Method for heat treating high strength cold rolled steel strip

Abstract

1. A method of heat treating cold rolled steel strip, comprising the steps of: soaking a cold rolled steel strip having a specific composition at (Ac3-60) for a specific time, thereby obtaining a cold rolled strip having an at least partially austenitized microstructure; cooling the resulting soaked steel strip to a temperature below Ms; heating the cooled steel strip to a temperature between Bs and Ms; Cooling the heated steel strip to ambient temperature A method comprising:

Description

[Technical field]

[0001] The present invention relates to a method for heat treating high strength cold rolled steel strip. [Background technology]

[0002] In the art, various types of cold rolled steel and manufacturing processes have been proposed to meet the requirements of automotive applications. For example, ultra-low carbon steel is used for automotive steel strips from the viewpoint of formability. This steel type exhibits a tensile strength of 280-380 MPa.

[0003] HSLA (high strength low alloy) steels contain micro-alloying elements that are hardened by a combination of precipitation and grain refinement.

[0004] Advanced high strength steels (AHSS), such as dual phase (DP) steels and transformation induced plasticity (TRIP) steels, are representative of the high ductility and strength steels currently used in the automotive manufacturing industry. DP steels have martensite in a ferrite matrix, which allows them to obtain tensile strengths of over 450 MPa and good cold formability.

[0005] In order to simultaneously achieve a high yield strength / tensile strength ratio and an even higher tensile strength, i.e., greater than 800 MPa, steels with complex (CP) microstructures containing ferrite, bainite, martensite and / or retained austenite have been developed. However, due to the difference in deformation ability between the ferrite, bainite or martensite structure and the retained austenite structure, these steels generally have poor stretch flange formability. Therefore, their use is limited to automotive parts that do not require high formability.

[0006] TRIP-type tempered martensitic steels (Q&P steels by quenching and partitioning) consisting of tempered martensite and retained austenite as matrix phases, and TRIP-type bainitic ferritic steels (TBF steels by austempering) consisting of bainitic ferrite and retained austenite as matrix phases have advantages (e.g., the ability to provide high strength due to hard tempered martensite and / or bainitic ferrite structure, and the ability to show good elongation because the matrix is ​​carbide-free), and fine retained austenite grains can be easily formed at the boundaries of lath-shaped bainitic ferrite in the bainitic ferrite structure. Therefore, carbide-free bainitic ferrite steels or tempered martensitic steels are expected to achieve good stretch flangeability due to the uniform fine lath structure. Due to the non-uniformity of hardness due to the presence of only a small amount of martensite in these microstructures, these steel types can achieve good deep drawability.

[0007] However, due to limitations of current continuous production lines, the expected beneficial combination of strength and ductility properties could not be obtained with currently available steel recipes. These limitations include, among others, that the reheating furnaces of current equipment of continuous annealing (CA) and continuous galvanizing (CG) lines are often only suitable for subjecting the steel strip to an intercritical or recrystallization heat treatment. For example, the maximum annealing temperature of some current annealing lines is limited to 890°C. Furthermore, the cooling rate of current CA / CG lines is limited within certain ranges. The overaging time available for many CA / CG lines is also limited, for example, this time period is less than about 160 seconds, which imposes a serious time limit to complete any desired transformation during overaging.

[0008] For example, WO2013 / 144373A1 discloses a cold-rolled TRIP steel with a matrix of polygonal ferrite, a specific composition including chromium and a specific microstructure, and a tensile strength of at least 780 MPa, which is said to enable production on conventional industrial annealing lines with overaging / austempered sections, i.e., for a relatively high overaging / austempering temperature, the austempering time can be less than 200 seconds.

[0009] EP2831296B1 and EP2831299 disclose TBF steels with tensile strengths of at least 980 MPa that can be produced on conventional production lines. However, the preferred overaging / austempering time is 280-320 seconds, which is too long to allow production on many conventional production lines. In other words, the bainite transformation rate is too slow to complete the bainite transformation in the limited time in the overaging section to obtain the desired microstructure on conventional production lines. Summary of the Invention [Problem to be solved by the invention]

[0010] It is an object of the present invention to provide a cold rolled steel strip, in particular for use in automotive applications, having a combination of desired high tensile strength and high total elongation (TE) properties at reasonable yield strength (YS), e.g. TS≧850 MPa, TE≧14% and / or YS≧500 MPa.

[0011] It is a further object of the present invention to provide a method for heat treating cold rolled steel strip to obtain the desired combination of properties as mentioned above, in particular a heat treatment which can be carried out using existing production lines or suitable alternatives.

[0012] It is another object of the present invention to provide high silicon cold rolled steel strip having a desirable combination of properties which can be produced on conventional industrial manufacturing lines.

[0013] Yet another object of the present invention is to provide a steel composition and heat treatment for high strength cold rolled steel strip that allows the bainite transformation to be completed in conventional production lines to obtain the desired microstructure. [Means for solving the problem]

[0014] With these considerations in mind, the present invention provides a method for heat treating cold rolled steel strip, comprising the steps of: The method comprises the steps of: a) soaking the cold rolled steel strip at (Ac3-60) or higher for a soaking time t2 of 1 to 150 seconds, thereby obtaining a cold rolled steel strip having an at least partially austenitized microstructure; b) cooling the soaked steel strip obtained from step a) to a temperature T4 of Ms~(Ms-200); c) heating the cooled steel strip obtained from step b) to a temperature between Bs and Ms; d) heating the heated steel strip at a temperature between Bs and Ms for a time t5 between 30 and 120 seconds; e) cooling the heated steel strip to ambient temperature. so that the steel strip comprises, in volume percent: Polygonal ferrite (PF) + acicular ferrite (AF) + higher bainitic ferrite (HBF): 20~55 Polygonal ferrite (PF): 0~50 Low bainitic ferrite (LBF) + partitioned martensite (PM): 25~70 Retained austenite (RA): 5~20 Martensite (M): 0~15 having a microstructure comprising The steel strip has the following composition, in mass %, C: 0.15~0.35 Mn: 1.50-4.00 Si: 0.50 to 2.00 Al: 0.01 to 1.50 P: less than 0.050 S: Less than 0.020 N: Less than 0.0080 In some cases, 0 <Cr≦1.00 0 <Cu≦0.20 0 <Ni≦0.50 0 <Mo≦0.50 0 <Nb≦0.10 0 <V≦0.10 0 <Ti≦0.10 0 <B≦0.0030 0 <Ca≦0.0050 0 <REM≦0.0100 One or more elements selected from Iron and unavoidable impurities: balance having The sum of (Si+Al) is ≧0.60, The method as described above, wherein REM is one or more rare earth metals.

[0015] The process of the present invention allows for the production of cold rolled steel strip having specific compositions and microstructures and combinations of properties desirable for automotive components requiring high strength, formability and weldability.

[0016] The present invention solves the problem of slow bainite transformation rate by introducing an appropriate amount of proeutectoid ferrite and controlling its morphology, by obtaining fine grains of austenite by controlling the maximum annealing temperature and time, and by using a modified quenching and distribution process in the production line.

[0017] The method according to the invention can be carried out using existing continuous annealing and galvanizing lines, within the limitations regarding maximum annealing section temperatures, cooling rate ranges and overaging time windows at the production rates typical of these production lines.

