably greater than 0.17%, to ensure the required strength and elongation. Increasing the carbon content results in increased steel strength, a greater amount of retained austenite, and a higher carbon content within the retained austenite. However, the weldability of the steel is significantly reduced when the carbon content exceeds 0.25%. For applications requiring welding, the carbon content is preferably between 0.15% and 0.25%, and more preferably between 0.17% and 0.23%. Silicon: 0.50-2.00% Silicon is a mandatory element in the composition of the steel according to the invention to obtain the microstructure described. Its main function is to prevent carbon from precipitating as iron carbides (most commonly cementite) and to suppress the decomposition of residual austenite. Silicon contributes to the strength property and suitable transformation behavior. Furthermore, silicon contributes to improved ductility, work hardenability, and formability. QCRC Ln / Lznz / E / YILI is used to draw the flange or wing by restricting austenite grain growth during annealing. A minimum of 0.50% Si is required to sufficiently suppress carbide formation. However, a high silicon content results in the formation of silicon oxides on the strip surface, which impairs surface quality, coatability, and workability. Furthermore, the Ac3 temperature of the steel composition increases with increasing silicon content. This can affect the feasibility of producing the steel strip using existing production lines due to the maximum achievable temperature in the annealing section. Therefore, the silicon content is 2.00% or less. Preferably, the Si content is in the range of 0.80–1.80% in view of wettability in combination with suppression of carbide formation and promotion of austenite stabilization. More preferably, 1.00–1.60%. Aluminum: 0.01 - 1.50% The primary function of aluminum is to deoxidize liquid steel before melting. A content of 0.01% Al or more is required for deoxidation. Additionally, aluminum has a silicon-like function in preventing carbide formation and stabilizing retained austenite. Al is considered less effective than Si and has no significant effect on strengthening. Small amounts of Al can be used to partially replace Si, adjust critical transformation temperatures and cooling rates for acicular ferrite (AF), and accelerate bainitic transformation kinetics. Al is added for these purposes. Therefore, the Al content is preferably greater than 0.03%.High levels of aluminum can increase the ferrite-to-austenite transformation point to levels incompatible with current facilities, making it difficult to obtain a microstructure where the main phase is a low-temperature transformation product. The risk of cracking during smelting increases with increasing aluminum content. Therefore, the upper limit is 1.50%, preferably 1.00%, and more preferably 0.70%. Regarding the relationship between the proportions of Si and Al, the composition meets the condition Si + Al > 0.60, preferably Si + Al > 1.00. Advantageously, the Al content is less than 0.5 times the Si content. Manganese: 1.50 - 3.00% Manganese is required to obtain the microstructure in the steel strip according to the invention, given its hardenability and the stabilization of the retained austenite. Mn also affects the formation of pro-eutectoid ferrite at higher temperatures and the transformation kinetics of bainitic ferrite. A certain amount of Si and / or Al is necessary to suppress carbide formation in the bainitic ferrite. The Ac3 temperature increases as the Si and Al content increases. Mn is also adjusted to balance the elevated Ac3 phase transformation point resulting from the presence of Si and Al. If the Mn content is 1.50% or less, the described microstructure is difficult to obtain. Therefore, it is necessary to add Mn at 1.50% or more. However, if Mn is present in an excessive amount, macrosegregation is likely to occur, resulting in the formation of unfavorable bands in the steels. Furthermore, Excessive amounts of Mn lead to slow bainitic transformation kinetics, resulting in too much fresh martensite and, consequently, impaired flange or wing stretch formability. Therefore, the Mn content is 3.00% or less, preferably 2.80% or less, and more preferably 1.80 < Mn < 2.60%. Phosphorus: <0.050% Phosphorus is an impurity in steel. It segregates at the grain boundaries and reduces workability. Its content is less than 0.050%, preferably less than 0.020%. Sulfur: <0.020% Sulfur is also an impurity in steel. S forms sulfide inclusions such as MnS, which initiates cracking and impairs the flange's formability due to steel stretching. The S content is preferably as low as possible, for example, below 0.020%, preferably below 0.010%, and more preferably less than 0.005%. Nitrogen: <0.0080% Nitrogen is another unavoidable impurity in steel. It precipitates as nitrides with microalloying elements and is present in solid solution to contribute to strengthening. Excess nitrides impair elongation, drawability, and flexibility. Therefore, advantageously, the nitrogen content is 0.0080% or less, preferably 0.0050% or less, and more preferably 0.0040% or less. The steel composition may comprise one or more optional elements as follows: Copper: 0-0.20% Copper is not required in some steel compositions, but it may be present. In some forms, depending on the manufacturing process, the presence of copper may be unavoidable. Copper below 0.05% is considered a residual element. Copper as an alloying element may be added up to 0.20% to facilitate the removal of high silicon scale formed during the hot rolling stage of manufacturing the starting steel strip, and to improve corrosion resistance when the cold-rolled steel strip is used as is without surface treatment, or in the case of a zinc-coated strip to improve wettability by molten zinc. Copper can promote bainitic structures, cause solid solution hardening, and precipitate out of the ferrite matrix as ε-copper, thus contributing to precipitation hardening.Copper also reduces the amount of hydrogen that penetrates the steel and therefore improves its delayed fracture characteristic. However, excessive amounts of copper cause a heat deficit. Therefore, when copper is added, its content is kept below 0.20%. Chromium 0 - 0.35%; Nickel 0 - 0.50%; Molybdenum 0 - 0.30% Chromium, nickel, and molybdenum are not required elements, but they may be present as residual elements in the steel composition. The permitted level of Cr, Ni, or Mo as a residual element is 0.05% for each. As alloying elements, they improve the hardenability of steel and facilitate the formation of bainite ferrite, and at the same time, they have a similar effectiveness that is useful for stabilizing retained austenite. Therefore, Cr, Ni, and Mo are effective for microstructural control. QCRC Ln / ίZРZ / B / YILI, thus creating oxysulfides that do not have a detrimental effect on ductility, as in the case of the elongated manganese sulfides that would form if Ca or REM were not present. This effect saturates when the Ca content exceeds 0.0050% or the REM content exceeds 0.0100%. Preferably, the amount of Ca, if present, is controlled to a value below 0.0030%, more preferably below 0.0020%. Preferably, the amount of REM, if present, is controlled to a value below 0.0080%, more preferably below 0.0050%. The