Patent Application
20240337000
The invention relates to a process for producing a cold-rolled flat steel product having a bainitic base microstructure, and to a correspondingly cold-rolled flat steel product having a bainitic base microstructure.
Processes for producing bainitic steel sheets are known from the prior art; cf. EP 2 707 514 B1, EP 3 024 951 B1.
Steels having a bainitic base microstructure with proportions of martensite and residual austenite are notable for a particularly good combination of strength and elongation at break, especially at relatively high carbon contents. At the same time, for high residual austenite contents, an average of about 1.5% by weight of silicon is usually included in the alloy, which increases the transition temperature to austenite and hence, in the case of a heat treatment, for example inline in a melt dip coating system, the temperatures for adjustment of the microstructure would have to be altered, i.e. the temperature would have to be increased compared to the conventional temperature, and hence higher costs would be incurred in the course of throughput. Furthermore, high silicon contents can distinctly worsen surface quality, coatability and weldability.
With regard to economic viability for production of flat steel products with bainitic base microstructure, especially with high tensile strengths and high elongation at break, there is a need for optimization.
It is thus an object of the invention to provide a process for producing a cold-rolled flat steel product with bainitic base microstructure, with which the disadvantages cited in the prior art can be overcome, and to specify a correspondingly produced cold-rolled flat steel product with bainitic base microstructure.
This object is achieved in a first aspect of the invention by a process having the features as described herein.
The invention provides a process for producing a cold-rolled flat steel product having a bainitic base microstructure, comprising the steps of:
This object is achieved in a second aspect of the invention by a cold-rolled flat steel product having the features as described herein.
The invention provides a cold-rolled flat steel product having a bainitic base microstructure consisting of, aside from Fe and unavoidable impurities from the production, in % by weight:
A bainitic base microstructure is thus understood to mean a microstructure having bainite in a proportion greater than the proportion of martensite and also greater than the proportion of residual austenite. In particular, the proportion of bainite is greater than the sum total of the proportions of martensite and residual austenite. The proportion of bainite in the microstructure is preferably greater than 50%.
All figures relating to contents of the alloy elements specified in the present description are based on weight, unless explicitly stated otherwise. All contents should therefore be regarded as figures in % by weight. The specified microstructure constituents are determined by evaluation of light microscopy or electron microscopy analyses and should therefore be regarded as the area proportions in area %, unless explicitly stated otherwise. An exception to this is formed by the microstructure constituent austenite or residual austenite, which is reported as a proportion by volume in % by volume, unless explicitly stated otherwise.
The invention provides a process for producing a cold-rolled flat steel product having a bainitic base microstructure, comprising the steps of:
This object is achieved in a second aspect of the invention by a cold-rolled flat steel product having the features described herein.
The invention provides a cold-rolled flat steel product having a bainitic base microstructure consisting of, aside from Fe and unavoidable impurities from the production, in % by weight:
A bainitic base microstructure is thus understood to mean a microstructure having bainite in a proportion greater than the proportion of martensite and also greater than the proportion of residual austenite. In particular, the proportion of bainite is greater than the sum total of the proportions of martensite and residual austenite. The proportion of bainite in the microstructure is preferably greater than 50%.
The cold-rolled flat steel product of the invention has a tensile strength of at least 1000 MPa, especially of at least 1100 MPa, preferably of at least 1200 MPa. Elongation at break A50 in the cold-rolled flat steel product of the invention is at least 9%, especially at least 11%, preferably at least 13%. Tensile strength and elongation at break A50 are determinable in a tensile test according to DIN EN ISO 6892-1.