[0018] The cold rolled steel strip may be Zn-coated, for example by hot dip galvanizing or electrogalvanizing. The hot dip galvanizing process may be easily integrated into the heat treatment according to the invention.

[0019] The terms used to describe the critical transformation temperatures of steels, as are familiar to those skilled in the art, are set forth below. Ae3: The temperature at which the transformation from ferrite to austenite or from austenite to ferrite occurs under equilibrium conditions. Ac3: The temperature at which the transformation from ferrite to austenite is completed during heating. Ac3 is usually higher than Ae3, but tends toward Ae3 as the heating rate tends to zero. In the present invention, Ac3 is measured at a heating rate of 3° C. / s. Ar3: The temperature at which austenite begins to transform into ferrite during cooling. Bs: The temperature at which the transformation from austenite to bainite begins during cooling. Bn: The nose temperature of the bainite transformation in the time-temperature transformation (TTT) curve of a steel. At this temperature, the transformation from austenite to bainite shows the fastest reaction rate. Ms: The temperature at which the transformation from austenite to martensite begins during cooling. Mf: the temperature at which the transformation from austenite to martensite is completed during cooling. A practical problem with Mf is that during cooling, the proportion of martensite approaches the maximum achievable amount only asymptotically. This means that it takes a very long time for the final martensite to form. Therefore, for practical reasons and in the context of the present invention, Mf is taken as the temperature at which 90% of the maximum achievable amount of martensite has been formed.

[0020] These critical phase transformation temperatures can be determined by dilatometer experiments, or the Ac3, Bs, Bn and Ms points of a steel according to the invention can be pre-calculated based on its composition using available software, e.g. JmatPro or the following empirical formula:

number

[0021] In this specification, all temperatures are given in degrees Celsius, all compositions are given in weight percent (wt%), and all microstructures are given in volume percent (vol%) unless otherwise specified. [Brief description of the drawings]

[0022] [Figure 1] FIG. 1 shows EBSD maps illustrating the microstructural characteristics of low temperature bainitic ferrite and / or partitioned martensite (FIG. 1a) and high temperature bainitic ferrite (FIG. 1b), respectively. [Diagram 2] FIG. 2 is a histogram of the misorientation angles of low-temperature bainitic ferrite and high-temperature bainitic ferrite. [Diagram 3] FIG. 3 illustrates a generally applicable time vs. temperature profile for an embodiment of the method according to the invention. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0023] After the compositional description, the steps and microstructure of the method according to the invention are presented.

[0024] composition Carbon: 0.15-0.35% A sufficient amount of carbon is necessary for strength and stabilization of the retained austenite, the latter providing the TRIP effect. With that in mind, to ensure the required strength and elongation, the amount of carbon is at least 0.15%, preferably at least 0.17%. Increasing the carbon content increases the strength of the steel, the amount of retained austenite, and the carbon content in the retained austenite. However, if the carbon content is higher than 0.25%, the weldability of the steel is significantly reduced. For applications requiring welding, the carbon content is preferably 0.15-0.25%, more preferably 0.17-0.23%.

[0025] Silicon: 0.50~2.00% Silicon is an essential element in the steel composition according to the invention to obtain the described microstructure. Its main function is to prevent carbon from precipitating in the form of iron carbides (most commonly cementite) and to inhibit the decomposition of the retained austenite. Silicon contributes to the strength properties and to the proper transformation behavior. Furthermore, silicon contributes to improving the ductility, work hardening and stretch flangeability by inhibiting the grain growth of austenite during annealing. A minimum of 0.50% Si is required to sufficiently inhibit the formation of carbides. However, a high silicon content leads to the formation of silicon oxide on the strip surface, which reduces the surface quality, coatability and workability. Furthermore, the Ac3 temperature of steels with this composition increases with increasing silicon content. This may affect the feasibility of producing the steel strip using existing production lines in terms of the maximum top temperature achievable in the annealing section. With these considerations, the silicon content is not more than 2.00%. From the viewpoints of suppressing carbide formation, promoting austenite stabilization and improving wettability, the Si content is preferably 0.80 to 1.80%, and more preferably 1.00 to 1.60%.

[0026] Aluminum: 0.01 to 1.50% The main function of aluminum is to deoxidize the molten steel before casting. At least 0.01% Al is required for deoxidation of the molten steel. In addition, aluminum has a function similar to silicon, preventing the formation of carbides and stabilizing the retained austenite. Al is considered to be less effective than Si. It has no significant effect on strengthening. Small amounts of Al can be used to partially replace Si and adjust the transformation temperature and critical cooling rate to obtain acicular ferrite (AF) and increase the bainite transformation rate. Al is added for these purposes. Therefore, the Al content is preferably at least 0.03%. High levels of Al may raise the ferrite to austenite transformation point to a level that is incompatible with current equipment, and as a result, it is difficult to obtain a microstructure in which the main phase is a low-temperature transformation product. As the Al content increases, the risk of cracking during casting increases. Taking this into account, the upper limit is 1.50%, preferably 1.00%, and more preferably 0.70%.

[0027] With regard to the relationship between the proportion of Si and the proportion of Al, the composition satisfies the condition Si+Al≧0.60, preferably Si+Al≧1.00. Advantageously, the content of Al is less than 0.5 times the content of Si.

[0028] Manganese: 1.50~4.00% Manganese is necessary to obtain the microstructure of the steel strip according to the invention in terms of hardenability and stabilization of the retained austenite. Mn also influences the formation of proeutectoid ferrite at high temperatures and the bainitic ferrite transformation rate. A certain amount of Si and / or Al is necessary to suppress the formation of carbides in the bainitic ferrite. The Ac3 temperature increases with increasing Si and Al content. Mn is also adjusted to balance the increased phase transformation temperature Ac3 as a result of the presence of Si and Al. If the Mn content is less than 1.50%, it is difficult to obtain the described microstructure. Therefore, Mn needs to be added at 1.50% or more. However, if Mn is present in an excessive amount, macrosegregation is likely to occur, resulting in unfavorable band formation in the steel. Moreover, an excessive amount of Mn slows down the bainitic transformation rate, which results in an excessive amount of fresh martensite, and as a result, the stretch flange formability is also reduced. Therefore, the Mn content is not more than 4.00%, preferably not more than 3.50%, more preferably 1.80≦Mn≦3.00%.

[0029] Phosphorus: <0.050% Phosphorus is an impurity in steel. It segregates at grain boundaries and reduces workability. Its content is less than 0.050%, preferably less than 0.020%.

[0030] Sulfur: <0.020% Sulfur is also an impurity in steel. S forms sulfide-based inclusions, such as MnS, which cause cracks and reduce the stretch flangeability of the steel. The S content is preferably as low as possible, for example less than 0.020%, preferably less than 0.010%, and more preferably less than 0.005%.

[0031] Nitrogen: <0.0080% Nitrogen is another inevitable impurity of steel. It precipitates as nitrides with microalloying elements and is present in solid solution, contributing to strengthening. Excessive nitrides reduce elongation, stretch flangeability and bendability. Therefore, advantageously, the nitrogen content is less than 0.0080%, preferably less than 0.0050%, more preferably less than 0.0040%.