rest of the composition of steel comprises iron and unavoidable impurities. The chemical composition of the steels according to the invention matches the capacity of conventional continuous production lines. Microstructure The cold-rolled steel strip or band that has been heat-treated according to the invention has a complex microstructure, comprising 20-55% polygonal ferrite (PF), acicular ferrite (AF), and upper bainitic ferrite (HBF), wherein PF is at most 45%, as well as 20-65% lower bainitic ferrite (LBF), 5-20% retained austenite (RA), and fresh martensite (M) in an amount of 0-20%. In this invention, the microstructures are functionally grouped in such a way that they can be observed using optical and scanning electron microscopy. Polygonal ferrite (PF) refers to ferrite formed during intercritical annealing or slow cooling at temperatures above Bs. Acicular ferrite (AF) refers to ferrite formed during cooling at temperatures between Bs and Ms. High-temperature bainitic ferrite (HBF) is bainitic ferrite formed during quenching at a temperature between Bs and Bn. Low-temperature bainitic ferrite (LBF) is bainitic ferrite formed during quenching at a temperature between Bn and Ms. Bainitic ferrite structure Bainitic ferrite (BF) forms during heat treatment when the quenching temperature is in the range between Ms and Bs. BF is present as plates with an ultrafine grain size. Carbide precipitation between the ferritic lamellae, which is known to be detrimental to ductility, is suppressed by alloying with Si and / or Al. Bainitic ferrite does not contain carbides, unlike conventional bainite, which does. Bainitic ferrite also differs from ferrite (proeutectoid) in that it has a low dislocation density. The carbide-free BF microstructures provide high strength due to the intermediate hard bainitic ferrite structure with a high dislocation density and supersaturated carbon content.The bainitic ferrite structure also contributes to the desired high elongation, as it contains no carbide and fine residual austenite grains may be present at the lattice-like boundaries of the bainitic ferrite. In the invention, bainitic ferrite is divided into two types: bainitic ferrite formed in a high temperature range between Bs and Bn, referred to as high bainitic ferrite (HBF), and bainitic ferrite formed in a low temperature range between Bn and Ms, referred to as low bainitic ferrite (LBF). HBF has an average aspect ratio (defined as the length of the minor axis divided by the length of the major axis) greater than 0.35, and LBF has an average aspect ratio less than 0.35 when the section A cross-section of the steel strip subjected to 3% Nital etching is observed using scanning electron microscopy with EBSD analysis. The reason for making this distinction is that bainitic ferrite formed in the higher temperature range above Bn (HBF) is similar to bainitic ferrite (AF) in grain size and shape, and it is difficult to distinguish HBF from AF using SEM. Like AF, HBF has a larger grain size, lower dislocation density, and is softer than LBF, and it contributes to increasing the elongation of the steel. On the other hand, LBF has higher strength than HBF due to its thinner plate size, which contributes to the strength of the steel strip and also improves formability. A feature of the high-strength steel strip according to the present invention is that the bainitic ferrite can have a composite microstructure that includes HBF and LBF. Therefore, a high-strength cold-rolled steel strip with high elongation can be obtained. To achieve a good balance between high strength and elongation, between 20 and 65% LBF is required. If LBF is present in a lower amount, the steel strip has insufficient strength. However, if LBF is present in an excessively large amount, the effects of the other ferrites (PF, AF, and HBF) and the retained austenite with respect to elongation may be compromised. Therefore, the LBF content is in the range of 20 to 65%, preferably 30 to 60%. The formation of HBF in the present invention is due to heating the strip through the latent heat produced by the bainitic transformation or by applying a hot-dip galvanizing process. The formation of HBF, if present, in the present invention allows the bainitic transformation kinetics to be accelerated, if necessary, so that the bainitic transformation can be completed within the limited time frame in the overaging section of an existing production line. Depending on the amount of PF and AF resulting from the soaking and cooling stages, the amount of HBF is controlled, such that the total amount of PF, AF, and HBF is from 20 to 55%, preferably from 25 to 50%. As described above, HBF has a function similar to that of PF and AF.If sufficient quantities of PF and AF have formed in the soaking and cooling sections, and in order to obtain steel strips with higher strength, the amount of HBF should be minimized to 0%. If the amount of PF and AF is insufficient, the amount of HBF may be increased. However, the amount of HBF should be controlled so that the total amount of PF, AF, and HBF is between 20% and 55%, preferably between 25% and 50%. Polygonal ferrite and acicular ferrite Proeutectoid ferrite is softer than bainitic ferrite and functionally increases the elongation of the steel strip. Alternatively, a certain amount of proeutectoid ferrite is introduced, and its characteristics are controlled to enhance the bainitic transformation kinetics, improve the stability of the retained austenite, and further increase elongation. Two types of proeutectoid ferrite can be produced using the invention during annealing, depending on the formation temperature. The ferrite phase formed when austenitizing at an intercritical temperature or during cooling at a high temperature above the Bs temperature in the slow-cooling section is polygonal or blocky, called polygonal ferrite (PF). It has been shown that QCRC Ln / Lznz / E / YILI: This type of ferrite increases elongation but decreases yield strength and formability in the presence of bainitic or martensitic phases. Ferrite formed at lower temperatures in the rapid-cooling section, between Bs and Ms, has an almost acicular shape and a smaller grain size than PF and is known as acicular ferrite (AF). It is similar to HBF in morphology but has a relatively lower number of dislocations. The presence of AF can increase elongation without sacrificing strength and formability. Since PF, AF, and HBF have a function similar to the tensile properties in steel according to the invention, three types of these ferritic microstructures may be present, or one or two of them. In order to ensure high elongation, the volume fraction of PF, AF, and HBF is 20% or more, preferably 25% or more. In any case, the total amount of PF, AF, and HBF must be controlled to be less than 55%, preferably less than 50%. If the content of these ferritic microstructures is too high and exceeds 55%, the final microstructure will not contain enough lower bainitic ferrite, and therefore the strength will be reduced. When PF is present in steel, its grain size, morphology, and distribution must be controlled. Steel strip can achieve even greater elongation by having a smaller PF grain size and a more dispersed PF distribution. According to this invention, when observed with SEM or an optical microscope, the PF structure is