After the casting of a melt with an alloy composition within the ranges specified to give a precursor, for example in a strand casting plant or combined casting and rolling plant, the precursor can be processed further directly, i.e. coming directly from the heat of casting, for example in the case of the combined casting and rolling plant, such that the precursor is kept at a temperature or, if required, reheated to a temperature, for example in an equalizing or reheating furnace in which maximum homogenization is assured and in which any precipitates formed are dissolved to the maximum possible degree. If the melt, for example, is cast in a strand casting plant to give a precursor, the cast and fully solidified strand is divided into multiple slabs of finite dimensions, and then the slabs are allowed to cool down to ambient temperature by natural cooling. The precursor or the slab is (then) reheated to a temperature, for example in a walking beam furnace or by other suitable means. Alternatively, the still-hot precursor or the still-hot slab can also be transferred without intermediate cooling into an equalizing or reheating furnace, for example.
The temperature in the reheating and/or in the holding of the precursor is at least 1100° C., especially at least 1140° C., preferably at least 1180° C., in order to ensure maximum dissolution of any precipitates present in the precursor. The reheating and/or holding temperature should not exceed 1350° C., in order to avoid partial melting and/or excessive scaling of the precursor. For environmental and economic reasons, the reheating and/or holding temperature is especially limited to a maximum of 1280° C.
Depending on the available plants and/or processing options, it is possible, and is thus specified as an option, that an intermediate hot rolling operation of the precursor can be implemented to give a flat intermediate product in one or more roll stands at a final rolling temperature between 950 and 1250° C.
A final rolling temperature for optional creation of the flat intermediate product of at least 950° C. is chosen in order to very reliably exploit a grain-refining effect of the recrystallization after the roll pass(es). For reasons of energy efficiency, in particular, a final rolling temperature for optional creation of the flat intermediate product of not more than 1250° C. is chosen.
The precursor or the optional flat intermediate product is hot-rolled in one or more roll stands with a final rolling temperature between 800 and 1000° C. to give a hot-rolled flat steel product.
A final rolling temperature for creating the hot-rolled flat steel product of at least 800° C. is chosen in order not to allow the forming resistance to rise too high. In order to avoid unwanted coarse-grained formation, the final rolling temperature for creation of the hot-rolled flat steel product is limited to a maximum of 1000° C. In particular, the final rolling temperature for creation of the hot-rolled flat steel product in a multi-stand hot-rolling/finishing train for assurance of a maximum austenite content is set to at least 850° C., and for assurance of recrystallization to preferably at least 880° C. In order to limit the amount of coolant required, in particular, final rolling temperatures up to a maximum of 950° C. are chosen in order to minimize recrystallization and grain growth between final rolling and winding, preferably up to a maximum of 930° C. The degree of hot rolling in the last pass or in the last hot rolling stand is preferably at least 10% in order to be able to establish a fine microstructure in the hot-rolled flat steel product.
The hot-rolled flat steel product is wound at a winding temperature between 400 and 650° C. The winding temperature must be at least 400° C. in order to prevent martensite formation. In order to limit the diffusion of alloy elements having oxygen affinity to the surface during the winding operation, the winding temperature is limited to a maximum of 650° C. The winding temperature may especially be at least 500° C. in order to create (considerable) ferrite contents in the microstructure, which enable good cold rollability, preferably with high degrees of cold rolling. In order to establish a particularly good surface, which, in conjunction with the low Si content within the aforementioned limits, can ensure a particularly broad joining spectrum (welding range), the winding temperature is especially chosen up to a maximum of 570° C.
The hot-rolled flat steel product (hot strip) may have a thickness between 1.5 and 10 mm.
The hot-rolled flat steel product may optionally be subjected to annealing at an annealing temperature between 500 and 900° C., especially up to a maximum of 800° C., preferably up to a maximum of 700° C. The optional annealing corresponds essentially to a standard process of annealing of hot-rolled flat steel products, and can especially lead to better cold rollability.