[0032] The steel composition may include one or more optional elements as follows:

[0033] Copper: 0~0.20% Copper is not required in steel embodiments having this composition, but may be present. In some embodiments, depending on the manufacturing process, the presence of Cu may be unavoidable. Less than 0.05% copper is considered a residual element. Copper as an alloying element may be added up to 0.20%. This facilitates the removal of scale formed during the hot rolling stage of producing the starting steel strip, and also improves corrosion resistance if the cold rolled steel strip is used as is without surface treatment, or improves wettability by molten zinc in the case of Zn-coated strip. Cu promotes the bainite structure, causes solid solution hardening, and may contribute to precipitation hardening since it precipitates as ε-copper from the ferritic matrix. Cu also reduces the amount of hydrogen that penetrates the steel, thereby improving delayed fracture properties. However, when added in excessive amounts, Cu causes high temperature brittleness. Therefore, when Cu is added, the Cu content is 0.20% or less.

[0034] Chromium: 0~1.00%; Nickel: 0~0.50%; Molybdenum: 0~0.50% Chromium, nickel and molybdenum are not essential elements, but may be present as residual elements in the steel composition. The acceptable level of Cr, Ni or Mo as residual elements is 0.05% each. As alloying elements, they have similar effects that improve the hardenability of the steel and promote the formation of bainitic ferrite while at the same time helping to stabilize the retained austenite. Thus, Cr, Ni and Mo are effective in microstructure control. To fully obtain this effect, it is preferred that the Cr, Ni or Mo content in the steel is at least 0.05%. However, if each is added in excess, the effect saturates and the bainite transformation rate becomes too slow to obtain the required microstructure in production lines where overaging time is limited. Therefore, the amount of Cr is limited to a maximum of 1.00%. Ni is simply used to reduce the hot brittleness characteristic when relatively large amounts of Cu are added. This effect of Ni becomes more pronounced when the Ni content is >[Cu(%) / 3]. The amount of Ni and Mo, if present, is limited to a maximum of 0.50% each.

[0035] Niobium: 0~0.100%; Vanadium: 0~0.100%; Titanium: 0~0.100% The acceptable levels of niobium, vanadium and titanium as residual elements are 0.005% each. One or more of niobium, vanadium and titanium can be added to refine the microstructure of the hot rolled intermediate and final products. These elements have a precipitation strengthening effect and may change the morphology of bainitic ferrite. They also contribute beneficially to the optimization of applications that depend on properties, such as ductility and bendability of the drawn edge. To obtain these effects, the lower limit of any of these elements, if present, must be adjusted to 0.005% or more. When the content of Nb, Ti and V exceeds 0.10%, respectively, the effect saturates. Therefore, when these elements are added, their contents are adjusted to 0.005% to 0.100%. Preferably, the upper limit is 0.050% or less for Nb and Ti, and 0.100% or less for V. The reason for this is that excessive addition leads to excessive carbide precipitation and reduced workability. Furthermore, from the viewpoints of workability and cost, the total of Ti+Nb+V preferably does not exceed 0.100%.

[0036] Boron: 0~0.0030% Boron is another optional element, and when added, it is adjusted to 0.0003% to 0.0030%. The acceptable level of B as a residual element is 0.0003%. The addition of boron improves the hardenability and also helps to increase the tensile strength. To obtain these effects of B, the lower limit needs to be 0.0003%, preferably 0.0005%. However, if too much B is added, the effect saturates. Advantageously, B is adjusted to 0.0025% or less, preferably 0.0020% or less.

[0037] In another preferred embodiment of the present invention, Ti and / or Nb and / or V and / or Ni and / or Cu and / or Cr and / or Mo and / or B are not added as alloying elements in order to reduce the cost of the final product while still obtaining a cold rolled high strength steel strip with the desired properties.

[0038] Calcium: 0~0.0050%; Rare Earth Elements (REM): 0~0.0100% Additionally, compositions according to the present invention may optionally contain one or two elements selected from Ca and rare earth metals (REM) in amounts consistent with treatments for MnS inclusion control. When present as residual elements, the acceptable level is 0.0005%. When added as alloying elements, Ca is controlled to a value not exceeding 0.0050% and REM is controlled to a value not exceeding 0.0100%. Ca and / or REM combine with sulfur and oxygen, resulting in: Oxysulfides Generate this Oxysulfides does not have a detrimental effect on ductility (as in the case of elongated manganese sulfides which form in the absence of either Ca or REM). This effect saturates at Ca contents higher than 0.0050% or REM contents higher than 0.0100%. Preferably, the amount of Ca, if present, is adjusted to a value less than 0.0030%, more preferably equal to or less than 0.0020%. Preferably, the amount of REM, if present, is adjusted to a value equal to or less than 0.0080%, more preferably equal to or less than 0.0050%.

[0039] The balance of the steel composition comprises iron and unavoidable impurities.

[0040] The chemical composition of the steel according to the invention is compatible with the capabilities of conventional continuous production lines.

[0041] Microstructure The cold rolled steel strip heat treated according to the present invention has a complex microstructure comprising 20-55% polygonal ferrite (PF), acicular ferrite (AF) and high bainitic ferrite (HBF), up to 50% PF, 25-70% low bainitic ferrite (LBF) and distributed martensite (PM), 5-20% retained austenite (RA) and fresh martensite (M) in an amount of 0-15%.

[0042] In the present invention, the microstructures are formed into functional groups as they are observable using optical and scanning electron microscopes. Polygonal ferrite (PF) refers to ferrite formed in intercritical annealing or during 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 ferrite formed during austempering at temperatures between Bs and Bn. Low temperature bainitic ferrite (LBF) is ferrite formed during austempering at temperatures between Bn and Ms. Partitioned martensite (PM) refers to martensite formed during fast cooling (quenching) and overaging (partitioning) heat treatment.

[0043] Bainitic ferrite and partitioned martensite PM is obtained during quenching and partitioning when the quenching stop temperature is between Ms and Mf and partitioning is performed at temperatures between the quenching stop temperature and Bn. BF is obtained by the transformation of untransformed austenite during partitioning (overaging). The amount of PM depends on the quenching temperature. The amount of BF is a function of the temperature and time of partitioning. It should be noted here that in this application the expression "partitioned martensite" is used instead of tempered martensite. Generally, in metallurgy, tempered martensite contains carbide precipitates resulting from tempering. In the modified quenching and partitioning process according to the invention, the formation of carbides is delayed during overaging due to the presence of Si and Al and due to the very short time of the partitioning process. As a result, carbon partitions from martensite to austenite, resulting in a carbon-rich retained austenite with higher stability, and the partitioned martensite does not contain carbides. BF is present in the form of plates with ultrafine grain size. PM has a similar substructure to BF, but with finer sized ferrite laths, resulting in a finer grain size of retained austenite. Precipitation of carbides between the ferrite laths is known to adversely affect ductility and is suppressed by alloying with Si and / or Al. Bainitic ferrite does not contain carbides, in contrast to conventional bainite, which contains carbides. Bainitic ferrite also differs from (proeutectoid) ferrite, which has a lower dislocation density.

[0044] The carbide-free BF and PM microstructures provide high strength, which is superior to the intermediate hardness structures that have high dislocation densities and contain supersaturated carbon. Quality Ferrite structure (intermediate har df The bainitic ferrite structure also contributes to the desired high elongation because the bainitic ferrite does not contain carbides and fine retained austenite grains may be present at the boundaries of the lath-like bainitic ferrite.