embedded equiaxally between the BF structures and is uniformly dispersed as smaller grains, whereas the morphological structure of PF in a conventional TRIP steel strip is elongated along a rolling direction. This morphological structure is considered to allow for uniform stress distribution during processing and enables maximum utilization of the TRIP effect of retained austenite. To obtain this morphological structure, the amount of PF formed during soaking should be 45% or less, preferably 10 to 40%.This modality is particularly suitable for steel compositions containing relatively high amounts of Mn, Al, and Si, where the ferrite recrystallization kinetics are slower. The PF is a partially recrystallized microstructure, which has a higher hardness than fully recrystallized ferrite. The presence of PF with a partially recrystallized microstructure is beneficial for local ductility. Advantageously, the grain size of the PF in the present invention is 10 µm or less, preferably 8 µm or less, and more preferably 5 µm or less. In one embodiment of the invention, it is preferable that the amount of PF be 0%. In this case, the total amount of AF and HBF is controlled so that AF + HBF is in a range of 20 to 55%, preferably 25 to 50%. Residual austenite Residual austenite (also known as retained austenite) refers to a region exhibiting an FCC (face-centered cubic) lattice phase in the final microstructure. Retained austenite enhances ductility, partly through the TRIP effect, which manifests as increased uniform elongation. The volume fraction of residual austenite is 5% or more, preferably 7% or more, to exhibit the TRIP effect. Below 5%, the desired level of ductility and uniform elongation will not be achieved. The upper limit is primarily determined by QCRC Ln / įZРZ / B / YILI the composition and processing parameters in a production line. For a given composition, the carbon content in the retained austenite becomes too low if the amount of retained austenite is too high. Then, the retained austenite is insufficiently stable and the local ductility (flange formability by stretching) could be reduced to an unacceptable level. Therefore, the upper limit of the retained austenite volume fraction is 20%, preferably 15%. The carbon concentration in the residual austenite impacts the TRIP characteristics. Retained austenite is effective in improving elongation, particularly when its carbon concentration is 0.90% or higher. If the carbon content is too low, the retained austenite is not stable enough to produce the TRIP effect. Therefore, advantageously, the carbon content in the retained austenite is 0.90% or higher, preferably 0.95% or higher. Although the carbon concentration in the retained austenite is preferably as high as possible, practical processing conditions generally impose an upper limit of approximately 1.6%. The carbon content and stability of the retained austenite can be adjusted by controlling the amount of ferrites. Martensite Martensite is newly formed in the final quenching section after austempering or tempering. It suppresses elongation at the yield strength and increases the work-hardening coefficient (n-value), which is desirable for achieving stable, neck-free deformation and uniform deformation in the final pressed part. Even with 1% fresh martensite in the final steel strip or slab, a tensile response and therefore pressing behavior comparable to that of conventional dual-phase steels can be achieved. However, the presence of fresh martensite will affect formability due to crack formation along the martensite and LBF / HBF interfaces. Therefore, the amount of fresh martensite should be controlled to 20% or less, preferably 15% or less. Carbides Carbides can be present as fine precipitates, which form during tempering when the overaging temperature is too high or the overaging time is too long, or as pearlite formed during quenching when the cooling rate is too slow. According to the invention, the microstructure of the invented steel is free of pearlite and carbide. "Free of pearlite" means that the amount of layered microstructure, including cementite and ferrite, is less than 5%. "Free of carbide" means that the amount of carbide is below the detection limit of standard X-ray measurements. Characterization of microstructures The microstructural components classified in the steel according to the invention as described above can be quantitatively determined using techniques described below. The volume fraction of the constituents is measured by equating the volume fraction with the area fraction and measuring the area fraction from a polished surface using a commercially available image processing program or other suitable technique. QCRC Ln / ίZΠZ / Β / YΙΛΙ Fresh martensite (FM), fresh M, austenite (AA), and pearlite can be distinguished using optical microscopy (OM) and / or scanning electron microscopy (SEM). When a sample etched with 10% aqueous sodium metabisulfite (abbreviated SMB) is characterized under OM, pearlite appears as dark areas, FM as tinted gray areas, and fresh martensite as light brown areas. When a sample etched with 3% Nital solution is characterized by SEM, FM appears as grains with a smoother surface that does not include retained austenite, and pearlite appears as a layered microstructure that includes both cementite and ferrite. The remaining microstructure appears as gray areas, characterized by plates or laths as ferritic substructures, in which AA is dispersed within the grains as white or pale gray areas, and carbides cannot be identified. This microstructural group is known as bainitic ferrite-like microstructure.It may include a mixture of HBF, LBF, AF, and another microstructure called partitioned martensite (PM). PM occurs when the quenching temperature is below the Ms point of the steel. Some martensite forms during rapid cooling, and then carbon partitioning occurs between the martensite and the retained austenite during quenching. PM ​​has a morphology similar to LBF and can coexist with LBF. AF, HBF, LBF, and PM in the bainitic ferrite-like microstructure cannot be clearly distinguished using OM and SEM because their morphologies are similar. In this invention, the bainitic ferrite-like microstructure is further separated into two distinct groups by Electron Backscatter Diffraction (EBSD). The first group consists of PM and LBF, and the second group consists of AF and HBF. From the measured EBSD data, the retained austenite can first be distinguished from the other microstructures by creating a partition of Fe(y) from Fe(a). Fresh martensite (M) is then separated from the bainitic ferrite-like microstructure by dividing the Fe(a) into a partition with a medium-high image quality (IQ) and a partition with a medium-low IQ. The low-IQ partition is classified as martensite, and the high-IQ partition is classified as the bainitic ferrite-like microstructure. The method for distinguishing the two groups is described below with reference to Figures 1a and 1b.In bainitic ferrite (high IQ partition), regions are identified that have an orientation difference of at least 15° in the tilt angle between adjacent structures. A region is considered to have the same crystal orientation and is defined as a bainitic plate in the present invention. For the bainitic plates thus detected, the diameter of a circle