The hot-rolled and optionally annealed flat steel product is subjected to cold rolling, wherein the cold rolling is conducted in one or more roll stands with a total degree of cold rolling of at least 30%. A total degree of rolling of at least 30% is required in order to specifically provide nucleation sites in the microstructure of the subsequent heat treatment, at which austenite grains can advantageously evolve. The overall degree of cold rolling may especially be at least 38%, preferably at least 45%, in order to break relatively long lines of perlite within the microstructure, if present in the microstructure, as a result of which cementite/ferrite interfaces present in particular can be distributed further within the microstructure, at which particularly good nucleation of austenite is possible during the subsequent heat treatment. The overall degree of cold rolling may be not more than 80%, especially not more than 70%.
The cold-rolled flat steel product (cold strip) may have a thickness between 0.5 and 4 mm.
According to the invention, controlled heat treatment of the cold-rolled flat steel product comprising austenitization at a temperature T_A between 800 and 950° C., quenching to a temperature T_B between 300 and 580° C., is effected in such a way that a bainitic base microstructure is established in the cold-rolled flat steel product. At a temperature T_A above 800° C., the thermodynamic driving force for austenite formation from cementite and ferrite is already extremely large, which can contribute to a rapid and desired austenitization. In order to achieve an energy-efficient process, the temperature T_A should be set at not more than 950° C. In particular, the temperature T_A may be set at not more than 900° C. in order to prevent significant austenite grain coarsening. The temperature T_A may preferably be set to not more than 875° C., at which the carbon in the microstructure has preferably not yet been distributed 100% homogeneously, which, in the subsequent quenching, can more preferably lead to faster bainite nucleation and hence to a faster bainitic transformation rate. The austenitization is followed by quenching to a temperature T_B between 300 and 580° C. The temperature T_B of at least 300° C. should be established such that the carbon content in the residual austenite can be redistributed, and not more than 40% martensite is formed. The temperature T_B may especially be at least 340° C., preferably at least 380° C., in order to degrade inhomogeneities in the carbon distribution such that barely any carbon accumulates at the bainite/residual austenite interfaces. The temperature T_B is set to a maximum of 580° C. in order to reliably avoid ferrite/perlite formation. The temperature T_B may especially be not more than 550° C., preferably not more than 510° C., in order to assure a high strength of the bainite.
For example, a first quench of T_A can be effected at a cooling rate dT_AB of at least 10 K/s to a temperature below the martensite starting point, in order to enable particularly easy nucleation of the bainite and to create a correspondingly fine bainite. Subsequently, the temperature is raised again to a temperature of at least 380° C., especially at least 450° C., in order to assure a particularly rapid redistribution of the carbon in the austenite.
By virtue of the obligatory elements chromium and nitrogen in the aforementioned contents, it is possible to specifically support the establishment of the desired microstructure during the heat treatment, and, in conjunction with carbon in the aforementioned contents, the nitrogen leads to a distinct increase in the rate of bainitic transformation, one reason being the formation of very fine chromium nitride that acts as a nucleator. Nitrogen can also distinctly reduce carbon segregation at the grain boundary. Since carbon slows the nucleation of bainite, it is thus also assumed that there will be an elevated nucleation rate at the grain boundaries.
At temperature Ac1, the microstructure begins to transform to austenite and is especially completely in austenitic form when the temperature Ac3 is exceeded. Bs bainite starting point, Bf bainite finishing point, Ms martensite starting point, and Mf martensite finishing point indicate the temperatures at which transformation to bainite or martensite commences or is complete. Ac1, Ac3, Bs, Bf, Ms and Mf are indices that are dependent on the composition (alloy elements) of the steel material used and can be inferred from what are called TTA and TTT diagrams. The required cooling rates can also be inferred from the TTT diagrams depending on the desired microstructure.
The alloy elements of the melt, or of the flat steel product, are specified as follows:
Carbon (C) contributes to hardness and, depending on its content, can delay ferrite formation and bainite formation, stabilize the residual austenite and reduce the Ac3 temperature. A content of at least 0.10% is required in order to achieve sufficient hardenability/hardness and strength. Bainite formation is too slow above a content of 0.30%. For improved weldability and establishment of a good ratio of force absorption and maximum bending angle in the bending test, and for progression of rapid bainite formation, the content may especially be set to a maximum of 0.25%, preferably to a maximum of 0.22%. For achievement of a higher strength level, the content may especially be set to at least 0.15% and, for establishment of a very good combination of hardenability and strength, preferably to at least 0.18%.