[0045] In the present invention, bainitic ferrite is divided into two types: bainitic ferrite formed at high temperatures between Bs and Bn (called high bainitic ferrite (HBF)) and bainitic ferrite formed at low temperatures between Bn and Ms (called low bainitic ferrite (LBF)). When the cross-sections of steel strips subjected to etching with 3% nital are observed by scanning electron microscopy accompanied by EBSD analysis, the average aspect ratio of HBF (defined as the length of the minor axis divided by the length of the major axis) is higher than 0.35 and the average aspect ratio of LBF is lower than 0.35. The reason for this distinction is that the grain size and grain shape of bainitic ferrite (HBF) formed at the higher temperature range above Bn are similar to AF, making it difficult to distinguish HBF from AF using SEM. Like AF, HBF has a larger grain size, lower dislocation density and softness than LBF and acts to increase the elongation of the steel. On the other hand, due to the fine plate size, the strength of LBF is higher than that of HBF, which contributes to the strength of the steel strip and also enhances its formability. PM has a similar microstructure to LBF, except that the size of the ferrite laths and the grain size of the retained austenite decrease with decreasing forming temperature. However, this change is gradual, so that LBF and PM cannot be clearly distinguished by SEM observation. In this invention, LBF and PM are grouped as one microstructure, since their contributions to the properties of steel are also similar.

[0046] A feature of the high strength steel strip according to the present invention is that the bainitic ferrite may have a complex microstructure comprising HBF and LBF+PM, thus resulting in a high strength cold rolled steel strip with high elongation and good formability.

[0047] To obtain a good balance between high strength and elongation, 25-70% LBF+PM is required. If LBF+PM is present in too small an amount, the strength of the steel strip will be insufficient. However, if LBF+PM is present in too large an amount, the effect of other ferrites (PF, AF and HBF) and retained austenite on elongation may be compromised. Therefore, the sum of LBF+PM is 25-70%, preferably 35-65%. PM formed in the quenching process may accelerate the BF transformation rate of untransformed austenite during overaging. To ensure that the bainite transformation is completed in the time available in a typical current production line, the amount of PM can be adjusted by controlling the quenching stop temperature below the Ms point of the steel. The lower the quenching stop temperature, the more PM will be formed. More PM is required for steels with higher alloying element contents.

[0048] The formation of HBF in the present invention is due to the heating of the strip due to the latent heat generated by the bainite transformation or due to the heating by applying the hot dip galvanizing process. The formation of HBF in the present invention allows the bainite transformation, if any, to be accelerated as required, so that the bainite transformation can be completed in a limited time in the overaging section of the 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 20-55%, preferably 25-50%. As mentioned above, HBF has a similar function as PF and AF. In case sufficient amount of PF and AF is formed in the previous section and to obtain a steel strip with higher strength, the amount of HBF needs to be minimized to 0%. In case the amount of PF and AF is not sufficient, the amount of HBF may be increased. However, the amount of HBF needs to be controlled such that the total amount of PF, AF and HBF is 20-55%, preferably 25-50%.

[0049] Polygonal ferrite and acicular ferrite Proeutectoid ferrite is softer than bainitic ferrite and functionally increases the elongation of the steel strip. A certain amount of proeutectoid ferrite is introduced and the properties of the ferrite are controlled to increase the bainite transformation rate and the stability of the retained austenite by increasing the carbon content in the retained austenite, further increasing the elongation. Using the present invention, two types of proeutectoid ferrite can be produced during annealing depending on the formation temperature. The ferrite phase formed during austenitization in the transformation interval temperature or during cooling at high temperatures above the Bs temperature in the slow cooling section becomes polygonal or block-like and is called polygonal ferrite (PF). This type of ferrite has been proven to increase elongation but reduce yield strength and formability. The ferrite formed at low temperatures in the fast cooling section at temperatures between Bs and Ms is almost needle-like and has a smaller grain size than PF and is called acicular ferrite (AF). It is morphologically similar to HBF but has a relatively smaller amount of dislocations. The presence of AF can increase elongation without sacrificing strength and formability.

[0050] Since PF, AF and HBF have similar functions on tensile properties in the steel according to the invention, three types of these ferritic microstructures can be present, or one or two of them can be present. To ensure high elongation, the volume fractions of PF, AF and HBF are 20% or more, preferably 25% or more. In any case, the total amount of PF, AF and HBF should be controlled to 55% or less, preferably 50% or less. If these ferritic microstructures are too abundant (more than 55%), the final microstructure will not contain enough low bainitic ferrite and therefore the strength will be reduced.

[0051] When PF is present in the steel, the grain size, morphology and distribution of PF must be controlled. The steel strip can have higher elongation by having a smaller grain size and a more dispersed distribution of PF. According to the present invention, when observed by SEM or optical microscope, the PF structure is equiaxed embedded between the BF structure and uniformly distributed as smaller grains, whereas the morphological structure of PF in conventional TRIP steel strip is elongated along the rolling direction. It is believed that this morphological structure can uniformly distribute the stress during processing and can maximize the TRIP effect of the retained austenite. To obtain this morphological structure, the amount of PF formed during soaking should be 50% or less, preferably 10-40%. This embodiment is particularly suitable for steel compositions containing relatively large amounts of Mn, Al and Si, where the recrystallization rate of ferrite is slow. PF is a partially recrystallized microstructure and has a higher hardness than recrystallized ferrite. The presence of PF with a partially recrystallized microstructure is beneficial for local ductility. Advantageously, the particle size of the PF according to the invention is less than or equal to 10 μm, preferably less than or equal to 8 μm, and more preferably less than or equal to 5 μm.

[0052] In one embodiment of the present invention, the amount of PF is preferably 0%. In this case, the total amount of AF and HBF is controlled so that AF+HBF is 20-55%, preferably 25-50%.

[0053] Retained austenite Residual austenite (also known as retained austenite) refers to regions that exhibit FCC phase (face-centered cubic lattice) in the final microstructure. Retained austenite enhances ductility, in part through the TRIP effect, which manifests itself as an increase in uniform elongation. To exert the TRIP effect, the volume fraction of retained austenite is at least 5%, preferably at least 7%. Below 5%, the desired level of ductility and uniform elongation is not achieved. The upper limit is determined mainly by the composition and the processing parameters of the production line. For a given composition, if the amount of retained austenite is too high, the carbon content of the retained austenite will be too low. In that case, the retained austenite will not be stable enough and local ductility (stretch flangeability) may decrease to unacceptable levels. Therefore, the upper limit of the volume fraction of retained austenite is 20%, preferably 15%.

[0054] The carbon concentration in the retained austenite affects the TRIP properties. The retained austenite is effective in improving the elongation properties, especially when the carbon concentration in the retained austenite is 0.90% or more. 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 more, preferably 0.95% or more. It is preferable that the carbon concentration in the retained austenite is as high as possible, with an upper limit of about 1.6% generally imposed by practical processing conditions. The carbon content and stability of the retained austenite can be adjusted by controlling the amount of ferrite.

[0055] Martensite Martensite (M) forms fresh in the final cooling section after austempering. Martensite reduces the yield point elongation and increases the work hardening exponent (n-value), which is desirable to achieve stability, neck-free deformation and strain uniformity in the final pressed part. Even with 1% fresh martensite in the final steel strip, the tensile response and therefore pressing behavior can be achieved comparable to conventional dual phase steels. However, the presence of fresh martensite impairs formability due to crack formation along the interface between martensite and LBF / HBF. Therefore, the amount of fresh martensite needs to be controlled below 15%, preferably below 10%.

