Heat Treatment of High-Strength Cold-Rolled Steel Strip

Abstract

1. A method of heat treating high strength cold rolled steel strip, comprising the steps of: a) soaking the cold rolled steel strip; b) cooling the soaked steel strip; c) heating the cooled steel strip; d) cooling the heated steel strip to ambient temperature The method includes the steps of: (a) forming a steel strip having a microstructure containing various ferrite, retained austenite, and martensite; (b) forming a steel strip having a microstructure containing various ferrite, retained austenite, and martensite; and (c) forming a steel strip having a microstructure containing various ferrite, retained austenite, and martensite. The main components of the steel composition include, in addition to iron, carbon, manganese, silicon, and aluminum.

Description

Technical Field

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

Background Art

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

[0003] HSLA (High-Strength Low-Alloy) steels contain micro-alloying elements. They 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 currently typical of the high-ductility and high-strength steels used in the automotive industry. In DP steels, since martensite exists in a ferrite matrix, a tensile strength exceeding 450 MPa and good cold formability can be obtained.

[0005] In order to simultaneously achieve a high yield strength / tensile strength ratio and an even higher tensile strength, i.e., exceeding 800 MPa, steels having a composite (CP) microstructure 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 are generally inferior in stretch flange formability. Therefore, their use is limited to automotive parts that do not require high formability.

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

[0007] However, due to the limitations of current continuous manufacturing lines, the beneficial combination of strength and ductility properties expected cannot be obtained with currently available steel recipes. These limitations include, among other things, that the current reheating furnaces of the facilities in continuous annealing (CA) lines and continuous galvanizing (CG) lines are often only suitable when the steel strip is subjected to transformation interval heat treatment 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 a certain range. The overaging time available for many CA / CG lines is also limited. For example, this time is less than about 160 seconds, which imposes a severe time limit for completing any desired transformation during overaging.

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

[0009] EP2831296B1 and EP2831299 disclose TBF steel having a tensile strength of at least 980 MPa, which can be produced even on a conventional production line. However, the preferred overaging / austempering time is 280 to 320 seconds, which is too long to enable production on most conventional production lines. In other words, the bainite transformation rate is too slow in the overaging section of a conventional production line to complete the bainite transformation within a limited time to obtain the desired microstructure. SUMMARY OF THE INVENTION PROBLEMS TO BE SOLVED BY THE INVENTION

[0010] An object of the present invention is to provide a cold-rolled steel strip having a combination of a desired high tensile strength and high total elongation (TE) characteristics at a reasonable yield strength (YS), for example, having TS≧850 MPa, TE≧14% and / or YS≧500 MPa, particularly a steel strip for use in automotive applications, or a suitable alternative.

[0011] A further object of the present invention is to provide a method for heat-treating a cold-rolled steel strip to obtain a combination of desired characteristics as described above, particularly a heat treatment that can be carried out using an existing production line or a suitable alternative.

[0012] Another object of the present invention is to provide a high-silicon cold-rolled steel strip having a combination of desired characteristics that can be produced on a conventional industrial production line.

[0013] Another object of the present invention is to enable the completion of bainite transformation on a conventional production line by means of a steel composition for a high-strength cold-rolled steel strip and its heat treatment to obtain a desired microstructure.

Means for Solving the Problems

[0014] In consideration of these, the present invention is a method for heat-treating a cold-rolled steel strip, wherein the method comprises the following steps: a) Soaking a cold-rolled steel strip at a temperature of (Ac3 - 60) or higher for a soaking time t2 of 1 to 150 seconds to obtain a cold-rolled steel strip having a microstructure that is at least partially austenitized; b) Cooling the soaked steel strip obtained in step a) to a temperature T4 between Bn and Ms; c) Heating the cooled steel strip obtained in step b) at a temperature between Bs and T4 for a time t5 of 30 to 300 seconds; d) Cooling the heated steel strip to ambient temperature and as a result, the steel strip has, in volume %, Polygonal ferrite (PF) + acicular ferrite (AF) + high bainitic ferrite (HBF): 20 - 55 Polygonal ferrite (PF): 0 - 45 Lower bainitic ferrite (LBF): 20 - 65 Retained austenite (RA): 5 - 20 Martensite (M): 0 - 20 and has a microstructure containing The steel strip has, in mass %, the following composition: 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 Optionally, 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 One or more elements selected from Iron and inevitable impurities: the balance having the total of (Si + Al) is ≥ 0.60, and REM is one or more rare earth metals, providing the above method.

[0015] The method of the present invention makes it possible to produce a cold-rolled steel strip having a specific composition and microstructure, and a desirable combination of properties for automotive parts that require 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, controlling its morphology, obtaining fine grains of austenite by controlling the maximum annealing temperature and time, and using a modified overaging process in the production line.

[0017] This method according to the present invention can be carried out within the limitations regarding the maximum temperature, cooling rate range, and overaging time window of the annealing section at the production speeds typical of these production lines, using existing continuous annealing lines and galvanizing lines.

[0018] The heat-treated cold-rolled steel strip can be Zn-coated, for example, by hot-dip galvanizing or electro-galvanizing. The hot-dip galvanizing process can be easily integrated with the heat treatment according to the present invention.

[0019] The terms used to describe the critical transformation temperatures of steel are well known to those skilled in the art and are described below. Ae3: The equilibrium temperature of the transformation from ferrite to austenite and from austenite to ferrite. Ac3: The temperature at which the transformation from ferrite to austenite is completed during heating. Ac3 is usually higher than Ae3, but tends to approach Ae3 because the heating rate tends to zero. In the present invention, Ac3 is measured at a heating rate of 3 °C / second. 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 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. The practical problem regarding Mf is that during cooling, the proportion of martensite only asymptotically approaches the maximum achievable amount. This means that the formation of martensite continues to very low temperatures. Therefore, for practical reasons, in the context of the present invention, Mf is considered the temperature at which 90% of the maximum achievable amount of martensite is formed.

[0020] These critical phase transformation temperatures can be determined by dilatometry experiments. Alternatively, the Ac3, Bs, and Ms points of the steel according to the present invention can be pre-calculated based on its composition using commercially available software, such as JmatPro, or the following empirical formula.

Equation

[0021] In this specification, unless otherwise specified, all temperatures are expressed in degrees Celsius, all compositions are indicated in weight percent (wt%), and all microstructures are indicated in volume percent (vol%).

Brief Description of the Drawings

[0022]

Figure 1

Figure 2

Figure 3

Embodiments for Carrying Out the Invention

[0023] After the description of the composition, the steps and microstructure of the method according to the present invention are shown.