having the same area as a bainitic plate is determined. The diameter of the equivalent circle of the bainitic plate is determined using the EBSD analysis photograph with a magnification factor of 3000. By fitting an ellipse to a bainitic plate, the aspect ratio (defined as the length of the minor axis divided by the length of the major axis) is also determined.Similarly, the diameters of the equivalent circles of all bainitic plates and the aspect ratios of the equivalent ellipses of all bainitic plates are measured in the measured area (approximately 100 by 100 pm), and the average values ​​are defined as the mean grain size of the bainitic plates and the mean aspect ratio of the bainitic plates in the present invention. The inventors have systematically studied the effect of annealing temperature on the microstructure of bainitic ferrite. The annealing temperature ranges from Ms - 200 to Bs. It has been found that the average size and average aspect ratio of the bainitic plates increase at QCRC Ln / ίZРZ / B / YILI as the tempering temperature increases. In particular, it is found that the aspect ratio of the bainitic plates changes sharply between samples austempered or tempered below 440°C, which is below Bn, and above 460°C, which is above Bn, of the steel composition used in the method according to the invention. Therefore, the critical average aspect ratio value of 0.35 is defined to divide the two groups of bainitic ferrite-like microstructure. The group consisting of LBF and PM has an aspect ratio of 0.35 or less, and the group consisting of HBF and AF has an aspect ratio greater than 0.35. In addition to the differences in morphology and size of the bainitic plates, the misorientation relationships between the intricate crystallographic plates of the HBF, AF, LBF, and PM groups are also different. The misorientation angle distribution in the steel according to the invention is shown in Figure 2. The peak at 60° is consistent with misorientations between neighboring grains, having the Kurdjumov-Sachs (KS / KS) ratio, which is caused by the axis-angle relationship 60°<111> and 60° <110> and corresponds to martensite. The peak at 53°–54° is due to grain misorientations obtained by phase transformations according to the Nishiyama-Wassermann and Kurdjumov-Sachs (NW / KS) relationship. According to the state of the art, (see A.-F. Gourgues, HM Flower, and TC Lindley, Materials Science and Technology, January 2000, vol. 16, p.(26-40), the acicular ferrite and upper bainite grow with Nishiyama-Wassermann relationships with the parental austenite phase, while the lower bainite and martensite consist of very intricate bundles that have Kurdjumov-Sachs relationships with the parental phase. In analogy with these results, the peak at 53-54°C is assumed to correspond to the formation of HBF and AF, and the peak at 60°C to the formation of LBF and PM. The peak at 53-54°C becomes more distinct and its peak height increases, but the peak height at 60°C decreases as the annealing temperature increases. In the present invention, the relative amounts of the HBF, AF group and the LBF, PM group can be determined by the ratio of the heights of the two peaks. Because some of the retained austenite is dispersed as a very thin film between the bainitic plates and cannot be detected by EBSD, the retained austenite fraction determined by EBSD is always less than the actual value. Therefore, an intensity measurement method based on X-ray diffraction (XRD) can be used as a conventional technique for measuring retained austenite content. The volume fraction of retained austenite is determined as a percentage of the steel strip thickness. The amount of cementite is also measured from this XRD analysis. A sample prepared from the steel strip is mechanically and chemically polished and then analyzed by measuring the integral intensity of each of the (200), (220) and (311) planes of fcc (face-centered cubic) iron and that of the (200), (211) and (220) planes of bcc (body-centered cubic) iron with an X-ray diffractometer using Co-Ka.The amount of retained austenite (RA) and the lattice parameter in the retained austenite were determined by Rietveld analysis. The C content in the retained austenite is calculated using the formula: C (% by weight) = (a[Á] - 3.572 - 0.0012 Mn% + 0.00157 Si% - 0.0056 Al%) / 0.033 QCRC Ln / ίZΖΠZ / Β / YΙΛΙ where a is the lattice parameter of the retained austenite in angstroms. Mechanical properties Cold-rolled steel strips or bands with the above microstructure and composition, and heat-treated according to the invention, have the following properties: The yield strength (YS) is at least 500 MPa; and / or the Tensile strength (TS) is at least 850 MPa; v / o the Total elongation (TE) is at least 14%. Preferably, the cold-rolled and heat-treated strip or band possesses all these properties. Stages of the method According to the method of the invention, a cold-rolled steel strip or band having the composition described above is heat-treated to obtain the microstructure and properties. The cold-rolled steel strip or band obtained by cold rolling is subjected to heat treatment as in a continuous annealing line. A typical process design is shown schematically in Figure 3. The cold-rolled steel strip or band is heated above the temperature (Ac3 - 60), for example, using a heating rate of at least 0.The steel strip is heated at a rate of 5°C / s, preferably in the temperature range of (Ac3 - 60) - (Ac3 + 20), typically to a predetermined austenitizing temperature T2, and held for a time period t2 within this temperature range (step a). It is then cooled, typically using two-step or stage cooling at controlled rates, to a temperature T4 in the range of Ms - Bn (step b). Following this, the steel strip is heat-treated in the range of T4 to Bs for tempering for a time t5 (step c), preferably at a temperature T5 in the range of T4 to Bn. Optionally, the steel strip is then heated to a temperature T6 in the range of Bn to Bs for a time period t6, which may be a hot-dip galvanizing treatment. Finally, the steel strip is cooled to room temperature (step d). The process parameters and functions in each step are described below. In the first step, the cold-rolled steel is soaked above (Ac3 - 60), within a temperature range of (Ac3 - 60) to (Ac3 + 20), for a soaking time t2 of 1-150 seconds to achieve a microstructure that is at least partially austenitic. Annealing at a higher temperature (Ac3 - 60) is necessary because the steel strip or band being heat-treated according to the invention needs to have the required amounts of low-temperature transformed phases, such as bainitic ferrite and retained austenite, which transform from high-temperature austenite, as well as a predetermined amount of ferrite. If T2 is greater than (Ac3 + 20), the austenite grains will grow, which influences the size and distribution of the retained austenite and also slows down the bainitic transformation kinetics later in the overaging process.An excessive amount of fresh martensite can form during the final cooling as a result of this incomplete bainitic transformation, leading to increased strength but low ductility and formability. Furthermore, a uniform austenite structure with larger grain sizes can suppress the formation of PF and AF in subsequent cooling. QCRC Ln / ίZРZ / B / YILI