[0056] carbide Carbides may be present as fine precipitates formed during austempering if the overaging temperature is too high or the overaging time is too long, or in the form of pearlite formed during cooling if the cooling rate is too slow. According to the invention, the microstructure of the steel of the invention is pearlite-free and carbide-free. By pearlite-free we mean that the amount of lamellar microstructure containing cementite and ferrite is less than 5%. By carbide-free we mean that the amount of carbides is below the detection limit of standard X-ray measurements.

[0057] Microstructural characterization The microstructural constituents classified in the steel according to the invention as described above can be quantitatively determined by the technique described below: The volume fractions of the constituents are measured by measuring the area fractions of polished surfaces, equating volume fractions with area fractions, using commercially available image processing programs or other suitable techniques.

[0058] PF, fresh M, RA and pearlite were examined by optical microscopy (OM) and RunThey can be distinguished using scanning electron microscopy (SEM). When samples etched with 10% sodium metabisulfite (abbreviated SMB) are characterized by OM, pearlite is observed as dark areas, PF is observed as light grey areas and fresh martensite is observed as light brown areas. When samples etched with 3% nital are characterized by SEM, PF is observed as smooth surface grains without retained austenite and pearlite is observed as lamellar microstructure containing both cementite and ferrite. The remaining microstructure is observed as grey areas characterized by plate or lath shaped ferritic substructure and RA is white or light grey areas distributed in the grains with no identifiable carbides. This microstructural group is called bainitic ferrite like microstructure. It may include mixtures of HBF, LBF, AF and PM. These microstructures cannot be clearly distinguished by OM and SEM due to their similar morphology.

[0059] In the present invention, the bainitic ferrite-like microstructure is further separated into two distinct groups by electron backscatter diffraction (EBDS). 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(α) to Fe(γ). Fresh martensite (M) is then separated from the bainitic ferrite-like microstructure by splitting Fe(α) into a partition with high average image quality (IQ) and a partition with low average IQ. The low IQ partition is classified as martensite, and the high IQ partition is classified as the bainitic ferrite-like microstructure. The method of distinguishing between the two types of groups is described below with reference to FIG. 1. In the bainitic ferrite (high IQ partition), regions with orientation differences between adjacent structures with tilt angles of 15° or more are identified. A region is considered to have the same crystal orientation and is defined in the present invention as a bainitic plate. For the bainite plate thus detected, determine the diameter of the circle having the same area as the bainite plate. Use the image of EBSD analysis with a magnification of 3000 to determine the diameter of the equivalent circle of the bainite plate. By fitting an ellipse to the bainite plate, the aspect ratio (defined as the length of the minor axis divided by the length of the major axis) is also determined. Similarly, measure the diameter of the equivalent circle of all bainite plates and the aspect ratio of the equivalent ellipse of all bainite plates in the measurement area (about 100 x 100 μm), and define the average value as the average grain size of the bainite plate and the average aspect ratio of the bainite plate of the present invention.

[0060] The inventors have systematically studied the effect of austempering temperature on the microstructure of bainitic ferrite. The austempering temperature is (Ms-200) to Bs. It has been found that the average grain size and the average aspect ratio of the bainitic plates increase with increasing austempering temperature. In particular, it is seen that the aspect ratio of the bainitic plates changes sharply between samples austempered below 440°C, which is less than the Bn of the steel having the composition used in the method according to the invention, and samples austempered above 460°C, which is more than the Bn. Therefore, a critical average value of the aspect ratio of 0.35 is defined to divide the two groups of bainitic ferrite-like microstructures. The aspect ratio of the group consisting of LBF and PM is less than or equal to 0.35, while the aspect ratio of the group consisting of HBF and AF is greater than 0.35.

[0061] In addition to the differences in bainite plate morphology and grain size, the complex crystallographic plate misorientation relationship between HBF, AF group and LBF, PM group is also different. The misorientation angle distribution of the steel according to the present invention is shown in Figure 2. The 60° peak is the misorientation between adjacent grains, which is consistent with the misorientation with Kurdjumov-Sachs (KS / KS) relationship, which is 60°. <111> and 60° <110> The peak at 53°-54° is due to the misorientation between grains obtained by phase transformation, which follows the Nishiyama-Wassermann and Kurdjumov-Sachs (NW / KS) relationships. According to the prior art (see A.-F. Gourgues, HM Flower, and TC Lindley, Materials Science and Technology, January 2000, Vol. 16, p. 26-40), acicular ferrite and upper bainite grow with the austenite matrix in the Nishiyama-Wassermann relationship, whereas lower bainite and martensite are composed of very complex packets that have the Kurdjumov-Sachs relationship with the matrix. Similar to these results, it is presumed that the peak at 53°-54° corresponds to the formation of HBF and AF, and the peak at 60° corresponds to the formation of LBF and PM. As the austempering temperature increases, the 53-54° peak becomes more distinct and the height of the peak increases, while the height of the 60° peak decreases. In the present invention, the relative amounts of the HBF, AF groups and the LBF, PM groups can be determined by the ratio of the heights of the two peaks.

[0062] The percentage of retained austenite measured by EBSD is always lower than the actual value, because some of the retained austenite is distributed between the bainite plates as a film with a very small grain size and cannot be detected by EBSD. Therefore, an intensity measurement based on XRD can be used, which is a conventional technique to measure the content of retained austenite. The volume fraction of retained austenite is determined at 1 / 4 thickness of the steel strip. The amount of cementite is also determined from this XRD analysis. The samples made from the steel strip are mechanically and chemically polished and then analyzed by an X-ray diffractometer using Co-Ka by measuring the integrated intensity of the (200), (220) and (311) planes of fcc iron and the integrated intensity of the (200), (211) and (220) planes of bcc iron. The amount of retained austenite (RA) and the lattice constant of the retained austenite were determined using Rietveld analysis. The C content in the retained austenite is calculated according to the following formula:

number

[0063] Mechanical properties Cold rolled steel strip having the above microstructure and composition and heat treated according to the invention has the following properties: A yield strength (YS) of at least 500 MPa; and / or A tensile strength (TS) of at least 850 MPa; and / or A total elongation (TE) of at least 14%; and / or has.

[0064] Preferably, the cold rolled and heat treated strip has all of these properties.

[0065] Method steps According to the method of the invention, the cold rolled steel strip with the above composition is heat treated to obtain the microstructure and properties. The cold rolled steel strip obtained by cold rolling is subjected to the heat treatment as in a continuous annealing line. A typical design of the method is shown diagrammatically in FIG. 3. The cold rolled steel strip is heated to a temperature above (Ac3-60), for example using a heating rate of at least 0.5° C. / s, preferably to a temperature between (Ac3-60) and (Ac3+20), usually to a predetermined austenitizing temperature T2, held at this temperature for a time t2 (step a), and then cooled to a temperature T4 below Ms, typically between Ms and (Ms-200), usually using a two-stage cooling with a controlled cooling rate (step b). The steel strip is then heated (step c), optionally with heating from a temperature below Ms, typically (T4-Ms), to a temperature above Ms, and treated for austempering at a temperature T5 between Ms and Bs, typically T4-Bn, for a time t5 (step d). The steel strip is then optionally heated to a temperature T6 between Bn and Bs for a time t6, which may be a temperature at which hot-dip galvanizing is possible. Finally, the steel strip is cooled to room temperature (step e). The process parameters and actions in each step are described below.