[0024] Composition Carbon: 0.15 to 0.35% A sufficient amount of carbon is necessary for strength and stabilization of retained austenite, and the latter provides the TRIP effect. Considering this, the amount of carbon is 0.15% or more, preferably 0.17% or more, to ensure the required strength and elongation. Increasing the carbon content increases the strength of the steel, the amount of retained austenite, and the carbon content in the retained austenite. However, when the carbon content is higher than 0.25%, the weldability of the steel significantly decreases. In the case of applications that require 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 present invention for obtaining the described microstructure. Its main function is to prevent carbon from precipitating in the form of iron carbide (most commonly cementite) and to suppress the decomposition of retained austenite. Silicon contributes to strength characteristics and appropriate transformation behavior. Furthermore, silicon contributes to the improvement of ductility, work hardening property, and stretch flange formability by suppressing the grain growth of austenite during annealing. A minimum of 0.50% Si is required to sufficiently suppress carbide formation. However, when the silicon content is high, silicon oxide is formed on the strip surface, deteriorating the surface quality, coating property, and workability. Furthermore, the Ac3 temperature of the steel having this composition increases as the silicon content increases. This may affect the feasibility of producing steel strips using existing production lines in terms of the maximum top temperature achievable in the annealing section. Considering these, the silicon content is 2.00% or less. Preferably, from the viewpoints of suppressing carbide formation, promoting austenite stabilization, and wettability, Si is 0.80 - 1.80%. More preferably, Si is 1.00 - 1.60%.

[0026] Aluminum: 0.01 - 1.50% The main function of aluminum is to deoxidize the molten steel before casting. For the deoxidation of molten steel, more than 0.01% Al is required. Furthermore, aluminum has a similar function to silicon, preventing the formation of carbides and stabilizing retained austenite. Al is considered to be less effective than Si. It has no significant impact on strengthening. A small amount of Al can be used to partially replace Si, adjust the transformation temperature and the critical cooling rate, obtain acicular ferrite (AF), and increase the bainite transformation rate. Aluminum is added for these purposes. Therefore, the Al content is preferably 0.03% or more. A high level of Al may increase the transformation point from ferrite to austenite to a level incompatible with the current equipment, resulting in difficulty in obtaining a microstructure where the main phase is a low-temperature transformation product. As the Al content increases, the risk of cracking during casting increases. Considering this, the upper limit is 1.50%, preferably 1.00%, more preferably 0.70%.

[0027] Regarding the relationship between the ratio of Si and the ratio of Al, the composition satisfies the condition of Si + Al ≥ 0.60, preferably Si + Al ≥ 1.00. Advantageously, the Al content is less than 0.5 times the Si content.

[0028] Manganese: 1.50 - 3.00% Manganese is necessary to obtain the microstructure of the steel strip according to the present invention from the viewpoints of hardenability and stabilization of retained austenite. Mn also affects the formation of primary ferrite at high temperatures and the bainitic ferrite transformation rate. To suppress the formation of carbides in bainitic ferrite, a certain amount of Si and / or Al is necessary. The Ac3 temperature rises as the contents of Si and Al increase. Mn is also adjusted to balance the increased phase transformation point Ac3 as a result of the presence of Si and Al. When the Mn content is less than 1.50%, it is difficult to obtain the described microstructure. Therefore, Mn needs to be added in an amount of 1.50% or more. However, when Mn is present in an excessive amount, macrosegregation is likely to occur, resulting in unfavorable band formation in the steel. Furthermore, an excessive amount of Mn slows down the bainite transformation rate, and as a result, the amount of fresh martensite becomes excessively large, and as a result, the elongation flange formability also decreases. Therefore, the Mn content is 3.00% or less, preferably 2.80% or less, more preferably 1.80 ≦ Mn ≦ 2.60%.

[0029] Phosphorus: < 0.050% Phosphorus is an impurity in the 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 the steel. S forms sulfide-based inclusions, such as MnS, generates cracks, and reduces the elongation flange formability of the steel. The S content is preferably as low as possible, for example, less than 0.020%, preferably less than 0.010%, more preferably less than 0.005%.

[0031] Nitrogen: < 0.0080% Nitrogen is another inevitable impurity in steel. It precipitates as nitrides containing microalloying elements and exists in solid solution to contribute to strengthening. Excessive nitrides reduce elongation, elongation 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 contain one or more optional component elements as follows.

[0033] Copper: 0 to 0.20% Copper is not necessary in embodiments of the steel having this composition, but may be present. In some embodiments, depending on the manufacturing process, the presence of Cu may be inevitable. Copper less than 0.05% is considered a residual element. Copper as an alloying element can be added up to 0.20%. This facilitates the removal of high-Si scale formed in the hot rolling stage for manufacturing the starting steel strip, and further improves corrosion resistance when the cold-rolled steel strip is used as it is without surface treatment, or improves wettability with molten zinc in the case of a Zn-coated strip. Cu promotes the bainite structure, causes solid solution hardening, and can contribute to precipitation hardening because it precipitates as ε-copper from the ferrite matrix. Cu also reduces the amount of hydrogen entering the steel, and as a result, improves the stress corrosion cracking properties. However, when an excessive amount is added, Cu causes high-temperature brittleness. Therefore, when adding Cu, the Cu content is 0.20% or less.

[0034] Chromium: 0 to 0.35%; Nickel: 0 to 0.50%; Molybdenum: 0 to 0.30% Chromium, nickel, and molybdenum are not essential elements but may exist as residual elements in the steel composition. The acceptable levels of Cr, Ni, or Mo as residual elements are 0.05% each. As alloying elements, they have a similar effect of improving the hardenability of the steel, promoting the formation of bainitic ferrite, and stabilizing retained austenite at the same time. Therefore, Cr, Ni, and Mo are effective in microstructure control. To fully obtain this effect, the content of Cr, Ni, or Mo in the steel is preferably at least 0.05%. However, when each is added in excess, the effect saturates, the bainite transformation rate becomes excessively slow, and the required microstructure cannot be obtained in a production line where the overaging time is limited. Therefore, the amounts of Cr and Mo are limited to a maximum of 0.35% and 0.30% respectively. Ni is simply used to reduce the high-temperature brittleness characteristics when a relatively large amount of Cu is added. This effect of Ni becomes prominent when the Ni content is > [Cu(%) / 3]. The amount of Ni, if present, is limited to a maximum of 0.50%.

[0035] Niobium: 0 to 0.100%; Vanadium: 0 to 0.100%; Titanium: 0 to 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 product and the final product. These elements have a precipitation strengthening effect and may change the morphology of bainitic ferrite. They also contribute beneficially to the optimization of properties, such as the ductility and bendability of the stretched edge, for applications that depend on these properties. To obtain these effects, the lower limit of any of these elements, if present, needs to be adjusted to 0.005% or more. When the content of each of Nb, Ti, and V exceeds 0.10%, the effect saturates. Therefore, when adding these elements, their content is adjusted to 0.005% - 0.100%. Preferably, the upper limit is 0.050% or less for Nb and Ti, and 0.100% or less for V. This is because excessive addition leads to excessive carbide precipitation and a reduction in workability. Furthermore, the total of Ti + Nb + V preferably does not exceed 0.100% from the viewpoints of workability and cost.