Silicon (Si) contributes to a further increase in hardenability/hardness and in strength via solid solution strengthening. In addition, it is also possible to use ferro silico manganese as alloying agent, which has a beneficial effect on production costs. In addition, depending on the content, suppression of cementite and hence stabilization of residual austenite is also possible. The use of chromium and nitrogen as obligatory elements makes it possible to dispense with very high silicon contents of about 1.50%, as customary in generic steels. Over and above a content of at least 0.40%, a first hardening effect arises, wherein a content of at least 0.60% in particular is established for a significant rise in strength. Preference is given to establishing a content of at least 0.80% in order to virtually completely suppress cementite formation and also to avoid excessively high martensite formation. Up to a content of not more than 1.20%, it is possible to establish a good surface coatable without difficulty and if required with a coating, especially with a zinc-based coating. Especially in the case of contents up to a maximum of 1.10%, as well as an improved surface quality, it is also possible to ensure and/or improve weldability.
Manganese (Mn) contributes to hardness and can greatly delay ferrite formation depending on the content. In order to suppress ferrite in the course of heat treatment, a content of at least 1.00% is established. In order not to limit weldability, the content is set to a maximum of 2.00%. In order to avoid procutectoid ferrite formation and to stabilize the residual austenite, the content may in particular be set at least at 1.10%. In order to improve elongation at break, the content may especially be set at a maximum of 1.80%, preferably at a maximum of 1.60%, in order to assure rapid bainite formation.
Chromium (Cr) contributes to hardness and can slow diffusive phase transformations during quenching, in particular ferrite. Chromium has a much smaller influence on bainite formation at lower temperatures. It is thus optimally suitable in order firstly to ensure a low critical cooling rate, but at the same time not to hinder bainite formation too much at low temperatures. In order to achieve a critical cooling rate low enough for avoidance of unwanted ferrite formation, a content of at least 0.50% is established. Up to a content of not more than 1.50%, it is possible to produce a good surface coatable without difficulty and if required with a coating, especially with a zinc-based coating. For good surface quality and improved weldability, the content may especially be limited to a maximum of 1.20%, preferably to a maximum of 1.00%. In order to stabilize the residual austenite even in the case of very long hold times in the bainite region, the content may especially be set to at least 0.60%, preferably to at least 0.70%.
Nitrogen (N), as austenite former, slows the critical cooling rate since nitrogen can suppress diffusive ferrite formation. In the course of holding within the bainite region, it is then possible in turn for very fine clusters and/or precipitates, in particular chromium nitrides, to form, which accelerate bainite formation at low temperatures. In addition, nitrogen reduces carbon oversaturation at grain boundaries and hence reduces unwanted chromium carbide formation, which can become very coarse, would form particularly along grain boundaries and hence would distinctly worsen ductility. For a significant effect, a content of at least 0.0030% is established. For good and problem-free castability of the melt/the steel, the content is limited to a maximum of 0.040%. If a content of at least 0.0070% in particular is established, the residual austenite can be stabilized against cementite formation, and bainite formation can be accelerated at a content preferably of at least 0.0090%, more preferably of at least 0.011%. In order to achieve improved weldability, the content may in particular be adjusted to a maximum of 0.030% and preferably to a maximum of 0.025%, such that chromium nitrides produced, for example, can be formed very finely.
Phosphorus (P) counts as an impurity in the broadest sense, which is entrained into the steel by iron ore and cannot be eliminated entirely in the industrial scale steelworks process. The content should be set as low as possible, with the content limited to a maximum of 0.10%. Adverse effects on formability can be reliably ruled out when the content is limited in particular to a maximum of 0.050% by weight, and for additional reduction of segregation effects preferably to a maximum of 0.030% by weight.