[0066] In the first step, the cold rolled steel is soaked at a temperature above (Ac3-60), for example between (Ac3-60) and (Ac3+20), for a soaking time t2 of 1 to 150 seconds, in order to achieve an at least partially austenitized microstructure. Annealing at a temperature above (Ac3-60) is necessary because the steel strip heat treated according to the invention requires a necessary amount of low-temperature transformation phases (e.g. bainitic ferrite and retained austenite) transformed from high-temperature austenite, and a certain amount of ferrite. If T2 is higher than (Ac3+20), the austenite grains will grow. This affects the grain size and distribution of the retained austenite, and also slows down the bainite transformation rate after the overaging process. Excessive amounts of fresh martensite formed during final cooling may form as a result of this incomplete bainite transformation, which increases strength but reduces ductility and formability. Furthermore, a homogeneous austenite structure with large grains is obtained, which suppresses the formation of PF and AF in the next cooling section, resulting in insufficient amount of ferrite within the current cooling schedule of the available production line. A homogeneous austenite structure with large grain size also leads to poor elongation of the steel strip. It has been observed that the homogeneity of austenite has a significant impact on the formation of PF and AF in the cooling section. When T2 is lower than (Ac3-60), PF is 50 %, and therefore the strength of the steel strip may be insufficient. On the other hand, the austenite formed may not be sufficient for the formation of LBF and retained austenite. Therefore, the annealing temperature needs to be (Ac3-60) or more, but advantageously (Ac3+20) or less, preferably (Ac3-50) to (Ac3+10). If t2 is longer than 150 seconds, the grain size of austenite and ferrite increases and the elongation decreases. If the annealing time t2 is shorter than 1 second, the reverse transformation to austenite may not proceed sufficiently and / or the carbides in the steel strip may not be sufficiently solid-dissolved. Therefore, the annealing time t2 is 1 second to 150 seconds, for example, 10 seconds to 120 seconds, preferably 1 second to 100 seconds.

[0067] In the next cooling step, the at least partially austenitized strip is cooled to a temperature T4 below Ms, typically Ms~ (Ms-200), the purpose of which is to adjust the amount of ferrite and partitioned martensite and to prevent the formation of pearlite.

[0068] In one embodiment of the present invention, the steel strip thus treated is directly cooled to a temperature T4 at a cooling rate of at least 15 ° C / s, so as to prevent the formation of pearlite. If the cooling rate is too slow, excessive amounts of ferrite may form, or even pearlite may form. Preferably, V4 is 20 ° C / s or more. The upper limit of the cooling rate is not particularly limited, as long as no temperature change occurs in the steel strip when cooling is stopped. In order to reduce the temperature non-uniformity of the steel strip, the cooling rate is preferably 100 ° C / s or less in standard equipment. If the average cooling rate exceeds 80 ° C / s in most available facilities, very large variations in the microstructure may occur both in the longitudinal direction and in the strip width direction. Therefore, a suitable cooling rate V4 is 15 to 80 ° C / s, preferably 20 to 70 ° C / s.

[0069] In another embodiment of the invention, in order to adjust the amount of ferrite and to homogenize the strip temperature, this cooling can be performed by two-stage cooling. This is compatible with most continuous annealing or hot-dip galvanizing lines currently in use that contain two connected cooling sections. First, the steel strip is cooled (called the slow cooling section) to a temperature T3 of 800-500°C, preferably 750-550°C, usually with a cooling rate V3 of at least 1°C / s, e.g., 2-15°C / s, preferably 3-10°C / s. Then, the steel strip is further cooled (called the fast cooling section) to a temperature T4, usually with a cooling rate V4 of at least 15°C / s, e.g., 15-100°C / s, preferably 20-70°C / s. Since the length of each section in the continuous annealing line is fixed, the cooling rates V3 and V4 for a given line speed can be controlled by adjusting the T3 temperature. The higher T3 is, the lower V3 is and the higher V4 is. During this cooling, some PF may be formed in the slow cooling section, and some AF may be formed in the fast cooling section. For a fixed line speed, the amount of PF formed in the slow cooling section mainly depends on T3, and the amount of AF mainly depends on V4. Therefore, T3 is selected in an appropriate range to adjust the amount of ferrite and prevent the formation of pearlite. If T3 is too low (e.g., less than 500 ° C), PF may be formed in an excessive amount in the slow cooling section, and AF may also be formed in an excessive amount in the fast cooling section, and further, if the resulting V4 is less than 15 ° C / s, pearlite may be formed. If T3 is too high (e.g., more than 800 ° C), PF may be formed insufficiently, and if the resulting V4 is too high, less AF will be formed. Therefore, T3 needs to be 800-500 ° C, preferably 750-550 ° C / s.

[0070] As mentioned above, in the conventional design of annealing line or galvanizing line, PF is obtained in the soaking step a) and the slow cooling section of step b), and AF is obtained in the fast cooling section of step b). To adjust the amount of ferrite, the soaking temperature T2 and the intermediate temperature T3 between the slow cooling section and the fast cooling section can be used. If a higher T2 is used, less PF is generated during soaking, then a lower T3 can be selected to obtain more PF in the slow cooling section and more AF in the fast cooling section. If a lower T2 is used, a sufficient amount of PF is generated during soaking, then a higher T3 can be selected to limit the amount of PF formed in the slow cooling section and the amount of AF formed in the fast cooling section.

[0071] After cooling to a temperature T4 below Ms, preferably Ms - (Ms-200), some amount of martensite is obtained. The lower the T4, the more martensite is formed. T4 is adjusted according to the steel composition to effectively accelerate the bainite transformation rate in the following partitioning process. For steels with a high content of alloying elements, a lower T4 is applied. If T4 is too high, the amount of PM formed will be insufficient. The bainite transformation of untransformed austenite cannot be completed in the overaging (partitioning) stage, and excessive fresh martensite may be formed in the following cooling process to ambient temperature. If T4 is too low, the PM will be excessively formed and the amount of retained austenite will be reduced. Therefore, T4 is preferably Ms - (Ms-200), more preferably (Ms-50) - (Ms-150). Since the amount of PM depends only on the T4 temperature, the steel strip is heated as fast as possible to the distribution temperature of Ms to Bs in order to make available for the bainite transformation the remaining time of the total time available in the overaged section. In practice, depending on the heating capacity of the production line and to facilitate the homogenization of the temperature of the steel strip, the total time t4 of step c) including any holding time is preferably not more than 10 seconds, more preferably not more than 5 seconds. Optionally, the heating step c) may include a short heating at a temperature below Ms, for example Ms to (Ms-200), for example (Ms-50) to (Ms-150).