cooling section, so that an insufficient amount of ferrite is obtained within the current cooling program on the available production line, and may cause insufficient elongation of the steel strip. It has been observed that the uniformity of the austenite has a great effect on the formation of PF and AF in the cooling section. If T2 is less than (Ac3 60), PF may form in an excessive amount exceeding 45%, so the steel strip may obtain insufficient strength. On the other hand, the amount of austenite formed may not be sufficient for the formation of LBF and retained austenite. Therefore, the annealing temperature should be higher than (Ac3 - 60), but advantageously should not exceed (Ac3 + 20), preferably in the range of (Ac3 - 50) to (Ac3 + 10).If t2 exceeds 150 seconds, the austenite and ferrite grain sizes increase, leading to less elongation. If the annealing time t2 is less than 1 second, the reverse transformation to austenite may not proceed sufficiently and / or the carbides in the steel strip may not have dissolved enough. Therefore, the annealing time t2 is from 1 second to 150 seconds, such as from 10 seconds to 120 seconds, preferably from 1 to 100 seconds. In a subsequent cooling stage, the at least partially austenitic band is cooled to a T4 temperature in the Bn-Ms range. The purpose of this cooling is to regulate the amounts of ferrite and bainitic ferrite and to prevent pearlite formation. Typically, this means that the cooling rate during quenching is high enough to prevent pearlite formation. In one embodiment of the invention, the treated steel strip is cooled directly to temperature T4 at a cooling rate V4 of at least 15°C / s to prevent pearlite formation. If the cooling rate is too low, excessive ferrite or even pearlite may form. Preferably, V4 is greater than 20°C / s. The upper limit of the cooling rate is not particularly restricted unless a temperature variation occurs in the steel strip when cooling is stopped. The cooling rate is preferably 100°C / s or lower in standard installations to reduce temperature inhomogeneity in the steel strip.Significant variation in the microstructure can occur in both the longitudinal and width directions of the strip when the average cooling rate exceeds 80°C / s in most available facilities. Therefore, a suitable cooling rate V4 is in the range of 15 to 80°C / s, preferably 20 to 60°C / s. In other embodiments of the invention, this cooling can be achieved by means of two-stage cooling to regulate the amount of ferrite and homogenize the strip temperature. This is suitable for most continuous annealing lines or hot-dip galvanizing lines that include two connected cooling sections, such as those currently in use. The steel strip is first cooled to a temperature T3 in the range of 800–500°C (referred to as the slow cooling section), preferably in the range of 750–550°C, typically at a cooling rate of V3 of at least 1°C / s, such as 2–15°C / s, preferably 3–10°C / s. Thereafter, the steel strip is further cooled to temperature T4. QCRC Ln / įZРZ / B / YILI (referred to as the rapid cooling section), typically at a cooling rate V4 of at least 15°C / s, such as 15–80°C / s, preferably 20–60°C / s. Since the length of each section in a continuous annealing line is fixed, the cooling rates V3 and V4 for a given line speed can be controlled by adjusting the temperature T3. The higher the T3, the lower the V3 and the higher the V4. During this cooling, some ferrite may form in the slow cooling section, and some pearlite may form in the rapid cooling section. For a fixed line speed, the amount of ferrite formed in the slow cooling section depends primarily on T3, and the amount of pearlite depends primarily on V4. Therefore, T3 is selected within a suitable range to adjust the amount of ferrite and prevent pearlite formation.If T3 is too low, for example, below 550°C, excessive PF may form in the slow cooling section, and excessive AF may also form in the rapid cooling section. Pearlite may even form if the resulting V4 is below 15°C / s. If T3 is too high, for example, above 800°C, insufficient PF may form, and less AF may form if the resulting V4 is too high. Therefore, T3 should be in the range of 800 to 500°C, preferably in the range of 750 to 550°C / s. As mentioned previously, ferrite (PF) can be obtained in the soaking stage (a) and in the slow cooling section (b), while ferrite (AF) is obtained in the rapid cooling section (b) in a conventionally designed annealing or galvanizing line. The soaking temperature (T2) and the intermediate temperature (T3) between the slow and rapid cooling sections can be used to regulate the amount of ferrite. If a higher T2 is used, less ferrite is produced during soaking; therefore, a lower T3 can be selected to obtain more ferrite in the slow cooling section and more ferrite in the rapid cooling section. If a lower T2 is used, a sufficient amount of ferrite is produced during soaking; therefore, a higher T3 is selected to limit the amount of ferrite formed in the slow cooling section and the amount of ferrite formed in the rapid cooling section. The quenching stop temperature T4 is between Bn and Ms to initiate the bainitic transformation. If T4 is too high, excessive HBF may be obtained during the subsequent tempering, and the strength of the steel strip or band may be insufficient. If T4 is too low, and in the absence of active heating, the latent heat produced by the bainitic transformation is insufficient to heat the steel strip or band to a temperature high enough for rapid bainitic transformation kinetics. Preferably, T4 is in the range between Bn and Ms + 50°C. In the next heat treatment stage c), the cooled strip is heat treated in the range between temperatures T4 and Bs, and preferably between T4 and Bn, for a time t5 in the range of 30–300 seconds, typically by heating and heat treating at a temperature T5 within this range. During this time, the austenite transforms into lower bainitic ferrite (LBF). If T5 is too low, the bainitic transformation is too slow, insufficient during overaging, and excessive amounts of fresh martensite may form during cooling after overaging. This increases strength but fails to provide the required elongation. Furthermore, carbon partitioning may be insufficient to QCRC Ln / Lznz / E / YILI stabilizes the retained austenite. If T5 is too high, there is a risk of obtaining too much HBF in the overaging section, which cannot provide the required strength. The most preferred range for stage c) is Bn-50 to Bn to achieve rapid bainitic transformation kinetics and obtain LBF. If the heat treatment time t5 is less than 30s, the bainitic transformation is incomplete, and not enough LBF is formed. Again, carbon partitioning is insufficient. If t5 is greater than 300s, there is a risk of carbide formation, thus decreasing the carbon content in the retained austenite. The maximum time for t5 is limited, among other things, by the total time available at a given production line speed. Preferably, t5 is in the range of 40 to 120 seconds. Subsequently, the heat-treated strip is cooled to ambient temperature, according to the production line's capacity, during which some fresh martensite may form. The steel strip is then cooled below 300°C at a cooling rate (V7) of at least 1°C / s, preferably at least 5°C / s, after which it is further