[0036] Boron: 0 - 0.0030% Boron is an element of another optional component and, when added, is adjusted to 0.0003% - 0.0030%. The acceptable level of B as a residual element is 0.0003%. The addition of boron improves hardenability during quenching and also helps increase the tensile strength. To obtain these effects of B, the lower limit needs to be 0.0003%, preferably 0.0005%. However, when excessive 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, in order to reduce the cost of the final product while still obtaining a cold-rolled high-strength steel strip having desired properties, 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.

[0038] Calcium: 0 - 0.0050%; Rare earth elements (REM): 0 - 0.0100% Furthermore, the composition according to the invention may optionally contain one or two elements selected from Ca and rare earth metals (REM) in an amount consistent with the treatment for MnS inclusion control. When present as residual elements, the acceptable level is 0.0005%. When added as alloying elements, Ca is adjusted to a value of 0.0050% or less and REM is adjusted to a value of 0.0100% or less. Ca and / or REM combine with sulfur and oxygen, and as a result, Oxysulfide is produced. This Oxysulfide has no adverse effect on ductility (such as the adverse effect in the case of the elongated manganese sulfide formed when neither Ca nor REM is present). This effect saturates when the Ca content exceeds 0.0050% or the REM content exceeds 0.0100%. Preferably, the amount of Ca, when present, is adjusted to a value of 0.0030% or less, more preferably 0.0020% or less. Preferably, the amount of REM, when present, is adjusted to a value of 0.0080% or less, more preferably 0.0050% or less.

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

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

[0041] Microstructure The cold-rolled steel strip heat-treated according to the invention has a composite microstructure containing 20 - 55% polygonal ferrite (PF), acicular ferrite (AF) and high bainitic ferrite (HBF), up to 45% PF, 20 - 65% low bainitic ferrite (LBF), 5 - 20% retained austenite (RA) and 0 - 20% fresh martensite (M).

[0042] In the present invention, the microstructures form functional groups so as to be observable using an optical microscope and a scanning electron microscope. Polygonal ferrite (PF) refers to ferrite formed during intercritical annealing or during slow cooling at a temperature exceeding Bs. Acicular ferrite (AF) refers to ferrite formed during cooling at a temperature between Bs and Ms. High-temperature bainitic ferrite (HBF) is bainitic ferrite formed during austempering at a temperature between Bs and Bn. Low-temperature bainitic ferrite (LBF) is bainitic ferrite formed during austempering at a temperature between Bn and Ms.

[0043] Bainitic ferrite structure Bainitic ferrite (BF) is formed during heat treatment when the austempering temperature is between Ms and Bs. BF exists in the form of ultrafine-grained plates. Precipitation of carbides between the laths of ferrite is known to have an adverse effect on ductility and is suppressed by alloying with Si and / or Al. Bainitic ferrite does not contain carbides, in contrast to conventional bainite containing carbides. Bainitic ferrite also differs from (primary) ferrite with a low dislocation density. The carbide-free BF microstructure provides high strength, which is due to an intermediate hard bainitic ferrite structure having a high dislocation density and containing supersaturated carbon. 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 exist at the boundaries of the lath-like bainitic ferrite.

[0044] In the present invention, the bainitic ferrite is divided into two types: the bainitic ferrite formed at a high temperature of Bs to Bn (referred to as high bainitic ferrite (HBF)) and the bainitic ferrite formed at a low temperature of Bn to Ms (referred to as low bainitic ferrite (LBF)). When the cross-section of a steel strip subjected to etching with 3% nital is observed with a scanning electron microscope accompanied by EBSD analysis, the average aspect ratio of HBF (defined as the length of the short axis divided by the length of the long 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 the bainitic ferrite (HBF) formed in a higher temperature range above Bn are similar to those of AF, and it is difficult to distinguish HBF from AF using SEM. Similar to 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 the formability.

[0045] The feature of the high-strength steel strip according to the present invention is that the bainitic ferrite can have a composite microstructure including HBF and LBF. Therefore, a high-strength cold-rolled steel strip having high elongation can be obtained. In order to obtain a good balance between high strength and elongation, 20 to 65% of LBF is required. If LBF is present in an amount less than this, the strength of the steel strip will be insufficient. However, if LBF is present in an excessively large amount, the elongation of other ferrites (PF, AF, and HBF) and retained austenite may be impaired. Therefore, LBF is 20 to 65%, preferably 30 to 60%.

[0046] The formation of HBF in the present invention is attributed to the heating of the strip by the latent heat generated by the bainite transformation, or to the heating by applying the hot dip galvanizing process. The formation of HBF in the present invention, if any, enables the acceleration of the bainite transformation rate as required, and as a result, the bainite transformation can be completed in a limited time in the over-aged section of the existing production line. Depending on the amounts 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 described above, HBF has the same function as PF and AF. When a sufficient amount of PF and AF is formed in the soaking and cooling sections, and in order to obtain a higher-strength steel strip, it is necessary to minimize the amount of HBF to 0%. When the amounts of PF and AF are insufficient, the amount of HBF can 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%.

[0047] Polygonal ferrite and acicular ferrite The proeutectoid ferrite is softer than the bainitic ferrite and functionally increases the elongation of the steel strip. On the other hand, introducing a certain amount of proeutectoid ferrite and controlling the properties of the ferrite can increase the bainite transformation rate, enhance the stability of the retained austenite, and further increase the elongation. Using the present invention, two types of proeutectoid ferrite can be produced during annealing according to the formation temperature. The ferrite phase formed during austenitization at the transformation interval temperature or the ferrite phase formed during cooling at a high temperature exceeding the Bs temperature in the slow cooling section becomes polygonal or blocky and is called polygonal ferrite (PF). This type of ferrite has been proven to increase elongation but decrease the yield strength and formability. The ferrite formed at a low temperature in the high-speed cooling section at the temperature of Bs - Ms is almost needle-shaped, has a smaller grain size than PF, and is called acicular ferrite (AF). Morphologically, it is similar to HBF, but the amount of dislocations is relatively small. The presence of AF can increase elongation without sacrificing strength and formability.

[0048] PF, AF, and HBF have similar functions with respect to the tensile properties in the steel according to the present invention. Therefore, there can be three types of these ferrite-based microstructures, or one or two of them may exist. 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 needs to be controlled to 55% or less, preferably 50% or less. When these ferrite-based microstructures are excessively numerous (more than 55%), the final microstructure does not sufficiently contain lower bainitic ferrite, and thus the strength decreases.