Sulfur(S) likewise counts as an impurity in the broadest sense and can be set to a content of not more than 0.050% in order to avoid any significant tendency to segregation and an adverse effect on formability as a result of excessive formation of sulfides (FeS; MnS; (Mn, Fe) S). The content is therefore limited in particular to a maximum of 0.020% by weight, preferably to a maximum of 0.0080%. In general, calcium is included in the alloy for desulfurization and adjustment of the S content depending on the Ca content.
The flat steel product may optionally contain one or more alloying elements from the group of (Al, V, Ti, Nb, Ni, Mo, W, Ca).
Aluminum (Al) may be included in the alloy as an optional alloying element, especially as a deoxidant, with a content of not more than 0.050%, and it is possible to include a content especially of at least 0.0010% in the alloy for reliable binding of any oxygen (O) present. Above a content of 0.050%, however, there is an elevated risk that (relatively coarse) aluminum nitride will form, and hence that nitrogen will be bound in an unwanted manner and cold rollability will also be worsened as a result. In particular, the content is limited to a maximum of 0.015%, in order to be able to reliably rule out the formation of aluminum nitride.
Vanadium (V) may be included in the alloy as an optional alloying element for grain refining with a content up to a maximum of 0.20%, such that there is especially no adverse effect on elongation at break. In order to achieve a desired effect of grain refining, it is especially possible to include a content of at least 0.0010% in the alloy.
Titanium (Ti) may be included in the alloy as an optional alloy element as a microsegregation element with a content up to a maximum of 0.010%, such that it is especially possible to rule out unwanted binding with nitrogen, which would form very hard and coarse titanium nitrides that could lead to embrittlement. In order to precisely adjust the free nitrogen content, it is possible to include a content of at least 0.0010% in particular in the alloy.
Niobium (Nb) may be included in the alloy as an optional alloy element for grain refining with a content up to a maximum of 0.10% in order in particular to avoid binding of nitrogen to give niobium nitride. In order to achieve a desired effect of grain refining, it is especially possible to include a content of at least 0.0010% in the alloy.
Nickel (Ni) as an optional alloy element, just like chromium, can improve transformation to austenite, increase strength and improve process stability in the course of a prolonged hold time during bainite formation, and so a content up to a maximum of 0.40% can be included in the alloy in order in particular not to slow bainite formation. For example, nickel is included in the alloy in conjunction with copper, since, when copper is added, nickel essentially suppresses the adverse effect of copper on hot rollability. For this purpose, for example, a copper content between 0.3*nickel and 0.7*nickel can be included in the alloy in order to avoid the iron-copper eutectic and hence the formation of a liquid phase at the surface on hot rolling. In order to achieve a desired effect of the aforementioned improvement, it is especially possible to include a content of at least 0.010% in the alloy.
Copper (Cu) may be included in the alloy as an optional alloy element for increasing hardness and strength with a content of not more than 0.80%, in order in particular not to worsen suitability for welding and hot rollability owing to low-melting Cu phases at the surface. In order to assure the strength-increasing effect, but also in order to improve resistance to atmospheric corrosion in the case of uncoated flat steel products, a content of at least 0.010% in particular may be included in the alloy.
Molybdenum (Mo) as an optional alloy element can increase strength and hardness, especially also in order to improve process stability, since molybdenum distinctly slows ferrite formation and has barely any effect on bainite formation in the temperature range between 300 and 580° C., such that a content up to a maximum of 1.00% can be included in the alloy. In particular, a content of at least 0.0020% may be included in the alloy, at which molybdenum-carbon clusters can form dynamically at the grain boundaries up to the level of ultrafine molybdenum carbides, which distinctly slow the mobility of the grain boundary and hence diffusive phase transformations. Moreover, the grain boundary energy can be reduced, which can in turn reduce the nucleation rate of ferrite.