cooled to ambient temperature. This cooling to ambient temperature may be forced or uncontrolled natural cooling. In one practical embodiment, the heat-treated steel strip is cooled to a temperature (T7) in the range of (Ms - 50) - Mf at a cooling rate (V7) in the range of 5.0 - 10.0 sC / s. Further cooling from T7 to ambient temperature is advantageously carried out at a cooling rate (V8) of 5.0 - 20.0 sC / s, more preferably 6.0 - 15.0 sC / s. In one embodiment of the invention, heat treatment step c) is preferably carried out at least partially by means of latent heat from the bainitic transformation. During this heat treatment step c), the lower bainite transformation produces latent heat. Since latent heat can be incompletely dissipated in the production line, the temperature of the steel strip automatically increases due to the accumulated latent heat. Therefore, the temperature of the steel strip being heat-treated can gradually increase due to the latent heat from the bainitic transformation. That is, the temperature T5 increases during the course of the heat treatment, particularly at longer t5 times. The upper temperature of T5 is not particularly limited; if the steel strip reaches temperatures above Bn, high-temperature bainitic ferrite will form. In one embodiment, the soaking step is carried out within an intercritical annealing temperature range of (Ac3 - 50) - (Ac3 + 10), preferably for a soaking time t2 of 1 - 100 seconds to ensure that a partially austenitic cold-rolled strip with a fine grain size is obtained. The fraction of PF formed at the soaking temperature is advantageously less than 40%. Advantageously, the heating stage, prior to the soaking stage, is carried out in two sub-stages, comprising heating a cold-rolled strip or band to a temperature T1 in the range of 680–740°C, preferably in the range of 700–720°C, at a heating rate V1 of 10.0–30.0 sC / s, preferably 15.0–25.0 sC / s; and further heating the cold-rolled strip or band from temperature T1 to the soaking temperature range at a rate of QCRC Ln / Lznz / E / YILI heating V2 of 0.5–4.0 sC / s, preferably 1.0–3.0 sC / s. During the slow heating from T1 to the soaking temperature T2, recovery and recrystallization of ferrite occur, as well as the dissolution of carbides and ferrite during the austenite transformation. T1 and V2 affect the progress of these processes, which in turn affect the austenite grain size and the homogeneity of the alloying element distribution within the austenite phase. Advantageously, the soaking time t2 is controlled, depending on the heating rate V2, to ensure the dissolution of all carbides and prevent a coarse austenitic grain size. In one embodiment, the method according to the invention comprises an additional heat treatment step between the heat treatment step c) and the cooling step d), in which the steel strip or band resulting from step c) is subjected to further heat treatment in the Bs-Bn range, preferably (Bs-50)-Bn, typically at a fixed temperature T6. The additional treatment time t6 is advantageously from 5 to 30 seconds, preferably from 10 to 20 seconds. This additional heat treatment increases the bainitic ferrite by forming high-temperature bainitic ferrite from the remaining austenite to complete the bainitic transformation and thereby further reduces the amount of martensite formed in the subsequent cooling section, thus improving the strength and ductility properties. The carbon is also further broken down in the retained austenite, making it more stable.When this additional heat treatment is applied to a given overaging section and therefore for a given total time span thereon, the time t5 is further reduced to meet the available time span, e.g., the sum of t5 + t6 is in the range of 30-300 s. In a preferred embodiment, this additional heat treatment comprises an integrated hot-dip galvanizing treatment, in which the steel strip or band resulting from step c) is coated with a coating based on Zn or Zn alloy. The steel strip or band that has been heat-treated according to the invention can be coated, advantageously with a zinc-based coating or a zinc alloy coating. Advantageously, the zinc-based coating is either galvanized or annealed after galvanizing. The Zn-based coating may comprise a Zn alloy containing Al as an alloying element. A preferred zinc bath composition contains 0.10 to 0.35% Al, the remainder being zinc and unavoidable impurities. Another preferred Zn bath comprising Mg and Al as principal alloying elements has the composition: 0.5–3.8% Al, 0.5–3.0% Mg, optionally a maximum of 0.2% of one or more additional elements; the remainder being zinc and unavoidable impurities. Examples of the additional elements include Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr, and B1. The coating, such as a protective coating of Zn or Zn alloy, can be applied in a separate step. Preferably, a hot-dip galvanizing step is integrated into the method according to the invention as explained above. Optionally, a tempered rolling treatment can be carried out on the annealed and zinc-coated strip according to the invention in order to refine the tensile properties and modify the appearance and surface roughness depending on the specific requirements resulting from the intended use. QCRC Ln / ίZΠZ / Β / YΙΛΙ Cold-rolled steel strip is typically manufactured according to the following general process. A steel composition is prepared as described above and cast into a slab. The cast slab is processed by hot rolling after being reheated to a temperature in the range of 1100 to 1300°C. The hot rolling of the slab is usually carried out on 5 to 7 stands to the final dimensions suitable for further cold rolling. The final rolling is typically performed under fully austenitic conditions above 800°C, advantageously 850°C or higher. The strip thus obtained from the hot rolling stages can be coiled, for example, at a coiling temperature of typically 700°C or lower. The hot-rolled strip is then pickled and cold-rolled to obtain a cold-rolled steel strip of the appropriate gauges.Preferably, the reduction in thickness during cold rolling is typically in the range of 30 to 80%. To reduce the rolling strength during cold rolling, the wound strip or the cold-rolled half-strip can be subjected to batch hot annealing. The batch annealing temperature should be in the range of 500–700°C. Thin slab casting, strip casting, or similar methods can also be used. In this case, it is acceptable for the manufacturing method to omit at least part of the hot rolling process. The invention also relates to a heat-treated, cold-rolled steel strip having a composition and microstructure as described above. The invention also resides in an article, such as a structural, engineering, or automotive component, produced from the cold-rolled and heat-treated strip according to the invention. Examples Steels with the compositions shown in Table 1 were cast into 25 kg ingots measuring 200 mm x 110 mm x 110 mm using vacuum induction. The following process program was used to manufacture 1 mm thick cold-rolled strips: • Reheating the ingots to 1225°C for 2 hours; • Rough or rough rolling of ingots from 140 mm to 35 mm; • Reheating the raw rolled ingots to 1200°C for 30 min; • Hot rolled from 35 mm to 4 mm in 6 passes; • Cooling of the exhaustion table: cool from the