[0049] When PF exists in the steel, it is necessary to control the particle size, morphology, and distribution of PF. The steel strip can have a higher elongation by having PF with a smaller particle size and a dispersed distribution. According to the present invention, when observed by SEM or optical microscope, the PF structure is equiaxially embedded between the BF structures and uniformly dispersed as smaller grains, while the morphological structure of PF in conventional TRIP steel strips extends along the rolling direction. This morphological structure is considered to be able to uniformly disperse the stress during processing and maximize the TRIP effect of retained austenite. In order to obtain this morphological structure, the amount of PF formed during soaking needs to be 45% or less, preferably 10 - 40%. This embodiment is particularly suitable for steel compositions containing a relatively large amount of Mn, Al, and Si with a slow recrystallization rate of ferrite. PF is a partially recrystallized microstructure and has a higher hardness than recrystallized ferrite. The presence of PF having a partially recrystallized microstructure is beneficial for local ductility. Advantageously, the particle size of PF in the present invention is 10 μm or less, preferably 8 μm or less, more preferably 5 μm or less.

[0050] 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 such that AF + HBF is 20 - 55%, preferably 25 - 50%.

[0051] Retained Austenite Retained austenite (also known as residual austenite) refers to regions in the final microstructure that exhibit an FCC phase (face-centered cubic lattice). Retained austenite partially enhances ductility through the TRIP effect, which manifests as an increase in uniform elongation. To exhibit the TRIP effect, the volume fraction of retained austenite is 5% or more, preferably 7% or more. If it is less than 5%, the desired levels of ductility and uniform elongation cannot be achieved. The upper limit is mainly determined by the composition and processing parameters of the production line. For a given composition, if the amount of retained austenite is excessively high, the carbon content of the retained austenite will be excessively low. In that case, the retained austenite is not sufficiently stable, and the local ductility (elongation flange formability) may decrease to an unacceptable level. Therefore, the upper limit of the volume fraction of retained austenite is 20%, preferably 15%.

[0052] The carbon concentration in retained austenite affects the TRIP characteristics. Retained austenite is effective in improving elongation characteristics, especially when the carbon concentration in the retained austenite is 0.90% or more. If the carbon content is excessively 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 be as high as possible, and an upper limit of about 1.6% is generally imposed by the actual processing conditions. By controlling the amount of ferrite, the carbon content and stability of the retained austenite can be adjusted.

[0053] Martensite Martensite is freshly formed in the final cooling section after austempering. Martensite suppresses the elongation at yield and increases the work hardening index (n-value), which is desirable to achieve stability, neck-free deformation, and strain uniformity in the final pressed parts. Even 1% fresh martensite in the final steel strip can achieve a tensile response and thus a press behavior equivalent to that of 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 20%, preferably below 15%.

[0054] carbide Carbides can exist as fine precipitates formed during austempering when the overaging temperature is excessively high or the overaging time is excessively long, or in the form of pearlite formed during cooling when the cooling rate is excessively slow. According to the present invention, the microstructure of the steel of the present invention does not contain pearlite (pearlite-free) and does not contain carbides. Not containing pearlite means that the amount of the lamellar microstructure containing cementite and ferrite is less than 5%. Not containing carbides means that the amount of carbides is below the detection limit of standard X-ray measurement.

[0055] Characteristic evaluation of microstructure The microstructure components classified as the steel according to the present invention as described above can be quantitatively determined by the techniques described below. The volume fraction of the components is measured by regarding the volume fraction as equivalent to the area fraction and measuring the area fraction of the polished surface using a commercially available image processing program or other appropriate techniques.

[0056] PF, fresh M, RA, and pearlite can be distinguished using an optical microscope (OM) and / or a scanning electron microscope (SEM). When a sample etched with a 10% aqueous sodium metabisulfite solution (abbreviated as SMB) is characterized by OM, pearlite is observed as a dark area, PF as a light gray area, and fresh martensite as a light brown area. When a sample etched with a 3% nital solution is characterized by SEM, PF is observed as smooth-surfaced grains that do not contain retained austenite, and pearlite is observed as a lamellar microstructure that contains both cementite and ferrite. The remaining microstructure is observed as a gray area characterized by plate- or lath-like ferrite substructures, and RA is dispersed as white or light gray areas within the grains, and no carbides that can be identified are present. This microstructural group is called a bainitic ferrite like microstructure. This may include a mixture of other microstructures called HBF, LBF, AF, and partitioned martensite (PM). PM is produced when the austempering temperature is below the Ms point of the steel. During rapid cooling, some martensite is formed, and then carbon partitioning occurs between the martensite and the retained austenite during austempering. PM exists in a similar form to LBF and may coexist with LBF. AF, HBF, LBF, and PM in the bainitic ferrite like microstructure have similar morphologies and thus cannot be clearly distinguished by OM and SEM.

[0057] In the present invention, electron backscatter diffraction (EBSD) is used to further separate the bainitic ferrite-like microstructures into two distinct groups. The first group consists of PM and LBF, and the second group consists of AF and HBF. From the measured EBSD data, retained austenite can be initially distinguished from the other microstructures by creating a partition of Fe(α) to Fe(γ). Next, fresh martensite (M) is separated from the bainitic ferrite-like microstructures by dividing Fe(α) into a partition with a high average image quality (IQ: image quality) and a partition with a low average IQ. The low IQ partition is classified as martensite, and the high IQ partition is classified as bainitic ferrite-like microstructures. A method for distinguishing between the two types of groups will be described below with reference to FIG. 1. In bainitic ferrite (high IQ partition), regions with an orientation difference where the tilt angle between adjacent structures is 15° or more are identified. A region is regarded as a region having the same crystal orientation and is defined as a bainite plate in the present invention. For the bainite plates detected in this way, the diameter of a circle having the same area as the bainite plate is determined. Using an image of an EBSD analysis at a magnification of 3000, the diameter of the equivalent circle of the bainite plate is determined. 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, the diameter of the equivalent circle of all bainite plates within the measurement region (about 100×100 μm) and the aspect ratio of the equivalent ellipse of all bainite plates are measured, and the average values are defined as the average grain size of the bainite plates and the average aspect ratio of the bainite plates of the present invention.

[0058] The inventors systematically studied the influence of the austempering temperature on the microstructure of bainitic ferrite. The austempering temperature is (Ms - 200) to Bs. It was found that the average grain size and the average aspect ratio of the bainite plates increase as the austempering temperature rises. In particular, a sharp change is seen between samples austempered at less than 440°C, where the aspect ratio of the bainite plates is less than Bn for the steel having the composition used in the method according to the present invention, and samples austempered at more than 460°C, where it is more than 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 0.35 or less, and the aspect ratio of the group consisting of HBF and AF is more than 0.35.