Tungsten (W) as an optional alloy element can act similarly to molybdenum, wherein a content up to a maximum of 1.00% can be included in the alloy. In order to be able to take any positive effect on hardness/hardenability, a content of at least 0.0010% in particular may be included in the alloy.
Calcium (Ca) as an optional alloy element in the melt can be included in the alloy as a desulfurizing agent and for controlled influencing of sulfide in contents up to a maximum of 0.0050%, which can lead to altered plasticity of the sulfides on hot rolling. Furthermore, the addition of Ca can also improve cold forming characteristics. The effect described may be effective over and above a content especially of at least 0.0001%, preferably of at least 0.0003%.
As well as iron, the flat steel product, as a result of the production, may contain one or more elements from the group of (O, H, As) as unavoidable impurities, which are not deliberately included as alloy elements.
Oxygen (O) is an unwanted impurity, but one which is generally unavoidable for technical reasons. The oxygen content is limited to a maximum of 0.0050%, especially to a maximum of 0.0020%.
Hydrogen (H), as the smallest atom at intermediate lattice sites in the steel, may be very mobile and lead to tears in the flat steel product. The possible hydrogen impurity is therefore reduced to a content of not more than 0.0010%, especially of not more than 0.0004%, preferably of not more than 0.0002%.
Arsenic (As) is an impurity which may be present in the flat steel product, where the content is limited to a maximum of 0.020%, in order to avoid adverse effects.
The alloy elements specified as optional may especially alternatively also be tolerated as impurities in contents below the minimum limits specified, without affecting, preferably not worsening, the properties of the flat steel product.
In one configuration of the process of the invention, in step h), the austenitization at T_A is conducted with a heating rate dTA of at least 1.0 K/s between 600 and 800° C. The controlled adjustment of the heating rate dTA in the range between 600 and 800° C. influences the formation of the austenite grain size inter alia, which is important since it affects the bainitic transformation rate, but also the final properties in the microstructure after heat treatment. The faster the heating within this range, the more austenite nuclei can form, which block one another, slow their growth, and hence lead to a fine austenite grain overall. Below 600° C., there are no relevant recrystallization processes, for example in continuous heat treatment plants; moreover, barely any cementite coarsening takes place, and so the heating rate is easily capable of following the technical circumstances below 600° C. In particular, the heating rate dTA is at least 2 K/s, preferably at least 2.5 K/s, such that a particularly fine austenite grain can form. For uniform heating, a heating rate dTA up to 50 K/s is helpful, but it is also possible in principle to choose higher heating rates dTA.
In one configuration of the process of the invention, in step h), after attainment of temperature T_A, the cold-rolled flat steel product is kept at temperature T_A for a hold time t_A between 1 and 300 s. In order to prevent grain coarsening, the hold time t_A of 300 s should not be exceeded. In particular, the hold time t_A may be chosen up to a maximum of 200 s in order, for example, to limit diffusion of unwanted trace elements such as phosphorus to the austenite grain boundaries.
In one configuration of the process of the invention, in step h), the quenching is effected in two stages, such that quenching is effected first to an intermediate temperature T_Z between 640 and 800° C. at a cooling rate dTZ of at least 0.50 K/s and then to the temperature T_B between 300 and 580° C. at a cooling rate dTB of at least 10 K/s, where dTB is greater than dTZ. It is thus possible for the cooling rate dTZ to correspond to a preliminary cooling and the cooling rate dTB to a rapid cooling. Even though quenching in two stages is not absolutely necessary for the final properties in the microstructure, it may nevertheless be economically viable for process-related reasons to provide for two-stage quenching. Firstly, the flat steel product, by virtue of the preliminary cooling, can be cooled down more gradually and homogenously. Secondly, the region or the distance with rapid cooling may frequently be limited in terms of plant technology, such that, in the case of provision of preliminary cooling to intermediate temperature, rapid cooling from intermediate temperature, the temperature T_B can be implemented much more easily. In order that the preliminary cooling does not lead to unwanted ferrite formation, the intermediate temperature T_Z and the cooling rate dTZ to the intermediate temperature must be high enough. The intermediate temperature is thus at least 640° C. in order to prevent coarse ferrite formation, and may especially be at least 700° C. in order, for example, to be able to completely suppress procutectoid ferrite formation. The cooling rate dTZ is at least 0.5 K/s in order to prevent coarse ferrite formation, especially at least 1.5 K/s in order, for example, to be able to completely suppress procutectoid ferrite formation. The cooling rate dTZ may be chosen up to a maximum 10 K/s or, if required, even more. The cooling rate dTB is at least 10 K/s in order to prevent ferrite formation. In particular, cooling rate dTB may be at least 20 K/s, in order, for example, also to be able to suppress complete bainite formation in the upper temperature range. The cooling rate dTB may be limited for economic reasons to a maximum of 200 K/s, especially to a maximum of 150 K/s.