[0072] In the subsequent heating step d), the cooled strip is heated at a temperature T5 between Ms and Bs, preferably below Bn, for a time t5 between 30 and 120 seconds. By heating to a temperature T5 in this range, the untransformed austenite is transformed into low bainitic ferrite (LBF) and carbon partitioning occurs in the previously formed martensite. If T5 is too low, the bainite transformation is too slow and the bainite transformation is insufficient during overaging, and fresh martensite may be formed in an excessive amount during cooling after overaging, which increases the strength but reduces the required elongation. On the other hand, the carbon partitioning may be insufficient to stabilize the retained austenite. If T5 is too high, excessive HBF is obtained in the overaged section and the required strength cannot be obtained. The preferred range of T5 is (Bn-50) to Bn in order to achieve a high bainite transformation rate. If the heating time t5 is less than 30 seconds, the bainite transformation is incomplete and the carbon partitioning in martensite and bainite is also insufficient. If t5 is greater than 120 seconds, there is a risk that carbides will start to form and the carbon content in the retained austenite will decrease. The maximum time for t5 is limited, among other things, by the total time available at a given speed of the production line. Preferably, t5 is between 40 and 100 seconds.

[0073] The temperature of the steel strip can be increased by the latent heat generated by the bainite transformation during overaging, so that when the steel strip reaches a temperature higher than Bn, a small amount of high temperature bainitic ferrite is formed.

[0074] The thus heated strip is then cooled to ambient temperature according to the performance of the production line, during which some fresh martensite may be formed. The steel strip is then cooled to below 300 ° C with a cooling rate V7 of at least 1 ° C / s, preferably at least 5 ° C / s, and then further cooled to ambient temperature. The cooling to ambient temperature can be forced cooling or uncontrolled natural cooling. In a practical embodiment, the heated steel strip is cooled to a temperature T7 of (Ms-50) to Mf with a cooling rate V7 of 5.0 to 10.0 ° C / s. Further cooling from T7 to ambient temperature is carried out with a cooling rate V8 of preferably 5.0 to 20.0 ° C / s, more preferably 6.0 to 15.0 ° C / s.

[0075] In one embodiment, to ensure obtaining a partially austenitized cold rolled strip with fine grain size, the soaking step is carried out at an interval annealing temperature between (Ac3-50) and (Ac3+10) with a soaking time t2 preferably between 1 and 100 seconds. The fraction of PF formed at the soaking temperature is advantageously less than or equal to 40%.

[0076] Advantageously, the heating step before the soaking step is carried out in two sub-steps. The two sub-steps include heating the cold rolled strip to a temperature T1 between 680 and 740°C, preferably between 700 and 720°C, at a heating rate V1 between 10.0 and 30.0°C / s, preferably between 15.0 and 25.0°C / s, and further heating from temperature T1 to the soaking temperature range at a heating rate V2 between 0.5 and 4.0°C / s, preferably between 1.0 and 3.0°C / s. During the slow heating from T1 to the soaking temperature T2, recovery and recrystallization take place in the ferrite and during the austenite transformation, the dissolution of carbides and ferrite takes place. T1 and V2 influence the progress of these processes, affecting the grain size of the austenite and the homogeneity of the distribution of the alloying elements in the austenite phase. Advantageously, the dissolution and dissolution of all carbides take place. coarse The soak time t2 is controlled according to the heating rate V2 to ensure avoidance of uneven austenite grain size.

[0077] In one embodiment, the method according to the invention comprises a further heating step between heating step d) and cooling step e), in which the steel strip resulting from step d) is subjected to an additional heating at a temperature between Bs and Bn, preferably between (Bs-50) and Bn, typically at a fixed temperature T6. The additional heating time t6 is advantageously between 5 and 30 seconds, preferably between 10 and 20 seconds. This additional heating increases the formation of the retained austenite into high-temperature bainitic ferrite to complete the bainitic transformation. As a result, the amount of martensite formed in the subsequent cooling section is reduced, which can improve strength and ductility. Carbon is also further distributed in the retained austenite, making it more stable. If this additional heating is applied to a given overaged section, and therefore for a given total time therein, the time t5 is further reduced to fill the available time. For example, the sum of t4+t5+t6 is between 30 and 120 seconds.

[0078] In a preferred embodiment, this additional heating comprises an integrated hot dip galvanizing process, and the steel strip resulting from step c) is coated with a Zn coating or a Zn alloy based coating.

[0079] The steel strip heat treated according to the invention may be provided with a coating, advantageously a zinc coating or a zinc alloy-based coating. Advantageously, the zinc-based coating is a galvanized or galvannealed coating. The Zn-based coating may comprise a Zn alloy containing Al as alloying element. A preferred zinc bath has a composition of 0.10-0.35% Al, the balance being zinc and unavoidable impurities. Another preferred Zn bath containing Mg and Al as main alloying elements has the following composition: 0.5-3.8% Al, 0.5-3.0% Mg, optionally up to 0.2% of one or more additional elements, the balance being zinc and unavoidable impurities. Examples of additional elements include Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr and Bi.

[0080] The coating, e.g. a protective coating of Zn or a Zn alloy, may be applied in a separate step, preferably a hot dip galvanizing step is integrated in the method according to the invention, as explained above.

[0081] Optionally, the annealed and zinc-coated strip according to the invention may be subjected to a temper rolling treatment in order to fine-tune the tensile properties and to modify the surface appearance and roughness according to the specific requirements arising from the intended use.

[0082] The cold rolled steel strip itself is usually produced according to the following general method: the above mentioned steel composition is prepared and cast into a slab. The cast slab is reheated at a temperature between 1100 and 1300 °C and then processed by hot rolling. Hot rolling of the slab is usually carried out in 5 to 7 stands to a final dimension suitable for further cold rolling. Finish rolling is usually carried out in a fully austenitic state above 800 °C, advantageously at or above 850 °C. The strip thus obtained from the hot rolling step can be coiled, for example at a coiling temperature usually below 700 °C. The hot rolled strip is pickled and cold rolled to obtain a cold rolled steel strip with the appropriate gauge. Preferably, the cold rolling reduction is typically between 30 and 80%. To reduce the rolling load during cold rolling, the coiled strip or the semi-cold rolled strip may be subjected to hot batch annealing. The batch annealing temperature should be 500-700℃.

[0083] Thin slab casting, strip casting, etc. may also be applied, in which case it is permissible for the production method to skip at least part of the hot rolling process.

[0084] The present invention also relates to heat treated cold rolled steel strip having the composition and microstructure as described above.

[0085] The invention also resides in articles, such as structural, engineering or automotive parts, manufactured from the heat treated cold rolled strip according to the invention. EXAMPLES

[0086] Steels having the compositions shown in Table 1 were cast using vacuum induction into 25 kg ingots with dimensions of 200 mm x 110 mm x 110 mm. Cold rolled strips of 1 mm thickness were produced using the following process schedule: The ingot was reheated at 1225°C for 2 hours. · The ingot was rough rolled from 140mm to 35mm. The rough rolled ingot was reheated at 1200℃ for 30 minutes. Hot rolled from 35mm to 4mm in 6 passes. Run-out table cooling: Cooling was performed from the finishing rolling temperature (FRT) (approximately 850-900°C) to 600°C at a rate of 40°C / sec. · Furnace cooling: The strips were transferred to a furnace preheated to 600°C and then cooled to room temperature to simulate the cooling process. · Pickling: The hot rolled strip was then pickled with HCl at 85°C to remove the oxide layer. Cold rolling: The hot rolled strip was cold rolled into 1mm strip. Heat treatment according to the invention: Using cold rolled sheets of suitable size, the annealing process was simulated using a continuous annealing simulator (CASIM). Samples for microstructural observation, tensile testing and hole expansion testing were machined from the thus treated strips.