[0059] In addition to the differences in the morphology and grain size of the bainite plates, the complex crystallographic plate orientation relationships between the HBF, AF groups and the LBF, PM groups also differ. The orientation difference angle distribution of the steel according to the present invention is shown in FIG. 2. The peak at 60° is the orientation difference between adjacent particles and coincides with the orientation difference having the Kurdjumov-Sachs (KS / KS) relationship, which is generated by the axis and angle relationships of 60° <111> and 60° <110> and corresponds to martensite. The peak at 53° - 54° is due to the orientation difference between particles obtained by a phase transformation according to the Nishiyama-Wassermann and Kurdjumov-Sachs (NW / KS) relationships. According to the prior art (see A.-F. Gourgues, H. M. Flower, and T. C. Lindley, Materials Science and Technology, January 2000, Vol. 16, p. 26-40), acicular ferrite and upper bainite grow in the Nishiyama-Wassermann relationship with the austenite parent phase, while lower bainite and martensite are composed of very complex packets having the Kurdjumov-Sachs relationship with the parent phase. Similarly to these results, it is estimated 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 peak at 53 - 54° becomes more distinguishable and the height of the peak increases, while the height of the peak at 60° 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.

[0060] A part of the retained austenite is dispersed as very fine-grained films between the bainite plates and cannot be detected by EBSD. Therefore, the ratio of the retained austenite measured by EBSD is always lower than the actual value. Thus, an intensity measurement method based on X-ray diffraction (XRD), which is a conventional method for measuring the content of retained austenite, can be used. The volume fraction of the retained austenite is determined at a thickness of 1 / 4 of the steel strip. The amount of cementite is also measured from this XRD analysis. Samples made from the steel strip are mechanically and chemically polished and then analyzed by measuring the respective integrated intensities of the (200), (220), and (311) planes of fcc iron and the (200), (211), and (220) planes of bcc iron using an X-ray diffractometer with Co-Kα. The amount of retained austenite (RA) and the lattice constant of the retained austenite are determined using Rietveld analysis. The C content in the retained austenite is calculated using the following formula: [Equation] (where a is the lattice constant (angstrom) of the retained austenite.) is calculated using.

[0061] Mechanical properties The cold-rolled steel strip having the above microstructure and composition and heat-treated according to the present 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 an elongation at total elongation (TE) of at least 14%; and / or has.

[0062] Preferably, the heat-treated cold-rolled strip has all of these properties.

[0063] Steps of the method According to the method of the present invention, the cold-rolled steel strip having the above composition is heat-treated to obtain a microstructure and properties. The cold-rolled steel strip obtained by cold rolling is subjected to heat treatment in the same manner as in a continuous annealing line. A typical design of the method is schematically shown in FIG. 3. The cold-rolled steel strip is heated at a temperature of (Ac3 - 60) or higher, for example, using a heating rate of at least 0.5 °C / second, preferably up to a temperature of (Ac3 - 60) to (Ac3 + 20), usually up to a predetermined austenitization temperature T2, held at this temperature for a time t2 (step a), then cooled to a temperature T4 of Ms to Bn, usually using two-stage cooling at a controlled cooling rate (step b). Then, the steel strip is heated and austempered at a temperature T5 of T4 to Bs, preferably T4 to Bn, for a time t5 (step c). Then, optionally, the steel strip is heated for a time t6 to a temperature T6 of Bn to Bs. This may be a hot-dip galvanizing treatment. Finally, the steel strip is cooled to room temperature (step d). Hereinafter, the process parameters and actions in each step will be described.

[0064] In the first step, in order to achieve a microstructure that is at least partially austenitized, the cold-rolled steel is soaked at a temperature of (Ac3 - 60) or higher, for example, at a temperature of (Ac3 - 60) to (Ac3 + 20), for a soaking time t2 of 1 to 150 seconds. Annealing at a temperature of (Ac3 - 60) or higher is necessary because the steel strip heat-treated according to the present invention requires a required amount of low-temperature transformation phases (for example, bainitic ferrite and retained austenite) that transform from high-temperature austenite and a predetermined amount of ferrite. When T2 is higher than (Ac3 + 20), the austenite grains grow. This affects the particle size and distribution of the retained austenite and also slows down the bainite transformation rate after the overaging process. An excessive amount of fresh martensite formed during the final cooling may be formed as a result of this incomplete bainite transformation, which increases the strength but decreases the ductility and formability. Furthermore, a uniform austenite structure with large grain size suppresses the formation of PF and AF in the next cooling section, and as a result, a sufficient amount of ferrite cannot be obtained within the current cooling schedule in the available production line. A uniform austenite structure with large grain size also causes insufficient elongation of the steel strip. It has been observed that the uniformity of austenite has a great influence on the formation of PF and AF in the cooling section. When T2 is lower than (Ac3 - 60), PF is It can be formed in an excess of more than 45%, and thus the strength of the steel strip may be insufficient. On the other hand, the formed austenite may not be sufficient for the formation of LBF and retained austenite. Therefore, the annealing temperature needs to be (Ac3 - 60) or higher, preferably (Ac3 + 20) or lower, and preferably (Ac3 - 50) to (Ac3 + 10). When t2 is longer than 150 seconds, the grain sizes of austenite and ferrite increase and the elongation decreases. When 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 dissolve sufficiently. Therefore, the annealing time t2 is 1 second to 150 seconds, for example, 10 seconds to 120 seconds, and preferably 1 second to 100 seconds.

[0065] In the next cooling step, the strip that has been at least partially austenitized is cooled to a temperature T4 between Bn and Ms. The purpose of this cooling is to adjust the amounts of ferrite and bainitic ferrite and to prevent the formation of pearlite. Usually, this means that the cooling rate during cooling is high enough to avoid the formation of pearlite.

[0066] In one embodiment of the present invention, the thus-treated steel strip is directly cooled to the temperature T4 at a cooling rate V4 of at least 15 °C / second, thereby preventing the formation of pearlite. If the cooling rate is excessively slow, an excessive amount of ferrite may be produced, or even pearlite may be produced. Preferably, V4 is 20 °C / second or more. The upper limit of the cooling rate is not particularly limited as long as there is no temperature change in the steel strip at the end of cooling. In order to reduce the temperature non-uniformity of the steel strip, the cooling rate is preferably 100 °C / second or less in standard equipment. If the average cooling rate exceeds 80 °C / second in most available facilities, very large fluctuations in the microstructure can occur in both the longitudinal direction and the strip width direction. Therefore, an appropriate cooling rate V4 is 15 to 80 °C / second, preferably 20 to 60 °C / second.