In one configuration of the process of the invention, in step h), after attainment of temperature T_B the cold-rolled flat steel product is kept at temperature T_B for a hold time t_B of at least 15 s. The hold time t_B at temperature T_B is at least 15 s. The longer the hold time chosen, the more completely the austenite can be transformed to bainite, although stabilized residual austenite may indeed be part of the bainite. The hold time t_B chosen may especially be at least 25 s, in order, for example, to be able to stabilize larger austenite regions, and preferably at least 35 s, in order, for example, to minimize the formation of fresh martensite, which would lead to embrittlement. The hold time t_B may be limited, for example, to a maximum of 100 s, although it may quite possibly also be higher if required and depending on the design of the plant.
In order that no excess martensite can form in the microstructure, and especially to prevent an excess martensite content in the microstructure of more than 40%, it is optionally the case that the temperature should not go below a temperature T_B_min of Ms−50° C., especially of Ms−25° C. If the temperature goes slightly below the Ms temperature within the specified limits, this can make it easier for bainite nucleation to occur, for example at martensite lancets, and hence lead to general acceleration of bainite formation. It is thus indeed possible to allow bainite formation to commence for a couple of seconds, then to cool below Ms briefly, or else directly to cool below Ms without prior bainite formation, and then return back to the range of temperature T_B between at least 300 and not more than 580° C., especially between at least 340 and not more than 550° C., preferably between 380 and 510° C., in order to conclude bainite formation. Depending on the selection of the alloy elements, it may be the case at particularly low carbon contents that the Ms temperature is greater than at least 300° C., especially at least 340° C., preferably at least 380° C., such that, in this case, during the hold time t_B, the temperature should not go below at least 300° C., especially at least 340° C., preferably at least 380° C.
The cold-rolled flat steel product of the invention, in one configuration, has a microstructure composed of martensite with a proportion between 0.5 and 40%, especially between 3 and 33%, preferably between 5 and 28%, residual austenite with a proportion between 5 and 22%, and a residue of bainite and unavoidable microstructure constituents. The proportion of bainite in the microstructure is especially at least 55%, preferably at least 60%. The proportion of martensite in the microstructure may be especially at least 8%, preferably at least 10%, and is preferably limited to a maximum of 26%, more preferably to a maximum of 24%. Further preferably, the martensite may have been completely or partly annealed. The proportion of residual austenite in the microstructure may especially be at least 8%, preferably at least 10%, and may especially be limited to a maximum of 20%, preferably to a maximum of 18%. It is quite possible for finely divided austenite and/or carbides to be part of the bainite and/or annealed martensite. Unavoidable microstructure constituents that may be present may be proportions in the form of ferrite, perlite and/or cementite, apart from bainite and martensite, which may be allowed up to a maximum of 10%, especially up to a maximum of 8%, preferably up to a maximum of 6%, more preferably up to a maximum of 4%. It is likewise possible for precipitation in the form of chromium nitrides to be present in the microstructure.