[0087] Dilatometry measurements were performed on cold rolled samples with dimensions 10 mm x 5 mm x 1 mm (length along the rolling direction). Dilatometry tests were performed with a Bahr dilatometer type DIL805. All measurements were performed according to SEP1680. The critical phase transformation points Ac3, Ms and Mf were determined from quenched dilatometry curves. Bs and Bn were predicted using the available software JmatPro10. The phase fractions during annealing for different process parameters were determined from dilatometry curves simulating the annealing cycle.

[0088] The microstructure was determined by optical microscopy (OM) and scanning electron microscopy (SEM) using commercially available image processing programs. The microstructure was observed at 1 / 4 thickness in the rolling direction and in the cross section normal to the rolling plane of the steel strip. The scanning electron microscope (SEM) used for EBSD measurements is a Zeiss Ultra55 machine 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) software OIM (Orientation Imaging Microscopy) Data Collection. EBSD scans were evaluated with TSL OIM Analysis software. The EBSD scan area was 100 × 100 μm in all cases, the step size was 0.1 μm, and the scan rate was about 80 frames / s.

[0089] The retained austenite was determined by XRD with a D8 Discover GADDS (Bruker AXS) using Co-Kα radiation according to DIN EN 13925. Quantitative determination of the phase ratios was carried out by Rietveld analysis.

[0090] Tensile testing - JIS 5 test specimens (gauge length = 50 mm, width = 25 mm) were machined from the annealed strips with the tensile direction parallel to the rolling direction. Room temperature tensile tests were performed on a Schenk TREBEL testing machine according to the NEN-EN10002-1:2001 standard to determine the tensile properties (yield strength YS (MPa), ultimate 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.

[0091] Using the representation of Figure 3, the process parameters are shown in Table 2. In CASIM, the steel strip is cooled to T4 in V4 and then heated to T5 within 5 seconds. The microstructure and tensile properties obtained are shown in Table 3. All steel compositions of the invention achieve the microstructure and tensile property requirements under the specified processing parameters. Steel A79 containing 0.5% Cr could not achieve the required elongation when T4 is above Ms because the bainite transformation is not sufficient (Examples 13 and 14). However, when T4 is below Ms, some PM is formed, so that the bainite transformation can be completed and the steel strip obtains the required properties (Examples 11 and 12).

[0092] [Table 1]

[0093] [Table 2]

[0094] [Table 3]

Claims

1. 1. A method for heat treating cold rolled steel strip, comprising the steps of: The method comprises the steps of: a) soaking the cold rolled steel strip at a temperature between (Ac3-50) and (Ac3+10) for a soaking time t2 between 1 and 150 seconds, thereby obtaining a cold rolled steel strip having an at least partially austenitized microstructure; b) cooling the soaked steel strip obtained from step a) to a temperature T4 of Ms to (Ms-200), either by directly cooling the soaked steel strip obtained from step a) to a temperature T4 at a cooling rate V4 of at least 15°C / s or by cooling the soaked steel strip obtained from step a) to a temperature T3 of 800-500°C at a cooling rate V3 of at least 1°C / s and then cooling from temperature T3 to temperature T4 at a cooling rate V4 of at least 15°C / s; c) heating the cooled steel strip obtained from step b) to a temperature of Bs to Ms; d) heating the hot steel strip at a temperature between Bs and Ms for a time t5 between 30 and 120 seconds; e) cooling the heated steel strip to ambient temperature. so that the steel strip comprises, in volume percent: Polygonal ferrite (PF) + acicular ferrite (AF) + high bainitic ferrite (HBF): 20-55 Polygonal ferrite (PF): 0 to 50 Low bainitic ferrite (LBF) + distributed martensite (PM): 25-70 Retained austenite (RA): 5-20 Martensite (M): 0 to 15 and having a pearlite-free microstructure, where "pearlite-free" means that the amount of lamellar microstructure containing cementite and ferrite is less than 5%; The steel strip has the following composition, in mass %, C: 0.15-0.35 Mn: 1.50-4.00 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 In some cases, 0<Cr≦1.00 0<Cu≦0.20 0<Ni≦0.50 0<Mo≦0.50 0<Nb≦0.10 0<V≦0.10 0<Ti≦0.10 0<B≦0.0030 0<Ca≦0.0050 0<REM≦0.0100 One or more elements selected from Iron and unavoidable impurities: balance having The sum of (Si+Al) is ≧0.60; The process wherein REM is one or more rare earth metals.

2. 2. The method according to claim 1, wherein step c) comprises heating the cooled steel strip obtained from step b) at a temperature T4 of between Ms and (Ms-200).

3. The method according to claim 1, wherein the total time t4 of step c) is 1 to 10 seconds.

4. A method according to any one of claims 1 to 3, wherein the heat-soaking time t2 is 1 to 100 seconds.

5. The method according to claim 1, wherein the cooling rate V3 is 2.0 to 15.0°C / sec and the temperature T3 is 750 to 550°C.

6. The method according to claim 1, wherein the cooling rate V4 is 20.0 to 70.0°C / sec and the temperature T3 is 750 to 550°C.

7. 7. The method according to any one of claims 1 to 6, further comprising, prior to step a), heating the cold rolled steel strip to a temperature of (Ac3-50) to (Ac3+10) at a heating rate of at least 0.5°C / s.

8. The method according to claim 1, further comprising, prior to step a), heating the cold rolled steel strip to a temperature T1 of 680-740°C at a heating rate V1 of 10.0-30.0°C / s, and further heating from temperature T1 to a temperature of (Ac3-50)-(Ac3+10) at a heating rate V2 of 0.5-4.0°C / s.

9. The method according to claim 8, wherein the heating rate V1 is 15.0 to 25.0° C. / sec, the temperature T1 is 700 to 720° C., and the heating rate V2 is 1.0 to 3.0° C. / sec.

10. The method according to any one of claims 1 to 9, wherein the heating in step d) is carried out at a temperature of Bn to (Ms+50).

11. The method of claim 10, wherein the heating in step d) is carried out for a time t5 of 40 to 100 seconds.

12. The heating in step d) is carried out at a temperature of Bn to Ms; 12. The method according to any one of the preceding claims, wherein the method comprises an additional heating step between step d) and step e) of heating the steel strip obtained from step d) at a temperature between Bs and Bn.

13. The method of claim 12, wherein the heating in the additional heating step is carried out for a time t6 of 5 to 30 seconds.

14. 14. The method of claim 12 or 13, wherein the additional heating step comprises a hot dip galvanizing process.

15. The method according to any one of the preceding claims, further comprising, after the heat treatment, a coating step of coating the heat treated steel strip with a protective coating.

16. The microstructure comprises, in volume percent: Polygonal ferrite (PF) + acicular ferrite (AF) + high bainitic ferrite (HBF): 25-50 Polygonal ferrite (PF): 10-40 Low bainitic ferrite (LBF) + distributed martensite (PM): 35-65 Retained austenite (RA): 7-15 Martensite (M): 0 to 10 and / or The method according to any one of claims 1 to 15, wherein the C content in the retained austenite (RA) is 0.90 wt% or more.

17. The resulting steel strip has the following properties: Yield strength (YS) of ≧500 MPa; Tensile strength (TS) of ≧850 MPa; Total Elongation (TE) ≧14% The method according to any one of claims 1 to 16, having at least one of the following characteristics:

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