[0067] In other embodiments of the present invention, to adjust the amount of ferrite and to equalize the strip temperature, this cooling can be carried out by two-stage cooling. This is suitable for most continuous annealing lines or hot-dip galvanizing lines that currently use two connected cooling sections. First, the steel strip is cooled to a temperature T3 of 800 to 500 °C, preferably 750 to 550 °C, at a cooling rate V3 of usually at least 1 °C / second, for example, 2 to 15 °C / second, preferably 3 to 10 °C / second (referred to as the slow cooling section). Then, the steel strip is further cooled to temperature T4 at a cooling rate V4 of usually at least 15 °C / second, for example, 15 to 80 °C / second, preferably 20 to 60 °C / second (referred to as the fast cooling section). Since the length of each section in the continuous annealing line is fixed, by adjusting the T3 temperature, the cooling rates V3 and V4 for a given line speed can be controlled. As T3 increases, V3 decreases and V4 increases. During this cooling, some PF can be formed in the slow cooling section and some AF can be formed in the fast cooling section. In the case of 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, to adjust the amount of ferrite and prevent the formation of pearlite, T3 is selected within an appropriate range. If T3 is excessively low (e.g., less than 500 °C), an excessive amount of PF can be formed in the slow cooling section, and an excessive amount of AF can also be formed in the fast cooling section. Further, if the resulting V4 is less than 15 °C / second, pearlite can be formed. If T3 is excessively high (e.g., over 800 °C), insufficient PF can be formed, and if the resulting V4 is excessively high, the amount of AF formed will be less. Therefore, T3 needs to be 800 to 500 °C, preferably 750 to 550 °C / second.

[0068] As described above, in the annealing line or galvanizing line of the conventional design, PF is obtained in the soaking process a) and the slow cooling section of process b), and AF is obtained in the fast cooling section of process 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. When using the higher T2, a smaller amount of PF is generated during soaking, and then by selecting the lower T3, a large amount of PF can be obtained in the slow cooling section and a large amount of AF can be obtained in the fast cooling section. When using the lower T2, a sufficient amount of PF is generated during soaking, and then by selecting the higher T3, the amount of PF formed in the slow cooling section and the amount of AF formed in the fast cooling section are restricted.

[0069] The cooling stop temperature T4 at which bainite transformation starts is Bn~Ms. If T4 is excessively high, excessive HBF may be obtained during the subsequent austempering process, and the strength of the steel strip may become insufficient. If T4 is excessively low and there is no active heating, the latent heat generated by bainite transformation is not sufficient to heat the steel strip to a temperature high enough for the high-speed bainite transformation rate. Preferably, T4 is Bn~(Ms + 50°C).

[0070] In the subsequent heating step c), the cooled strip is heated at a temperature of T4 to Bs, preferably T4 to Bn, for a time t5 of 30 to 300 seconds, typically to a temperature T5 within this range, and is heated by heating. During this time, the austenite transforms into low bainitic ferrite (LBF). If T5 is excessively low, the bainite transformation is excessively slow, the bainite transformation is insufficient during overaging, and an excessive amount of fresh martensite may be formed during cooling after overaging, which increases the strength but the required elongation cannot be obtained. On the other hand, carbon partitioning may be insufficient to stabilize the retained austenite. If T5 is excessively high, excessive HBF is obtained in the overaging section, and there is a risk that the required strength cannot be obtained. The most preferred range for step c) is (Bn - 50) to Bn in order to achieve a high bainite transformation rate and obtain LBF. If the heating time t5 is less than 30 seconds, the bainite transformation is incomplete and sufficient LBF is not formed. Also in this case, carbon partitioning is insufficient. If t5 exceeds 300 seconds, carbides begin to form and there is a risk that the carbon content in the retained austenite decreases. The maximum time of t5 is restricted, inter alia, by the total time available at a given speed of the production line. Preferably, t5 is 40 to 120 seconds.

[0071] The strip thus heated 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 at a cooling rate V7 of at least 1 °C / second, preferably at least 5 °C / second, and then further cooled to ambient temperature. Cooling to ambient temperature can be forced cooling or uncontrolled natural cooling. In an actual embodiment, the heated steel strip is cooled at a cooling rate V7 of 5.0 to 10.0 °C / second to a temperature T7 of (Ms - 50) to Mf. The further cooling from T7 to ambient temperature is preferably carried out at a cooling rate V8 of 5.0 to 20.0 °C / second, more preferably 6.0 to 15.0 °C / second.

[0072] In one embodiment of the present invention, the heating step c) is preferably carried out at least partially by the latent heat from the bainite transformation. During this heating step c), the lower bainite transformation generates latent heat. Since the latent heat can be incompletely dissipated in the production line, the temperature of the steel strip automatically rises due to the accumulated latent heat. Therefore, the temperature of the steel strip to be heated can gradually rise due to the latent heat of the occurring bainite transformation. That is, the temperature T5 rises during the heating process, especially at a longer time t5. The upper limit temperature of T5 is not particularly limited. When the steel strip reaches a temperature higher than Bn, high-temperature bainitic ferrite is formed.

[0073] In one embodiment, to ensure obtaining a partially austenitized cold-rolled strip having a fine particle size, the soaking step is carried out at a transformation interval annealing temperature of (Ac3 - 50) to (Ac3 + 10), preferably with a soaking time t2 of 1 to 100 seconds. The fraction of PF formed at the soaking temperature is preferably 40% or less.

[0074] 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 of 680 to 740 °C, preferably 700 to 720 °C, at a heating rate V1 of 10.0 to 30.0 °C / second, preferably 15.0 to 25.0 °C / second, and further heating from the temperature T1 to the range of the soaking temperature at a heating rate V2 of 0.5 to 4.0 °C / second, preferably 1.0 to 3.0 °C / second. During the slow heating from T1 to the soaking temperature T2, recovery and recrystallization occur in ferrite, and solid solution of carbide and ferrite occurs during austenite transformation. T1 and V2 affect the progress of these processes and affect the grain size of austenite and the uniformity of the distribution of alloying elements in the austenite phase. Advantageously, to ensure complete solid solution of all carbides and Coarse To avoid non-uniform austenite crystal grain size, the soaking time t2 is controlled according to the heating rate V2.

[0075] In one embodiment, the method according to the invention includes an additional heating step between the heating step c) and the cooling step d), and the steel strip obtained from step c) is subjected to additional heating at Bs to Bn, preferably (Bs - 50) to Bn, typically at a fixed temperature T6. The additional heating time t6 is advantageously 5 to 30 seconds, preferably 10 to 20 seconds. This additional heating increases the formation from retained austenite to high-temperature bainitic ferrite and completes the bainite transformation. As a result, the amount of martensite formed in the subsequent cooling section is reduced, and the strength and ductility can be improved. Carbon is also further distributed into the retained austenite and becomes more stable. When this additional heating is applied to a given overaging section, and thus applied within a given total time therein, the time t5 is further shortened to meet the available time. For example, the total of t5 + t6 is 30 to 300 seconds.

[0076] In a preferred embodiment, this additional heating includes an integrated hot-dip galvanizing process, and the steel strip obtained from step c) is coated with a Zn coating or a Zn alloy-based coating.