In one configuration of the cold-rolled flat steel product of the invention, bainite is also present in the form of lancet bainite, and the ratio of the fractions of lancet bainite to bainite is at least 60%, especially at least 65%, preferably at least 70%. This allows a finer and/or more fracture-resistant microstructure to be achieved. Furthermore, this can make a particular contribution to an increase in strength as a result of its fineness and also permits accelerated stabilization of residual austenite by virtue of particularly short diffusion pathways of the carbon from the bainitic ferrite to the austenite.
In a first working example, a melt A consisting of (in % by weight) C=0.217%, Si=0.98%, Mn=1.56%, Cr=0.8%, N=0.019%, P=0.01%, S=0.003%, Al=0.01%, Ti=0.003%, Nb=0.001%, Mo=0.005%, balance: Fe and unavoidable impurities was created, cast to a precursor, solidified and then divided in the form of slabs. Alloy elements that are unspecified here were present not in measurable contents and only as unavoidable impurities. The slabs were allowed to cool down to ambient temperature. The slabs were reheated/through-heated to a temperature of 1250° C. in a walking beam furnace, such that the microstructure of the precursor consisted entirely of austenite and all deposits that had formed in the microstructure in the course of strand casting were able to dissolve. After reheating, the slab was sent to a rolling train in which the slab was first intermediately hot-rolled in a (preliminary) stand in a reversing manner with a final rolling temperature of 1100° C. to give an intermediate flat product, and the intermediate flat product was then finish-/hot-rolled in a seven-stand finishing/hot-rolling train, for example, to give a hot-rolled flat steel product (hot strip) at a thickness of 3 mm, where the final rolling temperature was 890° C. and the degree of hot rolling in the last roll pass was 15%. Immediately after the last roll pass, the hot-rolled flat steel product was actively cooled to a winding temperature of 560° C. with water along a cooling zone. The coil was then cooled to ambient temperature. The hot-rolled flat steel product was cold-rolled in a five-stand cold rolling train, for example, with an overall degree of cold rolling of 50% to give a cold-rolled flat steel product (cold strip) at a thickness of 1.5 mm. Ten samples in the form of blanks were divided from the cold-rolled flat steel product, samples 1 to 10, and were subjected to further studies.
In a second working example, the same conditions as in the first working example for production of a cold-rolled flat steel product were observed, but with the difference that a melt B consisting of (in % by weight) C=0.221%, Si=1.0%, Mn=1.58%, Cr=0.8%, N=0.0092%, P=0.011%, S=0.004%, Al=0.032%, Ti=0.004%, Nb=0.001%, Mo=0.003%, balance: Fe and unavoidable impurities, was cast. Two samples in the form of blanks were divided from this cold-rolled flat steel product, samples 11 and 12, and were subjected to further studies.
Samples 1 to 12 were subjected to a heat treatment on laboratory scale with the defined specifications according to step h) in order to establish the desired bainitic base microstructure in the cold-rolled flat steel product. The individual parameters in step h) are listed in Table 1. Table 2 gives the microstructure and the corresponding properties. The samples marked with an asterisk* are in accordance with the invention.
Since the residual austenite is measured by diffractometry by volume, for example by XRD, this can quite possibly also be a constituent of bainite and/or martensite, and so the addition of the microstructure constituents may in some cases add up to more than 100%. Depending on the coarseness of the residual austenite, however, it may also be regarded as a separate microstructure constituent. n.d. means not determinable.
| Number | Date | Country | Kind |
|---|---|---|---|
| 10 2021 119 047.9 | Jul 2021 | DE | national |
This application is the United States national phase of International Patent Application No. PCT/EP2022/069774 filed Jul. 14, 2022, and claims priority to German Patent Application No. 10 2021 119 047.9 filed Jul. 22, 2021, the disclosures of which are hereby incorporated by reference in their entireties.
| Filing Document | Filing Date | Country | Kind |
|---|---|---|---|
| PCT/EP2022/069774 | 7/14/2022 | WO |