final rolling temperature (FRT) (approximately 850 to 900°C) to 600°C at a rate of 40°C / s; • Oven cooling: The strips are transferred to an oven preheated to 600°C and then cooled to room temperature to simulate the cooling process; • Pickling: Next, the hot-rolled strips were pickled in HCI at 85°C to remove the oxide layers. • Cold rolled: The hot rolled strips were cold rolled into 1 mm strips; • Heat treatment according to the invention: Cold-rolled sheets or plates of suitable size were used to simulate the annealing process using a continuous annealing simulator (CASIM). QCRC Ln / Lznz / E / YILI Samples were machined from the treated sheets for microstructure observations, tensile tests, and hole expansion tests. Dilatometry was performed on cold-rolled samples measuring 10 mm x 5 mm x 1 mm (length along the rolling direction). The dilatation tests were conducted using a Bahr DIL 805 dilatometer. All measurements were performed in accordance with SEP 1680. The critical phase transformation points Ac3, Ms, and Mf were determined from the quenched dilatometry curves. Bs and Bn were predicted using the available software JmatPro 10. The phase fractions during annealing for different process parameters were determined from dilatation curves simulating the annealing cycles. The microstructure was determined by optical microscopy (OM) and scanning electron microscopy (SEM) using commercially available image processing software. Microstructures were observed at % thickness in the cross-section of the rolling and normal directions of a steel strip. The scanning electron microscope (SEM) used for EBSD measurements was a Zeiss Ultra 55 equipped with a Field Emission Gun (FEG-SEM) and an EDAX PEGASUS XM 4 HIKARI EBSD system. EBSD scans were captured using TexSEM Laboratories (TSL) OIM (Orientation Imaging Microscopy) data collection software. The EBSD scans were analyzed using TSL OIM Analysis software. The EBSD scan area was 100 x 100 µm in all cases, with a step size of 0.1 µm and a scan rate of approximately 80 frames per second. Retained austenite was determined by XRD according to DIN EN 13925 on a D8 Discover GADDS (Bruker AXS) with Co-Kα radiation. Quantitative determination of phase proportions was performed by Rietveld analysis. Tensile tests – JIS5 test pieces (gauge length = 50 mm; width = 25 mm) were machined from the annealed strips so that the tensile direction was parallel to the rolling direction. Tensile tests at room temperature were performed on a Schenk TREBEL testing machine according to NEN-EN10002-1:2001 to determine the tensile properties (yield strength YS (MPa), tensile strength UTS (MPa), total elongation TE (%)). For each condition, three tensile tests were performed, and the average values ​​of the mechanical properties are reported. The process parameters are presented in Table 2, using the indications in Figure 3. In CASIM, the latent heat produced during the bainitic transformation is compensated for by the active cooling system. The temperature variation due to latent heat in the overaging section is simulated by temperatures T4, T5, and T6. Where T4 equals T5, t4 is 0. Where T4 is less than T5, t4 is 1 s, and subsequent heating to T5 occurs at a heating rate of 5–20°C / s. The heating rate from T5 to T6 is 5–10°C / s. The resulting microstructures and tensile properties are given in Table 3. All steel compositions are inventive except for A79 and meet the microstructure and tensile property requirements under various specified processing parameters. Steel A79 failed to achieve the required elongation, even though its tensile strength was sufficiently high. QCRC Ln / Lznz / E / YILI (examples 26 and 27). The reason is that A79 contains 0.5% Cr, which significantly slows the bainitic transformation kinetics, so the bainitic transformation and the partitioning of C between BF and the retained austenite could not proceed sufficiently. Consequently, the amount of martensite increases, but the amount of retained austenite and the C content in the retained austenite decrease, leading to a reduced TRIP effect. QCRC Ln / ίZΠZ / Β / YΙΛΙ Table 1. Compositions (in % by weight) and critical phase transformation points (in C) of steels QCRC Ln / Lznz / E / YILI Table 2. Process parameters €0 > =C / s co co co en n CO en co 03 en en 9.3 en co 11.3 11.3 11.3 cados en la Figura 3 ¿1 ¿Λ 91 91 SI SI 'C s CC s °C / s C 1 430 53 455 19 6 3 300 | 455 53 455 19 6.3 300 | 440 60.3 460 21.5 5.8 300 | 440 53 460 19 6 5 300 | 440 53 460 19 6.5 300 I 440 70.5 460 25 1 5 300 I 1440 53 460 19 6.5 300 II 450 53 460 19 6.5 300 440 53 460 19 6.5 300 | 440 70.5 460 25.1 5 300 440 60.3 460 21.5 5.8 300 440 53 460 19 6.5 300 I 440 53 460 19 6 5 300 | I 440 70.5 460 25.1 5 300 I 440 53 460 19 6.5 300 440 53 460 19 6.5 300 440 53 460 19 6.5 300 440 ' 53 460 19 6.5 300 I 440 53 460 19 6.5 300 I 440 53 460 19 6.5 300 σ c 2 (Λ ooooooooooo OO ooooooo -- o roceso O o co 04 o 440 440 I 440 I 440 4501 440 440 440 440 440 440 440 440 440 440 o ΟίΦ s del pi > ω 5 ω OO co O 38.5 CO o CO OOO co OOO o LO (XI O Φ E CO 1- o O § O oo I'- 099 o § 700 660 700 002 o I'- o 700 o I'- 660 oo 600 099 <0 Q. > <Λ p en m en LO co CO <O CO SC en en (XI 4.4 3.7 (XJ CO 6.4 <N ω co CO cO <0 98¿ cO (D un CO o co co co 2 cO co LO CO 73.5 un co cO (O CO CO CO CO 2 cO cT>CO CO CO CO cO CO (J 0 850 098 860 860 860 840 860 850 850 098 850 860 098 860 098 O CO 078 0S8 OO CO cO CO CO 850 098 | TI V2 po 0 I 17 20 | 720 1.28 | 720 1 24 I 720 11 I 720 1.34 νε i ozz | 720 1.55 | 720 1.34 720 1.65 720 1.34 I 720 1.3 [ 720 1.4 Alloy Code 9LV ] ω [ A50 | A50 A50 LSV I L9V ] L9V 15 AV 15 | VSV | ϊη A72 A73 I A73 | A74 Example (N un (O co CD o (XI CO cO CO r- co in OJ. QCRC Ln / ίΖΠΖ / Β / ΥΙΛΙ QCRC Ln / ίΖΠΖ / Β / ΥΙΛΙ Table 2. (continued) Table 3. Microstructures and Tensile Properties QCRC Ln / Lznz / Ε / ΥΙΛΙ QCRC Ln / ίΖΠΖ / Β / ΥΙΛΙ Table 3 (continued)

Claims

1. A method for heat treating a cold-rolled steel strip, characterized in that the method comprises the steps of: a) soaking a cold-rolled steel strip above (Ac3 - 60) for a soaking time t2 of 1-150 seconds, thereby obtaining a cold-rolled steel strip having a microstructure that is at least partially austenitic; b) cooling the soaked steel strip resulting from step a) to a temperature T4 in the range of Bn-Ms; c) heat-treating the cooled strip obtained in step b) in a temperature range of Bs-T4 for a time period t5 of 30-300 seconds; d) cooling the heat-treated steel strip to ambient temperature; such that the steel strip has a microstructure (in % vol.) comprising: QCRC Ln / ίZΖΠZ / Β / YΙΛΙ polygonal ferrite (PF) + acicular ferrite (AF) + upper bainitic ferrite (HBF): 20 - 55; wherein polygonal ferrite (PF): 0 - 45; lower bainitic ferrite (LBF): 20 - 65; retained austenite (RA): 5 - 20; martensite (M): 0-20; wherein the steel strip or band has a composition (in % by mass) comprising: C: 0.15-0.35; Mn: 1.50-3.00; Si: 0.50-2.00; Al: 0.01-1.50; P: less than 0.050; S: less than 0.020; N: less than 0.0080; where the sum (Si + Al) is 0.60; and optionally one or more elements selected from 0 < Cr < 0.35; 0 < Cu < 0.20; 0 < Ni < 0.50; 0 < : Mo < 0.30; 0 < : Nb < 0.10; 0 < : V < 0.10; 0 < Ti < 0.10; 0 < B < 0.0030; 0 < Ca < 0.0050; 0 < REM < 0.0100, where REM is one or more rare earth metals; and the remainder is iron and unavoidable impurities.