[0077] The steel strip heat-treated according to the present invention may be provided with a coating, preferably 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 include a Zn alloy containing Al as an alloying element. A preferred composition of the zinc bath contains 0.10 to 0.35% of Al, with the balance being zinc and inevitable impurities. Another preferred zinc bath containing Mg and Al as main alloying elements has the following composition: 0.5 to 3.8% of Al, 0.5 to 3.0% of Mg, including one or more additional elements up to 0.2% as optional components, with the balance being zinc and inevitable impurities. Examples of the additional elements include Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr, and Bi.

[0078] A coating, for example, a protective coating of Zn or a Zn alloy, can be applied in a separate process. Preferably, the hot dip galvanizing process is integrated in the method according to the invention, as described above.

[0079] Optionally, temper rolling is carried out on the annealed and zinc-coated strip according to the invention in order to finely adjust the tensile properties and to modify the surface appearance and roughness according to specific requirements resulting from the intended use.

[0080] The cold-rolled steel strip itself is usually produced according to the following general method. Prepare the above steel composition and cast it into a slab. After reheating the cast slab at a temperature of 1100 to 1300 °C, it is processed by hot rolling. Usually, hot rolling of the slab is carried out within 5 to 7 stands to the final dimensions suitable for further cold rolling. Usually, finish rolling is carried out in a fully austenitic state above 800 °C, preferably in a fully austenitic state above 850 °C. The strip obtained from the hot rolling process in this way 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 an appropriate gauge. Preferably, the reduction in cold rolling is typically 30 to 80%. In order to reduce the rolling load during cold rolling, the coiled strip or semi-cold-rolled strip may be subjected to hot batch annealing. The batch annealing temperature needs to be 500 to 700 °C.

[0081] Thin slab casting, strip casting, etc. can also be applied. In this case, it is allowed that the manufacturing method skips at least a part of the hot rolling process.

[0082] The present invention also relates to a heat-treated cold-rolled steel strip having the composition and microstructure as described above.

[0083] The present invention is also present in articles manufactured from the heat-treated cold-rolled strip according to the present invention, such as structural parts, engineering parts or automotive parts.

Examples

[0084] Steel having the composition shown in Table 1 was cast into a 25 kg ingot with dimensions of 200 mm × 110 mm × 110 mm using vacuum induction. A cold-rolled strip with a thickness of 1 mm was manufactured using the following process schedule. · The ingot was reheated at 1225 °C for 2 hours. · The ingot was rough-rolled from 140 mm to 35 mm. · The rough-rolled ingot was reheated at 1200 °C for 30 minutes. · Hot-rolled from 35 mm to 4 mm in 6 passes. · Run-out table cooling: Cooled from the finish rolling temperature (FRT) (about 850 - 900 °C) to 600 °C at a rate of 40 °C / second. · Furnace cooling: The strip was 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 pickled with HCl at 85 °C to remove the oxide layer. · Cold rolling: The hot-rolled strip was cold-rolled into a 1 mm strip. · Heat treatment according to the present invention: Using a cold-rolled sheet of appropriate size, the annealing process was simulated using a continuous annealing simulator (CASIM).

[0085] Samples for microstructure observation, tensile testing and hole expansion testing were machined from the sheet thus treated.

[0086] The expansion rate measurement was performed on cold-rolled samples with dimensions of 10 mm × 5 mm × 1 mm (height along the rolling direction). The expansion test was carried out using a Bahr dilatometer, model DIL805. All measurements were performed according to SEP1680. The critical phase transformation points Ac3, Ms, and Mf were determined from the quenched dilatometry curve. Bs and Bn were predicted using the available software JmatPro10. The phase fraction during annealing for different process parameters was determined from the expansion curves simulating the annealing cycle.

[0087] 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 cross-sections in the rolling direction and the rolling plane normal direction of the steel strip. The scanning electron microscope (SEM) used for EBSD measurement 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. The EBSD scans were evaluated using 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 approximately 80 frames / second.

[0088] The retained austenite was determined by XRD according to DIN EN13925 using D8 Discover GADDS (Bruker AXS) with Co-Kα radiation. Quantitative determination of the phase ratio was performed by Rietveld analysis.

[0089] Tensile test: JIS 5 test pieces (gauge length = 50 mm, width = 25 mm) were machined from annealed strips such that the tensile direction was parallel to the rolling direction. A room temperature tensile test was carried out using a Schenk TREBEL testing machine conforming to the NEN-EN10002-1:2001 standard, and the tensile properties (yield strength YS (MPa), ultimate tensile strength UTS (MPa), total elongation TE (%)) were determined. For each condition, three tensile tests were carried out and the average value of the mechanical properties was reported.

[0090] Using the display in Fig. 3, the process parameters are shown in Table 2. In CASIM, the latent heat generated during bainite transformation is offset by an active cooling system. The temperature change due to the latent heat in the overaging section is simulated by temperatures T4, T5, and T6. When T4 is equal to T5, t4 is 0. When T4 is lower than T5, t4 is 1 second, and the subsequent heating to T5 is at a heating rate of 5 - 20 °C / second. The heating rate from T5 to T6 is 5 - 10 °C / second. The obtained microstructures and tensile properties are shown in Table 3. All steel compositions except A79 achieve the requirements for microstructures and tensile properties under the various specified processing parameters with respect to the present invention. Steel A79 could not achieve the required elongation despite having a sufficiently high tensile strength (Examples 26 and 27). The reason for this is that A79 contains 0.5% Cr, which significantly slows down the bainite transformation rate, and as a result, the bainite transformation and the C partitioning between BF and retained austenite could not proceed sufficiently. As a result, the amount of martensite increases, but the amount of retained austenite and the C content in the retained austenite decrease, and the TRIP effect deteriorates.

[0091]

Table 1

[0092]

Table 2

[0093]

Table 3

[0094]

Table 4

[0095]

Table 5

Claims

1. A method for heat-treating a cold-rolled steel strip, comprising: the method comprising the following steps: a) soaking the cold-rolled steel strip at a temperature of (Ac3 - 60) or higher for a soaking time t2 of 1 to 150 seconds, thereby obtaining a cold-rolled steel strip having a microstructure that is at least partially austenitized; b) cooling the soaked steel strip obtained in step a) to a temperature T4 between Bn and Ms, wherein the soaked steel strip obtained in step a) is directly cooled to the temperature T4 at a cooling rate V4 of at least 15°C / second, or the soaked steel strip obtained in step a) is cooled to a temperature T3 of 800 - 500°C at a cooling rate V3 of at least 1°C / second and then cooled from the temperature T3 to the temperature T4 at a cooling rate V4 of at least 15°C / second; c) heating the cooled steel strip obtained in step b) at a temperature between Bs and T4 for a time t5 of 30 to 300 seconds; d) cooling the heated steel strip to ambient temperature whereby the steel strip, in volume %, polygonal ferrite (PF) + acicular ferrite (AF) + high bainitic ferrite (HBF): 20 - 55 polygonal ferrite (PF): 0 - 45 low bainitic ferrite (LBF): 20 - 65 retained austenite (RA): 5 - 20 martensite (M): 0 - 20 and has a microstructure that does not contain pearlite, where "does not contain pearlite" means that the amount of the lamellar microstructure containing cementite and ferrite is less than 5%, the steel strip having, in mass %, the following composition: 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 optionally, 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 one or more elements selected from iron and unavoidable impurities: the balance and the total of (Si + Al) is ≥ 0.60, where REM is one or more rare earth metals, the method.