2. The method according to claim 1, further characterized in that step a) comprises soaking a cold-rolled steel strip or band within a temperature range of (Ac3 - 60) - (Ac3 + 20), preferably within a temperature range of (Ac3 - 50) - (Ac3 +10), preferably for a soaking time t2 of 1 - 100 s.

3. The method according to claim 1 or claim 2, further characterized in that step b) comprises cooling the soaked steel strip or band of step a) to temperature T4 at a cooling rate sufficient to prevent pearlite formation.

4. The method according to any one of the preceding claims, further characterized in that step b) comprises a sub-step of cooling the soaked steel strip or band from step a) to a temperature T3 in the range of 800 - 500°C, preferably in the range of 750 - 550°C, at a cooling rate V3 of at least 1 sC / s, preferably at a cooling rate V3 of 2.0 - 15.0 sC / s, more preferably at a cooling rate V3 of 3.0 - 10.03 sC / s.

5. The method according to any one of the preceding claims, further characterized in that step b) comprises a sub-step of cooling the soaked steel strip or band from a temperature T3 in the range of 800 - 500°C, preferably in the range of 750 - 550°C, to T4 at a cooling rate V4 of at least 15°C / s, preferably at a cooling rate V4 of 20.0 - 60.0°C / s.

6. The method according to any one of the preceding claims, further characterized in that prior to step a) it additionally comprises heating a cold-rolled strip or band to a temperature above (Ac3 - 60), at a heating rate of at least 0.5°C / s, preferably comprising heating the cold-rolled strip or band to a temperature T1 in the range of 680 - 740°C, preferably in the range of 700 - 720°C, at a heating rate V1 of 10.0 - 30.0°C / s, preferably at a heating rate V1 of 15.0 - 25.0°C / s; and further heating the cold-rolled strip from temperature T1 to a temperature above (Ac3 60), preferably to the temperature range of (Ac3 - 60) - (Ac3 + 20), more preferably (Ac3 - 50) - (Ac3 + 10), at a heating rate V2 of 0.5 to 4.0°C / s, preferably at a heating rate V2 of 1.0 to 3.0°C / s.

7. The method according to any one of the preceding claims, further characterized in that step c) is carried out at least partially by latent heat produced by the transformation of bainite.

8. The method according to any one of the preceding claims, further characterized in that the heat treatment step c) is performed in the range of Bn - (Ms + 50), preferably for a time period t5 of 40-120 seconds.

9. The method according to any one of the preceding claims, further characterized in that it comprises an additional heat treatment step between heat treatment steps c) and d) of the steel strip or band resulting from step c) in the range of Bs - Bn, preferably (Bs-50) - Bn, preferably for a time period t6 of 5 to 30 seconds, more preferably for a time period t6 of 10 to 20 seconds. QCRC Ln / ίZΖΠZ / Β / YΙΛΙ 10. The method according to claim 9, further characterized in that the additional heat treatment step comprises a hot-dip galvanizing treatment.

11. The method according to any one of the preceding claims 1 to 9, further characterized in that after the heat treatment it further comprises a coating step of coating the heat-treated steel strip or band with a protective coating, preferably a Zn or Zn alloy coating.

12. The method according to any one of the preceding claims, further characterized in that the microstructure comprises in % by vol: polygonal ferrite (PF) + acicular ferrite (AF) + upper bainitic ferrite (HBF): 25 - 50; wherein polygonal ferrite (PF): 10-40; lower bainitic ferrite (LBF): 30 - 60; retained austenite (RA): 7-15; martensite (M): 0-15; and / or wherein the C content in the retained austenite (RA) is 0.90% by weight or more, preferably 0.95% by weight or more.

13. The method in accordance with any one of the preceding claims, further characterized in that the resulting steel strip or band has at least one, preferably all, of the following properties: Yield strength (YS) > 500 MPa; and / or Tensile strength (TS) > 850 MPa; and / or Total elongation (TE) > 14%.

14. A cold-rolled, heat-treated steel strip or band having a composition (in % by mass) characterized in that it comprises: C: 0.15-0.35; Mn: 1.50-3.00; Si: 0.50-2.00; Al: 0.01-1.50; P: less than 0.050; S: less than 0.020; N: less than 0.0080; wherein the sum (Si + Al) is 0.60; and optionally one or more elements selected from 0 < Cr < 0.35; 0 < Cu < 0.20; 0 < ; Ni < 0.50; 0 < ; Mo < 0.30; 0 < ; Nb < 0.10; 0 < ; V < 0.10; 0 < Ti < 0.10; 0 < B < 0.0030; 0 < Ca < 0.0050; 0 < REM < 0.0100, wherein REM is one or more rare earth metals; and the remainder is iron and unavoidable impurities; and a microstructure (in % by volume) comprising polygonal ferrite (PF) + acicular ferrite (AF) + upper bainitic ferrite (HBF): 20 - 55; wherein polygonal ferrite (PF): 0 - 45; lower bainitic ferrite (LBF): 20 - 65; retained austenite (RA): 5 - 20; martensite (M): 0 - 20. QCRC Ln / Lznz / E / YILI 15. The cold-rolled steel strip or band, heat-treated, according to claim 14, further characterized in that it has at least one, preferably all, of the following properties: Yield strength (YS) > 500 MPa; and / or Tensile strength (TS) > 850 MPa; and / or Total elongation (TE) > 14%.