2. The method according to claim 1, wherein step a) comprises soaking the cold-rolled steel strip at a temperature of (Ac3 - 60) to (Ac3 + 20).

3. The method according to claim 1, wherein step a) comprises soaking the cold-rolled steel strip at a temperature of (Ac3 - 50) to (Ac3 + 10).

4. The method according to claim 2 or 3, wherein in step a), the soaking of the cold-rolled steel strip is carried out during a soaking time t2 of 1 to 100 seconds.

5. The method according to any one of claims 1 to 4, wherein the temperature T3 is 750 to 550 °C.

6. The method according to any one of claims 1 to 5, wherein the cooling rate V3 is 2.0 to 15.0 °C / second.

7. The method according to any one of claims 1 to 6, wherein the cooling rate V3 is 3.0 to 10.0 °C / second.

8. The method according to any one of claims 1 to 7, wherein the cooling rate V4 is 20.0 to 60.0 °C / second.

9. The method according to any one of claims 1 to 8, further comprising, prior to step a), heating the cold-rolled steel strip to a temperature of (Ac3 - 60) or higher at a heating rate of at least 0.5 °C / second.

10. The method according to any one of claims 1 to 8, further comprising, prior to step a), heating the cold-rolled steel strip to a temperature T1 of 680 to 740 °C at a heating rate V1 of 10.0 to 30.0 °C / second, and further heating from temperature T1 to a temperature of (Ac3 - 60) or higher at a heating rate V2 of 0.5 to 4.0 °C / second.

11. The method according to claim 10, wherein the temperature T1 is a temperature of 700 to 720 °C.

12. The method according to claim 10 or 11, wherein the heating rate V1 is 15.0 to 25.0 °C / second.

13. The method according to any one of claims 10 to 12, wherein the temperature of (Ac3 - 60) or higher is a temperature of (Ac3 - 60) to (Ac3 + 20).

14. The method according to any one of claims 10 to 13, wherein the temperature of (Ac3 - 60) or higher is a temperature of (Ac3 - 50) to (Ac3 + 10).

15. The method according to any one of claims 10 to 14, wherein the heating rate V2 is 1.0 to 3.0 °C / second.

16. The method according to any one of claims 1 to 15, wherein step c) is at least partially carried out by the latent heat generated by bainite transformation.

17. The method according to any one of claims 1 to 16, wherein the heating in step c) is carried out at a temperature of Bn to (Ms + 50).

18. The method according to claim 17, wherein the heating in step c) is carried out for a time t5 of 40 to 120 seconds.

19. The heating in step c) is carried out at a temperature of Bn to T4, The method according to any one of claims 1 to 18, wherein the method includes an additional heating step of heating the steel strip obtained from step c) at a temperature of Bs to Bn between step c) and step d).

20. The heating in step c) is carried out at a temperature of Bn to T4, The method according to any one of claims 1 to 18, wherein the method includes an additional heating step of heating the steel strip obtained from step c) at a temperature of (Bs - 50) to Bn between step c) and step d).

21. The method according to claim 19 or 20, wherein the heating in the additional heating step is carried out for a time t6 of 5 to 30 seconds.

22. The method according to claim 21, wherein the time t6 is 10 to 20 seconds.

23. The method according to any one of claims 19 to 22, wherein the additional heating step includes a hot dip galvanizing treatment.

24. The method according to any one of claims 1 to 22, further including a coating step of coating the heat-treated steel strip with a protective coating after the heat treatment.

25. The method according to claim 24, wherein the protective coating is a Zn coating or a Zn alloy coating.

26. The microstructure is in volume %, Polygonal ferrite (PF) + acicular ferrite (AF) + high bainitic ferrite (HBF): 25 - 50 Polygonal ferrite (PF): 10 - 40 Low bainitic ferrite (LBF): 30 - 60 Retained austenite (RA): 7 - 15 Martensite (M): 0 - 15 including, and / or The method according to any one of claims 1 to 25, wherein the C content in the retained austenite (RA) is 0.90 wt% or more.

27. The method according to claim 26, wherein the C content in the retained austenite (RA) is 0.95 wt% or more.

28. The obtained steel strip has the following properties: Yield strength (YS) of ≧ 500 MPa; Tensile strength (TS) of ≧ 850 MPa; Total elongation (TE) of ≧ 14% The method according to any one of claims 1 to 27, having at least one of the above properties.

29. The obtained steel strip has the following properties: Yield strength (YS) of ≧ 500 MPa; Tensile strength (TS) of ≧ 850 MPa; Total elongation (TE) of ≧ 14% The method according to any one of claims 1 to 27, having all of the characteristics.

30. A heat-treated cold-rolled steel strip, wherein the heat-treated cold-rolled steel strip has the following composition in mass%: C: 0.15 to 0.35 Mn: 1.50 to 3.00 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 Optionally, 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 One or more elements selected from Iron and unavoidable impurities: the balance having, the total of (Si + Al) being ≧ 0.60, REM being one or more rare earth metals, the heat-treated cold-rolled steel strip having, in volume%, Polygonal ferrite (PF) + acicular ferrite (AF) + high bainitic ferrite (HBF): 20 - 55 Polygonal ferrite (PF): 0 - 45 Low bainitic ferrite (LBF): 20 - 65 Retained austenite (RA): 5 - 20 Martensite (M): 0 - 20 including and having a microstructure that does not contain pearlite, where "does not contain pearlite" means that the amount of the lamellar microstructure containing cementite and ferrite is less than 5%, the heat-treated cold-rolled steel strip.

31. The following characteristics: Yield strength (YS) of ≧ 500 MPa; Tensile strength (TS) of ≧ 850 MPa; Total elongation (TE) of ≧ 14% The heat-treated cold-rolled steel strip according to claim 30, having at least one of the characteristics.

32. The following characteristics: Yield strength (YS) of ≧ 500 MPa; Tensile strength (TS) of ≧ 850 MPa; Total elongation (TE) of ≧ 14% The heat-treated cold-rolled steel strip according to claim 30, having all of the characteristics.

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